understanding microstructure and …rk242cd0160/...i would also like to thank mike mcgehee for his...

UNDERSTANDING MICROSTRUCTURE AND CHARGE TRANSPORT IN

SEMICRYSTALLINE POLYTHIOPHENES

A DISSERTATION

SUBMITTED TO THE DEPARTMENT OF

MATERIALS SCIENCE AND ENGINEERING

AND THE COMMITTEE ON GRADUATE STUDIES

OF STANFORD UNIVERSITY

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS

FOR THE DEGREE OF

DOCTOR OF PHILOSOPHY

Leslie Hendrix Jimison

March 2011

http://creativecommons.org/licenses/by-nc/3.0/us/

This dissertation is online at: http://purl.stanford.edu/rk242cd0160

© 2011 by Leslie Hendrix Jimison. All Rights Reserved.

Re-distributed by Stanford University under license with the author.

This work is licensed under a Creative Commons Attribution-Noncommercial 3.0 United States License.

ii



http://purl.stanford.edu/rk242cd0160

I certify that I have read this dissertation and that, in my opinion, it is fully adequatein scope and quality as a dissertation for the degree of Doctor of Philosophy.

Alberto Salleo, Primary Adviser


Michael McGehee


Michael Toney

Approved for the Stanford University Committee on Graduate Studies.

Patricia J. Gumport, Vice Provost Graduate Education

This signature page was generated electronically upon submission of this dissertation in electronic format. An original signed hard copy of the signature page is on file inUniversity Archives.

iii

iv

Abstract Semiconducting polymers are a promising class of organic electronic materials, with

the potential to have a large impact in the field of macroelectronics. In this thesis, we

focus on understanding the relationship between microstructure and charge transport

in semicrystalline polythiophenes. A method is presented for the measurement of

complete pole figures of polymer thin films using an area detector, allowing for the

first time quantitative characterization of crystalline texture and degree of crystallinity.

Thin film transistors are used to measure electrical characteristics, and charge

transport behavior is modeled according to the Mobility Edge (ME) model. These

characterization methods are first used to investigate the effect of substrate surface

treatment and thermal annealing on the microstructure of polythiophene thin films,

and the effect of microstructural details on charge transport. Next, we investigate the

semicrystalline microstructure in confined polythiophene films. Pole figures are used

to quantify a decrease in the degree of crystallinity of films with decreasing thickness,

accompanied by an improvement in crystalline texture. Next, we investigate the

influence of the degree of regioregularity, molecular weight and the processing solvent

on microstructure (degree of crystallinity and texture) and charge transport in high

mobility P3HT thin films. Surprisingly, when processing conditions are optimized,

even a polymer with moderate regioregularity can form a highly textured film with

high charge carrier mobility. Finally, we use films of P3HT with engineered,

anisotropic in-plane microstructure to understand the importance and mechanism of

transport across grain boundaries in these semicrystalline films. Results from this

study provide the first experimental evidence for the application of a percolation

model for charge transport in high molecular weight semicrystalline polymer

semiconductors. Understanding how characteristics of the polymer as well as details

of the processing conditions can affect the film microstructure and device performance

is important for future materials design and device fabrication.

v

Acknowledgements I am extremely grateful for my years at Stanford and the people that have made my

PhD such a positive experience. I would like to first thank my advisor, Alberto Salleo.

Only now as I look back can I fully appreciate his incredible patience with me as a

graduate student. Through the years he has continued to be a wonderful mentor. He

has taught me so much, and I am grateful for the opportunities he has given me to be a

part of the scientific community. I would also like to thank my fellow Salleo group

members, past and present. They are both brilliant and fun. I would like to especially

thank a handful of members that have had a significant impact on my PhD. Thank you,

Ludwig, for being a phenomenal post doc and a great mentor (life coach). Thanks to

Rodrigo for being a constant source of inspiration – brilliant, always helpful and a

great friend. Finally, a huge, huge thanks to Jon. His contributions to my work and

daily life are truly unforgettable. Having a best friend as a lab partner made the hours

at the synchrotron (and everything else) so much more fun. I would also like to thank

Mike McGehee for his generous use of lab space my first year, and to Alberto’s

former colleagues at PARC, in particular Michael Chabinyc, who stepped up to offer

equipment and discussion in the group’s early years. So much of the research

presented here took place at the Stanford Synchrotron Radiation Laboratory. I am

extremely grateful for the facility and all of the staff that keep it running, and

especially grateful for staff member Mike Toney. His enthusiasm for science and

teaching is contagious, and this thesis would not have been possible without him. A

large part of what has made my time here so enjoyable is my wonderful friends in the

Melleo Office: thanks for the laughs, food and fun. Thanks to all of the staff in

McCullough and the Department of Materials Science and Engineering, especially

Jungmee Kim, Fi Verplanke and Elise McKone, for your continuous support. I would

also like to thank my reading committee for helping to improve and polish this final

document, and NSF and Toshiba for funding. I would like to thank all of my friends

and family both near and far (at Stanford, in SF, on Caltrain, in Ohio, NC and

Maryland). Without you all, I would be a mess. Special thanks to my fiancé Matt and

my parents Chuck and Holly for their undying support, encouragment and love.

vi

Table of Contents Abstract........................................................................................................................ ivAcknowledgements....................................................................................................... vList of Tables ................................................................................................................ xList of Figures.............................................................................................................. xi1 Introduction............................................................................................................ 1

1.1 Overview........................................................................................................... 1

1.2 Materials for Organic Electronics ..................................................................... 2

1.3 Poly(3-hexylthiophene) Microstructure............................................................ 5

1.4 Conclusions....................................................................................................... 7

2 Characterization Methods..................................................................................... 92.1 Electrical Characterization................................................................................ 9

2.1.1 Thin Film Transistors................................................................................. 9

2.1.2 Transistor Fabrication .............................................................................. 10

2.1.3 Measuring Device Characteristics ........................................................... 11

2.1.4 Temperature Dependent Measurements................................................... 14

2.1.5 Mobility Edge Model ............................................................................... 15

2.2 Microstructural Characterization with X-ray Diffraction ............................... 17

2.2.1 X-ray Diffraction with Synchrotron Radiation ........................................ 17

2.2.2 X-ray Diffraction Basics .......................................................................... 18

2.2.3 Synchrotron X-ray Diffraction Experiments............................................ 22

2.3 Pole Figures of Thin Films using X-ray Diffraction and an Area Detector:

Quantifying Texture and Relative Degree of Crystallinity...................................... 27

2.3.1 Introduction to Pole Figures..................................................................... 28

2.3.2 Measuring Pole Figures with Area Detectors .......................................... 29

2.3.3 Corrections ............................................................................................... 34

vii

2.3.4 Examples of Complete Pole Figures........................................................ 35

2.3.5 Quantitative Texture and Crystallinity Using Pole Figures ..................... 37

2.3.6 Towards Absolute Crystallinity ............................................................... 39

2.4 Conclusions..................................................................................................... 40

2.5 Calculations..................................................................................................... 41

2.5.1 Calculation I: Transformation of Coordinate System.............................. 41

2.5.2 Calculation II: Calculating the Maximum Angle of χ for a Local Specular

Diffraction Pattern................................................................................................ 44

2.5.3 Calculation III: Derivation of the Relative Degree of Crystallinity......... 45

3 Understanding Processing Conditions, Microstructure and Charge Transport in Thin Films of PQT-12............................................................................................ 48

3.1 Introduction..................................................................................................... 48

3.2 Experimental Details....................................................................................... 52

3.3 Results............................................................................................................. 52

3.3.1 Grazing Incidence X-ray Diffraction (GIXD) ......................................... 52

3.3.2 Specular Diffraction................................................................................. 54

3.3.3 Two-dimensional Grazing Incidence X-ray Diffraction.......................... 55

3.3.4 Pole Figures.............................................................................................. 56

3.4 Discussion ....................................................................................................... 58

3.4.1 Effect of Annealing.................................................................................. 58

3.4.2 Effect of Surface Treatment..................................................................... 62

3.4.3 Understanding Crystallization Behavior Using Pole Figures .................. 65

3.5 Conclusions..................................................................................................... 67

viii

4 Microstructure of P3HT Thin Films as a Function of Thickness: Characterizing Texture and Degree of Crystallinity .............................................. 69

4.1 Introduction..................................................................................................... 69


4.3 Results............................................................................................................. 70

4.4 Discussion ....................................................................................................... 76

4.4.1 Implications for Charge Transport........................................................... 82

4.5 Conclusion ...................................................................................................... 83

5 Microstructure of P3HT Thin Films as a Function of Thickness: Characterizing Texture and Degree of Crystallinity .............................................. 85

5.1 Introduction..................................................................................................... 85


5.3 Results............................................................................................................. 89

5.3.1 Effect of Solvent ...................................................................................... 90

5.3.2 Effect of Molecular Weight ..................................................................... 92

5.3.3 Effect of Regioregularity ......................................................................... 93

5.4 Discussion ....................................................................................................... 94

5.4.1 Effect of Processing Conditions and Polymer Chemistry on Film

Microstructure...................................................................................................... 94

5.4.2 The ME Model and the Extracted Parameters. ........................................ 96

5.5 Conclusions..................................................................................................... 98

6 Microstructure of P3HT Thin Films as a Function of Thickness: Characterizing Texture and Degree of Crystallinity ............................................ 100

6.1 Introduction................................................................................................... 100

6.2 Fabrication and Characterization of Engineered Microstructures ................ 102

6.2.1 Experimental Details.............................................................................. 102

ix

6.2.2 Results.................................................................................................... 103

6.2.3 Discussion .............................................................................................. 106

6.3 Using the Engineered Microstructures to Investigate the Role of Grain

Boundaries ............................................................................................................. 110

6.3.1 Experimental Details.............................................................................. 110

6.3.2 Results.................................................................................................... 111

6.3.3 Discussion .............................................................................................. 112

6.4 Conclusions................................................................................................... 115

7 Conclusions ......................................................................................................... 116Bibliography and References .................................................................................. 120

x

List of Tables

Table 3-1: Mobility extracted from TFT measurements for PQT-12 films as-spun and

annealed on OTS/SiO2 and untreated SiO2 substrates. ............................................ 49

Table 3-2: ME fitting parameters for PQT-12 films as-spun and annealed on

OTS/SiO2 and annealed on SiO2.............................................................................. 51

Table 3-3: Positions and FWHM of diffraction peaks in the GIXD spectra of the PQT-

12 films. ................................................................................................................... 53

Table 3-4: Positions and FWHM of diffraction peaks in the specular patterns of the

PQT-12 films............................................................................................................ 55

Table 3-5: Values calculated from the pole figures: r. DoC and contribution to

crystallinity of crystallites oriented within 1° of the substrate normal. ................... 57

Table 3-6: Relative degree of crystallinity values calculated from pole figures shown

in Figure 3-8. ........................................................................................................... 66

Table 4-1. Values extracted from the specular diffraction patterns of P3HT films on

OTS-treated SiO2 as a function of thickness: (100) peak position, corresponding

lattice spacing, and the FWHM of the (100) peak. .................................................. 71

Table 4-2. Values extracted from the specular diffraction patterns of P3HT films on

SiO2 as a function of thickness, including (100) peak position, corresponding lattice

spacing, and the FWHM of the (100) peak.............................................................. 71

Table 5-1. Details of the polymer samples and the parameters extracted by modeling

transfer characteristics at temperatures ranging from 80-100 K using the ME model.

.................................................................................................................................. 89

xi

List of Figures

Figure 1-1. Molecular structures of common organic semiconductors. Shown are a)

Pentacene b) TIPS-Pentacene c) F8T2 d) RR-P3HT e) PQT f) PBTTT. The R-

groups represent alkyl chains of varying length: 8 carbons for F8T2, 6 carbons for

P3HT, typically 12 for PQT and typically 10 to 16 for PBTTT. ............................... 5

Figure 1-2. a) Packing structure of P3HT, illustrating the (h00), (0k0) and (00l)

reciprocal lattice directions. The green arrows indicate the two directions of fast

charge transport, along the polymer chain and across overlapping π-orbitals. b)

Illustration of the semicrystalline structure in P3HT. The blue regions represent

ordered π-stacked molecules, the black lines represent disordered polymer. c) Top-

down view of a typical spin cast thin film. The π-π stacking direction lies

isotropically in the plane of the substrate. d) Side view of a typical P3HT spin cast

film, the alkyl stacking direction lies out of the plane of the substrate, indicated by

the black arrow........................................................................................................... 6

Figure 1-3. Summary of key characteristics of semicrystalline polymers investigated in

this work..................................................................................................................... 8

Figure 2-1. Geometry of a bottom gate, top contact organic thin film transistor, with

OTS surface treatment. The “top view” image below shows the definition of

transistor width (W) and length (L). ........................................................................ 10

Figure 2-2. Example transistor characteristics for a P3HT device a) Output

characteristics for Vg = -30 V (black), -20 V (dark grey) and -10 V (light grey.) b)

Transfer characteristics in the linear regime (Vd = -10 V) c) Transfer characteristics

in saturation regime (Vd = -50 V), shown on both a linear and a log scale. The Vth

and Ion/Ioff ratio are marked. ..................................................................................... 13

Figure 2-3. The proposed density of states as a function of energy used for the

Mobility Edge Model. The black dashed line marks the Mobility Edge (ME). The

shaded region represents charges existing in the channel for a given VG (and EF) and

a given temperature. ................................................................................................. 16

xii

Figure 2-4. Derivation of Bragg’s law in real space based on the requirement for

constructive interference of incoming diffraction beams. The blue triangle on the

right is a larger version of the blue triangle in the diagram on the left. The planes

shown have an interplanar spacing of d. .................................................................. 19

Figure 2-5. Relationship between incoming vector ko, outgoing vector k, and

scattering vector q. ................................................................................................... 20

Figure 2-6. Ewald sphere construction for a simple, two-dimensional square lattice.

The radius of the Ewald sphere is equal to (1/λ)...................................................... 21

Figure 2-7. Illustration of beam line 2-1, indicating location of the 3 sets of slits, the

ionization chamber and the orientation of θ and β................................................... 23

Figure 2-8. a) Specular diffraction geometry (left) where the reciprocal lattice point

(RLP) is shown in red, k0 is the incoming radiation beam, k is the diffracted beam,

and q is the scattering vector, oriented perpendicular to the substrate. Specular

diffraction probes only the crystallites oriented with 0.03° of the surface normal.

Diffracted beams that are collected by the point detector are shown in black (right)

b) Rocking curve geometry, where the scattering vector is set to qB, and the sample

is rocked about the θ axis (left). Rocking curves map the orientation of the chosen

reciprocal lattice direction in the film (right)........................................................... 24

Figure 2-9. Left: Grazing incidence X-ray diffraction geometry. Both ko, the

incoming radiation beam, and k, the diffracted beam, lie nearly in the plane of the

substrate; q is the scattering vector. ϕ scans are measured by rotating around the ϕ

(vertical) axis. The RLPs (not shown) for an isotropic in-plane texture form a ring in

the plane of the substrate. Right: Grazing incidence X-ray diffraction probes repeat

distances in the plane of the substrate. ..................................................................... 25

Figure 2-10: Geometry of grazing incidence X-ray diffraction with a two-dimensional

image plate detector. ................................................................................................ 27

Figure 2-11. a) Schematic of relevant geometry definitions. α is the incidence angle,

zSRF is the surface normal of the sample, ko is the incoming radiation wavevector, k

is the outgoing (diffracted) radiation wavevector, L is the detector-sample distance,

δ is the vertical angle (with respect to the incoming beam) of the diffracted beam, γ

xiii

is the horizontal angle (with respect to the incoming beam) of the diffracted beam,

xDRF and zDRF are the planar coordinates of the area detector, and θB is the Bragg

angle. b) A sphere of possible crystallite orientations for a certain Bragg reflection.

The radius of the orientation sphere is defined by the magnitude of qB. χ is defined

between the surface normal and the plane of the substrate...................................... 30

Figure 2-12: Ewald sphere and orientation sphere construction a) in grazing incidence

diffraction, where α=0 and b) in local specular diffraction condition, where α=θB. 32

Figure 2-13. Steps involved in pole figure measurement for the (200) Bragg reflection

of a thin film of annealed PQT. a) 2D-GIXD pattern (left) and intensity as a

function of χ (right). The region of radially integrated intensity is marked on the 2D-

GIXD pattern, the arc is centered around q200, with a width of Δq = ~0.2 Å-1. (The

values shown at low angle of χ are understood to be distorted.) b) 2D-Local

Specular curve, taken with the α=θB, such that specular condition is achieved locally

at the (200) Bragg reflection (left). Inset shows zoomed in region around the (200)

Bragg reflection. Intensity as a function of χ (right), as marked in the local specular

curve. c) High resolution rocking curve data, collected with a point detector (left).

Intensity data from the three measurements overlain (right). .................................. 34

Figure 2-14. One-dimensional pole figures of PNDI thin films a) annealed above Tm

(black) and annealed below Tm (dashed). b) Zoomed in region of the boxed region

in (a), highlighting the resolution limited peak in the pole figure for the annealed

PNDI film................................................................................................................. 36

Figure 2-15. Intensity mapped onto the (100) orientation sphere, forming three-

dimensional pole figures of PNDI films a) annealed below the Tm b) annealed above

the Tm. Top figures show the orientation sphere in perspective and the bottom

figures show the orientation sphere form a top-down view..................................... 37

Figure 2-16. a) Pole figure of PNDI, annealed above Tm. b) Resolution-limited peak at

χ=0°, arising from perfectly oriented crystallites. c) Pole figure of PNDI, annealed

above Tm, on a log-log plot. The slowly varying intensity, accounting for the

remainder of the crystalline material, is highlighted in red. When calculating the

xiv

degree of crystallinity, the grey and red regions of pole figure intensity are treated

separately. ................................................................................................................ 45

Figure 3-1. Unit cell of PQT-12, showing the reciprocal lattice vectors corresponding

to the π-π stacking repeat (010), the alkyl stacking repeat (h00) and the repeat along

the chain backbone (00l). ......................................................................................... 50

Figure 3-2. AFM images of annealed PQT-12 thin films a) on bare SiO2, topography

mode b) and phase mode c) on OTS-8 treated SiO2, topography mode and d) phase

mode. Images reproduced with permission form Ref. [87]. .................................... 51

Figure 3-3. a) Grazing incidence diffraction patterns of PQT-12 on OTS/SiO2,

annealed and as-spun (black, grey, respectively) and SiO2, annealed (dashed green).

b) Specular diffraction patterns of PQT-12 on OTS/SiO2, annealed and as-spun

(black, grey, respectively) and SiO2, annealed (dashed green)................................ 53

Figure 3-4. 2D-GIXD patterns for PQT-12 a) annealed on OTS/SiO2 b) as-spun on

OTS/SiO2 c) annealed on SiO2 and d) as-spun on SiO2. The plots on the right are

vertical slices of the intensity at qx,y=0. ................................................................... 56

Figure 3-5. Pole figures of PQT-12 annealed on OTS/SiO2 (black), as-spun on

OTS/SiO2 (grey) and annealed on SiO2 (green). a) Intensity is on a log scale and χ is

on a linear scale. b) Same data as in (a), but on a log-log scale............................... 57

Figure 3-6. Williamson-Hall plot (FWHM of (h00) peaks, versus h2) for PQT thin

films, annealed on bare SiO2 (circles) and annealed on OTS/SiO2 (squares).......... 60

Figure 3-7 Microstructure of PQT-12 a) as-spun on OTS/SiO2 b) as-spun on SiO2 c)

annealed on OTS/SiO2 and d) annealed on SiO2. Crystallite nucleation occurs on

initial deposition directly on or near the polymer/substrate interface (a,b). Growth

occurs on annealing (c,d). ........................................................................................ 64

Figure 3-8 Pole figures of PQT-12 after quenching, after a 1 minute anneal, and after

a full anneal. ............................................................................................................. 65

Figure 3-9. Microstructure of PQT thin films a) Directly after quenching from the

isotropic melt. b) After a short anneal. Arrows indicate crystallite growth. c) After a

full anneal................................................................................................................. 67

xv

Figure 4-1. P3HT films of varying thickness on a) OTS-treated SiO2 and b) on

untreated SiO2. Darker colors represent thicker films. Film thickness ranges from

approximately 6-100 nm. For thinner films, scans shown are slightly off-specular, to

avoid film thickness interference effects. ................................................................ 71

Figure 4-2. Estimated values of domain size in the direction perpendicular to the

plane of the substrate, extracted by fitting FWHM values of the (h00) peaks and

plotting against h2..................................................................................................... 72

Figure 4-3. 2D-GIXD images with corresponding (100) local specular curves

collected with a two-dimensional image plate detector for P3HT films of different

thicknesses, spun on OTS-treated silicon oxide (a-f) and untreated silicon oxide (g-

l). Film (a) was not used for quantitative calculations. ............................................ 73

Figure 4-4. Pole figures of P3HT films of different thicknesses on a) OTS-treated

SiO2 and b) untreated SiO2. Pole figures are normalized with respect to thickness.

The pole figure with the maximum intensity at χ=0 for each dataset is shown in

black. ........................................................................................................................ 74

Figure 4-5. a) Relative degree of crystallinity of P3HT films as a function of thickness.

Values were extracted from the integrated intensity of the pole figures. b) Integrated

intensity from bulk of the pole figure, attributed to crystallinity throughout the bulk

of the film, normalized with respect to film thickness. c) Integrated intensity from

resolution-limited peak, attributed to crystallinity at the interface, normalized with

respect to thickness of the interface layer (~ 9 nm). Intensities shown in (b) and (c)

are therefore comparable.......................................................................................... 75

Figure 4-6. a) Pole figures for a thick film on SiO2 (green) and a thick film on OTS/

SiO2 (blue). Films are approximately 100 nm thick and normalized for thickness.

The main difference in the pole figures is the absence of the resolution-limited peak

for the film on SiO2. b) Unnormalized pole figures for a thick film (ca. 100 nm) and

a thin film (ca. 7 nm) on oxide................................................................................. 78

Figure 4-7. Sketch of film microstructure for a thick film (top) and a thin film

(bottom.) Perfectly oriented crystallites nucleated off of the substrate interface are

xvi

drawn with black lines. There are more perfectly oriented crystallites when film

thickness is approximately equal to the crystallite coherence length. ..................... 81

Figure 4-8. One-dimensional complete pole figures of PBTTT thin films of varying

thicknesses: 85 nm (dark red), 50 nm (medium dark red), 23 nm (red), 14 nm

(dashed pink). The three thinnest films have the same crystallite orientation

distribution, with intensity varying only close to χ=0°. ........................................... 82

Figure 5-1. Pole figure data analysis. The solid circles represent the relative degree of

crystallinity of the films in this study. (The relative degree of crystallinity for P3HT,

84% RR, DCB was associated with considerable error due to low film thickness and

low intensity, and therefore this datum is not shown.) The open circles represent

integrated intensity of the resolution-limited peak, on an arbitrary scale. All data are

normalized with respect to thickness. ...................................................................... 90

Figure 5-2. Pole figures of P3HT films illustrating the effect of solvent on crystalline

texture. a) 84% RR, 130 kDa P3HT spun from DCB (dark green) and TCB (light

green). b) 97% RR, 158 kDa, spun from DCB (dark blue) and TCB (light blue). c)

97% RR, 64 kDa, spun from DCB (black) and TCB (grey). ................................... 92

Figure 5-3. Pole figures of P3HT illustrating the effect of molecular weight. a) P3HT,

97% RR, 158 kDa (blue) and 65 kDa (black) spun from DCB. b) P3HT, 97% RR

158 kDa (light blue) and 65 kDa (grey) spun from TCB......................................... 93

Figure 5-4. Pole figures of P3HT films illustrating the effect of RR. a) P3HT, high

molecular weight spun from TCB, 97% RR (light blue) and 84% RR (light green.) b)

P3HT, high molecular weight spun from DCB, 97% RR (dark blue) and 84% RR

(dark green.) ............................................................................................................. 94

Figure 6-1. Illustration of the steps involved in the directional crystallization of P3HT

in the presence of 1,3,5 trichlorobenzene. Starting from left: A P3HT TCB solution,

nucleation of TCB needles at cool end of the substrate, continued growth of TCB

needles, solidification of P3HT on TCB needles, and an anisotropic P3HT film

following sublimation of TCB. .............................................................................. 103

xvii

Figure 6-2. a) Optical microscopy image under crossed polarizers of a directionally

crystallized P3HT film on glass. Axis of polarization is aligned as indicated by the

arrows..................................................................................................................... 104

Figure 6-3. Tapping mode AFM images, topography mode (left) and phase mode

(right.) Box 1 highlights an area of equiaxed features, and Box 2 highlights an area

of elongated features. ............................................................................................. 104

Figure 6-4. a) 2D-GIXD image of a directionally crystallized P3HT thin film

illustrating the unique texture. b) Illustration of edge-on and face-on crystallites

present in a directionally crystallized film. c) High resolution grazing incidence X-

ray diffraction patterns, with the scattering vector oriented parallel (thin black) and

perpendicular (thick black) to the fibers. d) ϕ-scan of the (010) peak. ................. 106

Figure 6-5. Illustration of relevant dimensions and stacking structure along the fast

growth direction of TCB. ....................................................................................... 107

Figure 6-6. a) Microstructure of a directionally crystallized P3HT film. On the left,

the yellow arrow indicates the direction of the fiber axis and the polymer chain axis.

Edge-on crystallites are shown in blue and face-on crystallites are shown in grey.

