Combining electron-neutral building blocks with intramolecular conformational locks affords stable, high-mobility P- and N-channel polymer semiconductors

Article abstract of DOI:10.1021/ja303401s

Understanding the relationship between molecular/macromolecular architecture and organic thin film transistor (TFT) performance is essential for realizing next-generation high-performance organic electronics. In this regard, planar π-conjugated, electron-neutral (i.e., neither highly electron-rich nor highly electron-deficient) building blocks represent a major goal for polymeric semiconductors, however their realization presents synthetic challenges. Here we report that an easily accessible (minimal synthetic steps), electron-neutral thienyl-vinylene (TVT)-based building block having weak intramolecular S???O conformational locks affords a new class of stable, structurally planar, solution-processable, high-mobility, molecular, and macromolecular semiconductors. The attraction of merging the weak TVT electron richness with supramolecular planarization is evident in the DFT-computed electronic structures, favorable MO energetics, X-ray diffraction-derived molecular structures, experimental lattice coehesion metrics, and excellent TFT performance. TVT-based polymer TFTs exhibit stable carrier mobilities in air as high as 0.5 and 0.05 cm²/V?s (n- and p-type, respectively). All-TVT polymer-based complementary inverter circuitry exhibiting high voltage gains (～50) and ring oscillator circuitry with high f_osc(～1.25 kHz) is readily fabricated from these materials by simple inkjet printing.

Full text of DOI:10.1021/ja303401s

Article
pubs.acs.org/JACS
Combining Electron-Neutral Building Blocks with Intramolecular
Conformational Locks” Affords Stable, High-Mobility P- and N-
“
Channel Polymer Semiconductors
†
‡,#
†,⊥,#
‡
‡
†
Hui Huang, Zhihua Chen, Rocio Ponce Ortiz,
Jangdae Youn, Yong-Young Noh, Kang-Jun Baeg, Lin X. Chen, Antonio Facchetti,*
Christopher Newman, Hakan Usta, Sylvia Lou,
†
§
∥
†
,†
,
†
and Tobin Marks*
†
Department of Chemistry and the Materials Research Center, Northwestern University, Evanston Illinois 60208, United States
Polyera Corporation, 8045 Lamon Avenue, Skokie, Illinois 60077, United States
‡
§
∥
Department of Chemical Engineering, Hanbat National University, Daejeon, 305-719 (Republic of Korea)
Convergence Components & Materials Laboratory, Electronics and Telecommunications Research Institute (ETRI), Daejeon
3
05-350, Republic of Korea
⊥
Department of Physical Chemistry, University of Malaga, Malaga 29071, Spain
*
S Supporting Information
ABSTRACT: Understanding the relationship between mo-
lecular/macromolecular architecture and organic thin film
transistor (TFT) performance is essential for realizing next-
generation high-performance organic electronics. In this
regard, planar π-conjugated, electron-neutral (i.e., neither
highly electron-rich nor highly electron-deficient) building
blocks represent a major goal for polymeric semiconductors,
however their realization presents synthetic challenges. Here
we report that an easily accessible (minimal synthetic steps),
electron-neutral thienyl-vinylene (TVT)-based building block
having weak intramolecular S···O “conformational locks” affords a new class of stable, structurally planar, solution-processable,
high-mobility, molecular, and macromolecular semiconductors. The attraction of merging the weak TVT electron richness with
supramolecular planarization is evident in the DFT-computed electronic structures, favorable MO energetics, X-ray diffraction-
derived molecular structures, experimental lattice coehesion metrics, and excellent TFT performance. TVT-based polymer TFTs
2
exhibit stable carrier mobilities in air as high as 0.5 and 0.05 cm /V·s (n- and p-type, respectively). All-TVT polymer-based
complementary inverter circuitry exhibiting high voltage gains (∼50) and ring oscillator circuitry with high f (∼1.25 kHz) is
osc
readily fabricated from these materials by simple inkjet printing.
