Bithiophene-imide-based polymeric semiconductors for field-effect transistors: Synthesis, structure-property correlations, charge carrier polarity, and device stability

Article abstract of DOI:10.1021/ja107678m

Developing new high-mobility polymeric semiconductors with good processability and excellent device environmental stability is essential for organic electronics. We report the synthesis, characterization, manipulation of charge carrier polarity, and device air stability of a new series of bithiophene-imide (BTI)-based polymers for organic field-effect transistors (OFETs). By increasing the conjugation length of the donor comonomer unit from monothiophene (P1) to bithiophene (P2) to tetrathiophene (P3), the electron transport capacity decreases while the hole transport capacity increases. Compared to the BTI homopolymer P(BTimR) having an electron mobility of 10 ^-2 cm² V^-1 s^-1, copolymer P1 is ambipolar with balanced hole and electron mobilities of ～10^-4 cm² V^-1 s^-1, while P2 and P3 exhibit hole mobilities of ～10^-3 and ～10^-2 cm² V^-1 s^-1, respectively. The influence of P(BTimR) homopolymer M_n on film morphology and device performance was also investigated. The high M_n batch P(BTimR)-H affords more crystalline film microstructures; hence, 3- increased electron mobility (0.038 cm ² V^-1 s^-1) over the low M_n one P(BTimR)-L (0.011 cm² V^-1 s^-1). In a top-gate/bottom-contact OFET architecture, P(BTimR)-H achieves a high electron mobility of 0.14 cm² V^-1 s^-1, only slightly lower than that of state-of-the-art n-type polymer semiconductors. However, the high-lying P(BTimR)-H LUMO results in minimal electron transport on exposure to ambient. Copolymer P3 exhibits a hole mobility approaching 0.1 cm² V^-1 s^-1 in top-gate OFETs, comparable to or slightly lower than current state-of-the-art p-type polymer semiconductors (0.1-0.6 cm ² V^-1 s^-1). Although BTI building block incorporation does not enable air-stable n-type OFET performance for P(BTimR) or P1, it significantly increases the OFET air stability for p-type P2 and P3. Bottom-gate/top-contact and top-gate/bottom-contact P2 and P3 OFETs exhibit excellent stability in the ambient. Thus, P2 and P3 OFET hole mobilities are almost unchanged after 200 days under ambient, which is attributed to their low-lying HOMOs (>0.2 eV lower than that of P3HT), induced by the strong BTI electron-withdrawing capacity. Complementary inverters were fabricated by inkjet patterning of P(BTimR)-H (n-type) and P3b (p-type).

Full text of DOI:10.1021/ja107678m

ARTICLE
pubs.acs.org/JACS
Bithiophene-Imide-Based Polymeric Semiconductors for Field-Effect
Transistors: Synthesis, Structure-Property Correlations,
Charge Carrier Polarity, and Device Stability
†
†
‡
‡
§
Xugang Guo, Rocio Ponce Ortiz, Yan Zheng, Yan Hu, Yong-Young Noh, Kang-Jun Baeg,
,
†,‡
,†
Antonio Facchetti,* and Tobin J. Marks*
†
Department of Chemistry and the Materials Research Center, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208,
United States
‡
Polyera Corporation, 8045 Lamon Avenue, Skokie, Illinois 60077, United States
§
Department of Chemical Engineering, Hanbat National University, 16-1 Dukmyung-dong, Yuseong-gu, Daejeon, 305-719,
Republic of Korea
Convergence Components & Materials Laboratory, Electronics and Telecommunications Research Institute (ETRI), 138 Gajeongno,
Yuseong-gu, Daejeon 305-350, Republic of Korea
S Supporting Information
b
ABSTRACT: Developing new high-mobility polymeric semiconductors with good processability and excellent device environ-
mental stability is essential for organic electronics. We report the synthesis, characterization, manipulation of charge carrier polarity,
and device air stability of a new series of bithiophene-imide (BTI)-based polymers for organic field-effect transistors (OFETs). By
increasing the conjugation length of the donor comonomer unit from monothiophene (P1) to bithiophene (P2) to tetrathiophene
(
P3), the electron transport capacity decreases while the hole transport capacity increases. Compared to the BTI homopolymer
-
2 2
-1 -1
P(BTimR) having an electron mobility of 10 cm V
s , copolymer P1 is ambipolar with balanced hole and electron mobilities
2
-
4
2
-1 -1
-3
-2
-1 -1
of ∼10 cm V
s
, while P2 and P3 exhibit hole mobilities of ∼10 and ∼10 cm V
