Synthesis and dry-spinning fibers of sulfinyl-based poly(p-phenylene vinylene) (ppv) for semi-conductive textile applications

Article abstract of DOI:10.1039/c2jm30186e

Soluble, processable sulfinyl-based poly(p-phenylene vinylene) (PPV) precursors were synthesized, and the first sulfinyl-based PPV fibers were subsequently spun on a dry-spinning apparatus. 1-(Chloromethyl)-4-[(n- octylsulfinyl)methyl]benzene was synthesi

Full text of DOI:10.1039/c2jm30186e

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PAPER
Synthesis and dry-spinning fibers of sulfinyl-based poly(p-phenylene vinylene)
(ppv) for semi-conductive textile applications†
a
b
b
b
Emily B. Anderson, Denis Ingildeev, Frank Hermanutz, Alexandra Muller, Michael Schweizer
b
€
ab
*
and Michael R. Buchmeiser
Received 10th January 2012, Accepted 4th April 2012
DOI: 10.1039/c2jm30186e
Soluble, processable sulfinyl-based poly(p-phenylene vinylene) (PPV) precursors were synthesized, and
the first sulfinyl-based PPV fibers were subsequently spun on a dry-spinning apparatus. 1-
(Chloromethyl)-4-[(n-octylsulfinyl)methyl]benzene was synthesized in a 94% yield and polymerized via
an analogous approach to the Gilch reaction. Post-polymerization, the sulfinyl group may be thermally
1
eliminated from the backbone of the polymer to obtain conjugated PPV. H NMR, 13C NMR, and
FTIR confirmed that the polymer structure contained both sulfinyl- and pure PPV-based units.
Thermogravimetric analysis indicated that up to 60% elimination of sulfinyl groups occurred during the
polymerization reactions and before additional thermal treatment. Spinning dopes were prepared with
45 wt% polymer from 45.0 to 50.0 g of polymer in chloroform and had zero shear viscosities around 60
Pa s at 20 ꢀC. Fibers were dry-spun with and without tension and various jet draw ratios from 0 to 151%
to investigate changes in crystallinity, and X-ray diffraction patterns indicated enhanced orientation in
the fibers compared to the unprocessed polymers. The fluorescent, conjugated polymer fibers possessed
ꢀ
diameters less than 60 mm by SEM and remained soluble until thermal treatment at 150 C.
flexibility, strength, and reduced weight compared to other metal-
based conductive composites. Bendable, conductive clothing,
flexible keyboards and circuits, and electrical wiring from cheap,
light-weight, and robust organic materials could be realized.
Conductive textiles and clothing are used for antistatic appli-
cations,12 electromagnetic shielding (EMI) protection,12 bio-
monitoring and various sensing applications.13 For small and/or
rigid device applications, thinner fibers or wires may be required.
However, for clothing and semi-conductive fabrics, larger, flex-
ible, weavable fibers from industrial spinning processes are
desirable. While anti-electrostatic applications require a resis-
tivity of around 109 to 1013 U cmꢁ2 or shielding fabric applica-
tions require 102 U cmꢁ2, most polymers have a resistivity nearer
Introduction
Poly(p-phenylene vinylene) (PPV) and its derivatives are semi-
conductive, electroluminescent, conjugated polymers found in
organic light-emitting diodes (OLED)s,1–5 field-effect transistors
(FET)s,6,7 and photovoltaic cells.8–10 Conjugated polymers for
semiconductors have delocalized p-molecular orbitals traversing
the polymer backbone.1–5 Various structures include poly-
(acetylene), poly(p-phenylene) (PPV), poly(aniline), poly(pyrrole),
and poly(thiophene). PPV was one of the first conjugated poly-
mers used in electronics, and its conductivity is achieved with
either n or p doping.1,2 However, PPV-based semi-conductive
polymers are traditionally coated on inflexible, easily broken
substrates like glass.1 Many PPV derivatives are insoluble,
intractable, and infusible below their decomposition tempera-
tures, rendering processing difficult.1,2,11 Fabricating flexible
fibers from these materials, however, would permit use in various
new applications and decrease the fragility of devices.9 Effectively
spinning PPV-based materials into fibers would provide
to 1015 U cmꢁ2 12
Semi-conductive organic textiles are predomi-
.
nately valuable for their flexibility, stretching ability, and weight
compared to metal wires. For instance, a flexible, organic poly-
mer and semi-conductive textile could be used in current indus-
trial weaving processes while a metal fiber like copper is often
brittle. Conductive fibers have already been woven into skin-
tight clothing for long-term joint monitoring,13 because body
movement around joints like the knee requires stretching of the
skin up to 35 to 45% during normal motion.14 Conductive or
semi-conductive textile coatings or dyes can also be used,15 but
investigating the spinning of conjugated polymer fibers such as
PPV derivatives allows one to study the material properties and
the proof of concept for flexible, weavable semi-conductive
clothing from organic polymer materials.
