ABSTRACT ROSKOV, KRISTEN EKIERT. Engineered ...

ABSTRACT
ROSKOV, KRISTEN EKIERT. Engineered Organometallic Polymer and Hybrid Systems
Containing Nanoparticles and/or Poly(ferrocenylsilanes). (Under the direction of Professor
Richard Spontak.)
Formation of polymer nanocomposites is becoming an increasingly attractive and
facile means by which to combine the desirable properties of metals and metal oxides (e.g.,
electrical, magnetic, optical, and thermal) with those of polymers (e.g., flexible, lightweight
and tough). Incorporation of nanoscale objects such as spheroidal nanoparticles or elongated
nanorods into electrospun polymer nano/microfibers measuring from 50 nm to 1 μm in
diameter yields functional nanomaterials that can be used in various applications ranging
from data storage and conductive nanowires to nonwoven sensors, magnetic filters and drug
delivery patches. By aligning nanoscale objects in one-dimensional constructs, we expect
that desirable attributes arising from highly anisotropic electronic, optical, thermal, magnetic,
and catalytic properties can be realized. The objective of this study is to gain a better
fundamental understanding of how to controllably align and position nanoparticles and
nanorods within polymer nano/microfibers to generate unique properties. To achieve this
objective, we focus on four specific process strategies. In the first, superparamagnetic iron
oxide nanoparticles (SPIONs) are aligned into one-dimensional nanoarrays through the use
of magnetic field-assisted electrospinning. In this case, an electromagnet is positioned near
the Taylor cone of the suspension to be electrospun so that the magnetic field is oriented
perpendicular to the electric field. Transmission electron microscopy (TEM) is utilized to
ascertain the morphology of the resultant nanocomposite fibers and reveals that the SPION

nanoarrays persist intact beyond 1 μm. Since the magnetic field can be pulsed, the length of
the nanoarrays can be judiciously controlled. Magnetization hysteresis curves measured on a
superconducting quantum interference device yield saturation magnetization and mean
magnetic moment values. Secondly, gold nanorods (GNRs) varying in aspect ratio have been
flow-aligned in electrospun fibers, and the fibers have likewise been aligned to permit long-
range orientation order at both the nanoscale and macroscale. This is an important
consideration in the fabrication of devices spanning multiple size scales. The GNRs within
nano/microfibers exhibit excellent alignment with their longitudinal axis parallel to the fiber
axis. Optical absorbance spectroscopy measurements reveal that the longitudinal surface
plasmon resonance bands of the aligned GNRs are highly anisotropic, depending on
polarization angle, and that maximum absorption occurs when polarization is parallel to the
fiber axis. Lastly, blends of hydrophobic and hydrophilic polymers have been prepared to
control the spatial position of SPIONs within electrospun fibers on the basis of
thermodynamic compatibility. In this case, TEM confirms that a core-sheath nanostructure
naturally forms due to polymer-polymer phase separation and that the hydrophobic
nanoparticles are sequestered in one preferred phase. Lastly, a nanocomposite fiber is
created using only one entity, the organometallic polymer poly(ferrocenylsilane) (PFS)and its
crystalline structure is probed alone and in the presence of SPIONs. Block copolymer
cylindrical micelles of PFS-b-poly(isoprene) (PI) are crosslinked within an elastomeric
matrix of poly (vinylmethoxysilane) (PVMS) and found to maintain their crystalline structure
with the target application being nanowires in soft electronics.

© Copyright 2011 by Kristen Ekiert Roskov
All Rights Reserved

Engineered Organometallic Polymer and Hybrid Systems Containing Nanoparticles and/or
Poly(ferrocenylsilanes)
by
Kristen Ekiert Roskov
A dissertation or submitted to the Graduate Faculty of
North Carolina State University
in partial fulfillment of the
requirements for the degree of
Doctor of Philosophy
Chemical and Biomolecular Engineering
Raleigh, North Carolina
2011
APPROVED BY:
_______________________________ ______________________________
Richard Spontak Saad Khan
Committee Chair
________________________________ ________________________________
Michael Dickey Russell Gorga

DEDICATION
To my parents, Andrea and Bernie, for their steadfast love and support. You have taught me
life’s greatest lessons by example and that has shaped the woman I have become. I love you
so much.
ii

BIOGRAPHY
Kristen was born in Pittsburgh, PA in 1983 to Andrea and Bernie and is an only child.
After graduating from North Allegheny High School, she went on to the University of
Maryland and received a B.S. in Chemical Engineering in 2006. In the summer of 2005, she
was awarded a summer undergraduate research fellowship at the National Institute of
Standards and Technology in Gaithersburg, MD and was hired as an undergraduate research
assistant for the following year. This research experience at NIST resulted in four
publications and was the experience that sparked her interest in graduate school. From there,
she moved to North Carolina to attend graduate school in the department of Chemical and
Biomolecular Engineering at North Carolina State University. In 2007, Kristen was awarded
the National Science Foundation graduate research fellowship and in 2010 she was awarded
the North Carolina Space Grant Fellowship. Kristen is enthusiastic about mentoring and
tutoring girls and encouraging them to pursue a career in science. Kristen will be joining the
Ph.D. Professional Development Program with BASF upon graduation.
iii

ACKNOWLEDGMENTS
Foremost, I’d like to thank my advisor, Dr. Richard Spontak. You have taught me to
be the scientist I am today. The biggest lesson I have learned from you is that the outcome is
usually not what you expect, but discovering the ‘why’ is the best part. You see your
students in the best possible light and I have always strived to meet those expectations.
Thank you for everything.
I’d like to thank my mentors at NIST who motivated me to attend graduate school
and initially sparked my interest in research, namely Thomas Epps, Christopher Stafford,
Michael Fasolka, and Matthew Becker. I can truly say that you shaped the direction of my
life and career. Thank you for taking the time to invest in a SURF student!
My thesis committee, Saad Khan, Michael Dickey, and Russell Gorga, have been
extremely helpful and supportive. I’d also like to thank many other faculty at NCSU; Jan
Genzer, Saad Khan, Joe Tracy, Kirill Efimenko, Roberto Garcia, and Chuck Mooney. I have
learned so much from all of you and thank you for all your help and fruitful discussions we
have had. Thank you to the professors at other universities with whom I collaborated with;
specifically Ian Manners, Lyudmila Bronstein, and Amy Oldenburg.
My fellow PMG peers and officemates made the day-to-day life of graduate school
much more interesting and I feel very lucky to have made lifelong friends. I’d like to
specifically thank Arjun Krishnan, Pruthesh Varagantwar, Omer Gozen, Evren Ozcam, and
Anand Patel. To Arjun and Anand- I can’t imagine going through this experience without
either of you. Thank you for always answering my deluge of questions and lending a helping
iv

hand. Thank you to the undergraduate students that helped with this research; Raleigh Davis
and Kathryn Earley. To Sara for being my constant throughout this journey-I am so grateful
that we met and decided to live together at Cardinal Club! You have motivated me to be a
true friend, a better researcher, and a great cook. To Michelle, for constantly being there on
so many different levels and always listening. To Julie, for enthusiastically reading every
publication and always asking how my polymers are doing.
I’d like to thank all the support I have received from not only my parents but many
other relatives. My grandma Helen Roskov was a strong woman with so much love for her
family. I thank her for the unwavering encouragement she always gave me and wish she
could be here to see what I’ve achieved. I am so thankful for my ‘sister’ Lisa, who initially
suggested I become an engineer when I was young to which I emphatically replied ‘I don’t
want to drive trains!’ You are such a giving, thoughtful person and I’ve always strived to be
half as caring as you. To my goddaughter Sienna, I love you so much and am so excited to
see you grow. I have been so fortunate to have my godparents Aunt Dee Dee and my Uncle
Kenny. It’s not the same without you being here, we miss you terribly. Thank you to all my
other friends and family for being there for me during this journey. I am incredibly lucky.
v

TABLE OF CONTENTS
LIST OF TABLES ................................................................................................................... ix
LIST OF FIGURES .................................................................................................................. x
CHAPTER I: Introduction and Overview ............................................................................... 1
1.1 Introduction ........................................................................................................................ 1
1.1.1 Electrospinning ................................................................................................... 2
1.1.2 Control over Fiber Characteristics ...................................................................... 6
1.1.3 Electropsinning Fiber Material ......................................................................... 12
1.1.4 Imparting Functionality to Electropsun Fibers ................................................. 18
1.1.5 Applications ...................................................................................................... 23
1.2 One-Dimensional Magnetic Nanostructures .................................................................... 28
1.3 Organometallic Polymers ................................................................................................. 31
1.3.1 Overview ........................................................................................................... 31
1.3.2 Poly(ferrocenylsilanes) ...................................................................................... 33
Nomenclature .......................................................................................................................... 38
Figures..................................................................................................................................... 39
References ............................................................................................................................... 61
CHAPTER II: Long-Range Alignment of Gold Nanorods in Electrospun Polymer Nano/
Microfiber ...................................................................................................................... 76
Figures..................................................................................................................................... 87
References ............................................................................................................................... 91
CHAPTER III: Magnetic Field-Induced Alignment of Nanoparticles in Electrospun
Microfibers ...................................................................................................................... 94
Figures................................................................................................................................... 106
References ............................................................................................................................. 113
CHAPTER IV: Using Polymer Blend Morphology to Position Ligand-Functionalized
Nanoparticles in Electrospun Polymer Microfibers .............................................................. 116
Tables ................................................................................................................................... 133
Figures................................................................................................................................... 135
References ............................................................................................................................. 142
CHAPTER V: Nanostructured Organometallic Polymer Systems Containing
Poly(ferrocenylsilanes) ......................................................................................................... 145
5.1 Introduction .................................................................................................................... 145
5.2 Experimental .................................................................................................................. 149
5.2.1 Materials .......................................................................................................... 149
5.2.2 Synthesis of Specialty Polymers ..................................................................... 149
5.2.3 Preparation of Nanoparticles........................................................................... 150
vi

5.2.4 Preparation of Electrospun Fibers ................................................................... 151
5.2.5 Characterization of PFS Nanomaterials .......................................................... 151
5.3 Electrospinning and Characterization of PFS Homopolymers ...................................... 151
5.4 Phase behavior of binary blends of PFS in Elastomeric Matrices ................................. 156
5.4.1 PFDMS-b-PI/PI Blend by Solvent Casting ..................................................... 158
5.4.2 Cross-linking within Polyisoprene .................................................................. 158
5.4.3 Shell-Cross-Linking of Cylindrical PI-b-PFS Micelles .................................. 161
5.4.4 Cross-linked PVMS as the Matrix Polymer.................................................... 162
5.5 Conclusion ..................................................................................................................... 165
Tables ................................................................................................................................... 166
Figures................................................................................................................................... 171
References ............................................................................................................................. 185
CHAPTER VI: Conclusions and Future Work ..................................................................... 189
6.1 Conclusions .................................................................................................................... 189
6.1.1 Long-Range Alignment of Gold Nanorods in Electrospun Polymer
Nano/Microfibers ........................................................................................... 190
6.1.2 Magnetic Field-Induced Alignment of Nanoparticles in Electrospun
Microfibers ..................................................................................................... 190
6.1.3 Using Polymer Blend Morphology to Position Ligand-Functionalized
Nanoparticles in Electrospun Polymer Microfibers ....................................... 191
6.1.4 Nanostructured Organometallic Polymer Systems Containing
Poly(ferrocenylsilanes) ................................................................................... 191
6.2 Recommendations for future work ................................................................................. 192
6.2.1 Gold Nanorod Alignment through Electrospun Fiber Degradation ............... 192
6.2.2 Field Uniformity in Magnetic-Assisted Electrospinning ............................... 193
6.2.3 Poly(ferrocenylsilane) Cylindrical Micelles Oriented within Electropun
Fibers .............................................................................................................. 194
Figures................................................................................................................................... 197
References ............................................................................................................................. 201
APPENDIX ........................................................................................................................... 203
APPENDIX I: Responsive PET Nano/Microfibers via Surface-Initiated
Polymerization .......................................................................................................... 203
APPENDIX II: Generation of Functional PET Microfibers through Surface-Initiated
Polymerization .......................................................................................................... 224
APPENDIX III: Modification of Melt-spun Isotactic Polypropylene and Poly(lactic acid)
Bicomponent Filaments with a Premade Block Copolymer ..................................... 261
APPENDIX IV: Enhanced Biomimetic Performance of Ionic Polymer-Metal Composite
Actuators Prepared with Nanostructured Block Ionomers ....................................... 307
vii

APPENDIX V: Block Copolymer Self-Organization vs. Interfacial Modificationin Bilayered
Thin-Film Laminates ................................................................................................ 329
viii

LIST OF TABLES
Table 4.1. XRD characteristics of PEO powder and electrospun PEO/P2VP
microfibers. ................................................................................................... 133
Table 4.2. XRD characteristics of electrospun PEO/P2VP microfibers with SPIONs
measuring 18 nm in diameter and added at a concentration of 2.5 vol%. .... 134
Table 5.1. Characteristics of the (Co)Polymers Used in this Study. .............................. 166
Table 5.2. PFDMS Fiber Diameters............................................................................... 167
Table 5.3. X-ray diffraction bragg angle, d-spacing, and average crystallite size for 96
kDa poly(ferryocenydimethylsilane) powder and electrospun fibers. .......... 168
Table 5.4. Components utilized in the vulcanization of poly(isoprene). ....................... 169
Table 5.5. Solubility parameters of polymers relative to the solvent n-hexane. ............ 170
ix

LIST OF FIGURES
Figure 1.1. Schematic figure of basic electrospinning setup .............................................. 39
Figure 1.2. Photograph of the whipping motion of the instability region of the polymer jet
during the electrospinning process. .................................................................. 40
Figure 1.3. A method to produce aligned fibers. A) a schematic utilizing the parallel
grounded electrode collection system and B) the resulting aligned polymer
fibers ................................................................................................................. 41
Figure 1.4. SEM micrograph of beaded PEO fibers. .......................................................... 42
Figure 1.5. Reduction in bead density upon increase in polymer solution viscosity .......... 43
Figure 1.6. Ribbon-like fibers formed via electrospinning ................................................. 44
Figure 1.7. SEM image of anatase hollow fibers created via a co-axial electrospinning
setup.................................................................................................................. 45
Figure 1.8. TEM of PPX/Pd TUFT hybrid nanotubes after the pyrolysis of PLA template
fibers and inset is an electron diffraction pattern of Pd crystals ...................... 46
Figure 1.9. SEM micrograph of porous PLA fibers obtained via electrospinning and
subsequent swelling .......................................................................................... 47
Figure 1.10. SEM iamges of anatase nanofibers whose surfaces have been decorated with a)
gold and b) silver nanoparticles via photocatalytic reduction .......................... 48
Figure 1.11. A list of organosoluble polymers and their molecular structure ...................... 49
Figure 1.12. Schematic of polymer/inorganic composite nanofibers when a) inorganic ions
are incorporated into electrospun fibers followed by exposure to gas to
synthesize inorganic nanoparticles both inside and outside of the nanofiber and
b) when only the surface of nanofibers are modified with metal ions. ............ 52
Figure 1.13. TEM micrographs of multicomponent polymer electrospun fibers
demonstrating a) A core-sheath structure formed by a polymer blend b)
Lamellar structure formed by a phase-separated block copolymer c) Cylindrical
structure formed by a phase-separated block copolymer ................................. 53
Figure 1.14. TEM image of a) PAN/CNT composite nanofiber mat and b) demonstrating the
uniform distribution and alignment of CNTs witin a PAN fiber ..................... 54
Figure 1.15. Average Young’s modulus for electrosopun nylon-6 and nylon-6/O-MMT
nanocomposite single fibers vs. fiber diameter ................................................ 55
Figure 1.16. Demonstration that as the aspect ratio of gold nanorods increases, as does the
maximum optical absorbance and thus the color of the aqueous colloidal
suspension. ....................................................................................................... 56
Figure 1.17. Three-dimensional mineralized electrospun fibers mimicking the hierarchical
structure of bone ............................................................................................... 57
Figure 1.18. Left: magnetic induction map from two pairs of bacterial magnetite chains.
Right: A bright-field TEM image of a double chain of magnetite
magnetosomes. ................................................................................................. 58
x

Figure 1.19. Pyroloysis of UV cross-linked PS-b-PFEMS films with a) height-mode
scanning force microscopy, b) phase-mode, c) TEM images, and d) A
schematic of the morphology. Inset scale bars = 50 nm. ................................ 59
Figure 1.20. TEM micrographs of scarf-shaped PI-b-PFS co-micelles (scale bar
= 500 nm) ......................................................................................................... 60
Figure 2.1. TEM image of GNRs deposited from an aqueous suspension onto a carboncoated
TEM grid. The inset shows the distribution of measured aspect ratios of
the GNRs, which measure 49 nm long and 17 nm in diameter on
average.............................................................................................................. 87
Figure 2.2. SEM image of macroscopically-aligned electrospun PEO fibers containing
GNRs. ............................................................................................................... 88
Figure 2.3. Aligned GNRs in electrospun PEO nano/microfibers as functions of fiber
diameter and GNR volume fraction (��: (a) 40 nm and (�� = 0.006, (b) 50 nm
and (�� = 0.045, (c) 650 nm and (���= 0.035, and (d) 3000 nm and (��= 0.031.
A selected-area electron diffraction pattern of the corresponding sample is
included as an inset in (b). ................................................................................ 89
Figure 2.4. Absorbance spectra for (a) randomly oriented GNRs in a PEO film measuring
~500 �m thick at different polarization angles and (b) GNRs aligned within
electrospun PEOmicrofibers measuring ~200 nm in diameter at polarization
angles varying from 0° (parallel to the fiber axis n) to 90° (perpendicular to n).
In both cases, the data are color-coded and labeled in (a). ............................... 90
Figure 3.1. Schematic illustration of the magnetic field-assisted electrospinning setup used
in this study. Note the position of the electromagnet, which yields a magnetic
field that is perpendicular to the electric field employed during
electrospinning. .............................................................................................. 106
Figure 3.2. TEM images of randomly dispersed SPIONs in electrospun PCL microfibers
varying in SPION concentration (in vol%): (a) 0.5 and (b) 2.5. A TEM image
of SPIONs measuring 17.6 nm in diameter and drop cast from chloroform is
included in the inset of (a). The scalemarker in the inset corresponds to
50 nm. ............................................................................................................. 107
Figure 3.3. TEM images of magnetic field-aligned SPIONs, measuring 17.6 nm in
diameter, in PCL microfibers illustrating long, contiguous arrays in (a), and
shorter, pulsed arrays in (b). The scalemarker in the inset corresponds to 100
nm ................................................................................................................. 108
xi

Figure 3.4. Magnetization (M) hysteresis curves at 300 K as a function of the magnetizing
field strength (H) for SPIONs measuring 17.6 nm in diameter. In (a), the
hysteresis curves are measured from unembedded SPIONs ( ), as well as
randomly dispersed and magnetically aligned SPIONs in electrospun PCL
microfibers (blue and red, respectively). The inset shows magnetization
hysteresis curves recorded for the embedded SPIONS at low fields and ambient
temperature. In (b), the hysteresis curves from the SPIONs embedded in PCL
microfibers (see the corresponding diagrams) are fitted to Eq. 2 in the text
(solid lines) to discern the saturation magnetization and mean dipole moment
from each dataset. ........................................................................................... 109
Figure 3.S1. SEM image of SPION-containing PCL fibers electrospun in the presence of an
external magnetic field. The inset shows evidence of surface dimpling on a
large fiber. The scalemarker in the inset corresponds to 2 �m. ..................... 111
Figure 3.S2. EDS spectrum of a SPION-containing PCL fiber electrospun in the presence of
an external magnetic field. The elements responsible for the observed peaks are
labeled, and the x-ray energies associated with the K� and L lines of Fe are
identified by the blue lines. ............................................................................ 112
Figure 4.1. TEM image of SPIONs measuring 16.4 nm in diameter and drop cast from
chloroform. ..................................................................................................... 135
Figure 4.2. (a) SEM image of SPION-containing PEO/P2VP microfibers composed of 80
wt% PEO and electrospun from an 8.5 wt% solution in chloroform. (b) An
enlargement showing the surface of the microfibers included in (a). The inset
in (b) displays a SPION-rich bead, the scalemarker corresponds to
500 nm. ........................................................................................................... 136
Figure 4.3. TEM images of SPIONs measuring 16.4 nm in diameter and dropcast from
chlorform....................................................................................................... 138
Figure 4.4. XRD patterns acquired from PEO powder and electrospun microfibers
composed of PEO/P2VP at different PEO concentrations (labeled). ........... 138
Figure 4.5. XRD patterns acquired from electrospun microfibers composed of PEO/P2VP
with SPIONs (18 nm and 2.5 vol%) at different PEO concentrations
(labeled). ....................................................................................................... 139
Figure 4.6. Average PEO crystal size (t) extracted from XRD patterns and presented as a
function of PEO concentration parallel (circles) and perpendicular (triangles)
to the fiber axis for systems without (open symbols) and with (filled symbols)
SPIONs. Values measured from PEO powder are included (triangles). The
solid and dashed lines serve as guides for the eye. ....................................... 140
xii

Figure 4.7. Schematic illustration depicting the arrangement of polymer chains in a
core/sheath microstructure of PEO/P2VP (a) before and (b) after SPION
addition (not to scale). Addition of SPIONs promotes a reduction in crystal
size but a more parallel chain arrangement with respect to the fiber axis. ... 141
Figure 5.1. Schematic of PFEMS synthesis in which the noted molecules are a)
ferrocenophane b) ethylmethylsilaferrocenophane and c) poly
(ferrocenylethylmethylsilane). ...................................................................... 171
Figure 5.2. SEM micrographs of a) 15% PFDMS in THF:DMF without surfactant b) 15%
PFDMS in THF:DMF with surfactant c) 20% PFDMS in DCM with
surfactant and d) 18% PFPMS in THF:DMF with surfactant. All scale bars
represent 20 μm. ............................................................................................ 172
Figure 5.3. PS-b-PFS-b-P2VP lithographic template used for the preparationof Nanoscale
magnetic dots. a) Phase-separation of the triblock in the bulk. b) Hollow PFS
cylinders are formed after etching because it has a selective resistance. c)
Profile of the hollow PFS cylinders. (Used with permission by Jessica
Gwyther from Bristol University.) ................................................................ 173
Figure 5.4. XRD curves from 2Ɵ = 7 – 25 °for PFDMS powder, fibers, and fibers with
both 10 and 14 nm iron oxide nanoparticles. ................................................ 174
Figure 5.5. Schematic demonstrating the chains of PFDMS and corresponding d-spacing
between adjacent ferrocene units in a) powder form, b) in electrospun fibers,
c) in electrospun fibers with larger (~14 nm) iron oxide nanoparticles, and d)
in electrospun fibers with smaller (~10 nm) iron oxide nanoparticles. ........ 174
Figure 5.6. TEM micrographs of PFS-b-PI micelles dropcast from a 1 mg/mL hexane
solution onto a carbon-coated TEM grid with an average width of 14.9 nm and
lengths exceeding one micron. ...................................................................... 176
Figure 5.7. TEM micrographs of PFS-b-PI micelles blended at a ratio of a) 3:100 with PI
and b) 3:25 with PI. ....................................................................................... 178
Figure 5.8. Schematic of the vulcanization and micelle crosslinking technique. ........... 178
Figure 5.9. TEM micrographs of 1:1000 PFS-b-PI:PI vulcanized at ~120 °C for 5 housr
demonstrating a complete dissolution of the micelles with only iron
nanoparticles remaining. ............................................................................... 179
Figure 5.10. TEM micrographs of shell cross-linked PFS-b-PI with an average width of 43
nm and lengths exceeding 1.5 microns dropcast from a hexane solution. .... 180
Figure 5.11. TEM micrographs of PFS-b-PI micelles blended with a 5 wt% PI solution in
hexane and dropcast onto a carbon-coated TEM grid. ................................. 181
Figure 5.12. TEM micrographs of microtomed PFS-b-PI micelles blended with PVMS at
ratios of 1:300 and 1:600 at thicknesses of ~120 nm. ................................... 182
xiii

Figure 5.13. TEM micrographs shell cross-linked PFS-b-PI dropcast from a hexane
solution as a) a control, b) heated to 70 °C for 30 minutes and c) stirred for
several minutes mimicking conditions during the cross-linking of the PVMS
solution. ......................................................................................................... 182
Figure 5.14. SEM micrographs of a PVMS cross-linked film containing PFS-b-PI micelles
a) looking down the surface from the cross-section and b) of the fractured
cross-section where the scalebar of the inset refers to 200 μm. .................... 182
Figure 6.1. Dropcast PFDMS-b-P2VP cylindrical micelles onto a carbon coated TEM
grid from DMF. ............................................................................................. 182
Figure 6.2. SEM micrographs of 32 wt% P2VP fibers electrospun from 9:1 DMF:THF a)
without micelle addition and b) with PFDMS-b-P2VP micelles. ................. 182
Figure 6.3. TEM micrograph of 32 wt% P2VP fibers electrospun with PFDMS-b-P2VP
micelles. ........................................................................................................ 182
xiv

Introduction
Chapter I
Introduction and Overview
As technology has developed in the last century, it has become possible to investigate
matter on a smaller and smaller scale and to manipulate matter on an atomic and molecular
scale. This development has led to the creation of the buzz word “nanotechnology.” which
typically refers to structures with a size scale between 1 and 100 nanometers. If we
specifically look at the construction of soft materials, they can be created through two
different approaches: a “bottom-up” approach, where molecules self-assemble on the atomic
level; and a “top-down” approach, using patterned templates to achieve atomic order. Soft
nanomaterials viewed through either approach are interesting mainly due to the weak, non-
covalent bonding that occurs within them and the ability to use thermal energy to break and
re-form these bonds. Changes in morphology can be induced by thermal energy such as
temperature, pH, or other external triggers. The term “nanocomposite” refers to a multiphase
material where one of the phases is nano-scaled and the phases differ both in structure and
chemistry, thereby allowing the formation of multifunctional materials with very high surface
area-to volume ratios of the dispersed phase. In this chapter, the creation of polymeric
1

materials with functional properties by utilizing electrospinning as a fabrication tool and the
processing or organometallics will be discussed.
1.1.1 Electrospinning
History
Generally, electrospinning is the process of forming fibers through the application of
electrostatic forces. It is an extension of the electrostatic spray technique where a high
voltage is applied to a liquid jet in order to form small droplets. Electrospinning was initially
investigated by Lord Rayleigh in 1882 1 when he questioned the electric potential that would
be needed to overcome the surface tension of a drop. Then, in 1914, Zeleny et al.
investigated the behavior of fluid droplets at the end of glass capillaries. 2 Electrospray is
used to manufacture particles of varying size and is used in applications including mass
spectrometry, painting, and inkjet printing. 3 The electrospinning of plastics first appeared as
a patent in 1934 by Formhals 4 however it was decades later in the 1990’s when the Reneker
group revived interest in this technique 5 with a goal of forming ultrathin, continuous polymer
fibers. Other fiber forming techniques include drawing, 6 melt-blowing, 7 and multicomponent
fibers. 7 Though these last two in particular have very high productivity, only electrospinning
has the ability to form continuous fibers with nanoscale dimensions and with flexibility in
terms of polymer choice. Since then, the interest in electrospinning has exploded and more
than 100 polymers have been electrospun 8 due to the simplicity of the process and versatility
it affords.
2

An example of an electrospinning scheme is seen in Figure 1.1. First in the
eletrospinning process, a polymer solution is drawn into a syringe with a spinneret, or
metallic needle, and connected to a syringe pump that produces a steady and controllable rate
of discharge. A direct current power supply creates an electric field with high voltages
between the needle tip and grounded collection plate that typically ranges from 1-20 kV.
Surface tension is responsible for holding a droplet of polymer solution on the tip of the
spinneret and when an electrode is connected to the spinneret and a voltage is applied, a
charge develops on the surface of the liquid due to mutual charge repulsion. 9 As the voltage
is increased, the hemispherical shape of the droplet becomes elongated into a conical
structure defined as a “Taylor cone” 10 and as the intensity of the electric field increases
further, a critical threshold is passed where electrostatic forces overcome surface tension and
a charged jet is emitted from the tip of the Taylor cone. The ejected polymer solution
undergoes a whipping process, 11 as seen in Figure 1.2 12 between the spinneret tip and the
collection plate during which time the solvent is evaporated. The dry, highly porous,
randomly oriented polymer fibers are deposited on the collection plate.
Electrospraying, the formation of droplets, or bead defects – fibers with beads on a
string – can occur instead of continuous fibers depending on a large set of parameters. These
include polymer properties such as molecular weight, solubility, and glass-transition
temperature; solution properties such as concentration, viscosity, surface tension, electrical
conductivity; and processing parameters such as applied voltage, plate distance from the
spinneret, and solution flow. Even ambient properties such as relative humidity and solvent
3

vapor pressure can have an effect on the resulting fiber morphology. Thus, it is demonstrated
that although electrospinning appears to be a straight-forward process, it is actually
multivariate problem that requires a high degree of optimization in order to achieve a
successful result.
Modification of the Electrospinning Setup
Since the electrospinning setup is so simple, it is very easy to modify in order to
change the structures and morphology of the resultant fibers. In the past, a rotating collector
plate has been used to recreate both uniform and aligned fiber mats. 13 Arrays of multiple
needles 14 have also been utilized in order to increase the productivity of the electrospinning
process but potential problems with interference between needles also becomes an issue.
Modification of the spinneret needle can also be performed to form a co-axial setup 15 which
creates a single fiber with a core-sheath structure and enables the ability to electrospin
immiscible polymer blends. A multi-axial 16 spinneret can also be constructed where multiple
layers of a fiber can be created. In addition, by switching the typical placement of the anode
and cathode in the electrospinning setup, it is possible to draw the more polarizable
component to the surface. 17 Furthermore, the type of current, alternating or direct, has been
found to have a slight impact on the degree of orientation and fiber density. 18 Needleless
electrospinning has also been successful, and eliminates the burden of clogged needles, by
utilizing magnetic fluids, 19 a pointed needle and the use of a tungsten electrode, 20 and
conducting electrodes. 21
4

The straightforward electrospinning setup discussed earlier creates randomly-oriented
fiber mats, however, many applications, such as the fabrication of electronic or photonic
devices, 22 will require well-aligned, unidirectional fiber constructs. Thus, the macroscopic
alignment of electrospun fibers is one processing modification that has garnered a lot of
attention. Another approach involves utilizing a rotating drum at high rotating speeds, which
has been found to orient fibers along the winding direction 23 and also along the sharp edge of
a tapered, wheel-like disc. 24 Both of these techniques have a high throughput, but
electrostatic differences in attraction across the drum have led to different densities and
alignment that is not perfect. Split electrodes are another common geometrical configuration
utilized to produce aligned fibers. Additionally, uniaxial alignment can occur between two
conductive strips separated by a variable gap (typically of several centimeters) on the
collector plate as seen in Figure 1.3. 25 This technique allows for control over the aligned
fiber density, length of alignment, and ease of removal for imaging or processing. Finally,
four electrodes have also been used to form a cross-bar array of electrospun fibers, 26 a
morphology needed for device fabrication. As electrospinning gets more developed towards
scale-up and industrial applications, the ability to finely control the alignment of polymer
fibers will become imperative.
Other types of Electrospinning
Traditional solutions used for electrospinning must have a sufficient viscosity in order
to produce the appropriate surface tension and usually produces fibers with diameters less
than 1 �m. For higher solution viscosities, filming can occur on the needle and impede the
5

formation of fibers. An alternative is electrospinning from the melt,which often leads to
much higher fiber diameters, but also must be performed under vacuum, and necessitates
much larger electrode separations and thus electric fields. 27 Bowl electrospinning has a 40x
increase in production rate and utilizes a bowl filled with a polymer solution and a concentric
cylindrical collector. 28 Electrospinning polymers has also been proven to be effective using
supercritical CO2 as the solvent 29 where solvent-free systems, specifically for biocompatible
polymers, is needed. Another method involves blowing-assisted electrospinning, which is
the use of a hot air stream for systems with very high viscosity. This has proven successful
for hyaluronic acid, for example, a natural polysacchride with very high viscosity, in which
no solvent was needed to create fibers. 30 Some downsides of this technique, however, are
very large resultant fiber diameters, no increase in throughput, and a very large gas volume
that is needed in addition to the expense of modifying the setup. 31 All of these methods have
been formulated to adjust for limitations to traditional solution systems and maximize
functionalities and target applications.
1.1.2 Control over Fiber Characteristics
External Structure
Both the morphology and diameter of electrospun fibers depends on many of the
processing and solution variables listed earlier. One of the more prevalent problems is the
appearance of beads in the fibers, as seen in Figure 1.4. Of higher importance in the case of
unidirectional fiber arrays is that any disturbance to the smooth fiber axis will negate the
effect of alignment. Beads are often formed because surface tension drives the conversion of
6

the polymer jet into spherical drops 1,32 as a way to decrease surface area. However, the
forces acting against surface tension, specifically, electrostatic repulsion and viscoelastic
forces, endeavor to form smooth fibers. Thus, in order to damper the appearance of beads, it
is necessary to decrease the surface tension. This can be achieved by increasing the viscosity
of polymer solutions, typically by an increase in molecular weight or concentration. The
effect of this is clearly seen in Figure 1.5. Another way to eliminate beads is by the addition
of salts 33 to increase the net charge density or blending of solvents 34 to lower the surface
tension. Deitzel et al. have also determined that a threshold voltage exists 35 and when this is
exceeded, it acts to weaken the stability of the jet and thus more beads are formed. Thus, as
the above shows, there a multitude of ways to control bead density which is dependent on the
specific polymer/solvent system.
Fiber diameters in electrospun fibers are usually reported as a distribution and can be
controlled by several means. 36 As more concentrated polymer solutions are used, fiber
diameters become thicker. In addition, if the conductivity of a solution is increased, to
reduce bead defects for example, then the resulting fiber diameter decreases. A higher feed
rate will lead to thicker fibers. Rutledge et al. completed a comprehensive study on the
multitude of variables that will affect fiber diameter and found that it is a fine balance
between flow rate, the strength of the electric field, and the surface tension of the polymer
solution that mostly affects fiber diameter. 37 Additionally, while fibers are typically circular
in cross-section, other shapes can also occur. Ribbon-like structures with rectangular cross
7

sections 38 have been found at higher viscosities and attributed to a film developing on the
surface of the liquid jet and its subsequent collapse, as seen in Figure 1.6.
Internal Structure
One of the ways to impart functionality to fibers is by altering the internal structure.
“Nanotubes” are electrospun fibers with a hollow interior that are most often created from a
co-axial electrospinning setup and encompass the internal structure. This has applications in
nanofluidics, drug delivery, and vascular engineering. Ceramic nanotubes are created by
electrospinning two immiscible materials in a coaxial electrospinning setup in which the core
can later be removed by solvent extraction. Ceramic nanotubes can be formed by the
resultant calcination of the sheath at elevated temperatures. 39
The benefits of using the coaxial spinning setup to form nanotubes is that the
thickness of the sheath and inner diameter of the fibers can be varied by changing the
spinning conditions 40 as seen in Figure 1.7. 39,41 Although these hollow fibers are very
advantageous, they are also very fragile. In the past, this problem has been eliminated by the
introduction an inorganic sol-gel precursor into the spinning solution and the subsequent
formation of a gel network in the polymer sheath. 42 Nanotubules have also been created with
porous surfaces, which were found to have higher rates of enzymatic reaction and are
excellent candidates for applications requiring substrate penetration. 43 TiO2 nanotubes have
been further functionalized by controlling the hydrophobic and hydrophilic properties on
both the inner and outer surface and allowing adsorption of inorganic nanoparticles. 44 Fluid
flow through these nanotubes would open up avenues for a variety of applications and flow
8

inside carbonized hollow fibers has been demonstrated successfully with pressure-driven
water in nanofluidic devices consisting of ~40000 nanotubes. 45
In addition, electrospun fiber mats can be used as templates for the preparation of
hollow fibers by the tubes by fiber templates (TUFT) process. 46,47 In this process the
dissolvable electrospun fibers are coated with polymer, metals, or other materials. A
negative replica is then created by the selective extraction of the template fiber. Composite
fibers can then be created, for example, by coating poly(lactic acid)(PLA)/Pd(OAc)2 with
poly(p-xylylene) (PPX) and pyrolyzing the PLA resulting with hollow PPX fibers with Pd
nanoparticles on the interior 46 as seen in Figure 1.8. This production technique can be used
to create ceramic nanotubes where a co-axial setup is not necessary, and has been performed
with titania, 46 aluminum 48 and gold 49 just to name a few.
Internal Poylmer Morphology
A decade ago, the general goal of electrospinning was to establish which polymers
could be electrospun from which solvents and gain a better understanding of how different
solution/processing conditions affected the resultant morphologies. Now with this breadth of
knowledge available, scientists are asking how to create a given morphology or property.
Electrospinning is a fabrication technique that operates far from equilibrium. In fact,
the time scale between a volume element leaving the jet to being collected is ~0.01
seconds. 50 The evaporation of solvent and the elongation of the jet control both the internal
and surface structure of the resulting fiber during this time frame and the choice of solvent
can also affect the resulting fiber morphology, less volatile solvents may not completely
9

evaporate and residual amounts could encourage chain relaxation. The elongation of the jet
can result in crystalline polymers achieving chain orientation parallel to the fiber spinning
direction, and crystal formation in electrospun fibers can even mirror those in processes with
longer time scales, such as extrusion. 51 Similar to polymer thin films, annealing of fibers will
convert crystals to a more ordered phase 52 but also requires the fiber to be encapsulated by a
sheath material 53 to prevent flow.
While electrospinning typically produces fibers with smooth surfaces, the creation of
a porous surface resulting in an increase in surface area and is desirable for applications such
as tissue engineering, filtration, and catalysis. Fiber mats can also be exposed to a solvent
that induces swelling, once the swelling agent is removed then what remains is a fiber with a
larger diameter and pores. One way porous fibers can be created is by the electrospinning of
immiscible polymer blends. 54 For example, if a blended fiber is immersed in a selective
solvent, one phase can be etched out and leave pores. Blended fiber mats can also be
exposed to a solvent that induces swelling, once the swelling agent is removed then what
remains is a fiber with a larger diameter and pores (Figure 1.9). Porosity can also be created
by electrospinning fibers into a bath of liquid nitrogen, 55 which causes a phase separation of
solvent and polymer.
Although the structural characteristics of electrospun fibers are advantageous, the
bulk properties tend to lack functionality that is desired for more multifunctional
applications. One way of overcoming this problem is to create composite nanofibers,
allowing the incorporation of two chemically and physically different components which can
10

enhance the mechanical, 56 conductive, 57 and magnetic 58 properties, just to name a few.
However, encapsulated molecules often show reduced activity 59 when constrained in a
polymer matrix and may not always exist at the surface. For example, when antibacterial
biocides are blended with a polymer prior to electrospinning their efficacy is limited and
leave no potential to attack airborne pathogens. 60 In contrast, surface modification through
the covalent bonding of poly(quarternary ammonium) can create a permanent antibacterial
surface on a fiber. 61
Because polymer surfaces exhibit typically low surface energy, 62 they need to be
treated chemically or physically. Simple surface coating is usually the most straightforward,
chemical modification for this treatment. In simple surface coating, electrostatic interactions
or liquid-phase attachment are responsible for the deposition of conducting polymers or
nanoparticles to the surface of electrospun fibers. 46 Physical or chemical vapor depositions
can then be utilized to coat fiber mats with ceramics, polymers, or metals 48 as seen in Figure
1.10. 63 In addition, precursor polymers can be spun, 64 such as poly(acrylonitrile)(PAN), 65
and then subsequently carbonized to prepare ceramic fibers. Physical treatment techniques
include plasma, 66 the formation of ‘layers,’ 67 ultraviolet treatment, 60 mineralization, 66
etching, 68 or the inclusion of a composite material that is reactive. 69 Once there are
chemically-active groups present on the surface, covalent bonding, 70 immobilization, 71 and
electrostatic interactions 72 can be used to stabilize reactive groups to the surface of the fiber.
Modification of just the surface of polymer fibers, not the bulk, can open these materials up
11

for multifunctional applications where the fibers can interact with their environment such as
filtering of target molecules, 73 protective textiles, 74 tissue scaffolds, 75 and drug delivery. 76
Additional Properties of Electrospun Fibers
Strength is an important consideration for many of the applications targeted by
electrospun fiber scaffolds. Whereas tensile testing can easily be performed on macroscopic
nonwoven fibers due to their larger size, the physical manipulation of nanofibers becomes an
issue. One way to possibly alleviate this is by utilizing nano tensile testing on single fibers
which has been reported for PLA. 77 While inconsistent results have been collected for tensile
testing of a randomly oriented fiber mat, mechanical property trends have been found to exist
for oriented fibers collected on a rotating drum as a function of the velocity of the drum. 78
Nanoindentation is perhaps the best method to measure the elastic properties of a single
fiber. 79 In situ tensile testing has been performed combining atomic force microscopy
(AFM) and scanning electron microscopy (SEM) and it was determined that electrospun
fibers had an increased tensile stress and elastic modulus as compared to bulk samples. 80 A
uniform and thus comparative method of measuring the mechanical properties of both
electrospun polymer fiber mats and individual fibers is a characterization technique that
needs to be better established as this field progresses.
1.1.3 Electropsinning Fiber Material
One of the greatest benefits of electrospinning is the ability to easily tune the fiber,
solvent, and processing parameters in order to create the desired material. Several classes of
12

electrospinnable polymers will be briefly discussed as an overview and as a basis for the rest
of this document.
Types of Fiber-Forming Polymers
Many of the applications for electrospinning are biological-based, ranging from tissue
scaffolds 81 to wound dressing 82 to artificial blood vessels. 83 For this reason, the creation of
electrospun mats utilizing biopolymers or blends of biopolymers is highly sought after and
has broad application. Natural polymers tend to exhibit better biocompatibility when used in
biomedical applications. Protein fibers, for example, including collagen, gelatin, elastin, and
silk fibroin have been very well studied. 84 Although most of these materials generally lack
strength, silk is unique in that there are multiple sources (silkworm 85 vs. spider) 86 and the
prepared mats can be treated to induce conformational changes to beta-sheet structures to
enhance its mechanical properties. 87 Polysacchrides such as dextran, 88 chitosan, 89 and
cellulose acetate 34 have also been electrospun to form fiber scaffolds. In addition, many
types of proteins and enzymes such as lipase, 90 cellulase, 91 and bovine serum albumin
(BSA), 76 cannot be electrospun alone and thus are blended with biopolymers to create fibers.
Fibers containing enzymes are often termed “bioactive” and, surprisingly, demonstrate
increased enzymatic activity when immobilized in electrospun fibers. 76 Control over the
release rate can be accomplished by either coating the external surface of the fiber or by
coupling the enzymes to the fibers. 92
Synthetic biopolymers, on the other hand, offer the advantage to tailor the resultant
properties and are less expensive to fabricate. Hydrophobic polyesters such as poly(glycolic
13