s , respectively. The influence of
P(BTimR) homopolymer M on film morphology and device performance was also investigated. The high M batch P(BTimR)-H
n
n
2
-1 -1
affords more crystalline film microstructures; hence, 3ꢀ increased electron mobility (0.038 cm V
s
) over the low M one
n
2
-1 -1
P(BTimR)-L (0.011 cm V
s
). In a top-gate/bottom-contact OFET architecture, P(BTimR)-H achieves a high electron
, only slightly lower than that of state-of-the-art n-type polymer semiconductors. However, the high-
lying P(BTimR)-H LUMO results in minimal electron transport on exposure to ambient. Copolymer P3 exhibits a hole mobility
2
-1 -1
mobility of 0.14 cm V
s
2
-1 -1
approaching 0.1 cm V
semiconductors (0.1-0.6 cm V
s
in top-gate OFETs, comparable to or slightly lower than current state-of-the-art p-type polymer
2
-1 -1
s
). Although BTI building block incorporation does not enable air-stable n-type OFET
performance for P(BTimR) or P1, it significantly increases the OFET air stability for p-type P2 and P3. Bottom-gate/top-contact
and top-gate/bottom-contact P2 and P3 OFETs exhibit excellent stability in the ambient. Thus, P2 and P3 OFET hole mobilities
are almost unchanged after 200 days under ambient, which is attributed to their low-lying HOMOs (>0.2 eV lower than that of
P3HT), induced by the strong BTI electron-withdrawing capacity. Complementary inverters were fabricated by inkjet patterning of
P(BTimR)-H (n-type) and P3b (p-type).
’
INTRODUCTION
After extensive research efforts by industry, government, and
academia, several high-performance organic semiconductor classes
Solution-processable organic semiconductors are attractive
18
have emerged. With few exceptions, organic semiconducting
materials can be divided into two classes: small molecules (or
for their potential applicability in low-cost electronics and
compatibility with plastic substrates, thereby enabling mechani-
1
9,20
21
1
-4
oligomers)
and macromolecules, both offering distinct
cally flexible circuits.
Major applications of organic semicon-
advantages and disadvantages in terms of processability and
ductor-based electronics include organic light-emitting diodes
2,22
5
,6
7,8
device performance.
Organic small molecules having well-
(
OLEDs), organic field-effect transistors (OFETs), organic
9,10
defined chemical structures can be obtained in high purity levels
through conventional purification techniques such as chroma-
tography, sublimation, and recrystallization. Both vacuum eva-
poration and solution-based techniques can be employed to fabricate
photovoltaic (OPV) cells, and organic electrochromic devices
1
1,12
(
ECDs).
The key attraction of organic semiconductors versus
conventional inorganic-based materials is the possibility of fabricating
electronic devices by solution-based methodologies such as spin
13-16
19
coating and printing.
Thus, the pivotal prerequisite to advance
OFETs using small molecule filmsastheactivecomponent. Using
from emerging prototypes to widespread applications is to develop
high-performance soluble organic semiconductors with acceptable
performance and robust air stability.
Received: August 30, 2010
Published: January 5, 2011
17
r 2011 American Chemical Society
1405
dx.doi.org/10.1021/ja107678m
|J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
Figure 1. Structures of representative polymeric semiconductors for high-performance OFETs.
3
5
36
solution-based film growth techniques, substantial carrier mobi-
large charge carrier mobilities for F8T2, TS6T2, and CDT-
2
-1 -1 23
37
lities (up to 1.0 cm V
s
)
can be obtained with small
BTZ (Figure 1). The second widely used approach is to insert a
spacer between alkylated bithiophene units, as reported for
pBTTT, PQT, and PBThDTP (Figure 1). These regio-
symmetric polythiophenes exhibit excellent device performance
with hole mobilities as high as 0.6 cm V
Recently, Bao and Watson reported the high-mobility p-type
polymeric semiconductors PQTBTz-C12 and PhBT12 with
HH linkages in the polymer backbones and having hole mobi-
24
molecules due to their highly ordered film microstructures. Never-
theless, small molecule semiconductors generally lack good
solution film-forming properties due to the limited achievable
solution viscosities, presenting a challenge to widespread appli-
38
39
40
2
-1 -1
17,21
s
(pBTTT).