INTRODUCTION
organic semiconductors that can be processed by low-cost
methodologies, such as spin coating and inkjet printing, and
■
18
Organic semiconductors have attracted intense academic and
industrial research interest because of their potential for
fabricating various opto-electronic devices, including active
which function in ambient environments without significant
19
degradation. Therefore, the rational synthesis of new organic
semiconductors and understanding of the fundamental relation-
ship between molecular/macromolecular architecture, lattice
packing, and TFT performance are essential for next-generation
low-cost, high-performance organic electronics. Furthermore,
results to date indicate that such semiconductors must contain
planar π-conjugated backbones to promote close solid-state
π−π stacking and efficient carrier transport via strong
1
−4
matrix organic light-emitting diodes (OLEDs),
organic
and organic thin film transistors
OTFT devices are fundamental building blocks
5
−10
photovoltaics (OPVs),
OTFTs).
1
1−17
(
for organic integrated circuits that can be used as radio
frequency identification cards, memories, sensors, and drivers
for flexible displays. During the past several years, organic
semiconductors have been shown to exhibit performance
comparable to or exceeding that of amorphous silicon, the
20,21
intermolecular orbital overlap.
22
14
To achieve these critical advances, rigorously planar, weakly
electron-rich organic π-structures have been sought in the quest
for solution-processable, air-stable, high-mobility polymer
most common semiconductor used in display backplanes.
However, it is essential for realizing next-generation electronics
that new semiconductors having higher mobility and greater
current-carrying capacity than amorphous silicon be developed
from rational principles. Furthermore, to fully capitalize on the
attractions of organic electronic systems, it is critical to create
Received: April 9, 2012
Published: June 8, 2012
©
2012 American Chemical Society
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2
3−29
44,45
semiconductors.
Promising building blocks include
macromolecular planarity.
However, much evidence shows
thiophenes and other heteraromatics, functionalized to enhance
that the strongly electron-donating alkoxy groups destabilize
3
0,31
38
solubility and processability.
Depending on the electronic
the π-system HOMOs and significantly compromise the TFT
structure, p- or n-channel (hole or electron transporting,
respectively) polymers can be created, with the best-performing
typically being copolymers of electron-rich (for p-type) or
electron-poor (for n-type) units, combined with weakly
electron-rich cores. For the latter, the 2.2′-bithiophene core
46
performance and ambient stability. Furthermore, combining
alkoxy-thiophenes with electron-poor units favors low-bandgap,
low-mobility ambipolar, rather than unipolar n-type, semi-
47
conductors. Here we report a new design approach based on
the TVT building block (Chart 1) functionalized with alkoxyl
groups at the CC linkage, which successfully addresses the
above issues. We report that alkoxy-TVT derivatives, such as
M1 (Chart 2), offer four major attractions: (1) Two-step
synthesis from commercially available monothiophenes having
tunable alkoxy dimensions; (2) experimental and theoretical
substantiation that alkoxy groups at the CC linkage rather
than at heteroaromatic positions minimally destabilize the
HOMOs, thereby providing a versatile, electron-neutral
building block for both p- and n-type semiconductors; (3)
quantitiative X-ray diffraction and computational evidence that
the O···S interaction functions as a “conformational lock,”
increasing planarity as well as carrier mobility; (4) synthetic
flexibility allowing additional alkoxy group introduction to
further enhance solubility and processability. Thus, TVT-based
(T2; Chart 1) has afforded n- and p-channel copolymers with
Chart 1. Structures of 2.2′-Bithiophene (T2), 1,2-Di-(2-
thienyl)-ethene (TVT) p-Channel Polymers PBTTT-R and
P(TBTR-T2), and n-Channel Polymer P(NDI2OD-T2) (R =
2-octyldecyl)
3
2,33
34,35
naphthalenediimide
or thienothiophene
/ benzodithio-
3
6
polymer TFTs exhibit carrier mobilities in air as high as 0.5 and
phene fragments, respectively. Introducing other heteroatom
groups can be used to promote planarity and enhance
2
0
.05 cm /V·s (n- and p-type, respectively) with good ambient
3
7−43
stability. We also report that complementary all-TVT polymer
inverters with high voltage gains (∼50) and ring oscillator
circuits with high fosc(∼1.25 kHz) can be readily fabricated from
these materials by straightforward inkjet printing.