aLehrstuhl fu€r Makromolekulare Stoffe und Faserchemie, Institut fu€r
€
Polymerchemie, Universitat Stuttgart, Pfaffenwaldring 55, D-70550
Stuttgart, Germany. Fax: +49 (0)711-685-64050; Tel: +49 (0)711-685-
64075
b
€
€
Institut fur Textilchemie und Chemiefasern (ITCF), Korschtalstr. 26, D-
73770 Denkendorf, Germany. E-mail: michael.buchmeiser@ipoc.
uni-stuttgart.de; Fax: +49 (0)711-9340-185; Tel: +49 (0)711-9340-101
† Electronic supplementary information (ESI) available. See DOI:
10.1039/c2jm30186e
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In 1990, Friend et al. used the first conjugated polymer in
a single-layer OLED through spin-coating a soluble PPV
precursor and subjecting it to thermal treatment which converted
the polymer to its conjugated form.1,4 Spin-coating soluble PPV
precursors with thermally labile functionalities became a viable
processing method to make organic-based large area displays,
replacing vapor deposition.1 Symmetrical (Gilch2,3,16–19 and
Wessling2,20–23) and unsymmetrical (sulfinyl) monomers11,24–33
were used to prepare soluble PPV precursors. All three routes
involve a, a0-functionalized monomers with a polarizing group
(P) and a leaving group (L). One equivalent or a slight excess of
base deprotonates an a-hydrogen on the methylene group next to
the polarizer. A p-quinodimethane derivative is formed from
the loss of the leaving group and rapidly polymerizes
(Scheme 1).11,24,31–33 The p-quinodimethane is not isolable, but
in situ UV-vis spectroscopy allows one to follow its formation.
Whether the polymerization proceeds through an anionic or
radical reaction is debatable.11,32,33 However, addition of radical
scavengers like 2,2,6,6-tetramethyl-1-piperidinyloxy (TEMPO)
does impede the production of high molecular weight polymer.33
Rehahn et al. showed that oxygen injected into Gilch polymeri-
zations effectively reduces the polymer molecular weight.17,18
Increasing the amount of oxygen injected progressively lowered
the polymer molecular weight; this result indicates that a radical
mechanism is predominant in this system.17,18 After termination
with dilute, aqueous hydrochloric acid, the precursor polymers
are often dialyzed or extracted to remove base and/or low
molecular weight fractions. Thermal elimination of the polar-
izing groups then creates a mostly insoluble structure.32,33
However, polymer precursors must be soluble and processable
before thermal treatment. Unsymmetrical alkyl, alkyloxy, or
phenyl groups afford solubility in organic solvents, but these
PPV derivatives often have different luminescence and quantum
efficiency as well as different glass transition temperatures and
crystallinity unless thermal elimination of these groups can result
in conversion to pure PPV structure with no substituents.2
Processability of PPV precursors still remains a problem as
many methods produce large defect levels, low molecular order,
and thus low power conversion efficiency. Molecular order and
self-assembly are pertinent to device efficiency. Spin-coating, and
Langmuir–Blodgett techniques, as well as chemical vapor
deposition have served as the main routes to obtain structural
order.1,2 Many comprehensive reviews exist on the synthesis and
processing of PPV-based materials.1–3,19,34 However, a few recent
attempts have been made to fabricate fibers from water and
organic soluble progenitor PPVs by electrospinning, and thermal
treatment of the fibers resulted in conjugated polymers.35–43 For
example, Huang et al. achieved conjugated fibers through co-
axial electrospinning. The fibers produced had an MEH-PPV
sheath and a PVP core with diameters of a few nanometers.40
Huang et al. also fabricated core–sheath nanofibers using
poly(styrene)/poly(p-phenylene vinylene) (PS/PPV), poly(vinyl
alcohol)/poly(p-phenylene vinylene) (PVA/PPV) and poly-
(styrene)/tris(8-quinolinolato) aluminium (PS/Alq3). However,
the PS/PPV-based materials had bead-on-string morphology
rather than forming uniform fibers, and the PVA/PPV formed
helical fibers.39 Huang et al. also synthesized PPV-based, Fe3O4-
containing composite nanofibers through electrospinning
precursor polymers.38 Thermal treatment led to the conjugated
structure which was oriented onto photoconductor devices.38
Some key features of electrospinning include its non-viability
for fast, large-scale production compared to dry-spinning as well
as its provision of small fiber diameters.44 Industrial dry-spin-
ning, for example, can reach production of nearly 1500 m minꢁ1
compared to electrospinning fiber production which is closer to
30 m minꢁ1 44
. One of our main goals is to realize the fabrication
of flexible, PPV-based fibers using traditional spinning tech-
niques rather than electrospinning to provide greater alignment
though adjustment of the drawing rate of the conjugated poly-
mers for enhanced orientation and conductivity. Oriented PPV is
mechanically robust and highly crystalline.45–49 Water-cast films
of PPV precursors have been drawn and heated to achieve
conjugated PPV with crystalline order according to X-ray
diffraction, and they contained crystallites around 5 nm in size,
which comprised around 50% of the bulk of each film.45
Alternatively, there are a few patents on fiber-based photo-
voltaics based on coating techniques to achieve flexible photo-
voltaics.50–52 In 2010, Sariciftci et al. used dip-coating to
manufacture photovoltaic fibers on a poly(propylene) (PP) base
and achieved a current density of 0.27 mA cmꢁ2.9 PP fibers with
a 0.59 mm diameter were purchased and dip-coated with
conductive layers in DMSO and chlorobenzene solutions. Sar-
iciftci’s fibers represent an early attempt at flexible, PPV- and
poly(thiophene)-based photovoltaic fibers. However, coating
technique, adhesion of the coating to the fiber, and device
performance all need improvement.