1
8
41
42
cations in printed electronics. In contrast, polymers are poly-
disperse materials systems for which large-scale purification methods
are generally limited to reprecipitation or extraction (e.g., Soxhlet).
Trace impurities that are difficult to remove from polymers may
significantly compromise device performance. Furthermore,
OFET performance reproducibility from polymer batch to batch
can be problematic due to molecular weight and polydispersity
4
1
42
2
-1 -1
lities up to 0.3 cm V
s . The high mobilities are due to well-
organized solid state nanostructures, enabled by intermolecular
43
π-π interactions in PQTBTz-12 and by close intramolecular
4
4,45
sulfur-oxygen contacts in PhBT12.
2
5-27
variations.
Thus, dispersion in the polymer chain length
Although current state-of-the-art polymer semiconductors
2
1
may adversely affect self-organization during film formation,
exhibit hole and electron mobilities comparable to that of
46 14
resulting in lower OFET carrier mobility in comparison to small
amorphous silicon (a-Si) (BBL and P(NDI2OD-T2) in
14,31,46
2
molecules. However, polymeric semiconductors typically exhi-
Figure 1),
their technological application is hindered by
17
bit minimal grain boundaries and excellent film-forming proper-
poor air stability. P-doping by reaction with ambient O and
2
38,47
ties, which are essential for fabrication of large-area devices by
ozone results in decreased hole mobilities,
currents, and hence lower I /I ratios as well as a positive shift
For π-electron organic semicon-
ductors, the frontier molecular orbital (FMO) energies play a
crucial role in device stability. For n-type organic semiconduc-
tors, the electron affinity (EA) controls the resistance of con-
ducting organic radical anions to oxidative dopants such as O ,
H O, and ozone. Lowering the energy of the lowest unoccu-
increased off
1
8
printing.
on off
8,49
4
In the evolution of high-performance polymer semiconduc-
tors (Figure 1), the development of highly regioregular poly(3-
hexylthiophene) (rr-P3HT or P3HT) marked a milestone in the
of the threshold voltage.
15
2
8,29
2
-1 -1
OFET field.
was achieved by optimizing film microstructure and device
design.
A hole mobility greater than 0.1 cm V
s
2
3
0-33
50
Understanding the relationship between charge
2
transport and film microstructure provided chemists with funda-
pied molecular orbitals (LUMOs) has been shown to enhance
17,21,34
51-53
mental design rules for designing optimum OFET materials.
In contrast to regioregular poly(3-hexylthiophene), regioirregular
or regiorandom) poly(3-hexylthiophene) (ri-P3HT) exhibits
the air stability of n-type materials.
For p-type organic semi-
conductors, a relatively low ionization potential (IP) will result in
p-type doping, and thus, lowering the energies of the highest
occupied molecular orbitals (HOMOs) results in increased air
(
21
several orders of magnitude lower mobility and head-to-head
HH) linkages in ri-P3HT induce steric repulsion between
54
(
stability. Therefore, fine tuning of the semiconductor FMO
55,56
neighboring thiophene units, resulting in a twisted polymer back-
bone which eliminates three-dimensional structural ordering in the
solid state.
To maintain good processability and to avoid HH linkage-
induced steric repulsion, two common approaches have been
utilized in designing new materials. In the first, a bridging atom is
used to lock the conformation between adjacent aromatic rings,
thus enforcing a coplanar polymer backbone. This strategy yields
energies must be a major focus for improving OFET stability.
Diverse strategies have been used to enhance the device stability.