Different colors do not indicate different materials. On the right is a sketch of the

film showing anisotropic grain structure. The black arrow indicates the direction of

the long fiber axis and the polymer chain axis. b) Microstructure of a low-angle in-

line grain boundary, with bridging polymer chains. c) Microstructure of a fiber-to-

fiber grain boundary, with no bridging polymer chains......................................... 109

Figure 6-7. a) A parallel and a perpendicular TFT with directionally crystallized

P3HT as the active layer. ....................................................................................... 112

1

1 Introduction

1.1 Overview

The emergence of the field of organic electronics can be considered one of the most

exciting scientific developments in the past two decades. The field is based on the use of

semiconducting carbon-based polymers and small molecules in the active layers of

electronic and optoelectronic devices. Origins of the modern state-of-the-art can be traced

to the 1977 discovery that the conductivity of the doped polyacetylene can be tuned over

11 orders of magnitude.[1] This discovery introduced a new class of conjugated

conducting polymers and was awarded a Nobel Prize in 2000.[2]

Organic electronics are not being developed to replace traditional electronics, but

will instead serve as a complement and supplement to existing technology. While the soft

lattice of organic materials introduces significant disorder and places a limitation on their

charge carrier mobility, the primary parameter to measure electronic performance[3],

organic semiconductors have a number of unique benefits. Many of these materials can

be dissolved in common solvents to create semiconducting inks, eliminating the

requirement for vacuum deposition and high temperatures, making organic

semiconductors compatible with flexible substrates and roll-to-roll processing. The

scalable processing is suitable for fabrication of large-area electronics, or

macroelectronics, such as electronic paper and large-area photovoltaics. Moreover,

organic materials offer unique benefits, such as the ability to tune the electrical and

optical properties through simple modifications in the chemical structure. Their electrical

properties are sensitive to the environment, which can be exploited to fabricate

sensors.[4-6] Because they are carbon based many of these materials are biocompatible,

allowing for unique applications in the fields of biosensing and bioelectronics.[7]

2

Devices based on organic semiconductors are beginning to come to market. The

organic electronic technology closest to market is the organic light emitting diode

(OLED), which can be found in camera and mobile phone displays. Benefits of OLEDS

over traditional liquid crystal displays include a wide viewing angle and no need for a

backlight, making the overall display much thinner and more lightweight. A review of

recent advances in OLED technology can be found in Ref [8]. Organic photovoltaic

(OPV) technology is also currently being commercialized. Konarka, has patented Power

Plastic®, an organic photovoltaic material being developed for use as portable power and

power for remote locations. Konarka and the city of San Francisco have teamed up to

provide the city with bus stops that incorporate their organic photovoltaic material, Power

Plastic®, in their rooftops. The power generated is used for LED lighting, LED signage

and Wi-Fi routers.[9] Organic thin film transistors (OTFTs) are not as prevalent on the

market. Sony currently has a prototype of an OLED display powered by OTFTs. With all

organic components, the display is flexible and fully functioning at only 80 µm thick.[10]

These recently introduced applications are a testament to the numerous scientific

and technological advances in organic electronics since their 1977 introduction. However,

the field is far from its full potential. Organic electronic devices must be more reliable

with higher device-to-device reproducibility and operate at higher currents before

widespread incorporation as the active layer in macroelectronics. Much of these

improvements can be realized through intelligent materials design. The goal of the work

presented here is to characterize the structure of semicrystalline semiconducting polymers

and understand how the structure affects the electronic properties. Understanding the

effect of processing conditions and polymer chemistry on microstructure and charge

transport will help in the effort to establish rigorous design rules for new organic

materials with improved properties. Following a brief introduction to materials for

organic electronics, the rest of this chapter will focus on semicrystalline polythiophenes,

providing an overview of what we know and the questions that remain.

1.2 Materials for Organic Electronics

Typical organic materials are insulating, comprising single C-C bonds with sp3

hybridization. Conjugated materials, on the other hand, have alternating single and

3

double bonds (with sp2 hybridization), allowing for the delocalization of π-electrons into

a band. The system is unstable with respect to bond alteration (Peierls distortion),

creating a gap in the energy levels.[1] The resulting band gaps are typically between 2

and 3 eV (in the realm of visible light), giving these materials interesting electro-optical

properties. In theory, the π-electrons can delocalize along the entire conjugated molecule

to arbitrary length scales. In real materials however, chemical and structural defects limit

the length of the conjugated region to only a few nm. Molecular overlap between

neighboring conjugated molecules is also highly dependent on relative spacing and

orientation between molecules, both of which are controlled by molecular packing and

defects throughout the film.[11] In this way, structural order has a large affect on charge

transport, a reoccurring theme throughout this dissertation.

There are two categories of organic electronic materials: small molecules and

polymers. One of the most common small molecule semiconductors is pentacene,

shown in Figure 1-1(a). Polycrystalline pentacene films commonly have mobilities

over 3.5 cm2V-1s-1.[12] Mobilities of high quality single crystals are expected be as

high as 35 cm2V-1s-1.[13] While electrical performance is high, pentacene and other

molecular semiconductors are very insoluble, even in warm aromatic solvents.

Molecular semiconductors are therefore often vacuum deposited. Pentacene also suffers

from instability: in the presence of light and oxygen, it easily undergoes photooxidation,

resulting in an irreversible chemical change and deteriorating performance.[14] In an

effort to increase solubility, side groups can be added. The side groups influence the

molecular packing, creating more free volume in the crystal structure and increasing the

material’s entropy on dissolution. As a result, side-group substituted small molecules

can be solution processed. Anthony and coworkers functionalized pentacene with bulky

alkynyl silyl groups, resulting in a high mobility, solution-processable pentacene

derivative, TIPS-Pentacene, shown in Figure 1-1(b).[15] The placement of the

sidegroups also protects the molecule from photooxidation, increasing stability.

Solution processable small molecules are an active and promising area of research.[16-

19]

The second category of organic electronic materials is semiconducting polymers.

Many high performance semiconducting polymers are based on a fluorene or thiophene

4

repeating unit.[20] Side chains (usually alkyl groups) along the chain backbone

increase solubility. As a result, semiconducting polymers are extremely soluble in a

number of common solvents, allowing for easier processing than their small molecule

counterparts. Shown in Figure 1-1(c) is a fluorene bithiophene copolymer poly(9,9-

dioctylfluorene-co-bithiophene) (F8T2). F8T2 is not semicrystalline, but the chain

backbone can be oriented on a macroscopic scale when heated to a mesophase

temperature in the presence of an aligning influence (such as an anisotropically textured

substrate), reaching a mobility of 0.02 cm2V-1s-1.[21, 22]

The highest performing semiconducting polymers are the semicrystalline

polythiophene derivatives. The thiophene unit in the chain backbone is an electron-rich

planar aromatic heterocycle. The presence of flexible alkyl side chains renders the

polythiophenes discussed here extremely soluble. Regiorandom (RRa) poly(3-

hexylthiophene) was introduced first, with a field effect mobility of 1 × 10-5 cm2V-1s-1

and an amorphous microstructure.[23] By controlling the coupling of the monomers,

regioregular (RR) P3HT was synthesized, with much higher mobilities of 0.01 cm2V-1s-

1 to 0.1 cm2V-1s-1, Figure 1-1(d).[24, 25] The drastic increase in mobility was attributed

to the improved molecular packing and semicrystalline microstructure. More recently,

additional polythiophenes have been synthesized with higher mobilities and improved

stability, namely poly[5,5’-bis(3-dodecyl-2-thienyl)-2,2’-bithiophene] (PQT)[26] and

poly(2,5-bis-alkylthiophene-2-yl)theino[3,2-bb]thiophene (PBTTT)[27]. Structures of

PQT and PBTTT are shown in Figure 1-1(e,f). PQT has the same thiophene backbone

as P3HT, but with a lower density of side chains. The low density of side chains allows

the molecular stacks to interdigitate. After heating through a mesophase, the resulting

film is more crystalline than P3HT, with a mobility of 0.1 cm2V-1-s-1.[28, 29] PBTTT

maintains the low density of side chains, but incorporates a fused thiophene ring in the

backbone. This serves two purposes: to lower the HOMO level, thus creating a more

stable material, and to increase the rigidity of the backbone, resulting in improved

molecular packing. When heated above its mesophase, PBTTT forms a highly

crystalline film with mobilities up to 0.6 cm2V-1s-1.[27, 30]

5

Figure 1-1. Molecular structures of common organic semiconductors. Shown are a) Pentacene b) TIPS-Pentacene c) F8T2 d) RR-P3HT e) PQT f) PBTTT. The R-groups represent alkyl chains of varying length: 8 carbons for F8T2, 6 carbons for P3HT, typically 12 for PQT and typically 10 to 16 for PBTTT.

1.3 Poly(3-hexylthiophene) Microstructure

P3HT no longer holds the record for the highest-mobility semiconducting polymer,

but remains technologically relevant due to its widespread success in bulk heterojunction

solar cells[31]. Furthermore, the extensive literature on P3HT makes it an ideal model

material for in-depth investigations of fundamental microstructure-property relationships

in semicrystalline semiconductors. For this reason, P3HT is the focus of much of the

work presented here.

Crystallites of P3HT pack in a lamellar fashion with adjacent molecules stacking

face-to-face, allowing for the overlap of their π-orbitals, Figure 1-2(a). This results in

two directions of charge delocalization, or two directions of fast charge transport:

along the chains and along the overlapping π-orbitals. The alkyl layer between

molecular stacks is insulating, resulting in poor charge transport in this direction. As

mentioned above, P3HT forms semicrystalline films, consisting of crystallites

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approximately 10-20 nm in diameter, separated by amorphous grain boundary regions,

Figure 1-2(b). In thin films of P3HT, the alkyl stacking direction tends to lie out of the

plane of the substrate while the π-stacking direction lies isotropically in the plane of

the substrate. This is classically referred to as fiber-texture and illustrated in Figure 1-

2(c,d). We refer to crystallites with their alkyl stacking out-of-plane as “edge-on.”

When crystallites are oriented with their alkyl stacking direction in the plane of the

substrate, we call them “face-on.” In Figure 1-2(a), the alkyl stacking, chain backbone

and π-stacking directions are labeled, along with their Miller indices.[3] Miller indices

will be used to identify reciprocal lattice directions throughout this work.

Figure 1-2. a) Packing structure of P3HT, illustrating the (h00), (0k0) and (00l) reciprocal lattice directions. The green arrows indicate the two directions of fast charge transport, along the polymer chain and across overlapping π-orbitals. b) Illustration of the semicrystalline structure in P3HT. The blue regions represent ordered π-stacked molecules, the black lines represent disordered polymer. c) Top-down view of a typical spin cast thin film. The π-π stacking direction lies isotropically in the plane of the substrate. d) Side view of a typical P3HT spin cast film, the alkyl stacking direction lies out of the plane of the substrate, indicated by the black arrow.

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1.4 Conclusions

It is well established that details of the polymer chemistry and processing conditions

affect film microstructure, which in turn affects charge transport. However, the

semicrystalline nature of these films complicates the understanding of this relationship.

Within the crystallites, charges are expected to traverse relatively easily, but may be

affected by intracrystalline disorder. It is expected, however, that the disordered

boundaries between crystalline regions (the grain boundaries) serve as significantly larger

barriers to transport. Despite the advances that the field has made, there are some

questions that remain unanswered:

1) How crystalline are these materials? The degree of crystallinity in polymer

films is not well known. Without this knowledge, it has been difficult to address

the importance of the degree of crystallinity in regards to charge transport. Are

other microstructural details, such as crystalline orientation, more important?

Moreover, how can we control the degree of crystallinity and crystalline

orientation?

2) What is the role of grain boundaries? The structural details of the grain

boundaries and the mechanism for charge transport across grain boundaries are

not well understood. Are grain boundaries always detrimental to charge transport,

or could their influence be controlled?

In order to understand the relationship between microstructure and charge

transport, it is important to address the questions above. This thesis presents work that

was designed to assess in a quantitative fashion the effect of processing conditions on

crystalline quality, crystallite orientation and degree of crystallinity. We then correlate

the details of the microstructure with charge transport properties. We also use films with

an engineered in-plane orientation to investigate transport through grain boundaries.

Figure 1-3 summarizes the details of film structure under investigation in this work.

8

Figure 1-3. Summary of key characteristics of semicrystalline polymers investigated in this work.

Chapter 2 provides a detailed discussion of the characterization methods used in

this thesis. Thin film transistors and the Mobility Edge model are used to investigate and

analyse the electrical properties of the semiconducting polymers. X-ray diffraction is

used to study the film microstructure. In this chapter, we introduce a novel method for

measuring pole figures of thin polymer films, which allows for quantitative

characterization of crystalline texture and degree of crystallinity. Chapter 3 discusses a

study that investigates the effect of substrate surface treatment and thermal annealing on

the microstructure of PQT, and the effect of microstructural details on charge transport.

Chapter 4 investigates semicrystalline microstructure in confined polymer films. In this

chapter, pole figures are used to quantify the change in texture and degree of crystallinity

of P3HT films as a function of thickness and the resulting implications on charge

transport are discussed. Chapter 5 discusses structural and electronic characterization in

high mobility P3HT films. Here, the variables under investigation are the degree of

regioregularity, the molecular weight and the processing solvent. In Chapter 6, films with

engineered, anisotropic in-plane microstructures are used to understand the mechanism of

transport across grain boundaries in high molecular weight semicrystalline

semiconducting polymer films. Chapter 7 provides a summary of conclusions made

throughout this work.

Degree of

Crystallinity

9

2 Characterization Methods

This chapter discusses methods that are important for the characterization of

semicrystalline semiconducting polymers. Section 2.1 discusses electrical

characterization. The thin film transistor is introduced, including a description of

transistor fabrication and transistor measurements. This is followed by a discussion of the

Mobility Edge (ME) model, a charge transport model used to analyze charge transport

behavior of semicrystalline polythiophenes. Section 2.2 discusses microstructural

characterization using X-ray diffraction, including details of the beam lines and

diffraction geometry available at Stanford Synchrotron Radiation Lightsource. Section

2.3 introduces a novel method for measuring pole figures of weakly diffracting thin films,

developed for the quantitative characterization of crystalline texture and degree of

crystallinity. Section 2.4 concludes the chapter, and Section 2.5 includes calculations for

reference. The techniques introduced in this chapter are implemented in work presented

throughout this thesis.

2.1 Electrical Characterization

2.1.1 Thin Film Transistors

We use thin film transistors (TFTs) as a tool to study the electrical characteristics

of semiconducting polymers. The mobility, a measure of how easily charges traverse in a

material under a given electric field, is the primary figure of merit for comparing charge

transport between different films. Transistors are three terminal devices consisting of

source, drain and gate electrodes. The gate electrode modulates the charge density in the

transistor channel, controlling the resistance between the source and the drain electrodes.

Common thin film transistor (TFT) geometries include 1) top gate, bottom contact, 2)

bottom gate, bottom contact and 3) bottom gate, top contact, (contact refers to the source

and drain electrodes.) In this work, we use a bottom gate, top contact geometry, shown in

10

Figure 2-1. Bottom gate, top contact devices benefit from extremely easy processing

without the use of lithography. This geometry also allows for unperturbed film formation

and crystallization at the buried dielectric interface, where charge transport takes place.

When thermally grown SiO2 (on a Si wafer) is used as the dielectric, this surface is

extremely flat.

Figure 2-1. Geometry of a bottom gate, top contact organic thin film transistor, with OTS surface treatment. The “top view” image below shows the definition of transistor width (W) and length (L).

2.1.2 Transistor Fabrication

Device substrates (~1 cm by 1 cm) were cut from a highly doped silicon wafer

(serving as the common gate electrode) coated with 200 nm of SiO2 (serving as the gate

dielectric). Substrates were cleaned by ultrasonication in acetone (Sigma Aldrich) and

isoproponal (Sigma Aldrich), rinsed with deionized water and dried with a filtered air gun.

A 10 minute vacuum oven step removed any excess water. Prior to film deposition, the

cleaned surfaces were subjected to a UV Ozone treatment (Novascan, PSD-UV, UV

Surface Contamination System). At this point, the substrates were either ready for

polymer deposition or additional surface treatment.

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For the octadecyltrichlorosilane (OTS-18) surface treatment, a solution deposition

technique was used[32]. First, 40 µL of OTS-18 (Gelest, Inc) was added to 20 mL of

hexadecane (Sigma Aldrich) in a clean glass petri dish. The cleaned substrates were then

submerged for 10-20 minutes. On removal, the substrates were thoroughly rinsed with

heptane (Sigma Aldrich) to avoid OTS polymerization (evidenced by a cloudy white film

on the surface of the substrate), rinsed briefly in a stream of acetone and isopropanol and

dried with a filtered air gun. The resulting surface was very hydrophobic, with a water

contact angle > 90°.

To make the polymer solution, dried polymer was weighed in air, but transferred

into a nitrogen glovebox before dissolving in solvent. Typical solution concentration was

0.5-1.0 wt%. Polymer deposition was performed in the glovebox. For the majority of the

samples presented here, the polymer solution was spin cast onto the substrates using a

Laurell spincoater. Spin acceleration was set between 100 and 500 rpm. Final spin speed

was typically between 1000 and 1250 rpm. The total spin time was approximately 2

minutes. In some cases, to aid film coverage the solution was allowed to sit on the

substrate for a designated “gel time” before starting the spin coating process. To dropcast

films, the substrate was placed in a glass petri dish. Solution was added carefully to the

substrate and the cover of the petri dish was replaced, creating a saturated solvent

atmosphere, encouraging slow drying of the film.

Gold source and drain contacts were deposited via thermal evaporation through a

shadow mask, with a thickness ca. 80 nm. Typical channel lengths were between 50 and

500 µm. Channel width was defined by scribing the polymer and gold prior to

measurement using the probe tip. In addition to defining the W/L, this served to isolate

the device from the gate contact, reducing parasitic gate leakage current. Values for W/L

were kept between 3 and 30.

2.1.3 Measuring Device Characteristics

The organic thin film transistors presented here are p-type, meaning that

positively charged holes are the majority charge carrier. Transistors operate in an

accumulation mode. When the gate electrode is negatively biased, a large concentration

12

of carriers (dependent on the gate voltage) is accumulated in the transistor channel. When

there are enough carriers in the channel, the transistor turns on, and a current can be

measured when a voltage is applied between the source and the drain electrodes. Charge

carrier mobility is the primary figure of merit when comparing electrical performance.

During analysis, we assume application of the gradual channel approximation.[33]

At low gate voltages, the transistor operates in linear regime. The drain-source current

measured is dependent on the device geometry (W/L), the capacitance of the insulating

dielectric Ci (17.3 nF/cm 2 for 200 nm SiO2), the gate voltage Vg, the threshold voltage Vth,

the drain voltage Vd and the mobility µ of the mobility µ of the charge carriers, as shown

in Equation (2-1):

(Eq. 2-1)

When (Vd > Vg – Vth), the transistor saturates. Current in the saturation regime is no

longer dependent on the drain voltage, as shown in Equation (2-2):

(Eq. 2-2)

Typical output characteristics (Ids vs. Vd) and transfer characteristics (Ids vs. Vg) are shown

in Figure 2-2.

!

Ids = µCi

W

2L(Vg "Vt )

2

!

Ids = µW

LCi(Vg "Vth )Vd

13

Figure 2-2. Example transistor characteristics for a P3HT device a) Output characteristics for Vg = -30 V (black), -20 V (dark grey) and -10 V (light grey.) b) Transfer characteristics in the linear regime (Vd = -10 V) c) Transfer characteristics in saturation regime (Vd = -50 V), shown on both a linear and a log scale. The Vth and Ion/Ioff ratio are marked.

To characterize the devices, transfer characteristics were collected in vacuum and

dark (MMR Technologies probe station). Mobilities in the linear regime were determined

according to the relationship given in Equation (2-3)

(Eq. 2-3)

where mlin is the slope of the transfer curve at high Vg, with a Vd of -10 V, marked by the

dotted lines in Figure 2-2(b). We note that this is not a rigorous extraction of mobility in

that it ignores the dependence of mobility on Vg (due to the changing carrier

concentration with gate voltage[28, 34-36]), but it is a good estimation for devices used

in this study.

Mobilities in the saturation regime were determined according to Equation (2-4).

€

µ =2L ×msat

2

W ×Ci

(Eq. 2-4)

where msat represents the slope of Ids(1/2) vs. Vg at high Vg, marked by the dotted lines in

Figure 2-2(c). Identical mobilities extracted in the linear and saturation regimes assure

that the devices are well behaved. The shapes of the output curves Figure 2-2(a) were

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assessed to assure that devices were saturating and effect of contact resistance was

minimal.[37, 38]

The threshold voltage (Vth) can be estimated from transfer characteristics in the

linear regime. The x-intercept (B) of the extrapolated linear fit of the curve at high Vg,,

derived from Equation (2-1) .

(Eq. 2-5)

Alternatively, Vth is equal to the intercept of the linear fit of a transfer curve in

saturation when plotted as Isd1/2 vs. Vd. The Ion/Ioff ratio is best estimated from a transfer

curve in the saturation regime on a log scale, see grey line in Figure 2-2(c). The value for

Ion is defined as the value where the curve begins to plateau and the value for Ioff is

defined as the lowest current measured, where the subthreshold region (voltage below Vth)

meets the leakage current.

One common problem in organic semiconductors is degradation during device

operation due to bias stress. Bias stress is observed as a shift in the Vth in consecutive

transfer curves. This is equivalent to a decrease in the measured Ids for a given (Vg, Vds)

while the mobility remains unchanged. Bias stress in polymer semiconductors has been

attributed to trapped charge existing in the semiconductor or in the dielectric.[39, 40] The

effects of bias stress during device measurements can be avoided by using a pulse drain

and gate, rather than a continuous electrical bias (and continuous current.) Electrical

characterization was performed with either a Keithley 2612 Sourcemeter and Labtracer

software (for direct current measurements, when bias stress was neglible), or two

Keithley 2400 Sourcemeters and Labview software (for pulsed gate measurements to

minimize bias stress.)

2.1.4 Temperature Dependent Measurements

In order to measure the temperature dependence of the charge carrier mobility, the

samples were mounted on a Joule-Thompson refrigerator finger in a vacuum probe

station (MMR Technologies) using a thin layer of thermally conductive silicone paste for

sample adhesion and efficient heat exchange. The samples were cooled to 80 K and

!

Vth

= B "Vd

2

15

heated to room temperature in steps of 20 K. At every temperature, a transfer curve was

recorded by sweeping the gate from 0 to -60 V with a source-drain voltage of -10 V. Care

was taken to reposition the probes on the contacts when needed. Care was taken to assure

the threshold voltage shift during measurements was no significant. If the observed Vth

shift significant for consecutive transfer curves taken at different temperatures, pulsed

measurements were used.

2.1.5 Mobility Edge Model

In order to extract more information about the electrical characteristics of the

semiconducting polymers, we can fit transport behavior to a charge transport model.

There are a number of different charge transport models available for the analysis of

disordered semiconductors. Here, we use the Mobility Edge (ME) model. The ME model

describes electron (hole) mobility as transport in a band state with occasional trapping

and de-trapping events in and from an exponential tail of trap states, situated below

(above) a fixed energy EME, denoted as the mobility edge (ME) [28, 41, 42]. While the

assumption of band-like transport in crystalline organic semiconductors is reasonable, in

disordered polymers this concept may not apply. Nevertheless, structural evidence

obtained in previous studies (and confirmed in this work) proves that regioregular

polythiophenes P3HT and PQT organize in crystalline lamellae where charge is partially

delocalized[43]. Based on these observations, we use the ME model to describe our

temperature dependent transistor data and extract trap-free mobility of mobile charges.

We do not imply that the carriers above the mobility edge in P3HT are completely free,

but rather assume that they are much more mobile that the carriers below the mobility

edge. The ME model has been successfully applied to small molecular organics[44]

polycrystalline oligothiophenes[45] and semicrystalline polythiophenes[28, 41, 46, 47].

The ME model proposes a density of states (DOS) composed by two well

distinguished parts that exist above and below the mobility edge (defined as EME = 0). In

this simplified DOS, carriers have a mobility µ0 and zero, respectively, which in

combination can model the band-like transport and occasional trapping and detrapping

expected for these materials. Semicrystalline P3HT and other similar semicrystalline

polythiophenes form nanocrystalline lamellae structures parallel to the semiconductor-

16

insulator interface. The electronic interaction between adjacent lamella planes is small

enough such that transport can be assumed to happen in a two-dimensional density of

states within the ordered regions[41]. Hence, a constant density of states with a hole

mobility µ0 is assumed below the ME. Due at least in part to structural disorder (between

or within crystallites), states are found in the band gap. For simplicity the DOS in the gap

is often modeled as an exponential tail characterized by a total number of trap states Ntot

and a characteristic energetic width Eb. The complete DOS is defined as

for (Eq. 2-6a)

for (Eq. 2-6b)

where Dmob is the density of mobile states, Ntot is the trap density and Eb is the

exponential width of the energetic distribution. This is illustrated below in Figure 2-3.

Figure 2-3. The proposed density of states as a function of energy used for the Mobility Edge Model. The black dashed line marks the Mobility Edge (ME). The shaded region represents charges existing in the channel for a given VG (and EF) and a given temperature.

In the ME model, the measured effective mobility is the average mobility of all the

carriers, both trapped and free. With increasing |Vg|, the fraction of holes in mobile states

increases and so does the effective mobility.[35] The trend of increasing mobility with

increasing charge density is commonly observed in all semicrystalline polymers. When

!

"(E) = Dmob

!

E " 0

!

E > 0

!

"(E) =N

t

Eb

e#E /Eb

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17

temperature is decreased, fewer charges are thermally excited above the ME and the

effective mobility at constant Vg decreases. The ME model can be a means to link carrier

transport to details of the microstructure. Structural defects, at the grain boundaries or

within the grains, can act as trap states. In the case of polycrystalline silicon thin films,

broader trap distributions have been associated with the smaller grained material. Details

pertaining to fitting electrical characteristics to the ME model are beyond the scope of

this work and can be found elsewhere.[28, 48] Importantly, the parameters µo, Eb, Nt, and

Von can then be estimated by fitting the model to the current-voltage measurements of the

transistors at different temperatures.

2.2 Microstructural Characterization with X-ray Diffraction

2.2.1 X-ray Diffraction with Synchrotron Radiation

Thin films of semiconducting polymers can be difficult to characterize with

traditional characterization methods. The films are too unstable for SEM and TEM,

except under certain low-energy conditions. Atomic force microscopy has proved

compatible with the soft materials, but does not give structural information and typically

is limited to measuring topology and other materials properties of the top polymer-air

interface. In order to probe the bottom surface, the film must be delaminated, which may

alter the local morphology. In contrast, X-ray diffraction probes the entire sample, and

provides information about the bulk of the film. Diffraction experiments in this thesis

were performed at the Stanford Synchrotron Radiation Lightsource (SSRL). The facility

operates at an injection energy of 3 GeV, with a current between 100-500 mA.

Synchrotron radiation at SSRL has a low angular emittance of 10 nm-rad in the vertical

direction, allowing for measurements with a high angular resolution. With the use of

wigglers and undulators, the diffracting beam is highly collimated. The high intensity and

high collimation results in a high brightness, with brightness defined as the flux per unit

area of the radiation source per unit solid angle of the radiation cone. Conventional X-

rays, in comparison, have a much smaller brightness and larger angular divergence.