In this work, the first sulfinyl-based PPV fibers were fabri-
cated. Optimization of polymer sulfinyl precursor versus PPV
composition and rheological characteristics was accomplished
for a dry-spinning process. Fibers were spun at various jet stretch
ratios, and the unique orientations and crystallinities of these
fibers were analyzed thoroughly with X-ray diffraction and
Scheme 1 General polymerization mechanism for Gilch-, Wessling-, and sulfinyl-type monomers.11,24,31–33
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FTIR. Soluble sulfinyl-containing precursor polymers to
was then added as needed to facilitate precipitation and filtering
of salts from the product. The product was collected as a clear
oil. Then, 100 mL of octane or heptane was added and distilled
three times from the oil to remove tetrahydrothiophene. The
product was stored at 5 ꢀC. The yield of thioether was $98%. As
mentioned in previous literature, this crude liquid contains
a small amount of a,a-dichloro-p-xylene and other substituted
products that are removed in the subsequent step workup.25 1H
NMR (CDCl3, 200 MHz) d 0.89 (t, 3H), 1.27 (m, 8H), 1.55 (m,
4H), 2.40 (t, 2H), 3.69 (s, 2H), 4.58 (s, 2H), 7.25–7.39 (m, 4H)
ppm. 13C NMR (CDCl3, 100 MHz) d 14.15, 22.68, 28.90, 28.92,
29.21, 29.26, 31.47, 31.84, 35.95, 46.02, 136.09, 137.30, 137.71,
139.13 ppm. IR (ATR): n 2924, 2853, 2055, 1512, 1465, 1420,
1377, 1265, 1237, 1199, 1105, 1021, 885, 828, 770, 693, 560, 512,
463 cmꢁ1. GC-MS (EI, 70 eV): calcd m/z ¼ 284.14; found: 284.14.
poly(p-phenylene vinylene) (PPV) were prepared from the
corresponding
1-(chloromethyl)-4-[(octylsulfinyl)methyl]-2,5-
dimethylbenzene monomer. This sulfinyl-based monomer was
synthesized according to an adapted procedure by Vanderzande
et al.25 Polymerization of this progenitor was accomplished via
an analogous approach to Gilch and Wessling polymerizations
where the polarizer remained in the polymer structure to retain
solubility until thermal treatment. Introduction of this solubi-
lizing polarizer provided polymer processability for production
of fibers through a dry-spinning process at 45 wt% in chloroform.
Subsequent drawing and thermal elimination of the polarizer at
temperatures around 150–220 ꢀC provided the first sulfinyl-based
PPV fluorescent fibers with conjugated polymer structures.
Experimental
Synthesis of 1-(chloromethyl)-4-[(n-octylsulfinyl)methyl]benzene
Reagents and general methods
A 2000 mL, round-bottomed Schlenk flask was charged with
a Teflonꢀ stir bar, 37.0 g (0.13 mol) of 1-(chloromethyl)-4-
[(n-octylsulfanyl)methyl]benzene, and 830 mL of methanol. A
tellurium dioxide catalyst (3.12 g, 0.02 mol) was used to reduce
the oxidation time to 12 h. Then, 30 wt% hydrogen peroxide in
water (29.6 g, 0.26 mol) was added dropwise to the flask. The
reaction was allowed to proceed at ambient temperature for 12 h
before quenching with 400 mL of aqueous saturated NaCl
solution. The product was extracted with 500 mL ꢂ 3 of chlo-
roform, dried over magnesium sulfate, and filtered. Solvent was
removed under reduced pressure, and the yield of the resultant
white powder was 94%. The product contains a small amount of
All reagents were purchased from Aldrich or Merck and used
without further purification unless specified. Sulfinyl-based
monomers were synthesized according to adapted procedures by
Vanderzande et al.11,23,25 All monomer syntheses and polymeri-
zations were performed under nitrogen using standard Schlenk
techniques.
Synthesis of 1,4-bis(tetrahydrothiopheniomethyl)xylene
dichloride
A 1000 mL, round-bottomed Schlenk flask was charged with
a Teflonꢀ stir bar and 158 g (0.90 mol) of a,a-dichloro-p-xylene
followed by 315 mL of methanol. Tetrahydrothiophene (315 mL,
3.57 mol) was added, and the flask and contents were purged
with nitrogen for 1 h and kept under a nitrogen blanket for the
duration of the reaction. The reaction was allowed to proceed for
90 h at ambient temperature and subsequently precipitated in
1640 mL of cold acetone, collected, washed thoroughly with cold
acetone, and dried under reduced pressure at room temperature.
The yield was $98%, and the bissulfonium product was collected
asꢀa white, hygroscopic powder. The compound was stored at
5 C and dried under reduced pressure at ambient temperature
immediately before subsequent reactions to remove absorbed
water. 1H NMR (D2O, 200 MHz) d 2.06–2.31 (m, 8H), 3.24–3.52
(m, 8H), 4.45 (s, 4H), 7.50 (s, 4H) ppm. 13C NMR (D2O, 62
MHz) d 28.54, 42.91, 45.36, 130.73, 131.83 ppm. IR (ATR): n
3446, 3395, 3011, 2982, 2943, 1634, 1519, 1468, 1416, 1309, 1250,
1202, 1148, 1135, 1088, 1027, 954, 906, 880, 860, 801, 695, 596,
a,a-dichloro-p-xylene impurity,25 undifferentiable by H NMR
1
but evidenced by 13C NMR resonances around 130 ppm. If
necessary, recrystallization is completed in 1 : 1 hex-
ane : CH2Cl2.25 1H NMR (CDCl3, 200 MHz) d 0.87 (t, 3H), 1.16–
1.81 (m, 12H), 2.57 (t, 2H), 3.96 (m, 2H), 4.59 (s, 2H), 7.21–7.46
(m, 4H) ppm. 13C NMR (CDCl3, 62 MHz) d 14.48, 22.86, 22.99,
29.21, 29.37, 29.55, 32.09, 46.09, 51.48, 58.14, 129.54, 130.64,
130.78, 130.98, 137.99 ppm. IR (ATR): n 2959, 2917, 2848, 1513,
1467, 1445, 1421, 1267, 1212, 1138, 1105, 1076, 1022 (S]O),
1003, 931, 882, 849, 769, 721, 669, 604, 526, 455 cmꢁ1. GC-MS
(EI, 70 eV) calcd m/z ¼ 300.13; found: 300.13.