For example, pBTTT exhibits superior air stability versus P3HT
because the thienothiophene ring has greater resonance stabili-
zation than thiophene, resulting in enhanced localization and a
17
HOMO lowered by ∼0.1 eV. Also, the fluorene unit has been
inserted into the polythiophene backbone of F8T2 to enhance
3
6
stability. A donor-acceptor strategy is another effective
1
406
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
Figure 2. Structures of BTI-based homopolymers and copolymers for OFETs, where P(BTimR) and P(BTimR-BT) have been reported by us
60
previously.
4
8,57
approach,
and thus, silole-based polymer TS6T2 (Figure 1) has a
substantial solubility and crystallinity and in top-gate/bottom-contact
2
-1 -1
significantly lower HOMO energy (-5.2 eV) than that of P3HT
OFETs, and have a hole mobility approaching 0.1 cm V
s ,
2
(
V
-4.9 eV), a demonstrated mobility approaching 0.06 cm
similar to the performance of P3HT OFETs fabricated under the
-
1
-1
14
s
, and undergoes negligible degradation of OFET char-
same conditions. Furthermore, due to the low HOMO energies
acteristics after 2000 repeated on-off cycles under ambient
conditions. Although several polymeric semiconductors with
low HOMO energies have been synthesized, these materials
of P2 and P3, OFETs fabricated from them exhibit enhanced
device air stability. Among them, both bottom-gate/top-contact
3
6
(
BG/TC) and top-gate/bottom-contact (TG/BC) devices fab-
2
-1 -1
usually exhibit low hole mobilities (<0.01 cm V s ) due to the
energetic barrier for hole injection from the Fermi level of the Au
ricated from P3 exhibit negligible performance degradation
(
mobility, I /I ratio, V ) after exposure to ambient for
on off t
3
5,58,59
electrode.
HOMOs to enhance air stability and yet enable high carrier
Thus, developing materials with low-lying
9
months. We believe that these new polymers offer significant
attractions versus other state-of-the-art p-type polymer semicon-
17
mobility remains an important scientific challenge.
40,47
ductors which degrade faster
or in which enhanced air
We report here a new series of bithiophene-imide (BTI)-
based conjugated copolymers for OFETs created by sequentially
varying the conjugation length of the electron-donor block. Our
previous work demonstrated that the BTI building block is highly
planar and exhibits antiparallel BTI packing and close π-π stacking
in the solid state, and that polymers constructed from BTI units
can exhibit high solubility in common organic solvents. The BTI-
based homopolymers P(BTimR) (Figure 2) are highly crystal-
line and exhibit moderate electron mobility (0.011 cm V
while BTI-based copolymers P(BTimR-BT) exhibit a hole
mobility of 0.008 cm V
M = 1.8 KDa) and poor solubility. Here we first investigate
how the P(BTimR) homopolymer molecular weight affects FET
performance by synthesizing high (P(BTimR)-H) and low
P(BTimR)-L) molecular weight batches. The high molecular
38,61,62
stability is achieved by controlling atmospheric humidity.
The present results show that BTI is effective in tuning FMO
energies and that proper self-assembly of these polymers achieves
2
-1 -1
both high mobility (∼0.1 cm V
s ) and ambient stability.
60
’
RESULTS AND DISCUSSION
The materials design strategy in this study grows out of
previous research using BTI as a novel building block for con-
2
-1 -1
s ),
60
structing polymeric semiconductors. However, the n-type BTI
homopolymer P(BTimR) suffers from poor OFET air stability,
and the p-type bithiophene-containing BTI copolymer P(BTimR-
BT) has limited solubility and low mobility. Using monothiophene or
alkylated oligothiophenes as donor units, we show here that the
present approach affords BTI-based copolymers having higher
molecular weights and greater solubilities than P(BTimR-
BT). Furthermore, dilution of the BTI loading in the polymer
backbones results in polymers with tunable FMO energies, increased
hole mobilities, and far greater p-type OFET environmental air
stability.