mobility,
however this typically requires complex and
costly multiple-step syntheses. Weak covalent intramolecular
interactions, as in O···S contacts between alkoxyl substituents
and thiophenes, have been employed previously to promote
Chart 2. Structures of TVT-Based Monomers M1 and M2, Small Molecules SM1 and SM2, and Polymers P1−P4 (2OD = 2-
octadecyl)
1
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RESULTS AND DISCUSSION
chromatography (GPC) at 150 °C (P1: M
= 3.0 K, M = 3.4
n w
■
K; P2: M = 2.7 K, M = 2.9 K). Copolymers P3 and P4 are
Building Block and Polymer Synthesis and Character-
ization. The ethoxy-TVT building block M1 was synthesized
in two steps from thiophene (Scheme 1). Lithiation followed
n
w
readily accessed using various building blocks via Stille coupling
Scheme 2). P3 is soluble in chloroform and chlorobenzene,
and M = 18.6 K and M = 63.5 K. P4 is also soluble in
(
n
w
chlorinated solvents, and M = 12.4 K and M = 25.0 K.
Slow evaporation of an ethyl ether solution of M1 afforded
n
w
Scheme 1. Synthesis of Monomers M1 and M2
colorless crystals suitable for single crystal X-ray diffraction
(
XRD). The rigorous planarity of M1 is evident in the crystal
structure (Figure 1A), which contains cocrystallized, ∼97% of
the low-energy sp,sp conformer and ∼3% of a higher-energy
50,51
ap,ap conformer, similar to the situation in TVT crystals.
DFT computation shows that the sp,sp conformer is ∼1.15
kcal/mol lower in energy than the ap,ap conformer (Figures 1
and Supporting 1nformation). Importantly, the intramolecular
O···S distance in M1 is 2.690(2) Å, indicating a relatively strong
4
5,52
O···S interaction.
M1 crystallizes in a typical herringbone
packing motif with a π−π packing distance between
neighboring molecules of 5.137 Å (Figure S1, Supporting
1
nformation) and with no intermolecular O···S bonding. To
by addition of Cr(CO) and Et O·BF affords a chromium
further probe the importance of the O···S interaction in TVT
planarity, the corresponding structure in which both thiophene
rings are twisted from planarity, as found in M2, was computed
by fixing the dihedral angles (see Figure S1, Supporting
1nformation). These results indicate that the planar structure is
more stable by 8.3 kcal/mol. The effects of steric congestion on
the conformation of M2 can also be estimated by analyzing the
energetic differences between the twisted, the more stable
structure, and the less stable planar structure. As shown in
Figure S1, Supporting 1nformation, the M2 planar config-
uration, where the vinylene−thiophene dihedral angles are fixed
6
3
4
48
carbene intermediate, which is converted catalytically to an
olefin M1 in the presence of Pd(PPh ) with an overall yield
3
4
>
80%. For comparison, analogous M2 with n-propyl sub-
stituents can be prepared through a McMurry reaction from an
49
intermediate ketone, also prepared from thiophene. Next,
oligomers SM1 and SM2 were prepared via Stille coupling,
starting from M1 and M2, respectively (Scheme 2).
Homopolymers P1 and P2 were prepared via FeCl oxidative
3
polymerization of M1 and M2, respectively. Both P1 and P2
have modest molecular weights as evaluated by gel permeation
a
Scheme 2. Synthetic Routes to Small Molecule and Polymer Semiconductors
a
2OD = 2-octyldecyl.
1
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at 180°, is less stable than the twisted one by 6.7 kcal/mol. This
energetic difference is non-negligible, indicating significant
steric repulsions between the CH group and the S atom in M2.
2
However, this value is smaller than the energetic difference
found between M1 planar and twisted configurations (8.3 kcal/
mol) and points to the importance of the O···S conformational
lock in M1.