acid) (PGA), 93 PLA, 48 and poly(caprolactone) (PCL) 94 have all been easily electrospun.
Hydrophilic biodegradable polymers such as polyurethane (PU), 95 poly(vinyl alcohol)
(PVA), 36 poly(ethylene oxide) (PEO), 96 polydioxanone, 97 and polyphosphazene derivatives 98
have been electrospun for biomedical applications. Synthetic copolymers have also been
derived in order to combine desirable properties from two or more biopolymers such as
mechanical robustness, morphology, pore size, biodegradability, cell affinity, etc. 99
Poly(lactide-b-glycolide) (PLGA) is one copolymer that is especially well-studied and
consists of a glycolide and lactide. Depending on the ratio of glycolide/lactide, the
mechanical properties and biodegradability can be carefully tuned and have been tested for
use as a antibiotic delivery vehicle. 100
Water soluble polymers like PEO, PVA, poly(acrylic acid) (PAA),
poly(vinylpyrrolidone) (PVP), hydroxypropylcellulose (HPC) are highly desirable for
biomedical applications because the use of corrosive solvents can be avoided. PEO has been
particularly well studied due to the range of molecular weights available, its solubility in a
multitude of organic solvents in addition to water, 24,101 and its biocompatibility. The use of
PVA is also advantageous since its degree of hydrolysis or water solubility can be varied. 102
The reactive hydroxy groups also open up the possibility of reaction or modification. The
solubility of polymers in water can also be adjusted by the pH value, temperature, or the
addition of surfactants or solvents. 103 Although aqueous electrospinning eliminates the need
for toxic solvents, applications such as filtration or use in textiles is not possible and,
therefore, these mats are often crosslinked to improve water resistance. Although chemical
14

crosslinking is very effective, 104 cross-linking agents can alter the properties of the polymer
and lead to degradation. Thus thermal cross-linking has been investigated for blends of PVA
and cylcodextrin 105 and photo cross-linking with PVA derivatives. 106
There is a large list of organosoluble polymers as seen in Figure 1.11. 103 Organic
solvents are useful for the ability to control volatility, conductivity, pH, and polarity. Many
properties of organic solvents are not apt for industrial processes; namely flammability,
toxicity, and corrosiveness. While there are many organosoluble polymers, just a few will be
touched on in this chapter as we discuss the interesting properties that can result. First, PAN
can be electrospun from dimethylformamide (DMF) and converted into carbon fibers via
pyrolysis. 107 This creates a class of materials used as reinforcement in aerospace, industrial,
and aerospace applications. 108 PAN fibers have also been reinforced with metallic oxide
nanoparticles 109 and multi-walled carbon nanotubes. 110 Aliphatic poly(amide) (PA), or more
commonly known as nylon, can be electrospun into very thin and uniform fibers but often
require solvents that are very toxic. 111 Due to their high solvent and thermal stability, PA
electrospun mats are also being investigated for filtration applications. 112 Poly(ethylene
terephthalate) (PET), an aliphatic polyester, is known for its excellent structural and
mechanical strength, transparency, and resistance to many solvents and has been electrospun
in several applications. 113 The surface of PET can also be chemically modified to create a
more functional surface. Poly(vinylidene fluoride) (PVDF), a polymer known for its piezo-
electric properties, has been electrospun for applications in lithium ion polymer batteries. 114
Finally, PU fibers are of interest due to their high flexibility and shape-memory properties 115
15

and may have applications in wound dressings. This is simply a brief demonstration of the
breadth of properties and applications that can be accessed through organosoluble polymers.
Inorganic/Polymer Composite Fibers
Although electrospinning is often associated with polymer fibers, there have been
many reports of ceramic and inorganic fibers formed as well which have been touched upon
already. These fibers can be created in three ways: electrospinning sol-gel precursor
solutions, gas-solid reaction, or in situ photoreduction. The sol-gel method combines a
colloidal solution (sol) which acts as a precursor to an integrated network (gel) of network
polymers. 116 The main challenge of this technique is the creation of a viscosity that is
equivalent to that of a viscous polymer solution. The hydrolysis rate of a sol-gel precursor
can be controlled by changing the pH value or aging conditions. 103 The direct
electorspinning of viscous inorganic sols have successfully created fiber mats of
TiO2/SiO2, 117 PbZrxTi1-xO3, 118 and SiO2 119 fibers. A slightly different approach to the
formation of inorganic fibers is through the use of a series of mesoporous molecular sieves
for the electrospinning of sols and structure-directing agents. 120 This approach involves the
formation of ceramic fibers when the organic portion is removed by calcination at elevated
temperatures and has been applied to anatase fibers 25 and many other metal oxides. Non-
oxide ceramic nanofibers have also been prepared by electrospinning a sol solution
containing Novolac resin and tetraethylorthosilicate followed by conversion to SiC through
pyrolysis. 117 Electrospinning of inorganic and ceramic fibers may become more widespread
16

especially when considering applications requiring a strong structure, membranes, catalytic
supports, or actuators.
Gas-solid reactions were initially introduced to incorporate semiconductor
nanostructures into electrospun fibers. 121 This technique involves co-dissolving a metal salt
and polymer into one solvent and then electrospinning to obtain a polymer/metal salt
composite nanofiber which can then be exposed to a H2S gas at room temperature to
synthesize PbS nanoparticles in situ. The color of the composite quickly turns from white to
yellow after exposure to the gas. 122 Both 20 nm diameter spherical nanoparticles 123 and
nanorods have been obtained on the surface and internally through this technique (Figure
1.12). Chloride nanoparticles have also been obtained from exposure to hydrochloric acid
(HCL). 121 Although this technique consists of multiple steps, it is possible to create
nanofibers with metallic nanoparticles through the simple exposure to a gas.
Metal colloids can also be created through photochemical reduction. UV irradiation
of aqueous solutions of AgNO3 induces photo-oxidation of water and results in the formation
of silver atoms. 124 This same idea can be applied to polymer solutions where the composite
fibers are UV treated and have been demonstrated with CA/AgNO3 composite fibers 125
which resulted in the formation of silver nanoparticles as small as 5 nm were formed. Gold
colloids from HAuCl4 can be reduced under UV irradiation but require an organic stabilizer
such as PVP or PVA. 63 In addition, a photocatalyst of TiO2 is also required and these
applications could include chemical or biological sensing. Although in situ formation of
17

metallic nanomaterials can result in high and uniform loadings, they also usually require the
use of catalysis and external modifications in order to be successful.
1.1.4 Imparting Functionality to Electropsun Fibers
Functional, organic fibers can be created by combining two or more polymers with
different characteristics; this can encompass polymer blends, block copolymers, and
encapsulating functional materials into the electrospinning solution.
Multicomponent Polymer Systems
Both miscible and immiscible polymer blends can be electrospun in order to create
new morphologies of phase separated nanofibers. When two polymers phase separate into
discrete domains, one component can be selectively removed in order to form porous fibers.
This has been conducted with PLA/PVP, 54 PEO/silk, 126 and PGA/chitin, 127 When more
controlled phase separation occurs, core-shell structures 128 or semiconducting wires 129 can be
formed. A core-sheath structure is highly desirable due to the fact that it is possible to
combine polymers with two different sets of properties; i.e. the sheath can be insulating while
the core can be conductive. 128 In addition, the core can also be etched to form nanotubes 39
which may be beneficial for drug release 130 and sensors. 131 Core-sheath structures (Figure
1.13a) can be formed via careful solvent selection, 132 due to thermostatic/kinetic differences
between two polymers, 133 utilizing a co-axial spinning method, 134 or induced by an electric
field. 17 When melt polymer blends are electrospun, it is also possible for a chemical reaction
to occur. With the broad range of polymers and their respective properties, there are endless
combinations of blend fibers that can be created by varying the ratio, molecular weight, and
18

number of components. Molecular weight control of one component can be used to tune the
resulting properties of the fiber mat, for instance by changing the molecular weight of PVA
in a blend system it was possible to reduce the formation of beads. 135 This can be a favorable
process, as in the case of copolymerization reactions during melt electrospinning. 136 Some
benefits to using polymer blends to create nanostructured fibers are careful control over
microphase size based on both molecular weight and polymer:polymer ratio, a direct control
over hydrophobicity/phillicity based on polymer choice, and ease of setup.
Electrospun block copolymer systems have restricted phase separation and usually
micro-phase separate into domains with sizes less than 100 nm. 137 Poly(styrene)-b-
poly(isoprene) (PS-b-PI) has been found to phase separate in electrospun fibers into both a
cylindrical and lamellar morphology parallel to the fiber surface as shown in Figure 1.13b-
c. 138 A block copolymer/amphiphile system, PS-b-poly(2-vinyl pyridine) (P2VP)/
poly(diallyl phthalate), was also found to phase separate into elongated spherical domains in
thin films and the bulk. Fibers with controlled hydrophobicity or hydrophilicity can also be
achieved by electrospinning with block copolymers. 139 Electrospun block copolymer
systems are most applicable for biomedical applications where the decomposition and
biocompatibility rate of the block copolymer can be carefully tuned. 140,141 Although the
ability to harness the periodic micro-phases that can be created with an electrospun block
copolymer system is highly appealing, a significant challenge is the aspect of annealing. In
thin or bulk films, the shape of the system is not impacted when annealing occurs. With a
fiber however, flow can cause a loss of shape. This loss of shape has been avoided by Kalra
19

et al. where the fibers are coated in silicon prior to annealing. 53 The question that remains is
whether this process has scale-up capability but whether block copolymers remain a great
way to create multifunctional polymer nanofibers.
Encapsulation of Functional Materials
Incorporation of inorganic nanomaterials into polymer matrices can yield hybrid
polymer nanocomposites with synergistic properties. These materials are often added to the
polymer solution prior to electrospinning but can also be formed in vivo from precursors in
the polymer solution.
Carbon nanotubes (CNTs) are of particular interest to scientists due to properties such
as electrical, electronic, and thermal conductivity 142 as well as its high strength, flexibility,
and resilience. 143 These properties can be imparted to the nanofiber when CNTs are
combined with the polymer solution prior to electrospinning. It has been well documented
that CNTs orient with their long axis parallel to the electrospun fiber axis, as seen in Figure
1.14. 144 Both single-walled 145 and multi-walled 146 have been electrospun but often require
high concentrations 147 or use of a surfactant in order to get good dispersion.
Functionalization of CNTs can be accomplished through esterification 148 or oxidation 110 and
directly affects the mechanical properties, where functionalized CNT had twice the tensile
strength of unmodified CNTs. 148 This was most likely due to better dispersion and better
interaction between the nanotube and the polymer matrix. A theoretical model was
formulated by Cohen et al. and it was determined that CNTs align along the jet stream
immediately prior to being expelled from the spinneret due to the sink-like flow leading to
20

the jet. 149 The electrical conductivity of CNT polymer composites was also tested and found
to reach 10 -2 S cm -1 . 150 Indeed, the use of CNTs continues to be investigated as a reinforcing
agent.
Although many applications would benefit from encapsulated enzymes or other
bioactive species within a fiber matrix, a large challenge exists in maintaining their activity
post-electrospinning. They are usually sensitive to heat and can lose activity when exposed
to solvents or other chemicals. Bioactive agents such as hyaluronic acid, 151
glycosaminoglycan, 152 heparin, 153 and bone morphogenic proteins 154 have all been
incorporated into polymer fiber matrices without losing their activity. One target application
of these composite nanofibers is for the detection of molecules such as ammonia 155 or
glucose. 156
Other notable, functional materials that have been electrospun include nanoplatelets,
carbon black, graphene, and quantum dots. The most common nanoplatelet is
montmorillonite (MMT) and its purpose is usually that of a reinforcement in polymer
matrices to improve mechanical strength and thermal stability. 157 Fiber size plays an
important role in this, as Li et al. determined that when smaller Nylon/MMT 158 fibers had a
tensile strength than composite, fibers four times the size which was attributed to the
crystallinity of the nylon (Figure 1.15). Carbon black has also been encapsulated into
electrospun fibers to increase either electrical conductivity or the nanofiber modulus, 159 and
to sense strain. 160 Nanoparticles such as calcium carbonate are also being used for
biomedical applications such as bone regeneration. 161 The direct dispersion of graphene
21

prior to electrospinning has been used to enhance the optical absorption of poly(vinyl
acetate) (PVAc) by a factor of ten. 162 Quantum dots (QDs) are also of interest due to their
interesting optical properties because of a quantum confinement effect. In fact, the high-
voltage source in electrospinning can actually effect the passivation of the QD and suppress
deep-level emission. If the voltage is increased high enough, it may be able to align the ZnO
QDs and lead to ultraviolet photoluminescence. 163 Aggregation is a problem with these
systems and the addition of surfactant has been shown to allow a uniform distribution for
CdTe QDs in a PVP matrix. 164 Cellulose nanocrystals have also been electrospun with a
PVA nanofibers and were found to significantly increase the elastic modulus while acting to
reinforce the nanofibers. 165 Additionally, superhydrophobicity has been created by
electorspinning epoxy-siloxane modified SiO2 nanoparticles with a diameter of 200 nm in a
PVDF matrix. 166 Metal nanoparticles have also been used as localized heat sources inside
PEO fibers when irradiated with a laser tuned to the surface plasmon resonance of the noble
metal nanoparticle and complete melting was observed. 167
Although it has been determined that nanoparticles located in the interior of a
polymer fiber can show reduced activity, there is also great concern about the toxicity and
release of the nanoparticles into the environment. For this reason, the goal of many scientists
is to increase the functionality of a nanofiber mat without toxicity, thus dispersing
nanoparticles on the interior of the fiber surface has been achieved. Electrospun fibers with
noble metal nanoparticles such as gold, 168 silver, 169 and manganese acetate 170 have been
produced to attempt to achieve this result. Silver nanoparticles can impart antimicrobial
22

properties to polymer fibers 171 since they are positively charged and attract electronegative
bacteria. Gold nanomaterials, on the other hand, exhibit strong surface plasmon resonances
in the visible electromagnetic spectrum as a consequence of optically-driven coherent
oscillations of conduction electrons. 172 These resonances give rise to characteristic optical
absorption and scattering spectra, 173 yielding brightly colored, nanoparticle-containing
suspensions (Figure 1.16). Palladium nanoparticles have specific catalytic activity in
hydrogenation of dienes and olefins, and has been electrospun with PAN. 174 Magnetic
particles such as iron oxide have also been blended in electrospun polymer fibers 175 and have
also been found to orient parallel to the fiber axis, deflect under a magnetic field, and
improve the mechanical properties of the fibers. The formation of nanocomposites can even
change the inherent properties of the filler material; core-shell Fe/FeO nanoparticles have
been found to increase in shell thickness by a factor of 7.4% when electrospun. 176
1.1.5 Applications
The main attributes of electropun fiber webs are: high surface area to volume ratio,
porosity, and low mass. Many applications exist based on the ability to functionalize the
nanofibers as discussed earlier in this chapter. The types of applications and the required
fiber characteristics will be discussed in this section.
Tissue Engineering
Electrospun fiber mats are applicable for regenerative tissue due to the similarity to
the extracellular matrix (ECM), which provides the structural support to the cells in tissue
and organs. 177,178 Biocompatabile and biodegradable polymers can be used to make an
23

artifical ECM and seed stem or human cells onto it. When comparing cell proliferation on
both fiber and film of the same polymer, the fibers had a positive effect on cell growth. 179
Some challenges that exist, however, include the creation of a smooth fiber surface for cell
growth, 178 which can be created by altering the processing parameters. Porosity is another
important challenge, as it can cause the cells to create bridges over the pores in a fiber
scaffold. 180 Different tissues have drastically different tensile strengths, elastic moduli, and
elongational strength (i.e. skin vs. cartilage) and this has to be taken into consideration when
choosing a polymer. The polymer choice must allow the anchorage, migration, and
proliferation of cells in addition to matching the mechanical properties of the target tissue
and include many of the natural and synthetic biopolymers listed earlier.
Polymer blends are also utilized in order to combine properties; i.e. biodegradability
with mechanical robustness. 181 Additives such as calcium carbonate, phosphate, and
hydroxyapatite can increase functionality for tissue such as bone. 182 The types of cells that
have been seeded include mesenchymal stem cells, endothelial cells, neural stem cells,
fibroblasts, osteoblasts, etc. 103 One challenge that still remains is the creation of a synthetic
ECM as a hierarchical complex structure, to be fully functional the structure and growth of
arteries, 183 etc. may also need to be taken into account. The physical properties of the mat,
such as using aligned fibers, has been shown to guide the growth and orientation of
neurons. 184 A three-dimensional assembly of aligned, mineralized fibers 185 has been created
by Teo, et al. to mimic the hierarchical structure of bone and an image of these bundles can
be seen in Figure 1.17.
24

Catalysis
An important step in catalysis is the removal of catalyst after the reaction, to avoid
contamination in the end product. If catalyst are encapsulated within a nanofiber matrix, this
could be alleviate that problem by either dipping the nanofiber mat into the reaction solution
or have the solution pass through it in a reactor. 186 Polymers can be electrospun with
encapsulated monometallic or bimetallic nanoparticles such as Rh, Pt, Pd, Pd/Pt, or Rh/Pd for
catalysis in hydrogenation reactions. 174 In contrast, nanoparticles can be formed in vivo by
electrospinning solutions containing metal salts and reducing these agents at high
temperatures. Core-shell fibers created by the TUFT method for homogeneous catalysis are
able to achieve conversion in shorter times relative to conventional catalysts and can be used
several times without a loss of activity. 187 Carbon fibers decorated with Pt nanoparticles
were fabricated for electrocatalysis and had excellent activity and stability towards the
oxidation of methanol. 188 This demonstrated that carbon fibers loaded with noble metal
catalysts are advantageous due to the high conductivity and close contact between carbon
matrix/Pt particle, careful control over particle dispersion would be necessary to enhance this
application.
Drug Release
Electrospun fibers for drug delivery are advantageous for the ability to control the
release rate, easy production and polymers have little influence on the carrier. 189 In order to
be successful in this application, nanofiber mats need to fulfill several requirements. First,
they need to possess some robustness in order to prevent the drug from dissolving before it
25

eaches its target and allow for a controlled release if possible. The drug release could be
triggered by a stimulus and complete its release by a certain point. Nanoparticles of lipids
and biopolymers have been investigated for the purposes of drug transport and release. 190
Nanomaterials are also be used for locoregional therapy for immediate release to a target area
and is applicable for wound healing or tissue engineering. Nanofibers loaded with iron oxide
nanoparticles can be subjected to an external magnetic field which in turn causes heating of
the nanoparticles. This heat can be the stimuli needed for drug release. 76,191 A controlled
release is essential for many therapies and this is a functionality that many traditional
nanofbiers lack. By utilizing core-shell fibers the core can immobilize the drugs while the
shell controls the release rate out of the fibers. 131 Core-shell fibers can also be loaded with
proteins, enzymes, and growth factors for a controlled release. 141,192
Water & Air Filtration
Electrospun fiber mats are currently being investigated as inexpensive, low-energy
filtration membranes in order to provide clean water. The high volume to surface area and
nanosize pores make them ideal to remove contaminants like particles, organisms, and
hazardous chemicals from water. While the mechanical stability can be tuned with the
choice of polymer, the question still remains whether it can withstand a high number of water
cycles and still be effective. Recent reports of an electrospun nanofiber mat being used for
microfiltration have demonstrated that 99% of all particles that passed through were collected
and no failure was observed. 68,193 For the removal of heavy metals, the addition of heavy
metals on the surface, for example, NH3+ functional groups, 194 carboxylic acid, 195
26

poly(methacrylic acid), 73 has a higher success rate than encapsulated materials due to a
reduced activity. These membranes can also be imparted with antimicrobial properties
through the incorporation of silver nanoparticles, 196 sulfonated phenol groups, 197 and n-
halamine. 198 Water filters consisting of boehmite nanoparticles impregnated in electrospun
fibers was found to have an excellent removal efficiency of cadmium (II) from water. 199
Functionalizing the surface of fiber mats may be as simple as exposure to a basic solution;
antimicrobial properties have been found for PAN fibers with amidoxime groups when
exposed to a NH2OH solution. 200
Air filtration covers everything from dust to viruses to warfare stimulants. Nanofiber
membranes are already in use in air filtration media by several companies 201 based on their
high efficiency, ability to trap smaller particulates, and large surface area. Smaller fibers do
have a larger pressure drop when compared with macroscale fibers, which is one downside.
Much like water filtration, air filters can be chemically functionalized for detoxification
purposes against neurotoxins. 202 Likewise, ceramic fibers such as zinc titanate had an
excellent decomposition efficacy against nerve agents. 203 The future of filtering technology
looks to be in functional, antimicrobial and detoxifying membranes but it will be necessary to
absolutely determine that none of the incorporated additives leach out of the nanofibers
overtime.
Sensors
Detecting contaminants in the air can be a key feature in environmental strategy.
There are many air pollutants that are both low in concentration and low in size and
27

nanofiber template sensors would be greatly applicable for both of these characteristics. For
liquid applications, trace amounts of contaminants such as pharmaceuticals, cosmetics,
hormones, etc are possible carcinogens and can cause health problems. 204 Molecularly
imprinted nanoparticles in polymer fibers can detect very small amounts of propranolol in tap
water, 113 even though they are not on the surface. This same idea could be applied to
detecting other pharmaceuticals in waste water. Nanofiber mats can also be used as gas
sensors in areas with high air pollution. These are often ceramic materials, molybdic oxide
and tungsten oxide fibers have been shown to be successful in detecting NO2 and NH3 in
air. 155 Nanofiber sensors are superior to other methodologies due to their fast response time
and detection limit sensitivity.
1.2 One-Dimensional Magnetic Nanostructures
As technology advances, electronic devices are getting both smaller and lighter. In
order to realize the concept of nanodevices, a replacement has to be found for electrical
circuits and electrical wiring. One candidate for nanowires is actually deoxyribonucleic acid
(DNA); it has excellent self-assembly and has the potential for nanoscale patterning.
Conductivity can be imparted by replacing the imino proton on each base pair with a Zn 2+
ion. 205 Ferromagnetism has also been incorporated when DNA is in the presence of Cu 2+
ions, the spacing between the base pairs mimicks a ferromagnetic material. 206 Any polymer
with a saturated backbone of pendant metal complex has the potential to be of use in memory
applications.
28

One-dimensional nanostructures have very unique properties that act on two dimensions;
properties of the individual nanoparticles but also anisotropic properties that exist
collectively. A multitude of properties can exist when the unidirectional arrangements of
these materials exist within an organic matrix. The presence of an organic matrix helps these
materials, which usually tend to self-aggregate, keep their shape and approve their stability.
These are of interest for sensing, magnetomechanical systems, spintronic devices, and data
storage systems. When an external magnetic field is applied to metal-containing atoms such
as Co, Ni, and Fe, the spins of unpaired electrons become aligned. The term ‘magnetism’
refers to how materials act on a molecular level to an external magnetic field. The magnetic
susceptibility, �, is the effectiveness of an applied field to induce a magnetic dipole. It is a
ratio of the induced magnetization, M and the applied field: � = M/H. The behavior of
materials in a magnetic field can be broken up into the following: 207
A) Diamagnetism: This is a basic property of all materials which are weakly repulsed by
a magnetic field.
B) Paramagnetism: There are unpaired electrons in this material whose magnetic
moments will align in the same direction as the applied field.
C) Ferromagnetism: The electron spins are coupled into a parallel alignment that is
maintained by thousands of atoms in magnetic domains. The number of domains
depends on the size of the material; materials below a critical diameter of 14nm
consist of just a single domain. Above this critical diameter, the material will show
remanence or a persistent magnetization where some directional domains still exist
29

even when the magnetic field is changed and coercivity, or the need for a magnetic
field in the opposite direction to demagnetize the material. The bulk property for
larger materials is that many magnetic domains where the spins point in the same
direction and act cooperatively.
D) Antiferromagnetism: In this type of substance the electron spins of atoms are fixed in
an antiparallel alignment. These materials have a zero net magnetic moment.
E) Ferrimagnetism: A strong net dipole is present due to a range of spin moments
aligned in an antiparallel direction. The critical diameter of a single domain of
magnetite, a ferromagnetic material, is 128 nm.
F) Superparamagnetism: This occurs when the size of a ferro- or ferrimagnetic particle
is below the critical diameter and a stable magnetization is not possible due to
Brownian rotation, or thermal fluctuations, of the material. The particles mimic that
of a paramagnetic spin with a much higher moment. A critical temperature, called the
blocking temperature Tb, exists where the dipole moments are able to align and
couple to form a larger collective magnetization. Above this temperature, thermal
fluctuations are the dominant force. These materials usually do not show any
coercivity or remanence magnetization at room temperature.
The reason scientists want to align magnetic materials is due to the shape anisotropic
properties that exist which can lead to some interesting magnetic properties. Crystal
anisotropy is due to the shape of the material; it is easier to magnetize rods or cylinder-
shaped nanomaterials along the long axis as opposed to the short one. Some of the ways
30

magnetic unidirectional materials are being fabricated include a template-based approach, 208
channels in solids, 209 template through block copolymer micro-phase separation, 210
cylindrical polymer brushes, 211 biological 1-D templates such as bacterial chains, 212 (seen in
Figure 1.18) and electrospinning. 213 It is conducive to utilize an organometallic matrix
because even very small concentrations of magnetic materials, i.e. less than 1%, can impart
magnetic properties to the material as a whole. Multi-segmented alloys can even act as tiny
magnets in solution by magnetizing only one of its components, as in the case of nanowires,
and leading to self-assembled nanowires in an unidirectional arrangement. 214
Superparamagnetic nanoparticles have been studied as possible contrast agents in MRI.
Nanomaterials would offer improved sensitivity over the currently used contrast agents;
chains of iron oxide nanoparticles within a biopolymer layer have the ability to circulate,
target, and image tumors. 215 Although toxicity in the human body is a clear concern, it has
been demonstrated in rats that these nanomaterials can be quickly expelled from the brain
area into the circulatory system. 215,216 Harnessing spatial positioning over unidirectional
magnetic nanoparticles can lead to multifunctional materials in a variety of applications.
1.3 Organometallic Polymers
1.3.1 Overview
The first chapter focused on different ways to impart functionality to polymer fibers
and some of the drawbacks of these approaches while the second chapter investigated the
benefits of utilizing unidirectional magnetic materials. What if instead of having a
multicomponent system, it was possible to have a metal in the polymer center which still
31

imparted interesting properties while maintaining conventional polymer processing?
Organometallic polymers make this possible, these polymer contain a transition metal in the
main chain or more specifically has a metal-carbon σ or π bond. 217 The term ‘transition
metal’ refers to any element that has an incomplete d sub-shell or which can give rise to
cations with an incomplete d sub-shell. 218 The properties that make these transition metals
interesting, however is the ability to have multiple oxidation states which can be easily
controlled through electric currents. The first organometallic polymer was synthesized in
1955 as poly(vinylferrocene) and displayed reversible oxidation-reduction properties. 219 The
incorporation of metallic elements into polymer systems would allow different coordination
numbers and geometries and thus supply fascinating magnetic, optical, or catalytic
properties. Structurally, metallopolymers can be broken down into three categories: metals
incorporated directly into the polymer chain, π or σ-coordinated metals, and metallic moieties
pendant to the polymer backbone or in the side chains. 220 Some challenges that have existed
in the synthesis of organometallics include low molecular weights, oligomers, impurities, and
insolubility.
Other high-yield organometallic polymers are often synthesized through a step-
growth polymerization mechanism to create coordination chain polymers. Coordination
polyelectrolytes have been synthesized through the self-assembly of bisterpyridines and
metal ions where the polymers are water soluble and whose optical properties scan the visible
spectrum. 221 The first bimetallic polymer was synthesized containing both ruthenium and
silver and contains a ligand as a spacer. 222 Although organometallic polymers have been
32

synthesized through this coordination step-growth polymerization, there is often little control
over the final molecular weight or polydispersity. A more precise approach ca be achieved
through coupling two polymer blocks through a living organic polymerization and metal
coordination interactions. 223 A multitude of block copolymers can be created with a single
metal atom joining the two blocks as reported by Shubert et al. 224 Though this synthetic
method, one of the block components can be etched out leaving a regularly-ordered metal-
polymer block copolymer thin film as seen in Figure 5.19. It is even possible to create
metallogels, utilizing ligands and silver salts it is possible to form cages of coordination
complexes. Other favorable properties possible from metallogels include fluorescence
enhancement, 225 catalysis, 226 templating, 227 and electrical actuators. 228 The incorporation of a
metal atom into the main chain of a polymer leads to traditional polymeric properties while
maintaining the functionalities of metals including redox, magnetic, catalytic, and optical
activity.
1.3.2 Poly(ferrocenylsilanes)
In the early 1990’s, an organometallic polymer was fabricated by Manners et al. and
this was poly(ferrocenylsilanes) (PFS) with a high molecular weight and narrow
polydispersity through a ring-opening polymerization method 229 and this synthesis technique
has since yielded other strained monomers including bridging elements such as germanium,
tin, and phosphous. 230 PFS can be tuned to have either a semi-crystalline or amorphous
configuration due to the constituent groups on the silicon atom 231 where symmetrically
substituted constituents, R=R’=Me, impart crystallinity. The prevalence of iron in the main
33

chain conveys some interesting properties not present in non-metal containing polymers such
as redox-activity due to the reversible (Fe(II)/Fe(III) couple) 218 , the ability to be pyrolyzed
into a magnetic ceramic, 232 and semi- and photo-conductivity. 233 Gelable PFS derivatives
have been electrospun and crosslinked, and strain-induced buckling on electroactuation was
found to occur. 234
Block copolymers are macromolecules containing two or more types of repeat units
within long contiguous sequences, or “blocks,” of the same unit. These sequences are
covalently linked to form a single macromolecule. They can spontaneously self-organize or
microphase separate into a variety of ordered nanoscale morphologies such as lamellae,
hexagonally packed cylinders, spherical micelles on a body-centered-cubic lattice, or
complex bicontinuous morphologies like the gyroid. The type of morphology exhibited by
the block copolymer depends on the chemical attributes and lengths of the blocks, the
molecule’s architecture, temperature and overall chain length as well as presence of an
additive like a solvent, homopolymer or another block copolymer. 235 Due to their ability to
microphase separate, block copolymers constitute a versatile platform for a number of
existing and emerging technologies such as adhesives, membranes, drug delivery,
biomaterials and lithography. 236
The formation of metal containing block copolymers was possible due to the controllability
of the anionic ring opening polymerization of silicon-bridged ferrocenes. 237 Polymerization
must occur through sequential addition of monomers with decreasing end-group reactivity
such that PS ~ PI > PFS > poly(dimethylsiloxane) (PDMS) 237 . The first two PFS-containing
34

lock copolymers were PS-b-PFS and PFS-b-PDMS 238 and have allowed the synthesis of
block copolymers of PFS and PI, 239 poly(methylmetacrylate) (PMMA), 240 P2VP,
poly(ferrocenylethylmethylsilane) (PFEMS), 241 and the hybridization with polypeptides. 242
PFS block copolymers have been shown to self-assemble in the solid state and have
demonstrated spherical, cylindrical, lamellae, and gyroid morphologies. 243 The lengths of
poly(ferrodimethylsilane) PFDMS block copolymer micelles can elongate through the
addition of unimers which act to couple existing micelles. 244 Triblock copolymers can be
selectively etched so that only PFS is remaining as seen in Figure 3 which gives rise to many
lithographic applications. In the solid state, amorphous PFS with unsymmetrical constituents
on the silicon atom is utilized in order to allow phase separation to occur. Co-micelles with a
scarf-shaped architecture (Figure 1.20) have likewise been prepared with platelet micelles as
intiators. When a second PFS block copolymer is added, cylindrical micelle tassles are
grown from the already present micelles. 245 Segregated micelles can also be used as a
template for the formation of metallic particle and has been achieved with gold nanoparticles,
PbS quantum dots, and titania. 246 Banded light-emitting barcode structures with fluorescent
segments of a triblock-containing PFS were synthesized separated by nonemissive segments
of a diblock. 247
The periodic PFS domains can be converted to iron-rich clusters within a ceramic
domain by pyrolysis at temperatures about 600°C which can lead to the growth of single-
walled carbon nanotubes for soft lithography, 248 the formation of magnetic ceramics, 249 and
PFS-derived catalysts. 250 Solution self-assembly of diblock copolymers occurs when a block
35

selective solvent is utilized in order to induce segregation of the solvent incompatible ‘core’
surrounded by the solvent compatible ‘corona.’ Micelles of PFS-containing block
copolymers have produced micellar morphologies such as cylinders, 251,252 tubes, 252 fibers, 253
and tapes. 239 PFS experiments in the bulk utilize PFDMS, whose crystalline nature is
responsible for the growth of micelles. 254 Initial experiments with PFDMS-b-PDMS (with a
block ratio of 6:1 respectively) in hexanes demonstrated that cylinders with a PFS core and
PDMS corona were formed. 251 Imaging these organometallic, self-assembled structures with
transmission electron microscopy (TEM) is easy due to the contrast already inherent in the
sample owing to the iron-rich core. In order to form a core of PFDMS in cylindrical micelles
it is necessary to have a corona:core ratio of at least 5:1 239,251 but when this ratio reaches 12:1
it is more common for hollow, tube-like structures to form. 240,252 The length of the
cylindrical micelles can also be easily controlled; the addition of a small amount of a
common solvent will allow growth of the micelles and ultrasound or high temperatures will
cleave longer cylindrical micelles. The self-assembly of PFDMS-b-P2VP was investigated
in several alcoholic solvents and it was noted that cylindrical micelles were formed in
isopropanol but not ethanol. 255 The reason for this was attributed to the fact that isopropanol
is a better solvent than ethanol based on solubility parameters for PFDMS, which gives the
PFS chains time to rearrange and crystallize more easily. PFS-b-PI formed cylindrical
micelles in a PI selective solvent and the vinyl groups in the PI corona were utilized to
conduct a Pt(0)-catalyzed shell crosslinking reaction which was a precursor to making PFS
nanoceramics by pyrolysis with shape retention. 256 Nanocomposite self-assembled structures
36

consisting of PFS-b-poly(vinylmethylsiloxane) (PVMS) wormlike micelles reacted with
Ag[PF6] to create a one-dimensional array of silver nanoparticles encapsulated within the
worm-like micelles 257 allowing for highly oriented arrays of nanoparticles to be formed.
37

Nomenclature
AFM Atomic force microscopy
BSA Bovine Serum Albumin
CNT Carbon nanotube
DMF Dimethylformamide
DNA Deoxyribonucleic acid
ECM Extracellular matrix
HCL Hydrochloric acid
HPC Hydroxypropyl cellulose
MMT Montmoillonite
PA Poly(amide)
PAA Poly(acrylic acid)
PAN Poly(acrylonitrile)
PCL Poly(caprolactone)
PDMS Poly(dimethoxysilane)
PFDMS Poly(ferrodimethylsilane)
PFEMS Poly(ferroethylmethylsilane)
PFS Poly(ferrocenylsilane)
PGA Poly(glycolide)
PI Poly(isoprene)
PLA Poly(lactic acid)
PLGA Poly(lactide)-co-(glycolide)
PMMA Poly(methyl methacrylate)
PPX Poly(p-xylxlene)
PS Poly(styrene)
PU Poly(urethane)
PVA Poly(vinyl alcohol)
PVAc Poly(vinyl acetate)
P2VP Poly(2-vinyl pyridine)
PVDF Poly(vinylidene fluoride)
PVMS Poly(vinylmethylsiloxane)
PVP Poly(vinylpyrrolidone)
QD Quantum dot
SEM Scanning electron microscopy
TEM Transmission electron microscopy
TUFT Tubes by fiber templates
UV Ultraviolet
38

Figures
Figure 1.1. Schematic figure of basic electrospinning setup. 103
39

Figure 1.2. Photograph of the whipping motion of the instability region of the polymer jet
during the electrospinning process. 12
40

A.
Figure 1.3. A method to produce aligned fibers. A) a schematic utilizing the parallel
grounded electrode collection system and B) the resulting aligned polymer fibers. 25
B.
41

Figure 1.4. SEM micrograph of beaded PEO fibers.
42

Figure 1.5. Reduction in bead density upon increase in polymer solution viscosity. 258
43

Figure 1.6. Ribbon-like fibers formed via electrospinning. 38
44

Figure 1.7. SEM image of anatase hollow fibers created via a co-axial electrospinning
setup. 39
45

Figure 1.8. TEM of PPX/Pd TUFT hybrid nanotubes after the pyrolysis of PLA template
fibers and inset is an electron diffraction pattern of Pd crystals. 46
46

Figure 1.9. SEM micrograph of porous PLA fibers obtained via electrospinning and
subsequent swelling. 259
47

Figure 1.10. SEM iamges of anatase nanofibers whose surfaces have been decorated with a)
gold and b) silver nanoparticles via photocatalytic reduction. 63
48

Figure 1.11. A list of organosoluble polymers and their molecular structure. 103
49

Figure 1.12. Schematic of polymer/inorganic composite nanofibers when a) inorganic ions
are incorporated into electrospun fibers followed by exposure to gas to synthesize inorganic
nanoparticles both inside and outside of the nanofiber and b) when only the surface of
nanofibers are modified with metal ions. 123
52

a
53
100 nm
c
100 nm
Figure 1.13. TEM micrographs of multicomponent polymer electrospun fibers
demonstrating a) A core-sheath structure formed by a polymer blend 128 b) Lamellar structure
formed by a phase-separated block copolymer 260 c) Cylindrical structure formed by a phaseseparated
block copolymer. 260

Figure 1.14. TEM image of a) PAN/CNT composite nanofiber mat and b) demonstrating the
uniform distribution and alignment of CNTs witin a PAN fiber. 56
54

Figure 1.15. Average Young’s modulus for electrosopun nylon-6 and nylon-6/O-MMT
nanocomposite single fibers vs. fiber diameter. 158
55

Figure 1.16. Demonstration that as the aspect ratio of gold nanorods increases, as does the
maximum optical absorbance and thus the color of the aqueous colloidal suspension. 261
56

Figure 1.17. Three-dimensional mineralized electrospun fibers mimicking the hierarchical
structure of bone. 185
57

Figure 1.18. Left: magnetic induction map from two pairs of bacterial magnetite chains.
Right: A bright-field TEM image of a double chain of magnetite magnetosomes. 212
58

Figure 1.19. Pyroloysis of UV cross-linked PS-b-PFEMS films with a) height-mode
scanning force microscopy, b) phase-mode, c) TEM images, and d) A schematic of the
morphology. 262 Inset scale bars = 50 nm.
59

Figure 1.20. TEM micrographs of scarf-shaped PI-b-PFS co-micelles (scale bar = 500
nm). 245
60

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75

Abstract
CHAPTER II
Long-Range Alignment of Gold Nanorods in Electrospun Polymer
Nano/Microfibers*
In this study, a scalable fabrication technique for controlling and maintaining the
nanoscale orientation of gold nanorods (GNRs) with long-range macroscale order has been
achieved through electrospinning. The volume fraction of GNRs with an average aspect ratio
of 3.1 is varied from 0.006 to 0.045 in aqueous poly(ethylene oxide) solutions to generate
electrospun fibers possessing different GNR concentrations and measuring 40-3000 nm in
diameter. The GNRs within these fibers exhibit excellent alignment with their longitudinal
axis parallel to the fiber axis n. According to microscopy analysis, the average deviant angle
between the GNR axis and n increases modestly from 3.8 to 13.3° as the fiber diameter
increases. Complementary electron diffraction measurements confirm preferred orientation
of the {100} and {111} GNR planes. Optical absorbance spectroscopy measurements reveal
that the longitudinal surface plasmon resonance bands of the aligned GNRs depend on the
polarization angle and that maximum absorption occurs when the polarization is parallel to n.
*This chapter has been published in its entirety:
KE Roskov, KA Kozek, WC Wu, RK Chhetri, AL Oldenburg, RJ Spontak, JB Tracy. “Long-
Range Alignment of Gold Nanorods in Electrospun Polymer Nano/Microfibers.” Langmuir. Web
publication date: 8-11-11, DOI # 10.1021/la202106
76

Nanoparticle (NP) synthesis, characterization and self-assembly have been thoroughly
investigated over the past two decades primarily because of their novel size- and shape-
tunable functional properties, 1 as well as the wide variety of property enhancements they
impart to matrix materials, such as polymers. 2 As their size decreases, metal NPs can exhibit
physical and chemical properties (e.g., catalytic, 3 optical, 4 electronic, 5 and magnetic 6 ) that are
not observed in the bulk. Gold and silver NPs, for instance, are known to exhibit strong
surface plasmon resonances in the visible electromagnetic spectrum as a consequence of
optically-driven coherent oscillations of conduction electrons. 7 These resonances give rise to
characteristic optical absorption and scattering spectra, 8 yielding brightly colored, NP-
containing suspensions. Gold nanorods 9 (GNRs) have recently attracted considerable
attention due to their strong and uniquely tunable longitudinal surface plasmon resonance
(LSPR) along the longitudinal GNR axis. The LSPR wavelength of GNRs can be adjusted
from ~520 nm in the case of low-aspect-ratio (i.e., spherical) gold NPs to at least 1750 nm in
the near-infrared region of the electromagnetic spectrum for high-aspect-ratio GNRs. 10
Currently, GNRs with a LSPR at 800 nm are of particularly significant interest for in vivo
imaging of biological systems, because blood and tissue exhibit an absorbance minimum at
this wavelength. 11
When dispersed in a solvent, GNRs are randomly oriented, and they tend to remain so
when physically deposited on surfaces. If the distribution of GNR sizes and shapes is
sufficiently narrow so that the GNRs can be considered nearly monodisperse, they can
spontaneously self-assemble into ordered grains with different orientations. 12 Presently,
77

alignment of such grains over macroscale dimensions is a significant technological challenge,
but is highly desirable, 13 because the optical anisotropy 14 of GNRs at the LSPR can be
greater than ~250:1. 15 A scalable method for controlling and maintaining the orientation of
GNRs at the nanoscale, while fabricating macroscale structures that exploit GNR orientation,
is therefore needed for creating nanostructured composites with tunable, anisotropic optical
properties. Alignment-enhanced absorption of polymer/GNR composites has been previously
reported using different preparation strategies: thin-film stretching, 16 block copolymer-
templated organization, 17 directed nanoscale assembly, 18 ultrathin film confinement, 19 and
polymer fiber coating. 20 An ongoing challenge for polymer/GNR composites produced by
these methods is the consistent acquisition of long-range, scalable order of highly-aligned
GNRs along a common axis (rather than randomly oriented in a common plane). Here, we
describe a general technique for aligning GNRs within electrospun polymer nano/microfibers
with diameters ranging from 40 to 3000 nm. Our approach enables the hierarchical alignment
of fibers containing aligned nanorods over large, macroscopic dimensions. The resultant
polymer/GNR composites were characterized by electron microscopy, electron diffraction,
and optical absorbance spectroscopy.
The GNRs investigated here were synthesized by modifying a method originally
introduced by Nikoobakht et al. 21 Details of the GNR synthesis are provided as Supporting
Information. A custom-built electrospinning unit with an Al collection target was operated at
an electric potential of 10 kV with a plate distance of 15 cm and solution flow rates of 30-50
�L/min. Aligned electrospun fibers, rather than conventional, randomly oriented fiber mats,
78

were generated by electrospinning between two grounded electrodes separated by 2 cm. 22
Prior to electrospinning, the GNR stock suspension was warmed with a heat gun to
redissolve precipitated cetyltrimethylammonium bromide (CTAB). Poly(ethylene oxide)
(PEO: 1000 kDa, Sigma-Aldrich) was selected for its intrinsic hydrophilicity and, by
inference, its compatibility with the native CTAB coating on the GNRs. A high-molecular-
weight grade was chosen to facilitate fiber formation and impart mechanical robustness (for
handling purposes). While other matrix materials have not yet been explored, we anticipate
similar results would be obtained for comparably water-soluble polymers amenable to
electrospinning. The PEO was dissolved directly into the GNR stock suspension to yield a
maximum GNR volume fraction (�) of 0.045. Lower GNR volume fractions were
subsequently prepared by first dissolving the polymer in a separate deionized water solution
and then adding it to the GNR suspension. Polymer concentrations ranged between 3.0 and
4.5 wt % PEO. Immediately before electrospinning, the mixture was sonicated for 10 min to
ensure a uniform dispersion. Electrospun fibers were dried under vacuum for 24 h at ambient
temperature.
Specimens imaged by transmission electron microscopy (TEM) consisted of randomly
oriented fibers electrospun directly onto carbon-coated Cu grids. Images and selected-area
electron diffraction (SAED) patterns were acquired on a field-emission Hitachi HF2000
microscope operated at 200 kV. Corresponding scanning electron microscopy (SEM) images
were collected from specimens sputter-coated with 6 nm of Au/Pd on a JEOL 6400F field-
emission microscope operated at 5 kV. For optical characterization, a custom-built polarized
79