Polymerization of 1-(chloromethyl)-4-[(n-octylsulfinyl)methyl]
benzene
A 1000 mL, round-bottomed Schlenk flask fitted with an addi-
tion funnel was charged with a Teflonꢀ stir bar, 37.5 g of
1-(chloromethyl)-4-[(n-octylsulfinyl)methyl]benzene (0.13 mol),
and 510 mL of solvent, either NMP, DMSO, or sec-butanol. A
second 500 mL flask was used to prepare sodium tert-butoxide
(12.0 g, 0.13 mol) solutions in 90 mL of solvent which was added
to the addition funnel. The solutions were purged for 3 h under
nitrogen. The sodium tert-butoxide solution was added in one
portion, and the reaction was allowed to proceed for 1 h at 45 ꢀC
before pouring the solution into 500 mL of ice water and
quenching with 1.0 M HCl to neutral pH. The polymer was
washed with chloroform/water to remove any residual base, and
the organic phase was collected and reduced. The polymer was
precipitated into 900 mL of 1 : 1 hexane : diethyl ether and dried
531, 483 cmꢁ1
.
Synthesis of 1-(chloromethyl)-4-[(n-octylsulfanyl)methyl]benzene
A 2000 mL, round-bottomed Schlenk flask was charged with
a Teflonꢀ stir bar, 1,4-bis(tetrahydrothiopheniomethyl)xylene
dichloride (70.0 g, 0.20 mol), and 1300 mL of methanol. A second
1000 mL, round-bottomed flask was charged with 19.2 g of
sodium tert-butoxide (0.20 mol) and 500 mL of methanol. Then,
34.4 mL of n-octanethiol (0.20 mol) was added to the second
flask. After 1 h, the thiol and base solution were added to the first
flask. The reaction was allowed to proceed under nitrogen for
1 h. Methanol was removed under reduced pressure. Chloroform
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under reduced pressure at ambient temperature. The yield was
approximately 60%, partly due to elimination during the poly-
merization. Polymers prepared in polar aprotic solvents were
yellow-green solids due to significant elimination during the
polymerization that resulted in a large number of conjugated
units. Polymers from polymerizations in sec-butanol were white,
colorless solids. The products were stored at 5 ꢀC under nitrogen
transmittance threshold for all calculations was 0.05%, 0.0002
Abs. Mechanical textile properties were measured by means of
a Textechno tensile tester at 20 ꢀC using an average of the results
from 10 different fibers. The fineness was obtained by sonic
measurements. The gauge length was 25 mm. Rheological
measurements were implemented on a Rheometrics (SR 500)
rheometer equipped with parallel plate geometry and a Peltier
temperature control system. Shear rates were varied between
0.1 and 100 sꢁ1. The diameters of the plates were 25 mm, and the
gap between them was 1 mm. The viscoelastic properties of the
spinning dopes were studied by dynamic oscillatory experiments
at temperatures between 5 ꢀC and 30 ꢀC, which provide the
determination of the storage (G0) and the loss modulus (G00) as
well as of the complex viscosity. Using the principle of
frequency–temperature superposition, master curves were
obtained. The zero shear viscosity was calculated by the Carreau
model. Scanning electron micrographs (SEMs) were recorded
with a Zeiss Auriga with 1.74 or 2.40 kV. The samples were
sputtered with Au/Pd before analysis. Thermogravimetric anal-
yses (TGA) were conducted with a TGA-7 from Perkin-Elmer
operating with a 5 ꢀC minꢁ1 ramp under nitrogen from 30 ꢀC to
1030 ꢀC. Differential scanning calorimetry (DSC) was performed
on a DSC-7 from Perkin-Elmer operating with heat/cool/heat
cycles from ꢁ50 ꢀC to 80 ꢀC at 10 ꢀC minꢁ1. Thermal scans were
limited to this rangeꢀdue to elimination of sulfinyl groups that
occurred above 100 C. Midpoint Tg values are reported from
the second heating scans. Dynamic light scattering (DLS) was
conducted on a Zetasizer Nano-ZS by Malvern Instruments
GmbH with the model ZEN3600 and series number
MAL1009539. A glass cuvette with a round aperture was used to
analyze 1 mg mLꢁ1 polymer solutions at 25 ꢀC. The equilibration
time between runs was 2 min, and 3 runs were made with an
average of 15 measurements each and a 1 s delay between each
measurement. X-ray diffraction (XRD) measurements were
conducted on a Phillips PW 1050 vertical goniometer with the
controlling unit Phillips PW 1710 in the reflection mode with
a Cu–K source, an aluminium standard (38.5ꢀ 2q), and ADM
software. Pellet samples were made via pressing the polymers
under 125 bar of applied surface pressure and under 120 mbar of
applied vacuum, both concurrently for 5 min. The pellets were all
made with the same thickness of 2 mm.