The HOMO and LUMO energies of M1 and M2 obtained
from cyclic voltammetry (CV) and optical spectroscopy are
−
5.6/−2.3 eV and −5.8/−2.2 eV, respectively (Table 1; Figure
1
E and Figures S2 and S3, Supporting 1nformation). The
smaller M1 bandgap argues for greater π-system delocalization
than in M2, which is supported by the DFT computations
(
Figure S4, Supporting 1nformation). Note that the HOMO
53,54
and LUMO energies of TVT lie at −5.7/−2.4 eV,
respectively, only slightly different from those of M1 (Figure
1
E) and arguing that M1 can serve as a electron-neutral
building block for both n- and p-type π-semiconductors. In
contrast to electron-neutral T2, the corresponding planar 3,3′-
dialkoxy-2,2′-dithiophene T2(OR) (Figure 1D) with alkoxy-
2
substituted thienyl rings, has far higher-lying MOs by ∼1 eV
(
HOMO/LUMO = −6.2/−2.4 vs −5.2/−1.6 eV, respectively),
Figure 1. (A) Top and side views of the diffraction-derived molecular
structure of M1. Selected bond lengths (Å) and angles (°) are C1−C5,
portending very different TFT characteristics. From the DFT
calculations, the HOMO/LUMO energies of electron-neutral
T2 and TVT and the corresponding T2(OR)2 and M1
molecules are computed to lie at −5.4/−1.2 and −5.2/−1.6 eV
vs −4.7/−0.7 and −5.1/−1.3 eV, respectively, consistent with
the aforementioned experimental data (Figure 1E and Figure
S4, Supporting 1nformation). Not surprisingly, oligomer SM1
1
.491(4); C5−C5a, 1.340(6); C5−O1, 1.411(3); C6−O1, 1.455(2);
C1−C5−O1, 122.3(3);C1−C5−C5a, 122.1(4); O1−C5−C5a,
1
3
15.5(3). (B−D) Top and side views of the DFT//B3LYP/6-
1G**-derived molecular structures of M1, M2, and T2(OR) (R =
2
OC H ), respectively. (E) Experimental HOMO/LUMO energies for
2
5
the indicated systems.
(
−5.4/−3.5 eV) has a smaller bandgap than SM2 (−5.4/−3.1
eV), and this pattern extends to homopolymers P1 and P2
a
Table 1. Thermal, Optical Absorption/Emission, Electrochemical, Field-Effect Mobility (μe/h) , and Current I /I Ratio Data
on off
for Compounds M1-SM2 and P1−P4
a
c
c
Bottom-gate/top-contact structure if not stated otherwise. Heating/cooling cycle. Top-gate/bottom-contact structure.
1
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Figure 2. (A) Tapping mode AFM images (5 × 5 μm) and out-of-plane XRD scans of spin-coated P3 thin films without annealing: (B) after
annealing at 200 °C; (C) after annealing at 240 °C; and (D) in-plane (top) and out-of-plane (bottom) grazing incidence X-ray scattering of spin-
coated P3 thin films after annealing at 240 °C.
(
Table 1). The LUMO and HOMO energies of copolymers P3
S7, Supporting 1nformation), the P3 film morphology and
microstructure change on annealing. AFM images of P3 suggest
that the morphology varies in a way suggesting that defects
develop due to inadequate wetting during film formation,
thereby creating a nonoptimal interface with the substrate.
However, out-of-plane XRD reveals that P3 films become
significantly more crystalline on annealing (Figure 2). A single
family of Bragg reflections is observed without the obvious π−π
stacking distance; the observed d-spacing is ∼22.2 Å after 200
°C annealing and is reasonably assigned to a-axis (h00) lamellar
spacing. This distance is consistent with structural models in
which the polymer side chains are either interdigitated or
closely packed and tilted out of the molecular plane, implicating
a polymorph in which most molecules have an edge-on
orientation with respect to the substrate and with the π−π
stacking axis parallel to the substrate plane, thus favoring in-
and P4 are −4.0/−5.4 and −3.2/−5.1 eV, respectively,
essentially identical to those of analogous polymers P-
(
NDI2OD-T2) and P(TBTR-T2) (Chart 1; −3.9/−5.4 and
3
2,33,36
−
3.0/−5.2 eV),
indicating that M1 electronic properties
are similar to those of weak electron-rich 2,2′-bithiophene. Note
also that such low-lying P3 LUMO (−4.0 eV) and P4 HOMO
4
6
energies portend significant TFT ambient stability.