2
-1 -1
s , despite the low molecular weight
(
n
(
2
-1 -1
weight batch affords an electron mobility of ∼0.04 cm V
s
in bottom-gate OFETs, which is 3ꢀ larger than the mobility we
previously reported for P(BTimR) (here P(BTimR)-L). The
mobility of P(BTimR)-H measured in top-gate OFETs is up
2
-1 -1
to 0.19 cm V 4s , which is slightly lower than that of
1
P(NDI2OD)-T2. However, the high LUMO energy of the
The materials presented here are characterized by NMR, EA,
GPC, DSC, optical spectroscopy, and electrochemistry. XRD
and AFM are employed to characterize and analyze the micro-
structure and morphology of the spin-cast films, followed by FET
device fabrication, evaluation, and optimization. Then device air
stability over time is evaluated and discussed as a function of
macromolecular architecture and electronic structure. Finally,
functional complementary inverters are fabricated from these
materials by inkjet printing.
homopolymer prevents electron transport when the devices are
exposed to air. We next further develop BTI-based polymer
architectures by copolymerization with selected electron-donat-
ing oligothiophenes (P1-P3, Figure 2). It will be seen that using
the monothiophene unit as a donor yields copolymer P1 which
exhibits ambipolar behavior having balanced hole and electron
-
4
2
-1 -1
mobilities (4 ꢀ 10 cm V
s ). Using a TT linkage con-
taining bithiophene as a donor yields copolymers P2 which
-
3
2
-1 -1
exhibit a moderate hole mobility of 2 ꢀ 10 cm V
s
. Using
Synthesis of Monomers and Polymers. The synthesis of
the BTI-based monomers and polymers is depicted in Schemes 1
tetrathiophene as a donor affords copolymers P3, which exhibit
1
407
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
a
Scheme 1. Synthetic Routes to Comonomers for BTI Polymer Synthesis
a
Reagents and conditions: (i) n-RMgBr, Ni(dppp)Cl , ether, 12 h, 60 °C; (ii) n-BuLi, TMEDA, CuCl , ether, reflux to -78 °C, then to rt; (iii) n-BuLi,
2
2
Me
3
SnCl, THF, -78 °C to rt; (iv) NBS, CHCl
3
/HOAc, rt, 2 h; (v) PdCl
2
(PPh
3
)
2
2 3
, THF, 90 °C, 12 h; (vi) Br , CHCl /HOAc, rt, 4 h.
a
Scheme 2. Synthesis of BTI-Based Copolymers
a
Reagents and conditions: (i) Pd (dba) , P(o-tolyl) , THF, 80 °C.
2
3
3
and 2. To increase the solubility of the BTI-based polymers and
with the corresponding 2-trimethylstannyl-3-alkylthiophenes 5a
and 5b under Stille coupling conditions to afford monomer
precursors 7a and 7b in moderate yields (>50%). Monomers 8a
and 8b were then obtained in high yields (>90%) by reaction of
7a and 7b with bromine.
at the same time to retain close π-π stacking, tail-to-tail (TT)
52
linkages containing bithiophene units 2a and 2b were incor-
porated into the polymer backbones. The monomer precursors
63
2
a and 2b were synthesized following published procedures.
The precursors were easily converted to monomers 3a and 3b by
The syntheses of BTI-based polymers P1-P3 (Scheme 2)
were carried out using metal-catalyzed Stille polymerizations
with Pd (dba) /P(o-tolyl) as the catalyst in THF at 80 °C. The
polymerization times were varied to optimize the solubility and
processability of the product polymers. After polymerization was
complete, 2-(tributylstannyl)thiophene was added and the reac-
tion mixtures stirred for another 24 h to end cap the polymer
chains. The reactions were then quenched in methanol and the
product copolymers collected by filtration. Purifications were
n-BuLi lithiation of 2a and 2b, respectively, and then quenching
1
13
with trimethylstannyl chloride. H and C NMR spectroscopy
indicates that the crude monomers are already quite pure (>90%),
and purification of 3a and 3b is easily accomplished by recrys-
tallization from hexane at -78 °C to afford the target products in
high yields (>80%). The synthesis of key BTI building block 6
2
3
3
60
was first reported from this laboratory in 2008. The synthesis of
8
a and 8b starts from 6a and 6b, respectively, which were reacted
1
408
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
-
5
Figure 3. Optical absorption spectra of polymer P1-P3 solutions (left) in chloroform (1 ꢀ 10 M) and as pristine films (right) cast from
chlorobenzene (5 mg/mL).