Moreover, synchrotron radiation comprises white light: monochromators can be used to

18

select a specific wavelength, allowing measurement of both small and large repeat

distances in real space. The radiation is highly polarized in the plane of the synchrotron

(~95%). Synchrotron radiation allows for the adequate collection of diffracted intensity

from weakly diffracting organic thin films, with a high signal to noise ratio.

2.2.2 X-ray Diffraction Basics

X-ray diffraction is a powerful, nondestructive technique that is used to

characterize crystalline structure in a material. Here, the very basics of X-ray diffraction

are discussed, but the reader is referred to other texts for a more in-depth description[49,

50].

The physics behind X-ray diffraction is based on the coherent scattering of

radiation beams reflected from well-ordered lattice planes. A derivation of Bragg’s law,

which governs constructive interference of outgoing waves by accounting for phase

change, is illustrated in Figure 2-4. The two waves shown are reflected off adjacent

atomic planes. Constructive interference will occur when the outgoing waves have the

same phase. For incoming radiation with wavelength λ, incident angle θ, and interplanar

spacing d, it follows that

(Eq. 2-7)

Where Equation (2-7) is Bragg’s Law, and θB is the incident angle for which Bragg

condition is satisfied, referred to as the Bragg angle.

19

Figure 2-4. Derivation of Bragg’s law in real space based on the requirement for constructive interference of incoming diffraction beams. The blue triangle on the right is a larger version of the blue triangle in the diagram on the left. The planes shown have an interplanar spacing of d.

A convenient way to think about diffraction is by operating in reciprocal space.[3,

49] Briefly, every crystal lattice in real space has a corresponding lattice in reciprocal

space, which can be described using the reciprocal lattice vectors defined below:

€

b1 =a2 × a3a1 ⋅ a2 × a3

€

b2 =a3 × a1a1 ⋅ a2 × a3

€

b2 =a1 × a2a1 ⋅ a2 × a3

The general reciprocal lattice vector is defined as follows:

€

Ghkl = hb1 + kb2 + lb3 (Eq. 2-8)

where h, k and l are the Miller indices used to designate planes within a crystal. The

magnitude of Ghkl has a reciprocal relationship with the hkl interplanar spacing (dhkl):

(Eq. 2-9)

!

Ghkl

=1

dhkl

20

When working in reciprocal space, it is necessary to define the scattering vector,

q, as shown in Figure 2-5. Diffraction data throughout this thesis will be presented in

terms of q, rather than incident angle.

Figure 2-5. Relationship between incoming vector ko, outgoing vector k, and scattering vector q.

Bragg condition is satisfied when the scattering vector q, intersects a reciprocal

lattice point, whose position is defined by the general reciprocal lattice vector Ghkl. In

other words, following the geometry in Figure 2-5, the scattering vector that satisfies

Bragg condition can be defined as follows:

€

qB =4πλsin(θB ) = 2π Ghkl

Therefore, an alternative representation of Bragg’s law (Equation 2-7) is given in

Equation (2-9).

€

qB =2πdhkl

(Eq. 2-9)

The diffraction condition in reciprocal space can be illustrated using Ewald sphere

construction. Figure 2-6 shows a two-dimensional illustration of Ewald’s sphere for a

simple square lattice. Ewald’s sphere is useful for comparing different diffraction

geometries and will be referred to throughout this chapter. With a given reciprocal lattice,

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21

the terminus of vector ko/(2π) is placed on a reciprocal lattice point. Ewald’s sphere is

then drawn with its center at the origin of the ko/(2π) vector. The radius of Ewald’s

sphere is defined as 1/λ (=ko/(2π)). Any reciprocal lattice point (RLP) that lies on this

sphere can be connected by reciprocal lattice vector Ghkl, and is at the terminus of k(2π)

(which also has its origin at the center of Ewald’s sphere, as shown), and thus satisfies

Bragg condition.

Figure 2-6. Ewald sphere construction for a simple, two-dimensional square lattice. The radius of the Ewald sphere is equal to (1/λ).

Ewald’s sphere correctly demonstrates possible diffraction conditions, but the

diffraction intensity recorded on a detector is more complicated. Importantly, the

diffracted amplitude is scaled by the modulus squared of the structure factor |F(q)|2,

which at Bragg condition is defined according to Equation (2-10) below.

€

Fhkl = F(qB ) = fnei2πGhkl ⋅rn

n

NB

∑ (Eq. 2-10)

where fn is the atomic scattering factor, NB is the number of atoms is a unit cell (repeating

unit in a crystal lattice), and rn is the vector defining the position of an atom within the

unit cell. The structure factor reflects information about the arrangement of the electron

!

qB = (k " ko)B = 2#Ghkl

22

distribution within the unit cell in a crystal lattice. The elimination of certain diffraction

peaks due to interference is accounted for in the structure factor. There are additional

factors that modulate the collected intensity, such as temperature, but discussion of these

terms is beyond the scope of this thesis.

Various structural details such as crystal size, crystal shape, crystalline quality

and orientation distribution influence the shape of the reciprocal lattice feature. Much of

the work presented in this thesis is dedicated to accurately characterizing features in

reciprocal space in order to better understand the microstructure of semicrystalline thin

films.

2.2.3 Synchrotron X-ray Diffraction Experiments

2.2.3.1 Specular Diffraction and Rocking Curves

Beamline Details

Specular and rocking curve diffraction experiments are performed at SSRL beam

line 2-1. This is a high-resolution beam line, equipped with a point detector. For

experiments in this dissertation, beam line 2-1 was operated at 8 keV and photon

wavelength of ~0.155 nm. The beam defining slits between the incoming beam and the

sample, v0-gap and h0-gap, are set at 0.2 and 1.5 mm, respectively. A gas ionization

monitor is placed between the beam defining slits and the sample to track incoming X-ray

flux, which is used to normalize diffracted intensity. Between the sample and the detector,

the two vertical slits, v1-gap and v2-gap, are set at 1.2 and 1.0 mm respectively, with the

corresponding horizontal slits set wide to increase beam footprint and horizontal

acceptance angle. Due to the scattering slits at this beam line, the resolution is anisotropic.

The resolution in the θ direction is slightly less than 0.005°. The acceptance angle of the

diffracted beam in the β direction (rotation perpendicular to the beam) is larger,

approximately 0.8°. To reduce beam damage and air scattering background, samples are

placed in a chamber purged with helium during measurements. The sketch in Figure 2-7

illustrates beam line setup.

23

Figure 2-7. Illustration of beam line 2-1, indicating location of the 3 sets of slits, the ionization chamber and the orientation of θ and β. Measurements Details

Specular diffraction measurements at SSRL beam line 2-1 are analogous to θ-2θ

scans performed using a traditional diffractometer. These measurements are used to probe

periodicity existing out-of-the-plane of the substrate, with scattering vector q along the

substrate normal. The diffraction geometry is illustrated in Figure 2-8(a) below. The

incoming wave vector forms an angle θ with respect to the substrate, and the detector

forms an angle 2θ with respect to the incoming scattering vector. As shown, the q-vector

being probed exists normal to the substrate. As θ is increased, q moves through reciprocal

space. The high resolution in the θ direction means that the only crystallites measured in

the specular diffraction experiments are oriented with the scattering vector within 0.03°

of the surface normal. This diffraction geometry is especially useful for high-resolution

peak position and peak shape analysis, but since it probes only a fraction of the

crystalline material -- it does not give complete information about the film microstructure.

Rocking curves are a means of measuring the angular distribution of crystal

orientations. To measure a rocking curve, 2θ is set to satisfy qB for a chosen Bragg

reflection. At this point the sample normal is rocked relative to the beam in the θ

direction, but the magnitude of the scattering vector (defined by 2θ) is kept constant. As

the film is rocked, different crystallites are brought into diffraction condition, and the

resulting plot maps the population of crystallites at different orientations, (b).

Background scans are collected for q<qB and q>qB to remove intensity due to reflection

of the substrate. Diffracted intensity is normalized for beam footprint by multiplying the

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24

collected intensity by the corresponding θ. The angular range available for a θ-axis

rocking curve is limited to 0<θ<2θB, where 2θB is typically between 5-10°.

Figure 2-8. a) Specular diffraction geometry (left) where the reciprocal lattice point (RLP) is shown in red, k0 is the incoming radiation beam, k is the diffracted beam, and q is the scattering vector, oriented perpendicular to the substrate. Specular diffraction probes only the crystallites oriented with 0.03° of the surface normal. Diffracted beams that are collected by the point detector are shown in black (right) b) Rocking curve geometry, where the scattering vector is set to qB, and the sample is rocked about the θ axis (left). Rocking curves map the orientation of the chosen reciprocal lattice direction in the film (right).


Beamline Details

High-resolution grazing incidence X-ray diffraction is measured at SSRL beam

line 7-2 with a point detector. For experiments in this dissertation, beam line 7-2 is

operated at 8 keV and photon wavelength of ~0.155 nm. Beam defining slits, v1-gap and

h1-gap, are set at 0.3 and 0.25 mm. As in beam line 2-1, a gas monitor exists between the

beam defining slits and the sample, allowing for diffracted intensity to be normalized

with respect to incoming radiation. The diffracted beam passes through Soller slits and

additional beam defining slits v2-gap and h2-gap (set between 4-6 mm and 6 mm).

Measurement Details

When the incident angle of the incoming radiation beam is set at the critical angle

an evanescent X-ray wave propagates along the surface decaying exponentially into the

film. The depth that the evanescent wave penetrates can be modulated by changing the

incidence angle. Many groups have used this property to perform depth-dependent

25

grazing incidence diffraction studies of polymer films[51, 52]. In the work presented

here, grazing incidence diffraction patterns were used to measure the in-plane structure of

the bulk film. Incidence angle was optimized to increase scattering from the film and

reduce scattering from the substrate. The magnitude of the q vector is controlled by

changing the in-plane incidence angle and the detector angle. Bragg reflections collected

in grazing incidence geometry correspond to repeat distances existing within the plane of

the substrate. The geometry is shown in Figure 2-9 below.

For films with an anisotropic structure in the plane of the film (i.e. oriented

samples), beam line 7-2 can be used to perform ϕ scans. In a ϕ scan, the magnitude of

the in-plane scattering vector is set to satisfy a chosen Bragg reflection, qB, and the

substrate is rotated with respect to the ϕ axis (Figure 2-9). Diffracted intensity

corresponds to the crystalline orientation distribution in the plane of the film. Background

ϕ scans are measured and subtracted from the final intensity. As in the other experiments,

samples are placed in a chamber purged with helium to reduce beam damage and air

scattering background.

Figure 2-9. Left: Grazing incidence X-ray diffraction geometry. Both ko, the incoming radiation beam, and k, the diffracted beam, lie nearly in the plane of the substrate; q is the scattering vector. ϕ scans are measured by rotating around the ϕ (vertical) axis. The RLPs (not shown) for an isotropic in-plane texture form a ring in the plane of the substrate. Right: Grazing incidence X-ray diffraction probes repeat distances in the plane of the substrate.


Beamline details

Two-dimensional grazing incidence X-ray diffraction (2D-GIXD) measurements

are performed at SSRL beam line 11-3. Beam line 11-3 is equipped with a MAR345

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26

image plate detector, and operates at an energy of 12.7 keV, with a photon wavelength of

0.0978 nm. Detector distance, which determines the q-range available for measurement,

is kept near 400 mm, with the exact distance calibrated using a LaB6 standard. The in-

plane component of the scattering vector projected on the detector is qxy, and the vertical

component is ~qz. Resolution is dependent on sample size, pixel size (150 µm by 150

µm), sample to detector distance, and scattering angle and estimated to be approximately

0.1 Å-1.[53]

Measurement details

As with grazing incidence diffraction with a point detector, the dependence of X-

ray penetration depth on incident angle has been exploited to collect depth-dependent

diffraction patterns[52]. In the work presented here, the grazing incidence angle is

optimized to maximize diffracted intensity from the film. The chosen incident angle is

typically between 0.09° and 0.12°. This is above the critical angle of a typical polymer

sample (α = ~0.08°) and below the critical angle of the silicon substrate (α = ~0.14°)[54],

so we are probing the entire film thickness. At the chosen incident angle, the beam spills

over the entire sample (in the dimension parallel to the beam). Two-dimensional area

diffraction patterns display a large slice of reciprocal space, revealing features

corresponding to both in-plane and out-of-plane repeat vectors (Figure 2-10). However,

intensity projected on the detector is not a true map of reciprocal space. Distortion exists

at values of nonzero qz, and for this reason a vertical slice along qx,y=0 is similar to, but

not equivalent to, a specular diffraction scan. For this reason, as well as the low

resolution of the area detector, these diffraction patterns are not ideal for quantitative

measurements. The detailed geometry of a two-dimensional grazing incidence X-ray

diffraction, along with how to use this beam line to collect quantitative texture

information, will be discussed in the next section. As in the other experiments, samples

are placed in a chamber purged with helium to reduce beam damage and air scattering

background.

27

Figure 2-10: Geometry of grazing incidence X-ray diffraction with a two-dimensional image plate detector.

2.3 Pole Figures of Thin Films using X-ray Diffraction and an

Area Detector: Quantifying Texture and Relative Degree

of Crystallinity

Traditional X-ray diffraction techniques for collecting pole figures have limitations

with regard to diffraction geometry, sample exposure duration and ease of application. To

satisfy a need in the field for more thorough characterization of organic semiconductors’

semicrystalline microstructures, we developed a method for collecting quantitative pole

figures with an area detector and a synchrotron light source for thin films with isotropic

crystallographic orientation in the substrate plane (classically referred to as fiber texture).

In addition to providing information about the crystallographic texture, we can use pole

figures to quantify the relative degree of crystallinity, which is an important characteristic

of semicrystalline, semiconducting polymers. The technique is rapid and ideal for thin

films of materials that are sensitive to beam damage or diffract weakly.

28

2.3.1 Introduction to Pole Figures

A pole figure is a plot of the orientation distribution of a particular set of

crystallographic reciprocal lattice planes, providing a useful illustration of a material’s

texture. Pole figures allow for the quantitative characterization of film texture. The

integrated intensity of a pole figure is directly proportional to a film’s degree of

crystallinity.

Traditional pole figures of bulk samples can be collected in either reflection or

transmission mode. Pole figures collected in reflection mode utilize a symmetric

geometry introduced by Schultz[55-57]. In this technique, diffraction intensities are

collected using a point detector as the sample is rotated along two axes. Accurate

collection of intensity in the Schultz geometry is generally limited to within 85° of the

surface normal, due to distortions that arise at the substrate edge. Transmission

techniques[58-60] are feasible but require either intensity corrections or special sample

shapes, and in general cannot be used for thin films for which the substrate absorbs the

X-rays. Whether dealing with reflection or transmission geometries, collection with a

point detector across such a large slice of the reciprocal lattice space is time-intensive,

especially for thin films. With radiation-sensitive thin films of organic polymers,

extensive time is associated with increased risk of beam damage, which lowers the

diffracted intensity. The use of an area detector facilitates more rapid collection of

intensity across a large section of reciprocal space, greatly decreasing total acquisition

time. Area detectors are commonly used with transmission-based geometries, allowing

for the simultaneous collection of Debye-Scherrer rings from multiple Bragg

reflections[61, 62]. The use of flat area detectors in combination with reflection

geometries for thin films results in some image distortions, making quantitative analysis

difficult. Many research groups have reported useful, but for the most part only

qualitative, texture data of delicate thin films collected using grazing incidence

synchrotron radiation with an area detector.[43, 54, 63]

29

2.3.2 Measuring Pole Figures with Area Detectors

We present a method to measure X-ray diffraction pole figures for fiber textured

thin films using synchrotron radiation and three separate measurements, taking advantage

of the benefits of an area detector. The first measurement collects diffraction intensity

with an image plate in grazing incidence geometry. Using a grazing incidence X-ray

beam for diffraction greatly increases the signal-to-noise ratio by increasing the X-ray

path length through the film, allowing for accurate intensity collection from weakly

diffracting samples such as organic semiconducting thin films. The large beam footprint

also serves to spread the beam power across the sample, resulting in less destructive data

collection for radiation sensitive samples. However, raw intensity collected with a flat

detector in grazing incidence geometry is distorted [64, 65], and the detector image is not

a direct map of reciprocal space. This makes it difficult to use the intensity on the

detector for quantitative analysis. By appropriately combining data from grazing

incidence diffraction patterns with data from local-specular diffraction patterns (where

the proper choice of incidence angle prevents distortion near the Bragg reflection), we

obtain diffracted intensity that represents the true intensity of a Bragg reflection across

the entire span of polar angles (-90o to 90o). The third measurement uses a point detector

in specular geometry to collect intensity from crystallites perfectly oriented with respect

to the substrate. This measurement provides higher resolution information of the fine

structure near χ=0°, and allows us to be quantitative when comparing intensities between

different samples.

In the compilation of complete pole figures, we make the following assumptions:

(a) The sample has an isotropic crystallite orientation distribution in the plane of

the substrate (i.e. has a fiber texture) or fiber texture is artificially created by

rotating the substrate normal throughout the measurement. We do not assume that

fiber texture is perfect, but more importantly, that the film has cylindrical

symmetry.

(b) The detector sensitivity is not polarization-dependent.

(c) The film thickness to be probed does not exceed ∼100 nm, with a more exact

thickness requirement depending upon the incidence angle as well as the material

30

in question. This final assumption is not a rigorous requirement, but merely a

guide. If the film is significantly thicker than the X-ray penetration depth then

only the surface region is probed[51, 54].

(d) The incident angle is set to probe the entire film, i.e. it is set above the critical

angle of the polymer film and below the critical angle of the substrate[65].

The geometry of grazing incidence diffraction with a two-dimensional image

plate is shown below in Figure 2-11(a). Here, α is the incident angle, controlled by tilt of

the substrate relative to the incoming beam. The incoming and outgoing (diffracted)

scattering vectors are defined as ko and k, respectively. L is the detector-sample distance.

The scattering angle 2θ is determined by two detector angles defined by γ in the

horizontal direction, and δ along the vertical, following the relationship:

cos(2θ)=cos(γ)cos(δ)[65, 66].

Figure 2-11. a) Schematic of relevant geometry definitions. α is the incidence angle, zSRF is the surface normal of the sample, ko is the incoming radiation wavevector, k is the outgoing (diffracted) radiation wavevector, L is the detector-sample distance, δ is the vertical angle (with respect to the incoming beam) of the diffracted beam, γ is the horizontal angle (with respect to the incoming beam) of the diffracted beam, xDRF and zDRF are the planar coordinates of the area detector, and θB is the Bragg angle. b) A sphere of possible crystallite orientations for a certain Bragg reflection. The radius of the orientation sphere is defined by the magnitude of qB. χ is defined between the surface normal and the plane of the substrate.

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When compiling pole figures for a given reciprocal lattice vector qB, we want to

extract intensity I as a function of χ. The angle χ is defined as χ=cos-1(qzsrf/||q||), where

qzsrf is the z component of q in the surface reference frame, illustrated in Figure 10(b).

The radius of the orientation sphere is defined by the magnitude of the qB of the Bragg

reflection of interest. What we measure on the detector is the q component in the detector

reference frame. For nonzero values of α, Figure 2-11 illustrates the necessity of two

different coordinate systems: one with respect to the sample and one with respect to the

detector. In order to obtain the scattering intensity I as a function of χ, the components of

the scattering vector in the sample reference frame must be determined. For this, the

reader is referred to Section 2.5, Calculation I at the end of this chapter.

It is important to understand the distortion in intensity associated with grazing

incidence diffraction and a planar two-dimensional image plate. In the case of grazing

incidence, for most calculations in this work, the sample is assumed to be horizontal, with

α near zero. (In practice, α varies between 0.1 and 0.12°). It is helpful to consider the

sphere of crystallite orientations in reciprocal space and the Ewald sphere, (a). The radius

of this orientation sphere is defined by the qB of the Bragg peak of interest. The center of

this sphere corresponds to q=0 and is placed on the surface of the Ewald sphere at the

head of the incoming k vector. In the case where α=0, the surface reference frame and

detector reference frame are the same. From this construction, the Figure 2-12

intersection of these two spheres corresponds to all observable diffraction: only

reciprocal lattice features lying at the intersection of these two spheres will result in

intensity on the detector. The most vertical orientation that will result in observed

diffraction (corresponding to the minimum detectable χ) has a polar angle of θB. In other

words, rather than ranging [-90° to 0° to 90°], χ is constrained to a range of [-90° to -θB°

and θB° to 90°] in the grazing incidence geometry. The intensity displayed along a

vertical slice of the detector arises from crystallites with χ near θB. Interpretation of a

vertical slice as the equivalent of a specular diffraction scan is not correct, and therefore,

sometimes data within this range is not displayed at all[65].

32

Figure 2-12: Ewald sphere and orientation sphere construction a) in grazing incidence diffraction, where α=0 and b) in local specular diffraction condition, where α=θB.

Specular condition at the center of the pole cannot be reached in grazing

incidence, but we can travel through reciprocal space by changing the incident angle of

the incoming beam (defined by the sample tilt.) When the incident angle is set to the

Bragg condition, the orientation sphere intersects the Ewald sphere as shown in Figure 2-

12(b). In this work, these area diffraction patterns are referred to as local specular scans,

because specular condition is reached at a location determined by the incident angle. In

practice, the incident angle is rocked by 0.2° on either side of θB during the measurement

to assure that we are capturing the maximum peak intensity. As mentioned previously,

grazing incidence diffraction measurements introduce a horizon at qz=0, limiting the

collection of diffraction to values above qz=0. When the sample tilt is increased, the

sample physically blocks a larger amount of diffraction, placing a limitation on the

maximum available χ. When incident angle = θB, χmax is 60°, derived in Section 2.4,

Calculation II.

The first two measurements of the pole figures discussed in this work use data

collected in the geometries described above. First, we collect a 2D-GIXD pattern taken at

grazing incidence, Figure 2-13(a). Using WxDiff software[67] we extract intensity as a

function of χ as seen on the detector, keeping in mind that intensity within ±θB of the

surface normal is distorted from true specular condition. Care is taken to radially

integrate intensity within an arc with a width of Δq (typically around 0.2 Å-1) that

encompasses the entire peak, centered at qB, as shown in Figure 2-13(a). Appropriate

background subtractions are incorporated. Next, a local specular diffraction pattern is

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collected with an incidence angle equal to the θB of the Bragg reflection of interest. In the

same way, WxDiff software is used to extract intensity as a function of χ, keeping in

mind that diffraction is physically blocked for >60°, Figure 2-13(b).

The data in the I vs. χ plots extracted from the grazing incidence and local

specular two-dimensional diffraction patterns are not on an absolute scale. To create the

pole figure, the two datasets can be scaled against each other and appropriately merged

where the curves overlay (their local derivatives are equal.) The relevant angular range of

the grazing incidence data is [θB < χ < 90°] and [-90° < χ < -θB] and the relevant angular

range for the local specular data is [0 < χ < 60°] and [-60°< χ<0]. The appropriate angles

to stitch the two datasets exist between θB and 60° (or the negative equivalent). Within

this range, the exact choice of stitching angles is somewhat arbitrary and depends in part

on the signal to noise ratio of the two diffraction patterns.

At this point the pole figure accurately represents intensity from all of the

diffracting material and can be used for quantitative texture analysis. These pole figures

have been used to characterize crystallite orientation in CdSe nanorods[68] and

polythiophenes[69]. However, we now proceed one step beyond the procedure discussed

in Baker et al.[70] and perform a third measurement, collecting high-resolution intensity

information of the chosen Bragg reflection as a function of χ near the substrate normal

with a point detector, Figure 2-13(c). In the literature as a high-resolution rocking

curve[48, 71, 72]. Due to the geometry of the experimental setup, as well as time

constraint, intensity is only collected within approximately ±2° of χ=0°. The I(χ) data

extracted from the area diffraction patterns are scaled to meet the intensity data collected

from the point detector. The I versus χ dataset from the point detector replaces the inner

few degrees of the local specular pattern intensity. This scan is important for two reasons.

First, it allows us to collect fine structure near χ=0°. Often, there is a resolution limited

peak attributed to crystallites oriented exactly parallel to the substrate normal, not

captured with the poor resolution of the image plate. Second, the well-defined beam

footprint and therefore incident radiation allows us to be quantitative with the diffracted

intensity, normalizing samples with respect to thickness and comparing the relative

degree of crystallinity. Ultimately, this quantitative intensity should also allow for

calculation of absolute crystallinity, as discussed later. Figure 2-13(c) shows the data

34

correctly overlain: in the final pole figure, intensity from the 2D-GIXD is kept for higher

angles of χ, intensity from the local specular for intermediate angles, and intensity from

the high resolution rocking for the small angles of χ.

Figure 2-13. Steps involved in pole figure measurement for the (200) Bragg reflection of a thin film of annealed PQT. a) 2D-GIXD pattern (left) and intensity as a function of χ (right). The region of radially integrated intensity is marked on the 2D-GIXD pattern, the arc is centered around q200, with a width of Δq = ~0.2 Å-1. (The values shown at low angle of χ are understood to be distorted.) b) 2D-Local Specular curve, taken with the α=θB, such that specular condition is achieved locally at the (200) Bragg reflection (left). Inset shows zoomed in region around the (200) Bragg reflection. Intensity as a function of χ (right), as marked in the local specular curve. c) High resolution rocking curve data, collected with a point detector (left). Intensity data from the three measurements overlain (right).

2.3.3 Corrections

WxDiff software corrects collected data for the nonlinear relationship between

pixel position and q value on a flat image plate detector. WxDiff also incorporates a

correction for polarization inherent to synchrotron radiation, according to Equation (2-11)

below. Here, horizontal polarization ph is assumed to be ~95% for synchrotron radiation.

35

(Eq. 2-11)

An absorbance correction should be considered. The path an X-ray follows through a film

is longer at low exit angles than at high angles. If the absorbance of the film is high, this

can affect the diffracted intensity data. Assuming a uniform sample absorbance, the

expression for the correction is[49, 73]

(Eq. 3-12)

where µ is the extinction coefficient (or inverse of the absorption length), t is the film

thickness and α and δ are defined previously. For the case of thin polymer films (<~100

nm), the correction for absorption is negligibly small (<1%).

A Lorentz correction[66, 74, 75] is required in X-ray geometries that involve

either a rotating single crystal or a stationary powder sample. However, in the present

case of morphology quantification (as opposed to structural characterization), the data are

obtained from a single Bragg ring, which obviates the need for a Lorentz correction.