For closely related molecular materials, melting points are
informative indices of lattice cohesion. Note that the M1
melting point, 158−159 °C, is significantly higher than that of
55
1
,2-di-(2-thienyl)-ethene, 133−134 °C. Introducing alkyl or
alkoxyl substituents on conjugated backbones typically
depresses the melting points, as in E-1,2-bis(3-methoxy-2-
56
57
thienyl)ethene (mp 121−123 °C), which is substantially
lower melting than that of the 1,2-di-(2-thienyl)-ethene parent
structure. The high M1 melting point is reasonably attributable
to the aforementioned O···S supramolecular backbone rigid-
ification/planarization. Note that, in contrast to M1, compound
M2 is an oil at room temperature, doubtless due to the twisted
backbone and looser intermolecular π−π stacking, while the
melting point of oligomer SM1 is 220−221 °C, significantly
higher than that of oligomer SM2, 62−63 °C. The
thermogravimmetric analysis (TGA) data indicate that both
P3 and P4 are thermally stable up to 280 °C, with less than 1%
weight loss (Figure S5, Supporting 1nformation). The thermal
transitions of the present materials were further investigated by
DSC (Table 1, Figure S6, Supporting 1nformation). Oligomer
SM2 and polymers P1, P2, and P4 exhibit no detectable
endotherms/exotherms in the heating/cooling cycles, while
oligomer SM1 exhibits a single endothermic transition at 218
58
plane source → drain charge transport. The crystallinity of
the solution-cast P3 films depends on the annealing temper-
ature, with a major evolution in long-range order (Figure 2)
and increasing crystallinity on increasing the annealing
temperature from 200 to 240 °C. A significant increase in the
(100) reflection intensity and the appearance of higher order
reflections up to (300) is observed, while the d-spacing remains
unchanged, indicating far greater intragrain ordering. In-plane
microstructure was also analyzed by grazing incidence
synchrotron X-ray scattering to better understand π−π stacking
in the highly crystalline P3 films (Figure 2d). The 240 °C
annealed films exhibit a (010) feature consistent with a π−π
stacking distance (4.11 Å), slightly larger than in P(NDI2OD-
59
T2) (3.93 Å). The other features can be indexed as (00l)
reflections based on a tentative orthorhombic cell with the c
axis equal to the repeat unit (15.1 Å), consistent with the DFT
results (15.6 Å). All other things being equal, this substantial
domain ordering should favor charge transport.
°
C, tentatively attributable to backbone melting and an
exothermic recrystallization peak at 173 °C. Polymer P3
exhibits a broad endotherm at 220 °C and an exotherm at 200
°
C, arguing that microstructural order is enhanced by
annealing.
Thin Film Microstructural Characterization. AFM and
XRD images (Figure 2) indicate that, in contrast to P4 (Figure
Field-Effect Transistor Fabrication and Measure-
ments. Bottom-gate/top-contact (BGTC) and top-gate/
bottom-contact (TGBC) geometry TFTs were fabricated on
+
HMDS-treated p -Si/SiO and glass substrates, respectively.
2
1
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Data are summarized in Table 1 and Figure 3. The hole
mobilities of oligomer SM1 and homopolymer P1 are 3 × 10−3
Figure 4. Complementary inverters fabricated from the P3 (n-
channel) + P4 (p-channel) materials set. (A) Schematic electrical
layout of the inverters. (B) Optical image of the inkjet-printed droplets
and of the TFT structure used in the inverter. (C) Static switching
characteristics of an inkjet-printed inverter with a PMMA gate
dielectric. (D) Gain data for the inkjet-printed devices.
gate dielectric was spin-coated on top and via holes defined
by selective removal of dielectric layer portions with inkjetted
solvent. Figure 4 shows the voltage transfer characteristics of a
complementary inverter fabricated with P3 and P4 (p and n-
type, respectively) (W /L = 5.0 mm/20 μm, W /L =1.0 mm/
p
p
n
n
2
0 μm) at various supply voltages (VDD). Note that the static
Figure 3. (A) Transistor architectures employed in this study. TFT
transfer plots for polymers P3 in (B) BGTC and (C) TGBC
structures. (D) P4 in a BGTC structure at VSD = −100 V.