1
and in the absence of light until further use. H NMR (CD2Cl2,
200 MHz) d 0.86 (-CH3), 1.07–4.02 (-CH2-), 6.71–7.75 (-CH]
CH-, phenyl-H) ppm. 13C NMR (CDCl3, 100 MHz) d 14.07
(-CH3), 22.86, 22.60, 28.61, 28.84, 28.99, 29.16, 31.71, 38.91,
51.24, 57.80, 126.02–127.25, 127.33–130.76, 136.72–138.60 ppm.
IR (ATR): n 3026, 2955, 2922, 2854, 1610, 1513, 1465, 1421,
1377, 1304, 1262, 1103, 1037, 1021 (S]O), 964 (C]C),ꢀ857, 827,
723, 6ꢀ74, 638, 555, 529 cmꢁ1. DSC (H/C/H ꢁ50 to 80 C) Tg 39
to 42 C.
Dry-spinning
For the spinning dopes, 45 wt% polymer solutions were prepared
by first dissolving 45.0–50.0 g of polymer in chloroform at 10 wt%
at ambient temperature. Once the polymer was fully dissolved, the
solution was rotaevaporated slowly at ambient temperature at 100
mbar until a viscous solution at 45 wt% solids was obtained. The
polymer solution was extruded (2 or 5 cm3 minꢁ1) employing
a Fourne lab scale piston spinning device. The fibers were spun by
utilizing a spinneret with 32 holes and a hole diameter of 0.1 mm
(aspect ratio (L/D) 0.8) or a hole diameter of 0.08 mm (L/D 0.16).
The filaments were spun at 30 ꢀC, then stretched on air (gap) with
different jet draw ratios, dried, wound on an electronically
controlled godet, and finally processed for analysis. Fibers were
collected (1) without tension, (2) with tension and a jet draw of 0%,
and (3) with tension and a jet draw ratio of 151%. Fibers collected
without tension were collected on a godet at 28 m minꢁ1, which
was less than the extrusion speed of 31 m minꢁ1 and related to a jet
draw ratio of ꢁ10% due to filament shrinkage while drying. Fibers
collected with tension were wound onto a godet with a collection
speed of 31 m minꢁ1, which was the same as the extrusion speed.
Drawn fibers were also wound at 78 m minꢁ1, which was related to
a 151% jet draw ratio. Fibers were collected on bobbins and stored
at 5 ꢀC and in the absence of light before analysis.
Characterization
Results and discussion
A 200 MHz Bruker Avance and a 400 MHz Bruker Avance III
NMR spectrometer were both used for 1H and 13C NMR
experiments, operating at 50 and 100 MHz for 13C NMR
respectively. A Hewlett Packard 5890 series II GC-MS was
employed, equipped with a MDN-35 column with 30 ꢂ 0.25 mm
dimensions and 0.25 mm film thickness. Mass spectroscopy
Scientific rationale
To our knowledge, this is the first report of dry-spun fibers from
any sulfinyl-based PPV precursors. Previous reports exist of
electrospinning various other types of PPV precursors (not sul-
finyl) from dilute solutions on a small scale,35–43 but electro-
spinning is not as commercially viable for fast, large-scale
production as traditional wet- or dry-spinning processes.44
Although electrospinning allows for slight adjustment of the jet
drawing through the applied voltage and more recently the
advent of the collection of electrospun fibers on rotating drums,44
dry-spinning offers a wide range of fiber stretching capabilities
since the subsequent fibers are easily wound onto godets at
various jet drawing rates without complicated set-ups or without
issues related to residual charge. This subsequent stretching of
ꢀ
measurements were taken starting with a 45 C isotherm for 1
minute, a ramp to 300 C at 12 C minꢁ1, followed by a 300 C
isotherm for 5 min. For photographs under UV light, a UV
Desaga (Heidelberg) HP-UVIS high pressure Hg-lamp at 366 nm
(serial number 80266) was operated at 240 V and 140 W. FTIR
spectra were collected on a Nicolet 380L from Thermo Fisher
Scientific with a diamond ATR and/or a Perkin Elmer 2000
FTIR using KBr pellets. The range was 4000 to 370 cmꢁ1 for all
ꢀ
ꢀ
ꢀ
measurements with an average of
8 scans. The percent
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Scheme 2 Synthesis of sulfinyl-based poly(p-phenylene vinylene).
the fibers also has an effect on the fiber crystallinity, orientation,
and strength.
intermolecular interactions from the conjugated segments and
much greater crystallization capability, although having a high
molecular weight could also be a factor. These interactions or
entanglements provide a highly viscous solution of interwoven
polymer chains that allows the fiber to remain intact after
formation. The viscous, honey-like flow of the spinning dope
temporarily lowers in viscosity during spinning, which occurs
with applied pressure to the spinning dope and also shearing
forces located at the spinneret (Fig. 2).