The integrated intensity of a Bragg reflection is linearly proportional to the

diffracting material in the direction normal to the substrate. Thus, thickness normalization

is incorporated by dividing pole figure intensity by film thickness. Data are plotted as

intensity/nm.

2.3.4 Examples of Complete Pole Figures

Two examples of pole figures of the (100) Bragg reflection of two poly[N,N 0-

bis(2-octyldodecyl)-napthalene-1,4] films, (PNDI) films are shown in Figures 2-14 and

2-15. PNDI is a high-performance n-type semiconducting polymer that exhibits an

unusual face-on texture on casting[76]. Pole figures were used to investigate the change

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36

in this texture as a function of thermal treatment. Figure 2-14 shows the complete pole

figures in their one-dimensional form, clearly illustrating the change in texture and

crystallinity in PNDI when annealed below and above the melting temperature (Tm). In

the film annealed below Tm, there are two populations of crystallites: those with their

(100) reciprocal lattice vector oriented parallel to the plane of the substrate, giving rise to

peaks at χ=±90°, and crystallites oriented with the (100) reciprocal lattice vector along

the substrate normal, giving rise to a peak near χ=0°. When the film is annealed above Tm,

the population of crystallites oriented parallel to substrate normal decreases, while the

population of crystallites oriented out of the plane increases, and the peak at χ=0°

dominates. The pole figure of the film annealed above Tm shows a resolution-limited

peak near χ=0°, which was not visible in the other PNDI pole figures. It has been shown

that this peak arises due to perfectly oriented crystallites that have nucleated off the very

flat silicon substrate surface[71]. These crystallites are oriented parallel to the surface

normal, and the resolution limits of the equipment result in the delta-like shape of the

peak. This region is designated by the box in Figure 2-14(a) and is enlarged in Figure 2-

14(b).

Figure 2-14. One-dimensional pole figures of PNDI thin films a) annealed above Tm (black) and annealed below Tm (dashed). b) Zoomed in region of the boxed region in (a), highlighting the resolution limited peak in the pole figure for the annealed PNDI film

Another way to show the pole figure data is to superimpose diffraction intensity onto the

orientation sphere corresponding to the Bragg reflection of interest. Pole figure data of

37

the two PNDI films are shown in this form in Figure 2-15. This is the same data as

shown in Figure 2-14. When viewed from the top, these three-dimensional pole figures

resemble two-dimensional stereograms as often seen in the literature. The one-

dimensional pole figures are useful for easy comparison between samples, while the

orientation spheres better represent the three dimensional nature of reciprocal space, and

the assumed cylindrical symmetry.

Figure 2-15. Intensity mapped onto the (100) orientation sphere, forming three-dimensional pole figures of PNDI films a) annealed below the Tm b) annealed above the Tm. Top figures show the orientation sphere in perspective and the bottom figures show the orientation sphere form a top-down view.

2.3.5 Quantitative Texture and Crystallinity Using Pole Figures

Pole figures can be used to quantitatively compare the relative degree of crystallinity

between films, and quantify texture. The relative degree of crystallinity is calculated

according to the equation below.

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DoC∝ΔβΔθ[Ipeak − Ibase ]+ 2π sin(χ)I(χ)dχ0

π / 2∫ (Eq. 3-13)

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In the calculation, the resolution-limited peak (if it exists) and the slowly varying

intensity are treated separately. The shape of the resolution-limited function is governed

by details of the diffraction equipment used to collect the high-resolution rocking curve.

This intensity is addressed in the first part of Equation (2-13). Here, Ipeak and Ibase are the

maximum and minimum intensity of the resolution limited peak, as illustrated in Figure

2-14(b). The angular acceptance of the diffracted beam is reflected in the values Δβ and

Δ2θ, which is set by either vertical slits or a crystal analyzer. The slowly varying portion

corresponds to diffracted intensity collected on the two-dimensional image plate, and is

accounted for in the second part of Equation (2-13). To calculate the relative degree of

crystallinity, the intensity is integrated over the solid angle, which introduces the 2π and

sin(χ) terms. The sum of the integrated intensity of the slowly varying peak and the

resolution limited peak is directly proportional to the degree of crystallinity of the film.

For a detailed derivation of Equation (2-13), the reader is referred to Section 2.4,

Calculation III.

Using the above equation and the example PNDI data shown in Figures 2-14 and

2-15, PNDI annealed above the Tm is twice as crystalline as PNDI annealed below the Tm.

We can calculate crystallinity existing within designated ranges of χ to make quantitative

statements about film texture. In this case, we define crystallites with their (100) lamellar

stacking oriented within 30° degrees of the surface normal to have an edge-on texture (-

30° < χ < 30°) and crystallites with their lamellar direction within 30° of the in-plane

direction to be face-on (-90° < χ < -60°, and 60° < χ < 90°). According to these

definitions, 77% of crystallites in the PNDI films annealed below Tm are face-on, while

in films annealed above Tm, 94.6% of the crystallites are edge-on.

It is important to understand sources of error in the values of relative degree of

crystallinity. Error may arise during normalization with respect to thickness, if film

thickness is not constant across the sample and the film area where thickness is measured

(via Atomic Force Microscopy or X-ray reflectivity) is not the exact area surveyed by the

beam. This error, however, is expected to be small. A larger source of error may arise

during the process of stitching the high-resolution rocking curve and local specular curve.

If the counting time is too low during the collection of the high resolution rocking curves,

the intensity at the base of the curve will not be well-defined, due to the low signal to

39

noise ratio. Intensity fluctuations will be present, and the integrated intensity will

therefore be dependent on the exact point chosen as the stitching angle. This is more of a

problem when a crystal analyzer is used and the diffracted beam flux is low. Further error

can arise when defining the noise floor of the collected intensity. The signal to noise ratio

in both the high resolution rocking curves and the two-dimensional data sets can be

increased by increasing the count time, thus reducing the combined error, at the expense

of beam damage to the radiation sensitive samples. To minimize beam damage, samples

were enclosed in a helium chamber. In the work presented, care was taken to evaluate

the error in the relative degree of crystallinity introduced by the above variables. In most

cases this was under 10%.1

2.3.6 Towards Absolute Crystallinity

The current method of pole figure compilation is a powerful characterization

technique for semicrystalline thin films. Pole figures can be used to quantitatively

measure the relative degree of crystallinity and simultaneously collect information on

film texture. While it is useful to compare relative degrees of crystallinity, being able to

put these values on an absolute scale is ideal. Not only would this be the first diffraction

based method to measure the absolute degree of crystallinity that accurately accounts for

diffraction from misoriented crystallites in fiber textured films, but it would also allow

for more meaningful comparison across different sample sets and different materials.

There are two possible approaches for putting the relative values of crystallinity on

an absolute scale: we must either compare the integrated intensity of semicrystalline films

to the integrated intensity of a 100% crystalline film of the same material, or theoretically

calculate the intensity expected from a 100% crystalline film. The former strategy is near

impossible, in that long chain polymers are not able to form 100% crystalline films due to

both crystallization kinetics and polydispersity among molecular weights. Thin films of

oligomers based on the same molecular structure as the corresponding polymer tend to be

more crystalline, but often pack with a different unit cell, changing diffraction

1 For the one case where expected error exceeded 15%, the data was excluded and only the shape of the pole figure was analyzed.

40

characteristics and expected integrated intensity. We therefore choose the latter route, and

aim to calculate the expected integrated intensity. The integrated intensity for Bragg

reflection (hkl) depends on the square of the structure factor, Fhkl. The calculation of a

structure factor is an important step in the calculation of an expected integrated intensity.

The atomic positions are often difficult to predict in extremely disordered semicrystalline

polymers. Annealed films of poly(2,5-bis(3-akylthiophene-2-yl)thieno[3,2-b]thiophene))

(PBTTT) present sufficient order to make theoretical structure predictions possible. A

PBTTT unit cell with atomic positions has been proposed and is in agreement with other

theoretical structures. Future work will be focused on using these atomic positions to

calculate a structure factor with confidence. The remaining factors modulating the

integrated intensity include details related to the diffraction geometry, material absorption

(neglible with low density polymer thin films) and thermal factors. A complete

theoretical calculation of the integrated intensity for PBTTT can be achieved with the

above information. This would enable the measurement of absolute crystallinity to

PBTTT films, a very important step in understanding microstructure and charge transport

in organic semiconductors.

2.4 Conclusions

The methods discussed in this chapter will be used throughout this thesis to

characterize the electronic and structural characteristics of semiconducting,

semicrystalline thin films. The ability to quantify details of the electronic structure using

the ME model as well as details of the crystalline structure using X-ray diffraction pole

figures allows for very thorough investigation of the microstructural relationships

governing charge transport.

41

2.5 Calculations

2.5.1 Calculation I: Transformation of Coordinate System

In order to extract the pole figure for the given reciprocal lattice vector qB, the X-ray

scattering intensity, I, needs to be evaluated as a function of the angle χ between the

scattering vector q and the surface normal. Refer to Figure 2-11 in main text of

manuscript for geometrical definitions.

In an in-plane isotropic powder sample, the pole vector is parallel to the z-axis in the

sample reference frame (zSRF). The angle χ is then determined by: χ = cos-1(qzSRF / ||q||),

where qzSRF refers to the z-component of q in the sample reference frame. In order to

obtain the scattering intensity, I, as a function of χ, the components of the scattering

vector in the sample reference frame must be determined.

However, the q-vector can directly be measured on the detector (detector reference

frame). In general, q is given by:

where ek0 and ek are the unit vectors in the directions of the incident beam and the

scattered beam, respectively. In the detector frame as depicted in Figure 2-11, ek is

conveniently defined in spherical coordinates by the azimuthal angle γ and the elevation

angle δ:

As shown in Figure 2-11, L, the sample to detector distance, δ and γ are easily calculated

using the pixel position (x,z) on the detector.

42

The q vector in the detector frame, qDRF, is then given by:

However, for the calculation of I(χ), we need the representation of q in the surface

reference frame, qSRF. Relative to the detector frame, the sample frame is tilted

backwards by the angle α as expressed by the rotation matrix Rα:

Therefore qSRF can be calculated using the inverse of the matrix Rα by:

and thus, the z-component qzSRF in particular is given by:

The absolute value of the scattering vector, |q|, is easily determined from the total

scattering angle 2θ. In the non-grazing geometry, 2θ is easily determined from the !

qzSRF

= k[cos(")sin(#) + sin(")(1$ cos(#)cos(%))]

43

position of a spot on the detector (x,z), as follows: With tan2(2θ) = (x2+z2)/L2 and Bragg’s

law, one finds:

In grazing geometry, i.e., for incidence angles below the material’s critical angle, a small

correction has to be made for the ek0 unit vector because the scattering results from an

evanescent wave field travelling in the sample surface:

The derivation of the qzSRF expression derived here for the non-grazing case is completely

analogous. However, the correction is very small (α is typically 0.1°) and can be safely

neglected for most purposes and the equations above can be used unchanged. With the

expressions for qzSRF and |q|, the pole figure can then be constructed from the pixels on

the screen I(x,z) by calculating I(χ), with χ = cos-1(qzSRF / ||q||)).

44

2.5.2 Calculation II: Calculating the Maximum Angle of χ for a Local Specular

Diffraction Pattern

At the maximum (nearest to horizontal) angle of χ, the diffracted beam is grazing over the

sample, meaning that the z-component of our outgoing (diffracted) k-vector in the surface

reference frame is zero, kz,oSRF = 0. Because the sample is tilted at the angle α=θB, the

incoming k-vector in the surface reference frame is kz,iSRF=-ksin(α)=-ksin(θB). Therefore,

it follows that

45

2.5.3 Calculation III: Derivation of the Relative Degree of Crystallinity

Adapted from notes by Michael F. Toney

Figure 2-16. a) Pole figure of PNDI, annealed above Tm. b) Resolution-limited peak at χ=0°, arising from perfectly oriented crystallites. c) Pole figure of PNDI, annealed above Tm, on a log-log plot. The slowly varying intensity, accounting for the remainder of the crystalline material, is highlighted in red. When calculating the degree of crystallinity, the grey and red regions of pole figure intensity are treated separately.

The pole figures measured using the method presented here can be used to derive the

relative degree of crystallinity of a thin film sample, according to the following

relationship in Equation (2-13) in the text, given below.

€

DoC∝ΔβΔθ[Ipeak − Ibase ]+ 2π sin(χ)I(χ)dχ0

π / 2∫ (Eq. 2-13)

where

is the angle q forms with the surface normal.

= horizontal resolution, typically 0.02 radians, ~.6° (set by horizontal slits in the

beam line)

= vertical angular resolution, is typically .001 radians (~.03°) when defined by the

vertical slits in the beam line or .0002 radians (~.01°) when defined by a crystal analyzer.

The first part of the equation accounts for the perfectly oriented fraction, highlighted in

grey in Figure 2-16(b), above. The second part of the equation accounts for “slowly

varying” intensity highlighted in red, Figure 2-16(c).

46

where fp and fb refer to the fraction of crystallites corresponding to the peak intensity and

the fraction of crystallites corresponding to the base intensity, respectively, and g refers

to the fraction of crystallites corresponding to the slowly varying intensity.

The pole figures are one-dimensional plots, but they represent two-dimensional intensity

of the Bragg reflection. The degree of crystallinity is proportional to the fraction of

perfectly oriented crystallites and the integration of g(χ) across the magnitude of q.

€

DoC∝ ( f p − fb ) + 2π g(χ)sin(χ)dχ0

π2∫

From pole figures, we collect the measured intensity Imeas, which is the following:

€

Imeas = R(Ω) f (χ +Ω)d∫ Ω

where is the solid angle and R( ) is the instrument resolution. Taking into account

the two populations of crystallites (perfectly oriented and slowly varying orientation), the

equation can be rearranged as follows:

€

Imeas = ( f p− fb )R(Ωχ= 0) + R(Ω)dΩg(θ +∫ Ω)

where

€

R(Ωχ= 0) =1 if

€

χ ≤Δθ2

€

R(Ωχ= 0) = 0 otherwise.

and

€

R(Ω)dΩg(θ +∫ Ω) =

€

ΔθΔβg(θ)

!

f (") = ( f p # fb )$(") + g(")

47

Thus, total crystallinity can be calculated from the complete pole figures according to the

equation below.

€

DoC∝ (Ip − Ib ) +2π

ΔθΔβsin(χ)Im (χ)dχ0

π2∫

Which is equivalent to Equation (2-13).

48

3 Understanding Processing Conditions,

Microstructure and Charge Transport in Thin Films

of PQT-12

In this chapter, we aim to develop a more thorough understanding of how surface

chemistry and thermal annealing affect the microstructure of thin films of poly[5,5’-

bis(3-dodecyl-2-thienyl)-2,2’-bithiophene] (PQT-12), a semicrystalline polythiophene,

and in turn how the microstructure is related to the film’s electronic properties. X-ray

diffraction is used to collect information about the crystalline microstructure, which is

then related to the electronic structure and the field-effect mobility. By increasing the

understanding of complex relationships that control TFT performance, we will be able to

move towards the establishment of design and processing rules for optimized organic

electronic devices.

3.1 Introduction

The most widely studied semicrystalline polymer is P3HT. Previous studies have

shown that the field-effect mobility of carriers in P3HT depends on the nature of the

semiconductor-dielectric interface,[27, 77, 78] the molecular weight of the polymer,[79-

83] and the conditions used for film formation such as solvent and drying time.[84] Most

TFTs are fabricated using spin coating; this coating method is known to form films with

microstructures that can be kinetically limited due to the short drying time. In general,

hydrophobic gate dielectric layers produce higher performance TFTs than those

fabricated with hydrophilic dielectrics.[32, 85] As introduced in Chapter 1, the

microstructure of P3HT is semicrystalline with a fiber texture. The out-of-plane texturing

has been shown to depend on the molecular weight, the substrate and the casting

solvent.[77-79, 81, 82] Typically, in crystallites of P3HT, the side chains are considered

to be mostly disordered[86].

49

PQT-12 has the same conjugated backbone as P3HT (Figure 3-1) but generally

has a higher field-effect mobility.[26, 29, 87, 88] In PQT-12 the density of the 12 atom

alkyl side chain along the backbone is half that of P3HT. Unlike P3HT, during thermal

annealing PQT-12 undergoes a thermal transition that allows for reordering in thin films

and in powder form. More specifically, PQT-12 shows two thermal transitions in

differential scanning calorimetry data that have been attributed to phase transitions of the

crystal to liquid crystal (~120 °C) and the liquid crystal to melt (~150 °C)[88]. Mobility

as a function of surface treatment and thermal processing is shown in Table 3-1.

PQT-12 µ (cm2/Vs)

As Cast on OTS/SiO2 0.004 Annealed on OTS/SiO2 0.1 As Cast on SiO2 No measurable field-effect Annealed on SiO2 0.0002-.0005

Table 3-1: Mobility extracted from TFT measurements for PQT-12 films as-spun and annealed on OTS/SiO2 and untreated SiO2 substrates.

Based on molecular models and low-resolution diffraction data, the basic

polythiophene stacking structure has been assumed, Figure 3-1. In contrast to P3HT, the

peak positions of the methylene stretches measured by infrared spectroscopy show that

the side chains are relatively well ordered after annealing and are only slightly more

disordered prior to annealing, supporting the fully extended side chains Figure 3-1.[89]

Based on density considerations, the side chains are interdigitated and the backbones are

π-π stacked.[90] This simple packing model is consistent with most of the reported

experimental data and likely captures the relevant features of the packing that can be used

to understand the crystalline order in the films.

50

Figure 3-1. Unit cell of PQT-12, showing the reciprocal lattice vectors corresponding to the π-π stacking repeat (010), the alkyl stacking repeat (h00) and the repeat along the chain backbone (00l).

Transport in PQT-12 TFTs has been extensively examined and used to infer

details about the nature of the electronic structure of the films. Charge transport, as

characterized by field-effect mobility in TFTs, depends dramatically on the chemistry of

the dielectric/semiconductor interface and on thermal treatments of the films (Table 3-

1).[40, 87, 91] Depending on the specific processing conditions, mobility can vary over

many orders of magnitude. The temperature dependences of mobility in PQT-12 and

P3HT have been modeled with a mobility edge (ME) model.[28, 92] A discussion of the

ME model and its parameters can be found in Chapter 2.

For details of the electrical measurements of PQT-12 presented here and their

modeling, see Ref. [28]. By applying the ME model to PQT-12, it was found that thermal

annealing did not significantly affect the crystalline mobility µ0 or the total trap density,

while it had a profound effect on the trap energy distribution (Table 3-2). Upon

annealing, the gap state distribution was found to have a tighter energy distribution, as

demonstrated by the decrease in Eb. The dielectric surface chemistry on the other hand

had an effect on both µ0 and Eb. Fabricating devices on bare SiO2 rather than OTS-treated

SiO2 caused a decrease in µ0 and a broadening of the trap state distribution (i.e., an

increase in Eb) and an increase in Ntot.

51

PQT-12 µ0(cm2/Vs) Eb(eV) Ntot(cm-3)

As Cast on OTS/SiO2 1-4 50 5.5×1020 Annealed on OTS/SiO2 1-4 34 6.5×1020

Annealed on SiO2 0.05 50 1.7×1021 Table 3-2: ME fitting parameters for PQT-12 films as-spun and annealed on OTS/SiO2 and annealed on SiO2.

The origins of the changes in electronic structure remain poorly understood. For

instance, phase contrast in tapping mode atomic force microscopy has been reported to

show that PQT-12 films annealed on bare SiO2 look identical to films annealed on OTS-

coated SiO2 (Figure 3-2)[87, 88]. AFM characterizes only the top surface and gives no

insight into the large change in mobility and ME parameters. Here, we connect the

previously described results from the ME model with processing variables and

corresponding changes in the microstructure of the PQT-12 thin films studied by X-ray

diffraction.[48]

Figure 3-2. AFM images of annealed PQT-12 thin films a) on bare SiO2, topography mode b) and phase mode c) on OTS-8 treated SiO2, topography mode and d) phase mode. Images reproduced with permission form Ref. [87].

52

3.2 Experimental Details

Poly[5,5’-bis(3-dodecyl-2-thienyl)-2,2’-bithiophene] (PQT-12) was obtained from

Xerox Research Center, Canada. Solutions were spin cast from dilute solutions in 1,2-

dichlorobenzene, as discussed in Chapter 2. For films spun on OTS-treated SiO2, OTS

was deposited via solution as discussed in Chapter 2. Annealed films were heated to

140°C on a hot plate under nitrogen and cooled to room temperature by switching off the

hot plate.

3.3 Results

3.3.1 Grazing Incidence X-ray Diffraction (GIXD)

Shown in Figure 3-2(a) are the high-resolution grazing incidence diffraction data

taken with a point detector for PQT on OTS/SiO2, both as-spun and annealed, and PQT

on bare SiO2, annealed. Peak positions and values of FWHM are shown in Table 3-3. No

features in grazing incidence diffraction were observed for the as-spun film on SiO2. For

all three samples, the most intense peak appears at a q-value corresponding to a distance

of ~3.7 Å. The high resolution of the grazing incidence patterns allows us to deconvolute

the main peak into two peaks occurring at qx,y = 1.68 and 1.71 Å-1. This peak has

considerably lower intensity for the PQT film spun on untreated SiO2. Two additional

peaks occur at qx,y = 1.22 and 1.4 Å-1. A broad feature spanning from qx,y = 1.2 Å-1 to qx,y

= 2.0 Å-1 is also present in all three diffraction patterns; because this feature is

substantially wider than the other peaks, it is likely due to scattering from disordered

regions in the film. Interestingly, the as-spun and the annealed films on OTS/SiO2 have

very similar grazing incidence diffraction patterns. The only notable change brought

about by thermal annealing is the appearance of a weak alkyl stacking peak at qx,y =

0.36 Å-1. This peak is also visible for the annealed film on SiO2. The d-spacing of the π-π

stacking repeat in PQT-12 shows a sample-to-sample variation between films of about

0.2 Å prior to annealing, but it is the same for all samples after annealing.

53

Figure 3-3. a) Grazing incidence diffraction patterns of PQT-12 on OTS/SiO2, annealed and as-spun (black, grey, respectively) and SiO2, annealed (dashed green). b) Specular diffraction patterns of PQT-12 on OTS/SiO2, annealed and as-spun (black, grey, respectively) and SiO2, annealed (dashed green).

PQT-12 Index q (Å-1) FWHM (Å-1)

As Cast on OTS/SiO2 1.219 0.024 1.37 0.122 1.512 010 1.6508 0.028 010 1.705 0.05 1.82 0.01

Annealed on OTS/SiO2 1.219 0.018 1.353 0.08 1.512 010 1.6508 0.095 010 1.71 0.04 1.82 0.1

Annealed on SiO2 1.219 0.018

1.353 0.09 1.512 010 1.651 0.08 010 1.710 0.04 1.82 0.1

Table 3-3: Positions and FWHM of diffraction peaks in the GIXD spectra of the PQT-12 films.

!"#$"%&#'!"#$%$&!

()*'+"'()&!!! (,+"'

()&!!!

!"#$"%&#'!"#$%$&!

!""#$%#&'()*+,+-).(

!/(+01"'()*+,+-).(

!""#$%#&'(+-).(

54

3.3.2 Specular Diffraction

Shown in Figure 3-3(b) are the high-resolution specular diffraction data taken

with a point detector for PQT on OTS/SiO2, both as-spun and annealed, and PQT on bare

SiO2, annealed. The positions and FWHMs of the diffraction peaks in the specular

direction for PQT films spun on OTS/SiO2, and PQT films spun on SiO2 are summarized

in Table 3-4. Only a very weak shoulder near the expected (100) peak position was seen

for the PQT film as-spun on SiO2. For sake of clarity, these data are not shown. In the

specular diffraction of the as-spun film on OTS/SiO2, a peak at qz = 0.305 Å-1

corresponding to the first order alkyl stacking (100) barely emerges above the

background. Because the background (due to the reflectivity from the polymer/substrate

interface) drops sharply with diffraction angle the (200) and (300) peaks are well defined.

These latter two peaks give an average d-spacing of 19.8 Å. A weak peak at qz = 1.55 Å-1

is likely due to diffraction from amorphous material. When the film is annealed on

OTS/SiO2, all peaks move to higher qz, increase in intensity and decrease in width, giving

an average d-spacing of 17.5 Å-1, which is in agreement with the previously determined

value ~18 Å-1 from lower-resolution diffraction.[26, 29] The disorder peak at qz = 1.5

Å−1 is no longer present. Finally, the (200) peak is a convolution of two peaks, one

located at the q value where the corresponding peak was found in the as-spun material

and the other at a slightly larger value of q. Although this observation indicates that the

(100) peak should be composed of two peaks, we cannot resolve them due to the close

proximity of the peaks and the high background scattering at low q. A corresponding

very weak shoulder also appears to be present near the (300) peak. For the PQT annealed

on SiO2 three orders of (h00) peaks are visible, with positions very close to those

observed for the annealed film on OTS/SiO2. The surface of the PQT film on SiO2 was

particularly smooth, and finite thickness fringes are visible. By measuring the fringe

spacing, Δ, the thickness of the annealed PQT-12 film on oxide can be estimated

according to the relationship (2π)/Δ=t. The thickness obtained from the finite thickness

fringes (33.6 nm) is in good agreement with this thickness obtained from AFM (40 nm).

55

PQT-12 Index q (Å-1) FWHM (Å-1)

As-spun on OTS/SiO2 100 0.305 0.0435 200 0.637 0.069 300 0.949 0.057

Annealed on OTS/SiO2 100 0.359 0.013 200 0.719 0.023 300 1.08 0.037

Annealed on SiO2 100 0.354 0.0127 200 0.726 0.0321 300 1.09 0.04

Table 3-4: Positions and FWHM of diffraction peaks in the specular patterns of the PQT-12 films.

3.3.3 Two-dimensional Grazing Incidence X-ray Diffraction

Grazing incidence X-ray diffraction with an area detector (2D-GIXD) was used to

look at a larger slice of reciprocal space, Figure 3-4(a-d). For PQT as-spun on OTS/SiO2,

two orders of the (h00) peak appear along the qz axis (Figure 3-4(b)). As mentioned

previously, the intensity profile projected along this axis is not true specular condition.