inverter characteristics exhibit negligible bias hysteresis and
high voltage gains (∼50) above V = −50 V. The inverting
point (V ) of the complementary inverter is reached when
DD
inv
both the p- and n-channel transistors are operating in the
saturation region and is expressed by eq 1:
and 3 × 10− cm /V·s, respectively, while SM2 and
homopolymer P2 based on the extensively twisted M2 building
block are both TFT inactive. Films of P3 exhibit n-type
behavior, and the electron mobility without annealing is 2.3 ×
6
2
62
p
n
Th
β
n
VDD + V + V
Th
β
p
−3
2
Vinv =
10
cm /V·s, consistent with the low crystallinity evident in
β
n
1
+
Figure 2a. However, annealing at 200 °C increases the mobility
β
p
(1)
2
to ∼0.2 cm /V·s, among the highest reported to date for an n-
60
type polymer in a BGTC TFT architecture (Figure 3A), while
where the β = (W/L)μFETCi is a design factor to adjust the p-
and n-channel currents of transistors, and the superscripts p
and n denote the semiconductor type. The Vinv of our device
was slightly shifted in the negative direction, by ∼ −15 V (at
a TGBC device affords an increased electron mobility of ∼0.5
2
cm /V·s in ambient (Figure 3B). Note that the high film
crystallinity assayed by the XRD data closely parallels the TFT
results. Note also that the electron mobility falls to 1.5 ×
0 cm /V·s on annealing the films at 240 °C, although the
crystallinity is slightly increased as judged by the XRD results.
That increased crystallinity does not necessarily correlate with
increased mobility is reasonable since transport across grain
V
= −70 V) from the ideal switching position at 1/2 V ,
DD
DD
−3
2
1
and the W /L vs W /L ratio was modified to 5:1 for
p p n n
enhancing β compared to β . This was mainly attributed to the
p n
higher μFET,e (∼10× ) of the P3 (n-type) OFET than that of P4
(p-type) OFET, since the other parameters, such as V and C ,
DD
i
35
boundaries can significantly limit organic TFT performance.
were equal to both p- and n-type OFETs. It is noted that the
All of the present TFTs exhibit significant ambient stability, in
accord with the low-lying LUMOs. After eight months in
laboratory air, the ambient mobility of the TFTs with TGBC
factor β must be carefully designed to optimize the perform-
ance of complementary circuits. The inverting voltage (V ) is
inv
slightly shifted in the negative direction with respect 1/2 VDD
due to the slight VTh and mobility differences between the p-
and n-type materials. For further integration of P3- and P4-
based OTFTs, the dynamic switching characteristics in the
present circuits were next evaluated using a complementary
ring oscillator (L = L = 10 μm; W = W = 10 mm). A top view
2
structures falls from 0.11 to 0.05 cm /V·s. The maximum hole
2
mobility of copolymer P4 in a TGBC device, 0.05 cm /V·s, is
achieved after thermal annealing in air at 130 °C for 1 h (Figure
3
). The mobility of P4 is moderate compared to other
61
copolymers based on benzodithiophene building blocks.
n
p
n
p
Organic Complementary Circuits. Complementary inverter
and ring oscillator circuits were next fabricated from P3 and P4
by inkjet printing (Figure 4). For monolithic integration of the
p- and n-channel OTFTs, the P3 and P4 solutions were
sequentially printed onto photolithography patterned Au
bottom-contact electrodes. The poly(methylmethacrylate)
of inkjet-printed features of P3 and P4 and the completed
polymer ring oscillator are shown in Figure S8, Supporting
1nformation. The oscillation frequencies (f ) of this circuit
osc
exhibit a strong dependence on VDD, and the highest fosc of
∼1.25 kHz is obtained at V = 100 V. Note that the device
DD
speed would be further increased by reducing gate to source
1
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63
and drain electrode overlap capacitances, and by using
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CONCLUSIONS
■
132, 5394−5404.
These results demonstrate a versatile, electron-neutral, planar
π-electron building block offering both “conformational locks”
and appropriate HOMO/LUMO energetics to enforce π-
system planarity, high carrier mobility, solution processability,
and ambient stability for both p- and n-type organic
semiconductors. TVT-based polymer TFTs exhibit carrier
mobilities in air as high as 0.5 and 0.05 cm /V·s (n- and p-
type, respectively). Using a n- and p-type variant of the same
material permits fabrication of high-performance complemen-
tary inverters (voltage gains ≈ 50) and ring oscillators (f
.25 kHz) by simple inkjet printing. Furthermore, these results
provide fundamental insight into the interrelationships between
molecular/macromolecular architecture and organic TFT
performance and durability.