Synthesis and solution behavior of sulfinyl-based PPVs
In order to synthesize sulfinyl-based PPV, first monomer
synthesis was achieved via adapted procedures from Vander-
zande et al. (Scheme 2).25
Polymerization was conducted in an inert atmosphere to
produce a high molecular weight polymer, suitable for fiber
spinning. Some elimination of polarizing groups to conjugated
poly(p-phenylene vinylene) occurred during polymerization in
polar aprotic solvents, evident from the polymer’s yellow color,
1H NMR spectra (Fig. S1†), and 13C NMR spectra. sec-Butanol
was alternatively used as the solvent for polymerization to
produce a white, non-conjugated polymer, which formed clear
polymer films when cast from chloroform (Fig. 1). However, the
polymer synthesized in sec-butanol did not form homogeneous,
viscous solutions in chloroform at high weight percent solids,
possibly due to having fewer intermolecular interactions or lower
molecular weight. The polymer formed in polar aprotic solvents
like NMP contained some amount of elimination, but this
polymer proved to have optimal rheology for fiber spinning
(Fig. 2). This result is possibly due to favorable p–p stacking
Thermal analysis
Thermogravimetric analyses gave a rough estimate of the
amount of sulfinyl groups on the backbone of sulfinyl-based
PPVꢀs. Thermal eliminatiꢀon of sulfinyl groups occurred between
ꢀ
110 C (onset) and 240 C for all polymers (TGA Ramp 5 C
minꢁ1, N2) (Fig. 3). The polymer synthesized in polar aprotic
solvents possessed elimination of approximately 60 wt% of
sulfinyl groups as estimated via TGA under nitrogen. Thermo-
grams indicated a 40% weight loss of polymer during the first step
from 113 to 240 ꢀC. The polymer possessed 14 wt% char at
1028 ꢀC. The polymer synthesized in sec-butanol underwent
much less elimination during polymerization, approximately
29 wt% of sulfinyl groups had been eliminated before thermal
analysis based on the 71 wt% loss of polymer after the first step in
the thermogram. This polymer also displayed a much lower char
ꢀ
yield, 8 wt% at 1028 C.
DSC indicated that the polymer with 29 wt% elimination
possessed a Tg of 42 ꢀC, and the polymer with 60 wt% elimination
possessed a glass transition of 39 ꢀC. Thermal scans were
completed from ꢁ50 to 80 ꢀC to avoid elimination of sulfinyl
groups. A melting point was not visible for these polymers, and
from 110–250 ꢀC elimination occurred and overlapped any
possible melting transition. Once converted into PPV, the poly-
mer appeared to degrade before it melted, making a clear
determination of a melting point difficult.
Dynamic light scattering (DLS) and viscosity
Attempts to analyze the molecular weights of these polymers by
size-exclusion chromatography (SEC) in chloroform resulted in
back-pressure problems even after filtering samples due to strong
intermolecular interactions between the polymer chains. Even
though viscous, visually homogeneous spinning dopes could be
derived in chloroform at high weight percent solids, solutions still
contained well-dispersed aggregates. Other types of substituted
Fig. 1 Films cast from CHCl3 of sulfinyl-based PPV with varied levels of
elimination and prepared from polymer derived from polymerization
reactions in (a) polar aprotic solvents and (b) sec-butanol.
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Fig. 2 Rheology of a sulfinyl-based PPV spinning dope at 20 ꢀC, indicating shear thinning behavior, which is a prerequisite for fiber processing.
PPVs in the literature, such as poly(2,3-diphenyl phenylene
vinylene)s with hexyl and decyl side chains, also displayed
significant aggregation in solvents such as chloroform and
toluene at dilute concentrations.53 DLS indicated that large
aggregates occurred in 1 mg mLꢁ1 chloroform-based sulfinyl-
containing PPV solutions that were 102 to 190 nm in diameter.
Filtering these solutions with a 0.2 mm filter resulted in
a re-equilibration of sizes to around 7–12 nm for both polymers,
although approximately 1% of 100 nm range aggregates were still
in the solutions. Solvents other than chloroform did not result in
full dissolution of the polymer once dry or provided only cloudy
or aggregated polymer (by DLS analysis) in solution, including
solutions in DMSO, NMP, THF, DMF, and DMF with LiBr or
oxalic acid. Both polymers with 29 and 60% elimination by TGA
possessed large aggregates in chloroform, NMP, NMP with 0.01
to 0.1 M LiBr, DMF, DMF with 1 mM oxalic acid, and THF.
Aggregates were more stable and did not change significantly in
size over the course of the DLS measurements for solutions in
chloroform and NMP due to better solubility in these solvents
compared to DMF and THF. Aggregation apparent by DLS
analyses indicated that SEC values for molecular weight in
chloroform were not entirely accurate and are thus not reported.
However, a high molecular weight, soluble polymer with strong
intermolecular interactions is often best for dry or wet spinning
fibers due to its creation of a viscous spinning dope, and an
indication of sufficient molecular weight for dry or wet fiber
spinning can be derived from the viscosity of the sulfinyl-based
PPV spinning dopes which were prepared with 45 wt% polymer
with 45.0 to 50.0 g of polymer in chloroform and displayed zero
ꢀ
shear viscosities around 60 Pa s at 20 C (Fig. 2).
Fiber spinning and scanning electron microscopy (SEM)
Molecular interactions through chain entanglements or supra-
molecular forces provided homogeneous spinning solutions with
viscoelastic flow behavior, or shear thinning (Fig. 2). The poly-
mer having some elimination of polarizer was also considered
optimal to achieve less shrinkage of the fiber since removal of the
polarizing groups could be achieved thermally after fiber
formation.
Fibers were dry-spun with 2 and 5 cm3 minꢁ1 volume flow
from viscous, 45 wt% solutions of polymer in chloroform using
45.0–50.0 g of polymer. Fiber diameters of 45–50 mm could be
prepared using a spinneret with a hole diameter of 0.1 mm (L/D
0.8). Thinner fibers were also prepared with a hole diameter of
0.08 mm (L/D 0.16) and 32 holes with drawing. Fiber stretchings
from 0 to 151% were achieved by winding the fibers that extruded
at 31 m minꢁ1 onto a godet providing low tension or a negative jet
draw at 28 m minꢁ1, providing tension at 31 m minꢁ1, or
providing tension and stretch at 78 m minꢁ1. These experiments
Fig. 3 Thermograms of polymers with varied sulfinyl content; polymer synthesized in (A) polar aprotic solvents versus in (B) sec-butanol.