Weaker peaks along the qx,y axis are also visible. The most intense of these features is a

peak near qx,y = 1.7 Å−1; this scattering vector corresponds to the π-π stacking distance

between two cofacial polymer chains (~3.7 Å). A broad ring due to diffraction from

disordered material is visible at qx,y ~1.5 Å−1 as well. Figure 3-4(a) shows the 2D-GIXD

pattern from the same film after a 20 minute anneal at 140°C. The (100) peak intensifies,

shifts to higher qz and sharpens, losing some of the broadening that is present in the as-

spun material. The (200) peak becomes considerably more intense and a (300) peak is

clearly visible. As before, all (h00) peaks lie along the nominal qz axis. More peaks along

the qx,y axis are present as well. The peak located at qx,y=1.7 Å−1 increases in intensity

upon annealing, while the disorder peak at q=1.5 Å−1 disappears. Several mixed index

peaks off the qz and qx,y axes appear as well. It should be noted that the mixed-index

peaks and the (h00) peaks are not aligned horizontally, showing that the lattice is not

orthorhombic.

The 2D-GIXD patterns of PQT-12 films annealed and as-spun on untreated SiO2

are shown in Figure 3-4 (c,d). Trends similar to those seen for the films on OTS/SiO2 are

56

observed. The pattern of the as-spun film shows two (h00) peaks along the qz axis and a

very faint peak along the qx,y axis. On annealing, the (h00) peaks sharpen and intensify,

revealing the presence of a third order (h00) peak. The peaks along qx,y intensify and

mixed index peaks appear.

Figure 3-4. 2D-GIXD patterns for PQT-12 a) annealed on OTS/SiO2 b) as-spun on OTS/SiO2 c) annealed on SiO2 and d) as-spun on SiO2. The plots on the right are vertical slices of the intensity at qx,y=0.

3.3.4 Pole Figures

By appropriately combining data from high-resolution rocking curves, 2D-GIXD

and 2D-Local Specular curves we constructed pole figures of the (200) Bragg

reflection[70], as discussed in Chapter 2. Specular and grazing incidence X-ray

diffraction is useful for extracting accurate positions and FWHMs. However, these

geometries probe intensity along a small slice of reciprocal space, collecting diffraction

only from crystallites oriented parallel to the substrate normal. In disordered

semicrystalline polymer films, a significant amount of the crystallites may be slightly

misoriented. Pole figures for PQT films on OTS/SiO2 as-spun and annealed, and PQT

annealed on SiO2, are shown in Figure 3-5. Intensity is shown on a log-linear scale (a)

57

and a log-log scale (b). For all three films, the maximum intensity occurs near χ=0°. The

intensity near χ=50° is due to a mixed index peak that occurs at the same magnitude of q.

This intensity is difficult to substract and is therefore not taken into account in any

quantitative calculations. Looking at the log-log plot in Figure 3-5(b), it can be seen that

both of the films on OTS/SiO2 have a resolution-limited peak. This peak is attributed to

perfectly oriented crystallites that have nucleated off of the very flat dielectric

semiconductor interface. The annealed film on SiO2 does not have a resolution-limited

feature at χ=0°. However, both of the annealed films on OTS/SiO2 and on SiO2 are nearly

overlapping at high angles. The pole figures can be used to calculate the relative degree

of crystallinity of the films, normalized for thickness. The values for relative degree of

crystallinity are given in Table 3-5. Pole figures can also be used to quantify film texture

by calculating the contribution of crystallinity from various angular ranges. The

contributions to the film crystallinity from crystallites with their (h00) direction oriented

within 1° of the surface normal are also given in Table 3-5.

Figure 3-5. Pole figures of PQT-12 annealed on OTS/SiO2 (black), as-spun on OTS/SiO2 (grey) and annealed on SiO2 (green). a) Intensity is on a log scale and χ is on a linear scale. b) Same data as in (a), but on a log-log scale.

PQT-12 r. DoC within 1°

As-spun OTS/SiO2 0.4 25% Annealed OTS/SiO2 1.0 65% Annealed on SiO2 0.5 7%

Table 3-5: Values calculated from the pole figures: r. DoC and contribution to crystallinity of crystallites oriented within 1° of the substrate normal.

58

3.4 Discussion

3.4.1 Effect of Annealing

As-spun films of PQT-12 on bare SiO2 and on OTS/SiO2 exhibit the typical thin-

film texture of polythiophenes where the alkyl side chains stack in the direction normal to

the substrate, indicated by the (h00) peaks lying along the qz axis of the 2D-GIXD

patterns in Figure 3-4[43]. On annealing, regardless of the surface treatment, the

crystalline quality of the films improves. This is illustrated by the more intense and well-

defined peaks along the qxy and qz axes on the 2D-GIXD for the film on OTS/SiO2 after

annealing, and the appearance of these peaks for the film on SiO2. The appearance of new

off-axis peaks indicates the formation of three-dimensional crystallites where there is

registry of the backbones rather than the unregistered layered structure typical of liquid-

crystalline materials. This organization has been suggested previously based mainly on

the molecular structure[29]. A three-dimensional structure is supported by the infrared-

absorption frequencies of the methylene alkyl side chains that suggest that they are in a

nearly all in trans conformation.[89] We were unable to index all of the observed peaks

simultaneously with a primitive monoclinic unit cell that is consistent with the molecular

dimensions of the repeat unit of PQT-12. We therefore believe that the unit cell is likely

triclinic if there is a unique crystalline form. It is also possible that there are two

polymorphs leading to the observed pattern. The d-spacing of the alkyl stacking changes

after annealing without a substantial change in the peak positions along qx,y, suggesting

that there may be multiple polymorphs with only small variations in overall structure.

Determination of the true molecular packing in the unit cell from the diffraction data

presented here requires a combination of molecular simulation and structure factor

calculations that is beyond the scope of this work.

Some information about the change in molecular packing after annealing can be

inferred from the data without a detailed packing structure. Significant changes are

observed in the specular diffraction patterns of PQT-12 on annealing: the (h00) peaks

intensify, shift to larger qz and narrow. The shift to larger qz of all the Bragg peaks after

annealing corresponds to a decrease in interplanar spacing. Such densification has been

observed in other rigid rod polymers after thermal annealing[93]. The molecular packing

59

in the as-spun films represents a kinetically limited structure that is dictated by the time

required for evaporation of the solvent. It is likely that the flexible alkyl side chains are

disordered in the film prior to annealing because there are no mixed indexed peaks. By

heating the films into a liquid-crystalline mesophase, the side chains become mobile,

allowing them to shift into a more closely stacked arrangement upon cooling. Due to their

sparse spacing along the backbone, the side chains are able to interdigitate into an

ordered crystal upon annealing to maximize the film density.[26, 89, 93] Interestingly, in

some annealed films the (h00) peaks can be fit with a superposition of two sets of (h00)

peaks at different values of q (Figure 3-3(b)). The presence of two d-spacings in the

annealed films suggests that the two distances correspond to two polymorphic crystalline

structures of PQT-12. One of the two structures corresponds to that of as-spun PQT-12

and the other is assumed to be the lower-energy configuration obtained after annealing.

The difference in d-spacing between the initial domains and the annealed ones is

relatively small (~1 Å) and could easily be caused by small changes in the ordering of the

alkyl side chains (e.g., the tilt of the side chains relative to the backbone or the extent of

interdigitation.)

The width of the lamellar (h00) peaks can be used to estimate the size of the

crystalline domains in the lamellar direction using a Williamson-Hall plot (FWHM (Δq)

vs. h) (Figure 3-6)[94] According to established broadening models, a constant value of

Δq with respect to h would indicate that peak broadening is dominated by crystallite size

in the diffraction direction considered. The FWHMs of the (h00) peaks of the as-spun

films of PQT-12 on OTS show no systematic dependence on order. In this case, the

Scherrer equation [49] can be applied directly to the extracted FWHM value, giving a

crystallite size of approximately 16 nm in the film thickness direction. In contrast, the

widths of the peaks in the annealed films show a systematic increase with order. A linear

relationship between Δq and h indicates that part of the broadening is due to variation in

interplanar spacing between adjacent grains (nonuniform strain)[49]. On the other hand,

a linear relationship between Δq and h2 indicates that part of the broadening originates

from paracrystalline disorder: a variation in the interplanar spacing within grains that is

cumulative.[49] The intercept B at h=0, is the equivalent FWHM extrapolated to the

point where all broadening is due to crystallite size. This FWHM can be related to the

60

crystallite coherence length Lc along the diffraction direction (normal to the substrate

surface) using the Scherrer Equation (3-1): [49]

(Eq. 3-1)

Figure 3-6. Williamson-Hall plot (FWHM of (h00) peaks, versus h2) for PQT thin films, annealed on bare SiO2 (circles) and annealed on OTS/SiO2 (squares).

In our case, linear fits of Δq vs. h2 yield intercepts that are positive and correlate

well with physical parameters. The linear fit is better for the PQT-12 film annealed on

OTS/SiO2 than for that of the PQT-12 film annealed on bare SiO2. The estimated

crystallite size for the annealed film on bare SiO2 using this intercept is approximately 50

nm, which is in fair agreement with the thickness extracted previously from finite-

thickness fringes and AFM measurements. For the annealed PQT-12 film on OTS/ SiO2,

the h2=0 intercept provides an estimate of the crystallite dimension in the direction

normal to the substrate of 70 nm. AFM measurements on the same film confirmed a

thickness of approximately 100 nm. The extrapolated dimensions of the coherence

lengths in the annealed films strongly suggest that coherent domains present span the

entire thickness of the film, or nearly so, from the substrate interface to the air interface.

Crystallite growth in the direction normal to the substrate surface does not explain

how annealing improves charge transport in TFTs. Crystallite growth perpendicular to the

61

substrate is a characteristic of the entire film, but charge transport in thin film devices

occurs in-plane within approximately 1 nm of the dielectric interface[35]. Improvements

in charge transport are expected to be reflected by noticeable changes in the

microstructure in the plane of the film. However, the grazing incidence scans for the as-

spun and annealed films of PQT-12 are identical. There are no large changes in either

intensity or peak width of the π-π stacking peak. This implies that on annealing there is

little change in the crystallite coherence length within the bulk of the film in the plane of

the substrate.

The (200) pole figures provide information about the texture and crystallinity of

the film, further elucidating the changes in microstructure on annealing. Comparing the

as-spun and annealed PQT-12 film on OTS/SiO2, we see that both films have a resolution

limited central peak. This intensity arises due to crystallites that nucleate off of the very

flat dielectric/polymer interface, which is the interface important for charge transport. On

annealing, the intensity of the pole figure at lower angles increases, indicating an increase

in the overall crystallinity and a tightening of the crystallite orientation distribution. The

as-spun film is 40% as crystalline as the annealed PQT-12 film. Moreover, for the

annealed film on OTS/SiO2, more of its crystallinity can be attributed to crystallites

oriented within 1° of the surface normal. For the as-spun film, the perfectly oriented

crystallites account for a lower fraction of the crystallites present.

Thus the effect of annealing can be summarized as an increase in the population

of perfectly oriented and highly oriented crystallites, accompanied by a decrease in the

grain-to-grain out-of-plane misorientation of these crystallites (texture improves). In

addition, the crystallinity increases on annealing, with the amount of disordered material

in the film decreasing accordingly. This is seen by integrating pole figure intensity and by

the disappearance of the scattering peak due to disorder in the specular diffraction pattern.

Better packing is indicated by the appearance of mixed index peaks in the 2D-GIXS

patterns.

The X-ray diffraction data correlates well with the TFT mobility and ME

parameters. The overall mobility correlates with the fractional percent of perfectly

oriented crystallites at the dielectric interface. According to the ME model, annealing

does not influence crystalline mobility or the total trap density but tightens the trap

62

energy distribution (Table 3-2). X-ray diffraction indicates that crystallites do not grow

in the plane of the substrate on annealing, therefore the number of traps (Ntot) and the

crystalline domain size in the plane of charge transport are roughly unchanged. However,

reducing the distribution of crystallite misorientation about the surface normal by

increasing the amount of perfectly and near-perfectly oriented crystallites, as well as

reducing the amount of disordered material between the grains (i.e. increasing the

crystallinity), is expected to lead to a reduction in the number of deeper traps, which is

equivalent to a tightening of the trap distribution (Eb).

3.4.2 Effect of Surface Treatment

As mentioned previously, the final microstructure in thin films formed via spin

coating is kinetically limited. Annealing can help give polymer chains the mobility

needed to rearrange into a more thermodynamically desirable structure. It has been

shown that some surface treatments may have the same effect. The low interface energy

due to the alkyl chains of the OTS monolayer can provide enhanced mobility of the

polymer chains during the deposition process, resulting in a more crystalline film[95]. It

is also possible that the interaction between the alkyl chains of the polymer and the

methyl terminations of the OTS chains helps the crystallization of the polymer. It should

be noted, however, that our OTS layers did not display any diffraction in GIXD, strongly

suggesting that they are disordered. This is expected for OTS layers deposited via

solution. Interestingly, molecular dynamics calculations have indicated that disordered

OTS monolayers have a smoother energetic landscape and better promote the nucleation

of perfectly edge-on crystallites, compared to crystalline OTS monolayers[96]. By

comparing the diffraction scans from a PQT-12 film on bare SiO2 and the corresponding

scans from a film on OTS/SiO2, we investigate the effect that the low surface-energy

OTS monolayer has on a PQT-12 thin film.

For the PQT-12 as-spun films, the presence of the OTS monolayer promotes

crystallization and reorganization of the polymer. Indeed, both the as-spun and annealed

PQT-12 films on SiO2 show weaker diffraction intensities compared to those spin-coated

on OTS/SiO2. The high-resolution specular scan for the as-spun PQT-12 film on SiO2

63

was near featureless, while the as-spun on the OTS-treated sample exhibits three orders

of (h00) peaks.

Features in the specular and 2D-GIXD patterns after annealing indicate that the

crystallites in the PQT-12 films on OTS/SiO2 and SiO2 have a similar texture and overall

structure (Figures 3-3 and 3-4). Both scans show three orders of (h00) peaks and no

detectable disorder peak. However, the difference in intensity in the specular scans

implies a difference in the films’ diffracting material with regard to crystallite size,

details of the crystalline texture, or the degree of crystallinity. As mentioned previously,

the crystallite size can be estimated at approximately 50 nm in the film thickness

direction for the annealed film on SiO2. This is not significantly different from the

estimated thickness of the crystallites for the OTS-treated annealed film (70 nm). The

intensity difference therefore does not arise entirely from the surface-normal coherence

length of the crystallites. Instead, the difference must arise from details in the film texture

or the relative percent crystallinity. By comparing the calculated relative degree of

crystallinity, the annealed film on OTS/SiO2 (the most crystalline film in the dataset) is

twice as crystalline as an annealed film on SiO2. Moreover, the annealed film on SiO2

does not have the resolution-limited peak attributed to the perfectly oriented crystallites.

Thus, for this film, less of the crystallites present are oriented such that their diffraction

will be collected in the high-resolution specular geometry. The details of the texture and

the lower degree of crystallinity results in the large difference seen in the two specular

diffraction patterns for the annealed films on OTS/SiO2 and SiO2. We hypothesize that

the PQT-12 films on OTS owe their higher degree of crystallinity and increased

population of perfectly oriented crystallites to the faster surface diffusion of the

molecules on the low-energy surface, encouraging crystallization. The microstructural

information is summarized in Figure 3-7.

64

Figure 3-7 Microstructure of PQT-12 a) as-spun on OTS/SiO2 b) as-spun on SiO2 c) annealed on OTS/SiO2 and d) annealed on SiO2. Crystallite nucleation occurs on initial deposition directly on or near the polymer/substrate interface (a,b). Growth occurs on annealing (c,d).

The XRD characterization is in agreement with electrical characterization and

fitting parameters obtains using the ME model (Table 3-2). First, unannealed films of

PQT-12 on untreated SiO2 barely show any field effect, in agreement with the poor

crystallinity of these films. The lower percent crystallinity in PQT-12 on bare SiO2

compared to OTS/SiO2 implies the presence of a larger fraction of disordered material on

the untreated dielectric. Furthermore, we do not see a resolution limited peak, implying

that if perfectly oriented crystallites are present at the dielectric interface, their intensity is

low and overwhelmed by diffraction from the bulk film.

A larger fraction of misoriented crystallites at the interface, compared to the other

films, would correspond to a larger Eb. An additional reason for the decrease in µ0 may be

related to the assumptions of the ME model. The ME model assumes that the DOS is

continuous throughout the film, which is dependent on there being enough crystallites in

the film to form continuous paths from source to drain. In PQT-12 on bare SiO2, the

crystallinity (and therefore crystallite density) is lower than on OTS/SiO2. We speculate

that on the untreated dielectric the crystallite density is below the percolation limit. Thus,

in order to cross the TFT channel charge must always travel through amorphous regions.

As a result the apparent crystalline mobility in the ME model is an average value between

the true crystalline mobility and the mobility in the amorphous regions.

65

3.4.3 Understanding Crystallization Behavior Using Pole Figures

The sketch in Figure 3-6 summarizes the crystallization behavior that we have

been able to elucidate using X-ray diffraction and pole figures of as-spun and annealed

PQT thin films. The OTS surface treatment increases polymer mobility at the interface

and encourages the formation of perfectly oriented crystallites. On annealing, the

crystallites grow through the thickness of the film and slightly rearrange, improving their

orientation with respect to the substrate and each other. On untreated SiO2 the crystallites

also grow on annealing, but we were not able to collect evidence for crystallite

rearrangement. Extracted crystallite domain size is on the order of film thickness,

implying that the majority of the crystallites present grow through the thickness of the

film. To gain further insight into crystallite growth during annealing, we measured pole

figures of crystallites in transient states. Shown in Figure 3-8 are pole figures of PQT-12

on OTS/SiO2 corresponding to three different states of crystallization. The first pole

figure (dashed light blue) was collected from a film that was quenched from the isotropic

melt state. The second pole figure (blue) was collected from the same film following a 1

minute anneal. The third pole figure (black) was collected from a film that was annealed

for 30 min (the same data that was discussed in previous sections). Relative DoC

calculated from the pole figures is given in Table 3-6.

Figure 3-8 Pole figures of PQT-12 after quenching, after a 1 minute anneal, and after a full anneal.

66

Table 3-6: Relative degree of crystallinity values calculated from pole figures shown in Figure 3-8.

The quenched film has a very low degree of crystallinity (1% as crystalline as the

PQT-12 film fully annealed on OTS/SiO2), but retains some characteristics of the

annealed films, including a population of perfectly oriented crystallites at the interface,

giving rise to a resolution-limited peak. Interestingly, this film has a surprisingly high

mobility of 0.02 cm2/Vs[28], much higher than the as-spun PQT-12 film on OTS/SiO2

(0.004 cm2/Vs). This supports the hypothesis that the remaining crystallites exist at the

dielectric interface. After a brief anneal of 1 min, the overall film crystallinity is still low,

but the shape of the pole figure indicates that the crystallites present are well oriented

with respect to the substrate. In fact, the shapes of the reannealed and annealed pole

figures are identical, indicating identical crystalline orientation distributions. We expect

that on further annealing, the pole figure of the original annealed PQT-12 film on OTS

would be reproduced. While the quenched film appears to have a broader crystallite

orientation distribution, this is probably an artifact of the low intensity, which is close to

the background noise level. The evolution of film microstructure is shown in Figure 3-9.

This series of pole figures tells us that crystallites at the substrate interface can

survive the isotropic melt temperature and serve as seed crystals for subsequent epitaxial

crystal growth during annealing. Such seed crystals have been observed before for thin

films of polyethelene[97], and provided the ability to clone single crystals through

thermal cycling. As discussed previously, we found that annealing an as-spun PQT film

involves crystallite growth in addition to crystallite rearrangement. We believe that the

difference arises due to the starting microstructure of the films. While the relative

crystallinity in an as-spun film (40%) is higher than the quenched film (1%), the structure

in the as-spun film is kinetically limited. Thus, in an as-spun film, the energy provided on

annealing allows for subtle rearrangements (improvement in texture) and subsequent

PQT-12 on OTS/SiO2 r. DoC

Quenched 0.01 Annealed, 1 min 0.13

Annealed 1.00

67

through-thickness crystal growth (increase in crystallinity). In contrast, the interface layer

of crystallites retained in a quenched film has already experienced an annealing process

and the crystallites are pinned at the interface; further annealing results only in one-

dimensional crystal growth without a change in texture.

Figure 3-9. Microstructure of PQT thin films a) Directly after quenching from the isotropic melt. b) After a short anneal. Arrows indicate crystallite growth. c) After a full anneal.

3.5 Conclusions

Charge transport in semicrystalline polymeric semiconductors is governed by many

different phenomena at different levels. At the mesoscale, charge transport is a property

of the microstructure of the polymeric thin film: it is governed by the complex interplay

between crystallites and disordered regions in the film. The microstructure of the film and

hence the mobility of charge carriers depends strongly on how the semiconducting film is

processed.

In this work, the microstructure of the crystallites in PQT-12 thin films was

thoroughly characterized by synchrotron-based X-ray diffraction. On annealing the as-

spun films, crystalline structure improved for films on SiO2 and OTS/SiO2. For PQT-12

films on OTS, specular diffraction indicated considerable growth and ordering in the out-

of-plane direction of the crystallites, while high resolution GIXD showed no significant

change in the in-plane ordering or coherence length. Pole figures were able to provide

68

details of the texture and the degree of crystallinity. The annealed film on OTS/SiO2 was

the most crystalline in the dataset. Moreover, both the as-spun and annealed films on

OTS/SiO2 showed a resolution limited peak attributed to crystallites at the dielectric

interface, which was not visible for the annealed film on SiO2. The OTS monolayer

promotes film crystallization and the formation of perfectly oriented crystallites.

Microstructural findings were in line with TFT electrical performance and electronic

structure parameters extracted using the ME model. Importantly, the complete pole

figures emphasized the importance of perfectly oriented crystallites at the dielectric

interface for charge transport. By characterizing crystallinity in a film that was quenched

from the melt and the same film after a 1 minute anneal, we found that after an initial

annealing step, these interface-nucleated crystallites are very stable. They can act as seed

crystals to reproduce the crystallite orientation distribution of an annealed bulk film.

Since this is the interface important for charge transport, mobilities remained high

through the quenching, reannealing and annealing process, despite the drastic changes in

crystallinity in the bulk of the film. This property may be exploited in a useful way

during device processing.

X-ray diffraction is a very powerful characterization tool that has proven to be

successful in enhancing our understanding of charge transport mechanisms in polymeric

semiconductors. In order to successfully correlate carrier mobility with the crystalline

microstructure of the film, one must take care to probe all of the diffracting material:

specular and in-plane diffraction patterns must be analyzed in combination with pole

figures. X-ray diffraction, however, is limited to characterization of the crystalline

regions. Therefore, when transport is limited by the disordered regions, other techniques

must be employed to gain a complete understanding of the structure-property

relationships in this class of materials.

69

4 Microstructure of P3HT Thin Films as a Function of

Thickness: Characterizing Texture and Degree of

Crystallinity

In this chapter, we use X-ray diffraction to investigate the microstructure of confined

films of high molecular weight P3HT. Pole figures show that the crystallinity of P3HT

increases as the film thickness increases. Simultaneously, the texture improves as the film

thickness decreases. An OTS surface treatment has a local effect of enhancing

crystallization of perfectly oriented crystallites. We find that texture drastically improves

when the film thickness is equal to the crystalline coherence length. Results here clearly

illustrate that the microstructural details influences by film thickness should be

considered when attempting to understand and optimize film microstructure for charge

transport.

4.1 Introduction

It is well accepted that the electronic performance of semicrystalline polymers is

highly dependent on the film microstructure. Since charge transport in a thin film

transistor takes place at the buried dielectric interface, this is a particularly important

region of the film to understand[35]. It has been qualitatively demonstrated that the this

interface is different from the bulk, but details of the substrate’s influence, including how

far into the bulk of the film the influence of the substrate persists, are not well

understood[71, 98]. Moreover, crystallization behavior of long polymer molecules has

been a topic of intense research for years[99-103] but knowledge of crystallization

behavior specific to semicrystalline polythiophenes is far more limited.

Here, we discuss the effect of film confinement on the microstructure of P3HT

films cast onto OTS-treated and untreated silicon oxide substrates. We address this topic

by measuring pole figures for films of varying thicknesses.[70] We pay particular

attention to how the degree of crystallinity and texture changes as a function of film

70

thickness. The relationship between microstructure and thickness is expected to become

interesting as the film thickness reaches the crystallite size[104]. Pole figures can be used

to quantify the relative degree of crystallinity in a film as well as film texture. Moreover,

by collecting information from perfectly oriented crystallites, pole figures serve as an

interface sensitive characterization tool. By varying film thickness, we are able

investigate the influence of the substrate on the local and bulk microstructure.

Information extracted from static films of different thicknesses can provide insight into

crystallization kinetics. While this work is motivated by implications on charge transport

in thin film transistors, and therefore focused on the structure at the buried interface,

studying microstructure of thicker films (>100 nm), is directly relevant for other devices

such as diodes[105] and solar cells[31].


Films were deposited by spinning dilute solutions of P3HT (MW 158 kDa) in 1,2-

dicholorobenzene. Thicknesses were controlled by varying solution concentration: 2.0,

1.0, 0.5, 0.25, 0.125 and 0.0625 wt %. Spinning recipe was kept constant for all samples

and consisted of a prolonged “gel” time of approximately 60 s (defined as the period

between solution deposition and initiation of spin coating), spin acceleration of

approximately 50 rpm, and two-minute spin, at 1250 rpm. Films were visibly dry after

the conclusion of spinning. Film thicknesses for the thicker films were measured using

atomic force microscopy in tapping mode, while thicknesses of the thinner samples were

extracted from X-ray reflectivity measurements

4.3 Results

Specular diffraction patterns for the P3HT films of various thicknesses are shown

in Figure 4-1. Film thicknesses range from approximately 6 nm to just over 100 nm.

Regardless of surface treatment, peak intensity decreases as thickness decreases. Three

orders of the (h00) peak are visible for nearly all of the films on OTS-treated silicon, and

for the thicker films on untreated silicon. There is an increase in the FWHM of the (100)

71

peak as the thickness is decreased, as listed in the Table 4-1 and Table 4-1. For the

thinnest films, on both OTS-treated silicon and untreated silicon, there is a slight shift in

peak position to lower q for the (100) peaks as indicated in the table. A similar shift is

seen in the position of the (200) peaks (data not shown). This indicates slight expansion

in the lattice spacing as the thickness becomes increasingly thin.

Figure 4-1. P3HT films of varying thickness on a) OTS-treated SiO2 and b) on untreated SiO2. Darker colors represent thicker films. Film thickness ranges from approximately 6-100 nm. For thinner films, scans shown are slightly off-specular, to avoid film thickness interference effects.