UV-vis spectrophotometer, housing a fiber-coupled white light source (Photon Control),
broadband linear polarizer, light collection fiber coupled into a USB spectrometer
(Lightspeed Technologies), and a series of collimation and collection lenses, was employed.
Absorbance spectra were acquired by orienting each specimen normal to the collimated white
light beam, which was passed through a ~2 mm diameter aperture and the polarizer.
The GNRs employed in this study have an average width and length of 17 ± 6 and 49 ± 10
nm, respectively (Figure 1), thereby yielding an average aspect (length/diameter) ratio of 3.1.
Electrospinning is becoming established as an important method for preparing polymer
nano/microfibers, because it is straightforward to perform and offers the flexibility to tune
fiber characteristics by controllably varying solution and/or processing parameters. Aligned
fibers generated by electrospinning between two grounded electrodes yield a freestanding,
oriented mat measuring 10 cm wide and 3 cm or more long (Figure 2). Here and in
subsequent discussion, n refers to the fiber axis direction. The fiber thickness is governed by
the PEO concentration when all other electrospinning parameters are held constant. For
instance, the nanofibers shown in Figure 3a with an average diameter of ~50 nm were
electrospun from an aqueous suspension with � = 0.006 and a PEO concentration of 3.2 wt
%. At this relatively low GNR loading in PEO nanofibers, the GNRs orient with their
longitudinal axis parallel to n. To quantify the extent of alignment along n, the average
deviant angle (), determined from the angular difference between the GNR orientation
and n, has been measured for 150-300 GNRs embedded within electrospun PEO fibers of
varying thickness. For nanofibers such as those shown in Figure 3a, = 3.8°. Due to the
80

nanoscale diameter of the PEO fibers and their CTAB surface coating, 9 very few GNR
aggregates are observed at this low loading level, for which the interparticle distance is
typically much longer than the length of the GNR. The theory proposed by Bates and
Frenkel 23 for hard-rod fluids predicts that GNRs with aspect ratios < 7 should exhibit random
orientation in solution, insofar as the GNRs do not interact with each other over large
distances. According to our measurements, it follows that oriented, well-separated GNRs are
aligned by forces other than those arising from interparticle interactions.
During electrospinning, the viscous polymer/GNR solution at the tip of the syringe can be
modeled 24 as a fluid cone leading to a jet, where the charged polymer solution is emitted.
Sink-like flow emerges at the apex of the cone, and streamlines form due to the rapid
decrease in area. In suspensions, the GNRs are initially randomly oriented but begin to align
along the streamlines leading to the jet. Since the Reynolds number, defined as the ratio of
inertial (drag) to viscous forces, is much less than unity at this point, the translational
velocity component dominates, 25 in which case the center of each GNR experiences the same
velocity as the local fluid velocity. As the polymer solution leaves the jet, the GNRs are
expected to be oriented for the most part along the fiber spinning direction. Indeed, the
alignment of significantly longer carbon nanotubes 26 and CdS nanorods with an aspect ratio
of ~20 27 have been experimentally confirmed in electrospun fibers, but we are unaware of
prior studies of nanorods with shorter aspect ratios. When the GNR loading is increased to Φ
= 0.045 (Figure 3b), the GNRs remain highly oriented, because their orientation is
predominately dictated by the local velocity profile. For an increased fiber diameter of ~650
81

nm and Φ = 0.035 (Figure 3c), the velocity profile is forced to extend across the diameter of
the microfiber. Consequently, the GNR longitudinal axes remain highly aligned with n, but
increases modestly to 8.8°. When the fiber diameter is further increased to ~3000 nm
and Φ = 0.031 (Figure 3d), the degree of GNR alignment decreases, as verified by an
accompanying increase in to 13.3°.
Two other key morphological observations warrant mention. The first is that long GNRs
are more highly oriented with respect to n than are short ones. For example, in a fiber
measuring ~80 nm in diameter, when measurements of are sorted according to the
GNR length, the shortest 26% of the GNRs possess that is 54% greater than for
all GNRs in the fiber. Secondly, the fiber diameter is one of the main factors that determine
the degree of GNR alignment within electrospun fibers. That is, within fibers of nearly
constant diameter, the extent of alignment does not appear to depend strongly on Φ over the
range of GNR concentrations investigated. These two results imply that the degree of
nanoscale GNR orientation over macroscale dimensions is primarily controlled by the flow
field introduced by the polymer and experienced by the GNRs during electrospinning,
coupled with the GNR aspect ratio.
Selected-area electron diffraction (SAED) performed on a segment of microfiber
measuring 200 nm in diameter and containing ~20 GNRs yields the pattern included as an
inset in Figure 3b. Previous studies have shown that GNRs possess a faceted crystal structure
with a [110] growth direction, {111} end facets, and {100} side facets. 28 Analysis of the
SAED pattern in Figure 3b confirms the existence of a face-centered cubic lattice with
82

preferred orientations identified by the {100} and {111} reflections, which is consistent with
results reported elsewhere. 21 These reflections appear over a limited angular range rather than
as single spots, which indicates a distribution of GNR orientations. This result has been
analyzed in terms of the truncated Herman’s orientation function (P2), 29 written as
��
P2 � 3 cos2 � �1
. (1)
2
Here, � is the azimuthal angle extending from 0 to 2� around the circular SAED pattern, and
cos 2 � =
I(�)cos 2 � (�)d�
, (2)
� I(�)d�
where I(�) is the scattering intensity that varies along a circular trace of constant radius.
��
Limiting values of
��
P2 are 1.0 for perfect alignment along n, 0.0 for random dispersions and
�0.5 for perpendicular alignment relative to n. The present analysis yields P2 = 0.73 for the
aligned GNR-containing PEO microfibers under investigation, which further confirms that
the GNRs are highly oriented within the microfibers.
Gold nanorods are of particular contemporary interest because of their anisotropic optical
properties. For observing the optical anisotropy of an ensemble of GNRs, alignment of the
GNRs is required. Furthermore, oriented GNRs may have controllable end-to-end coupling
between their LSPRs (in contrast to the distribution of relative orientations found in isotropic
dispersions
of GNRs), 30 especially at high GNR loading levels. Several independent reports of the optical
properties of GNRs aligned with electric fields 31 or in stretched polymers 16 have established
that linearly polarized light oriented with the electric field parallel to the GNR long axis
83

excites the LSPR but does not excite the transverse surface plasmon. Here, we show similar
results for aligned GNRs in oriented electrospun PEO microfibers.
Optical absorbance spectra have been collected using the custom-built UV-vis
spectrophotometer described earlier. As a benchmark, spectra acquired at different
polarization angles from GNRs randomly dispersed in a PEO film measuring ~500 �m thick
are presented in Figure 4a and reveal the existence of clearly discernible and angle-
independent LSPR peaks near ~520 and ~850 nm. These signature features are likewise
observed for GNRs dispersed in water, confirming that the bulk, semicrystalline PEO has
little effect on GNR orientation. The absorption spectra of aligned GNRs in electrospun PEO
microfibers measuring 200 nm in diameter (Figure 4b) strongly depend on the polarization
angle. As expected, the LSPR peak at 804 nm is most pronounced when the polarizer is
parallel to n (and to the GNR longitudinal axis) at 0° but vanishes when the polarizer is
perpendicular to n at 90°. Similarly, the absorbance in the 500-600 nm region decreases
when the polarizer is parallel to n at 0° because the transverse surface plasmon is not excited.
The position of the LSPR peak indicates a redshift of 40 nm relative to a random GNR
dispersion in water (cf. Figure S-1 in the Supporting Information). The analogous LSPR peak
in the random dispersion of GNRs in a PEO film is centered at 847 nm (Figure 4a), which
corresponds to a further red shift of ~40 nm. Shifts in the LSPR wavelength may arise from
several sources, such as interparticle coupling, 32 differences between the refractive indices of
PEO and water, and the crystallinity of PEO.
84

The absorbance spectra in Figure 4b have been processed to remove background
contributions from the glass substrate and from optical scattering of the polymer fibers.
Spectra for the GNR-containing and control fiber specimens prior to subtraction are included
for examination in Figure S-2 of the Supporting Information. In contrast to polymer thin
films, polymer fibers contribute a scattering signal that is significantly greater than the
absorbance of the GNRs to extinction, which results in absorbance spectra from the oriented
GNRs that are noisier. All spectra are smoothed using a 17-point Savitzky-Golay numerical
procedure. 33 An aligned PEO microfiber mat without GNRs is selected as an appropriate
control specimen for Figure 4b. The difference in specimen density accompanying the
incorporation of GNRs is taken into account by Beer’s law:
A = ��z, (3)
where A denotes the absorbance, μ is the extinction coefficient and z is specimen thickness.
In Figure 4a, the absorbance spectra of randomly dispersed GNRs in a PEO thin film display
a minimum near 650 nm. It immediately follows that �GNR ≈ �c at this wavelength, thereby
yielding AGNR/zGNR ≈ Ac/zc, where the subscripted c represents the control specimen without
GNRs. The thicknesses of the control and GNR-containing specimens can therefore be
related by their absorbance values at 650 nm, and control spectra have been subtracted from
the spectrum for the GNR-containing fibers. After this correction, the spectra are normalized
to zero and unity. We note, however, that �GNR > �c at 650 nm, because the GNRs have a
small, but non-zero, absorbance at 650 nm. Consequently, the background correction in
Figure 4b does not completely remove spectral contributions originating from polymer fiber
85

scattering, which is responsible for the peak present at ~600-700 nm. The optical anisotropy
of the GNRs could be further increased by improving the alignment of the GNRs within
polymer microfibers and the parallel orientation of the microfibers.
In conclusion, we have demonstrated a scalable method for controlling and maintaining the
nanoscale orientation of GNRs with long-range macroscopic order over a distance of several
centimeters. Here, GNRs with an aspect ratio of 3.1 exhibit excellent alignment with their
longitudinal axes parallel to n for electrospun polymer nano/microfibers with diameters of
40-600 nm, and they maintain substantial alignment in microfibers measuring up to 3000 nm
in diameter. While fiber diameter is found to play a crucial role in GNR alignment, GNR
concentration can be varied with no discernible impact on the net degree of alignment.
Electron diffraction measurements of the aligned GNRs confirm preferred orientation of the
{100} and {111} GNR planes. Optical absorbance spectroscopy measurements performed on
microscopically aligned GNRs in macroscopically aligned electrospun fibers demonstrate
that the LSPR bands are polarization dependent and display maximum absorption when the
polarizer is parallel to n.
Acknowledgement.
This work was supported by the National Science Foundation (CBET-0967559 and a
Graduate Research Fellowship for K. E. R.) and startup funds from North Carolina State
University. We thank Prof. Jesse Jur for helpful discussions and Prof. Benjamin Wiley for
assistance with preliminary optical measurements.
86

Figures
Figure 2.1. TEM image of GNRs deposited from an aqueous suspension onto a carboncoated
TEM grid. The inset shows the distribution of measured aspect ratios of the GNRs,
which measure 49 nm long and 17 nm in diameter on average.
87

Figure 2.2. SEM image of macroscopically-aligned electrospun PEO fibers containing
GNRs.
88

Figure 2.3. Aligned GNRs in electrospun PEO nano/microfibers as functions of fiber
diameter and GNR volume fraction (��: (a) 40 nm and (�� = 0.006, (b) 50 nm and (�� =
0.045, (c) 650 nm and (���= 0.035, and (d) 3000 nm and (��= 0.031. A selected-area
electron diffraction pattern of the corresponding sample is included as an inset in (b).
89

Figure 2.4. Absorbance spectra for (a) randomly oriented GNRs in a PEO film
measuring ~500 �m thick at different polarization angles and (b) GNRs aligned within
electrospun PEO microfibers measuring ~200 nm in diameter at polarization angles varying
from 0° (parallel to the fiber axis n) to 90° (perpendicular to n). In both cases, the data are
color-coded and labeled in (a).
90

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Abstract
CHAPTER III
Magnetic Field-Induced Alignment of Nanoparticles in Electrospun
Microfibers
We report on the facile and switchable alignment of superparamagnetic iron oxide
nanoparticles measuring ~18 nm in diameter in electrospun microfibers. Application of a
magnetic field perpendicular to the electric field employed during electrospinning yields
polymeric microfibers with nanoparticles aligned in one-dimensional arrays, thereby
providing control over when and where the nanoparticles align. According to electron
microscopy, the length over which alignment is desired can be judiciously selected, thereby
making these nanomaterials excellent candidates for nanotechnologies requiring nanoscale
alignment on-demand. Concurrent alignment of the electrospun fibers using established
procedures provides a viable route to organic/inorganic materials possessing anisotropic
properties that reflect multiscale alignment.
Introduction
Previous efforts to achieve nanoparticle alignment in polymeric matrices have relied on a
variety of process strategies, such as three-dimensional superlattices, 1 deoxyribonucleic acid 2
or block copolymer 3 templating, surface nanolithography, 4 electrostatic desalting transition, 5
and surface chemical modification. 6 Electrospinning is a rapidly developing fabrication
technique that produces solid polymeric fibers with diameters ranging from several tens of
94

nanometers up to several microns, a high surface-area-to-volume ratio, and the potential for
porosity at multiple length scales. 7 As such, it provides an attractive approach to the
formation of nano/microscale polymeric fibers containing nanoparticles that are spatially
restricted. Further control over nanoparticle positioning can be achieved within such fibers
by post-crystallization of the polymer matrix, 8 the electric field employed during
electrospinning, 9 development of a polymeric nanostructure via self-assembly, 10 or strategic
use of coaxial electrospinning. 11 These methodologies and others, however, depend on the
intrinsic nature of the polymer matrix and/or the processing conditions associated with
electrospinning, and are not intrinsically switchable. Here, we postulate that control over the
location and alignment of nanoparticles within electrospun fibers can be achieved through the
straightforward use of an external electromagnetic field applied during electrospinning. As
with electrospinning, several design issues, such as the strength of the magnet and its
orientation/position with respect to the onset of the polymer jet at the Taylor cone, warrant
consideration. Unlike other approaches intended to align nanoparticles during
electrospinning, an external magnetic field can be applied and removed at any time, thereby
permitting switchable alignment that can even be pulsed. As a consequence of switchable
alignment, the magnetic properties of the nanoparticles can be correspondingly changed from
superparamagnetic to ferromagnetic, 12 which may allow for future read-write capability.
Existing nanotechnologies require aligned nanoparticle arrays in surface-enhanced Raman
scattering (SERS) substrates for use in ultrasensitive analytical tools, 13 micromechanical
95

sensors, 14 nanoscale barcodes, 15 and protein separation. 16 Prior studies have shown that
magnetic fields can be applied during the electrospinning of polymer/nanoparticle systems to
improve alignment of the fibers, but not of the nanoparticles within the fibers. Parallel-
positioned permanent magnets with field strengths ranging from 25 to 120 mT have been
placed, for instance, on the collector plate during the electrospinning of polymer fibers
containing Fe3O4 nanoparticles to align the electrospun fibers. 17 Deflection of entire fibrous
mats generated by electrospinning and containing magnetite nanoparticles in the presence of
a magnetic field has also been observed. 9 In addition, magnetic fields have also been used to
control the spatial arrangement of nanoscale objects within bulk, not electrospun, polymer
matrices. Magneto-polymer nanocomposite particles measuring 200 nm in diameter are, for
instance, reported 18 to align in hydrogel nanocomposites under a magnetic field of 1.5 T.
Even multi-wall carbon nanotubes dispersed in a monomer precursor are found 19 to align
when exposed to an external magnetic field during matrix polymerization. The use of an
external magnetic field has likewise been found to have an effect on biological systems by
organizing cell rods, seeded on magnetically susceptible fiber bundles, into three-
dimensional tissue constructs. 20 In the spirit of these prior observations, the objective of the
present work is to create one-dimensional, periodic arrays of magnetic nanoparticles in
electrospun polymer fibers by applying a magnetic field, in conjunction with an electric field,
during electrospinning.
96

Experimental
ε-Polycaprolactone (PCL) with a molecular weight of 80 kDa according to the
manufacturer was provided by Solvay (Warrington, England). Reagent-grade chloroform for
electrospinning was obtained from Fisher Scientific (Fairlawn, NJ). Iron oxide nanoparticles
with an average diameter of 17.6 nm were prepared by thermal decomposition of iron oleate
in the presence of oleic acid as the capping agent in high boiling hydrocarbons (docosane and
eicosane) according to the procedures described elsewhere. 21,22 X-ray diffraction data
confirmed that as-synthesized nanoparticles are crystalline and contain mostly wüstite (Fe(1-
x)O) and some spinel (most likely Fe3O4), 21 but at the same time, they are superparamagnetic
iron oxide nanoparticles (SPIONs) according to magnetic measurements. 23 The resultant
SPIONs were precipitated with a mixture of acetone and hexane and then separated by
centrifugation and suspended in chloroform at concentrations between 10 and 50 mg/mL.
Suspensions for electrospinning were prepared by first dissolving PCL in chloroform and
then adding appropriate amounts of the SPION suspension to yield 5 wt% PCL at three
different SPION concentrations: 0.5, 1.0 and 2.5 vol%.
The in-house electrospinning unit, operated at 8 kV, employed a syringe pump and an Al
collection target. The separation distance was 12 cm, and the solution flow rate varied from
10 to 35 �L/min. A horseshoe electromagnet (Edmund Scientific, Tonawanda, NY) was
connected to a 6 V battery to generate a 26 mT magnetic field, which corresponds to the most
uniform field as measured by a DC magnetometer (Alphalab Inc., Salt Lake City, UT). As
schematically depicted in Figure 3.1, the magnet was positioned parallel to and below the
97

syringe needle. Electrospun fibers were collected with and without the magnetic field applied
onto the Al plate, as well as onto carbon-coated transmission electron microscopy (TEM)
grids taped to the plate. To discern the spatial distribution of SPIONs within the PCL fibers,
TEM was performed on a field-emission Hitachi HF2000 microscope operated at 200 kV.
The magnetic behavior of aligned 7 fiber mats was examined on a Quantum Design MPMS
superconducting quantum interference device (SQUID) for fibers electrospun from PCL
dissolved in a 30 mg/mL SPION suspension. Magnetization curves were recorded at 300 K
from fiber mats possessing an average mass of 1.5 mg. The magnetizing field strength for
embedded samples ranged from �20 to +20 kOe, whereas saturation was reached at 6 kOe as
the field strength varied from �80 to 80 kOe for unembedded SPIONs.
Results and Discussion
The polymer solution concentration, separation distance and voltage used to electrospin
PCL fibers have been judiciously selected to yield microfibers with diameters ranging from
100 to 500 nm so that their interior morphologies could be interrogated by TEM. Field-
emission scanning electron microscopy (SEM) results (provided in the Supporting
Information) confirm that these PCL microfibers with and without SPIONs possess an
average diameter of ~300 nm. The microfibers exhibit slight dimpling on their surface, and
very few bead defects containing large SPION aggregates (most likely formed by
nanoparticles with insufficient surface functionalization in the suspension reservoir) develop.
In the absence of an applied magnetic field, SPIONs measuring 17.6 nm in diameter appear
to be randomly distributed throughout the PCL microfibers, as seen in the TEM images
98

provided in Figure 3.2 for microfibers differing in SPION concentration: 0.5 vol% (Figure
3.2a) and 2.5 vol% (Figure 3.2b). A TEM image of the SPIONs drop cast from suspension is
included in the inset of Figure 3.2a for reference. Since these SPIONs are coated with oleic
acid which is fairly compatible with PCL, relatively little aggregation is observed in the PCL
microfibers. Few small SPION aggregates, not evident in these images but found in other
specimens, appear as spheroidal clusters with no preferential orientation. Complementary
energy-dispersive spectroscopy analysis (data provided as Supporting Information) reveals
that few, if any, SPIONs reside on the surface of the microfibers.
An electromagnet is selected for the present study since it generates a magnetic field only
when an electric current is flowing, in which case the magnetic field can be immediately
terminated by stopping the current. Since the objective of this study is to create aligned
arrays of magnetic nanoparticles on-demand and since many types of nanoparticles with
various degrees of magnetic susceptibility exist, we have elected to use superparamagnetic
nanoparticles due to their ability to respond to an external magnetic field but remain
nonresponsive when the field is removed. 24 In the field-induced alignment experiments, the
degree to which SPIONs respond to an externally applied magnetic field governs the extent
to which alignment will occur. For a spherical particle, the magnetic moment (mp) is related
to particle size by 25
��
mp � �d 3
mMS
6
99
(1)

where dm is the particle diameter and MS denotes the saturation magnetization. On the basis
of Eq. 1, very small nanoparticles may have an insufficient magnetic moment to respond to
the applied magnetic field and undergo noticeable alignment. When an external magnetic
field is applied, however, attractive dipolar interactions arise between adjacent SPIONs.
Aggregates of SPIONs develop to maximize their magnetic moment in response to the
magnetic field and align along the magnetic field lines created by the electromagnet. Under
these conditions, the anisotropic field-induced magnetic dipolar interactions promote the
formation of highly elongated aggregates (arrays), rather than random distributions, of
SPIONs.
The syringe portrayed in Figure 3.1 is positioned within several centimeters of the open
end of the horseshoe electromagnet (where the magnetic field lines are the most
concentrated). The optimal distance is ~1.5 cm, which suggests that this position identifies
where the magnetic field is the most uniform. The distance between the poles of the magnet
is 2 cm, with one pole placed directly under the Taylor cone (reasons for which are discussed
later) and the other located 2 cm closer to the collection plate. Although the horseshoe shape
of the electromagnet increases the strength of the magnet, we recognize that the field is not
uniform. This issue is currently being addressed by using electromagnets varying in physical
shape and field uniformity. Another factor influencing the ability of the SPIONs to align in
electrospun PCL microfibers is their concentration in the suspension reservoir. The attractive
forces between two SPIONs whose dipole moments are aligned must be sufficient to
overcome the matrix viscosity for particle alignment to ensue, and the magnitude of such
100

forces depends on the interparticle distance, which relates to the particle concentration. 26 A
sensitivity analysis for 17.6 nm SPIONs reveals that a minimum SPION concentration of ca.
0.05 vol% is required for discernible alignment in the present study. At higher
concentrations, the interparticle distance decreases and greater attraction, or chaining,
between neighboring SPIONs occurs. The maximum SPION concentration is set by other
considerations, namely, suspension conductivity and viscosity, as well as undesirable SPION
aggregation.
Examples of aligned SPIONs measuring 17.6 nm in diameter in electrospun PCL
microfibers at a concentration of 0.5 vol% are presented in Figure 3.3. In all cases, the
electromagnet is connected first, followed by the high-voltage power supply for
electrospinning. The 6 V battery remains outside the electric field to prevent the possibility
of a short circuit. In images such as the one displayed in Figure 3.3a, the SPIONS remain
aligned in one-dimensional arrays parallel to the fiber surface over lengths exceeding 1.5 �m
insofar as the electromagnet remains active. These long arrays are visible in most microfibers
of a mat collected on a TEM grid, and they appear on multiple grids used to collect the same
sample. In Figure 3.3b (including the inset), the effect of the magnetic field is not constant
(most likely due to the reasons discussed below), which results in shortened, but nonetheless
highly aligned, SPION arrays of variable thickness. Although such arrays constitute the
predominant feature when the magnetic field is applied, some randomly distributed SPIONs
remain. Fine adjustment of both electrospinning conditions and solution properties has been
unable to resolve this lack of position uniformity within the microfibers. Possible reasons for
101

this shortcoming are three-fold: (i) the magnetic field produced by the horseshoe-shaped
magnet is inherently not uniform; (ii) instabilities in the jet, 27 coupled with fiber whipping
that occurs between the Taylor cone and the collection plate; and (iii) nanoparticles residing
in the Taylor cone, as previously reported. 1 All of these considerations could result in fibers
that vary in position with respect to the applied magnetic field. In both cases, the strength of
the magnetic field, and thus the uniformity and contiguity of SPION alignment,
simultaneously vary. One way to enhance the spatial characteristics of SPION arrays is by
aligning the microfibers between electrodes instead of generating a random mat of
microfibers on a collection plate. 28 However achieved, improved control over the position of
the fiber within the field is ultimately expected to promote uniform SPION alignment that
can be pulsed or otherwise induced on-demand.
Unlike ferromagnetic materials, magnetization of the superparamagnetic SPIONs under
investigation only occurs under an external magnetic field. When this field is removed, the
net magnetization of the SPION dispersion becomes close to zero. The magnetization (M)
behavior of SPIONs measuring 17.6 nm in diameter is presented as a function of magnetizing
field strength (H) at 300 K in the SQUID hysteresis plot shown in Figure 3.4a. The value of
MS for the unembedded SPIONs is 38.8 emu/g, whereas that for SPIONs encapsulated within
electrospun PCL microfibers is sharply lowered to 0.77�0.94 emu/g. This nontrivial
reduction in MS for the SPIONs residing in electrospun PCL microfibers with and without an
external magnetic field is due largely to the decreased mobility of the SPIONs, a result that
has been previously reported for embedded nanoparticles. 29 If the matrix viscosity is
102

sufficiently high, only the magnetic moment of a SPION can rotate in the presence of an
external field, but the SPION itself may be unable to do so. In light of these results, the
applied magnetic field must be positioned in the vicinity of the Taylor cone formed during
electrospinning (where solvent remains and the matrix viscosity is relatively low). This
location has been selected here to maximize the ability of the external magnetic field to
induce SPION alignment in electrospun PCL microfibers.
According to Figure 3.4, the SPIONs alone and embedded in PCL exhibit a
superparamagnetic loop characterized by a minimum coercivity and a remanent
magnetization at room temperature. The mean magnetic moment (�) and MS are extracted
from each hysteresis curve by fitting the data to the Langevin function, 30 viz.,
M � MS coth �H �� ��
�� ���
��k
BT��
k ��
��
��
BT
�� (2)
�� �H ��
where kB is the Boltzmann constant and T denotes absolute temperature. Regressed values of
��
MS and � for unembedded SPIONs (Figure 3.4a) are 38.8 emu/g and 3.37 10 �17 emu,
respectively. Similarly, from Figure 3.4b, the values of MS and � determined for randomly
dispersed SPIONs are 0.94 emu/g and 1.26 10 �16 emu, respectively, whereas those for
magnetically-aligned SPIONs are 0.77 emu/g and 1.37 10 �16 emu, respectively. The inset
in Figure 3.4a displays magnetization hysteresis curves recorded at relatively low fields and
reveals that the differences in coercivity and remanent magnetization are, for the most part,
negligible at ambient temperature, which is important for applications conducted at ambient
103

conditions. Such differences are, however, expected to be more apparent at low temperatures.
In fact, previous studies 31 conducted at low temperatures (ca. 5 K) report that arrays of
SPIONs possess a noticeably higher remanence and coercivity than randomly distributed
nanoparticles. No such relation is evident from our results at ambient temperature. This may
be due to the fact that the interparticle distance in our systems is ~5 nm, in which case little,
if any, magnetostatic coupling occurs. 26 Thus, even though the SPIONs are arranged in
unidirectional arrays, their magnetic behavior is still consistent with individual particles
rather than a connected nanowire. It must be recognized, however, that the values of MS are
mass-normalized by multiplying the mass ratio of SPION:PCL in the suspension prior to
electrospinning by the mass of the fiber mat (~1.5 mg). Therefore, the corresponding MS
values are sensitive to the SPION populations present. With the possibility of aggregated
SPION clusters forming in suspension or during electrospinning, the calculated SPION mass
in a given specimen may not be sufficiently accurate to permit direct comparison. In contrast,
values of � are not sensitive to the mass present and reveal an interesting trend that is
amenable to computational modeling: the mean magnetic moment increases with increasing
SPION alignment.
Conclusions
Nanoscale alignment of SPIONs with a mean diameter of 17.6 nm has been achieved
through the use of magnetic field-assisted electrospinning. In this case, an electromagnetic
field of 26 mT is applied perpendicular to the electric field required for electrospinning,
resulting in microfibers that contain one-dimensional arrays of SPIONs. In the present
104

system investigated, a SPION concentration of greater than 0.05 vol% is required to induce
discernible alignment. The arrays thus produced can extend beyond 1 �m under a constant
magnetic field, or they can be controllably shortened by pulsing the magnetic field. The latter
finding suggests that it may be possible in the future to write information to individual
microfibers on the basis of array lengths. Moreover, alignment of the electrospun microfibers
through the use of various established methods can yield nanocomposites with multiscale
(i.e., nanoscale and macroscale) anisotropic properties. Complementary SQUID
measurements at ambient temperature reveal that the saturation magnetization is significantly
lower for SPIONs in electrospun fibers with or without magnetic field-induced alignment
than for unembedded SPIONs. Improving alignment, however, appears to increase the mean
magnetic moment, which provides evidence that nanoscale alignment of SPIONs affects their
intrinsic physical properties.
Acknowledgments
This work was supported by a National Science Foundation Graduate Research
Fellowship (K. E. R.), a North Carolina Space Grant, and the NATO Science for Peace
Program (L. M. B) and Indiana University Faculty Research Program (L. M. B.). We thank
Professor Frank Tsui (University of North Carolina at Chapel Hill) for use of his SQUID, and
Mr. Matt Wolboldt for insightful discussions
105

Figures
Figure 3.1. Schematic illustration of the magnetic field-assisted electrospinning setup
used in this study. Note the position of the electromagnet, which yields a magnetic field that
is perpendicular to the electric field employed during electrospinning.
106

Figure 3.2. TEM images of randomly dispersed SPIONs in electrospun PCL
microfibers varying in SPION concentration (in vol%): (a) 0.5 and (b) 2.5. A TEM image of
SPIONs measuring 17.6 nm in diameter and drop cast from chloroform is included in the
inset of (a). The scalemarker in the inset corresponds to 50 nm.
107

Figure 3.3. TEM images of magnetic field-aligned SPIONs, measuring 17.6 nm in
diameter, in PCL microfibers illustrating long, contiguous arrays in (a), and shorter, pulsed
arrays in (b). The scalemarker in the inset corresponds to 100 nm.
108

Figure 3.4. Magnetization (M) hysteresis curves at 300 K as a function of the
magnetizing field strength (H) for SPIONs measuring 17.6 nm in diameter. In (a), the
hysteresis curves are measured from unembedded SPIONs ( ), as well as randomly
dispersed and magnetically aligned SPIONs in electrospun PCL microfibers (blue and red,
respectively). The inset shows magnetization hysteresis curves recorded for the embedded
SPIONS at low fields and ambient temperature. In (b), the hysteresis curves from the
SPIONs embedded in PCL microfibers (see the corresponding diagrams) are fitted to Eq. 2 in
the text (solid lines) to discern the saturation magnetization and mean dipole moment from
each dataset.
109

Supporting Information
Samples for examination by field-emission scanning electron microscopy (SEM) were
prepared by electrospinning fiber mats directly onto silicon wafers that could be inserted into
the specimen chamber of a JEOL 6400F field-emission microscope, which was operated at
an accelerating voltage of 5 kV. To reduce charging effects, the fibers were sputter-coated
with a thin conductive layer (~6 nm) of Au/Pd at ambient temperature. Energy-dispersive
spectroscopy (EDS) was performed with an Oxford Isis x-ray detector mounted in a Hitachi
S3200 variable-pressure electron microscope and equipped with an ultrathin window for light
element analysis.
An example of a PCL fiber mat containing SPIONs is presented in Figure 3.S1. While the
fibers vary considerably in size, the average diameter of fibers acquired under different
experimental conditions is about 300 nm. Some of the fibers appear to possess a dimpled
surface, which is evident in the inset (a large fiber is selected for display to facilitate
scrutinization of the fiber surface). A representative EDS spectrum, as well as an
enlargement of a portion thereof, of a PCL fiber (with 2.5 wt% SPIONs measuring 17.6 nm
in diameter) subjected to a magnetic field during electrospinning is provided in Figure 3.S2.
Elements detected in this and related spectra include C from the fibers, Si and O from the
silicon wafer, and Au and Pd from the conductive coating. Peaks corresponding to the K�
and L lines of Fe (in the vicinities of 6.4 and 0.7 keV, respectively) are absent, indicating that
the SPIONS are, for the most part, embedded within the fibers and do not reside, to an
appreciable extent, on the fiber surface.
110

Figure 3.S1. SEM image of SPION-containing PCL fibers electrospun in the presence of
an external magnetic field. The inset shows evidence of surface dimpling on a large fiber.
The scalemarker in the inset corresponds to 2 �m.
111

Figure 3.S2. EDS spectrum of a SPION-containing PCL fiber electrospun in the
presence of an external magnetic field. The elements responsible for the observed peaks are
labeled, and the x-ray energies associated with the K� and L lines of Fe are identified by the
blue lines.
112

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115

CHAPTER IV
Using Polymer Blend Morphology to Position Ligand-Functionalized
Abstract
Nanoparticles in Electrospun Polymer Microfibers
Blends of hydrophobic and hydrophilic polymers have been prepared to discern the
feasibility of controlling the spatial location of superparamagnetic iron oxide nanoparticles
(SPIONs) within electrospun microfibers on the basis of thermodynamic compatibility.
Although inorganic nanoparticles tend to aggregate, a carefully designed polymer matrix can
be used to position nanoparticles and achieve unique optical, electronic and/or magnetic
properties. In the present study, poly(ethylene oxide) (PEO) and poly(2-vinyl pyridine)
(P2VP) phase-separate in electrospun microfibers and form discrete, but diffuse, dispersions
at low concentrations of the minority component. At higher concentrations, a core-sheath
structure naturally develops wherein the hydrophobic SPIONs are sequestered, according to
electron microscopy. X-ray diffractometry confirms that the (120) reflection increases in
intensity after SPIONs are added, suggesting that the PEO chains become increasingly
aligned along the fiber axis and that the nanoparticles reside at the PEO/P2VP interface affect
PEO crystallization. We thus achieve control over SPION positioning within these
microfibers without requiring post-modification.
116

Introduction
The formation of polymer nanocomposites has become a preferred means by which to
combine the highly desirable electrical, magnetic, optical and/or mechanical properties of
metals and metal oxides with the flexibility, light weight and processability of inexpensive
and tough polymers. 1 Incorporation of stimuli-responsive nanoscale objects into
homogeneous polymer matrices generally yields functional materials that are suitable for
various applications such as data storage, conductive nanowires, nonwoven sensors,
magnetic filters, and nanoactuation. 2 Metallic nanoparticles, in particular, tend to aggregate
and reduce the population of surface atoms because of the driving force to decrease their total
surface energy. 3 Harnessing control over the spatial location of nanoparticles within a
polymer matrix may enable favorable attributes 4 associated with their electronic, optical,
magnetic, and catalytic properties. One way by which to attain this objective is to employ a
heterogeneous polymer matrix composed of either a macrophase-separated polymer blend 5 or
a microphase-separated block copolymer. 6 Independent studies have demonstrated 7 that
judicious manipulation of enthalpic and entropic contributions to the system free energy by
systematic variation in nanoparticle size, concentration and surface chemistry can be
exploited to position nanoparticles within (micro)phase domains or along interfaces on the
basis of thermodynamic considerations.
Another, process-oriented, means by which to control nanoparticle positioning is
through physical confinement, 8 which can be readily achieved by spin-casting or
electrospinning. Electrospinning is an emerging fabrication technique that produces solid
117

fibers with diameters ranging from several tens of nanometers up to several microns, a high
surface-area-to-volume ratio and high porosity from both polymer solutions and melts. 9 It is
an appealing fiber-spinning strategy due to the relatively straightforward setup and its ability
to tune fiber properties on the basis of both processing and solution variables. A polymer
solution or melt is slowly ejected from a spinneret which is connected to a high voltage
source. As a voltage is applied, a charge develops on the surface of the liquid due to mutual
charge repulsion 10 and as these electrostatic forces exceed the surface tension of the polymer
solution or melt, a charged jet is emitted. This jet has a conical structure defined as a “Taylor
cone” 11 and undergoes a whipping process, 12 during which time the solvent evaporates and is
collected as a dry, randomly oriented polymer fiber mat on the collection plate.
Since the fiber-forming polymer often lacks functionality that is desired for growing
multifunctional applications, functionality can be imparted by several different routes, such
as addition of a second polymer, 13 spinning a block copolymer system, in situ growth of core
or surface functionalities through the incorporation of solution or melt precursors, 14 co- 15 and
multi-axial 16 spinnerets where multiple layers of a fiber can be created, or post-electrospun
surface modification 17 of the fiber surface.
Both miscible and immiscible polymer blends can be electrospun to create new
morphologies of phase-separated nano/microfibers. 18 When two polymers phase-separate
into discrete domains, one component can be selectively removed to yield nanoporous fibers,
which has been reported for polymer pairs such as poly(lactic acid)/poly(vinylpyrrolidone) 19 ,
poly(ethylene oxide)/silk, 20 poly(ethylene oxide)/poly(methyl methacrylate) 21 and
118

poly(glycolic acid)/chitin. 22 When phase separation yields more contiguous biphasic
morphologies, core-sheath structures 23 can develop. A core-sheath microstructure is highly
desirable since it is possible to combine polymers with two vastly different property sets
(e.g., the sheath can be insulating while the core can be conductive 23 ). In addition, the core
may also be selectively removed to form nano/microtubes, 24 which would be beneficial for
time-controlled drug release 25 and sensory applications 26 . Core-sheath fiber microstructures
are formed by several different protocols: careful solvent selection, 27 thermostatic/kinetic
polymer contrast, 28 co-axial electrospinning, 29 or electric field-induced. 21 Although these
specialized procedures have been developed to generate core-sheath microstructures, we
hasten to point out that such microstructures rarely form naturally. Using polymer blends to
spontaneously form core-sheath microstructures during electrospinning is particularly
appealing since the chemistry, molecular weight and concentration of the polymer
constituents can be independently adjusted.
The primary objective of this work is to generate nanocomposite nano/microfibers
containing spatially positioned superparamagnetic iron oxide nanoparticles (SPIONs) by
electrospinning. Control over nanoparticle positioning has been previously achieved within
such fibers by post-crystallization of the polymer matrix, 30 alteration of the electric field
employed during electrospinning, 31 development of a polymeric nanostructure via self-
assembly, 32 or strategic use of coaxial electrospinning. 33 Here, we utilize a blend of
hydrophobic and hydrophilic polymers to discern the feasibility of controlling the spatial
location of SPIONS within electrospun fibers on the basis of thermodynamic compatibility.
119

Experimental
Materials
Poly(2-vinyl pyridine) (P2VP) with a molecular weight of 200 kDa according to the
manufacturer was obtained from Scientific Polymer Products (Ontario, NY), and
poly(ethylene oxide) (PEO) with a molecular weight of 200 kDa according to the
manufacturer was provided by Sigma-Aldrich (Milwaukee, WI). Reagent-grade chloroform
for electrospinning was purchased from Fisher Scientific (Fairlawn, NJ) and used as-
received. Iron oleate, oleic acid and the various solvents used in the nanoparticle synthesis
were acquired from TCI America (Portland, OR) and also used without further purification.
Preparation of Nanoparticles
The SPIONs with diameters ranging from 12 to 24 nm were prepared by thermal
decomposition of iron oleate in the presence of oleic acid as the capping agent in high boiling
hydrocarbons (docosane and eicosane), according to the procedures described elsewhere. 34,35
Previous X-ray diffractometry (XRD) data confirmed 34 that these as-synthesized SPIONs are
crystalline and contain mostly wüstite (Fe(1-x)O) and some spinel (most likely Fe3O4).
According to magnetometry measurements, 36 they are superparamagnetic. The resultant
SPIONs were precipitated with a mixture of acetone and hexane, separated by centrifugation
and suspended in chloroform at concentrations between 10 and 50 mg/mL (Figure 4.1).
120