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Fig. 6 SEM images of crystalline platelet formation in fibers spun from
sulfinyl-based PPV synthesized in polar aprotic solvents, spinneret
0.10 mm (L/D 0.8) hole diameter, volume flow 2 cm3 minꢁ1, jet stretch
ratio 0%.
flow of 2 cm3 minꢁ1 allowed slightly more time for packing of the
polymer and slow evaporation of chloroform, but orientation in
the fibers occurred before drawing and resulted in more irregu-
larly shaped fibers. However, these first spinning trials prove that
these PPV-based polymers are viable for use in fiber-spinning
processes.
X-ray diffraction and fiber crystallinity
X-ray diffraction (XRD) measurements indicated a significant
increase in crystallinity and orientation of the polymers after
spinning and drawing compared to bulk, unspun, unprocessed
polymer samples (Fig. 7). The polymer with 60 wt% elimination
by TGA possessed an XRD profile with a broad amorphous
peak around 23ꢀ 2q and low structural order, or a more amor-
phous structure. Despite its large amorphous character, the
X-ray diffraction pattern for this unprocessed polymer did
Fig. 4 Sulfinyl-based poly(p-phenylene vinylene) fibers: (a) dry-spin-
ning, (a) and (c) fibers under ambient light, (b) and (d) fibers under UV
light; fibers collected with tension on godets rotating at 31 m minꢁ1
.
resulted in fibers with 60, 51, and 32 mm diameters, respectively.
Fluorescent, conjugated polymer fibers were collected (Fig. 4 and
5). Reducing spinneret hole size, column aspect ratio, and fiber
diameters resulted in brittle fibers that fractured easily when
removing the dried fibers from the godets. Their brittle nature
was due in part to their porous structure (Fig. 6), and elongations
at break of both the fibers and the films drawn from the spinning
dopes were not greater than 0.95 to 1.80% at 20 ꢀC with an
average of 10 different samples and a gauge length of 25 mm.
During the spinning process, a shorter column aspect ratio
(L/D 0.16) of the spinneret holes and faster spinning speed of
5 cm3 mꢁ1 allowed slightly less time for orientation of polymer
chains into large crystals. The spinneret with larger holes of
0.1 mm (L/D 0.8) resulted in dense fibers with few interior pores
and a crystalline structure (Fig. 5 and 6). This oriented crystalline
packing may be advantageous to improve fiber properties if
controlled properly and directed in the drawing direction of the
fiber. This longer, wider spinneret column and slower volume
Fig. 7 X-ray diffraction patterns of the dry-spun fibers versus the
unprocessed polymer before spinning; Al standard peak 38.5ꢀ 2q.
Fig. 5 SEM images of fibers spun from sulfinyl-based PPVs and spun under different conditions: spinneret: (a) 0.10 mm (L/D 0.8) hole diameter and (b
and c) 0.08 mm (L/D 0.16) hole diameter; volume flow: (a) 2 cm3 minꢁ1 and (b and c) 5 cm3 minꢁ1; jet stretch ratio (a) 0%, (b) ꢁ10%, and (c) 151%.
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from 24 to 79%. Comparing the fibers spun with tension and 0%
jet draw ratio and those with tension and 151% jet draw ratio, the
percent crystallinity was similar, 79 versus 78% respectively, but
ꢀ
the average crystallite size nearly doubled from 226 to 430 A.
FTIR analysis
After fiber formation, the fibers remained soluble and could
be reused. After a post-fiber formation thermal treatment at
ꢀ
150–220 C, the fibers were no longer soluble since the polymer
structure had been converted to insoluble, conjugated PPV via
elimination of the sulfinyl groups. Elimination of sulfinyl groups
ꢀ
was apparent with the application of temperature over 150 C
from the appearance of an IR peak at 965 cmꢁ1 due to the
formation of vinylene C]C. Fiber spinning was conducted at
much lower temperatures around 30 ꢀC to prevent premature
elimination. Although comparison of the ratios of the areas of
the S]O region at 1044–1022 cmꢁ1 versus the C]C region at
965 cmꢁ1 can only be interpreted qualitatively, these values
indicated that no significant elimination or increase in C]C
occurred specifically during fiber spinning and drawing (Table 1).
After processing, there was a shift in the S]O peak in the IR
spectrum from 1044 cmꢁ1 (amorphous) to 1022 cmꢁ1 (crystalline)
due to an increase in crystallinity. However, one must keep in
mind that FTIR involves analysis of a surface and depends on
surface orientation of polymer chains, and X-ray diffraction
scans the entire volume of the sample.55 The height intensity
ratios of S]O amorphous versus crystalline peaks are seen in
Table 1. The intensity of the S]O 1044 cmꢁ1 peak for the
unprocessed polymer powder was much higher than its
1022 cmꢁ1 S]O shoulder. In contrast, all fiber samples possessed
the opposite, a greater S]O 1022 cmꢁ1 peak intensity than S]O
1044 cmꢁ1 peak intensity. Poly(ethylene terephthalate) (PET), for
example, also exhibits changes in its FTIR spectrum due to
crystallization during drawing, and these changes in PET’s IR
spectrum are associated with a greater population of trans versus
gauche conformers in crystalline PET.55 Thus, this shift in the
FTIR spectrum may confirm the morphology change present
upon fiber spinning as seen in the X-ray diffraction patterns. The
S]O peak at 1022 cmꢁ1 can be associated with conformations
Fig. 8 X-ray diffraction patterns of fibers collected without tension and
thus a negative jet draw ratio of ꢁ10%, with tension but a 0% jet draw
ratio, and with tension with a jet stretch ratio of 151%; Al standard peak
38.5ꢀ 2q.