Sample Thickness (nm) q(100) (Å-1) d(100) (Å) FWHM (Å-1)

105 0.388 16.5 0.050 40 0.388 16.6 0.050

15.4 0.385 16.2 0.050 9.1 0.385 16.5 0.065 4.1 0.373 17.8 0.098

Table 4-1. Values extracted from the specular diffraction patterns of P3HT films on OTS-treated SiO2 as a function of thickness: (100) peak position, corresponding lattice spacing, and the FWHM of the (100) peak.

Sample Thickness (nm) q(100) (Å-1) d(100) (Å) FWHM (Å-1)

107 0.381 16.5 0.050 43 0.377 16.6 0.050 18 0.387 16.2 0.050 12 0.382 16.5 0.065 8.7 0.377 16.6 0.098 7.6 0.354 17.8 0.077

Table 4-2. Values extracted from the specular diffraction patterns of P3HT films on SiO2 as a function of thickness, including (100) peak position, corresponding lattice spacing, and the FWHM of the (100) peak.

72

The crystalline coherence length can be estimated from the specular diffraction

patterns using an approach based on a Williamson-Hall Plot and the Scherrer equation

[94] as discussed in Chapter 3. Values extracted using this approach are shown in Figure

4-2 below. The line of best fit was achieved when the FWHM values were plotted against

h2, indicating that the majority of peak broadening arises due to paracrystalline

disorder[3]. For film thickness greater than 20 nm, the domain size in the direction

normal to the substrate can be estimated as approximately 10 nm for films on OTS, and

approximately 12 nm for films on oxide, as shown in the plots below. Domain size

decreases significantly when films are less than 10 nm thick. A coherence length near 8-

10 nm for film thicknesses greater than 20 nm was confirmed using a more rigorous

approach based on the Warren-Averbach Fourier transform peak shape analysis technique,

discussed elsewhere[106]. These values for domain size agree with reports in the

literature (based on atomic force microscopy) [54, 79, 107].

Figure 4-2. Estimated values of domain size in the direction perpendicular to the plane of the substrate, extracted by fitting FWHM values of the (h00) peaks and plotting against h2.

As stated earlier, specular diffraction patterns probe only a small slice of

reciprocal space and miss details of the texture. In disordered semcrystalline polymer

films, diffraction from slightly misoriented crystallites can be significant. In order to look

at a larger slice of reciprocal space, grazing incidence diffraction patterns were collected

with a two-dimensional image plate, shown in Figure 4-3. In the series of 2D-GIXD

images, it is clear that the arcing associated with the family of (h00) peaks along the

vertical axes increases with increasing film thickness. This implies a qualitative

improvement in fiber texture with decreasing film thickness. The arcing is also visible in

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73

the Local Specular images shown on the right of the 2D-GIXD images. The trend seems

to be similar for both P3HT films on SiO2 and OTS/SiO2.

Figure 4-3. 2D-GIXD images with corresponding (100) local specular curves collected with a two-dimensional image plate detector for P3HT films of different thicknesses, spun on OTS-treated silicon oxide (a-f) and untreated silicon oxide (g-l). Film (a) was not used for quantitative calculations.

To be more quantitative, the 2D-GIXD images, 2D local specular images, and

high resolution rocking curves were combined to form pole figures of the (100) Bragg

reflection, as discussed in Chapter 2.

74

Figure 4-4. Pole figures of P3HT films of different thicknesses on a) OTS-treated SiO2 and b) untreated SiO2. Pole figures are normalized with respect to thickness. The pole figure with the maximum intensity at χ=0 for each dataset is shown in black.

The pole figures shown in Figure 4-4 are normalized for thickness. For all samples, the

intensity is highest at χ=0° and quickly decreases as the value of χ increases. This

confirms the texture typical of spin cast P3HT films and other semicrystalline

polythiophenes: the majority of crystallites are oriented with the (100) repeat direction

lying out of the plane of the substrate. Regardless of surface treatment, the pole figures

reveal a similar trend in texture as a function of thickness. As thickness is decreased, the

width of the pole figures also decreases, indicating an improvement in film texture.

The central resolution-limited peak is present in almost all of the pole figures. As

mentioned previously, it has been shown that this intensity arises from perfectly oriented

crystallites nucleated at the very flat dielectric interface[71]. It is interesting to note here

that in general, the thinner the film, the higher the intensity at χ=0° after thickness

normalization. However, the maximum intensity occurs at a thickness of approximately 9

nm, when film thickness is the same as crystallite size, regardless of surface treatment

(the pole figures with the maximum intensity at χ=0° are highlighted in black Figure 4-4).

We can use the pole figures to quantitatively compare the degree of crystallinity

between films of different thicknesses for the two different surface treatments. The

relative degree of crystallinity calculated from the complete pole figures for P3HT films

on OTS-treated SiO2 and P3HT films on untreated SiO2 are shown in Figure 4-5.

75

Regardless of surface treatment, the crystallinity decreases with increasing thickness. It

should be emphasized that the values listed here are not absolute values of crystallinity

but instead values corresponding to the relative degree of crystallinity. The most

crystalline film in each dataset is given a value of 1, against which other films are

compared. The thinnest film is approximately 20% as crystalline as the thickest film. The

crystallinity of the thickest film is expected to range between 20-50%, but this value is

not well known. [77, 108]

Figure 4-5. a) Relative degree of crystallinity of P3HT films as a function of thickness. Values were extracted from the integrated intensity of the pole figures. b) Integrated intensity from bulk of the pole figure, attributed to crystallinity throughout the bulk of the film, normalized with respect to film thickness. c) Integrated intensity from resolution-limited peak, attributed to crystallinity at the interface, normalized with respect to thickness of the interface layer (~ 9 nm). Intensities shown in (b) and (c) are therefore comparable.

In addition to providing a means to measure the degree of crystallinity, pole figures

simultaneously capture details of the texture (perfectly textured versus the remainder). It

is interesting to break up the pole figures according to texture, and monitor the

contributions to crystallinity. We attribute the resolution-limited intensity at χ=0° to

crystallites nucleated at the dielectric interface. The rest of the pole figure intensity, with

the slowly varying intensity at nonzero values of χ, is attributed to crystallites nucleated

elsewhere (such as off existing crystallites) and is herein referred to as the “bulk”

76

crystallinity. While not perfectly oriented, crystallites contributing to the bulk

crystallinity are still textured, as illustrated by the majority of the intensity within a few

degrees of χ=0°. The contribution of this intensity is shown in Figure 4-5(b). This

intensity is normalized for thickness, since we believe these crystallites exist throughout

the film thickness.

This intensity shown in Figure 4-5(c) is normalized by the expected thickness of

the interface layer size (~9 nm), or if thickness for the films less than 10 nm. For P3HT

on oxide or OTS, the intensity attributed to perfectly oriented crystallites is independent

of thickness, except when film thickness is equal to the characteristic crystalline

coherence length, near 9 nm Figure 4-5(c)). At this point there is a drastic increase in

intensity, or an increase in the population of crystallites that are perfectly oriented with

respect to the substrate. It should also be noted that the intensity is higher for films on

OTS-treated SiO2 compared to films on untreated silicon oxide, regardless of thickness.

4.4 Discussion

Pole figures were used to investigate the influence of interfaces on the

microstructure of P3HT films and the overall change in P3HT microstructure as a

function of thickness. With these pole figures we were able to confirm a decrease in

degree of crystallinity with decreasing thickness. In this work, pole figures of P3HT films

on oxide and OTS indicate an increase in texture, or edge-on orientation of the

crystallites, as a function of thickness. This trend has been documented previously for

P3HT [52, 54, 98, 109], but in a qualitative fashion. Joshi et. al. and Porzio have used

incident angle depth profiling[52, 54, 110], which is illustrative of overall trends, but the

intensity cannot be quantified and the resolution is poor.

We are able to monitor the intensity arising due to the presence of crystallites that

exist at the dielectric interface. The near perfect orientation of these crystallites, with

their (h00) reciprocal lattice vector oriented parallel to the surface normal, allows us to

extract information about the buried interface of the films. Resolution-limited peaks have

77

been observed before, most notably in the pioneering work by Kline, et al.[71] However,

these studies were limited to intensity collected from crystallites oriented within only a

few degrees of the surface normal. While the authors were able to be quantitative with the

diffracted intensity, without complete pole figures, they could not relate the information

to overall film texture, or the relative degree of crystallinity of the film. A layer of

increased order at the dielectric interface in P3HT is confirmed by Joshi, et al[110] and

Porzio, et al[52]. However, it should be noted that an ordering influence of the buried

interface is not always observed, and increased order at the free surface has been

observed in other systems [111, 112].

By isolating the intensity of the resolution-limited peak, we can study the effect of

substrate choice on the interface crystallinity. Regardless of thickness, P3HT films on

OTS have more perfectly oriented crystallites than films on oxide, as seen in Figure 4-

5(b), implying that OTS promotes the nucleation of strongly textured crystallites, as

suggested previously in Chapter 3[71]. This relationship is illustrated by directly

comparing pole figures of the thickest film on OTS with the thickest film on oxide, both

near 100 nm Figure 4-6(a). As argued in Chapter 3, the effect of OTS could be due to

enhanced polymer mobility on the layer of flexible alkyl chains compared to polymer

mobility on an untreated SiO2 surface. Enhanced polymer mobility will allow for a more

energetically favorable microstructure. It is also worth noting the OTS monolayers in

these studies were deposited via submersion in dilute solution. While contact angle on the

OTS-treated silicon with water was measured to be between 90° and 100°, assuring

complete coverage of OTS on silicon, the monolayers are most likely significantly

disordered. Molecular dynamics simulations have been used to show that disordered

monolayers, compared to crystalline monolayers of OTS, provide a smoother energetic

landscape, promoting the formation of perfectly edge-on crystallites[96]. In Figure 4-

6(a), the resolution-limited peak is barely visible for the film on SiO2 and this could be

interpreted as a complete absence of perfectly oriented crystallites. However, by plotting

the thinnest film on oxide (approximately 6 nm) and the thickness film on oxide

(approximately 100 nm) on the same scale, not normalized for thickness, we can see that

the perfectly oriented peak is most likely not absent, but simply overwhelmed by the bulk

crystallites, Figure 4-6(b).

78

Figure 4-6. a) Pole figures for a thick film on SiO2 (green) and a thick film on OTS/ SiO2 (blue). Films are approximately 100 nm thick and normalized for thickness. The main difference in the pole figures is the absence of the resolution-limited peak for the film on SiO2. b) Unnormalized pole figures for a thick film (ca. 100 nm) and a thin film (ca. 7 nm) on oxide.

Perhaps the most interesting and important result obtained from the complete pole

figures is the thickness dependence of the perfectly oriented crystallites at the dielectric

interface. For film thicknesses larger than 15 nm, the relative crystallinity at the interface

remains constant, regardless of surface treatment. When film thickness is on the order of

crystallite size as extracted from specular diffraction patterns (ca. 9-10 nm), there is a

drastic increase in the population of perfectly oriented crystallites. This confirms that the

interface microstructure is localized and does not propagate into the bulk of the film. The

increase in intensity when film thickness is near 10 nm is a new observation that we do

not yet understand. One explanation is that when films are this thin, every nucleation

event is within 10 nm of the interface. Bulk nucleation is nearly completely thwarted, and

the orientation of every crystallite is strongly influenced by the very flat polymer-

dielectric interface. This is supported by the drastic decrease in integrated intensity from

the bulk crystallites for a film thickness less than 10 nm. Furthermore, it has been found

for systems of polystyrene[113] and poly(ethylene terephthalate)[111] that the free

surface of a polymer film (within 5 to 10 nm) possesses enhanced liquid-like mobility.

When film thickness is larger than the coherence length of a crystallite, this would serve

to enhance formation of the misoriented crystallites. However, when film thickness is less

than or equal to 10 nm, the mobile layer enhances crystallization of the perfectly oriented,

79

dielectric nucleated crystallites. When the film is decreased further, below the

characteristic crystalline coherence length, overall film ordering suffers from decreased

polymer mobility, and crystallites are not able to form completely, thus intensity drops.

The bulk crystallinity is monitored by the calculating the integrated intensity of

the slowly varying portion of the pole figure. These crystallites, while still textured, are

not perfectly oriented with respect to the substrate. The relationship of the bulk

crystallinity and film thickness is shown in Figure 4-5(b). This intensity is normalized

with respect to thickness, and thus is comparable to the intensity of the resolution-limited

peak, Figure 4-5(c). From this figure, we can see that the bulk crystallinity of the film is

not affected by the surface treatment on the substrate. As discussed in the previous

paragraph, the substrate has a strong influence on the perfectly oriented crystallites, but

this influence is quickly lost (within ~10 nm of the dielectric interface.) One may expect

that the relationship between crystallinity and thickness can be attributed to the increasing

influence of the disordered layers that exist at the film interfaces. The higher the volume

fraction of the film attributed to the disordered interface layer, the lower the overall

crystallinity. However, complete pole figures tell us that for semicrystalline P3HT, this

cannot be the operating mechanism. P3HT has an interface layer of extremely well

ordered crystallites. In variance with flexible polymer chains, modeling experiments have

shown that rather than becoming more disordered at a flat interface, the stiff portions of a

semiflexible chain prefer to orient parallel to the surface of the substrate, enhancing

polymer chain orientation and crystalline order.[114]

The crystallization behavior and therefore resulting trends in crystallinity of the

P3HT films in this study are affected by a number of important factors that must be

considered when attempting to understand the crystallization mechanisms. Films are spun

from warm solutions to minimize any crystallization prior to deposition. Since P3HT

films are semicrystalline directly following deposition, it can be assumed that the

majority of crystallization occurs during the spin coating process. During spin coating,

films are subjected to shear forces from initial fluid flow and subsequent solvent

evaporation[100]. Spin speed was kept constant during the deposition of the different

films, so it’s influence can be ignored[84]. In addition, ordering and crystallization may

occur after initial deposition, while a significant amount of residual solvent molecules

80

remain in the film. The films were not subjected to any subsequent processing such as

thermal or solvent annealing, so rearrangement after complete solvent evaporation is

minimal.

In the system at hand, the presence of a surface greatly decreases the energy for

crystallization and heterogeneous nucleation should therefore be much more favorable

than homogeneous nucleation[95] In our explanation, we rule out contributions from

crystallites nucleated far away from any surface. We believe that the perfectly ordered

crystallites are nucleated on solution deposition at the dielectric interface, and continue to

grow during the time period between solution deposition and initialization of spinning

(which was purposely kept to approximately 60 s.) When the substrate begins to spin, the

film is subjected to extensive fluid flow and complex shear forces. Bulk crystallites

nucleate off of the existing perfectly oriented crystallites or the available dielectric

interface, and the film crystallizes from solution in a one-dimension fashion (from the

dielectric interface towards the free surface.) These bulk crystallites are not coherent with

the crystallites nucleated off the dielectric and are therefore not perfectly oriented. Due to

the one-dimensional growth, they are still influenced by the dielectric interface and

therefore maintain the fiber texture. The degree of fiber texture is expected to decrease

with increasing distance from the dielectric interface, as new crystallites are nucleated off

increasingly misoriented existing crystallites. As the film thickness decreases, the bulk

crystallites constitute a smaller percentage of the film, and overall film texture is

improved. This relationship is qualitatively illustrated in for a thin (ca. 10 nm) and thick

(greater than 10 nm) film in Figure 4-7. It is interesting to note that similar one-

dimensional growth has been modeled using the Avrami equation[115] of polymers[95]

and PCBM thin films[116]. The increasing degree of crystallinity seen in the thicker

films may be due to the difference in solution viscosity. As stated in the experimental

section, film thickness was controlled by solution concentration. The thicker films are

spun from more concentrated solutions, which are in turn more viscous. Fluid flow in

viscous solutions is decreased, and the solvent evaporation is slower. Therefore, the

thicker the film, the longer the drying time, with a longer drying time leading to a more

crystalline film[77, 117].

81

Figure 4-7. Sketch of film microstructure for a thick film (top) and a thin film (bottom.) Perfectly oriented crystallites nucleated off of the substrate interface are drawn with black lines. There are more perfectly oriented crystallites when film thickness is approximately equal to the crystallite coherence length. oriented crystallites when film thickness is approximately equal to the crystallite coherence length.

It is interesting to compare the P3HT dataset with annealed films of PBTTT of

varying thickness due to the fundamental differences in their crystallite formation.

PBTTT has a stiffer chain backbone, which, combined with a lower density of side chains,

allows for the interdigitation of alkyl chains and a crystallite that is fully coherent in

three-dimensions[27, 89]. Pole figures of PBTTT films of different thicknesses, spun on

OTS-treated SiO2 are shown in Figure 4-8. Unlike P3HT, the texture of three thinnest

films of PBTTT (50, 23 and 14 nm) is not thickness dependent. The unannealed P3HT

films discussed previously have a crystalline domain size perpendicular to the substrate

of approximately 10 nm, while crystallines domains of annealed PBTTT, have a

coherence length of approximately 30 nm or larger and often extend throughout the

thickness of the film (as extracted from X-ray diffraction measurements and suggested by

the terraces visible via AFM[72, 89]). This through-thickness coherence in PBTTT

results in less of a difference in average orientation between crystallinity at the interface

and crystallinity far from the interface, so texture is not thickness-dependent. When the

film thickness is significantly larger, ca. 80 nm, crystallites no longer grow through the

film thickness and the pole figure width increases.

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Figure 4-8. One-dimensional complete pole figures of PBTTT thin films of varying thicknesses: 85 nm (dark red), 50 nm (medium dark red), 23 nm (red), 14 nm (dashed pink). The three thinnest films have the same crystallite orientation distribution, with intensity varying only close to χ=0°.

4.4.1 Implications for Charge Transport

It is well documented that TFT mobility of semicrystalline polythiophenes is

higher when the surface is treated with an OTS monolayer[32, 48, 107]. It is has been

shown that the origin of the effect is at least partially structural[118], which is supported

here by substrate surface dependence on interface crystallinity. We find that an OTS

monolayer encourages the nucleation of highly oriented crystallites at the substrate

interface, which is the important interface for charge transport in a bottom-gate thin film

transistor. It has also been suggested that the OTS monolayer serves to mask interfacial

traps present on a silicon oxide interface and prevent contamination[119].

Pole figures have shown that with decreasing thickness, there is an improvement

in crystalline texture and a decrease in degree of crystallinity for P3HT thin films. Most

importantly for charge transport, which takes place at the dielectric interface, we have

found that when the thickness is on the order of crystallite coherence length, there is a

drastic increase in the fraction of perfectly oriented crystals. A number of studies have

documented the trend of better crystalline texture resulting in better electronic

performance[48, 84, 120, 121]. Better overall texture of crystallites in the film

corresponds to better alignment between neighboring crystallites and therefore easier

transport across grain boundaries. Unfortunately, to date we have not been able to

measure a trend in TFT mobility for the P3HT films investigated here.

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Studies of thickness dependence of P3HT TFT mobility have been published. Joshi

et. al measured no dependence on thickness for P3HT TFT mobility, and attributed the

lack of a relationship to the significant portion of amorphous material present at the

interface, dominating transport[122]. Jia et al. found an increasing transistor channel

current with increasing thickness due the presence of significant bulk conductance[123].

Sandberg et al. found a lower mobility for monolayer films and attributed the poorer

performance both to strain introduced by the proximity of the interfaces to the entirety of

the film and decreased dimensionality of a percolating network for charge transport[124].

The lattice expansion we observed for the thinnest P3HT films may indicate a similar

presence of interfacial strain. However, since our thin films still consist of multiple

molecular layers, the argument of reduced percolation should not apply, as there can be

transport across the layers.

4.5 Conclusion

Understanding the influence of interfaces on the microstructure of semicrystalline

polythiophenes and the dependence on microstructure on film thickness is important if

we want to understand variables that affect film formation, and in turn how to control

charge transport. In the work presented here, we chose to study P3HT films of varying

thicknesses deposited on OTS-treated SiO2 and untreated SiO2 using pole figures. Pole

figures provide a quantitative means to characterize film texture and degree of

crystallinty. Moreover, since perfectly oriented crystallites are known to nucleate off the

dielectric interface, pole figures provide information about polymer morphology at the

buried interface.

We confirmed a decreasing overall crystallinity with decreasing thickness for

P3HT films on OTS and P3HT films on oxide. The fiber texture of the P3HT films

improves when thickness is decreased, for both OTS and oxide. By looking closely at the

components contributing to crystallinity, perfectly oriented crystallites and bulk

crystallites, we were able to better understand the influence of thickness and the influence

of the substrate. The bulk crystallinity dramatically decreases with decreasing thickness

and is nearly independent of surface chemistry. On the other hand, the interface

84

crystallinity is nearly thickness independent, except when thickness approaches crystallite

coherence length. At this point, there is a dramatic increase in the population of perfectly

oriented crystallites, both on OTS and oxide. In general, there are more perfectly oriented

crystallites on a surface treated with OTS than an untreated surface.

The discovery of a thickness that encourages perfectly oriented crystallites may be

important for optimizing charge transport in thin film transistors. General trends in

texture and degree of crystallinity will help to establish rigorous processing guidelines for

the fabrication of organic electronic devices optimized for good charge transport

properties. Interestingly, the change in crystallinity of P3HT is so pronounced with

changing thickness that it renders a strict definition of “standard P3HT” degree of

crystallinity meaningless. Degree of crystallinity is not a static material property, but a

very volatile characteristic that changes drastically as a function of the film processing

parameters.

Future work in this area will involve using pole figures to investigate different

materials systems with fundamentally different crystallization habits such as PBTTT, and

introducing variables such as thermal and solvent annealing. In-situ diffraction

characterization of P3HT and PBTTT films during thermal and solvent annealing

processes would provide more insight into crystallization behavior. The effect of surface

treatment can also be studied in more detail. The quality of the OTS layer, which is easily

modified by using different deposition techniques, has been shown to result in different

microstructures of pentacene[125]. Pole figures could provide more insight into the

interface and bulk crystallization habits.

85



Crystallinity

In this chapter we use pole figures to characterize the effect of processing solvent,

molecular weight (MW) and regioregularity (RR) on the microstructure of P3HT films.

We then correlate microstructure characterization with electronic parameters extracted

using the ME model. By combining these two characterization tools we help clarify the

complex relationship between processing, microstructure, and charge transport in

semicrystalline P3HT films. We find that when processing conditions are optimized, even

a polymer with moderate RR of 84% can have a high room temperature mobility (0.07

cm2/Vs).

5.1 Introduction

The relatively high mobility of P3HT relies on the ability to control both the micro-

and macromolecular parameters of the polymer. Since the early P3HT field effect

transistors,[23] large improvements have been made in device performance through

modifications of polymer chemistry and film processing. With the synthesis of

regioregular P3HT[24, 25] came a large improvement in device mobility from 1 × 10-5 to

1 × 10-2 cm2V-1s-1 [126]. It was concluded that this improvement was due to an increased

order of the polymer chains, in terms of film crystallinity and quality of the individual

crystallites.[126] Sirringhaus et al. later observed that in some cases, the degree of

regioregularity of the P3HT polymer may influence crystallite orientation, or film texture.

Grazing incidence diffraction revealed that films of P3HT with a moderate degree of

regioregularity of 81% and a MW of 28 kDa spun from a dilute solution in chloroform

86

comprise crystallites arranged with their π-stacking direction oriented normal to the plane

of the film in a “face-on” fashion. When the P3HT regioregularity was increased to 96%

and the Mw decreased to 11 kDa, the crystallites were oriented in an “edge-on” fashion,

with the π-stacking direction oriented parallel to the plane of the substrate. The observed

structural difference was correlated with a difference in mobility of three orders of

magnitude: 0.1 cm2V-1s-1 (96% RR P3HT) compared to 1 × 10-4 cm2V-1s-1 (84% RR

P3HT). The better electronic performance in the high RR P3HT was attributed to the fact

that “edge-on” crystallites have two directions of charge transport in the plane of the

substrate, which is favorable for charge delocalization and transport across the thin film

transistor channel. Charge modulation spectroscopy measurements corroborated this

hypothesis.

Further studies regarding the dependence of the mobility on other factors such as

molecular weight (MW), methods of film formation (spin-coating, drop-casting, dip-

coating)[127] and choice of solvent have focused almost exclusively on polymers with a

high RR (>94%). Kline et al.[80, 128] attributed the increase in mobility with molecular

weight (ranging from 2 × 10-6 cm2 V-1s-1 for P3HT with a molecular weight of 4.0 kDa to

9× 10-2 cm2 V-1s-1 for P3HT with a molecular weight of 36.5 kDa) to details of the

semicrystalline grain structure. The higher molecular weight polymers formed films with

a nodular structure, with the possibility of tie-molecules connecting adjacent crystalline

regions, while the low molecular weight films formed isolated crystallites with abrupt

grain boundaries, hindering charge transport. Interestingly, XRD results in this study

suggested a higher degree of crystallinity for films of the low molecular weight polymer.

Zen et al.[82] explained the increase of mobility (from µsat of 10-7 cm2V-1s-1 to 10-3 cm2V-

1s-1) with molecular weight (from 2.2 kDa to 19 kDa) using a different model of film

microstructure, in which P3HT films consist of sparse crystalline regions embedded in an

amorphous matrix. As a result, charge transport was controlled not by grain boundaries,

but by the details of polymer chains existing in the amorphous regions. Within these

regions, spectroscopy experiments revealed that lower molecular weight films have a

more twisted backbone conformation than higher molecular weight films, resulting in

intrachain transport barriers. Zhang et. al investigated a similar range of molecular weight

and found that ultrathin dropcast P3HT films had a nanofibrillar structure. There was a

87

linear increase in both the mobility (ca. 5.0 × 10-5 to 0.01 cm2V-1s-1) and the width of the

fibrils (10 nm to 30 nm) when the molecular weight was increased from 2.4 to 7.5 kDa.

When the MW was further increased past 10 kDa, both the fibril width and mobility

remained constant[83]. Chang et al.[79] evaluated the evolution of the mobility over a

large range of molecular weight and processing solvents and observed a similar increase

of mobility with increasing molecular weight as long as drying time was sufficiently long

(achieved by using a high boiling point solvent and/or drop cast deposition). The authors

distinguish three major regimes in the dependence of mobility on MW: 1) low MW (<22

kDa) and low mobility, in which short polymer chains back-fold into the lamellae,

leaving chain ends behind and creating highly defective crystals; 2) intermediate MW

(25-52 kDa) and sharp increase in mobility, in which the polymer chains are long enough

to create crystalline lamellae free of any chain end defects, with a dramatic increase of

the charge carrier mobility as a consequence; and 3) high MW (>52 kDa) and slow rise in

mobility by one order of magnitude, where once the chain length exceeds the natural

contour length (defined as 30-60 nm for well-ordered lamellar stacked P3HT crystallites)

the polymer molecules fold into neighboring lamellae, thereby increasing the

connectivity between crystalline regions and improving intercrystalline charge transport.