Preparation of Nano/microfibers
Suspensions for electrospinning were prepared by first dissolving predetermined
quantities of PEO and P2VP (ranging from 0 to 100 wt% PEO) in chloroform and then
adding a target mass of a SPION suspension to yield polymer concentrations ranging from 2
to 8 wt% at three different SPION concentrations: 0.5, 1.0 and 2.5 vol%. The in-house
electrospinning unit, operated at 8 kV, employed a syringe pump operated at flow rates
between 10 and 35 �L/min, and an Al collection target maintained at a separation distance of
12 cm. Electrospun nano/microfibers were collected as unoriented nonwoven mats on the Al
plate, as well as on carbon-coated transmission electron microscopy (TEM) grids adhered to
the plate.
Characterization of Nano/microfibers
To discern the blend morphologies of the electrospun nano/microfibers and the spatial
distribution of SPIONs contained therein, TEM was performed on a field-emission Hitachi
HF2000 microscope operated at 200 kV. To enhance the contrast between the two polymer
phases, the P2VP has been selectively stained with the vapor of 0.5% OsO4(aq) for 7 min.
Corresponding scanning electron microscopy (SEM) images were collected from specimens
sputter-coated with 6 nm of Au/Pd on a JEOL 6400F field-emission microscope operated at 5
kV. Complementary XRD was performed at 2θ angles ranging between 5 and 30° in 0.01°
increments on a Rigaku Smartlab diffractometer with CuKα radiation at a wavelength (�) of
0.1541 nm. Zero-shear viscosity measurements were performed on a TA Instruments AR-G2
121

heometer over shear stresses extending from 0.05 to 50 Pa on 40 mm 2° steel cone plates in
soft bearing mode.
Results and Discussion
Blends of hydrophobic and hydrophilic polymers have been prepared and electrospun to
discern the feasibility of controlling the spatial location of SPIONs within electrospun
nano/microfibers due to preferential compatibility. Both P2VP and PEO are specifically
chosen for this purpose and likewise satisfy two additional criteria: they are of comparable
molecular weight and they are both soluble in chloroform, which is used for electrospinning.
The blends range in concentration from 0 to 100 wt% PEO. Due to differences in mass
density (1.15 g/cm 3 for P2VP and 1.21 g/cm 3 for PEO), the concentration of each polymer
solution is appropriately adjusted to yield electrospun nano/microfibers of comparable
diameter (~1200 nm), thereby eliminating fiber size from consideration in morphological
analyses. Fibers of this size will hereafter be referred to as microfibers. Nonpolar solubility
parameters for PEO, P2VP and chloroform are reported 37 as 15.6, 21.7 and 17.8 MPa 1/2 ,
respectively. On the basis of these values alone, it follows that PEO is more compatible with
chloroform and will remain solvated longer during electrospinning than P2VP. In the event
that a core-sheath microstructure developed, one could reasonably presume that P2VP should
form the core and PEO should form the surrounding sheath (and crystallize).
Field-emission SEM confirms that these PEO/P2VP microfibers with and without
SPIONs possess an average diameter of ~1200 nm and relatively smooth surfaces with very
few bead defects, as evidenced by the images displayed in Figures 4.2a and 4.2b. It is
122

interesting that the defects tend to contain large SPION aggregates (cf. the inset of Figure
4.2b), which most likely formed in the suspension reservoir by nanoparticles with insufficient
surface functionalization. Examples of TEM images acquired from microfibers varying in
composition without SPIONs establish that the PEO phase is indiscernible at low
concentrations (suggesting that the two homopolymers may actually be at least partially
miscible under the conditions employed). Unlike previous studies that have reported 38
nanoparticle-induced changes in nanoscale morphologies, these morphologies are largely
retained when SPIONs at a concentration of 2.5 vol% are added, as in Figure 4.3. Since the
polymers are phase-separated at some blend compositions, the hydrophobic nanoparticles can
be anticipated to preferentially reside in one phase. According to the TEM images of PEO-
rich blends presented in Figure 4.4, the SPIONs, measuring 18 nm in diameter, are largely
sequestered within the P2VP phase, which has been selectively stained and thus appears
more electron-dense (dark). The SPIONs do not adopt a particular alignment and appear to
be randomly distributed throughout the P2VP phase. Since these SPIONs are coated with
oleic acid, which induces mutually hydrophobic interactions between alkyl chains, 39
relatively little nanoparticle aggregation is observed in these microfibers, regardless of blend
composition. Very few small SPION aggregates, not evident in the images shown in Figure
4.3 but occasionally observed in other specimens, appear as spheroidal clusters. The diameter
and concentration of SPIONs has been systematically varied from 12 to 24 nm and from 0.5
to 2.5 vol%, respectively, with no significant effect on their host phase or the accompanying
123

lend morphology. For these reasons, we only present images of SPIONs measuring 18 nm
in diameter at a concentration of 2.5 vol%.
The phase morphology of polymer blends subjected to nonequilibrium processing such as
electrospinning depends sensitively on thermodynamic (e.g., interfacial tension, solubility
and molecular weight) and kinetic (viscoelasticity and molecular mobility) factors. 40
Nanostructure formation within a polymer fiber during electrospinning is estimated to be on
a time scale of ~10 ms due to concurrent solvent evaporation and jet elongation. 41 In the
present PEO/P2VP blends, thermodynamic driving forces seemingly dominate, especially at
intermediate to high concentrations of PEO, even though P2VP in chloroform possesses a
lower viscosity than PEO (discussed later). Such incompatibility is evident from the
differences in solubility parameters described earlier. At low PEO concentrations (e.g., 20
wt% PEO), the absence of two distinct phases is suggestive of phase mixing. One must
exercise caution in drawing this conclusion since the stained P2VP is more electron-dense
than PEO, in which case discrete PEO domains may not be visible within the P2VP matrix.
Analysis of images such as Figure 4.3a confirm that there is no significant change in optical
density along the length of the microfibers. As the concentration of PEO is increased to 30
wt% (Figure 4.3b), seemingly unconnected PEO domains develop along the edge of the
microfibers, thereby forming the onset of core-sheath microstructure. At this blend
composition, the low viscosity of the blend in chloroform allows PEO chains to migrate to
the microfiber-air interface as the chloroform evaporates. In response, the P2VP chains
remain along and thus form the core.
124

To quantify the extent to which the PEO and P2VP are segregated, the interfacial
thickness (�) corresponding to the change in optical density across the interfacial normal has
been measured from as many as 10 different regions in up to 3 different images. We
recognize that these values are only estimates, as the interfacial curvature has not been
considered, but they can provide insight into the incompatibility between PEO and P2VP
upon electrospinning. At 30 wt% PEO, � is approximately 32 nm, which corresponds to a
very diffuse interface. 42 As the concentration of PEO is increased to 40 wt%, P2VP forms the
major phase with small pockets of PEO along the fiber surface. It is interesting to note that
these PEO domains, which range from 175 to 400 nm in length, are oriented perpendicular to
the fiber spinning direction. This orientation may be a consequence of solvent evaporation as
the microfiber solidifies. Although the morphology is somewhat different from what is seen
at the lower PEO concentration, � is comparable at ~29 nm, confirming that the degree of
phase segregation is similar. When the concentration of PEO is further increased to 50 wt%
(Figure 4.3c), two sizes of P2VP domains reside within an apparent PEO matrix. The larger
domain sizes measure ~200 nm across, whereas the smaller ones are ~60 nm, and � increases
noticeably to ~52 nm, which is indicative of enhanced phase mixing. At 60 wt% PEO (Figure
4.3d), the PEO possesses sufficient mobility to migrate in large part to the microfiber-air
interface and � ≈ 25 nm is similar to what it was at lower PEO concentrations. In the case of
the fractal-like core-sheath microstructure formed at 70 wt% PEO (Figure 4.3e), the
increased viscosity and correspondingly low mobility of PEO prevents an intact core-sheath
structure from forming and phase separation is not sharply delineated, with a very large � >
125

100 nm. At 80 wt% PEO (Figure 4.3f), well-defined domains of P2VP with � ≈ 26 nm are
visible within the PEO matrix. Even at this low concentration of hydrophobic P2VP, the
majority of the SPIONs prefer to reside within this phase.
In summary, the SPIONs prefer to reside in the hydrophobic P2VP phase. Since they are
coated with oleic acid, relatively little aggregation is observed in the microfibers.
Comparison of Figures 4.3 and 4.4 suggests that the addition of SPIONs improves the extent
of phase separation between PEO and P2VP without affecting the morphology that develops
at a given blend composition. In this regard, the SPIONs act as a decompatabilizer between
the two homopolymers. Such behavior has been previously reported 4 due to incompatability
between the surface ligand and the host polymer phase. Some of the morphologies seen in the
microfibers are attributed to differences in viscosity. To ascertain the effect of SPIONs on the
viscosity of PEO and P2VP in chloroform, we have measured the zero-shear viscosity (�) of
the following systems in chloroform at ambient temperature: 8 wt% P2VP and 2 wt% PEO,
both with and without 1 wt% SPION. Without nanoparticles, the PEO solution possesses a
higher viscosity (0.420 Pa-s) than the P2VP solution (0.201 Pa-s) despite the 4x higher
polymer concentration in the latter. Incorporation of the SPIONs consistently lowers � in
both systems and causes their viscosities to become comparable: 0.083 Pa-s for PEO and
0.106 Pa-s for P2VP. A nanoparticle-induced reduction in viscosity has been previously
reported for polymer melts, 43,44 and molecular dynamics simulations 45 have confirmed that a
viscosity decrease upon nanoparticle addition is characteristic of repulsive systems. The
solubility parameter of the oleic acid at 15.6 MPa1/2 is sufficiently different from that
126

corresponding to the host P2VP, which is consistent with a repulsive system and an
accompanying reduction in �. This observation also implies that the high-molecular-weight
P2VP chains responsible for � tend to physisorb on the nanoparticle surface and contribute
much less to �. 43
To elucidate the chain packing within the morphologies observed in the electrospun
microfibers of PEO/P2VP blends, we have used XRD to examine the crystalline structure of
microfiber mats with and without added SPIONs. The reference XRD pattern of PEO powder
clearly displays the two signature peaks of the monoclinic lattice of crystalline PEO at 19.06°
2� (Peak I) and 23.22° 2� (Peak II). 46 Peak I corresponds to the (120) reflection, while peak
II refers to the (112) and (032) reflections. 47 The average crystallite size (�) can be calculated
by the Scherrer equation, given by 48
� � K�
�cos�
where K is the shape factor (= 0.9 in the present case), β is a measure of line broadening (i.e.,
full-width at half the ��maximum
intensity) and � denotes the Bragg angle. The corresponding
d-spacing is determined by the Bragg equation: 49
d � n�
2sin�
in which n is the integral diffraction order equal to unity here. Pure PEO consists of large,
well-defined crystals
��
(with � = 28.27 nm from Peak I) wherein the intensity of Peak I is
~20% higher in intensity than Peak II. Values of � and the Peak I/Peak II ratio (I/II) extracted
from the scattering profiles provided in Figure 4.4 for electrospun microfibers at different
127
(1)
(2)

lend compositions are listed in Table 4.1. When PEO is electrospun, the positions remain
nearly the same for both Peaks I and II, but peak broadening yields an increase in � and thus
a reduction in �. This result has been previously reported 40 for electrospun systems since the
timescale of crystal formation during electrospinning is much shorter 31 than that in a bulk
system, such as the powder produced upon polymerization. A decrease in the I/II peak ratio
to 0.93 means that the (120) reflection is stronger, which occurs when the PEO polymer
chains preferentially orient along the axis of elongational flow during electrospinning. In this
scenario, the PEO lamellae lie perpendicular to the fiber axis. 50
As the P2VP content in the fiber increases (without SPIONs), several trends become
apparent from the profiles displayed in Figure 4.4. The first is that the positions of Peaks I
and II do not change within experimental error, which translates into a crystal d-spacing that
is independent of blend composition. This observation is consistent with a PEO phase that
remains relatively pure (unmixed) in the presence of P2VP at PEO concentrations as low as
40 wt%. At lower concentrations, the peaks shift and become distorted, indicating enhanced
phase mixing as inferred earlier. The values of t extracted from Peak I systematically
decrease from ~28 to ~19 nm with increasing P2VP loading, indicating that the length of the
crystals along the fiber axis decrease as the reservoir of PEO within the microfiber is lowered
and the domains of PEO within the microfibers become smaller and more dispersed. In
contrast, the width of the crystals oriented normal to the fiber axis fluctuates modestly
between about 9 and 13 nm. Values of the I/II peak ratio are always less than unity in
electrospun microfibers varying in blend composition, in which case the crystalline lamellae
128

prefer lying perpendicular to the fiber axis in electrospun microfibers. Other than the diffuse
interfaces discussed earlier, there is no evidence from XRD that the PEO and P2VP interact
to any discernible extent. 51
Addition of SPIONs measuring 18 nm in diameter at a concentration of a 2.5 vol% to the
electrospun PEO/P2VP microfibers yields the XRD patterns presented in Figure 4.4.
Pertinent metrics extracted from these profiles are provided in Table 4.2. In microfibers
composed of 100 wt% PEO, the SPIONs promote a significant reduction (by about 30%) in
crystal size along the fiber axis and almost no change in either crystal d-spacing or crystal
size normal to the fiber axis. In addition, the I/II peak ratio is less than unity (0.46), which
confirms that the PEO chains, as well as their innate ability to crystallize, are strongly
affected by the presence of SPIONs within PEO. This dependence is not evident from the
TEM images in Figure 4.3 since the images only show PEO/P2VP blend morphologies with
SPIONs. Upon reducing the PEO concentration to 80 wt%, the longitudinal crystal size
increases, but nonetheless remains lower than that in the absence of SPIONs. Further
decreasing the PEO concentration results in a systematic reduction in longitudinal crystal
size, as was seen in PEO/P2VP microfibers without SPIONs. These results, as well as the
corresponding transverse crystal sizes, are shown as a function of blend composition in
Figure 4.6 and confirm that, even in the PEO/P2VP blends, the SPIONs affect the crystal
characteristics of PEO. Recall from TEM images such as those in Figure 4.3, however, that
the SPIONs are preferentially sequestered within the P2VP phase.
129

The influence of SPIONs on PEO in electrospun microfibers composed of PEO/P2VP
blends can be explained in one of two ways. Firstly, a fraction of added SPIONs, while
thermodynamically attracted to the P2VP phase due to the presence of the oleic acid ligands,
remains kinetically trapped in the PEO phase during microfiber formation. Alternatively,
some SPIONs reside along the PEO/P2VP interface. In both cases, the SPIONs contact PEO
chains and affect their ability to enter into three-dimensional registry and form crystals. To
decide which situation is more likely, we turn our attention to the I/II peak ratio, which is
always less than unity in PEO/P2VP systems (except for pure PEO) without SPIONs.
Confinement due to P2VP at any loading level appears to favor the formation of PEO
lamellae oriented perpendicular to the fiber direction (transverse). With SPIONs, even pure
electrospun PEO exhibits a peak ratio that is less than unity. As the concentration of PEO is
reduced in this series, however, the peak ratio increases beyond unity so that the PEO
lamellae are oriented parallel to the fiber direction (longitudinal). This change in lamellar
orientation implies that the PEO phase is nearly pure (since the presence of SPIONs in PEO
promote the opposite orientation), in which case the SPIONs promote decompatibilization, as
concluded earlier. Since the PEO chains adopt a crystal orientation that is reminiscent of pure
PEO but remain affected by SPIONs preferentially located within the P2VP phase, it
immediately follows that SPIONs located along the PEO/P2VP interface are most likely
responsible for the observed changes in the PEO crystal structure. This scenario is
schematically illustrated in Figure 4.7. At 20 wt% PEO, the XRD pattern in Figure 4.5 shows
a mostly amorphous scattering signature, which reflects the low loading level of semi-
130

crystalline PEO. A broad Peak I and virtually nonexistent Peak II reveal very small crystals,
which is consistent with the trend observed in Figure 4.6. In this case, phase mixing between
PEO and P2VP and the existence of SPIONs within PEO hinder the crystallization of PEO
chains.
Conclusions
Blends composed of a hydrophobic (P2VP) and a hydrophilic (PEO) polymer have been
modified with SPIONs and electrospun into microfibers to discern the feasibility of
controlling the spatial location of SPIONs on the basis of thermodynamic compatibility and
blend morphology. At blend compositions favoring phase separation into dispersed domain
or core-sheath morphologies, the SPIONs tend to locate within the P2VP phase. Analysis of
the PEO crystallinity demonstrates that the addition of P2VP to electrospun microfibers
without SPIONs generally reduces the size of PEO crystals and changes the orientation of
PEO lamellae from longitudinal to transverse due most likely to confinement effects. In the
presence of SPIONs, a similar, but more pronounced reduction in crystal size is observed
with increasing P2VP concentration. In addition, the orientation of PEO lamellae changes
from transverse to longitudinal, suggesting that SPIONs located along the PEO/P2VP
interface influence PEO chain packing and crystallization.
Acknowledgments
This work was supported by a the NC Space Grant and the National Science Foundation
Graduate Research Fellowship (K. E. R.), a National Aeronautics and Space Administration
131

Graduate Fellowship (K. E. R.), and the NATO Science for Peace Program (L. M. B) and
Indiana University Faculty Research Program (L. M. B.).
132

Tables
Table 4.1: XRD characteristics of PEO powder and electrospun PEO/P2VP microfibers.
System Peak Angle
(° 2�)
PEO Powder
100 wt% PEO
Fiber Mat
80 wt% PEO
Fiber Mat
60 wt% PEO
Fiber Mat
40 wt% PEO
Fiber Mat
133
d-spacing
(nm)
��
(nm)
I 19.06 0.47 28.27
II 23.22 0.38 12.27
I 19.06 0.47 28.23
II 23.25 0.38 11.22
I 19.08 0.47 26.56
II 23.22 0.38 10.17
I 19.21 0.46 23.24
II 23.40 0.38 12.78
I 19.07 0.47 19.20
II 23.34 0.38 9.40

Table 4.2: XRD characteristics of electrospun PEO/P2VP microfibers with SPIONs
measuring 18 nm in diameter and added at a concentration of 2.5 vol%.
System Peak Angle
(°2�)
100 wt% PEO
Fiber Mat
80 wt% PEO
Fiber Mat
60 wt% PEO
Fiber Mat
40 wt% PEO
Fiber Mat
20 wt% PEO
Fiber Mat
134
d-spacing
(nm)
��
(°)
��
(nm)
I 19.27 0.46 0.28 19.62
II 23.52 0.38 0.74 11.56
I 19.13 0.46 0.35 23.56
II 23.31 0.38 0.76 11.27
I 19.16 0.46 0.40 20.79
II 23.35 0.38 0.72 11.85
I 19.07 0.47 0.59 14.09
II 23.36 0.38 1.19 7.19
I 19.61 0.45 7.30 1.14
II — — — —

Figures
Figure 4.1. TEM image of SPIONs measuring 16.4 nm in diameter and drop cast from
chloroform.
135

a.
b.
Figure 4.2. (a) SEM image of SPION-containing PEO/P2VP microfibers composed of 80
wt% PEO and electrospun from an 8.5 wt% solution in chloroform. (b) An enlargement
showing the surface of the microfibers included in (a). The inset in (b) displays a SPION-rich
bead, the scalemarker corresponds to 500 nm.
136

Figure 4.3. TEM images of SPIONs, measuring 18 nm in diameter at a concentration of 2.5
vol%, in biphasic microfibers composed of PEO/P2VP at different PEO concentrations
(labeled). Inset scalebars all represent 200 nm and all other scalebars represent 500 nm.
137

16 18 20 22 24 26 28 30
2��(degrees)
138
80% PEO
100% PEO
PEO Powder
60% PEO
Intensity (a.u.)
40% PEO
Figure 4.4. XRD patterns acquired from PEO powder and electrospun microfibers composed
of PEO/P2VP at different PEO concentrations (labeled).

16 18 20 22 24 26 28 30
2��(degrees)
139
80% PEO/NP
80% PEO/NP
40% PEO/NP
60% PEO/NP
Intensity (a.u.)
20% PEO/NP
Figure 4.5. XRD patterns acquired from electrospun microfibers composed of PEO/P2VP
with SPIONs (18 nm and 2.5 vol%) at different PEO concentrations (labeled).

� (nm)
30
25
20
15
10
5
0
100
Longitudinal
Transverse
80
Figure 4.6. Average PEO crystal size (t) extracted from XRD patterns and presented as a
function of PEO concentration parallel (circles) and perpendicular (triangles) to the fiber axis
for systems without (open symbols) and with (filled symbols) SPIONs. Values measured
from PEO powder are included (triangles). The solid and dashed lines serve as guides for the
eye.
140
60
40
PEO concentration (wt%)
20

a.
b.
Figure 4.7. Schematic illustration depicting the arrangement of polymer chains in a
core/sheath microstructure of PEO/P2VP (a) before and (b) after SPION addition (not to
scale). Addition of SPIONs promotes a reduction in crystal size but a more parallel chain
arrangement with respect to the fiber axis.
141

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144

CHAPTER V
Nanostructured Organometallic Polymer Systems Containing
Poly(ferrocenylsilanes)
5.1 Introduction
An organometallic polymer contains a transition metal in the main chain or more
specifically has a metal-carbon σ or π bond. 1 The term ‘transition metal’ refers to any
element that has an incomplete d sub-shell or which can give rise to cations with an
incomplete d sub-shell. 2 The properties that make these transition metals interesting,
however, are their ability to have multiple oxidation states that can be easily controlled
through electric currents. The first organometallic polymer was synthesized in 1955 as
poly(vinylferrocene) and displayed reversible oxidation-reduction properties. 3 The
incorporation of metallic elements into polymer systems allow different coordination
numbers and geometries and thus supply valuable magnetic, optical or catalytic properties.
Structurally, metallopolymers can be divided into three categories: metals incorporated
directly into the polymer chain, π or σ-coordinated metals, and metallic moieties pendant to
the polymer backbone or in side chains. 4 Some challenges that have thwarted the synthesis of
organometallics include low molecular weights, oligomers, impurities, and insolubility.
In the early 1990s, Manners et al. 5 fabricated poly(ferrocenylsilane) (PFS) with a high
molecular weight and narrow polydispersity through a ring-opening polymerization method,
145

and this synthesis technique has since yielded other strained monomers including bridging
elements such as germanium, tin and phosphous. 6 Poly(ferrocenylsilane)s can be tuned to
adopt either a semi-crystalline or amorphous state as a consequence of the constituent groups
on the silicon atom 7 where symmetrically substituted constituents, R=R'=Me, impart
crystallinity. The prevalence of iron in the main chain conveys some interesting properties
not present in non-metal containing polymers such as redox-activity due to the reversible
(Fe(II)/Fe(III) couple) 2 , the ability to be pyrolyzed into a magnetic ceramic, 8 and semi- and
photo-conductivity. 9
Block copolymers are macromolecules containing two or more types of repeat units
within long contiguous sequences, or “blocks,” of the same unit. These sequences are
covalently linked to form a single macromolecule. They can spontaneously self-organize, or
microphase-separate, into a variety of ordered nanoscale morphologies: lamellae,
hexagonally packed cylinders, spherical micelles on a body- or face-centered cubic lattice, or
complex bicontinuous morphologies (e.g., the gyroid). 10 The type of morphology exhibited
by the block copolymer depends on the chemical attributes and lengths of the blocks, the
molecular architecture and temperature, as well as the presence of an additive such as a
solvent, 11 homopolymer, 12 nanoparticle, 13 or another block copolymer. 14 Due to their ability
to microphase-separate, block copolymers constitute a versatile platform for a number of
existing and emerging technologies such as adhesives, membranes, drug delivery,
biomaterials and nanolithography. 15
146

The formation of metal-containing block copolymers was possible due to the
controllability of anionic ring-opening polymerization of silicon-bridged ferrocenes. 16
Polymerization must occur through sequential addition of monomers with decreasing end-
group reactivity such that poly(styrene) (PS) ~ poly(isoprene) (PI) > PFS >
poly(dimethylsiloxane) (PDMS). 16 The first two PFS-containing block copolymers, 17 PS-b-
PFS and PFS-b-PDMS, have allowed the synthesis of block copolymers of PFS and PI, 18
poly(methylmetacrylate) (PMMA), 19 poly(2-vinyl pyridine) (P2VP),
poly(ferrocenylethylmethylsilane) (PFEMS), 20 and hybridization with polypeptides. 21 These
metallated block copolymers have been shown to self-assemble in the solid state and have
ordered into spherical, cylindrical, lamellar, and gyroid morphologies. 22 Triblock copolymers
can be selectively etched so that only PFS is remaining, as seen in Figure 5.3 which gives
rise to many lithographic applications. In the solid state, amorphous PFS with asymmetrical
constituents on the silicon atom can be used to promote phase separation.
Periodic PFS domains can be converted to iron-rich clusters within a ceramic domain by
pyrolysis at temperatures of about 600°C, which can lead to the growth of single-walled
carbon nanotubes for soft lithography, 23 the formation of magnetic ceramics 24 and PFS-
derived catalysts. 25 Solution self-assembly of diblock copolymers occurs when a block-
selective solvent is utilized to induce segregation of the solvent-incompatible ‘core’
surrounded by the solvent-compatible ‘corona.’ Micelles of PFS-containing block
copolymers have produced micellar morphologies such as cylinders, 26,27 tubes, 27 fibers, 28 and
tapes. 18 Investigations performed in the bulk typically utilize poly(ferrocenyldimethylsilane)
147

(PFDMS), whose crystalline structure is responsible for the growth of micelles. 29 Initial
experiments of PFDMS-b-PDMS (with a block ratio of 6:1, respectively) in hexanes
demonstrated that cylinders with a PFS core and a PDMS corona formed. 26 Imaging these
organometallic, self-assembled structures by transmission electron microscopy (TEM) is
facilitated due to the inherent contrast arising from the iron-rich core. Forming a core of
PFDMS in cylindrical micelles requires a corona:core ratio of at least 5:1, 18,26 but when this
ratio reaches 12:1, hollow, tube-like structures most likely form. 19,27 The length of the
cylindrical micelles can also be controlled: the addition of a small amount of a common
solvent will allow growth of the micelles, whereas ultrasound or high temperatures will
cleave long cylindrical micelles and yield short ones instead. Self-assembly of PFDMS-b-
P2VP has been investigated in several alcohols, and cylindrical micelles are observed to form
in isopropanol but not ethanol. 30 The reason for this is attributed to the fact that isopropanol
is a better solvent for PFDMS than ethanol on the basis of solubility parameters. Thus, the
PFS chains remain solvated longer and have time to rearrange and crystallize more
efficiently. In addition, PFS-b-PI block copolymers self-organize into cylindrical micelles in
a PI-selective solvent, and the vinyl groups in the PI corona were subjected to a Pt(0)-
catalyzed crosslinking reaction, which was a precursor to making PFS nanoceramics with
shape retention by pyrolysis. 31 Nanocomposite self-assembled structures consisting of PFS-b-
PVMS wormlike micelles react with Ag[PF6] to create a one-dimensional array of silver
nanoparticles encapsulated within the worm-like micelles, 32 thereby allowing for alignment
of highly oriented nanoparticle arrays.
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5.2 Experimental
5.2.1 Materials
Characteristics of all the polymers employed here, including their chemical constitution,
number-average molecular weight (Mn), polydispersity index (PDI), and composition (�PFS),
are listed in Table 5.1. Reagent-grade chloroform and dimethylformamide (DMF) were
obtained from Fisher Scientific (Fairlawn, NJ). Tetrahydrofuran (THF) was purchased from
Fisher Scientific (Fairlawn, NJ), and dichloromethane (DCM) was acquired from Sigma-
Aldrich (Milwaukee, WI)
5.2.2. Synthesis of Specialty Polymers
All organometallic polymers were synthesized in the Manners laboratory at Bristol
University specifically for this research. Due to this, molecular weights and PDI’s are not
consistent between samples, and extremely limited amounts of polymer prohibit
determination of solution properties. A month-long visit to Bristol University was undertaken
in 2008 to prepare low-molecular-weight poly(ferrocenylpropylmethylsilane) (PFPMS) and
poly(ferrocenyl-dimethylsilane)-b-polyisoprene(PFS-b-PI).
Dilithioferrocene·tetramethylethylenediamine was previously prepared in the Manners
laboratory via the synthesis by Rider. 20 All parts of the synthesis were conducted under
nitrogen since ferrocene can self-combust. Dilithioferrocene·tetramethylethylenediamine
(15.4 g) was suspended in diethyl ether (500 mL) and the suspension was chilled to -78°C in
an acetone/dry ice bath. Distilled dichloroethylmethylsilane (8.4g) was added dropwise via
149

canulla tubing to the flask, which was allowed to warm to ambient temperature for 4 h.
During this heating, the color of the material changed from orange to dark red. The solvent
was then evaporated and the solids were dried under dynamic vacuum overnight. The solids
were then dissolved in hexanes and recrystallized twice at -55°C to remove any impurities.
This was then followed with two sublimation cycles to eliminate any reacted ferrocenophane.
In an inert glove box, 250 mg of the ethylmethyl-silaferrocenophane was dissolved in 5 mL
of THF to which 1.5 μL of n-butyllithium was added (1.6 M solution in hexanes). The
polymerization was allowed to proceed for 30 min and was terminated by degassed
methanol. The solid was then precipitated into methanol and dried overnight under dynamic
vacuum, resulting in PFEMS, as depicted in Figure 5.1. Thermal ring-opening
polymerization, as described above, is diverse in that it allows a range of silicon-substituted
functional groups. Thus, PFDMS and poly(ferrocenylphenylmethylsilane) (PFPMS) can be
prepared by the same synthetic scheme with different starting monomers.
5.2.3. Preparation of Nanoparticles
The thermal decomposition of iron oleate with oleic acid occurred through thermal
decomposition and resulted in iron oxide nanoparticles with diameters ranging from 10-14
nm. 33,34 These as-synthesized nanoparticles are crystalline according to x-ray diffraction
(XRD) and contain mostly wüstite (Fe(1-x)O) and some spinel (most likely Fe3O4), 33 and are
superparamagnetic iron oxide nanoparticles (SPIONs) according to magnetic
measurements. 35 The resultant SPIONs were precipitated with a solvent mixture and then
150

separated by centrifugation and suspended in chloroform at concentrations between 10 and
50 mg/mL
5.2.4. Preparation of Electrospun Fibers
Suspensions for electrospinning were prepared by first dissolving PFS in either 9:1
THF:DMF or DCM and then adding between 0.5 – 4% Triton-X 405 surfactant. The in-house
electrospinning unit, operated at 8 kV, employed a syringe pump and an Al collection target.
The separation distance was held constant at 12 cm, and the solution flow rate varied from 10
to 35 �L/min. Electrospun fibers were collected on the Al plate, as well as on carbon-coated
transmission electron microscopy (TEM) grids adhered to the plate.
5.2.5. Characterization of PFS Nanomaterials
Transmission electron microscopy was performed on a field-emission Hitachi HF2000
microscope operated at 200 kV. Complementary scanning electron microscopy (SEM)
images were collected from specimens sputter-coated with 6 nm of Au/Pd on a JEOL 6400F
field-emission microscope operated at 5 kV. X-ray diffractometry (XRD) was performed on
a Rigaku Smartlab diffractometer with CuKα radiation at 2� angles ranging from 5 to 30° in
0.01° increments at a wavelength (�) of 0.1541 nm.
5.3 Electrospinning and Characterization of PFS Homopolymers
Functional materials can be created through polymer nanocomposites in order to combine
different sets up properties. One way these materials can be formed is through
electrospinning, the application of electrostatic forces to a polymer solution to create
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nano/micro-fibers with characteristics that can be tailored by processing, solution, and
polymer properties. 36 Functionality must be added to the polymer however; often through
the use of dispersed metallic nanoparticles, 37 surface modification, 37 or post-spinning
annealing techniques. 38
The shortcomings of all of these modification techniques, however, are pre- and/or post-
processing steps, as well as aggregation issues associated with metallic materials. A
straightforward solution to these disadvantages is to combine the favorable properties of
metallic and polymer materials into one molecular species: an organometallic polymer. The
goal of this section is to lay the groundwork for electrospinning organometallic polymers and
investigate how the processing conditions during electrospinning influence the resulting fiber
morphology.
As previously mentioned, PFS can be synthesized to exist in either a semi-crystalline or
amorphous state due to the constituent groups on the silicon atom. 7 Symmetrically
substituted constituents, R=R'=Me, tend to impart crystallinity. The electrospinning of
polyferrocenylalkylsilanes was first reported by Chen et al. 39 in the case of PFDMS with a
molecular weight of 95.6 kDa in a 9:1 volume ratio of THF:DMF. Initially, replication of
their solution/processing conditions with PFDMS at 95.6 kDa did not result in similarly
smooth fibers. Concentrations of 15-20% PFDMS in a 9:1 volume ratio of THF:DMF yielded
fibers with a large population of elongated beads along with a wide distribution of fiber
diameters. At concentrations higher than 15 wt% DMF, which was increased to raise the
conductivity of the solution, the PFDMS precipitated out of solution. The addition of
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surfactants have been shown to lower the bead density in electrospun fibers by reducing
surface tension and increasing conductivity. 40 The surfactant, Triton X-405 (polyethylene
glycol tert-octylphenyl ether), was added at concentrations ranging from 0.5 – 4 wt%. At
concentrations higher than 2%, however, the surfactant segregated into domains within the
polymer solution. For PFS systems, the addition of surfactant successfully eliminated the
majority of bead defects and also resulted in a higher mean fiber diameter, as seen in Table
5.2. Electrospinning PFDMS in DCM was also successful, which resulted in a more uniform
distribution of fiber diameters, as well as residual bead. The PFPMS was electrospun from a
9:1 THF:DMF solution with 2% Triton X-405 to delineate differences for the amorphous
polymer. Representative SEM images of all the PFS fibers are displayed for comparison in
Figure 5.2. Fibers were very heavily beaded, even with the addition of surfactant. Widely
varying fiber diameters are responsible for the large standard deviations. A broad range of
fiber diameters result from ‘splaying,’ wherein the axial forces at the tip of the spinneret are
greater than the longitudinal force. For both crystalline and amorphous PFS, a narrow
concentration window exists for spinnability. This may be due, in part, to the low molecular
weights of the polymers employed, ~96 kDa. If the concentration is too low, the electrostatic
forces dominate, and the jet breaks up into droplets or electrospraying ensues. On the other
hand if the concentration is too high, viscous forces dominate, and the polymer solution
begins to solidify at the tip of the syringe with little to no spraying. Another point of interest
is that relatively high concentrations (~20 wt%) are required to produce fibers. This tells us
that below a critical threshold (~18 wt% in the case of PFDMS), the polymer chains produce
153

a relatively dilute solution. As the concentration increases and the distance between chains
decreases, strong interactions between ferrocene groups lead to an abrupt increase in
viscosity to the point where these interactions dominate over electrostatic forces.
To elucidate the arrangement of chains comprising the observed morphologies, XRD has
been used to explore the crystalline structure of the powders and fibers. As a reference point,
XRD analysis of PFDMS powder reveals a strong principal peak at 13.95° 2�, which
corresponds to an interplanar distance (d-spacing) of 0.63 nm. This value is comparable to
that reported 42,43 previously for melt- and solid-state polymerized PFDMS and is indicative
of the distance between adjacent planes of iron atoms, since this atom possesses a large
scattering length. The average crystallite size (�) was obtained from the Scherrer equation,
which can be expressed as 41
� � K�
�cos�
and found to be 5.70 nm. Here, K represents the shape factor (= 0.9), β is the full-width at
half the maximum intensity, �� and � is the Bragg angle. In addition, two smaller peaks located
at 20.09° and 23.39° 2� in Figure 5.4 were also found to exist and are representative of
tetragonal/hexagonal chain packing and indicative of a secondary crystalline phase in the
PFDMS powder. When PFDMS is electrospun, the d-spacing associated with the primary
peak increases slightly, most likely due to distortion of the chains as they undergo flow. The
value of � corresponding to the principal peak, however, increases substantially from 5.74 to
17.20 nm, which is contrary to what is normally observed in electrospun fibers. 42 As the
154
(1)

polymer chains are stretched and become elongated, packing between chains is marginally
compromised and the ferrocene groups are pulled apart allowing for larger crystal sizes (see
Figure 5.5).
Organometallic nanocomposites are often fabricated by physically blending metal or
metal oxide nanoparticles with polymer matrices. In this study, we investigate the effect of
blending SPIONs with PFDMS to ascertain if any coordination takes place between the two
sources of iron. A suspension of 10 nm SPIONs in chloroform was blended with PFDMS and
subsequently electrospun. Information regarding the subsequent crystalline structure is
provided in Table 4 and indicates a negligible increase in d-spacing by 0.01 nm and a
decrease in �. While the SPIONs have little effect on chain packing, the reduction in t
suggests that the SPIONs may be positioned between ferrocene groups so that the chains
wrap around the nanoparticle, as depicted in the schematic diagram displayed in Figure 5.5.
A consequence of this proposed explanation is that a larger SPION is expected to promote a
decrease in d-spacing, since the local radius of curvature due to the SPION would be
reduced. Increasing the SPION diameter to 14 nm does yield a decrease in d-spacing, which
is lower than that of pure PFDMS fiber. Images collected from these systems confirm that
little, if any, SPION aggregation occurs in these systems, which provides additional support
that individual SPIONs are coordinated with the PFDMS chains.
The effect of asymmetrical silicon atom substitution on fiber crystallinity in PFPMS has
also been investigated in powder and fiber form. The powder displays an amorphous halo in
XRD, which is not surprisingly on the basis of previous reports. 43 Of keen interest here are
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the crystalline peaks that become evident in XRD patterns collected from electrospun
PFMPS fibers. These peaks evince that electrospinning can induce crystalline order in an
otherwise amorphous polymer. Two broad peaks are visible at 12.23° and 16.71° 2�,
corresponding to interplanar distances of 0.72 and 0.53 nm, respectively. Corresponding
values of � are 20.10 and 20.50 nm, respectively. A d-spacing of 0.69 nm has been
previously found 44 for the intrachain distance between ferrocene groups in PFEMS. The
present system contains a phenyl ring rather than an ethyl group, in which case a larger d-
spacing would be consistent. In addition, the interchain distance between ferrocene units is
measured to be 0.52 nm, which agrees favorably with our value of 0.53 nm. These results
demonstrate that shear-induced crystallinity can be imparted to otherwise amorphous PFS
polymers and that chain alignment occurs within PFMPS electrospun fibers.
5.4. Phase behavior of binary blends of PFS in Elastomeric Matrices
Blending an AB diblock copolymer with homopolymer A (hA) can lead to improved
emulsifying qualities 45 and interfacial elasticity 46 , since the copolymer effectively behaves as
a macromolecular surfactant 47 . The blocks of such binary blends can form either dry or wet
brushes depending on the value of MhA/MA (hereafter denoted α), where Mi denotes the
molecular weight of species i (i = hA or A). When α is large, a ‘dry brush’ forms due to the
entropic penalty that arises due to the inability of the homopolymer to penetrate the dense
brush created by the copolymer block 48 . In extreme cases, this drives the system to
macrophase-separate into discrete domains of homopolymer and copolymer. When α is of the
order unity or smaller, a ‘wet brush’ develops wherein homopolymer molecules penetrate
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and swell the copolymer brush 49 . The morphologies of these miscible microphase-separated
blends can be controllably altered by modifying 12 α, the blend composition, the copolymer
composition, and the thermodynamic incompatibility (χN), which relates to temperature.
Block-selective solvents serve as the low-molecular-weight limit of homopolymers and
ensure that the compatible block is fully wetted and swollen. Within a block-selective
solvent, diblock copolymer molecules aggregate so that the solvent-incompatible blocks are
sequestered within a core surrounded by a swollen corona of the solvent-compatible block.
Depending on the copolymer composition, the solvated corona typically possesses a larger
volume than the collapsed core and is responsible for the curvature induced at the core-
corona interface 31 .
Although the formation of cylindrical micelles is not unusual, PFS micelles are unique in
that their length is driven by crystallization, in which case it is possible to ‘grow’ the
cylinders by thermal treatment. Such crystal-stabilized nanostructures can lead to
temperature-responsive properties. If a specimen containing elongated crystalline micelles is
heated and the melting point of PFS is surpassed, the micelles will melt and then reform
classical morphologies. In the case of conductive cylinders composed of PFS block
copolymers, heating would eliminate conductive pathways so that the cylinders could serve
as a thermal sensor. In this section, we investigate the behavior of PFS-b-PI micelles in two
different elastomeric matrices that are compatible with the PI corona and how cross-linking
affects the copolymer morphology.
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5.4.1 PFDMS-b-PI/PI Blend by Solvent Casting
According to Manners and co-workers, 31 PFS-b-PI micelles form cylinders in an
isoprene-selective solvent, such as hexane, and they have reported that the lengths of these
micelles range from 100 nm to 1 μm. The copolymer used here was synthesized by living
anionic polymerization in THF at ambient temperature. The molecular weight and block ratio
were determined using GPC and 1 H NMR. Hexane was used to assemble one-dimensional
micellar structures, the morphology of which is governed by varying the PI:PFS block ratio.
When the blocks of PI are much smaller than those of PFS, the block copolymers form tape-
like lamellae. When the length of the PI block far exceeds that of PFS, however, cylinders
form with a crystalline PFS core. 50 Furthermore, these micelles can be grown epitaxially
through the addition of more polymer, which serves to emulate a living polymer growth
reaction. 51 Cylindrical micelles develop by dissolving a PFDMS-b-PI block copolymer in
hexane at elevated temperatures. In the present work, cylinders form when the sample is
heated to 60°C for 30 min and then allowed to cool at ambient temperature. The lengths of
the resultant cylindrical micelles depend on cooling time 51 and can exceed 1 �m with some
end-to-end alignment, as seen in Figure 5.6. Note that the micelles, measuring about 14.9 nm
across, appear to be clustered together with little separation between individual cylinders.
5.4.2. Cross-linking within Polyisoprene
Organic “nanowires” are useable in a wide range of plastic electronic materials such as
solar cells and disposable circuits. 52 Traditional disadvantages of plastic electronics arise due
to high costs and difficult processability. Micelles composed of PFS-b-PI block copolymers
158

have many characteristics that make them favorable candidates for nanowires: a conductive
core, insulting sheath, robust mechanical properties, nanoscale dimensions, and lengths
exceeding 2 �m. The goal of the present work is to crosslink these ‘nanowire’ micelles into
an elastomeric matrix to determine their feasibility for soft electronics. Polyisoprene is a
natural choice for the matrix polymer due to the high degree of cross-linking that can be
achieved because of the large number of unsaturated bonds available in the system. These
bonds can be reacted with sulfur atoms to connect neighboring molecules via a 6-8 chain
sulfur ‘bridge’. This fundamental reaction, known as vulcanization, 53 constitutes the standard
chemical reaction to cross-link natural rubber and has been widely used in the rubber
industry for over 100 years. One complication of this reaction is, however, the high
temperature (~170°C) needed for this reaction to proceed. Since the melting point of PFS is
128°C, the cylinders would evolve at temperatures required for vulcanization.
For tests aimed at blending micelles with low concentrations of PI, 1 wt% PI was added
to the micelle suspension in hexane and allowed to slowly dissolve quiescently. The polymer
suspension was subsequently dropcast onto a TEM grid for direct imaging. In Figure 5.7a, a
TEM image establishes that, at a ratio of 3:100 PFS-b-PI:PI, the dimensions of the cylinders
remain unaffected by the addition of PI. While the cylinders form large aggregates, they are
now separated by an average distance of ~57 nm, which is attributed to the interaction of the
PI matrix with the corona of the cylinders. At the solution concentration employed, the
resultant PI films is insufficiently thin to remain stable and consequently dewets to forms
discrete islands. 54 When the ratio of PFS-b-PI:PI is increased to 3:25 (Figure 5.7b), similar
159

ehavior is observed, and a dewetting pattern reminiscent of spinodal dewetting 55 is uniform
across the grid. In this case, the cylinders are separated further, with an average intermicellar
distance of 94 nm, and are less aggregated than at 3:100. Taken together, these results
confirm that the addition of PI does not affect micelle morphology, but interactions between
PI with the coronae of the micelles promotes greater separation and less aggregation of the
micelles.
Self-assembled PFDMS-b-PI cylinders that had been permitted to grow for 24 h have
been added to a PI solution to disperse conductive organometallic cylinders into a cross-
linked rubber matrix, as schematically depicted in Figure 5.8. The PI was cross-linked
according to straightforward vulcanization with the materials listed in Table 5.4. The
benchmark reaction was allowed to proceed under nitrogen at 110, 120 and 130°C without
micelle addition for 3-4 h, until the surface was no longer tacky. Based on these results, 120
°C was chosen as the reaction temperature since it was still less than the melting point of the
PFS and yielded cross-linked PI. In the subsequent analysis, PI was added to the micelle
suspension in hexane at a ratio of 1:1000 PFS-b-PI:PI and then placed in a vacuum oven.
Due to the solvent that was not previously present, the cross-linking time was increased to 5
h to allow solvent evaporation and subsequent cross-linking to occur. The resultant elastomer
was cryomicrotomed to produce electron-transparent sections so that the internal structure of
the micelles cross-linked in PI could be examined by TEM. The TEM images displayed in
Figure 9 demonstrate that cylindrical micelles are no longer observed and what remains
appear as iron particles ranging in size from 100-500 nm. Possible reasons for PFS-b-PI
160