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3
ASSOCIATED CONTENT
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*
S
Supporting Information
2011, 23, 568−582.
(17) Usta, H.; Facchetti, A.; Marks, T. J. Acc. Chem. Res. 2011, 44,
Monomer/polymer synthesis and characterization, H and 13C
NMR spectra, XPS spectra, UV−vis spectra, DFT calculations,
DSC and TGA traces, OTFT fabrication details, OTFT device
data, and thin-film characterization (XRD and AFM) for
different polymer deposition conditions. This material is
available free of charge via the Internet at http://pubs.acs.org.
1
5
(
01−510.
18) Katz, H. E. Chem. Mater. 2004, 16, 4748−4756.
19) Bobbert, P. A.; Sharma, A.; Mathijssen, S. G. J.; Kemerink, M.;
(
de Leeuw, D. M. Adv. Mater. 2012, 24, 1146−1158.
20) Salleo, A.; Kline, R. J.; DeLongchamp, D. M.; Chabinyc, M. L.
Adv. Mater. 2010, 22, 3812−3838.
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3898.
(
(
AUTHOR INFORMATION
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(
22) Beaujuge, P. M.; Frechet, J. M. J. J. Am. Chem. Soc. 2011, 133,
Corresponding Author
a-facchetti@northwestern.edu; t-marks@northwestern.edu
2
(
0009−20029.
23) Zhan, X.; Facchetti, A.; Barlow, S.; Marks, T. J.; Ratner, M. A.;
Author Contributions
These authors contributed equally.
Wasielewski, M. R.; Marder, S. R. Adv. Mater. 2011, 23, 268−284.
#
(24) Bronstein, H.; Chen, Z.; Ashraf, R. S.; Zhang, W.; Du, J.;
Durrant, J. R.; Shakya Tuladhar, P.; Song, K.; Watkins, S. E.; Geerts,
Y.; Wienk, M. M.; Janssen, R. A. J.; Anthopoulos, T.; Sirringhaus, H.;
Heeney, M.; McCulloch, I. J. Am. Chem. Soc. 2011, 133, 3272−3275.
Notes
The authors declare no competing financial interest.
(25) Beaujuge, P. M.; Reynolds, J. R. Chem. Rev. 2010, 110, 268−320.
ACKNOWLEDGMENTS
(26) McCulloch, I.; Heeney, M.; Chabinyc, M. L.; De Longchamp,
D.; Kline, R. J.; Colle, M.; Duffy, W.; Fischer, D.; Gundlach, D.;
Hamadani, B.; Hamilton, R.; Richter, L.; Salleo, A.; Shkunov, M.;
Sparrowe, D.; Tierney, S.; Zhang, W. Adv. Mater. 2009, 21, 1091−
■
We thank AFOSR (FA9550-08-1-0331) and Polyera Corp. for
support of this research at Northwestern University. We thank
the NSF-MRSEC program through the Northwestern Materials
Research Center (DMR-1121262) for support of character-
ization facilities. R.P.O. acknowledges funding from the
European Community’s Seventh Framework Programme
through a Marie Curie International Outgoing Fellowship
1
(
109.
27) Allard, S.; Forster, M.; Souharce, B.; Thiem, H.; Scherf, U.
Angew. Chem., Int. Ed. 2008, 47, 4070−4098.
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grant agreement 234808). Use of the Advanced Photon
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31) Marks, T. J. MRS Bull. 2010, 35, 1018−1027.
32) Chen, Z.; Zheng, Y.; Yan, H.; Facchetti, A. J. Am. Chem. Soc.
Source was supported by the U. S. Department of Energy,
Office of Science, Office of Basic Energy Sciences, under
Contract No. DE-AC02-06CH11357.
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33) Yan, H.; Chen, Z.; Zheng, Y.; Newman, C.; Quinn, J. R.; Dotz,
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