accomplished by Soxhlet extraction using different solvent sequences,
depending on the polymer solubility. Due to the low volume
fraction of solubilizing alkyl chains in P1, “swallow-tail”-like
branched side chains (1-octylnonyl) were used in this polymer,
which then exhibits acceptable solubility in warm chlorinated
solvents (>5 mg/mL). Note however that incorporation of linear
The results show good reproducibility in the synthesis of these
BTI-based polymers. The number-average molecular weight
(M ) and polydispersity (PDI) were 4.0 kDa and 2.12, respec-
n
tively, for P(BTimR)-L, which is comparable to our previous
60
result of M = 3.6 kDa and PDI = 2.20. By increasing the
n
polymerization temperature from 60 to 80 °C and recovering the
polymer using chloroform as solvent after hexane extraction,
60
chains or branched 2-octyldodecyl chains results in insoluble or
poorly soluble P1 derivatives. P2a has a high volume fraction of
solubilizing chains, and the major fraction can be extracted with
hot chlorobenzene. While polymer P2a shows decent solubility
in hot chlorobenzene, it is barely soluble at room temperature,
which significantly affects the processability and film quality in
the corresponding OFET devices (vide infra). The other polymers
exhibit excellent solubility in chlorinated solvents. The fraction of
P3b used for OFET fabrication can be extracted with hexane due
to the large number of long alkyl substituents. The identity and
purity of all polymers is supported by elemental analysis (EA)
P(BTimR)-H was obtained with M = 7.2 kDa and PDI = 1.98.
n
However, polymerization in DMF and toluene as cosolvents
at 80 °C results in insoluble P(BTimR)-I, which could not be
further processed.
Polymer Optical Properties. The optical absorption spectra
of polymers P1-P3 in solution and as thin films are shown in
Figure 3, and relevant data are summarized in Table 1. All
polymers exhibit large oscillator strengths in the visible region
ranging from λmax = 502 to 576 nm in solution and from λmax
=
558 to 574 nm for the thin films. The corresponding optical band
gaps are estimated from the red absorption edge and fall within a
small range for all the present polymers, both in solution
(1.80-1.81 eV) and in the solid state (1.81-1.85 eV). The
band gaps of P1-P3 are smaller than that of rr-P3HT (1.90-
1
and high-temperature H NMR spectra in 1,1,2,2-tetrachlor-
oethane-d (see Supporting Information for details). The EA
2
demonstrates high purity with error lower than 0.5% for C, H,
and N for all polymers except P2a. The EA of P2a gives 68.12%
for C, 7.99% for H, 1.84% for N, and 3.60% for ash. The ash is due
to the Pd residues from catalyst. Adjusting the numbers to an ash-
free formulation gives 70.66% for C, 8.29% for H, and 1.91% for
N, which is comparable to theoretical values. The reason for
contaminating Pd in the P2a sample is likely due to the low
solubility of P2a at room temperature, which prevents filtration
of the P2a solution through a 0.22 μm PTFE filter used to
remove the Pd particles. However, all other polymer solutions
can be easily filtered through the filter before sending for
elemental analysis. This could lead to the significant amount of
Pd in the P2a sample but not in other samples.
64,65
2.0 eV)
and the BTI homopolymer P(BTimR) (2.02 eV),
which can be attributed to the donor-acceptor interactions in
these BTI-based copolymers. Although the band gaps of the new
polymers are within a small range, it is interesting to note that the
opt
E
s of P2 and P3 are slightly smaller than that of P1 due to the
g
stronger donating ability of the bithiophene donor units of P2
and the tetrathiophene units of P3 vs that of the monothiophene
unit of P1. When considering polymers with the same backbone,
those having shorter side chains have a stronger tendency to
aggregate in solution. In chloroform, the λmax of P2b is 45 nm
blue-shifted in comparison to that of P2a and the λmax of P3b is
37 nm blue-shifted vs that of P3a. However, note that the thin
film absorption spectra are almost identical (Δλmax = 1 nm). On
going from solution to the solid state, these polymers exhibit
different bathochromic shifts which reflect different levels of
backbone planarization and interchain π-π interactions. Although
P(BTimR)-H has a higher molecular weight than P(BTimR)-L,
both batches show identical optical absorption profiles, indicat-
ing that saturation of the conjugation length is already present in
P(BTimR)-L.