micelle disassociation are two-fold: (i) the temperature necessary to induce PI cross-linking
was too close to the melting point of the micelles, or (ii) the micelles were not sufficiently
stable to survive the long cross-linking time of the PI. Based on these observations, cross-
linking must occur at significantly lower temperatures, which effectively eliminates PI as the
choice of matrix polymer. In addition, shell-cross-linked PFS-b-PI cylinders will be used to
produce a stable micelle capable of surviving cross-linking.
5.4.3 Shell-Cross-Linking of Cylindrical PI-b-PFS Micelles
To enhance the stability of the cylindrical micelles, the peripheral vinyl groups on the PI
corona chains were cross-linked via a Pt(0)-catalyzed hydrosilylation reaction, which was
conducted using a vinyldimethylsiloxy-terminated polydimethylsiloxane with triethoxysilane
in the presence of Pt(0) catalyst in hexane at ambient temperature. Subsequent 1 H NMR
analysis of the product indicated that all vinyl groups were consumed. The success of the
hydrolsilylation cross-linking reaction was monitored by transferring micelles from a PI-
selective solvent (hexane) to a neutral solvent (THF). 56 In the case of non-cross-linked
samples, the aggregates disassociated in the neutral solvent. When an insufficient amount of
cross-linker or catalyst was used and the shell-cross-linking reaction was incomplete, the
micelles swelled substantially. Completion of the reaction resulted in no discernible swelling
upon transfer to THF and no post-reaction growth. The only change in cylindrical
morphology, as determined by TEM, is that the previously rigid cylinders appear to become
more flexible. The shell-cross-linked micelles used here are 43 nm wide and range in length
up to 1.5 μm, as evidenced by Figure 5.10.
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5.4.4. Cross-linked PVMS as the Matrix Polymer
Poly(vinyl methyl siloxane) (PVMS) was chosen because it satisfied several conditions;
solubility in hexane, an elastomer with a low Tg (~150 K), solubility parameter compatibility
with the PI corona of the cylindrical micelles and the solvent hexane (Table 5.5), the ability
to form a cross-linked network at temperatures below the Tg of the micelles, and the potential
for interesting interactions between the vinyl group and the corona. PVMS was synthesized
in house using an equilibrium-controlled ring opening polymerization (ROP) and step-growth
polymerization methods. Initially, vinyl methyl dichlorosilane was dissolved in a dilute
aqueous solution of HCl to form a mixture of cyclic silicone structures and short oligomeric
vinyl methyl siloxane chains. In the second step, the cyclic structures were separated by
vacuum distillation. The short-chain silanols were then condensed until the desired polymer
chain was reached. This occurred under mild basic conditions resulting in high yields of
PVMS, 95% and greater.
Compatibility between the matrix polymer and the micelle corona is important in order
to assist in good dispersion and maintaining their cylindrical shape. A dilute solution of 5
wt% PVMS containing PFS-b-PI micelles was blended in order to investigate whether the
cylindrical mcielles could hold their shape when dispersed within PVMS, the matrix polymer
and dropcast on a carbon-coated TEM grid for TEM analysis. As evidenced by Figure 5.11,
we see an uneven film of PVMS that is very thick in some areas and thin in others with
thicker PVMS islands in these sections. We see that the majority of the diblock copolymer
micelles exists within the thick PVMS sections, and even in the thinner areas that the
162

micelles are surrounded by an additional ‘corona’ of PVMS. One would expect high
interactions between PI and PVMS based on the fact that both contain reactive groups. This
is evidence of excellent compatibility between the PVMS and the micelles. The most
important result is the cylindrical shape is intact at lengths greater than 750 nm with little
aggregation. In order to cross-link PVMS and form a network, a typical alkoxy condensation
route is utilized by reacting the silanol-terminated PVMS with tetraethoxysilane (TEOS) as
the cross-linker and a tin catalyst with a molar ratio of TEOS:PVMS:catalyst of 1:7:1 in
order to produce films with the highest elasticity. The PVMS solution can be placed on glass
slides and cured either at 60 °C or at room temperature for 30 minutes.
Shell cross-linked micelles at a concentration of 5 mg/mL of hexane were added to the
PVMS:TEOS mixture prior to catalyst addition at ratios of PFS-b-PI:PVMS of 1:150, 1:300,
and 1:600 and cured at both ambient temperature and 60 °C for 30 minutes. Films were
microtomed below the Tg of PVMS (-128 °C). After probing the interior of the PVMS films,
we see the appearance of cylindrical micelles, spherical micelles, and iron clusters, as
depicted in Figure 12. As far as the structure of any remaining micelles, there is not an
appreciable difference of the morphology in the ambient vs. 60°C cured films. This
demonstrates that the micelles are not affected by the curing temperature, which is well
below the Tg. Lengths of cylindrical micelles range from 20-400 nm, which is a drastic
reduction from the length of PFS-b-PI in solution which exceeds one micron. The reason for
this is explored by exposing by heating and agitating the micelle solution at conditions that
imitate the crosslinking reaction conditions.
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Shell-cross-linked PFS-b-PI were dropcast from hexane onto a carbon coated TEM grid
were found to have an average length of 810 ± 300 nm. Some affinity exists between
micelles both in an end-to-end and side-to-side fashion. When the solution is hand-agitated,
similar to the conditions by which its necessary to mix the catalyst into the PVMS solution,
we see that the average length is decreased to 620 ± 290 nm. Likewise less affinity exists
between the micelles. Some of the cylindrical micelles also collapse into spherical micelles,
as seen in Figure 5.13b. When the solution is heated in an oven at 60 °C for 15 minutes, we
see little change in the average length (840 ± 380). However the largest change is that we see
networks of micelles formed and much more curvature. These results are indicative that
heating the micelle solution prior to cross-linking may achieve a more continuous network
throughout the PVMS. These results are summed up in Figure 5.13.
Complementary SEM images were also taken on both the surface of the PVMS films and
also the cross-sections. Of interest primarily was the 1:150 phase-separated film. An image
of the the surface shows large protrusions from the film, edx reveals these are largely
composed of silicon. Due to the fact that the PVMS solution was not stirred prior to cross-
linking, it is not unusual that some parts of the solution cross-linked quicker before the
matrix did as a whole and exist as flocs in the matrix. These parts are protruding from the
surface of the film. In addition, cross-sections of these films were taken by exposing the
films to liquid nitrogen and then immediately fracturing. Figure 5.14 shows the SEM
micrographs of changes in surface density as well as the fractured surface.
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5.5. Conclusions
In this work PFS is investigated as electrospun fibers and PFS-b-PI cylindrical micelles
are cross-linked within an elastomeric PVMS matrix. In order to produce uniform, beadles
fibers from electrospinning it is necessary to include the surfactant Triton-X in order to
reduce the surface tension in the solution. Electrospinnning PFDMS demonstrates an increase
in interplanar distance between parallel ferrocene chains due to the elongational forces during
electrospinning. Adding in small iron oxide nanoparticles (~10 nm) further increases this
interplanar distance since the polymer chains orient around the nanoparticles. In addition,
although the completely polymer PFMPS is completely amorphous, we see electrospining-
induced crystallinity. Cylindrical micelles of PFS-b-PI can be shell cross-linked and
successfully maintain their shape when cross-linked in a PVMS matrix. Although the length
sof these molecules decrease from greater than 1.5 microns in solution to less than 400 nm
when dispersed in the PVMS matrix this demonstrates a great starting point for soft nanowire
devices utilizing the crystalline core of PFS block copolymers.
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Tables
Table 5.1: Characteristics of the (Co)Polymers Used in this Study
Polymer Mn (kDa) PDI ΦPFS
Poly(ferrocenyldimethylsilane) 95.6 1.42 1.00
Poly(ferrocenylmethylphenylsilane) 109.6 1.12 1.00
Poly(ferrocenylethylmethylsilane) 38.0 1.10 1.00
Poly(2-vinylpyridine)-bpoly(ferrocenylethylmethylsilane)
20.0 1.17 0.20
Poly(2-vinylpyridine)-bpoly(ferrocenydimethylsilane)
20.0 1.16 0.15
Poly(ferrocenyldimethylsilane)-bpoly(isoprene)
32.0 1.03 0.14
Poly(ferrocenyldimethylsilane)-bpoly(isoprene)
(cross-linked corona)
56.8 1.11 0.09
Poly(2-vinyl pyridine) 78.5 — —
Poly(isoprene) 80.0 1.01 —
Poly(ε-caprolactone) 80.0 — —
Poly(vinyl methoxysilane) 30.0 — —
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Polymer Solvent
15%
PFDMS
15%
PFDMS
20%
PFDMS
18%
PFPMS
Table 5.2: PFDMS Fiber Diameters
9:1 v:v
THF:DMF
9:1 v:v
THF:DMF
167
2 vol %
Surfactant
Mean
Diameter (nm)
St Dev
Yes 1320 980
No 1800 400
Dichloromethane Yes 610 220
9:1 v:v
THF:DMF
Yes 1260 540

Table 5.3: X-ray diffraction bragg angle, d-spacing, and average crystallite size for 96 kDa
poly(ferryocenydimethylsilane) powder and electrospun fibers
PFDMS Powder
PFDMS Fiber
PFDMS Fiber
+ 14nm NP
PFDMS Fiber
+ 10nm NP
Angle
(°2Ɵ)
168
d-spacing
(nm)
��(nm)�
13.95 0.63 5.74
20.09 0.44 19.90
23.39 0.38 6.69
13.55 0.65 17.20
18.55 0.47 16.10
25.62 0.35 2.70
13.62 0.65 1.69
23.43 0.37 30.50
13.55 6.53 2.47

Table 5.4: Components utilized in the vulcanization of poly(isoprene)
169

Table 5.5: Solubility parameters of polymers relative to the solvent n-hexane.
Component
170
Solubility
Parameter (δ)
PI 16.5
PVMS 15.1
Hexane 14.9

Figures
Fe
tmeda
Fe
BuLi
Fe
a.
Si + BuLi Fe
Li
. tmeda +
Li
Si
Li
Bu
+ Si
Fe
Figure 5.1: Schematic of PFEMS synthesis in which the noted molecules are a)
ferrocenophane b) ethylmethylsilaferrocenophane and c) poly (ferrocenylethylmethylsilane).
171
Si
Cl
Cl
2 LiCl
BuLi
Fe
Fe
c.
Si
b.
Si
n

a.
c.
b.
d.
Figure 5.2. SEM micrographs of a) 15% PFDMS in THF:DMF without surfactant b) 15%
PFDMS in THF:DMF with surfactant c) 20% PFDMS in DCM with surfactant and d) 18%
PFPMS in THF:DMF with surfactant. All scale bars represent 20 μm.
172

a
.
c
b.
Figure 5.3. PS-b-PFS-b-P2VP lithographic template used for the preparation of Nanoscale
magnetic dots. a) Phase-separation of the triblock in the bulk. b) Hollow PFS cylinders are
formed after etching because it has a selective resistance. c) Profile of the hollow PFS
cylinders. (Used with permission by Jessica Gwyther from BristolUniversity.)
173
PFS hollow cylinders
Etched
P2VP
matrix
Etched
PS
holes

10 15 20 25
2�
Figure 5.4. XRD curves from 2θ = 7 – 25 °for PFDMS powder, fibers, and fibers with both
10 and 14 nm iron oxide nanoparticles.
174
Fiber + 14 nm NP
Fiber + 10 nm NP
PFDMS Fiber
PFDMS Powder
Intensity

a
.
b.
c.
d.
Figure 5.5 Schematic demonstrating the chains of PFDMS and corresponding d-spacing
between adjacent ferrocene units in a) powder form, b) in electrospun fibers, c) in
electrospun fibers with larger (~14 nm) iron oxide nanoparticles, and d) in electrospun fibers
with smaller (~10 nm) iron oxide nanoparticles.
175

Figure 5.6. TEM micrographs of PFS-b-PI micelles dropcast from a 1 mg/mL hexane
solutiononto a carbon-coated TEM grid with an average width of 14.9 nm and lengths
exceeding one micron.
176

a.
Figure 5.7. TEM micrographs of PFS-b-PI micelles blended at a ratio of a) 3:100 with PI and
b) 3:25 with PI.
177

Add micelles and apply
heat
Figure 5.8. Schematic of the vulcanization and micelle crosslinking technique.
178

Figure 5.9. TEM micrographs of 1:1000 PFS-b-PI:PI vulcanized at ~120 °C for 5 hours
demonstrating a complete dissolution of the micelles with only iron nanoparticles remaining.
179

Figure 5.10. TEM micrographs of shell cross-linked PFS-b-PI with an average width of 43
nm and lengths exceeding 1.5 microns dropcast from a hexane solution.
180

Figure 5.11. TEM micrographs of PFS-b-PI micelles blended with a 5 wt% PI solution in
hexane and dropcast onto a carbon-coated TEM grid.
181

Figure 5.12. TEM micrographs of microtomed PFS-b-PI micelles blended with PVMS at
ratios of 1:300 and 1:600 at thicknesses of ~120 nm.
182

a.
c.
b.
Figure 5.13. TEM micrographs shell cross-linked PFS-b-PI dropcast from a hexane solution
as a) a control, b) heated to 70 °C for 30 minutes and c) stirred for several minutes
mimicking conditions during the cross-linking of the PVMS solution.
183

a.
b.
Figure 5.14. SEM micrographs of a PVMS cross-linked film containing PFS-b-PI micelles a)
looking down the surface from the cross-section and b) of the fractured cross-section where
the scalebar of the inset refers to 200 μm.
184

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188

6.1 Conclusions
CHAPTER VI
Conclusions and Future Work
In this dissertation, we determined how to controllably align and position metal-
containing nanomaterials within electrospun polymer fibers, develop control over
electrospinning parameters to create organic-inorganic interactions, and maximize
functionalities of the organometallic filler materials. Firstly, we demonstrated the concept of
hierarchical alignment of fibers containing aligned gold nanorods over macroscopic
dimensions through electrospinning . We hypothesized that control over the location and
alignment of nanoparticles within electrospun fibers can be achieved through the
straightforward use of an external electromagnetic field applied during electrospinning.
Likewise, thermodynamic incompatible polymer blends were prepared in order to control the
spatial location of superparamagnetic iron oxide nanoparticles within electrospun fibers.
Lastly, organometallic polymers were utilized to create functional nanocomposites that
married the highly desirable properties of metals with those of polymers. This approach is
one that can be applied to a wide variety of both filler materials and polymer matrix systems,
which will broadly contribute to the field of material science.
189

6.1.1 Long-Range Alignment of Gold Nanorods in Electrospun Polymer Nano/Microfibers
The nanoscale orientation of GNRs was achieved with scalable, macroscopic order a
distance of several centimeters. Here, GNRs with an aspect ratio of 3.1 exhibit excellent
alignment with their longitudinal axes parallel to the fiber axis n for electrospun polymer
nano/microfibers with diameters of 40-600 nm, and they maintain substantial alignment in
microfibers measuring up to 3000 nm in diameter. Loading of GNRs can be varied with no
discernible impact on the net degree of alignment while fiber diameter has a more direct
correlation. Electron diffraction measurements of the aligned GNRs confirm preferred
orientation of the {100} and {111} GNR planes. Optical absorbance spectroscopy
measurements performed on macroscopically aligned electrospun fibers containing aligned
GNRs demonstrate that the longitudinal surface plasmon resonance bands are polarization
dependent and display maximum absorption when the polarizer is parallel to n.
6.1.2 Magnetic Field-Induced Alignment of Nanoparticles in Electrospun Microfibers
Nanoscale alignment of superparamagnetic iron oxide nanoparticles (SPIONs) with
diameters ranging from 12 to 18 nm has been achieved through the coupling of an external
magnetic field with the electric field in the electrospinning process. Microfibers containing
one-dimensional arrays of SPIONs extending beyond one micron in length result from an
perpendicular electromagnetic field of 26 mT. In the present system investigated, a SPION
concentration of greater than 0.05 vol% is required to induce discernible alignment.
Moreover, alignment of the electrospun microfibers through the use of various established
methods can yield nanocomposites with multiscale (i.e., nanoscale and macroscale)
190

anisotropic properties. Complementary superconducting quantum interference device
(SQUID) measurements reveal that the saturation magnetization is significantly lower for
SPIONs in electrospun fibers with or without magnetic field-induced alignment than for
unembedded SPIONs. The mean magnetic moment increases with improving alignment,
however,, which demonstrates that nanoscale alignment of SPIONs affects their intrinsic
physical properties.
6.1.3 Using Polymer Blend Morphology to Position Ligand-Functionalized Nanoparticles
in Electrospun Polymer Microfibers
Blends of hydrophobic (P2VP) and hydrophilic polymers (PEO) have been prepared
to discern the feasibility of controlling the spatial location of SPIONs within electrospun
fibers on the basis of thermodynamic compatibility. In this case, a core-sheath structure
naturally forms with the hydrophobic SPIONs sequestered in one preferential phase at 30
wt% PEO. For both a pure PEO solution and a pure P2VP solution a decrease in zero shear
viscosity was found upon the addition of SPIONs. Interfacial distances ranged from 26 – 109
nm demonstrating diffuse interfaces between the two immiscible polymers. X-ray diffraction
(XRD) confirms that PEO crystals becomes more aligned upon SPION addition suggesting
that the nanoparticles may reside at the interface between the PEO and P2VP phases.
6.1.4. Nanostructured Organometallic Polymer Systems Containing Poly(ferrocenylsilanes)
Lastly, poly(ferrocenylsilanes) (PFS) are investigated as the fiber-forming polymer in
an electrospinning system. Uniform fibers are proven to be formed from the addition of a
surfactant to a solution of polyferrodimethylsilane (PFDMS) in dichloromethane.
191

Electrospun PFDMS fibers demonstrates an increase in interplanar distance between parallel
ferrocene chains due to the elongational forces during electrospinning when compared to the
pure powder. The addition of small iron oxide nanoparticles (~10 nm) further increases this
interplanar distance where we hypothesize that the polymer chains orient around the
nanoparticles. In addition, although the polymer poly(ferromethylphenylsilane) (PFMPS) is
completely amorphous, we see electrospining-induced crystallinity. Cylindrical micelles of
PFS-b-PI can be shell cross-linked and successfully maintain their shape when cross-linked
in a poly(vinyl methoxysilane) (PVMS) matrix. Although the length sof these molecules
decrease from greater than 1.5 microns in solution to less than 400 nm when dispersed in the
PVMS matrix this demonstrates a great starting point for soft nanowire devices utilizing the
crystalline core of PFS block copolymers.
6.2 Recommendations for future work
6.2.1 Gold Nanorod Alignment through Electrospun Fiber Degradation
Utilization of polymer matrices to align gold nanorods can produce macroscale
alignment but a method to self-assemble gold nanorods in solution or on a substrate without
utilizing a hierarchical structure is needed. In this work the average deviant angle between
the GNR axis and the fiber axis for small fibers (

a 35-46% decrease in molecular weight. 2 This same principal could be applied to PEO
nanofibers at various wavelengths and times to see if it would be possible to completely
degrade the PEO matrix. This would require aligned fibers with very small fiber diameters;
in addition another variable that could be explored is the effect UV exposure on a monolayer
of fibers versus multilayers of fibers. Optical absorbance spectra could be collected at
various times of UV exposure to remove the troublesome polymer scattering in the spectra
and provide a more precise way to monitor GNR alignment. Dissolution of PEO by water or
a solvent such as chloroform could also be attempted to see if aligned GNRs could result. If
either of these techniques were successful it could be possible to align GNRs onto a substrate
used for biomedical imaging. 3
6.2.2 Field Uniformity in Magnetic-Assisted Electrospinning
In this work, alignment of SPIONs in electrospun fibers was achieved through the use
of an external u-shaped electromagnet in conjunction with the electric field during
electrospinning. However, alignment was not perfect due to several imperfections; namely,
whipping of the jet and a lack of magnetic field uniformity due to the shape of the
electromagnet used. An electromagnet utilizing a hollow cylinder, wrapped in magnetic wire
connected to a voltage source would provide a much more uniform field and allow the
spinneret needle to be placed ‘inside’ the magnet. In order to achieve a magnetic field
equivalent to that of the electromagnet, 26 mT, hundreds of coils would have to be placed
over the hollow cylinder. Another follow-up experiment for this work would be to determine
a way to decrease the interparticular distance between SPIONs in the electrospun fibers so
193

that they act as a single ‘nanowire’ and not as individual particles. This could lead to some
very interesting superparamagnetic properties; such as a larger remanence magnetization and
coercivity for aligned samples. 4 This could be investigated by using particles that are
surface-functionalized with a different ligand, by altering the length of the ligand on the
nanoparticle surface, or look at another type of magnetic nanoparticle such as cobalt
nanoparticles. 5
6.2.3 Poly(ferrocenylsilane) Cylindrical Micelles Oriented within Electropun Fibers
The focus of this thesis has been on controlling the spatial location of nanoscale
objects within electrospun fibers in order to create functional nanocomposites. A natural
extension would be to spin PFS block copolymer cylindrical micelles to see if they orient
within electorspun fibers and investigate the resulting properties. Self-assembly of PFS-b-
poly(2-vinyl pyridine) (P2VP) has been proven to form cylindrical micelles in alcohol
solvents 6 such as ethanol and isopropanol. If these micelles could be formed in a solvent
with a higher dielectric constant, such as dimethylformamide (DMF), then it would be
suitable for electrospinning. The self-assembled cylinders could then be blended with P2VP
to explore the interactions between a diblock and homopolymer binary blend behave under
longitudinal forces during electrospinning. The time scales at which self-assembly occurs
can be investigated by electrospinning both assembled and unassembled micelles. If the
assembled micelles do in fact stay intact during electrospinning, this will be an indication of
their robustness and also may stabilize the electrospun fibers much like carbon nanotubes 7 .
194

The location of the micelle within the fiber (on the surface, random, oriented) will also give
information about the diblock/hompolymer interactions.
PFDMS-b-P2VP, at a concentration of 2 mg/mL, was dissolved in DMF, a P2VP
selective solvent, and after 24 hours formed cylinders with diameters ranging from 100 -
1000 nm, which can be seen in the TEM images in Figure 6.1. This has not been previously
reported in literature. P2VP (200 kDa) was electrospun from DMF at initial concentrations
ranging from 14-20%. At lower concentrations, very thin fibers were formed but with a large
proportion of beads. Electrospinning at 20% produced thin fibers, with a mean diameter of
260 ± 170 nm, and few bead defects. To ensure this criteria was met, instead of increasing
the concentration of P2VP, which has been shown to reduce bead defects 8 , the solvent was
changed to volume ratio of 9:1 DMF:ethanol. The addition of ethanol reduces the surface
tension of the jet and thus also reduces bead defects. Initial concentrations of 21% P2VP
prepared in a 9:1 DMF:ethanol solution was blended with 0.8 mg of PFDMS-b-P2VP in
DMF. The solution was gently shaken to disperse the two systems. The addition of the
micelles did not disturb the solution during electrospinning, fibers without beads were
formed with a mean diameter of 300 ± 190 nm, very similar to the pure P2VP solution.
Energy-dispersive x-ray spectroscopy (EDX) results confirmed that no iron was present on
the surface and scanning electron microscopy (SEM) revealed that the fiber surfaces did not
appear different before and after micelle addition (Figure 6.2). It was not possible to see the
micelles within the fibers through transmission electron microscopy (TEM), even when the
fibers were very small and more transparent to the electron beam as seen in Figure 6.3. To
195

emedy this, the electrospun fibers were aligned and coated with 5 nm of Pt/Pd, and then
cured within an epoxy matrix for microtoming. Fiber cross-sections of ~120 nm were cut on
a room temperature microtome and the internal structure was probed via TEM. It appears as
if the micelles have stayed intact during the electrospinning process, (Figure 6.4), however
the structures depicted are not oriented parallel to the fiber axis (since cross-sections would
only present the top face of the cylinders, or just a sphere). The micelles are likely
orthogonal to the fiber axis. In order to probe this result further, higher concentrations of
micelles should be included in the solution to be electrospun and very small fiber diameters
should be prepared. It would then be beneficial to microtome perpendicular to the fiber
surface, since we would expect the cylinder-like rods to be aligned to the fiber axis, as is the
case with CNTs. 9 In conclusion, extending the topic of this dissertation to the case of block
copolymer micelles within electrospun fibers would be an excellent inclusion and would
further probe the dynamics between a diblock copolymer/hompolymer blend in an
electrospun fiber.
196

Figures
Figure 6.1. Dropcast PFDMS-b-P2VP cylindrical micelles onto a carbon coated TEM grid
from DMF.
197

a.
b.
Figure 6.2. SEM micrographs of 32 wt% P2VP fibers electrospun from 9:1 DMF:THF a)
without micelle addition and b) with PFDMS-b-P2VP micelles.
198

Figure 6.3. TEM micrograph of 32 wt% P2VP fibers electrospun with PFDMS-b-
P2VP micelles.
199

Figure 6.4. TEM micrographs of microtomed P2VP fibers containing PFDMS-b-P2VP
micelles.
200

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201

APPENDIX
202

APPENDIX I
Responsive PET Nano/Microfibers via Surface-Initiated Polymerization
Abstract
A. Evren Özçam, Kristen E. Roskov, Jan Genzer and Richard J. Spontak
Poly(ethylene terephthalate) (PET) is one of the most important thermoplastics in
ubiquitous use today due to its mechanical properties, clarity, solvent resistance and
recyclability. In this work, we functionalize the surface of electrospun PET microfibers by
growing poly(N-isopropylacrylamide) (PNIPAAm) brushes through a chemical sequence that
avoids PET degradation to generate thermoresponsive microfibers that remain mechanically
robust. Amidation of deposited 3-aminopropyltriethoxysilane, followed by hydrolysis, yields
silanol groups that permit surface attachment of initiator molecules, which can be used to
grow PNIPAAm via "grafting from" atom-transfer radical polymerization. Spectroscopic
analyses performed after each step confirm the expected reaction and the ultimate growth of
PNIPAAm brushes. Water contact-angle measurements conducted at temperatures below and
above the lower critical solution temperature of PNIPAAm, coupled with adsorption of Au
nanoparticles from aqueous suspension, demonstrate that the brushes retain their reversible
thermoresponsive nature, thereby making PNIPAAm-functionalized PET microfibers
suitable for filtration media, tissue scaffolds, delivery vehicles, and sensors requiring robust
microfibers.
203

Introduction
Electrospinning is an emerging fabrication technique capable of generating solid polymer
fibers that range from tens of nanometers to several microns in diameter. Such
nano/microfibers are of fundamental and technological interest due to their high surface-to-
volume ratio. During wet electrospinning, a polymer solution of sufficiently high viscosity
and conductivity is subjected to an electric field. When the electrostatic forces overcome
surface tension, a charged jet emitted from the tip of a Taylor cone 1 undergoes a whipping
action 2 wherein the solvent evaporates, and is subsequently collected as a dry, randomly
oriented fiber mat on a grounded collector plate. This process strategy is appealing due to the
simple setup required and the ability to tailor fiber characteristics with relative ease. 3
Although the morphology of electrospun nano/ microfibers is desirable, they tend to lack the
functionality that is sought in contemporary applications. One way to overcome this
deficiency is by developing multicomponent nano/ microfibers, in which the fiber-forming
polymer is modified with one or more species designed to enhance targeted properties. 4
Surface-active compounds added to the polymer solution prior to electrospinning may,
however, remain trapped within the resultant fiber upon solidification and thus exhibit
substantially reduced activity. 5 While antibacterial biocides incorporated in this fashion lose
much of their efficacy, 6 quaternary ammonium species covalently bonded to as-spun fibers
can create a permanent antibacterial surface. 7 Alternatively, polarizable antibacterial
copolymers co-dissolved with the fiber-forming polymer can be brought to the fiber surface,
where they remain anchored in place, by the electric field during electrospinning. 8 Recently,
204

Agarwal et al. 9 have surveyed chemical routes by which to modify and functionalize the
surface of electrospun nanofibers for diverse applications ranging from functional textiles,
catalyst supports and ion-exchange membranes to drug delivery and tissue engineering.
Polymers such as poly(ethylene terephthalate) (PET), which is widely known for its
mechanical strength, transparency and solvent resistance, tend to possess a hydrophobic
surface and a low surface energy, 10 in which case electrospun nano/microfibers require post-
treatment so that chemically-active species are positioned on the fiber surface. Methods by
which to achieve such surface functionalization include UV treatment, 6 mineralization, 11
core-shell formation, 12 chemical vapor deposition 13 or inclusion of reactive compounds. 14
Once these chemically-active groups are available, covalent bonding, 15 immobilization 16 or
electrostatic interactions 17 can be used to introduce functional moieties to the fiber surface
without adversely affecting the bulk fiber properties. While surface modification could
permit the use of electrospun PET 18 nano/microfibers in filtration media, 19 protective
textiles, 20 tissue scaffolds, 21 and drug-delivery vehicles, 22 most of the modification
approaches listed above purposefully or inadvertently promote PET degradation. Thus, the
conditions by which surface modification is conducted must be monitored carefully to avoid
compromising the bulk properties of PET.
Grafting polymer brushes represents an alternative approach by which to modify and
control the surface properties of materials. 23 Numerous studies reported on surface-initiated
grafting on surfaces of various geometries with a plethora of different monomers by
employing numerous polymerization routes. Poly(N-isopropylacrylamide) (PNIPAAm) is
205

solely considered because of its thermoresponsive nature 24 (it possesses a lower critical
solution temperature, Tc, in water at ≈32°C). Prior efforts to polymerize styrene 25 and
NIPAAm 26 on flat PET surfaces have relied on different means of activating the PET surface
(e.g., saponification, plasma treatment and aminolysis) for the purpose of attaching initiators.
The major drawback of such treatments, however, is that they may increase surface
roughness by degrading PET, which is of concern with regard to electrospun PET
nano/microfibers. Independent studies 27 have confirmed that 3-aminopropyltriethoxysilane
(APTES) can be used to functionalize the surface of PET via amidation with negligible
degradation of the parent PET material. Unlike short alkyl amines (which can diffuse into
and react throughout, and thus weaken, PET 28 ), the bulky triethoxysilane group on APTES
hinders diffusion, changes its chemical nature upon amidation and creates a barrier by
restricting the diffusion of other APTES molecules. Moreover, since the ethoxysilane groups
of APTES are exposed at the polymer/air interface, hydrolysis of triethoxysilane yields
silanol groups that facilitate initiator attachment.
Thermoresponsive PNIPAAm brushes on electrospun fibers have been recently reported.
For instance, Brandl et al. 29 describe the synthesis of a copolymer of 2-hydroxyethyl
methacrylate (HEMA) and methyl methacrylate (MMA) and its post-polymerization
modification with 2-bromoisobutyrylbromide to prepare a macroinitiator. They claim that
electrospinning of the macroinitiator and subsequent polymerization of NIPAAm results in
thermoresponsive polymer brushes. The disadvantage of this technique is that the location of
the "active" initiator group in the fiber depends on the dielectrophoretic forces, polarizability
206

contrast and surface tension of the comonomers, which invariably reduces the concentration
of "active" initiator centers on the fiber surface. The presence of ester groups between the
butyrylbromide group of the initiator and hydroxyethyl group of HEMA likewise yields
hydrolyticly unstable bonds at pH values greater than 8 and lower then 5. Furthermore, the
presence of HEMA comonomer on the macroiniatitor may result in swelling and absorption
of NIPAAm monomer by the electrospun fiber in the aqueous polymerization medium.
Similarly, Fu et al. 30 have synthesized and electrospun a copolymer of 4-vinylbenzylchloride
and glycidyl methacrylate. Subsequent modification of the electrospun fibers with sodium
azide produces azide surface groups that can be coupled with alkyne-functionalized
PNIPAAm chains via a click reaction to generate PNIPAAm surface chains that affect the
wettability of the fibers. Such grafting of PNIPAAm brushes on electrospun fibers would be
necessarily low because of the sparse population of active azide groups on the fiber surface
and the accompanying steric hindrance caused by the "grafting to" polymerization technique.
In this work, we report a robust and universal way of preparing functional and
thermoresponsive PNIPAAm brushes that are covalently attached to electrospun PET fibers.
First, electrospun PET microfibers are modified with APTES to create surface-bound
hydroxyl groups for the attachment of [11-(2-bromo-2-methyl)propionyloxy]
undecyltrichlorosilane (BMPUS), which serves as an ATRP initiator for the polymerization
of NIPAAm. Several analytical techniques are employed to (i) characterize the properties of
as-spun and post-modified PET microfibers and (ii) follow the polymerization of NIPAAm
via ATRP. In addition, we investigate the thermoresponsive nature of PNIPAAm-decorated
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PET microfibers by attaching Au nanoparticles at temperatures above and below the Tc of
PNIPAAm.
Experimental
Food-grade recycled PET flakes were kindly supplied by the United Resource Recovery
Corp. (Spartanburg, SC). The HFIP was obtained from Oakwood Products Inc. (Estill, SC),
and anhydrous toluene, 2-chlorophenol, APTES, NIPAAm, copper I bromide (CuBr), copper
II bromide (CuBr2), and N,N,N',N',N"-pentamethyldiethylenetriamine (PMDETA) were all
purchased from Sigma-Aldrich and used as-received. The citrate-stabilized Au
nanoparticles 29 (diameter = 16.9±1.8 nm) and BMPUS initiator 30 were prepared as described
earlier. The PET flakes were dissolved in HFIP at different concentrations and electrospun at
ambient temperature and 10 kV to generate microfibers varying in diameter. Thin films of
PET measuring 12 and 180 nm thick, as discerned by ellipsometry (v.i.), were spun-cast at
25ºC on silicon wafers from 0.5 and 3.0% (w/w) solutions, respectively, in 2-chlorophenol.
Microfiber mats and thin films were stored under vacuum for at least 48 h prior to use to
remove entrapped solvent.
The APTES was deposited on the PET microfibers and thin films by exposing the
samples to 1% (v/v) APTES/anhydrous toluene solutions for 24 h at ambient temperature,
followed by sonication in toluene for 10 min to remove loosely adsorbed APTES molecules.
The ethoxysilane groups of the surface-anchored APTES molecules were hydrolyzed in
acidic water (pH ≈ 4.5-5.0) for 6 h at ambient temperature, and then the fiber mats were
washed with copious amount of water. After drying the samples under reduced pressure,
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BMPUS was deposited on the PET-SiOH surfaces by established protocols. 31 The PNIPAAm
brushes were subsequently grown from PET-SiOH surfaces by ATRP of NIPAAm, as
described elsewhere. 29 Briefly, 6.30 g NIPAAm was dissolved in a mixture of 4.86 g
methanol and 6.30 g water in an argon-purged Schlenk flask, and oxygen was removed via
three freeze-thaw cycles. After removal of oxygen, PMDETA (0.56 g), CuBr (0.16 g) and
CuBr2 (0.016 g) were added to the solution prior to an additional freeze-thaw cycle. The
Schlenk flask was tightly sealed and transferred to an argon-purged glove box. Microfiber
mats and thin films of PET were submersed in the solution for 8 h at ambient temperature,
after which they were removed, promptly rinsed with methanol and deionized water, and
then sonicated in deionized water for 10 min.
The thickness of the thin PET films was measured by variable-angle spectroscopic
ellipsometry (J.A. Woollam) at a 70º incidence angle before and after each modification step
to discern the PNIPAAm brush height. Surface chemical composition was monitored by XPS
performed on a Kratos Analytical AXIS ULTRA spectrometer at a take-off angle of 90º. The
FTIR analysis of the PET microfibers was conducted in transmission mode on a Nicolet 6700
spectrometer after embedding the microfiber mats in potassium bromide pellets. For each
sample, 1024 scans were acquired after background correction at a resolution of 4 cm -1 .
Resultant XPS and FTIR spectra were analyzed using the Vision and Omnic Specta software
suites, respectively. The thermoresponsive behavior of PET and PET-PNIPAAm microfibers
was interrogated by measuring the WCA at different temperatures via the sessile drop
technique on a Ramé-Hart Model 100-00 instrument. As-spun and modified PET microfibers
209

were coated with ≈16 nm of Au, and their diameter and surface morphology were examined
by field-emission SEM performed on a JEOL 6400F electron microscope operated at 5 kV.
Results and Discussion
The diameters of electrospun PET microfibers, prepared according to the protocol
provided in the Experimental section and measured by scanning electron microscopy (SEM),
are 450±100, 800±200 and 1200±300 nm for 6, 8 and 10% (w/w) solutions, respectively, of
PET in hexafluoroisopropanol (HFIP). The surfaces of unmodified PET microfibers
consistently appear smooth (cf. Figure 1) with some slight dimpling observed occasionally
along the fiber axis. Microfibers modified with thermoresponsive PNIPAAm brushes have
been generated in a sequence of four steps, which are depicted schematically in Figure 1.
Briefly, APTES molecules are attached to the PET surface via aminolysis between PET and
the primary amine of APTES. Next, the ethoxysilane groups on APTES are hydrolyzed to
generate silanol groups for BMPUS attachment. Finally, PNIPAAm brushes are grown
directly from the PET microfiber surface. A second SEM image displaying PET microfibers
modified with PNIPAAm brushes is included for comparison in Figure 1 to demonstrate that
these microfibers appear marginally rougher than the as-spun microfibers at the end of the
modification and brush growth process. The difference in microfiber morphology is almost
indiscernible and the PNIPAAm brushes on spin-coated PET films on silicon wafers appears
smooth, combination of these verifies that the brush is uniformly distributed on the surface of
the microfibers. Below, we provide a detailed assessment of each of the steps in this
polymerization sequence.
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In Figure 2, Fourier-transform infrared (FTIR) spectra are presented for three materials:
(a) as-spun microfibers (PET), (b) APTES-modified PET microfibers following hydrolysis
(PET-SiOH) and (c) PET microfibers with PNIPAAm brushes (PET-PNIPAAm). The
appearance of new peaks located at 1650 cm -1 (amide I band) 1550 cm -1 (amide II band),
1470 cm -1 , and 3300 cm -1 in Figure 2b is due to the formation of secondary amide groups,
thereby confirming the grafting of APTES to the PET microfiber surface. Previous reports 27
regarding the surface modification of PET with APTES could not detect the amidation
reaction via FTIR due to a very low signal-to-noise ratio. Detection of these groups by FTIR
in the present work is attributed to the large surface area afforded by the microfibers.
Successful attachment of APTES can also be inferred from the surface properties of modified
microfibers upon exposure to acidic water, which promotes hydrolysis of the ethoxysilane
groups to silanol groups. Resulting changes in static water contact angle (WCA) and
specimen thickness are measured on flat PET films spun-cast on silicon wafer.
Corresponding values of WCA for films of PET-SiOH and PET after hydrolysis are 50�0.8˚
and 71�0.8˚, respectively, whereas that for untreated PET is 75�0.2˚. In addition, the results
of X-ray photoelectron spectroscopy (XPS) measurements shown in Figure 3a reveal the
existence of a small N1s, Si2s and Si2p peaks at 400, 150 and 100 eV, respectively. These
peaks correspond to 0.6 atom% N and 1.1 atom% Si from the hydrolyzed APTES on the
PET-SiOH surface. In the next step, BMPUS molecules are attached to the PET-SiOH
surface (cf. Figure 1) to serve as initiator centers for the "grafting from" polymerization of
NIPAAm.
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Subsequent growth of PNIPAAm brushes from the initiator centers at the fiber surface is
verified by the FTIR and XPS spectra presented in Figures 2c and 3b, respectively. The
characteristic secondary amide IR vibrations located at 1650, 1550, 1470, and 3300 cm -1 are
the most pronounced for PET-PNIPAAm microfibers. In addition, the appearance of a
relatively intense N1s peak at 400 eV in Figure 3b indicates an elevated concentration of N,
which is consistent with the presence of PNIPAAm brushes. Quantitation of this spectrum
yields the following atomic concentrations: 76.8�0.4% C, 11.6�0.5% N and 11.6�0.3% O.
These values agree favorably with theoretical concentrations (75.0% C, 12.5% N and 12.5%
O) obtained from the chemical structure of PNIPAAm. The high-resolution C1s spectra
included in the insets of Figures 3a and 3b likewise demonstrate that the PNIPAAm brushes
cover the PET film surface. In Figure 3a, the spectrum displays peaks at 289.0 and 286.6 eV
corresponding to O-C=O and C-O functionalities, respectively. These signature peaks for
PET disappear upon growth of the PNIPAAm brushes, which are responsible for a new peak
at 287.8 eV (N-C=O groups) and a shoulder at 286.1 eV (C-N bonds). 32 Since the XPS
fingerprint for PET is lost upon PNIPAAm brush growth, it can be inferred that the thickness
of the dry brushes is at least the probe depth of XPS (~10 nm). According to ellipsometry
measurements of PET-PNIPAAm films on silicon wafer, the dry thickness of the PNIPAAm
brush after a polymerization time of 30 min is ≈40 nm, which, assuming an average grafting
density of 0.45 chains/nm 2 , corresponds to a molecular weight of ≈48 kDa. 33 Although the
microfibers possess a curved surface, we contend that, on the basis of the brush gyration
212

diameter (≈40 nm) relative to the microfiber diameter (600-1200 nm), the thickness of the
PNIPAAm brush does not differ substantially from that produced on a flat film.
The thermoresponsiveness of the PNIPAAm brushes grown on PET microfibers is first
evaluated with WCA experiments performed successively above and below the Tc of
PNIPAAm, as shown in Figure 4. The WCA of unmodified PET microfibers at 25˚C (Figure
4a) is �125˚, which is higher than that of a flat PET film (75˚) because of the "rough" nature
of the microfiber mat. Despite this increase in surface roughness, the size of the water droplet
on the surface of unmodified PET microfibers does not change during the course of the
measurement, and the measured WCA remains constant. This result also verifies that no
significant evaporation of water takes place during the course of the WCA measurement
because liquid evaporation during WCA measurement may sometimes reduce the apparent
contact angle values due to pinning of the contact line. In Figure 4b, the WCA of the
unmodified PET microfibers at 60˚C is �124˚ and likewise does not change, which suggests
that water evaporation is negligible. Cycling the specimen between these two temperatures in
Figures 4c and 4d yields comparable results, confirming that the PET surface stays
hydrophobic. Measured WCA values of PET-PNIPAAm microfibers, on the other hand,
display significantly different behavior. At 25˚C (Figure 4a), the WCA is also �125˚ when
the water droplet is initially placed on the microfiber surface, but quickly decreases to 0˚ in
just over 40 s as the water is wicked by the hydrophilic PNIPAAm brushes on the surface of
the microfibers. When the temperature is increased beyond Tc of PNIPAAm to 60˚C (Figure
4b), the water droplet is not strongly affected by the microfiber due to the increased
213

hydrophobicity of the PNIPAAm chains, and the WCA is �124˚. Repetition of these
measurements upon thermal cycling in Figures 4c and 4d confirm that the thermorespon-
siveness of PNIPAAm brushes on PET microfibers is reversible with no evidence of
hysteresis.
A second probe of the thermoresponsive nature of PNIPAAm brushes on PET
microfibers employs Au nanoparticles as tracers. Previous studies 29,34 have established that
Au nanoparticles attach to PNIPAAm chains via hydrogen bonding between the citrate
groups present on the nanoparticle surface and the amide groups on PNIPAAm. To discern
the extent to which the PNIPAAm brushes could bind Au nanoparticles, electrospun PET
microfibers have been submerged in a 0.05 mg/ml suspension of Au nanoparticles in
deionized water for 24 h at the same two temperatures examined in Figure 4, i.e., 25 and
60ºC. Images acquired by SEM after drying the fibers reveal that the nanoparticle loading on
the surface of PET-PNIPAAm microfibers is significantly higher at 25ºC (Figure 5a) than at
60ºC (Figure 5b). This difference is attributed to the thermoresponsiveness of the PNIPAAm
chains, which are hydrophilic and swell in water at temperatures below Tc, but become
hydrophobic and collapse in water at temperatures above Tc. As a result of such temperature-
driven swelling or contracting of the brush, the number of NIPAAm units available for
attachment of the Au particles increases or decrease, respectively, which, in turn, governs the
concentration of Au nanoparticles bound to PNIPAAm.
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Conclusions
In this study, we have demonstrated that the surface of electrospun PET microfibers can
be modified via amidation of the amine group on APTES with the ester group on PET to
permit further chemical modification ultimately resulting in the growth of polymer brushes
by ATRP. Step-by-step examination of the PET surface during the modification sequence,
along with quantitative analysis whenever possible, verifies expectations, and establishes the
sequence as a straightforward and viable route for PET microfiber functionalization. The
thermoresponsive behavior of the PNIPAAm brushes on PET microfibers has been
investigated using both contact angle measurements to determine the nature of the modified
PET surface and Au nanoparticle tracers to determine the extent of brush swelling at
temperatures below and above the lower critical solution temperature of PNIPAAm in water.
Surface functionalization of electrospun PET microfibers using this approach and PNIPAAm
in particular yields mechanically robust and highly porous mats that are temperature-
sensitive, which means that they are suitable candidates for diverse technologies as
responsive filters, scaffolds, delivery vehicles, and sensors.
Acknowledgments: This work was supported by the United Resource Recovery Corporation
and the National Science Foundation through a Graduate Fellowship (K. E. R.).
215