Despite subtle differences, the phenomenological observations by Chang and Kline are

very similar and also supported by Verilhac et al.[127] The authors concluded that both

the improvement of crystalline quality and the increase of the inter-crystalline charge

transport pathways are important factors in the realization of P3HT thin films with high

carrier mobilities. However, the relative contribution of both effects in the overall charge

carrier mobility is unclear.

Several studies investigate the effect of solvent choice on the mobility of

semicrystalline P3HT[77, 117]. In these studies, the intensity of the Bragg reflection

corresponding to the lamellae (100) was used to estimate the degree of crystallinity of

P3HT films and the crystallite orientation distribution. While an increase in the apparent

degree of crystallinity, seen in higher boiling point solvents and longer drying times, was

generally correlated with an increase in carrier mobility, misorientation of crystallites in

the film can also play an important role in measured intensity and charge transport

88

properties. Thus, both the film crystallinity and crystalline texture (orientation) should be

characterized in order to fully understand microstructure and charge transport.

Here, we continue the ongoing investigation on the importance of molecular weight

and solvent choice on the microstructure and electronic structure of P3HT and

reintroduce as a variable the degree of regioregularity. We are only interested in

semicrystalline P3HT films with a high mobility (µ > 0.01 cm2V-1s-1) and therefore limit

our studies to P3HT samples with relatively high degrees of regioregularity (>80%) and

high molecular weights (>60 kDa) and cast films from solutions made with relatively

high boiling point solvents. Thin film transistors are used to collect temperature-

dependent transfer (Ids vs. Vg) characteristics, which are then fit to a two-dimensional

Mobility Edge (ME) model, as discussed in Chapter 2.[28] The use of this model allows

us to deconvolute the relative contribution of traps and mobile states in determining the

effective mobility measured in the TFT. We characterize the crystalline texture and

degree of crystallinity using pole figures, also discussed in Chapter 2[68]. By correlating

this quantitative structural characterization method with TFT modeling, we are able to

understand how processing and polymer structure affect the performance of the P3HT

thin films. In particular, we address the effect of regioregularity, molecular weight and

choice of processing solvent on film structure (including the degree of crystallinity,

crystallite orientation distribution and crystalline quality) and the subsequent effects of

these microstructural details on the carrier trap density, trap energetic distribution and

room temperature field effect mobility.


P3HT with different degrees of regioregularity and molecular weights were

synthesized by Merck Chemicals and their characteristics are summarized in Table 5-1.

Solutions of P3HT were prepared at concentrations of approximately 10 mg/mL in either

1,2-dichlorobenzene (DCB) or 1,2,4 trichlorobenzene (TCB) and stirred overnight at a

temperature of 90°C in order to attain good solubilization of all polymers. Films were

cast from heated solutions. Thin film transistors were fabricated as discussed previously.

89

All substrates were treated with an OTS monolayer. Room temperature and variable

temperature mobilities were measured as discussed previously. Polymer films for XRD

were deposited on OTS treated silicon.

5.3 Results

Room temperature mobilities are shown in Table 5-1, alongside parameters

extracted using the mobility edge model (µ0, Eb, Nt). P3HT with a moderate

regioregularity of 84% has a surprisingly high charge carrier mobility of 0.02 cm2V-1s-1

when cast from DCB. Moreover, when the same material is spun from TCB, the mobility

increases to 0.07 cm2V-1s-1. Two P3HT samples have a high mobility of 0.1 cm2V-1s-1: 97%

RR, 158 kDa P3HT, spun from TCB and 97% RR, 64 kDa P3HT, spun from DCB.

RR (%) MW (kDa) Solvent Thickness (nm) µTFT(cm2V-1s-1) µ0(cm2V-1s-1) Eb(meV) Nt(cm-2)

97 158 DCB 15 0.02 0.8 39 3.6 ×1013 97 158 TCB 18 0.10 1.4 26 2.2 ×1013 97 64 DCB 25 0.10 4.1 26 6.3 ×1013 97 64 TCB 15 84 130 DCB 10 0.01 0.4 38 4.1 ×1013 84 130 TCB 12 0.07

Table 5-1. Details of the polymer samples and the parameters extracted by modeling transfer characteristics at temperatures ranging from 80-100 K using the ME model.

Pole figures were normalized with respect to film thickness, with the assumption

that film thickness did not vary enough to cause large changes to the degree of

crystallinity or texture. However, as discussed in Chapter 4, even the small changes in

thickness as seen here (from ~10 to 25 nm) may have a significant effect on film

microstructure. We proceed with data analysis and discussion with this in mind. The

relative degrees of crystallinity of all films are shown in Figure 5-1. Also shown in

Figure 5-1 are the integrated intensities, on an arbitrary scale, of the resolution-limited

peaks.

90

Error bars account for error arising from error in the measurement of film

thickness and the definition of the noise floor at high χ. As mentioned previously, this

resolution peak is attributed to crystallites nucleated off the dielectric interface, the

interface important for charge transport. Pole figures are shown in Figure 5-2, 5-3 and 5-

4. The qualitative shape and quantitative relative degree of crystallinity of the pole

figures are discussed in the following sections. To simplify the discussion of the different

samples, we will use the following notation: P3HT [RR, MW, Solvent].

Figure 5-1. Pole figure data analysis. The solid circles represent the relative degree of crystallinity of the films in this study. (The relative degree of crystallinity for P3HT, 84% RR, DCB was associated with considerable error due to low film thickness and low intensity, and therefore this datum is not shown.) The open circles represent integrated intensity of the resolution-limited peak, on an arbitrary scale. All data are normalized with respect to thickness.

5.3.1 Effect of Solvent

To process the P3HT films, we made solutions using two different solvents with

similar molecular structure but different boiling points: 1,2 dichlorobenzene (DCB) with

a boiling point of 178°C and 1,2,4 trichlorobenzene (TCB) with a boiling point of 214 °C.

The higher boiling point translates to a longer film drying time during the spinning

process (approximately 1 minute for films spun from DCB at 1000 rpm, compared to a

drying time in excess of two minutes for TCB under the same spinning conditions.) In

comparison to other common solvents used to process P3HT, the boiling points of both

DCB and TCB are relatively high; allowing us to focus our study on high mobility

P3HT[77, 117]. Figure 5-2(a-c) compare pole figures of films spun from DCB and TCB

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for three sets of materials: 97% RR, 64 kDa P3HT, 97% RR, 158 kDa P3HT, and 84%

RR, 130 kDa P3HT. For the polymers with the higher molecular weights, the use of a

higher boiling point solvent results in films with a tighter crystallite orientation

distribution, as illustrated by the width of the pole figures. For the intermediate molecular

weight P3HT, the change in texture is not as pronounced. Interestingly, processing high

molecular weight P3HT with TCB does not increase the crystallinity, despite the longer

time for polymer rearrangement (see Figure 5-1). Nor is there a significant difference in

the degree of crystallinity for the intermediate molecular weight P3HT between the two

solvents. The intensity from the resolution-limited peak is always larger for a film

processed with DCB compared to its TCB counterpart.

The room temperature mobilities of P3HT [97, 158, DCB] and P3HT [97, 158,

TCB] were 0.02 and 0.1 cm2V-1s-1, respectively. The mobility of the P3HT [97, 64, TCB]

was 0.1 cm2V-1s-1. The room temperature mobilities measured for the P3HT [84, 130,

DCB] and P3HT [84, 130, TCB] were 0.01 and 0.07 cm2V-1s-1, respectively. According

to parameters extracted from the ME model, the total number of traps in the P3HT [97,

158, DCB] film decreases slightly (from 3.6 × 1013 cm-2 to 2.2 × 1013 cm-2) when

processed with TCB (P3HT [97, 158, TCB]). At the same time, the crystalline mobility,

µ0, increases by a factor of two (0.8 to 1.4 cm2V-1s-1 for P3HT [97, 158, DCB] and P3HT

[97, 158, TCB], respectively). Hence, the remainder of the five-fold increase in mobility

can be attributed to the decrease in the energetic distribution of the traps from

approximately 39 meV for the samples spun from DCB to 26 meV for the samples spun

from TCB, which allows the Fermi level (EF) to move closer to the ME.

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92

Figure 5-2. Pole figures of P3HT films illustrating the effect of solvent on crystalline texture. a) 84% RR, 130 kDa P3HT spun from DCB (dark green) and TCB (light green). b) 97% RR, 158 kDa, spun from DCB (dark blue) and TCB (light blue). c) 97% RR, 64 kDa, spun from DCB (black) and TCB (grey).

5.3.2 Effect of Molecular Weight

We investigated high regioregularity (97%) P3HT films with different molecular

weights: an intermediate molecular weight of 64 kDa and a high molecular weight of 158

kDa. As mentioned previously, higher molecular weights were chosen to assure that this

study remains focused on P3HT films with high room temperature mobilities. Figure 5-3

shows pole figures obtained from P3HT films with different molecular weights (a) P3HT

[97, 64, DCB] and P3HT [97, 158, DCB] and (b) P3HT [97, 64, TCB] and P3HT [97,

158, TCB]. When spun from the lower boiling point solvent (DCB), the pole figure of the

64 kDa P3HT reveals a better textured film, with less intensity at higher angles of χ.

When the same materials are processed with the higher boiling point solvent (TCB), the

pole figures are nearly identical. The P3HT, 64 kDa has slightly more intensity at low

angles of χ. There is no trend in degree of crystallinity with molecular weight. However,

the integrated intensity from the resolution-limited peak is always higher for the 64 kDa

film compared to its 158 kDa counterpart. P3HT [97, 64, DCB] has the highest

resolution-limited intensity in this study by a factor of two.

The room temperature mobility of P3HT [97, 64, DCB] was measured as 0.1

cm2V-1s-1. When molecular weight is increased to 158 kDa and processing solvent kept

the same, the mobility drops to 0.02 cm2V-1s-1 (P3HT [97,158, TCB]). According to

parameters extracted from the ME model, P3HT [97, 64, DCB] has a higher density of

trap states (6.3 × 1013 cm-2) compared to P3HT with P3HT [97, 158, DCB], (3.6 × 1013

cm-2). In contrast, P3HT with a MW of 64 kDa, spun from DCB has a tighter energetic

distribution of traps (Eb = 26 meV), than P3HT with a molecular weight of 158 kDa spun

from DCB (Eb = 39 meV). Between these two polymers, there is a five-fold difference in

room-temperature field-effect mobility. Interestingly, P3HT [97, 64, DCB] has a higher

mobility, despite the larger number of traps.

93

Figure 5-3. Pole figures of P3HT illustrating the effect of molecular weight. a) P3HT, 97% RR, 158 kDa (blue) and 65 kDa (black) spun from DCB. b) P3HT, 97% RR 158 kDa (light blue) and 65 kDa (grey) spun from TCB.

5.3.3 Effect of Regioregularity

To investigate the effect of regioregularity, we included in this study P3HT with a

RR of 84%, which is only a moderate decrease in RR, again assuring that this study is

focused on semicrystalline P3HT films with high room temperature mobilities. The 84%

RR has a molecular weight of 130 kDa, which we consider comparable to the higher

molecular weight of 158 kDa for the high regioregularity P3HT. Figure 5-4 shows pole

figures of the two high molecular weight P3HT films spun from the two solvents: P3HT

[84, 130, TCB], P3HT [97, 158, TCB] and P3HT [84, 130, DCB], P3HT [97, 158, DCB].

When spun from the lower boiling point solvent, the higher RR results in a slightly

tighter orientation distribution, although both of these pole figures are broader than the

others in this study. It should be noted that the P3HT [84, 130, DCB] film spun very thin,

corresponding to low unnormalized intensity. Thus, the absolute intensity of this pole

figure after thickness normalization is expected to be associated with significant error

(due to both the low intensity and large error in measured thickness). For this reason, the

relative degree of crystallinity is associated with significant error (>30%) and is not

stated. Qualitative observations about the shape, however, will hold. When spun from

TCB, the pole figures of the 84% RR and 97% RR P3HT films are nearly identical. The

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94

pole figures reveal no trend in degree of crystallinity as a function of regioregularity.

However, the two 84% RR samples have the two lowest integrated intensities of the

resolution-limited peak in this study. Between the two low regioregularity samples, the

integrated intensity was slightly higher for the film spun from DCB. The difference in

film thickness between the low regioregularity samples and the higher regioregularity

samples makes accurate comparison difficult.

Figure 5-4. Pole figures of P3HT films illustrating the effect of RR. a) P3HT, high molecular weight spun from TCB, 97% RR (light blue) and 84% RR (light green.) b) P3HT, high molecular weight spun from DCB, 97% RR (dark blue) and 84% RR (dark green.)

5.4 Discussion

5.4.1 Effect of Processing Conditions and Polymer Chemistry on Film

Microstructure

Pole figures allow us to quantitatively compare details of the texture between

films, as well as the relative degree of crystallinity. As stated previously, all films in this

study are well textured, with the (100) peak lying out of plane of the substrate. The pole

figure intensities are all the same order of magnitude. Moreover, the resolution-limited

peak is present in all pole figures, indicating that a population of perfectly oriented

crystallites exists at the buried dielectric interface.

Subtle changes in the shapes and intensities of the pole figures provide insight

into the effect of processing solvent, MW and RR. In all cases, the use of a higher boiling

point solvent (TCB instead of DCB), results in narrower pole figure, with less intensity at

the wings. We attribute this effect to the difference in drying time associated with the two

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95

solvents and not the slight change in polarity. (At 25°C, 1,2, DCB has a dielectric

constant of 9.93, compared to 2.24 for 1,2,4 TCB.) Indeed, the results shown can be

reproduced by spinning films from the same solvent (DCB) and increasing the drying

time by controlling the spinning atmosphere. However, from our experience, processing

with a higher boiling point solvent, while keeping the spinning atmosphere inert proved

to be a more reproducible way to control drying time. In all cases, the longer drying time

results in a tighter crystallite orientation distribution (i.e. fewer misoriented crystallites).

The effect is more pronounced in the films of the higher molecular weight P3HT.

Interestingly, the use of TCB does not have a consistent effect on the degree of

crystallinity. P3HT [97%, HMW, TCB] has a significantly lower degree of crystallinity

than P3HT [97%, HMW, DCB], for reasons we do not completely understand. This may

be due to variation in the film thickness. One explanation may be that the longer drying

time results in higher quality crystallites, at the expense of crystalline volume: defects

such as chain ends are expelled into the grain boundary regions between crystallites,

instead of being incorporated in the crystal lattice. The result of this process is a film with

more ordered crystallites compared to a film spun from DCB, but with a larger volume of

disordered material and therefore a smaller overall degree of crystallinity. Regardless, the

pole figures are direct evidence that processing with TCB lets P3HT crystallites rearrange

with an overall more perfect orientation with respect to the substrate. The improvement in

crystallite-substrate alignment results in better alignment between neighboring crystallites.

Evidence that TCB can improve the crystallite alignment, but does not improve the

degree of crystallinity, provides insight into the kinetics of P3HT film formation: short-

range polymer crystallization happens fast (<1 minute), and is near completion after the

drying time allotted by DCB; the additional time associated with processing a film with

TCB, however, allows for subsequent arrangement of the polymer crystallites with

respect to the substrate and each other. Further evidence that the longer crystallization

time allows for better order within the crystallites is provided by parameters extracted

using the ME model, as discussed in the next section.

In regards to crystallite orientation distribution, a lower molecular weight has a

similar effect to raising the boiling point of the solvent. For films spun from DCB, the

polymer with the lower molecular weight has a tighter crystallite orientation distribution.

96

When spun from TCB, the effect is smaller but the intermediate molecular weight does

result in more crystallites oriented at low χ. We expect that the shorter chain lengths in

the 64 kDa P3HT (ca. 150 nm compared to ca. 350 nm for 158 kDa) are more mobile,

allowing a film with more oriented crystallites to be formed in less time. There is not a

clear trend in overall crystallinity between the molecular weight, as seen in previous

studies. According to these studies, we are operating in a regime where crystallinity

plateaus as a function of molecular weight. However, there are trends in the resolution-

limited intensity. P3HT [97, 64, DCB] has the highest integrated intensity from the

resolution-limited peak, or the most perfectly oriented crystallites of all of the sample.

When RR is kept high, the 64 kDa film has more resolution-limited intensity compared to

the 158 kDa film, whether processed from DCB or TCB.

When the films are processed from DCB, an increase in the degree of RR

tightened the crystallite orientation distribution. When processed from TCB, the effect

was no longer present. In this work, P3HT with a moderate RR formed highly textured

films. P3HT [84, 130, TCB] was as crystalline as the P3HT [97, 158, TCB]. However,

whether processing with DCB or TCB, increasing the RR resulted in an increase in the

intensity attributed to the perfectly oriented crystallites.

5.4.2 The ME Model and the Extracted Parameters.

The parameters extracted from the ME model characterize details of the electronic

structure of the film (the total number of traps Nt and their energetic distribution, Eb) and

the crystalline mobility, µ0. By correlating the parameters extracted by fitting device

characteristics with the ME model and the structural information obtained using X-ray

diffraction, we attempt to gain insight into the nature of the traps.

P3HT [97, 64, DCB] has the largest trap density in this study (6.3 × 1013 cm-2),

nearly twice the density found in the higher molecular weight films. This film is one of

the most crystalline films in the study: the lack of correlation between trap density and

the amount of crystalline material implies that charge traps are not uniformly distributed

within the completely amorphous regions. The apparent dependence of total trap density

on molecular weight (with lower molecular-weight polymer films having a larger trap

97

density) suggests that charge traps may be associated with chain end defects. It has been

shown previously that the chemical nature of P3HT chain ends alone does not affect

transport.[80] Therefore we hypothesize that the traps are rather structural defects

associated with the chain ends, such as packing irregularities within crystallites or at the

crystalline boundaries. In an ordered region of the polymer, an incorporated chain end

would be analogous to an edge dislocation in a crystal. The incorporation of chain ends

into crystallites has been previously observed in high MW P3HT.[117, 127, 129]

Two films in this study, P3HT [97, 158, DCB] and P3HT [84, 130, DCB] have

larger values of Eb (ca. 40 meV), compared to the rest of the films. In these films, charge

carrier traps span a larger energetic range. Interestingly, pole figures of these materials

have a larger width, with more intensity in the wings. When the high MW P3HT is spun

from TCB or when the MW is decreased, the energetic distribution of the traps decreases

to below 30 meV, as does the pole figure intensity at large χ. In both cases, we find that a

tighter crystallite orientation distribution is associated with a smaller energetic spread of

charge carrier traps. We expect a tighter crystallite orientation distribution will allow for

polymer chains to bridge with less bends or twists, and therefore present a smoother

energetic landscape to the charge carrier.

In addition to a lower Eb, there is an increase in the crystalline mobility of the

high RR P3HT when either the MW is decreased or the boiling point of the solvent is

increased. The crystalline mobility that we extract using the ME model is directly related

to the π-orbital overlap within the crystallites, and therefore trends in µ0 are good

indicators of trends in crystalline quality. Thus, when organization time is limited (as

with a lower boiling-point solvent) and the molecular weight is high, P3HT forms films

with more misoriented crystallites of relatively poorer crystalline quality, resulting in

deeper traps and low crystalline mobility. However, when allowed sufficient time for

reorganization (which is achieved by increasing the boiling point of the solvent or

decreasing the molecular weight), the decrease in Eb and increase in µ0 results in a 5-fold

increase in room temperature mobility attributed to an increase in texture and crystalline

quality.

The highest mobility in this study of 0.1 cm2V-1s-1 was obtained from two films:

P3HT [97, 64, DCB] and P3HT [97, 158, TCB]. The films have a similar texture and a

98

similar value of Eb, but differ in crystalline mobility and total number of traps, illustrating

the complex interplay between details of the film microstructure and polymer chemistry.

It appears that a high mobility can be achieved with only a moderately high crystalline

mobility, as long as the film has a smaller trap density, as seen in the case of P3HT [97,

158, TCB]. Alternatively, a high mobility can also be achieved in a film with a large

density of traps, as long as the film has higher quality crystals, with a high crystalline

mobility. As has been shown, however, one cannot hope to improve crystalline mobility

indefinitely by using progressively shorter polymers, as eventually intergrain transport is

suppressed and the mobility drops[80, 130, 131]. A higher crystalline mobility can be

obtained by molecular design. For example, when side-chain density is low enough to

allow for interdigitation, the crystalline mobility increases, leading to the exceptionally

high room-temperature mobility of PBTTT[47].

Results throughout this study indicate that polymers with different characteristics

can be used to produce high quality films, as long as processing conditions, in particular

time allowed for crystallization, are optimized. Along this vein, we remark that the

moderate RR (84%) P3HT, with a high molecular weight, drop cast from TCB has a

surprisingly high room temperature mobility of 0.07 cm2V-1s-1. Unfortunately, the device

behavior did not follow the ME model. For the 84% RR P3HT spun from DCB, the lower

RR results in a lower crystalline mobility (and a lower crystalline quality), but the

polymer maintains the same number of charge traps with the same energetic distribution

as the higher RR sample. Even without ME parameters for the 84% RR film spun from

TCB, we can conclude that when deposited from a high boiling point solvent, a high MW

P3HT with a moderate RR forms a semicrystalline, fiber-textured film that is

microstructurally very similar to a film spun from high RR P3HT, resulting in a high

room temperature mobility.

5.5 Conclusions

We studied the structural and electronic properties of P3HT thin films while

varying the molecular weight, regioregularity and solvent used for deposition, with an

99

aim of better understanding the microstructural effects on charge transport. We used pole

figures to characterize the microstructure of P3HT films, paying particular attention to

the orientation distribution of the crystalline portions of the films, as well as the relative

degrees of crystallinity. The mobility edge (ME) model allowed us to extract parameters

related to crystalline mobility, energetic trap distribution and total number of charge

carrier traps. Results indicate that the effect of a higher boiling-point solvent improves

crystallite orientation, but does not involve significant growth of crystallinites during

drying, as indicated by a decreasing degree of crystallinty. The lower boiling point

solvent results in more perfectly oriented crystallites. More intermediate molecular

weight P3HT forms higher quality crystals, with a tighter orientation distribution and

more perfectly oriented crystallites.

In general, trap density is not correlated with crystallinity, but there is an apparent

relationship between the energetic distribution of traps and the crystallite orientation

distribution, with a better texture corresponding to a tighter distribution of traps.

Importantly, we found that there are a number of ways to achieve a high mobility in

P3HT films. A large number of charge carrier traps can be balanced by a large crystalline

mobility. Likewise, a moderate crystalline mobility can be balanced by a small number of

charge carrier traps. Most surprisingly, P3HT films with a moderate RR can still have a

high mobility if enough time is allowed for crystallization and crystallite rearrangement.

All three scenarios result in room temperature mobility near 0.1 cm2V-1s-1. The ability to

relax the criteria of near perfect regioregularity allows for increased flexibility for organic

chemists in materials design and more options when choosing polymeric semiconductors

for devices. For instance, low RR P3HT might be more environmentally stable owing to

its increased ionization potential. Understanding the relative importance of different

aspects of the design and fabrication processes contributes to the fundamental science of

film formation and charge transport, and can also be very important in the design of

polymer semiconductor based electronics.

100



Crystallinity

Thus far, the work presented in this thesis has involved the use of X-ray diffraction

to characterize the quality and orientation of crystalline regions within a film. We have

shown that in order to understand charge transport in semicrystalline polythiophenes, it is

important to characterize crystalline quality, crystalline texture and the overall degree of

crystallinity. In regard to charge transport, it is generally assumed that charges traverse

distances within a crystal easily, while the grain boundaries offer some barrier to

transport. However, the mechanism of charge transport across grain boundaries, as well

as their overall importance in film performance, is not well understood. It has been

hypothesized that in films of high molecular weight P3HT, long polymer chains connect

neighboring grains with low misorientation, resulting in percolative charge transport.[41]

In this chapter, we use P3HT films with engineered microstructures to experimentally

verify the role of percolative charge transport in high molecular weight semicrystalline

polythiophenes.

6.1 Introduction

As organic semiconductors approach commercialization, there is a need to better

understand the relationship between charge transport and microstructure, in particular to

identify the inherent bottlenecks to charge transport. A common method to investigate the

effect of specific microstructural features on electrical performance is to introduce known

or controlled defects. In noncrystalline, glassy polymeric semiconductors, such as poly-

101

9,9’-dioctyl-fluorene-co-bithiophene (F8T2), studies have focused on interchain versus

intrachain transport, where anisotropic transport was realized by aligning the backbone of

the polymer molecule on rubbed polyimide.[22] The ratio of the mobility measured along

the backbone alignment direction to that measured across was approximately four. In

semicrystalline and polycrystalline materials, charge transport is most likely dominated

by grain boundary effects, although the exact mechanism is not understood. In small-

molecule organic semiconductors, isolated grain boundaries in neighboring crystals have

been probed with field-effect transport measurements.[132, 133] It was shown that most

performance benchmarks, including activation energy, threshold voltage, and field-effect

mobility, were indeed affected by the presence of the grain boundaries. Similar studies of

grain boundaries in semicrystalline polymers are more difficult: crystalline domains form

on length scales too small to allow devices to be made across a single, isolated grain

boundary. Furthermore, measurement interpretation is complicated by the unknown chain

orientation between neighboring grains. As a result, considerably less is known about the

transport bottlenecks in these semicrystalline films. It has been previously suggested that

polymer chains bridging neighboring grains can provide electrical pathways across grain

boundaries and reduce the associated transport barrier.[134, 135] If bridging polymer

chains are forced to undergo sharp bends or twists as they transition from one crystalline

domain to the other, there may be localized states associated with the breaks in

conjugation, slowing transport. On the other hand, polymer chains that bridge crystallites

with minimal distortion will provide a fast intergrain transport path.[41] Fastest transport,

thus, should arise when there is small misorientation between neighboring grains with

respect to polymer-chain axis: boundaries between grains with low misorientation angles

should have lower transport barriers than boundaries between grains with high

misorientation angles. There have been previous attempts to understand charge transport

using oriented films of polythiophenes.[136-139] However, these attempts have suffered

from either poor, nonuniform surface coverage, anisotropy in dielectric roughness, poor

crystallinity, or limitations due to crystallite orientation. In this work, the problem of

transport across grain boundaries in semicrystalline polymers is addressed quantitatively

using a directional crystallization technique to make anisotropic thin films of poly(3-

hexylthiophene) (P3HT) with controlled grain-boundary placement and orientation. Part I

102

presents the directional crystallization and characterization of the film microstructure and

morphology. Part II discusses transistor fabrication and charge transport measurements.