In order to investigate the effect of P(BTimR) molecular
weight on film morphology and device performance, different
polymerization procedures and postpurification conditions were
employed by varying reaction temperature, reaction solvent, and
the solvent sequence of Soxhlet extractions. Thus, three batches of
polymers with different molecular weight and processability were
obtained (see Supporting Information): P(BTimR)-L (low molec-
ular weight batch), P(BTimR)-H (high molecular weight batch),
and P(BTimR)-I (insoluble batch). Among them, P(BTimR)-L
was synthesized and purified under the exactly same conditions
The thin films of the present polymer family exhibit both an
absorption maximum and a shoulder at longer wavelength
than in solution, similar to the vibronic profiles seen in regiore-
60
reported by some of us before. Note also that two batches were
synthesized for each of P1-P3 under the same conditions and
that comparable results were achieved (Supporting Information).
2
9
gular polythiophene spectra. This indicates highly ordered
1
409
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
Table 1. Optical and Electrochemical Properties of Polymers P1-P3 vs those of P3HT, P(BTimR), and P(BTimR-BT)
ARTICLE
opt
onset
onset
λ
max abs sol
λ
shoulder abs sol
λ
max abs film
λ
shoulder abs film
E
g
E
ox
E
HOMO
e
E
red
E
LUMO
f
a
a
b
b
c
d
d
polymer
P3HT
(nm)
(nm)
(nm)
(nm)
(eV)
(V)
(eV)
(V)
(eV)
447
535
517
NA
578
607
551
524
569
603
564
615
1.90
2.02
1.91
0.33
1.48
1.08
-5.13
-6.28
-5.88
NA
NA (-3.23)
-3.47
h
P(BTimR)
P(BTimR-
-1.33
-1.76
-3.04
h
BT)
P1
527
576
531
618
638
640
646
558
572
573
574
615
630
632
1.85
1.81
1.82
0.75
0.65
0.64
-5.55
-5.45
-5.44
-1.47
-1.58
-1.54
-3.33
g
(-3.70 )
P2a
P2b
-3.22
g
(-3.64 )
-3.26
g
(-3.63 )
g
P3a
P3b
539
502
633
632
1.82
1.83
0.60
0.58
-5.40
-5.38
NA
NA
NA (-3.58 )
g
649
573
b
NA(3.55 )
a
c
f
-5
Solution absorption spectra (1 ꢀ 10 M in chloroform). Thin film absorption spectra from pristine film cast from 5 mg/mL chlorobenzene solution.
d
þ e
onset
Optical band gap estimated from the absorption edge of the as-cast thin film. Electrochemically determined vs Fc/Fc . E
= -(Eox
þ 4.80).
HOMO
onset
g
opt
h
ELUMO = -(Ered
þ 4.80). Values in parentheses calculated according to ELUMO = E
g
þ EHOMO
.
Data from ref 60 for P(BTimR) and
P(BTimR-BT), in which R = 2-octyldodecyl.
48
three-dimensional solid state structures. Note that thermal
annealing of the pristine films under N at 150 °C does not alter
2
the optical absorption profiles. It is worth comparing the
absorption spectra and optical band gaps of the bithiophene-
containing P2 and P(BTimR-BT) polymers (structures shown
in Figure 2). The band gap of P2 is slightly smaller than that of
opt
60
P(BTimR-BT) (ΔEg = 0.1 eV, Table 1), which implies that
P2 maintains the same backbone conformation as P(BTimR-
BT) and has a greater number of repeat units due to the higher
molecular weight afforded by the enhanced solubility of P2 vs
P(BTimR-BT) and the electron-releasing capacity of the alkyl
1
7
chains on the bithiophene comonomer unit of P2. Since the
optical band gaps of all BTI copolymers and P3HT are similar,
the electron-deficient imide group apparently lowers both the
HOMO and the LUMO energies; this will be investigated by
cyclic voltammetry in the next section. Lower HOMOs are
generally predictors of greater air stability in OFETs.