Figures
Figure A1.1. Sequence of surface modification steps employed in this study to functionalize
electrospun PET microfibers with thermoresponsive PNIPAAm brushes. The steps require
deposition and amidation of APTES (a), followed by hydrolysis of the ethoxysilane groups
on APTES to form silanol groups (b), which permit attachment of BMPUS (c) and
subsequent ATRP of NIPAAm to yield PNIPAAm brushes (d). The top and bottom SEM
images display PET and PET-PNIPAAm microfibers, respectively.
216

Figure A1.2. FTIR spectra of (a) as-spun PET, (b) PET-SiOH and (c) PET-PNIPAAm
microfibers. Spectra arranged in the same order in the ex 35 ded views reveal the appearance of
peaks associated with the formation of secondary amide moieties (dotted lines; see text for
assignments).
217

Figure A1.3. XPS spectra of (a) PET-SiOH microfibers and (b) PET-PNIPAAm microfibers.
The survey scans confirm the presence of N upon amidation of PET by APTES in (a) and
PNIPAAm brush formation in (b). The high-resolution insets show the C1s peak (� 285 eV)
before (a) and after (b) PNIPAAm brush growth.
218

Figure A1.4. Cyclic WCA measurements of as-spun PET ( ) and PET-PNIPAAm ( )
microfibers at temperatures (in ºC) below and above the Tc of PNIPAAm: (a) 25, (b) 60, (c)
25, and (d) 60. The error bars correspond to one standard deviation in the data.
219

Figure A1.5. SEM images acquired from PET-PNIPAAm microfibers exposed to aqueous
suspensions of Au nanoparticles at temperatures (labeled) below and above the Tc of
PNIPAAm. The illustrations in the insets (not drawn to scale) portray the conformation of
the PNIPAAm brush at each temperature.
220

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223

APPENDIX II
Generation of Functional PET Microfibers through Surface-Initiated
Abstract
Polymerization
Ali Evren Özçam, Kristen E. Roskov, Richard J. Spontak and Jan Genzer
In this study, we report on a facile and robust method by which poly(ethylene
terephthalate) (PET) surfaces can be chemically modified with functional polymer brushes
while avoiding chemical degradation. The surface of electrospun PET microfibers has been
functionalized by growing poly(dimethylaminoethyl methacrylate) (PDMAEMA) and
poly(2-hydroxyethyl methacrylate) (PHEMA) brushes through a multi-step chemical
sequence that ensures retention of mechanically robust microfibers. Polymer brushes are
grown via "grafting from" atom-transfer radical polymerization after activation of the PET
surface with 3-aminopropyltriethoxysilane. Spectroscopic analyses confirm the expected
reactions at each reaction step, as well as the ultimate growth of brushes on the PET
microfibers. Post-polymerization modification reactions have likewise been conducted to
further functionalize the brushes and impart surface properties of biomedical interest on the
PET microfibers. Antibacterial activity and protein resistance of PET microfibers
functionalized with PDMAEMA and PHEMA brushes, respectively, are demonstrated,
224

thereby making these surface-modified PET microfibers suitable for filtration media, tissue
scaffolds, delivery vehicles, and sensors requiring mechanically robust support media.
Introduction
Electrospinning produces solid polymer fibers with diameters ranging from several tens
of nanometers up to several microns. These nano/microfibers possess a high ratio of surface
area to volume and are typically organized in high-porosity mats that are suitable for use in a
broad range of applications involving (but not limited to) filters, 1 sensors, 2 nanocomposites, 3
tissue engineering scaffolds, 4 drug delivery vehicles, 5 and energy storage media. 6 During
electrospinning, a polymer solution or melt with acceptable viscosity and conductivity levels
is subjected to an electric field acting between a syringe needle and a collector plate. When
electrostatic forces overcome the surface tension of the liquid at the tip of the needle, a
charged polymer solution/melt jet is emitted from the resulting conical structure known as
the Taylor cone. 7 The jet undergoes a whipping process during which any solvent present
evaporates, and the polymer is commonly collected as a dry, randomly oriented fiber mat on
the grounded collector plate. 8 Electrospinning represents an appealing and facile means of
nano/microfiber production due to its relatively straightforward setup and the ability to tune
independent fiber properties on the basis of both solution/melt characteristics and process
parameters.
Although the structural features of electrospun nano/microfibers are advantageous, the
bulk properties of such fibers tend to lack the (multi)functionality that is needed for many
technologies. One way to overcome this problem is to create composite nano/microfibers by
225

incorporating chemically and/or physically different species (molecules or nanoparticles) into
the fibers to enhance, for example, mechanical, 9 electrical, 10 magnetic, 11 or optical 12
properties. Because functional species intended for use on polymer surfaces often exhibit
reduced surface activity when incorporated in a polymer matrix prior to electrospinning, they
may not always locate at the surface where their functionality is desired. 13 For instance, when
antibacterial biocides are added to a polymer prior to electrospinning, their efficacy is greatly
compromised, and they may become unable to attack airborne pathogens. 14 Sun et al. 15 have,
however, established that polarizable peptide-containing copolymers added to a polymer
prior to electrospinning can be brought to the surface of nano/microfibers by the applied
electric field, thereby resulting in fibers that are concurrently electrospun and
biofunctionalized. Alternatively, the surface of electrospun nano/microfibers can be modified
through the covalent bonding of poly(quarternary ammonium), which likewise creates a
permanent antibacterial surface. 16
Because polymer surfaces typically possess low surface energy, they must be pretreated
chemically or physically to obtain an active surface suitable for subsequent
functionalization. 17 Physical methods by which to activate a polymer surface include plasma
treatment, 18 'layer' formation, 19 UV treatment, 14 mineralization, 18 etching, 20 or inclusion of a
reactive composite material. 21 Once chemically-active groups reside on the surface, covalent
bonding, 22 immobilization, 23 and electrostatic interactions 24 can be used to attach reactive
groups to the fiber surface. Modification of only the fiber surface can make commodity and
engineering plastics in particular suitable for applications wherein the fibers interact with
226

their environment, such as molecular filtration, 25 protective textiles, 26 tissue scaffolds, 27 and
drug delivery. 5 A synthetic polymer that shows particular promise in this regard is
poly(ethylene terephthalate) (PET), and electrospun PET microfibers have already been
considered in applications that benefit from the mechanical strength, transparency, and
solvent resistance of PET. 28 As with most organic polymers, however, PET does not possess
good adhesion and wetting properties because of its inherently low surface energy (42
mJ/m 2 ). Application of electrospun PET microfibers as functional materials therefore
necessitates alteration of their surface properties without compromising their bulk
characteristics.
Modification of PET surfaces has been conducted by a variety of chemical methods,
including chemical treatment (e.g., hydrolysis, 29,30,31 reduction, 31,32,33 aminolysis, 30,32,33,34
glycolysis, 31 polyelectrolyte deposition, 35 surface graft polymerization after surface
activation 34,36 ) and physical modification (e.g., plasma, 37,38 ultraviolet/ozone, 38,39 flame, 38
corona treatment, 38,40 electrical discharge, 41 ion beam bombardment, 42 and laser treatment 43 ).
Since most of these surface modification techniques involve, purposefully or inadvertently,
polymer degradation, careful selection of experimental conditions is imperative for the
successful surface modification of PET microfibers without degrading the bulk polymer and
its desirable mechanical properties. Grafting polymer brushes on surfaces represents an
attractive approach by which to modify and control the surface properties of materials.
Surface-initiated graft polymerization has been performed successfully on flat surfaces with a
variety of monomers and polymerization methods. 44 Specifically, atom transfer radical
227

polymerization (ATRP) has been employed 45,46 extensively because (i) its controlled
implementation does not require ultrapurification of the chemicals used, and (ii) it can be
used to polymerize numerous functional monomers, such as N-isopropylacrylamide
(NIPAAm), 47 2-(dimethylamino)ethyl methacrylate (DMAEMA) 48 and 2-hydroxyethyl
methacrylate (HEMA), 49 as well as others.
Graft polymerization on PET surfaces has been reported by Roux and Demoustier-
Champagne 36 and Bech et al. 50 (for styrene), as well as Farhan and Huck 51 (for NIPAAm).
These efforts employ various means of attaching surface initiators for "grafting from"
polymerization. For instance, Roux and Demoustier-Champagne 36 have attached free-radical
polymerization initiators to the surface of PET via electrostatic interactions and covalent
bonding after surface activation by saponification and oxidation. Farhan and Huck 51 and
Bech et al. 50 have alternatively attached ATRP initiators after activating the PET surface by
plasma treatment and aminolysis, respectively. The major drawback of these approaches —
viz., saponification, aminolysis, and plasma treatment — is that they often induce severe
degradation of PET and promote a roughened surface topography. Because of the
nano/micrometer dimensions of electrospun PET fibers, it is paramount that material
degradation and surface roughening must be minimized. In this work, we have elected to
functionalize PET microfiber surfaces by means of 3-aminopropyltriethoxysilane (APTES).
Bùi et al., 32 Fadeev and McCarthy, 52 and Xiang et al. 53 have demonstrated that the primary
amine group in APTES inserts into the PET chain via an amidation reaction with negligible
degradation to bulk PET. In this reaction mechanism the triethoxysilane groups of APTES
228

are exposed at the air interface, and subsequent hydrolysis of the ethoxysilane units yields
silanol groups on the PET surface. These groups are suitable as attachment points for an
ATRP initiator, such as [11-(2-bromo-2-methyl)propionyloxy] undecyltrichlorosilane
(BMPUS). We use this approach to grow PDMAEMA and PHEMA brushes on electrospun
PET microfibers and a battery of analytical techniques to characterize the properties of
electrospun PET microfibers before and after surface polymerization of DMAEMA and
HEMA via ATRP. We also investigate the post-polymerization modification of the
corresponding PDMAEMA and PHEMA brushes via quarternization and fluorination,
respectively, and demonstrate the antibacterial and protein-resistance properties of these
functionalized PET microfiber mats.
Experimental
Materials
Food-grade recycled PET flakes were kindly supplied by United Resource Recovery
Corporation (Spartanburg, SC). Anhydrous toluene, 2-chlorophenol, methanol, iodomethane,
iodopropane, iodobutane, bromoethane, bromopropane, bromobutane, trifluoroacetic
anhydride (TFAA), fibrinogen from human plasma, 1X-PBS buffer (0.137 M NaCl, 0.0027
M KCl, and 0.0119 M phosphates), APTES, DMAEMA, HEMA, Cu(I) bromide (CuBr),
Cu(II) bromide (CuBr2), Cu(I) chloride (CuCl), Cu(II) chloride (CuCl2), bipyridine, and
N,N,N',N',N"-pentamethyldiethylenetriamine (PMDETA) were all purchased from Sigma-
229

Aldrich (St. Louis, MO) and used as-received. Hexafluoroisopropanol (HFIP) was obtained
from Oakwood Products Inc. (Estill, SC).
Brush Growth/Modification
The PET flakes were dissolved in HFIP at different polymer concentrations and
electrospun at 10 kV to generate microfibers possessing different diameters. Thin PET films
measuring 12 and 180 nm thick were likewise spin-coated on silicon wafers from 0.5 and 3
wt% solutions, respectively, in 2-chlorophenol. The latter specimens allowed us to follow
each modification step by measuring film thickness increments associated with the various
chemical modification steps and protein adsorption. Fiber mats and thin films were kept
under vacuum for at least 48 h prior to use to remove entrapped solvent. The initiator for
ATRP was BMPUS, synthesized as described earlier 46 and deposited on the surface of PET
microfibers and films after activation of the surface with APTES. 54 Polymer brushes
composed of either PDMAEMA or PHEMA were subsequently grown from the BMPUS-
decorated PET surfaces by ATRP according to established protocols. 48 For instance, 10.09 g
HEMA was mixed with 6.81 g of methanol, 1.88 g of water and 0.63 g of bipyridine in an
Ar-purged Schlenk flask, and oxygen was removed via three freeze-thaw cycles. After
removal of oxygen, CuCl (0.18 g) and CuCl2 (0.01 g) were added to the solution and 1 more
freeze-thaw cycle was performed. This ATRP solution was transferred to a tightly sealed
Schlenk flask, which was stored in an Ar-purged glove box. Fiber mats and films decorated
with BMPUS were submerged in the ATRP solutions for 6 and 8 h to produce PDMAEMA
and PHEMA brushes, respectively. After removal from the ATRP solution, the samples were
230

insed promptly with methanol and deionized water, and then sonicated in deionized water
for 10 min.
The PDMAEMA brushes grown on PET microfibers and silicon wafers were
quarternized with iodomethane, iodopropane, iodobutane, bromoethane, bromopropane, and
bromobutane in acetonitrile at 60˚C for �20 h. An excess amount of quarternization agents
was added to the glass vial containing PDMAEMA-modified PET microfiber mats and
acetonitrile to yield fully quarternized (q) PDMAEMA brushes. In similar fashion, the
PHEMA brushes were fluorinated with TFAA to discern the effect of fluorinated PHEMA
(fPHEMA) on protein adsorption. In this case, TFAA was used to bind fluorinated moieties
to the hydroxyl terminus of the HEMA pendant group. All reactions were conducted at
ambient temperature in the gas phase, and the samples were washed with copious amounts of
ethanol and water and dried under reduced pressure before protein adsorption experiments. In
both of the post-polymerization modification reactions listed above, bare (i.e., brush-free)
PET microfibers were immersed in the post-polymerization modification reaction mixtures as
controls.
Material Characterization
The thicknesses of the PET films deposited on silicon wafer were measured with
variable-angle spectroscopic ellipsometry (VASE, J.A. Woollam) at an incidence angle of
70º (between the beam and the surface normal) before and after each modification step to
measure the approximate PDMAEMA and PHEMA brush thicknesses on the PET
microfibers. In addition, the thickness of the polymer brushes after quarternization,
231

fluorination, and protein adsorption was also measured with VASE to determine the extent of
these post-polymerization modification steps. The surface chemical composition of modified
microfibers was measured after each modification step by X-ray photoelectron spectroscopy
(XPS) performed on a Kratos Axis Ultra DLD spectrometer at a take-off angle of 90º (under
these conditions the probing depth of XPS is estimated 55 to be ≈9-10 nm). Fourier transform
infrared (FTIR) spectroscopy was utilized to monitor chemical changes that occurred on the
surface of the PET microfibers after modification. Spectra were recorded on a Nicolet 6700
spectrometer after embedding microfiber mats in KBr pellets for analysis in transmission
mode, and resulting data were analyzed by the Omnic Specta software. For each sample,
1024 scans were collected at a resolution of 4 cm -1 . As-spun and surface-modified PET
microfibers were coated with ≈8 nm of Au, and their diameter and surface morphology were
examined by field-emission scanning electron microscopy (SEM) performed on a JEOL
6400F electron microscope operated at 5 kV.
Antibacterial Activity
The PET microfibers decorated with a qPDMAEMA brush were subjected to
antibacterial testing using a modified ASTM standard (E2149-01 Standard Test Method for
Determining the Antimicrobial Activity of Immobilized Antimicrobial Agents under Dynamic
Contact Conditions). Here, E. coli, a model gram-negative bacteria, was grown in a Lauria-
Bertani (LB) medium overnight to yield a bacteria population of 5 x 10 8 according to UV-Vis
spectrophotometry. After serial dilutions, the modified PET microfiber mats (measuring �1
cm 2 in area) were incubated in a suspension containing 3 x 10 5 bacteria in sterile conical
232

tubes at 37�C while being shaken at 300 rpm for 1 h. The resultant suspension was then
diluted with LB medium to a desired concentration and spread on L-agar plates. The L-agar
plates were incubated at 37°C for 18 h. Each surviving cell developed into a distinct bacterial
colony and provided information regarding bacterial activity. The number of viable cells was
measured in terms of colony forming units (CFUs) on each plate.
Protein Resistance
A 0.1 mg/ml solution was prepared at the isoelectric point of fibrinogen (FIB, at pH =
5.5) by dissolving FIB in 1X-PBS buffer solution (0.2% NaN3 was added to the buffer to
prevent bacterial growth). The solution was passed through a 0.2 �m filter, and adsorption
studies of FIB were conducted by incubating substrates in protein solution for 16 h at
ambient temperature. Both fluorinated and unmodified PHEMA brushes on PET microfibers
and films were tested alongside bare PET microfibers and films. After incubation, samples
were washed thoroughly with deionized water, dried under reduced pressure and stored in
glass vials for further characterization. The thickness of the adsorbed FIB layer was
measured by VASE on flat samples (PET films on silicon wafers), and the corresponding
nitrogen surface concentration was measured by XPS to ascertain the amount of adsorbed
FIB.
Results and Discussion
The diameters of electrospun PET microfibers, prepared according to the protocol
provided in the Experimental section and measured by SEM, are 450±100, 800±200 and
1200±300 nm for 6, 8 and 10 wt% solutions, respectively, of PET in HFIP. The surfaces of
233

unmodified PET microfibers appear consistently smooth, as is evident in Figure 1. Functional
PDMAEMA and PHEMA brushes have been grown on PET microfibers in the sequence of
four steps reported earlier 54 for PNIPAAm brushes. Briefly, APTES molecules are attached
to the PET surface via aminolysis between PET and the primary amine of APTES. Next, the
ethoxysilane groups on APTES are hydrolyzed to generate silanol groups for BMPUS
attachment. Finally, PDMAEMA and PHEMA brushes are grown directly from the PET
microfiber surface via ATRP.
As reported previously by Bùi et al. 32 and Fadeev and McCarthy, 52 the primary amine
group in APTES reacts with the ester functionality in PET by forming an amide bond via
aminolysis (cf. Figure 1). Both studies claim that aminolysis of PET with APTES does not
degrade bulk PET as opposed to aminolysis of PET with short alkyl amines, since the latter
can diffuse into a PET fiber and react all the way through, and thus weaken, the fiber. 30,33
The presence of bulky triethoxysilane group on APTES molecules hinders the diffusion of
APTES into PET by increasing the size of the molecule, changing the solubility of the alkyl
amine to which it is attached and creating a protective surface layer (which serves to impede
the diffusion of other APTES molecules). While neither Fadeev and McCarthy 52 nor
Howarter and Youngblood 56 could detect the formation of amide groups on PET due to the
small population of amide groups available on their flat samples, the presence of amide
groups on PET-SiOH microfiber surfaces has been directly confirmed via FTIR analysis by
Ozcam et al. 54 as a result of the enlarged surface area afforded by electrospun PET
microfibers. The appearance of new peaks at 1650 (amide I band), 1550 (amide II band),
234

3300, 1470, and 3300 cm -1 are attributed to the formation of secondary amide groups on the
surface of PET microfibers. Attachment of APTES, followed by hydrolysis of the exposed
triethoxysilane groups in acidic water (pH ≈ 4.5-5.0), yields further reactive silanol groups,
as indicated by both a decrease in water contact angle (WCA) and associated thickness
measurements performed on flat PET films on silicon wafers. For instance, the WCAs of
APTES-modified PET and virgin PET films are 50˚�0.8˚ and 71˚�0.8˚, respectively, after
exposure to acidic water, in contrast to the WCA of native PET (75˚�0.2˚). In addition, XPS
spectra collected from PET-SiOH microfibers confirm the existence of a small nitrogen N1s,
Si2s and Si2p peaks at 400, 150 and 100 eV, respectively, which correspond to 0.6 atom%
nitrogen and 1.1 atom% silicon present on the PET-SiOH surface.
Attachment of BMPUS molecules to the PET-SiOH surface as initiator centers for
"grafting from" polymerization of DMAEMA and HEMA, followed by the ATRP conditions
described in the Experimental section, results in the formation of dry brushes measuring ≈50
and ≈45 nm thick, respectively, as determined by VASE analysis of thin spin-coated PET
films on silicon wafers. Here, we assume that the thicknesses of brushes grown on silicon
wafers is comparable to that produced on the microfibers, since the microfibers are relatively
large and possess negligible curvature on the size scale of the brushes. Examples of PET
microfibers after brush growth are presented in Figure 2 and verify that the chemical
reactions undertaken have no discernible effect on microfiber morphology. Corresponding
FTIR spectra of the polymer brushes, along with spectra acquired from electrospun PET and
PDMAEMA and PHEMA brushes grown directly on silicon wafers, are provided in Figure 3.
235

The spectra of PDMAEMA and PHEMA brushes on silicon wafer are included to point out
the chemical changes that occur on the PET microfibers. Careful comparison of these spectra
confirms that PDMAEMA and PHEMA brushes grew from the surface of PET microfibers.
The appearance of new stretching vibrations located at 2770 and 2820 cm -1 for the
PDMAEMA brush (blue line in Figure 1a), for instance, reflects the C-H bond of the –
N(CH3)2 group of PDMAEMA. Likewise, the increase in peak intensity at �3400 cm -1 for the
PHEMA brush (blue line in Figure 1b) is a consequence of the broad –OH peak originating
from PHEMA.
The chemical compositions of PDMAEMA and PHEMA brushes grown from the surface
of PET microfibers have been assessed by XPS, and resulting values are listed in Table 1.
The theoretical values of these compositions are calculated on the basis of the number of
atoms present on each repeat unit of both polymers and the assumption that the brush
thickness is larger than the probing depth of XPS (�10 nm). Representative XPS spectra of
PET microfibers with grafted PDMAEMA and PHEMA brushes are plotted for comparison
in Figure 4. Bare PET exhibits 2 ionization peaks, one for carbon (at 285 eV) and the other
for oxygen (at 536 eV). Corresponding surface concentrations, computed from the areas
under these curves, are 73.2�0.4 (71.4) and 26.8�0.4 (28.6) atom% for carbon and oxygen,
respectively, which agree favorably with the theoretical values provided in parentheses. On
one hand, growth of a PDMAEMA brush on the surface of PET microfibers is responsible
for the appearance of the nitrogen peak at 400 eV (cf. Figure 4b). Quantitation of such XPS
spectra results in surface concentrations of 73.6�0.5, 7.6�0.1 and 18.8�0.5 atom% for
236

carbon, nitrogen and oxygen, respectively. Introduction of a PHEMA brush, on the other
hand, does not change the number or the position of the peaks recorded in XPS spectra.
Instead, the relative peak areas are affected so that the surface concentrations become
70.3�0.5 and 29.7�0.5 atom% for carbon and oxygen, respectively. These concentration
values measured experimentally for PDMAEMA and PHEMA brushes grown on PET
microfibers are in good quantitative agreement with the theoretical values determined on the
basis of the chemical structures of the individual species (cf. Table 1).
Ellipsometry measurements of flat PET films on silicon wafer, in conjunction with
independent XPS measurements conducted on PET microfiber surfaces, suggest that the
polymer brushes completely cover the PET surfaces, since the dry thicknesses of the brushes
exceed the probing depth of XPS. The characteristic XPS "fingerprint" of PET disappears
from the high-resolution spectra (displayed in the insets of Figure 4) after growing the
PDMAEMA and PHEMA brushes. Introduction of the peak corresponding to the C-N bond
at 286.1 eV serves to broaden the peaks at 285.0 (the C-C bond) and 286.6 eV (the C-O
bond) for the PDMAEMA brush grown on PET microfibers. In the case of the PHEMA
brush, the intensity of the peak located at 286.6 eV is larger than that at 285.0 eV relative to
the XPS spectrum of bare PET. The peak at 290.0 eV, which corresponds to the O-C=O
groups of acrylates, is present for both PDMAEMA and PHEMA brushes grown on the PET
microfibers. 57
Post-polymerization modification reactions have been performed on the PDMAEMA and
PHEMA brushes grafted to the surface of PET microfibers to introduce antibacterial and
237

protein resistance properties, respectively. One of the potential applications of PDMAEMA
after quarternization of the dimethylamino groups on the DMAEMA repeat unit is as an
antibacterial material. Polymer chains quarternized with alkyl halides possess positive
charges and hydrophobic alkyl chains, which induce cation exchange and penetration through
the bacterial cell membrane, respectively. These result in disruption of membrane integrity
and death of bacterial cells. 58 Antibacterial properties of quarternary ammonium compounds
(QACs) have been reported earlier in solution 59 and on solid surfaces. 60,61,62 The latter has an
important advantage over free QACs because they are covalently attached to substrates,
which, in turn, permit repeated use with limited biocidal release to the environment. 63
Quarternization of the PDMAEMA brush grown on the surface of PET microfibers has been
achieved with alkyl bromides differing in length to yield a polycationic brush. Conversely,
the PHEMA brush can be modified to resist protein adsorption, 64 which remains a significant
challenge in biomedical applications involving artificial implants. Adsorption of biomass on
the surface of functional materials degrades the functionality over time. Biomass
accumulation begins with protein adsorption and denaturation on any surface with which
proteins come in contact. Protein adsorption on various surfaces has been studied extensively
over the past several decades, and the incorporation of ethylene glycol and fluorinated units
into polymeric coatings have been found to be among the most effective at reducing the
propensity for protein adsorption.
In the present study, several quarternization agents differing in alkyl length —
iodomethane, iodopropane, iodobutane, bromoethane, bromopropane, and bromobutane —
238

have been used to introduce positive charges into the PDMAEMA brush grown on PET
microfibers and generate a polycationic qPDMAEMA brush with antibacterial properties.
Similarly, TFAA has been used to fluorinate the –OH groups of the PHEMA brush so that
the effect of fPHEMA on the protein resistance of functionalized PET microfiber mats can be
probed. The morphologies of PET microfibers after these post-polymerization modification
reactions are visible in Figure 5 and verify that a microfibrous network of the electrospun
mat is retained. Chemical modification of PDMAEMA and PHEMA brushes grown on
silicon wafers (not on PET film) with quarternization and fluorination agents, respectively,
results in multiple changes in the FTIR spectra. For example, the stretching vibrations
located at 2770 and 2820 cm -1 for the PDMAEMA brush in Figure 6a are attributed to the C-
H bond of the –N(CH3)2 group. These peaks disappear completely after quarternization.
Water absorbed by the more hydrophilic qPDMAEMA brush is responsible for the
appearance of the peak located at �3400 cm -1 . New peaks reflecting the formation of C-CO-
CF bonds (at 1789 cm -1 ) and C-F bonds (at 1224 and 1157 cm -1 ) likewise appear after
conversion of PHEMA to fPHEMA in Figure 6b. Moreover, the –OH groups of PHEMA are
consumed during the fluorination reaction with TFAA, and the peak located at �3400 cm -1
disappears. It is noteworthy that the FTIR spectra collected from PDMAEMA and PHEMA
brushes grown on PET microfibers before and after quarternization (PDMAEMA) and
fluorination (PHEMA) possess the same characteristic peaks in Figure 6, thereby providing
evidence that the brushes grown on the PET microfibers are functionalized.
239

The chemical compositions of brushes grown on PET microfibers and silicon wafers after
quarternization and fluorination reactions have been measured by XPS. Examination of the
corresponding XPS spectra (data not shown) reveals the appearance of new peaks for iodine
or bromine after quarternization of PDMAEMA and fluorine after fluorination of PHEMA.
Quantitation of these spectra yields the surface concentrations listed in Table 2 for
qPDMAEMA (with alkyl bromides) and Table 3 for fPHEMA. Interestingly, the
concentration of bromine in the qPDMAEMA brush grown on PET microfibers is �50%
greater than that in the qPDMAEMA brush on silicon wafer. While this difference was
initially attributed to the adsorption or absorption of alkyl bromides on/in PET, XPS spectra
obtained from bare PET microfibers exposed to the quarternization medium for the same
reaction time reveal no existence of bromine. Therefore, we propose that this difference is a
result of the curved nature of the PET microfibers, which apparently possess a lower steric
hindrance (due to the higher surface area) for the quarternization reaction as compared to a
flat surface. The same trend is also observed for the quarternization reaction of PDMAEMA
brushes with alkyl iodides. In contrast, the extent of fluorination of the PHEMA brush grown
on PET microfibers and silicon wafers is similar, and these values are in agreement with
those reported earlier for TFAA-modified PHEMA brushes. 64 This observation, which differs
from the results obtained for qPDMAEMA brushes, may be due to the smaller size of TFAA
relative to the alkyl halides and the gas-phase reaction of TFAA (wherein TFAA molecules
can diffuse through the brush and completely react with all available –OH groups without
restriction). Arifuzzaman et al. 64 have demonstrated that the surface concentration of TFAA-
240

modified PHEMA brushes on silicon wafer does not change as a function of XPS take-off
angle, which evinces that (i) PHEMA brushes react homogeneously throughout the XPS
probing depth and (ii) the gas-phase reaction of TFAA is quantitative with PHEMA brushes
irrespective of the substrate on which they are grown. We hasten to add that exposure of bare
PET microfibers to TFAA did not alter the surface composition of the bare PET, according to
XPS analysis.
The presence of QACs endows the surface of the PET microfibers with antibacterial
properties due to the presence of cationic groups that disrupt cell membranes and induce
bacterial lysis. 65 In the case of gram negative bacteria such as E. coli, the phosphate groups
of lipopolysaccharide molecules located in the outer bacterial membrane are stabilized by
divalent cations, which would otherwise strongly repel each other, via bridging and
neutralizing. Bacteria lose their natural counterions and their outer membrane is destabilized
upon interacting with QACs due to the electrostatic compensation of these charges with the
cationic charges of the QACs. Thus, the release of counterions from the outer cell wall
initiates the death of the bacteria. 63,65,66 Quarternization of the PDMAEMA brush on PET
microfibers with alkyl bromides differing in methylene length produces string (quenched)
polycationic brushes on the microfiber surface. As reported earlier, 60,61,67 the antibacterial
efficacy of a QAC depends on the extent of quarternization, as well as the length of the alkyl
chain in the quarternization agent. The extent of quarternization dictates the number of
positive charges available to interact with the bacterial membrane, whereas the length of the
alkyl chain affects the antibacterial efficacy by governing the depth of penetration through
241

the cell wall. In general, antibacterial efficiency increases as the length of the alkyl spacer is
increased, but deteriorates after 6 methylene units. 61 As seen in Figure 7a, the presence of a
polycationic qPDMAEMA brush on PET microfibers provides antibacterial properties
against E. coli as the number of CFUs on the agar plates with qPDMAEMA-modified
microfibers is lower than on those with PDMAEMA-modified microfibers. In addition, the
antibacterial efficiency of the qPDMAEMA brush increases substantially with increasing
alkyl length of the quarternization agent from bromoethane to bromobutane, as indicated by
the results provided in Figure 7b.
The resistance of PHEMA and fPHEMA brushes to protein adsorption on flat surfaces
has been recently investigated, and the presence of PHEMA 49 and fPHEMA 64 brushes has
been found to reduce protein adsorption, depending on the graft density and molecular
weight of the brush. In this work, we only examine the protein resistance of a PHEMA brush
grown on PET microfibers before and after fluorination with TFAA. Several different
samples including PHEMA brushes on PET microfibers, fPHEMA brushes on PET
microfibers, PHEMA brushes on silicon wafer, fPHEMA brushes on silicon wafer, bare PET
microfibers and bare PET microfibers exposed to TFAA have all been incubated in FIB
solution for 16 h. Adsorption of FIB on flat substrates has been monitored by measuring the
brush thickness with VASE (on silicon wafers) and the surface nitrogen concentration (due to
FIB) with XPS (on fibers and silicon wafers). Comparing the FIB layer thickness on flat
surfaces reveals that the presence of a PHEMA brush dramatically reduces protein
adsorption. According to the results presented in Figure 8, the thickness of FIB plummets
242

from �4 nm to almost 0 nm after in the presence of a PHEMA brush, and this reduction is
corroborated by XPS data that confirm a corresponding decrease in nitrogen concentration
from 15.1 to 0.6 atom%. The amount of FIB adsorbed on spin-coated PET film is comparable
to that adsorbed on bare silica wafer, but the concentration of adsorbed FIB on bare PET
microfiber (with and without TFAA treatment) is noticeably lower than that on spin-coated
PET film. Introduction of a PHEMA brush on PET microfibers effectively prevents FIB
absorption, but quaternization of the brush does not appear to afford further improvement. In
contrast, fluorination of the PHEMA brush grown on silicon wafer slightly improves protein
resistance, as discerned from both thickness and XPS results.
Conclusions
In this work, we have demonstrated that the surface of electrospun PET microfibers can
be controllably modified via the amidation reaction of the amine group on APTES with the
ester group of PET and subsequent growth of functional polymer brushes (PDMAEMA and
PHEMA) by ATRP. Post-polymerization modification of these brushes has been conducted
by quarternization (PDMAEMA) and fluorination (PHEMA) reactions. None of these brush-
growing or post-functionalization reactions have any discernible deleterious effect on the
morphology of the PET microfibers and, by inference, their robust mechanical properties.
The improved antibacterial efficacy of quarternized PDMAEMA brushes and protein
resistance of (fluorinated) PHEMA brushes grown on PET microfibers are established. These
functional microfiber mats are suitable for use as affinity filters, antibacterial clothing and
responsive sensors. Specifically, we envisage that these modified PET microfiber mats can
243

e employed as multi-use filters for water purification applications, in which case the
stability of such brushes must be ascertained as a function of pH and temperature. In this and
related technologies, the ability of surface-modified microfibers to withstand environmental
stresses is of paramount importance, which is why we have elected to use electrospun PET
microfibers and why we have
chosen a chemical reaction route that does not compromise the mechanical robustness of
PET.
Acknowledgments
This work was supported by the United Resource Recovery Corporation and the National
Science Foundation through a Graduate Fellowship (K. E. R.).
244

Tables
Table A2.1. Compositions of PET microfiber surfaces with grafted PDMAEMA and
PHEMA brushes from XPS analysis..
PET
microfiber
PDMAEMA
brush
PHEMA
brush
Species analyzed
245
Concentration (atom%)
Carbon Nitrogen Oxygen
Theoretical 71.4 � 28.6
Experimental 73.2�0.4 � 26.8�0.4
Theoretical 72.7 9.1 18.2
Experimental 73.6�0.5 7.6�0.1 18.8�0.5
Theoretical 66.7 � 33.3
Experimental 70.3�0.5 � 29.7�0.5

Table A2.2. Compositions of PET microfiber surfaces and silicon wafer modified with
grafted PDMAEMA brushes after quarternization.
Species analyzed
qPDMAEMA on PET microfiber
with bromoethane
qPDMAEMA on silicon wafer
with bromoethane
Bromoethane on PET microfiber
(control)
qPDMAEMA on PET microfiber
with bromopropane
qPDMAEMA on silicon wafer
with bromopropane
Bromopropane on PET microfiber
(control)
qPDMAEMA on PET microfiber
with bromobutane
qPDMAEMA on silicon wafer
with bromobutane
Bromobutane on PET microfiber
(control)
246
Concentration (atom%)
Carbon Nitrogen Oxygen Bromine
73.9 5.3 17.3 3.5
71.2 7.2 19.2 2.4
71.6 � 28.4 �
74.6 5.7 16.5 3.2
72.4 6.7 18.2 2.7
73.7 � 26.3 �
74.0 5.2 17.3 3.5
73.5 6.4 17.8 2.3
72.6 � 27.4 �

Table A2.3. Compositions of PET microfiber surfaces and silicon wafer modified with
grafted PHEMA brushes after fluorination.
Species analyzed
247
Concentration (atom%)
Carbon Oxygen Fluorine
fPHEMA on PET microfiber 54.4 25.0 20.6
fPHEMA on silicon wafer 53.2 24.6 22.2
TFAA-treated PET microfiber
(control)
74.2 25.3 0.5

Figures
Figure A2.1. Synthetic strategy for growing functional polymers on electrospun PET
microfibers. The deposition and subsequent hydrolysis of APTES is followed by attachment
of BMPUS, which provides access to a variety of "grafting from" reactions. The fiber color
and corresponding reactive species are color-matched. The SEM image shows bare
(unmodified) PET microfibers electrospun from HFIP.
248

Figure A2.2. SEM images acquired from electrospun PET microfibers with (a) PDMAEMA
and (b) PHEMA brushes. The solution concentrations used to generate the PET microfibers
was 8 wt%.
249

Figure A2.3. FTIR spectra collected from systems with (a) PDMAEMA and (b) PHEMA
brushes. Spectra correspond to bare electrospun PET microfibers (dotted black lines),
brushes grown on silicon wafers (solid black lines) and brushes grown on the PET
microfibers (solid blue lines).
250

Figure A2.4. XPS spectra collected from (a) electrospun PET microfibers, as well as PET
microfibers functionalized with (b) PDMAEMA and (c) PHEMA brushes. A high-resolution
carbon-edge spectrum is included in the inset of each panel.
251

FigureA2.5. SEM images acquired from electrospun PET microfibers with postfunctionalized
(a) qPDMAEMA (using bromobutane) and (b) fPHEMA brushes. The
solution concentrations used to generate the PET microfibers were 8 wt%.
252

Figure A2.6. FTIR spectra collected from systems with (a) PDMAEMA and (b) PHEMA
brushes. Spectra correspond to original and post-functionalized brushes grown on silicon
wafers (black and red, respectively), as well as original and post-functionalized brushes
grown on PET microfibers (blue and green, respectively).
253

Figure A2.7. In (a), photographs of E. coli colonies on L-agar plates containing bare PET
microfibers, as well as PET microfibers modified with a PDMAEMA brush and postquaternized
with bromoethane, bromopropane and bromobutane (labeled) after an incubation
period of 18 h at 37°C. The dependence of the number of colony forming units (CFUs) on
the length of the alkyl bromide used is evident in (b). For reference, the CFU corresponding
to the bare PET microfibers is indicated by the red line.
254

Figure A2.8. Surface nitrogen concentration (left ordinate, red bars) and fibrinogen thickness
(right ordinate, blue squares) for various systems containing silicon wafer, electrospun PET
microfibers, PHEMA brushes and post-functionalized fPHEMA brushes.
255

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260

APPENDIX III
Modification of Melt-spun Isotactic Polypropylene and Poly(lactic acid)
Abstract
Bicomponent Filaments with a Premade Block Copolymer*
Sara A. Arvidson, Kristen E. Roskov, Jaimin J. Patel, Richard J. Spontak,
Russell E. Gorga and Saad A. Khan
While numerous studies have investigated the effect of adding a block copolymer as a
macromolecular surfactant to immiscible polymer blends, no such efforts have sought to alter
the properties of melt-spun bicomponent core-sheath filaments with a nonreactive
compatibilizing agent. In this study, we examine the effect of adding poly[styrene-b-
(ethylene-co-butylene)-b-styrene] (SEBS) triblock copolymer to core-sheath filaments
consisting of isotactic polypropylene (iPP) and poly(lactic acid) (PLA). Incorporation of the
copolymer into blends of iPP/PLA is observed to reduce the size scale of phase separation.
Interfacial slip between molten iPP and PLA layers is evaluated by rheology under steady-
shear conditions. Addition of SEBS to the PLA sheath during filament formation reduces the
tendency of PLA sheaths to crack prior to iPP core failure during tensile testing. In reversed
filament configurations, the copolymer does not hinder the development of molecular
orientation, related to fiber strength, during fiber spinning. Electron microscopy reveals that
the copolymer molecules form unique, highly nonequilibrium morphologies under the
spinning conditions employed here.
261