indicate a small fraction of crystallites that may have acted as
nucleating agents for larger crystals during spinning and
drawing, but the amorphous area was too large for calculation of
the percent crystallinity. This broad, amorphous peak evident in
Fig. 7 for the unprocessed polymer was also apparent in X-ray
diffraction patterns of the fibers, but the fiber profiles displayed
new, additional peaks that indicated enhanced orientation and
crystallization with fiber spinning (Fig. 7 and 8).
Crystallite size and percent varied with the type of spinneret as
well as the stretch ratio. The percent crystallinity was calculated
based on an equatorial X-ray scan of the anisotropic fiber
samples, and thus the values displayed in Table 1 depend only on
the concentration of crystallinity in an equatorial slice and not
the entire halo.54 Thus, these values do not represent an absolute
percent crystallinity throughout the sample, but they are
comparable to each other. Increasing the tension during the
collection of fibers may have led to the appearance of more
crystallites of a smaller size. Therefore, the average crystallite size
decreased when comparing fibers collected with and without
ꢀ
tension from 365 to 226 A, but the percent crystallinity increased
Table 1 Molecular and supramolecular analyses of drawn fibers as indicated via IR and XRD
Spinneret hole
aspect ratio
(L/D)
Extrusion Drawing
rate
Average Relative
FTIR intensity ratioe FTIR total
Spinneret hole
diameter (mm)
Volume flowa rate
Jet stretch crystallite crystallinityd S]O 1022 cmꢁ1
/
area ratioe
S]O/C]C
c
(cm3 minꢁ1
)
(m minꢁ1
)
(m minꢁ1
)
ratiob (%) size (A) (%)
S]O 1044 cmꢁ1
ꢀ
Unprocessed powder
0.1
0.08
0.08
0.08
—
0.8
0.16
0.16
0.16
—
2
5
5
5
—
8
28
31
78
—
0
—
—
41
25
79
78
0.67
1.26
1.53
1.43
1.41
3.82
5.48
6.13
5.22
5.82
8
31
31
31
415.2
364.8
226.0
430.4
ꢁ10
0
151
a
Spinning rates (As) were calculated based on the equation As ¼ ((V/t)n)/(Lnr2p), where the corrected volume per minute (V/t) times a correction
b
factor n is divided by the area times Ln, the number of spinneret holes. Jet stretch ratio (D) or percent stretch is calculated from the equation D ¼
c
((Ad ꢁ As)/As) ꢂ 100, where Ad is the drawing rate and As is the extrusion rate. Average crystallite size (t) was calculated based on Scherrer’s
formula, t ¼ (Kl)/(Bcos(qB)). B is the width at half height (integral breadth) of the reflection at 2q and is corrected for the internal broadening of
the equipment (ꢃ0.1ꢀ) for Lorentzian curves. K is the Scherrer factor dependent on crystallite shape and is equal to 1. The wavelength is l, and qB
d
is the Bragg angle. Crystallinity was calculated from the intensities of crystalline versus amorphous peaks via the equation percent crystallinity ¼
(SIc/(SIc + Ia)) ꢂ 100, where Ic and Ia represent the crystalline and amorphous peak intensities. This analysis was based on a peak separation of the
e
equatorial X-ray scan on an anisotropic sample and orientation was not accounted. FTIR absorbance ratios were calculated based on the peak
area for the S]O peaks at 1044–1022 cmꢁ1 versus the C]C peak at 965 cmꢁ1, and the absorbance height ratio was taken for the S]O (crystalline)
at 1044 cmꢁ1 versus the S]O (amorphous) at 1022 cmꢁ1. The height ratio was used due to overlap of the two S]O peaks.
11858 | J. Mater. Chem., 2012, 22, 11851–11860
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pertaining to crystalline polymer versus the peak at 1044 cmꢁ1
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Conclusions
18 T. Schwalm and M. Rehahn, Macromol. Rapid Commun., 2008, 29,
207–213.
The first sulfinyl-based poly(p-phenylene vinylene) fibers were
fabricated, and their unique fiber crystallinities and orientations
were examined as related to various jet draw ratios. Soluble,
processable poly(p-phenylene vinylene) was synthesized using
a sulfinyl-based monomer, and both polymer composition and
solution rheology for dry spinning were optimized. Spinning
dopes with 45.0 to 50.0 g of polymer were prepared in chloroform
and spun on a dry-spinning apparatus with various jet stretch
ratios from 0 to 151% stretch. Fluorescent fibers resulted. Fiber
diameters were under 60 mm by SEM, analyses of X-ray profiles
indicated enhanced orientation in the fibers compared to the
unprocessed polymers. Future experiments will aim to make
these fibers more uniform with enhanced mechanical properties
from the addition of additives or co-spinning. Additionally,
doping of these PPV fibers with various dopants will be examined
to achieve optimal conductivities. The thermal gradient needed
to allow the fiber structure to remain intact after elimination
of all sulfinyl groups will also be evaluated. Finally, the
PPV-progenitor described here is currently used to prepare
PPV-coated poly(ester)s.
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