6.2 Fabrication and Characterization of Engineered

Microstructures

In order to study grain boundaries, we chose to create thin films with a controlled

grain boundary structure. The film crystallization process was adapted from a previous

publication[108, 140], but the work presented here is the first to thoroughly investigate

the microstructure using X-ray diffraction and to study electrical behavior using TFTs.

6.2.1 Experimental Details

P3HT (Mn 24 500 g mol-1, polydispersity index (PDI) of 2.63, >97% RR) was

used as received from Merck Chemicals. TCB was used as received from Sigma Aldrich.

The directional crystallization process was adapted from descriptions in publications by

Brinkmann and De Rosa [108, 140]. The process is illustrated in Figure 6-1. In this

technique, 1,3,5-trichlorobenzene (TCB), a small-molecule solvent that is solid at room

temperature, acts first as a solvent and second as a substrate for polymer epitaxy. Thin

films of P3HT were cast onto glass or silicon substrates from a 3 wt% solution of the

polymer in chlorobenzene. The films in this work were made on OTS-treated treated SiO2,

but the process can be replicated on bare SiO2 and quartz substrates. The polymer films

are placed on the warm end of a thermal gradient on aluminum stage (~75°C.)

Approximately 10 mg of TCB was added on top of the warm polymer film. As the TCB

melts and begins to dissolve the film, a glass coverslip is placed on top, creating a

sandwich structure. At this point the film is completely dissolved. The P3HT:TCB

solution is then undercooled by moving the substrate to a region of the temperature

gradient that is below the melting temperature of TCB (Tm = 55°C). The crystallization of

TCB is induced by cooling one edge of the substrate via contact with a metallic heat sink

103

(such as tweezers or a spatula). The TCB crystals form in long needles in the plane of the

substrate, and initially expel the P3HT into the remaining solution. As the TCB continues

to solidify, the concentration of P3HT in solution increases. At its solubility limit, the

P3HT precipitates out of solution, with the TCB needles acting as a substrate for chain

alignment. The TCB is subsequently sublimed away in a vacuum oven at a slightly

elevated temperature (~40°C) over several hours. On removal from the vacuum oven, the

coverslip can be easily lifted from substrate. On both the substrate and the coverslip, an

anisotropic thin film of P3HT is left behind, with oriented areas limited by the size of the

substrate.

Figure 6-1. Illustration of the steps involved in the directional crystallization of P3HT in the presence of 1,3,5 trichlorobenzene. Starting from left: A P3HT TCB solution, nucleation of TCB needles at cool end of the substrate, continued growth of TCB needles, solidification of P3HT on TCB needles, and an anisotropic P3HT film following sublimation of TCB.

6.2.2 Results

An optical micrograph of an anisotropic, directionally crystallized film of P3HT is

shown below in Figure 6-2. The film changes color upon rotation between crossed

polarized filters, indicating birefringence arising from the macroscopic alignment of

polymer chains.

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104

Figure 6-2. a) Optical microscopy image under crossed polarizers of a directionally crystallized P3HT film on glass. Axis of polarization is aligned as indicated by the arrows.

Atomic force microscopy (AFM) images are shown in Figure 6-3. There are large

areas of continuous films, with an anisotropic fiber structure. Fiber widths range from 10

to 250 nm, with an average width of 130 nm. Overall film thickness is approximately 80

nm. Across the fibers, films have a root mean square (RMS) roughness of 16.4 nm; along

the fibers, films are smoother, with an RMS roughness of 1.4 nm. The contrast in phase

mode reveals periodicity along the fibers of approximately 40 nm. Two types of features

are visible in the phase mode image. Box 1 highlights equiaxed features existing within

fibers, while Box 2 highlights elongated features reminiscent of lamellar grains.

Figure 6-3. Tapping mode AFM images, topography mode (left) and phase mode (right.) Box 1 highlights an area of equiaxed features, and Box 2 highlights an area of elongated features.

!"#$%#

105

Two-dimensional grazing incidence X-ray diffraction (2D-GIXD) survey patterns

(Figure 6-4(a)) show a strong (010) π-π stacking peak along the vertical cut (nominally

qz axis). Weaker (h00) peaks are also visible along the vertical cut, but with considerable

arcing. Two orientations of high-resolution GIXD were performed, one with the substrate

oriented with the fibers perpendicular to the scattering vector, q, and one with the fibers

oriented parallel to q (Figure 6-4(c)). The two prominent peaks in these scans are

indexed as the (100) (alkyl stacking with a d-spacing of 16.52 Å), and the (010) (π-π

stacking with a d-spacing of 3.82 Å) peaks, indicating significant in-plane anisotropy. A

Φ-scan (in-plane azimuthal angle) of the (010) peak illustrates the degree of in-plane

orientation of the crystallites having the π-π stacking direction in the plane of the

substrate (Figure 6-4(d)). The azimuthal full-width at half-maximum (FWHM) of the

Bragg peak is approximately 10°, with a small isotropic component. A Φ -scan of the

(100) peak indicates that these crystallites show similar in-plane orientation. With the q

vector parallel to the fiber, a weak peak appears near q=1.52 Å-1, and has been indexed as

the c-axis repeat of the P3HT molecule.[135]

106

Figure 6-4. a) 2D-GIXD image of a directionally crystallized P3HT thin film illustrating the unique texture. b) Illustration of edge-on and face-on crystallites present in a directionally crystallized film. c) High resolution grazing incidence X-ray diffraction patterns, with the scattering vector oriented parallel (thin black) and perpendicular (thick black) to the fibers. d) ϕ-scan of the (010) peak.

6.2.3 Discussion

The optical micrograph shows an anisotropic film of P3HT, with fiber-like

features. The film is dark magenta when polarization is along the fibers, but near

colorless when polarization is oriented perpendicular to the fibers, indicating that the

majority of the chain backbones are oriented parallel to the fiber axis. This is also parallel

to the fast growth axis, or the c-axis of the TCB needles. It has been proposed[108] that

one reason directional crystallization of P3HT with TCB works so well is because a

lattice matching epitaxy that exists between the cP3HT and cTCB, as illustrated in Figure 6-

5. The similar dimensions along the polymer chain and the TCB needle combined with

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107

the unidirectional growth front influences chain orientation.

Figure 6-5. Illustration of relevant dimensions and stacking structure along the fast growth direction of TCB.

The AFM reveals a film with a unique, anisotropic film microstructure. The AFM

height mode image shows fiber like features, and the phase mode image, while not giving

direct information about crystalline structure, suggests the presence of an anisotropic

grain structure. We believe that the periodicity of the boundaries along the fibers arises

from stacked crystallites reminiscent of the lamellar crystallization seen in many

semicrystalline polymers[141-143].

X-ray scattering confirms the semicrystalline nature of the films, and reveals a

pronounced in-plane film texture. Typical spin-cast thin films of P3HT are highly

textured out of plane, with alkyl-stacking repeat direction lying out of the plane of the

substrate and the resulting (h00) peaks appearing along the qz axis, as discussed in

Chapter 1.[8] The π-π stacking direction lies preferentially in the plane of the substrate,

but with no in-plane texture. In striking contrast, in directionally crystallized thin films, a

significant amount of the crystallites have their π-π stacking direction oriented out of the

plane of the substrate. However, characteristic features of spin-cast-like texture, (h00)

peaks along the qz axis (close to vertical cut in Figure 6-4(a)), are still visible. There are

therefore two populations of crystallites in the directionally crystallized films, as shown

in Figure 6-4(b): those with the π-π stacking direction lying in the plane of the substrate

(‘‘edge-on’’ crystallites) and those with the π-π stacking direction lying out of the plane

of the substrate (‘‘face-on’’ crystallites). Without the structure factor of crystalline P3HT,

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108

we are not able to accurately determine the relative populations of each orientation.

However, since the (010) peaks are always very strong relative to the (100) peaks in the

2D-GIXD patterns of P3HT, we conclude that the majority of crystallites adopt a face-on

orientation.

We speculate that once a crystallite is formed in a fiber, the same crystallographic

orientation will be preserved for a certain distance along the fiber, until a perturbation in

the crystallization process induces a change in the polymer habit plane. AFM does not

allow us to assign an orientation to the crystallites observed. We believe, however, that

edge-on crystallites would have the π-π stacking direction in the plane of the film. The

overlapping π orbitals result in a driving force for crystallization in the plane of the

substrate, leading to more elongated structures (the stacked lamellar crystallites highlight

in Figure 6-3, Box 2), reminiscent of the fibrils often observed in polythiophene thin

films.[29, 135] Face-on crystallites, on the other hand, have the alkyl stacking direction

in the plane of the film, with a corresponding lower driving force for in-plane

crystallization, leading to the more equiaxed features in AFM image (Figure 6-3, Box 1).

The high-resolution GIXD experiments confirm the crystalline anisotropy in the

plane of the substrate. The intense (010) and (100) peaks that appear when probing the

direction perpendicular to the fiber, in the plane of the substrate, confirm the presence of

both edge-on and face-on grains, with the chain axes oriented parallel to the fiber axis.

The drastic intensity drop for these peaks when probing along the fiber direction indicates

that there are very few crystallites with the π-π or alkyl stacking directions along the fiber.

This texture is consistent with the appearance of a chain-backbone peak where the

scattering vector is parallel to the

fiber direction.

109

Figure 6-6. a) Microstructure of a directionally crystallized P3HT film. On the left, the yellow arrow indicates the direction of the fiber axis and the polymer chain axis. Edge-on crystallites are shown in blue and face-on crystallites are shown in grey. Different colors do not indicate different materials. On the right is a sketch of the film showing anisotropic grain structure. The black arrow indicates the direction of the long fiber axis and the polymer chain axis. b) Microstructure of a low-angle in-line grain boundary, with bridging polymer chains. c) Microstructure of a fiber-to-fiber grain boundary, with no bridging polymer chains.

A summary of the microstructure of directionally crystallized P3HT thin films is

shown in Figure 6-6. From the sketch it is clear that we have not eliminated grain

boundaries, but have controlled to a certain extent their type and where they are located

in the film. Along the fibers, we have low angle grain boundaries, where chains meet

end-to-end, Figure 6-6 (b). We know there will be some disorder present, but it is also

reasonable to assume that the presence of bridging polymer chains along these boundaries

is promoted, especially since this is also along the growth direction. We refer to these as

“in-line” grain boundaries. The bridging polymer chains in these grain boundaries should

increase electrical connectivity. Across the fibers, while we maintain the low angle

misorientation, neighboring grains are not oriented in a way that promotes bridging

110

polymer molecules, Figure 6-6(c). There are a number of possible variations in the grain

boundary microstructure, including edge-on to edge-on, face-on to face-on, and edge-on

to face-on. While two adjacent edge-on crystallites may have overlapping π-orbitals if

they are perfectly aligned, this is not likely. Any shift, tilt, or rotation of the crystallites

will greatly decrease the π–π overlap and create a barrier for charge transport.[11]

Adjacent face-on molecules already suffer from the low conductivity of the alkyl stacking

direction.[43] Additional disorder introduced at the edge of the fiber will further increase

the transport barrier. Adjacent fibers of different crystallite orientations, similarly, have

an insulating layer of alkyl chains between crystallites, with poor electrical connectivity.

Most likely, due to the one-dimensional growth of the fiber like features, the fiber-to-

fiber grain boundaries host extreme disorder and resemble amorphous material. In any of

the above possibilities, a polymer chain will not be able to cross the boundary without

substantial bending or twisting, thus decreasing the electrical connectivity of neighboring

grains in the direction perpendicular to the fibers.

6.3 Using the Engineered Microstructures to Investigate the

Role of Grain Boundaries

The anisotropic grain structure in these directionally crystallized films makes them

a unique characterization tool for charge-transport studies. By studying charge transport

along the fibers (with in-line grain boundaries) and across the fibers (with more

disordered grain boundaries), we can better understand the mechanism of charge

transport across grain boundaries in semicrystalline thin films. Specifically, we aim to

investigate whether the bridging polymer chains promote charge transport.

6.3.1 Experimental Details

To fabricate the thin film transistors, directionally crystallized films of P3HT

were grown on heavily doped silicon substrates (serving as the common gate electrode)

111

with a 200 nm dielectric, treated with OTS. Top-contact thin-film transistors were

prepared in one of two orientations: either with fibers parallel to the channel length

(parallel devices) or with fibers perpendicular to the channel length (perpendicular

devices). Channel lengths were kept long (>200 µm) to minimize effects of contact

resistance. Care was taken to measure devices only in areas of continuous film coverage,

as observed via optical microscopy. The device characteristics obey standard field-effect-

transistor equations.[144] Room temperature and variable temperature measurements

were taken as described in Chapter 2.

6.3.2 Results

In parallel devices, the long axis of the fiber is parallel to the channel length,

while in perpendicular devices, the long axis of the fiber is perpendicular to the channel

length. There is a reproducible room-temperature mobility anisotropy between the

parallel and perpendicular device orientations, as shown in Figure 6-7. The average

mobility of the parallel devices is 3 × 10-3 cm2V-1s-1, while the average mobility of the

perpendicular devices is 2 × 10-4 cm2V-1s-1. The best mobilities of each orientation differ

by a factor of 20. Note that measurements shown in Figure 6-7(b) were taken on

different devices and different days. Data are spaced along the x-axis for clarity (the

values are arbitrary.)

112

Figure 6-7. a) A parallel and a perpendicular TFT with directionally crystallized P3HT as the active layer. b) Room temperature mobilities for parallel devices and perpendicular devices. Data are space along the x-axis for clarity. c) Temperature dependent mobilities for parallel and perpendicular devices. d) Activation energies of parallel and perpendicular devices, as well as a spin cast film of P3HT.

The temperature dependence of the field-effect mobility was measured between

100 and 340 K (Figure 6-7). The mobility anisotropy is consistent down to low

temperatures. Assuming an Arrhenius-like relationship, where

€

µ∝e−EA

kT , values

corresponding to an activation energy EA for charge transport can be extracted, shown in

Figure 6-7. Also shown is the activation energy for a neat film of the same P3HT, spun

from DCB (µ = 0.1 cm2/Vs). While the mobility anisotropy is maintained down to low

temperatures, there is no anisotropy in activation energy within error. Errors bars for the

parallel devices are smaller than the markers.

6.3.3 Discussion

In a semicrystalline film with a mixture of phases, mobility follows the relation

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113

where µeff is the effective mobility extracted from the transfer characteristics, µg is the

mobility within the crystalline grains, µb is the mobility within the grain boundary

regions[41] (which may contain amorphous or disordered polymer), and Lg and Lb refer

to the total lengths of the grains and the grain boundaries, respectively. Such an

expression is derived from a model where the film is broken down into a series of

resistors. Mobility within the grains is most likely anisotropic due to the difference in

charge transport along the conjugation direction and across overlapping π-π orbitals.[41]

There may be some amount of disorder within the grains; however, it is safe to assume

that µg>>µb. In fine-grained materials such as the directionally crystallized films of P3HT,

the effective device mobility is dominated by the mobility in the grain boundaries, µb.

The devices in this work were fabricated such that in parallel devices, charge

transport occurs predominately along the fiber direction, whereas in perpendicular

devices, transport occurs across the fibers. In this way, we believe the mobility anisotropy

reflects the anisotropic transport barriers presented by the different types of grains

boundaries. With the grain structure presented in Figure 6-7 in mind, transport can be

described in the two devices as follows. In a parallel device, charges easily traverse the

low-angle “in-line” grain boundaries with more bridging polymer chains relatively easily.

In a perpendicular device, a charge traversing the channel has to cross every fiber-to-fiber

grain boundary. The device geometry will allow the charge to diffuse along a fiber,

across a series of relatively low transport barriers, but will not allow the charge to bypass

higher fiber-to-fiber grain boundaries. These fiber-to-fiber grain boundaries are the

boundaries that are most likely more disordered, presenting large barriers to transport,

and thus the low mobility. It is interesting to note that due to the shape of the crystallites,

in a parallel device there are on average three times as many boundaries in the current

direction as in a perpendicular device (25 boundaries µm-1 vs. 8 boundaries µm-1) as

estimated from AFM images. The higher average measured mobility of the parallel

devices in spite of the higher number of grain boundaries is further evidence that grain

boundaries along a fiber provide significantly easier charge transport when compared to

114

cross-fiber boundaries.

The similar activation energies for the two device orientations can be explained

by considering the common bottleneck to charge transport. Single fibers do not bridge the

entire sample, and thus fiber discontinuities and collisions, along with the other defects

mentioned in the previous paragraph, will render fiber-to-fiber hopping events necessary,

even in the parallel devices. As a result, charges in a parallel device will not be limited to

traversing only across grain boundaries along the fibers, but will be forced to cross fiber-

to-fiber grain boundaries. Both device orientations have the same transport bottleneck,

fiber-to-fiber transport, leading to nearly identical activation energies. The larger number

of fiber-to-fiber barriers in the perpendicular device lowers the mobility.

A Boltzmann-weighted hopping model was also used to characterize transport in

these films.[145] Electrical data presented here was used as input. The model supported

the hypothesis that it is the anisotropic grain structure, rather than the presence of defects

(such as fiber breaks and discontinuities) that cause the anisotropic mobility. The model

also confirmed that it is reasonable that fiber discontiuities account for the similar

temperature dependencies (i.e. activation energies) in mobility for both the parallel and

perpendicular devices.

This data presented here support the percolation model for charge transport in high

molecular weight semiconducting polymers. The presence of bridging polymer chains

existing across in-line, low angle grain boundaries (seen in Figure 6-7(a,b)) results in a

lower barrier to transport along the fibers, leading to a higher mobility for parallel

devices, compared to perpendicular devices. However, by limiting the dimensionality of

the percolating network so drastically, without the ability to completely eliminate defects,

we have in turn hindered charge transport compared to a neat film of spin cast P3HT. In a

spin cast P3HT film, there is a higher fraction of edge-on crystallites, which has proven

to be the better texture for charge transport[43], accompanied by a more complex

network of percolating pathways. In an in-plane isotropic film, a charge has so many

possible pathways that it can completely avoid very disordered grain boundaries. This is

further evidenced by the lower activation energy in comparison to directionally

crystallized films (Figure 6-7).

115

6.4 Conclusions

In this chapter, we have investigated the role of grain boundaries in semicrystalline

semiconducting polymer films. The films were made using a directional crystallization

technique in the presence of TCB, which acted first as a solvent and then as a substrate

for epitaxy. In order to understand the unique grain structure, the films were thoroughly

characterized using X-ray diffraction and atomic force microscopy. The polymer chain

axis was oriented parallel to the fiber axis, thus creating an anisotropic grain boundary

structure. Along the fibers, grains were likely connecting with bridging polymer chains.

Across the fibers, grain boundaries did not have the possibility of bridging polymer

chains. Measuring charge transport mobilities in two different directions allowed us to

study the charge-transport properties of grain boundaries between different orientations

of crystallites. Boundaries along the fiber provide a small barrier to charge transport

when compared to fiber-to-fiber grain boundaries. In these ‘‘in-line’’ grain boundaries, it

is likely that relatively straight polymer molecules provide an easy intergranular charge

transport path. We believe that fiber-to-fiber grain boundaries act as large transport

barriers because intergranular chains cannot exist without sharp bends or twists, thus

decreasing electrical connectivity. This work helps to solidify the hypothesis that charge

transport across low-angle grain boundaries is easier only in the direction parallel to the

polymer backbone, due to bridging molecules. Unlike their inorganic and small molecule

counterparts[17, 146], long chain polymer molecules have a way of making low-angle

grain boundaries relatively benign by incorporating bridging molecules with little bends

or twists that increase electrical connectivity betweens grains. Thus, charge transport in

semicrystalline polymers should not be thought of as completely two-dimensional, but

instead as transport through a series of one-dimensional pathways of low angle grains

with bridging polymer chains. The results presented here are the first experimental

evidence supporting the use of a percolation model for charge transport. Keeping this

transport mechanism in mind, optimization of polymer semiconductors’ microstructure in

electronic devices should therefore not focus solely on the elimination of grain

boundaries, but also include efforts to control grain boundary placement and relative

grain orientation.

116

7 Conclusions

Experiments presented in this thesis were designed with the common goal of

understanding how polymer chemistry and processing conditions affect microstructure

and how details of the microstructure, including structure at the grain boundaries, affect

electronic structure and charge transport.

The crystalline microstructure and electronic structure of disordered semicrystalline

films can be difficult to characterize accurately and quantitatively. The characterization

methods used in this work were discussed in Chapter 2. Importantly, we introduced for

the first time a method for measuring pole figures of weakly diffracting thin films. This

allowed for the quantitative characterization of crystalline texture and relative degree of

crystallinity. Previous estimates of the degree of crystallinity in the field of

semicrystalline polymers have been only qualitative. The ME model allowed us to

correlate details of the microstructure with field-effect mobility and details of the

electronic structure.

Chapter 3 discussed the use of X-ray diffraction to understand the effect of

processing conditions on the microstructure of PQT. Specifically, dielectric surface

treatment and post processing thermal treatments were explored. We found that an OTS

treatment results in a more crystalline film and encourages the formation of perfectly

oriented crystallites compared to bare SiO2. On annealing the as-spun film, crystalline

regions rearrange and grow. Pole figures were also used to show that interface crystallites

can survive above the melting temperature and act as seed crystals for subsequent crystal

growth. Furthermore, as long as the crystalline structure is maintained at the interface,

mobility can be high. In all cases, the presence of perfectly oriented crystallites correlated

with good charge transport in the bottom gate field-effect transistors, as previously

suggested.[147]

117

In Chapter 4, the crystalline texture and degree of crystallinity of P3HT films was

characterized as a function of thickness. We found that texture is strongly dependent on

thickness, as is film crystallinity. Thinner films have better texture, but a lower degree of

crystallinity. Pole figures allowed us to break the crystallinity into two components:

crystallinity existing at the interface, and crystallinity existing everywhere else. The

interface crystallinity is independent of thickness. This is in line with the apparent local

effect of the dielectric surface treatment in P3HT films: OTS promotes the formation of

perfectly oriented crystallites (as we saw with the PQT films) but the remaining

crystallinity is unaffected. P3HT has a small crystalline coherence length of

approximately 10 nm. When the crystallites grow through the film thickness, as in thin

films of PBTTT, texture is independent of thickness. Interestingly, when P3HT film

thickness was on the order of crystalline coherence length, there was a drastic increase in

the population of perfectly oriented crystallites, which we expect to have a beneficial

effect on charge transport in thin film transistors. The degree of crystallinity in the films

varied significantly as a function of thickness: the thinnest film (~7 nm) was less than 20%

as crystalline as the thickest film (100%). Such a drastic change as a function of such one

variable questions the validity of a strict degree of crystallinity of “standard P3HT”. The

property should instead be considered a volatile characteristic that is highly dependent on

processing parameters.

In Chapter 5, we investigated the effect of molecular weight, processing solvent and

regioregularity on the microstructure and electronic structure of P3HT, remaining

focused on high mobility P3HT. This study emphasized the complex relationship

between microstructure and charge transport. In the P3HT samples studies here, charge

transport corresponded to trends in film texture, particularly the presence of crystallites

with perfect texture, rather than the degree of crystallinity. Importantly, we found that

given enough time to crystallize (which can be a achieved with the use of a higher boiling

point solvent), even P3HT with a moderate regioregularity can form a film with a high

mobility near 0.1 cm2V-1s-1. With proper optimization of processing conditions, material

design constraints (such as near perfect regioregularity) can be relaxed.

Much of this thesis was spent characterizing crystalline regions of semicrystalline

films, but this is only part of the story. The structure of grain boundaries and mechanism

118

of charge transport across them must also be considered. We used P3HT films where we

controlled the type and placement of grain boundaries to experimentally validate the

hypothesis of percolative transport in high molecular weight polymer semiconducting

films.[28] Indeed, results confirmed the presence of bridging polymer chains across low

angle grain boundaries, providing experimental evidence for a charge transport model

proposed previously by Kline and coworkers, where longer chains form grain boundaries

with many bridging polymer chains, resulting in lower transport barriers.[135]

Importantly, unlike their inorganic and small molecule counterparts, polymer

semiconductors can offer a way to make grain boundaries more benign with the

incorporation of bridging polymer chains.

Understanding how characteristics of the polymer as well as details of the processing

conditions can affect the device performance is important for the optimization of

materials design and device fabrication. The work presented here outlines some

parameters that should be considered. Results show that molecular weight should be kept

relatively high in order to maintain grain boundaries with bridging polymer chains and

the constraint on RR can be relaxed if processing conditions are optimized. Furthermore,

film thickness is a critical parameter for controlling film texture and films should be

deposited in such a way to encourage the formation of perfectly oriented crystallites at

the dielectric interface. Importantly, the method for pole figure collection introduced in

this work allows for the first time the quantitative characterization of the effect of

processing conditions on a film’s degree of crystallinty. This method is general and will

allow for more thorough microstructural characterization of other organic electronic

materials.

Future work should be focused on translating the results presented here to newer

materials such as PBTTT and novel high mobility n-type polymers.[76, 148] A catalog of

high mobility p-type and n-type polymers will allow for the fabrication of

semiconducting polymer complementary logic circuits. In addition, before OTFTs can be

integrated into flexible electronics, high mobilities must be achieved using flexible

dielectrics instead of the rigid SiO2 substrates used here. Due to the surface roughness,

flexible dielectrics will discourage the formation of perfectly oriented crystallites. These

crystallites have proven important for charge transport in OTFTs. Engineers should

119

therefore find a way to encourage the formation of highly textured films on imperfect

surfaces, or compensate for their absence with higher structural quality within the

crystallites present. In addition to the development and commercialization of

macroelectronics such as displays and large area photovoltaics, effort should be devoted

to the realization of novel applications that take advantage of the unique characteristics of

semiconducting polymers. Devices that exploit the bandgap tunability, sensing

capabilities and biocompatibility of these materials could prove extremely useful, making

large impacts in the fields of smart packaging and bioelectronics.

120

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understanding microstructure and …rk242cd0160/...i would also like to thank mike mcgehee for his...

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