Introduction
Bicomponent filament spinning involves the co-extrusion of two polymers (or polymer
blends) from the same spinneret into a single filament that consists of both starting materials.
This type of spinning can be configured in a variety of cross-sectional geometries such as
core-sheath, side-by-side, segmented pie, islands in the sea, or trilobal. 1 In this study, we only
consider further the core-sheath arrangement. The benefits of co-spinning two polymers
include a reduction in the cost associated with a single-step process, a net increase in the
performance of fibers and webs as derived from the desirable characteristics of the
constituent polymers, and suppression of an unfavorable rheological behavior (e.g., spinning
a polymer behaving as a Newtonian liquid with a small fraction of a desirable polymer
exhibiting Maxwell properties). 2 Moreover, since the viscoelasticity of the sheath polymer
dictates the mechanics of melt flow during coaxial spinning even if its flow rate and viscosity
are lower than those of the core polymer, core-sheath bicomponent fibers may be
significantly thinner than would otherwise be achieved by spinning the polymers
individually. 2 Lastly, while the bicomponent spinning of two polymers permits retention of
component properties, comparable spinning of pre-mixed polymer blends generally results in
properties that lie intermediate between those of the individual components, thus
compromising the net properties of the fibers. 3
Depending on the thermodynamic compatibility of the polymers involved, bicomponent
spinning can result in filaments or fibers (these words are used interchangeably throughout
text) that fail by splitting when subjected to an external mechanical force. 4 While splittable
262

fibers may be desired in the manufacture of synthetic suede and leather, technical wipes, and
some filtration applications, 5 good adhesion between the individual species comprising
bicomponent fibers is required for maintaining mechanical integrity, developing sutures, or
improving chemical or flammability resistance. 6,7 At the interface separating immiscible
polymers, a relatively low population of entanglements can lead to fiber delamination at
temperatures below the melting temperatures (Tms) of both components. Above the Tm of
each component, subjecting immiscible polymers to external forces may result in “slip” of
one molten polymer across the other. Evaluating the presence of slip and, by inference, the
adequacy of interfacial chain entanglements in the melt, remains a nontrivial task. Zhao and
Macosko 8 have deduced that slip of multilayered samples occurs when a drop in viscosity
accompanies an increase in the number of layers and, hence, interfacial contact area. Jiang 9
has similarly interpreted that a viscosity below the theoretical viscosity discerned from the
reciprocal rule of mixtures (R-ROM) for layered high-density polyethylene/polystyrene
(HDPE/PS) is indicative of slip. Slip has also been proposed as the cause for the viscosity
discontinuity encountered while shearing layered HDPE/PS filled with tracer particles and
observing the layers in situ with confocal microscopy. 10 Park et al. 11 have likewise
investigated multilayer slip by using a sliding plate rheometer equipped with a camera.
Block copolymers, composed of two or more long, covalently-linked sequences of
chemically-dissimilar repeat units, may be used to lower interfacial tension along polymer-
polymer interfaces and thus improve adhesion by promoting chain entanglements. The
apparent results of such compatibilization are a net reduction in slip during extrusion and
263

delamination in formed fibers. 12 Generally speaking, compatibilization is considered to be
effective if an added block copolymer reduces the size scale of phase domains and/or
enhances the mechanical properties of a blend. 3,13 The stabilizing efficacy and spatial
segregation of premade block copolymer molecules along polymer-polymer interfaces has
been previously addressed by Wei et al. 22 and more recently by Gozen et al. 23 Alternatively,
reactive compatibilization of polymer blends can be induced through the use of species that
react along the polymer-polymer interface to form block copolymers in situ. 24 In this spirit,
blends of isotactic polypropylene (iPP) and poly(lactic acid) (PLA) have been modified with
a maleic anhydride (MA)-grafted copolymer, which has been shown 25 to increase the impact
strength of the resultant alloy. Similarly, PP-g-MA has also been reported 26 to reactively
crosslink core-sheath and side-by-side bicomponent fibers composed of nylon-6 and iPP. The
concept of compatibilizing polymer-polymer interfaces that are not blended but rather
contacted, such as those involved in core-sheath fibers or laminates, with premade block
copolymers has largely been ignored. To the best of our knowledge, no prior studies
investigating the core-sheath compatibilization of bicomponent fibers by the addition of a
block copolymer to one component have been reported.
During large-scale mechanical deformation, melt-spun bicomponent fibers composed of
iPP and PLA are observed to possess interfacial voids that extend up to millimeters in length
along the fiber axis. These voids are attributed to inherently poor interfacial adhesion
between iPP and PLA. 4 In this work, we endeavor to compatibilize and thus improve the
mechanical properties of these two non-blended polymers by incorporating a triblock
264

copolymer during melt spinning. In addition to filament property assessment, the
morphology of the block copolymer in filaments is compared to that formed during melt
mixing without extrusion to elucidate the effect of high-shear spinning on block copolymer
structuration. Electron microscopy reveals that unique, nonequilibrium copolymer
morphologies are generated during bicomponent filament spinning. We also discuss the
application of steady-shear rheological methods to evaluate slip at layered iPP-PLA
interfaces and the effect of adding the block copolymer.
Experimental
Materials and Specimen Preparation
The iPP and PLA were provided by Sunoco Chemicals (Pittsburgh, PA; CP360H) and
NatureWorks (Minnetonka, MN; 6202D), respectively. A premixed compound containing a
poly[styrene-b-(ethylene-co-butylene)-b-styrene] (SEBS) triblock copolymer with 18.6 wt%
styrene (according to the manufacturer) and an overall molecular weight of 67 kDa
(according to independent size exclusion chromatography) was supplied by Kraton Polymers
(Houston, TX). While the identity of the compounding material is proprietary, it is midblock-
selective, which indicates that it is more aliphatic than aromatic in nature. We do not,
however, discount the possibility that the compounding species (hereafter referred to as the
"midblock extender") may be partially unsaturated, unlike the EB midblock of the
copolymer. Reagent-grade dichloromethane (DCM) was purchased from Mallinckrodt
Chemicals (Phillipsburg, NJ) and used as-received.
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Single and bicomponent filaments were melt-spun at various aspirator pressures on the
Partners’ Pilot Spunbond line located in the Nonwovens Cooperative Research Center at
North Carolina State University. All filaments examined here were spun with a total mass
throughput of 0.4 g/hole-min at a fiber composition of 50/50 (w/w) core/sheath. In select
cases, 5 wt% of the PLA was replaced with the SEBS copolymer. In these instances, the PLA
and copolymer were melt-compounded and co-extruded. Confluence of the molten PLA (or
PLA + SEBS) and iPP occurred in the spin pack. Below the spin pack, bicomponent
filaments were directed through the quench zone to an attenuation zone, where the aspirator
pressure controlled the air velocity around the fibers and effectively the fiber spinning
velocity. Non-bonded fibers were collected immediately following extrusion so that the as-
spun fiber morphology could be examined. Fiber diameters were used to calculate the
"spinning velocity" (V) at the point where the fibers solidified according to
V � Q
�A c
where Q is the mass flow rate of polymer per spinneret hole, ρ is the fiber mass density, and
Ac is the cross-sectional �� area of the fiber. For comparison with the bicomponent fibers,
corresponding polymer blends were prepared using a Haake-Buchler HBI System 90 twin-
screw melt mixer operated at 185°C. Binary iPP/PLA blends were immersed in DCM for
~100 h under constant agitation at ambient temperature to selectively dissolve the PLA for
morphological evaluation.
266
(1)

Specimen Characterization
X-ray diffraction (XRD) studies were conducted on a Bruker D-5000 diffractometer
(Madison, WI) equipped with a Highstar area detector and using CuKa radiation (λ = 0.1542
nm) at 40 kV and 30 mA. Resultant 2-dimensional XRD patterns were normalized with
respect to an empty sample holder and analyzed with the Bruker General Area Detector
Diffraction System (GADDS) software. Mechanical testing of fibers was conducted at
ambient temperature on an Instron Model 5544 extensiometer (Norwood, MA) fitted with a 5
N load cell. Single filaments with a gauge length of 28.6 mm were strained at a constant
crosshead speed of 25.4 mm/min. Data were analyzed with the Bluehill v.2 software
package, and a constant volume cylinder was assumed to calculate true stress. Representative
stress-strain curves were obtained for specimens prepared at different fiber configurations
and pressures after at least 10 trials. Optical microscopy of bicomponent fibers was
performed on a Mach-Zehnder type interference microscope by Aus
Jena (Jena, Germany) with polarized light (λ= 546 nm), and digital images were collected on
a CCD camera for birefringence analysis. Both scanning and transmission electron
microscopies (SEM and TEM, respectively) were employed to explore the morphologies of
the fibers and blends prepared here. After using DCM to dissolve the PLA from melt-
processed iPP/PLA blends, the remaining iPP matrix was sputter-coated with ~10 nm of Au
and subsequently analyzed by SEM performed in an environmental FEI XL-30 microscope
operated under high vacuum at 5 kV. The average size and standard deviation of the pores
267

introduced by extracting PLA were discerned by measuring the diameter of 100 pores using
the ImageJ software suite.
For complementary TEM examination of the block copolymer morphology, fibers were
conformally sputter-coated with 30 nm of Au (as a barrier layer to avoid fiber
contamination), embedded in epoxy and microtomed at ambient temperature on a Leica
UltraCut 7 with a diamond knife. Resultant sections were stained for 7 min with the vapor of
0.5% RuO4(aq), which is a selective stain for the phenyl groups on the S blocks of the SEBS
copolymer. Cross-sectional TEM images were acquired on a field-emission Hitachi HF2000
microscope operated at an accelerating voltage of 200 kV. Average microdomain sizes and
their corresponding standard deviations were determined by measuring 50-100 features of
interest, unless otherwise noted, using ImageJ. The zero-shear viscosities of the iPP and PLA
homopolymers with and without copolymer, as well as multilayered samples thereof, were
measured at 185°C under nitrogen on a TA Instruments AR-G2 rheometer equipped with
parallel plates measuring 25 or 40 mm in diameter. Multilayered specimens consisted of 1, 2,
4, or 8 alternating layers of iPP and PLA. Dynamic rheology was likewise performed on the
same instrument operated at 185°C with 25 mm plates and a 1 mm gap. Differential scanning
calorimetry (DSC) was conducted on a TA Instruments Q2000 model calorimeter calibrated
to an indium standard. Scans were carried out at heating rates of 10°C/min under 50 mL/min
N2 purge with samples of approximately 10 mg in standard aluminum pans.
268

Results and Discussion
Bicomponent iPP/PLA Fibers
Lipscomb 27 predicts that, for situations wherein the core polymer undergoes a greater
increase in viscosity during cooling than the sheath polymer, the core polymer bears more of
the spinline tension, which we hasten to add can serve to enhance crystallization and
molecular orientation. Similarly, Kikutani 6 reports that the solidification temperature and
viscosity disparities in co-spun polymers constitute the main factors influencing the "mutual
interactions" of the constituent polymers. This conclusion is interpreted to relate to the ability
of one polymer to direct or influence the crystallization and molecular orientation of the other
polymer. On one hand, we find that the crystallization of PLA is not strongly influenced by
its location (i.e., core or sheath) in bicomponent fibers, as seen in Figure A3.1. The iPP
crystal morphology, on the other hand, is sensitively affected by the fiber configuration. Our
observation that iPP tends to crystallize in the sheath but not in the core, coupled with the
results of Lipscomb 27 and our previous work 28 on the quiescent and stress-induced
crystallization of iPP, suggests that iPP crystallization is due to exposure to the quench air
while in the sheath and is not stress-induced while in the core. The presence of the iPP
mesomorphic phase for most spinning velocities indicates an absence of high tension on the
iPP core when co-spun with a PLA sheath (this may reflect the comparable melting
temperatures and zero-shear viscosities of iPP and PLA, as listed in Table A3.1). In marked
contrast, Kikutani et al. 10 have demonstrated that co-spinning iPP with poly(ethylene
terephthalate) (PET) into bicomponent iPPsheath/PETcore fibers results in sheaths that do not
269

solidify in the draw zone and fibers that could not be drawn as finely as either constituent
polymer. Since the difference in melting points between iPP and PET is significantly greater
(~100°C) than that between iPP and PLA, the reason why this problem is not encountered
here is because the melt viscosities and thermal transitions of iPP and PLA are sufficiently
similar so that neither polymer significantly influences the rheological properties or
solidification behavior of the other, compared to each polymer when spun individually.
According to the results displayed in Figure A3.1, an increase in aspirator pressure is
generally accompanied by a reduction (that is significant from 0 to 10 psi) in fiber diameter
for both single-component and bicomponent fibers. At a given aspirator pressure, the
diameters of bicomponent fibers are often marginally lower than the fiber diameters of either
polymer spun individually. This observation, along with the inherent thermodynamic
incompatibility between iPP and PLA (discussed further in the next section), implies that the
iPP/PLA interface may slip at high spin speeds, which occur at high shear rates. We return to
address the occurrence of slip between molten iPP and PLA, as determined by rheology,
later. Addition of the SEBS copolymer to PLA in bicomponent fibers is expected to influence
the iPP/PLA interface and, by inference, the fiber morphology. Figure A3.2 shows the fiber
diameters of bicomponent fibers co-spun with and without copolymer. As is evident in
Figures A3.1 and A3.2, an increase in aspirator pressure generally promotes a decrease in
fiber diameter. Compared to iPP spun alone, finer fibers result from co-spinning an iPP
sheath around a PLA core (Figure A3.2a), but the addition of copolymer appears to have a
non-systematic effect on fiber diameter. With PLA as the sheath (Figure A3.2b),
270

icomponent fibers are slightly smaller in diameter at moderate spin speeds, and the addition
of copolymer has little statistical effect on fiber diameter over the conditions examined.
Compatibilized iPP/PLA Blends
The interfacial energy between iPP and PLA is related to the surface energy of each polymer
according to Antonoff's rule:
� iPP�PLA � � iPP � � PLA (2)
From Eq. 2, the interfacial energy between iPP and PLA is estimated to be about 9
��
dyn/cm. 33,34 On the basis of the infinite-molecular-weight surface energies of the copolymer
constituents (included in Table A3.1), addition of a SEBS block copolymer to an iPP/PLA
blend lowers the interfacial tension at both the iPP-EB interface (3-5 dyn/cm) and the PS-
PLA interface (~1 dyn/cm). Although Antonoff’s rule is overly simplistic for many systems,
it has been shown 35 to accurately describe the interfacial energy between PS and PLA. The
interfacial energy provides a measure of thermodynamic incompatibility between two
dissimilar species and contributes to the enthalpic portion of the system free energy. As such,
it relates 36,37,38 directly to the Flory-Huggins � interaction parameter and, by extension, to the
corresponding difference in solubility parameters between the two species. The solubility
parameters of all polymer species of interest here are also listed for comparison in Table
A3.1. Solubility parameters are often used to estimate the compatibility of two polymers
insofar as the species are relatively non-polar and mixing is endothermic. While there is no
general guideline for predicting polymer-polymer miscibility from solubility parameters
alone, 30 polymer pairs with nearly identical solubility parameters are more likely to be
271

mutually miscible than those with even modestly different solubility parameters, regardless
of their chemical constitution. 39 On the basis of their solubility parameters, the EB midblock
of the copolymer should be compatible with iPP, whereas the S endblocks of the copolymer
are not expected to show much preference for either iPP or PLA. We cannot comment much
on the compatibility of the midblock extender in the copolymer, but it stands to reason that,
since it is mixed with the EB midblock, it, too, will be compatible with iPP.
To discern if the SEBS copolymer compatibilizes iPP/PLA blends, we examine the
morphologies of blends composed of 50 wt% iPP. Figure A3.3a displays a cross-sectional
SEM image of a melt-mixed iPP/PLA blend after removal of the dispersed PLA phase upon
selective solvent exposure in DCM at ambient temperature. The diameter of the PLA
domains (appearing as pores) is measured to be 5.0 ± 2.9 μm. Replacing 5 wt% of the PLA
with SEBS is found to reduce the PLA domain diameter to 1.4 ± 1.3 μm (cf. Figure A3.3b),
which confirms that the copolymer effectively compatibilizes the iPP/PLA blends, as
anticipated from the thermodynamic considerations discussed above. It immediately follows
that, if the copolymer molecules can locate along the interface between iPP and PLA, they
should reduce slip (if it exists) during extrusion and improve the adhesion between iPP and
PLA. It is of interest to note here that dissolution of the precompounded SEBS in DCM at a
concentration of 5 wt% results in the formation of a cloudy solution that appears to remain
stable to the unaided eye for several months at ambient temperature. We believe that the
observed solution opacity arises from the self-organization of copolymer molecules into
swollen micelles (spherical microemulsions) or vesicles 40 that are sufficiently large to scatter
272

light and capable of encapsulating the midblock extender, which is most likely incompatible
with DCM. A detailed account of this solution nanostructure is, however, beyond the scope
of the present work.
Multilayer Melt Rheology
Various polymer pairs such as PS/poly(methyl methacrylate), PS/HDPE, PS/iPP,
PE/fluoropolymer, and iPP/nylon-6 exhibit incompatibility-induced slip during rheological
testing of layered specimens. 8,9,11 Layered specimens are more representative of the present
bicomponent fibers than are blends, because the two polymer species are melt-contacted
along an artificially-introduced interface rather than along multiple interfaces that develop in-
situ due to thermodynamic instability. Generally speaking, "slip," or negative viscosity
deviation, tends to increase with increasing polymer-polymer incompatibility. To determine
whether iPP and PLA undergo slip during extrusion, rheological testing has been conducted
on alternating multilayers of iPP and PLA. As illustrated in Figure 4, these sandwich
structures are first positioned on the bottom flat plate of the rheometer, and then the top plate
of the rheometer is lowered until it contacts the top polymer layer. We have measured the
shear viscosity of molten polymer multilayers as a function of interfacial area, which is
controllably manipulated by increasing the number of alternating layers, as well as the size of
the plates. The zero-shear viscosities of these multilayers are presented as a function of
interfacial area in Figure A3.4. To put these results into perspective, a melt-spun
bicomponent fiber with an average inner diameter of 20 μm and measuring 1 m long would
possess an interfacial area of approximately 0.6 cm 2 . Despite the greater incompatibility
273

etween iPP and PLA as compared to other polymer pairs, no significant decrease in
viscosity is detected with increasing interfacial area up to an interfacial area of ~90 cm 2 ,
thereby indicating that slip between iPP and PLA is not detected by steady shear rheology at
stresses below ~1 kPa.
One reason for the absence of slip in this work may be due to the low shear rates
accessible with a parallel-plate rheometer relative to capillary rheometers or fiber extrusion.
Slip first becomes apparent, for instance, at interfacial areas between 34 and 152 cm 2 at shear
stresses above 1 kPa for multilayers composed of iPP and PS. 8 Inertial forces cause the
polymer to flow from the plates at stresses above 1 kPa, thereby resulting in poor data
quality. Smaller parallel plates exacerbate this problem. Alternatively, we consider the
manner by which previous studies concluded the existence of slip. While viscosities are
reported, 8,9 the deviation from predicted "no-slip" conditions is often small. Due to the large
variation in viscosity encountered in each sample, we have chosen to repeat each
measurement 7 times in this work. Jiang et al. 9 have reported the viscosity of a PS/HDPE
bilayer with an interfacial area of about 5 cm 2 to be 11% lower than the reciprocal rule of
mixtures (R-ROM). On the basis of this comparison, they conclude this deviation is
representative of slip. To discern the validity of this criterion, we apply several rules of
mixtures to the data in Figure 4. The first is the R-ROM, which has been used to predict the
viscosity of multilayered specimens composed of 2 polymers (indicated by subscripts 1 and
2). It is given by
��
1
� � �1 �
�1 �2 �
�2 �i �i 274
(3)

where the subscripted i corresponds to the interfacial layer, which has few polymer
entanglements and potentially lower viscosity. 41 In Eq. 3, � is the measured viscosity of the
multilayer, , �1 and �2 are the volume fractions of polymers 1 and 2, respectively, and �1 and
�2 are the viscosities of polymers 1 and 2, respectively, measured individually. Since no
difference is evident in Figure 4 up to an interfacial area of ~90 cm 2 , this interfacial layer
contribution is neglected.�
Alternatively, other rules of mixtures should be considered. For example, the standard rule
of mixtures (S-ROM) predicts viscosity according to
��
� � � 1 � 1 � � 2 � 2
Moreover, a logarithmic rule of mixtures (L-ROM) has also been proposed for polymer
blends, viz,
14 14 14 42
log� � � 1 log� 1 � � 2 log� 2 (5)
Other rules include additional terms that attempt to explain deviations from the three
provided in Eqs. 3-5
��
and are not considered here. The multilayer viscosity averaged over all
the interfacial areas in Figure 4 is 754 Pa-s, whereas those predicted by Eqs. 3-5 are 714 (R-
ROM), 746 (S-ROM) and 729 Pa-s (L-ROM). Thus, the S-ROM most closely predicts the
viscosity of our multilayered samples, while the R-ROM is found to give the poorest
prediction of the three. The standard errors of the PLA and iPP viscosities used in the ROM
predictions are 9 and 3%, respectively, which suggests that measured multilayer viscosities
deviating modestly (by 11% according to Jiang et al. 15 ) from predicted values are not
statistically significant in systems composed of iPP and PLA. The relatively large variation in
275
(4)

measured multilayer and pure-component viscosities encountered here is indicative of the
difficulty in confirming the existence of interfacial slip between iPP and PLA on a parallel-
plate rheometer.
Fiber Mechanical Properties
The interface, which is largely responsible for the mechanical properties of bicomponent
fibers (as well as fibers derived from polymer blends), is known 4 to fail during the
mechanical drawing of bicomponent fibers spun from iPP and PLA. To reduce the inherent
incompatibility between iPP and PLA, the SEBS copolymer is compounded with PLA before
bicomponent fibers are co-spun. When PLA with and without copolymer is spun as the core
(Figure A3.5a), the tenacity of the bicomponent fibers increases with increasing aspirator
pressure up to a level near 2 cN/dtex. Incorporation of the copolymer has no systematic
effect, positive or negative, on these results. Recall from Figure A3.1 that the PLA and iPP
are both semi-crystalline in this fiber configuration. If the fibers are spun with an iPP core
and PLA sheath, a similar dependence on aspirator pressure is seen (Figure A3.5b), but the
maximum level attained is slightly lower for fibers with PLA only and lower still (but within
experimental uncertainty) for fibers with a mixture of PLA+SEBS. In this configuration, the
PLA is semi-crystalline, but the iPP is, for the most part, mesomorphic. Recall that, while the
EB midblock of the SEBS copolymer is compatible with iPP, the S endblocks are not
expected, on the basis of solubility parameters alone, to be strongly iPP- or PLA-compatible.
Blending incompatible polymers into bicomponent fibers is expected to decrease fiber
strength due to enlarged interfacial area and reduced interfacial strength. Yet, addition of the
276

SEBS copolymer to PLA as the core (Figure A3.5a) or sheath (Figure A3.5b) does not
compromise the mechanical properties of the fibers, indicating that the copolymer (with
midblock extender) and PLA are not strongly incompatible.
At the lowest aspirator pressure (fiber spinning velocity) in Figure A3.6, addition of
copolymer into the PLA sheath is observed to increase substantially the strain at which
rupture of the sheath occurs. As the aspirator pressure is increased, however, the strain at
which PLA ruptures does not appear to be dependent on copolymer addition. Note that not all
fibers undergo failure of the PLA sheath at lower strains than the iPP core. Rather, in these
cases, a single catastrophic failure signals nearly simultaneous rupture of both polymers.
Such events are accompanied by an abrupt increase in the true strain at break of the fibers.
By adding the SEBS copolymer to PLA, the frequency at which the PLA sheath ruptures
prior to failure of the iPP core is reduced at most aspirator pressures (cf. Figure A3.6), which
indicates that the SEBS-modified PLA is capable of undergoing greater extension prior to
failure. In the event that the PLA sheath does not fail before the iPP core (i.e., the sheath and
core co-rupture) in any of the tested specimen, the frequency in Figure 6 is shown as 0%.
While the Tg (Figure A3.7) and melt viscoelasticity (Figure A3.8) of PLA are not greatly
affected by the incorporation of SEBS, the elastomeric nature of the copolymer, coupled with
its intrinsic ability to self-organize into nanostructural elements, may serve to improve the
elasticity of the PLA sheath, which, in turn, promotes an increase in the strain at which the
PLA sheath ruptures. In Figure A3.7, addition of the SEBS copolymer alters the Tg of PLA
by just over 1°C, which is considered to be within instrumental uncertainty. The melt
277

frequency (�) spectra displayed in Figure A3.8 show little variation upon addition of the
SEBS copolymer to PLA. In both cases, the dynamic loss modulus (G") scales as � 1.0 in the
terminal region, which agrees with the behavior of entangled homopolymers (� 1 ). The
analogous scaling behavior of the dynamic storage modulus (G'), which is � 2 for entangled
homopolymers, changes upon copolymer addition from � 1.9 to � 1.6 . While this subtle
reduction in slope may signify a copolymer-induced change in the molten structure, it is
reasonable to expect that such a change might be more apparent in the solid state (where the
styrenic endblocks are glassy) than in the melt. The copolymer-induced reduction in the
frequency of PLA rupture at low strains may also indicate improved adhesion along the
iPP/PLA interface, which would promote greater stress transfer to the iPP core and which
would permit the PLA to support a higher stress before rupturing.
Previously, some of us have introduced 4 a facile optical method for estimating the
molecular orientation of bicomponent core/sheath fibers from true stress-true strain curves
when the refractive indices of the core and sheath polymers significantly differ. Figure A3.9
shows the progression of matching a fiber with an iPP core and PLA sheath to liquids of
known refractive index so that the birefringence of the sheath, which is related to molecular
orientation and fiber strength, can be determined. These images confirm that the conformal
PLA sheath completely wets the iPP core. The liquid surrounding the fiber must possess a
refractive index (RI) that is similar to those of the two species in the fiber to yield clear
fringes (dark bands) that can be followed through each interface. Because the refractive
indices of the iPP core and PLA sheath are considerably different (1.504 and 1.542,
278

espectively, at 20°C for unoriented samples), no single liquid can provide clear fringes for
both the core and sheath simultaneously, which explains why the fringes blur at the core-
sheath and/or sheath-liquid interface in each of the images in Figure A3.9. In addition, while
the interface of unstrained fibers appears continuous for the fibers in this study, birefringence
measurements would be difficult or altogether impossible for strained iPPcore/PLAsheath fibers
exhibiting voids at the polymer-polymer interface. Therefore, using the method described
elsewhere, 4 we estimate the molecular orientation of PLAcore/iPPsheath fibers with and without
added SEBS copolymer using stress-strain curves.
The strain shift of single-component and bicomponent fibers, which relates proportionally
to molecular orientation, is presented as a function of spinning velocity in Figure A3.10.
Here, spinning velocity, rather than aspirator pressure, is used to facilitate comparison of
birefringence from fibers produced by other spinning methods, such as melt spinning or
electrospinning, that do not require an aspirator pressure. Bicomponent fibers exhibit
markedly increased strain shift relative to iPP and PLA alone, and addition of the SEBS
copolymer to the PLA core does not inhibit molecular orientation in the iPP sheath. This
observation is consistent with previous studies 4,6 reporting that the sheath component dictates
flow mechanics for core/sheath fiber extrusion, as well as the ultimate properties of
bicomponent fibers, since the sheath component experiences the greatest stress in both the
melt and solid phases. Spinning with an iPP sheath and a PLA (or PLA+SEBS) core serves to
focus the spinline stress over a smaller cross-sectional area, which results in higher molecular
orientation in the iPP sheath. While the sheath component may control flow mechanics, the
279

core polymer is crucial to achieve the synergistic effect of enhanced molecular orientation in
the sheath polymer beyond that which can be realized by spinning either polymer
individually. The master stress-strain curve for PLA required in this methodology has been
developed for PLA alone, not with SEBS copolymer. Therefore, we do not report strain shifts
for iPPcore/PLAsheath fibers with and without SEBS to avoid attributing effects introduced by
SEBS to changes in PLA molecular orientation.
Nonequilibrium Copolymer Morphologies
The morphologies of triblock copolymers, such as the SEBS copolymer employed in this
study, have been the subject of numerous studies. Most commercial triblock copolymers are
designed as thermoplastic elastomers with dispersed glassy microdomains embedded in, and
connected to, a continuous, rubbery matrix. 46 Addition of a midblock-selective oil, 47
tackifying resin 48 or homopolymer 49 to a triblock copolymer can yield the same
morphologies observed in diblock copolymer/homopolymer blends, 50,51 in which case
comparable design rules can be sensibly presumed. It immediately follows that, on the basis
of the copolymer composition, the proprietary midblock extender added to the copolymer
used here promotes a spherical or cylindrical morphology if the extender is largely or
marginally miscible, respectively, with the EB midblock. Incorporation of the copolymer into
PLA adds another level of complication, as the copolymer must now partition between the
midblock extender and PLA, as well as interact with iPP along the iPP-PLA interface. Under
ideal equilibrium conditions of slow solvent evaporation or melt mixing, followed by
extensive solvent or thermal annealing, the resultant copolymer morphology may be
280

complex, depending on the individual strengths of 6 different binary interactions (assuming
that the EB midblock and midblock extender can each be treated as a single species). If the
rapid melt processing of the bicomponent fibers modified by the copolymer is now
considered, highly nonequilibrium morphologies can be reasonably expected.
We begin with an overview of the morphologies of the SEBS copolymer in different
fibers. Figure A3.11 shows a series of TEM images acquired from fibers varying in
configuration, but spun at the same aspirator pressure (15 psi). In Figure A3.11a, a relatively
low-magnification image of a bicomponent fiber composed of an iPP core and PLA sheath
demonstrates that (i) the conformal Au coating around the edge of the fiber prevents swelling
of the fiber with epoxy resin, which reacts with the vapor of RuO4(aq); (ii) the iPP-PLA
interface is clearly differentiated in cross-section; and (iii) the iPP core is lightly stained by
the vapor of RuO4(aq). An image of a PLA fiber containing 5 wt% SEBS copolymer is
provided in Figure 11b and reveals that the copolymer molecules, selectively stained with the
vapor of RuO4(aq), are present in the form of discrete and aperiodic nanostructures that most
closely resemble dispersed tubules. The existence of tubules suggests that the midblock
extender is either not highly compatible with the EB midblock of the copolymer or present at
sufficiently high concentration to preclude appreciable solubilization within the copolymer
matrix. Moreover, assuming that the copolymer is uniformly dispersed in the PLA prior to
melt spinning, the styrenic endblocks of the copolymer do not appear to be very compatible
with PLA.
281

In the TEM images displayed in Figures A3.11c and A3.11d, the copolymer morphologies
formed in bicomponent iPPcore/(PLA+SEBS)sheath and (PLA+SEBS)core/iPPsheath fibers are
evident. In both cases, the copolymer forms dispersed nanostructures that remain distributed
throughout PLA rather than accumulating along the iPP-PLA interface. This tendency is
consistent with the property measurements provided earlier and explains why the copolymer
does not significantly alter the breaking strength of the bicomponent fibers investigated here.
Although thermodynamic considerations indicate that the copolymer should migrate to the
iPP-PLA interface, the timescale associated with fiber spinning is on par with or faster than
that required for the diffusion of individual copolymer molecules from the PLA phase to the
interface. To complicate matters further, the copolymer molecules are assembled into
nanostructures that are less driven (due to a concentration gradient) and slower to diffuse to
the interface where they are needed to compatibilize the core and sheath. Careful
examination of the nanostructures observed in Figures A3.11c and A3.11d reveals that the
copolymer nanostructures tend to orient along a common direction (which may be normal to
the iPP-PLA interface) and, more importantly, that the copolymer molecules appear to self-
organize into tubules, not cylinders, in PLA co-spun with iPP. [Stained cylinders appear the
most electron dense (darkest) along their centerline, whereas tubules are darkest along their
periphery, due to thickness considerations in projection.] While bicomponent block
copolymers have been previously reported 52,53 to form nanotubes, such morphologies
normally require an additional driving force, such as crystallization, to do so.
282

While it is intriguing that in all the core/sheath fiber configurations examined the
copolymer molecules most often form tubules within PLA, more exotic morphologies, such
as concentric tubules and spheres in tubules (what we term "peas in a pod"), are also
observed, as evidenced by the TEM images provided in Figures A3.12a and A3.12b,
respectively. Examples of the "peas in a pod" morphology are also highlighted in Figures
A3.11c and A3.11d. A qualitatively similar sphere-in-cylinder morphology has been
observed 54 in a thin film of an ABC triblock copolymer swollen with a good, neutral solvent.
To the best of our knowledge, however, this morphology has not been previously reported for
a blend of an ABA triblock copolymer dispersed in a C homopolymer. These highly non-
classical morphologies are schematically depicted in Table A3.2, and relevant dimensions
identified in the illustrations and measured from TEM images are included. To put these
dimensions into perspective, the unperturbed gyration diameter of the blocks comprising the
SEBS copolymer are calculated from the freely-jointed chain model 55 and the known block
lengths. Since the EB midblock most likely adopts a looped or bridged conformation (which
is rigorously true only at equilibrium), its molecular weight is halved so that it can be treated
as a tail, a chain tethered at only one end, in similar fashion as the S endblocks. Block
gyration diameters, estimated with the assumptions that (i) the statistical segment lengths of
S and EB are comparable (~0.7 nm) 56,57 and (ii) the EB midblock consists of equal fractions
of E and B, are ~4 nm for the S block and ~15 nm for (half) the EB block. The number of
unperturbed blocks (N) corresponding to the measured dimensions identified in the diagrams
shown in Table A3.2 is included in the same table and reveals several important features.
283

Analysis of the tubule walls generally yields N ≈ 2, which corresponds to an endblock
bilayer. This finding is consistent with the EB midblocks extending into both the tubular core
and the surrounding matrix. In the case of a single tubule (morphology A in Table A3.2), the
measured internal diameter also results in N ≈ 2 (i.e., 2 EB blocks) for tubules formed in the
(PLA+SEBS)sheath configuration, but varies considerably from about 1 to 3 in the
(PLA+SEBS)core configuration. This variation may reflect differences in the level of spin-line
stress experienced by the copolymer molecules, as well as stochastic swelling of the EB
midblock by the midblock extender or PLA. Similar variation is evident in the EB-rich
regions of the concentric tubule morphology (morphology B in Table A3.2). Within the
PLA+SEBS core, a single internal tubule appears highly swollen with N ≈ 8, but the distance
between the internal and external walls yields N ≈ 1, which implies that the EB midblocks
extending from the internal and external tubule walls are interdigitated into a monolayer,
rather than bilayered. Lastly, in the "peas in a pod" morphology (morphology C in Table
A3.2), the diameter of the internal spheres (which appear circular but, in a few cases, as
shells) is significantly larger than two S endblocks (N ≈ 4). These enlarged spheres are
presumed to be the result of partial PLA incorporation, although the possibility of other
kinetically trapped species cannot be outright disregarded. The distances between
neighboring spheres (which align along the direction of the tubule) and between the spheres
and tubule wall in both the core and sheath fiber configurations yields N ≈ 1, which strongly
suggests that the spheres are most likely connected together, as well as to the tubule walls, by
bridged EB midblocks, as schematically depicted in Figure A3.13. Such molecular
284

connectivity within the nanostructures present might serve to reinforce the PLA phase and
improve its elasticity, as deduced from the flow and mechanical properties discussed earlier.
It is comforting that these unique morphologies, although nonequilibrium in nature, tend to
obey classical chain packing behavior.
In addition to establishing the presence of uncommon copolymer morphologies in
bicomponent fibers, the images presented in Figures A3.11b, A3.11c, A3.12a and A3.12b
confirm that the S endblocks are not in direct contact with PLA. Rather, the EB midblock
forms a contiguous coronal layer around the S-rich features. Such isolation helps to explain
why the copolymer does not preferentially migrate to and accumulate along the iPP-PLA
interface where it can promote compatibilization. To discern the extent to which the SEBS
copolymer is dispersed within PLA prior to melt spinning, we have examined the melt-
compounded PLA+SEBS mixture. Large structures such as those portrayed in Figure A3.12c
are evident, indicating that the two materials are not thoroughly mixed. These
macrostructures, measuring on the order of hundreds of nanometers across, also appear
tubular (the figure displays a central "ring" and what appears to be an end cap). Close
examination of their walls reveals an organized copolymer nanostructure. Complementary
inspection of the as-received copolymer with midblock extender (Figure A3.12d) confirms
the existence of an irregular morphology that resembles the nanostructured walls of the
macroscale tubules in Figure A3.12c. Thus, we conclude that the copolymer retains some of
its as-received nanostructure after being melt-blended with PLA, indicating that the
285

compounding temperature and mechanical mixing employed here were insufficient to
molecularly disperse the copolymer within PLA.
On a side note, it is strangely curious, though, that the stainable, but minor, styrenic
component in the as-received compound with the midblock extender appears as the matrix in
Figure A3.12d, which suggests that the morphology may be more complicated than styrenic
spheres or cylinders as previous anticipated. Although this figure is representative of entire
sections of the as-received copolymer, it is likewise possible that the as-received copolymer
is heterogeneous at macroscopic length scales. Such macroscale heterogeneity may be largely
responsible for the unique nanoscale tubular morphologies reported here in Figures A3.11
and A3.12, as they could be the result of subjecting the parent morphologies in the
heterogeneous PLA+SEBS mixture to very high shear and extensional flow through the
spinneret so that they became distorted and rearranged into lower-energy nanostructural
elements. This nonequilibrium formation mechanism seems more plausible than conventional
self-organization of the copolymer from a disordered state and would help to explain the
variety of nanostructures observed on the basis of local copolymer composition and
diffusion.
Conclusions
The strategy of copolymer-induced blend compatibilization has been applied to melt-spun
bicomponent fibers that bring together iPP and PLA in contact along a single interface
separating the core from the sheath. Although the two polymers are incompatible, there is no
evidence of measurable slip at the molten iPP-PLA interface according to steady-shear
286

heology. While the addition of a SEBS copolymer to a melt-mixed iPP/PLA blend serves as
a compatibilizer by reducing the size of PLA domains arising from macrophase separation,
incorporation of the copolymer into PLA prior to co-spinning does not drastically improve
the properties of bicomponent iPP/PLA fibers, which implies that the copolymer molecules
are unable to concentrate along the iPP-PLA interface during spinning. Despite an absence of
copolymer along the fiber interface, the strain at which PLA as the sheath ruptures increases
sharply at low spinning pressure, and the number of fibers that undergo failure of the PLA
sheath prior to failure of the iPP core is reduced by up to 30%, upon addition of copolymer.
Thus, although the copolymer does not modify the iPP-PLA interface, it does affect the host
PLA by improving its elasticity through the formation of unique copolymer nanostructures
that include single tubular, concentric tubular and "peas in a pod" morphologies. Careful
analysis of such unexpected morphologies confirms the existence of nanostructural
dimensions capable of accommodating local connectivity through midblock bridging, which
consequently allows these highly-elastic, EB-rich copolymer dispersions to rubber-toughen 58
the PLA.
Acknowledgments
Financial support of this work has been provided by the Nonwovens Cooperative
Research Center at North Carolina State University and the National Science Foundation
through a Graduate Research Fellowship (K. E. R.). We thank Dr. Behnam Pourdeyhimi for
insightful and fruitful discussions, as well as Christina Tang and Alina Higham for analytical
assistance.
287

Tables
Table A3.1. Physical properties of the polymers employed in this study.
Polymer
Tm 29 (°C)
288
Zero-shear
viscosity at
185°C
(Pa-s)
Surface free
energy 30 at 20°C
(dyn/cm)
Solubility
parameter 31,32
(cal/cm 3 ) 1/2
Ipp 165 a 901 a 30.1 7.4
PLA 161 a 591 a 39.3 12.1
SEBS copolymer — 6980 a — —
Atactic polystyrene — — 40.7 9.5
Low-density polyethylene — — 35.3 8.0
Polybutylene — — 33.6 7.8
a property was measured in the present study.

Table A3.2. Dimensions (in nm) of the SEBS copolymer morphologies observed in melt-
spun bicomponent fibers.
A B C
Morphology A Morphology B Morphology C
Core/sheath fiber d t di do d ds s g
configuration a (n) (n) (n) (n) (n) (n) (n) (n)
PLA + SEBS 51 ± 31 8 ± 4 126 225 56 ± 16 16 ± 6 10 ± 4 17 ± 4
in core (50) (50) (1) (1) (12) (32) (32) (32)
PLA + SEBS 30 ± 10 8 ± 2 70 100 42 ± 15 14 ± 4 13 ± 4 16 ± 5
in sheath (50) (50) (2) (2) (13) (56) (56) (56)
a n denotes the number of features measured.
289

Figure A3.1. Filament diameter and polymer morphology as functions of aspirator
pressure and fiber configuration, denoted by core (c) and sheath (s). While all filaments are at
least partially amorphous, only crystalline or otherwise ordered morphologies are identified
by symbols if present: �-crystalline iPP (filled squares), mesomorphic iPP (filled triangles)
and crystalline PLA (open squares). Filaments spun at the same aspirator pressure with
different configuration are connected by solid lines (color-coded and labeled).
290

Figure A3.2. Filament diameter as a function of spinning pressure for different fiber
configurations: (a) iPP and iPP in the sheath, and (b) PLA and PLA in the sheath.
Homopolymers, bicomponent filaments without copolymer and bicomponent filaments with
copolymer are color-coded black, blue and red and are correspondingly labeled. Error bars
indicate the standard error in the data from ~10 trials, and the solid lines serve to connect the
data.
291

Figure A3.3. Cross-fracture SEM images of melt-mixed blends: (a) 50/50 w/w iPP/PLA
and (b) 50/47.5/2.5 iPP/PLA/SEBS. The PLA is selectively removed upon immersion in
DCM for ~100 h at ambient temperature.
292

Figure A3.4. Zero-shear viscosities of iPP/PLA multilayers in different test
configurations: PLA on the bottom with either 25 ( ) or 40 ( ) mm plates, or iPP on the
bottom with 25 mm plates ( ). Included are the individual viscosities of iPP ( ) and PLA (
). The error bars denote the standard error in the data, and the dashed line corresponds to the
S-ROM prediction. The uncertainty on the prediction is based on the standard deviation of
the component iPP and PLA viscosities used to calculate the viscosity of the multilayers.
Included are schematic illustrations of the iPP/PLA multilayers: (a) PLA (blue) on bottom,
(b) iPP (red) on bottom and (c) 8 alternating layers.
293

Figure A3.5. Dependence of tenacity of iPP/PLA filaments on aspirator pressure for two
different fiber configurations identified by the location of the iPP: (a) sheath and (b) core.
Open and filled symbols identify specimens with and without added SEBS copolymer,
respectively. Error bars indicate the standard error in the data from ~10 trials.
294

Figure A3.6. (left axis) True strain at which the sheath of iPPcore/PLAsheath bicomponent
filaments with ( ) and without ( ) added SEBS copolymer fail as a function of aspirator
pressure. (right axis, red) Frequency of PLA sheath failure prior to iPP core rupture with ( )
and without ( ) added copolymer. The solid lines serve to connect the data of systems
without copolymer.
295

Figure A3.7. Series of DSC thermograms collected at a cooling rate of 10°C/min from
iPPcore/PLAsheath bicomponent filaments with and without SEBS copolymer added to the PLA
phase at two aspirator pressures (labeled). The location of the PLA Tg in each system is
marked.
296

Figure A3.8. Frequency (�) spectra of the dynamic moduli (G', circles; G", triangles)
measured from PLA (filled symbols) and PLA blended with 5 wt% SEBS (open symbols) at
185°C and a stress of 5 Pa (in the linear viscoelastic limit). Shown for comparison are the
scaling relationships expected for G' and G" in the low-� (terminal) region for entangled
homopolymers.
297

Figure A3.9. Optical micrographs of PPcore/PLAsheath bicomponent filaments (see
schematic diagram) digitally recorded under polarized light to measure the birefringence
according to ref. [43]. The measured birefringence can be related to molecular orientation
and, thus, mechanical properties, as described elsewhere. 44,45 The refractive index (RI)
provided on each image corresponds to that of the liquid surrounding the fiber.
298

Figure A3.10. Dependence of the strain shift of iPP ( ), PLA ( ) and PLAcore/iPPsheath (
) filaments on spinning velocity. Included are results ( ) from bicomponent filaments into
which 2.5 wt% SEBS copolymer is added to the PLA phase. The lines serve as guides for the
eye.
299

Figure A3.11. Cross-sectional TEM images of filaments in different configurations. In (a),
an iPPcore/PLAsheath bicomponent filament coated in Au and embedded in epoxy (labeled) is
shown. A PLA filament modified with 5 wt% SEBS copolymer is presented in (b). In (c) and
(d), iPPcore/(PLA+SEBS)sheath and (PLA+SEBS)core/iPPsheath bicomponent filaments are
displayed. The styrene-rich copolymer nanostructures appear dark due to selective staining.
In (c) and (d), the arrowheads identify a unique nanostructure: tubules with internal spheres
("peas in a pod").
300

Figure A3.12. Series of TEM images showing unexpected morphologies of the SEBS
copolymer after being compounded with PLA and melt-spun into bicomponent filaments. In
(a), single tubules and concentric tubules are evident, whereas the "peas in a pod"
morphology is clearly seen in (b). Large-scale structures that appear vesicular with ordered
copolymer walls are visible in PLA melt-compounded with the SEBS copolymer prior to cospinning
(c). The morphology of the as-received copolymer with midblock extender is
provided for reference in (d).
301

Figure A3.13. Schematic illustrations of the unexpected SEBS morphologies observed in
melt-spun bicomponent filaments composed of iPP and PLA: (a) single tubules with
bilayered S (red) walls, (b) tubules swollen with either the midblock extender or entrapped
PLA (green), (c) concentric tubules that are connected by bridged EB midblocks (blue), (d)
concentric tubules swollen with either the midblock extender or entrapped PLA, (e) equally
sized and spaced S spheres in a single tubule ("peas in a pod"), and (f) equally sized and
spaced hollow S spheres in a single tubule.
302

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306

APPENDIX IV
Enhanced Biomimetic Performance of Ionic Polymer-Metal Composite
Abstract
Actuators Prepared with Nanostructured Block Ionomers*
Ionic polymer-metal composites (IPMCs) represent an important class of stimuli-
responsive polymers that are capable of bending upon application of an electric potential.
Conventional IPMCs, prepared with Nafion ® and related polyelectrolytes, often suffer from
processing challenges, relatively low actuation levels and back relaxation during actuation. In
this study, we examine and compare the effects of fabrication and solvent on the actuation
behavior of a block ionomer with a sulfonated midblock and glassy endblocks that are
capable of self-organizing and thus stabilizing a molecular network in the presence of a polar
solvent. Unlike Nafion ® , this material can be readily dissolved and cast from solution to yield
films that vary in thickness and exhibit enormous solvent uptake. Cycling the initial chemical
deposition of Pt on the surfaces of swollen films (the compositing process) increases
theextent to which the electrodes penetrate the films, thereby improving contact along
thepolymer/electrode interface. The maximum bending actuation measured from IPMCs
*This chapter has been published in its entirety:
PH Vargantwar, KE Roskov, TK Ghosh, RJ Spontak.. “Enhanced Biomimetic Performance of Ionic Polymer-
Metal Composite Actuators Prepared with Nanostructured Block Ionomers.” Accepted to Macromolecular
Rapid Communications 9/2011.
307

prepared with different solvents is at least comparable, but is often superior, to that reported
for conventional IPMCs, without evidence of back relaxation. An unexpected characteristic
observed here is that the actuation direction can be solvent-regulated. Our results confirm
that this block ionomer constitutes an attractive alternative for use in IPMCs and their
associated applications.
Introduction
Electroactive polymers (EAPs) constitute a rapidly growing genre in the blossoming field
of stimuli-responsive polymers [1] and afford a wide variety of important material-related
advantages over their inorganic counterparts. [2] Some of these benefits include light weight,
low cost, mechanical robustness, facile processability, and straightforward scalability.
Macromolecules classified as EAPs are further categorized according to the mechanism by
which they undergo actuation upon application of an electrical potential. In the case of
electronic EAPs, the mode of actuation depends on whether an applied field either promotes
attraction of compliant, oppositely-charged surface electrodes to compress a low-modulus
polymer (the electrostatic mechanism) or induces a change in molecular polarization and,
hence, lattice dimensions of a semi-crystalline polymer (the electrostrictive mechanism). [3]
Examples of electrostatic and electrostrictive EAPs are dielectric elastomers [4] and
ferroelectric polymers, [5] respectively. Actuation can likewise occur by an ionic mechanism
wherein ionic species migrate from one electrode to another and, by doing so, induce a
change in the size and/or shape of a solvated polymer. Although systems developed with
carbon nanotubes [6] fall into this category, we only consider further here ionic polymer-metal
308

composites (IPMCs), wherein ion transport is accompanied by solvent diffusion and
sufficient nonuniform polymer swelling to trigger a bending motion. [7,8] Due to their inherent
similarity to biological muscle, IPMCs are often referred to as artificial muscle, and their
electrical actuation is considered biomimetic. They can be used in motion-control devices [9]
(e.g., micropumps, grippers, camera apertures, and catheters) and are alternatively suitable
for energy harvesting [10] and vibration sensing. [11] Since the transport of ionic species
necessitates a solvent-swollen hydrophilic polymer, most conventional IPMCs consist of a
polyelectrolyte such as Nafion ® , a sulfonated fluorocarbon-based copolymer. In this study,
we report how, with appropriate preparation, a block ionomer can provide comparable, if not
superior, electroactuation performance.
Block ionomers derive from block copolymers, which have become increasingly
ubiquitous in fundamental and technological endeavors due largely to their unique ability to
self-assemble into soft nanostructures. [12] Within this broad classification, linear ABA
triblock copolymers and higher-order multiblock copolymers with at least one midblock
provide the added, and often desirable, benefit of forming a contiguous network consisting
predominantly of bridged midblocks that physically connect the copolymer nanostructure. [13]
In this study, we focus on a block ionomer derived from a network-forming copolymer in the
presence of a midblock-selective, low-volatility polar solvent. Consider, for illustrative
purposes, an ABA triblock copolymer preferentially solvated with a B-selective solvent
under equilibrium conditions. In the limit of high solvent concentrations, copolymer
molecules self-organize into spherical micelles, each with an A core and a B corona. [14]
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While the long-range order of these micelles may become compromised, the network
connecting the micelles remains intact unless the copolymer concentration is reduced below
the critical gel concentration. Within a midblock-solvated network, micelles composed of
glassy or semi-crystalline endblocks serve as network-stabilizing cross-link sites, thereby
making such elastomeric gels suitable for a wide range of applications, including dielectric
elastomers, [15] pressure-sensitive adhesives [16] and microfluidic substrates. [17] Recently, we
have provided [18] evidence to show that this soft nanostructure design can be used in
conjunction with a midblock-sulfonated pentablock ionomer (PBI) to generate highly
responsive and stable IPMCs. In this work, we more closely explore the performance of these
IPMCs under different preparation conditions and in different solvents.
The PBI employed here is a poly[p-t-butyl styrene-b-(ethylene-alt-propylene)-b-(styrene-
co-styrenesulfonate)-b-(ethylene-alt-propylene)-b-p-t-butyl styrene] ionomer, the chemical
structure of which is displayed in Fig. A4.1a. While previous studies [19] have shown that
endblock-sulfonated styrenic triblock copolymers can yield IPMCs, it must be recognized
that the endblocks in the current system are incompatible with polar solvents and therefore
form the glassy microdomains required for network stabilization, as schematically depicted
in Fig. A4.1b. In the presence of a polar solvent, the sulfonated midblock of each molecule
becomes highly swollen in a solvent-rich matrix and, for the most part, adopts either a
bridged or looped conformation. Under equilibrium conditions (achieved, for instance, by
mixing the copolymer and solvent together, followed by increasing the temperature to form a
solution and then decreasing the temperature to promote copolymer self-assembly),
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endblock-rich micelles surrounded by a rubbery shell uniformly disperse throughout the
solvent-rich matrix. The present study, however, utilizes a nonequilibrium process strategy
(detailed in the Experimental Section) to generate nanostructured mesogels, [20] which retain
morphological features of the solvent-free copolymer. The existence of a bridged-midblock
network connecting, and stabilized by, glassy endblock-rich microdomains can be confirmed
by dynamic rh
acquired at ambient temperature in the linear viscoelastic regime are shown in Fig. 1c for the
PBI selectively swollen with two hydrophilic solvents used here: glycerol (GLY) and
ethylene glycol (EG) at their maximum solvent uptake levels (~450 and ~500 wt% GLY and
EG, respectively). Two signature features evident in this figure are consistent with a
physically cross-linked network: [21] (i
(confirming the load-bearing capability of the network), and (ii) G' is weakly (if at all)
and nearly parallel to G".
In general, IPMCs are fabricated from a polyelectrolyte membrane whose surfaces are
coated with metals such as Pt, Au or Ag. [7,22,23] The IPMCs generated here begin with water-
swollen PBI films that are subjected to Pt chemical deposition to metalize the inner surface,
followed by surface electroding with Ag. Since the electrodes must establish a large contact
area with the polyelectrolyte to achieve satisfactory capacitance, the electrode morphology
constitutes a critical design consideration. Good interfacial adhesion between the electrode
and polyelectrolyte is required for an efficient electric double layer, which, in turn, governs
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the storage capacity (and actuation strain) of IPMCs. [24] During conventional IPMC
fabrication, a polyelectrolyte membrane is subjected to surface roughening and mechanical
prestrain to improve the adhesion and distribution of the electrode metal along the
polymer/metal interface. Since PBI swells extensively in deionized water at elevated
temperatures and then collapses in the tetraamine platinum(II) chloride hydrate solution used
for Pt electroding (due to salting-out), no such surface pretreatment is required in the present
study. This feature, coupled with our ability to vary the number of Pt coatings applied,
permits precise control over the distribution of Pt to desired depths within the membrane,
which dictates the effective interfacial area and, in so doing, governs the actuation
performance, as discussed later. On the basis of the findings reported by Lu et al., [25] we first
apply one or more Pt inner-surface electrode layers, followed by an Ag surface electrode
layer, [23] to improve actuation performance. The distribution of Pt in the IPMC after each Pt
deposition cycle has been probed by scanning electron microscopy (SEM) and energy-
dispersive spectroscopy (EDS), and results obtained after 3 consecutive cycles are presented
in Figs. A4.2a-c. The cross-sectional SEM image displayed in Fig. A4.2a reveals that the PBI
membrane is sharply delineated from the top and bottom electrode layers composed of Pt and
Ag. The Pt layer is not readily observed after a single coating, but it becomes increasingly
evident upon deposition cycling.
Corresponding elemental x-ray maps acquired by EDS are shown in Figs. A4.2b-c for 3
Pt deposition cycles and confirm the spatial distribution of Pt (Fig. A4.2b) and Ag (Fig.
A4.2c). It is important to recognize that, although Ag is restricted to the outer surface, Pt
312

migrates into the membrane upon deposition cycling. This aspect of the IPMC electrodes is
more clearly seen in the EDS line scans provided for 1, 2 and 3 cycles in Fig. A4.2d. As the
and the concentration of Pt within the membrane correspondingly increases. In contrast,
electrode metals are found [26] to penetrate to only 10-20 μm in Nafion ® . The Ag surface
electrodes measuring 30-40 μm thick reduce the surface resistance to 2 �, which ensures that
the potential along the entire length of the actuator is relatively uniform during testing. Recall
that the electrodes are sequentially applied to both surfaces of water-swollen IPMCs. If
desired, the water can be replaced by other polar solvents, such as GLY and EG, that are
miscible in water. Regardless of the nature of the solvent, the free acid form of the ionomer is
exchanged with Li ions, which promote substantially increased solvation. Application of an
electric potential across the thickness of the IPMC results in bending actuation, as illustrated
by the schematic diagram presented in Fig. A4.1b and confirmed by the real-time sequence
of superimposed images (discussed later) shown in Figs. A4.3a and A4.3b. While several
mechanisms have been proposed [27] to explain such actuation, one of the most widely
accepted models is based on the microelectromechanical (MEM) model proposed by Asaka
and Oguro. [28] Recent neutron imaging and fluorescence microscopy studies performed by
Park et al. [8] and Park et al., [7] respectively, provide direct experimental evidence of this
mechanism in IPMCs composed of Nafion ® .
The MEM mechanism requires a potential-driven spatial redistribution of counterions and
solvent molecules within the nanostructured ionomer membrane of the IPMC. In the
313

cantilever test configuration, the electric potential causes solvated cations to migrate towards
the cathode, thereby increasing the electro-osmotic pressure of the ionic clusters along that
side of the specimen. As solvent enters the clusters, the internal strain in the polymer matrix
swells the interfacial boundary layer separating the cathode and membrane, and a
solvent/counterion gradient forms across the membrane. Due to depletion of solvent
molecules and counterions at the anode, the boundary layer separating the anode and
membrane contracts. Concurrent swelling and contraction along opposite boundary layers
and about the neutral plane results in the IPMC bending towards the anode. The rate at
which, as well as the extent of which, bending occurs depends sensitively on several material
characteristics, including the nanostructure and composition of the membrane, the quality of
the solvent, and the nature of the neutralizing cation. The two most commonly employed
polyelectrolytes used in IPMC actuators are Nafion ® (and chemically-related analogs) and
endblock-sulfonated triblock copolymers. Unlike Nafion ® , [29] the PBI introduced here affords
(i) facile processing in solvents without sacrificing useful mechanical or conductive
properties, (ii) considerably higher solvent absorption levels, (iii) tunable morphologies in
terms of nanostructural dimensions and anisotropy, and (iv) precise control over electrode
morphology. Solvent processing can be used to tailor membrane dimensions, whereas solvent
uptake determines the ionic conductivity of IPMCs and their actuation efficacy. Since most
prior studies of IPMCs rely exclusively on the use of Nafion ® , control over membrane
morphology provides a largely unexplored route by which to alter ion and solvent transport
and, thus, the rate and extent of actuation. While endblock-sulfonated triblock copolymers
314

have, in fact, been used [19] to a limited extent as IPMC membranes, introduction of a polar
solvent plasticizes the glassy endblocks and compromises the mechanical properties of such
IPMCs at the solvent concentrations needed for satisfactory conductivity. This trade-off is
not encountered in the current PBI membrane design.
The electromechanical performance of IPMC actuators is quantified by measuring
cantilever deflection relative to zero load. In this work, measurements have been conducted
at electric potentials ranging from 0.5 to 9 V, depending on the solvent employed. Although
most of the results presented here correspond to organic (GLY and EG) solvents, water is
used initially to permit comparison with other IPMCs, as shown in Fig. A4.3c. Bending
deflections generated here with a dc potential of 1 V are marginally greater than those
achieved with Nafion ® actuated with an ac potential of 1 V at 0.5 Hz. Note, however, that the
PBI investigated here yields actuation levels that are more than 3x greater than an IPMC
derived from an endblock-sulfonated triblock copolymer subjected to a dc potential of 1 V,
and more than 6x greater than an IPMC fabricated from Nafion ® (exchanged with Na ions)
under a dc potential of 2 V. [30] While the PBI membrane requires more time (~30 s) to
achieve full actuation and the Nafion ® -based IPMC actuates to its maximum in ~5 s, the
latter also shows considerably greater evidence of back relaxation, in which the IPMC
subjected to continued electrical bias changes its actuation direction (i.e., towards the
original, zero-load position) after reaching a maximum forward deflection. Back relaxation
has been previously reported [26] to depend on the nature of the ionic moiety present on the
polyelectrolyte backbone and molecular-level relaxation processes associated with the
315

polymer chains comprising the membrane. Actuation tests using water as the conductive
solvent are limited by the relatively low normal boiling point and electrolysis potential of
water. To avoid these limitations and reduce the extensive likelihood of oxidizing the
electrode metals, we have elected to use GLY and EG in the remainder of this study. Their
normal boiling points are 290 and 197°C, respectively, whereas their corresponding
electrolysis potentials are ~10 V for GLY and ~4 V for EG. One drawback to using organic
liquids instead of water is that their viscosities are significantly higher (934 cP for GLY and
16 cP for EG, relative to 1 cP for water, at 25°C), which is expected to result in markedly
slower actuation rates.
A sequence of digital images acquired during actuation of a GLY-containing IPMC
fabricated from the PBI (with 3 Pt electrode coatings) and exposed to an electric potential of
7 V is displayed in Fig. A4.3a. This sequence demonstrates that the active length (L) of the
IPMC (not gripped at the clamps) undergoes significant bending deflection upon actuation,
the extent of which is measured in terms of the dimensionless parameter �L, where � is the
curvature (equal to the reciprocal of the radius of curvature) of the IPMC cantilever strip.
Values of �L extracted from digital images such as those in Fig. 3a confirm that bending
proceeds to a plateau level and that the magnitude of bending increases with (i) the number
of Pt coatings applied and (ii) the applied potential up to 7 V. At higher potentials (9 V, not
shown here), the actuation level decreases due presumably to observed loss of solvent
through the nanoporous Ag electrode during actuation. Included for comparison in this figure
are actuation kinetics of IPMCs produced with Nafion ® exchanged with different ionic
316

species and actuated in GLY at 2 V. As in Fig. A4.3c, Nafion ® -based IPMCs actuate more
quickly than those containing the PBI membrane, but the Nafion ® systems also exhibit
considerable back relaxation while the potential remains applied. Back relaxation, which
hinders control of actuator motion and peak force, is not evident in the PBI-containing
IPMCs over the course of the actuation measurement, but relaxation back to zero deflection
ensues upon removal of the potential. It is of further interest that the IPMCs fabricated from
Nafion ® with K ions bend, for the most part, towards the cathode following a small initial
deflection toward the anode. Similar behavior, however with no initial bending towards the
anode, is observed with the PBI membrane (and Li ions) swollen with EG, as evidenced by
the images shown in Fig. A4.3b and the extracted data presented in Fig. A4.3e.
The sequence of images provided in Fig. A4.3b verifies that the IPMC undergoes
substantial bending toward the cathode upon electrical actuation. As far as we are aware, this
is the first time that the direction of bending actuation is solvent-regulated, holding all other
parameters constant. Although the precise reason for this unexpected behavior is not known
at this time, we hypothesize that it is a consequence of the chemical interaction between EG
molecules and the functional groups on the PBI chains or the mobile Li ions, and is not due
to viscosity differences. As with GLY in Fig. A4.3d, the extent of actuation increases to a
plateau level as the electrical potential is increased to 1.0-1.5 V, beyond which a reduction in
bending is observed most likely for the same reasons offered to explain the maximum voltage
encountered with GLY. When the voltage is removed, no back relaxation occurs, further
confirming that solvent quality plays a critical role in the fabrication of PBI-containing
317

IPMCs with targeted actuation properties. To compare the transport properties of the data
presented in Figs. A4.3d and A4.3e, we apply the MEM framework [28] to account for bending
as a function of time (t), but recognize that other models are likewise available for this
purpose. According to this theory (originally developed for water-swollen IPMCs), the
curvature of an initially flat membrane can be conveniently written as
� � �
1� exp ��
Dm 2 Dmt /W 2 � ( ��
(1)
where � for a given specimen under actuation is a constant that depends on several
parameters, �� including the thickness (W), equilibrium solvation and modulus of the
membrane, as well as the diffusivity (Dm), density and transfer coefficient of the solvent, and
the applied current density.
Regression of Eq. 1 to the data yields the solid lines included in Figs. A4.3d and A4.3e and
the values of Dm provided as a function of applied voltage in Fig. A4.3f. For comparison, the
value of Dm for water in a Nafion ® -based IPMC at 25°C is ~4.5 x 10 -6 cm 2 /s, [28] whereas
those for GLY and EG in the PBI tend to be lower, ranging from about 3 x 10 -6 cm 2 /s for EG
to 1 x 10 -7 cm 2 /s for GLY.
Maximum values of �L measured from this study are compared with those from previous
efforts in Fig. A4.4 and immediately reveal that IPMCs containing the PBI membrane
swollen with GLY are capable of substantially larger actuation levels at higher voltages. This
figure also provides insight into the dependence of actuation on the number of Pt deposition
cycles. We hasten to point out that if the number of Pt cycles is increased to 4, the actuation
properties deteriorate as a percolated network of Pt develops and permits uninterrupted
318

current flow. In addition, the results displayed in this figure confirm that the direction of
actuation can be controlled by judicious selection of the solvent employed. Thus,
incorporation of the PBI membrane into an IPMC permits tunable bending actuation and
direction at high solvent concentrations without sacrificing the mechanical integrity of the
IPMC. The novel multicomponent design afforded by the PBI membrane can provide
unprecedented access to next-generation IPMCs with tailorable morphologies and, hence,
ion/solvent transport properties and (electro)mechanical properties for motion-control
devices, designer sensors and bio-inspired applications such as artificial cilia [31] and
wormlike robots. [32]
Experimental
The PBI used here was generously supplied by Kraton Polymers (Houston, TX). The
sizes of the blocks (cf. Fig. A4.1a) in the non-sulfonated parent copolymer were
approximately 15 (tbS), 10 (EP) and 28 (S) kDa. As reported by the manufacturer, the
midblock was about 57 mol% sulfonated, and the ion exchange capacity (IEC) of the PBI
was 2.0 meq/g polymer. Reagent-grade tetrahydrofuran (THF), GLY, EG, and lithium
chloride were purchased from Fisher Scientific and used as-received. Tetraamine
platinum(II) chloride hydrate was obtained from Aldrich. The PBI was dissolved in THF to
form 2 wt% solutions, which were subsequently cast into Teflon molds to form films that
were annealed at 60°C for 24 h under vacuum. These films, measuring ~400 μm thick, were
swollen with GLY or EG at 60°C overnight, and the resulting materials were analyzed on a
Rheometrics Mechanical Spectrometer (RMS 800) equipped with 8 mm parallel plates
319

separated by a 1-2 mm gap. The linear viscoelastic (LVE) limits at ambient temperature were
determined by performing dynamic strain sweeps from 0.5 to 10% strain amplitude at a
frequency of 1 rad/s. Frequency spectra were acquired at a strain amplitude of 1% in the LVE
region. Fabrication of IPMCs required the formation of electrodes on the top and bottom
surfaces of solvent-cast films. This involved two steps: initial compositing and surface
electroding. In the first step, the top and bottom Pt electrodes were formed on a water-
swollen PBI film by electroless deposition of Pt, as described elsewhere. [25] This procedure
was repeated up to 4 times. Surface electroding resulted in the formation of Ag electrodes on
top of the Pt electrodes by the silver mirror reaction. [23] As the reaction proceeded, discrete
Ag particles deposited on the film surfaces and formed a contiguous layer. The dried
membrane/electrode assembly thus obtained was examined by scanning electron microscopy
(SEM) conducted in a Hitachi S3200 variable-pressure instrument operated at 20 kV in the
presence of He at a pressure of 100 Pa. Samples were frozen in liquid nitrogen and cross-
fractured to differentiate the layered morphology without the use of a conductive coating. An
Oxford Isis EDS system was used to measure the spatial distribution of the electrode metals.
For electromechanical characterization, fabricated IPMC films with electrodes were swollen
in water, followed by immersion in a 0.5 M lithium chloride solution for 24 h at 60°C to
promote Li ion exchange. Each film was subsequently submersed in a predetermined organic
solvent at 60°C for 12 h to produce a mesogel (which retains morphological characteristics of
the PBI prior to introduction of solvent) and then cut into pieces measuring 3 mm x 25 mm.
One end of the cut strip was clamped, and a free (active) length (L) measuring 1.9 to 2.0 cm
320

long was continuously monitored as a predetermined electric potential was applied through
the clamps. Recorded digital footage was analyzed by the Matrox Inspector software to
determine the level of bending actuation from the radius of curvature along L.
Acknowledgments
This study was funded by the National Science Foundation, and K. E. R. gratefully
acknowledges support from a National Science Foundation Graduate Fellowship. We thank
Professor S. A. Khan for use of his laboratory facilities.
321

Figures
Figure A4.1. (a) Chemical structure of the PBI block ionomer possessing p-t-butyl styrene
[tbS] endblocks separated from the styrene-co-styrenesulfonate [S(sS)] midblock by flexible
ethylene-alt-propylene [EP] linkages. (b) Schematic illustration of the idealized microphaseseparated
morphology of a midblock-swollen PBI film — the copolymer shading matches
that in (a) — and the resulting actuation mechanism of IPMCs prepared therefrom. (c)
Frequency spectra of the PBI selectively swollen with GLY (open symbols) and EG (filled
symbols). The dynamic storage (G', circles) and loss (G", triangles) moduli are labeled.
322

Figure A4.2. (a) SEM cross-sectional image and corresponding elemental x-ray maps of (b)
Pt and (c) Ag after 3 Pt deposition cycles. Line scans collected from cross-sectional x-ray
maps for Pt (light line) and Ag (dark line) are included in (d) for IPMCs subjected to 1, 2 and
3 Pt cycles.
323

Figure A4.3. Digital image sequences of PBI-based IPMCs acquired during electroactuation
with (a) GLY and (b) EG at 7 and 1 V, respectively. (c) Actuation-induced bending
deflection of Li-exchanged, hydrated PBI ( ) at 1 V. For comparison, deflection results
reported for hydrated IPMCs composed of an endblock-sulfonated triblock copolymer ( ) at
1 V [19] and Li-exchanged Nafion® (dashed line) at 1 V and 0.5 Hz[19] and Na-exchanged
Nafion® ( ) at 2 V[30] are included. (d) Bending actuation as a function of time for IPMCs
consisting of GLY-solvated PBI (labeled with the cation and applied voltage) and Nafion®
(labeled with the cation in parentheses) at 2 V.[26] (e) Bending actuation as a function of
time for IPMCs consisting of EG-solvated PBI (labeled with the cation and applied voltage),
as well as Nafion® ( , ) and Flemion® ( ) (each labeled with the cation and applied voltage
in parentheses).[26] In both (d) and (e), the effect of Pt coating on bending actuation is
shown for PBI-based IPMCs ( , labeled with the number of Pt cycles) at 7 and 1 V,
respectively. The dashed lines in (d) and (e) serve to connect literature data, whereas the
solid lines correspond to regressions of Eq. 1 to data collected here. The bending direction of
the IPMC is identified by the background shading: anode (white) or cathode (gray). (f)
Values of Dm extracted from the regression analyses in (d) and (e) and presented as a
function of electric potential for PBI-based IPMCs solvated with GLY (circles) and EG
(triangles) and subjected to 2 (filled) and 3 (open) Pt deposition cycles. The solid and dashed
lines in (f) serve to connect the data.
324

Figure A4.4. Maximum bending actuation achieved for IPMCs fabricated from the PBI
investigated in this work. Shown here are the IPMCs with GLY (filled symbols) and EG
(open symbols), each labeled with the number of Pt deposition cycles. Error bars correspond
to the standard error. The cross-hatched region identifies the range of actuation levels
achieved for conventional IPMCs under similar test conditions. Background shading is the
same as in Fig. 3, and the solid lines serve to connect the data.
325

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328

Abstract
APPENDIX V
Block Copolymer Self-Organization vs. Interfacial Modificationin
Bilayered Thin-Film Laminates*
Arif O. Gozen, Jiajia Zhou, Kristen E. Roskov, An-Chang Shi,
Jan Genzer and Richard J. Spontak
Block copolymers remain one of the most extensively studied and utilized classes of
macromolecules due to their extraordinary abilities to (i) self-assemble spontaneously into a
wide variety of soft nanostructures and (ii) reduce the interfacial tension between, and thus
compatibilize, immiscible polymer pairs. In bilayered thin-film laminates of immiscible
homopolymers, block copolymers are similarly envisaged to stabilize such laminates.
Contrary to intuition, we demonstrate that highly asymmetric block copolymers can
conversely destabilize a laminate, as discerned from macroscopic dewetting behavior, due to
dynamic competition between copolymer self-organization outside and enrichment at the
bilayer interface. The mechanism of this counterintuitive destabilization is interrogated
through complementary analysis of laminates containing mixtures of stabilizing/destabilizing
*This chapter has been published in its entirety
AO Gozen, J Zhou, KE Roskov, AC Shi, J Genzer, RJ Spontak. Gozen, Arif O. "Block copolymer
self-organization vs. interfacial modification in bilayered thin-film laminates." Soft matter 2011, 7,
3268.
329

diblock copolymers and time-dependent Ginzburg-Landau computer simulations. This
combination of experiments and simulations reveal a systematic progression of
supramolecular-level events that establish the relative importance of molecular aggregation
and lateral interfacial structuring in a highly nonequilibrium environment.
Introduction
One of the most important technological challenges in the development of inexpensive
polymeric materials with tailored properties is compatibilization of two or more immiscible
homopolymers. 1 The demand for such polymeric systems continues to grow as the need for
lightweight, processable and mechanically robust materials increases in response to efforts
aimed at conserving natural resources and reducing energy consumption. 2 Most polymer
pairs are inherently immiscible due to unfavorable thermodynamics: 3 the enthalpy of mixing
is typically endothermic and the entropy of mixing is often negligibly small.
Compatibilization requires a reduction in interfacial tension, which, in turn, is achieved
through a variety of interfacial modification strategies. 4 One effective route is reactive
blending, 5,6 which relies on the chemical coupling of dissimilar macromolecules at the
polymer/polymer interface. The result is the in-situ development of block (or graft)
copolymer molecules, which, when localized at the interface, serve to lower the interfacial
tension (thereby reducing the size scale of phase separation), improve fracture toughness 7
and restrict droplet coalescence. 8 While this approach is broadly applicable to a wide range
of polymers and ensures that copolymer molecules reside at the interface where they are
330

needed, it can suffer from undesirable side reactions, as well as low yields since the reactive
macromolecules must meet at the interface for the reaction to proceed.
An alternative to reactive blending relies on the physical addition of a premade block
copolymer to an incompatible polymer blend. 9 In this case, the copolymer molecules must be
allowed to diffuse to the developing polymer/polymer interface. While this can be largely
accomplished through the use of plasticizing agents or high shear fields, a fraction of the
copolymer population remains inevitably in one or both of the homopolymer phases. In this
case, the copolymer molecules in the bulk seek to reduce unfavorable contacts with the
surrounding matrix by self-assembling into nanostructured domains that protect their
incompatible elements. Alone, diblock copolymers can organize spontaneously into periodic
nanostructures ranging from spheres or cylinders of the minority component (on a body-
/face-centered cubic or hexagonal lattice, respectively) in a matrix of the majority component
to bicontinuous channels or lamellae. 10 In the presence of a solvent or homopolymer,
aperiodic bicontinuous 11,12 and network 13 morphologies may also develop. Because of the
propensity for block copolymers to self-organize, melt blending of incompatible
homopolymers must be conducted in such fashion to keep the population of copolymer
molecules remaining in one or both homopolymers relatively low; otherwise, trapped
copolymer molecules form aggregates that resemble micelles, which prevent the interface
from attaining its maximum strength and hinder compatibilization.
Similar phenomena likewise occur during the copolymer-induced stabilization of
bilayered thin-film laminates composed of two molecularly thin homopolymer films. Such
331

laminates are important to the development of advanced protective coatings, 14 solar cells 15
and waveguide assemblies, 16 in which case a detailed understanding of the molecular-level
processes governing stabilization is required. In addition, the planar arrangement of such
laminates provides a convenient test platform for the exploration of material and/or design
variations in systematic fashion 17 while avoiding changes in interfacial curvature that would
occur in bulk systems due to compatibilization. Without an added copolymer, a laminate may
rupture either in a single layer or in both layers, 18,19 forming circular holes or other complex
dewetting patterns, 20 when heated above their glass transition temperatures. Here, we only
consider the cases where (i) the melt viscosity of the substrate layer is much larger than that
of the top layer so that the substrate may be considered solid-like relative to the top layer; 21
and (ii) destabilization proceeds by the nucleation and growth of rimmed holes that
eventually impinge. 18 In the absence of interfacial slip, 22 the hole diameter (D) varies linearly
with time (t), and the hole growth (dewetting) rate (dD/dt) depends on the ratio of the
dewetting force to the friction caused by viscous dissipation. The magnitude of dD/dt affords
a relative measure of interfacial stability and can, along with the mechanism of dewetting, be
controlled by varying material parameters such as the thickness 23 or molecular weight 24 of
the top layer (cf. Fig. A5.1), as well as adding a species that modifies the nature of the
interface. 24,25
We prepared thin-film laminates from two polystyrene (PS) homopolymers with number-
average molecular weight ( M n) values of 30 and 50 kDa (PS30 and PS50, respectively), in
conjunction with a poly(methyl methacrylate) homopolymer with
��
332
��
M n = 243 kDa

(PMMA243), all from Pressure Chemical, Inc. (Pittsburgh, PA). To the PS homopolymers,
we added poly(styrene-b-methyl methacrylate) diblock copolymers (Polymer Source, Inc.,
Dorval, Quebec, Canada) varying in molecular symmetry: S10M50, S50M54 and S50M10,
where each numerical designation denotes the block molecular weight (in kDa). These
materials, as well as solvent-grade toluene (Sigma-Aldrich, St. Louis, MO), were used as-
received. Each PMMA243 substrate measuring 50±2 nm thick from ellipsometry was spun-
cast at a speed of 2000 rpm onto a silicon wafer from a 1.35 wt% solution in toluene.
Similarly, 1.55 wt% solutions of PS30 and PS50 in toluene were spun-cast at the same speed
with and without added copolymers onto glass. Each film measured 60±2 nm thick and was
floated off on deionized water and then deposited on a PMMA243 substrate to form a
bilayered laminate. All laminates were dried for 24 h at ambient temperature and
subsequently annealed at 180°C under nitrogen. Dewetting kinetics were monitored in
reflection mode with an Olympus BX60 optical microscope equipped with a Mettler heating
stage and a computer-interfaced CCD camera. Transmission electron microscopy (TEM;
Hitachi HF2000, 200 kV) was performed on laminates prepared on silica-coated grids, and
atomic force microscopy (AFM; Park Systems XE-100) was conducted in non-contact mode
on specimens before and after selective removal of the top PS30 layer in 1-chloropentane. 26
Dewetting rates are presented as a function of copolymer concentration for laminates
composed of PS30 (Fig. A5.1A) and PS50 (Fig. A5.1B) top layers on PMMA243. In the
absence of copolymer, measured values of dD/dt are 310±17 and 145±4 μm/h, respectively,
confirming that a larger melt viscosity due to increased molecular weight of the top layer
333

educes dD/dt and thus improves stability. 24 Similarly, these homopolymer dewetting rates
can be used to distinguish copolymer-induced stabilization (or destabilization) by discerning
whether dD/dt is lower (or higher) than that of the unmodified PS/PMMA243 interface. The
dewetting rates with the addition of the asymmetric S50M10 and nearly symmetric S50M54
copolymers (Fig. 1A) verify that both copolymers tend to promote stabilization, with
S50M10 being more effective than S50M54. This comparison reveals that, with their long
styrenic block and short methacrylic block, S50M10 molecules are less likely to form
aggregates in PS30 and therefore diffuse to the interface where they physically intercalate the
two homopolymers and reduce interfacial tension. An increase in copolymer concentration
further improves the stability of PS30 due to a larger population of copolymer molecules
available for interfacial modification. At very low concentrations, however, the S50M54
copolymer is found to induce destabilization due to the presence of the longer, more PS-
incompatible methacrylic block. In stark contrast to the general behavior of these
copolymers, incorporation of the S10M50 copolymer fully destabilizes the PS30 layer over
the entire copolymer concentration range examined (with no discernible concentration
dependence). The molecular-level mechanism responsible for this unexpected and
counterintuitive result is described and discussed below. Similar trends are evident in Fig. 1B
for laminateswith PS50. Note, however, that in this case the dewetting rates measured for
laminates with the S10M50 copolymer decrease with increasing concentration and approach
that of the neat PS50.
334

Comparison of the dewetting rates achieved by adding the S50M10 and S10M50 block
copolymers (with identical molecular weights) in Fig. A5.1 indicates that S50M10 brings
about stabilization, whereas S10M50 enhances destabilization. The mechanism by which the
S50M10 copolymer improves the compatibility of the immiscible interface has already been
discussed, but the behavior of the S10M50 copolymer is more complex, as illustrated in Figs.
A5.1C-E. We hypothesize that the incompatibility between either PS30 or PS50 and the
S10M50 molecules, each possessing a short styrenic block and a relatively long methacrylic
block, is sufficiently high to induce the spontaneous formation of aggregates with
methacrylic cores and styrenic shells, as evidenced by the inset of Fig. A5.1B. These
aggregates measure 30±6 nm in diameter and resemble crew-cut micelles 27 (depicted in Fig.
A5.1C). Because they are far from equilibrium, the copolymer molecules are likewise
capable of adopting more complex shapes (e.g., vesicles or toroids). As the system evolves,
aggregates and individual copolymer molecules migrate (at different rates), and ultimately
fuse, to the interface, as portrayed in Fig. A5.1D. Existence of partially fused, as well as
intact, aggregates along the interface causes increases in interfacial roughness and, hence,
area, which promote an increase in free energy and destabilization of the top layer. In this
case, the in-plane distribution of copolymer aggregates and molecules is not uniform.
Eventual dissolution of aggregates into brush patches (cf. Fig. A5.1E) is expectedly related to
the incompatibility between the styrenic matrix and the methacrylic block, which is greater in
the PS50 than in the PS30 laminates. This consideration explains why dD/dt is (within
335

experimental uncertainty) independent of copolymer concentration in Fig. A5.1A, but
decreases noticeably with increasing copolymer concentration in Fig. A5.1B.
Evidence supporting our proposed mechanism can be attained from time-dependent
Ginzburg-Landau computer simulations, which were performed with the assumption that the
PMMA243 substrate layer can be treated as a PMMA-attractive surface to simplify and
accelerate the calculations. A free energy functional proposed 28 previously for AB diblock
copolymer/homopolymer blends is incorporated into the Cahn-Hilliard equation to model
short-time system dynamics. The spatiotemporal behavior of two asymmetric copolymers
configured to emulate the S10M50 and S50M10 molecules, one with 17% A (A17B83) and
the other with 83% A (A83B17), is considered in a homopolymer A matrix in terms of (i) an
order parameter (�) that reflects the local copolymer concentration and (ii) local height
variations that provide a measure of roughness and, hence, structuring. The time-dependent
variation of � in the z direction, where z is normal to the A/B interface (Fig. A5.3A), shows
that the concentrations of both copolymer molecules along the interface (z = 0) increase as
the system evolves. A striking difference between the two species is that the A17B83
molecules extend from the interface as organized aggregates (evidenced by the fluctuations
in �), whereas the A83B17 molecules do not. A 2D image of a lateral simulation near the
interface for a laminate with A17B83 molecules is provided in the inset of Fig. A5.2A, and
reveals the existence of a complex copolymer morphology loosely reminiscent of the
nanostructure in Fig. 1B. In contrast, a corresponding image of the A83B17 molecules
displays significantly less lateral structuring. Root-measured (rms) roughness values
336

extracted from such 2D simulations aregiven in Fig. A5.2B and confirm that the A17B83
molecules are more organized, especially near the interface, than the A83B17 molecules,
which is consistent with our proposed mechanism.
Experimental AFM measurements of the interfacial roughness of the PS30/S10M50
laminate after selective removal of the PS30 layer are included for comparison in the inset of
Fig. A5.2B and indicate that, at short times, the roughness discerned from both dry (i.e.,
dewetted) and wet (i.e., not dewetted) regions on the PMMA243 surface increases as
dewetting proceeds, in agreement with simulation results (at different time scales). This
increase in roughness is attributed to the attachment and partial fusion of copolymer
aggregates along the interface. At longer times, the roughness decreases as copolymer
aggregates meld into the PMMA243 substrate. This process is observed and expected to be
faster (and more complete) for dry regions exposed to surface tension than for wet regions
subjected to lower interfacial tension. Existence of interfacial copolymer structuring due to at
least partially fused aggregates is verified by the TEM images presented in Figs. A5.3A and
A5.3B for laminates containing 0.15 and 0.75 wt% S10M50, respectively, after 6 min at
180°C. The dark features on the dry PMMA243 surface distinguish stained styrenic moieties
and serve to indicate copolymer-rich interfacial regions. At the low copolymer concentration
(Fig. A5.3A), discrete features possess diameters up to 35 nm, which is consistent with the
size of copolymer aggregates measuring ≈4Rg, where Rg denotes the copolymer gyration
radius (≈7 nm). At the higher concentration (Fig. A5.3B), these features are irregularly
337

shaped and possess a broad size distribution extending up to several hundred nanometers
across.
According to experimental observations and simulation results, self-assembly of the
S10M50 copolymer molecules occurs rapidly, resulting in the formation of micelle-like
aggregates that migrate to and roughen the polymer/polymer interface, consequently
destabilizing the top PS layer. In contrast, the mirrored S50M10 copolymer behaves in
largely opposite fashion: individual copolymer molecules diffuse to and meld with the
interface, where they stabilize the laminate. To discern the relative importance of these
competitive molecular-level mechanisms, we have prepared laminates containing mixtures of
these two copolymers and measured the dewetting rates, which are presented in Fig. A5.4. At
1 and 2 wt% S50M10, the destabilization mechanism dominates. Here, the population of
S50M10 molecules is insufficient to modify the polymer/polymer interface, whereas the
remaining (and more numerous) S10M50 chains favor self-assembly over interfacial
modification. Between 3 and 5 wt% S50M10, however, destabilization at low concentrations
precedes stabilization. Stabilization is achieved to different extents by having as little as 10
wt% S50M10 in the S10M50/S50M10 mixture. As seen in Fig. A5.4, using block copolymer
mixtures rather than single copolymers to tune stabilizing/ compatibilizing efficacy provides
an unexplored route to achieving property control from the ground up. Such control must
consider the complex interplay between block copolymer self-assembly and interfacial
modification under highly nonequilibrium conditions. Our results using a planar test
338

configuration elucidate a molecular-level mechanism responsible for this interplay, which is
of critical importance to the contemporary development of tailored polymeric materials.
Acknowledgments
Graduate fellowships for A.O.G. and K.E.R. were provided by Dade-Behring, Inc. and
the National Science Foundation, respectively. J.Z. and A.C.S. were supported by the Natural
Science and Engineering Council of Canada. The computer simulations were made possible
by the facilities of the Shared Hierarchical Academic Research Computing Network
(SHARCNET: www.sharcnet.ca) and Compute/Calcul Canada.
339

Figures
Figure A5.1. Dewetting rates (dD/dt) presented as a function of copolymer concentration for
laminates with (A) PS30 and (B) PS50 as the top PS layer (labeled in A). Diblock
copolymers defined in the text and added to the top layer are likewise color-coded and
labeled. Solid lines connect the data, and the dashed line delineates top-layer stabilization
from destabilization. Error bars denote one standard deviation in the data. The TEM image in
(B) shows the ill-defined and faint S10M50 nanostructure that develops upon spin-casting a
PS50/S10M50 film with 0.75 wt% S10M50 onto glass, followed by floated transfer onto a
PMMA243 substrate layer spun-cast on a silica-coated grid. In this image, styrenic units are
stained with the vapor of RuO4(aq) so that unstained methacrylic moieties appear light.
Examples of aggregates resembling micelles are circled, whereas more complex shapes are
identified by arrowheads. Included are schematic illustrations of the mechanism by which
asymmetric S10M50 block copolymers destabilize a bilayered laminate: (C) copolymer
molecules self-organize into aggregates (shown here), as well as more complex
nanostructural elements (cf. the inset in B) upon initial casting; (D) copolymer aggregates
and chains in the melt diffuse to the polymer/polymer interface where they adsorb and
eventually undergo fusion, promoting an increase in interfacial roughness; and (E) additional
aggregates form (depending on the available copolymer reservoir) and continue to migrate to
and meld with the interface to form copolymer brushes.
340

Figure A5.2. In (A), the concentration-based order parameter (�) presented as a function of
distance from the A/B polymer interface (at z = 0) at different simulation times (τ, labeled)
for A17B83 and A83B17 copolymer molecules (labeled and defined in the text) at a
copolymer concentration of 1.0 wt%. A pair of 2D lateral simulation images near the A/B
interface at τ = 40 is displayed for both copolymers (labeled) in the inset of (A). Values of the
rms roughness (in lattice units, l.u.) extracted from simulation images such as those provided
in (A) are provided for the A17B83 ( ) and A83B17 ( ) copolymers as a function of τ, in
(B). Included in the inset of (B) are experimental rms roughness values measured by AFM of
wet (not dewetted, ) and dry (dewetted, ) interfacial regions of a laminate after selective
removal of the PS30/S10M50 top layer. The solid lines connect the data.
341

Figure A5.3. TEM images acquired from dry regions of annealed laminates with
PS50/S10M50 top layers on PMMA243 at two S10M50 concentrations (in wt%): (A) 0.15
and (B) 0.75. Styrene-containing features remaining on the PMMA243 substrate after
dewetting appear electron-opaque (dark).
342

Figure A5.4. Dewetting rates presented as a function of total copolymer concentration for
PS50/PMMA243 bilayered laminates with and without mixtures of the asymmetric S50M10
and S10M50 block copolymers at different mixture compositions (color-coded, labeled and
expressed in w/w S10M50/S50M10). Solid lines connect the data, and the error bars denote
one standard deviation in the data.
343

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