þ
Figure 4. Cyclic voltammograms of BTI copolymer thin films (Fc/Fc
redox couple was used as internal standard). P3HT is shown for comparison.
electrochemical potentials are reported vs SCE (Figure 4), which
67
has an energy of -4.44 eV below the vacuum level. The HOMO
energies of P1-P3 were determined from the oxidation onset of the
Polymer Electrochemical Properties. It is known that or-
ganic semiconductors must have appropriate HOMO or LUMO
energies to facilitate hole (p-type) or electron (n-type) injection
from the source electrode and to optimize charge transport in the
onset
ox
CV curves and calculated according to eq 1 (E
potential vs Fc/Fc ).
=onsetoxidation
þ
onset
EHOMO ¼ - ðEox
þ 4:80ÞðeVÞ
þ 4:80ÞðeVÞ
ð1Þ
ð2Þ
ð3Þ
5
1
FET channel as well as device environmental stability. The electro-
chemical properties of the BTI-based polymers were investigated
onset
ELUMO ¼ - ðEred
as thin films in anhydrous acetonitrile under N at a scan rate of
0 mV/s using tetrakis(n-butyl)ammonium hexafluoropho-
sphate ((n-Bu) NPF ) as the supporting electrolyte. Platinum
electrodes were used as both the working electrode and the
counter electrode, and a Ag wire was used as the pseudoreference
electrode. Polymer films were drop cast onto the platinum
working electrode from a 0.2% (w/w) chlorobenzene solution. A
ferrocene/ferrocium (Fc/Fc ) redox couple was used as the
internal standard and assigned an absolute energy of -4.80 eV vs
vacuum. The peaks of the Fc/Fc redox couple are relatively
poorly resolved in some traces, which could be due to the film
thickness. The thick film may lead to inefficient penetration of
electrolyte and ferrocene through the polymer films. The poor
resolution of peaks can result in large experimental errors; thus,
the data reported in Table 1 are estimated from at least two runs for
each polymer sample to minimize the experimental errors. All
2
opt
5
ELUMO ¼ Eg þ EHOMOðeVÞ
4
6
All of the present BTI-based copolymers exhibit detectable
quasi-reversible oxidation waves. The onset oxidation potentials
þ
(vs Fc/Fc ) are þ0.75, þ0.65, þ0.64, þ0.60, and þ0.58 V for
P1, P2a, P2b, P3a, and P3b, respectively (Table 1). As expected,
a change of alkyl substituent length has a negligible influence on
the polymer electrochemical properties. However, increasing the
conjugation length of the donor from monothiophene (P1) to
bithiophene (P2) reduces the onset of the oxidation potential
significantly by ∼0.1 V, and further increasing the donor con-
jugation length from bithiophene (P2) to tetrathiophene (P3)
decreases the oxidation potential onset further by ∼0.05 V. The
more facile oxidation on going from P1 to P3 is due to the increased
electron-donating capacity of tetrathiophene vs monothiophene,
þ
6
6
þ
1
410
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
Table 2. Bottom-Gate/Top-Contact (BG/TC) OFET Performance for Polymers P1-P3, P3HT, P(BTimR), and P(BTimR-BT)
Measured in Ambient Conditions
2
-1 -1 c
polymer
casting solvent
DCB
T
annealing (°C)/time (h)
μ (cm V
s
)
V
t
(V)
Ion/Ioff
a
3.6 ꢀ 10 ((2.2 ꢀ 10-4) [h]
-4
3
P1
150/2
-95
100
-25
-60
-10
-5
5 ꢀ 10
-
4
4
4
.6 ꢀ 10 ((6.5 ꢀ 10-5) [e]
10
2.4 ꢀ 10 ((4.8 ꢀ 10-5) [h]
-
4
5 ꢀ 10
2 ꢀ 10
5 ꢀ 10
3
6
4
P2a
P2b
P3a
P3b
CB
150/2
200/2
150/2
150/2
300/2
270/2
120/0.5
240/2
180/0.5
1.7 ꢀ 10 ((1.6 ꢀ 10-4) [h]
-3
CHCl
CHCl
CHCl
CHCl
CHCl
CHCl
CHCl
DCB
3
3
3
3
3
3
3
6.6 ꢀ 10 ((2.0 ꢀ 10-4) [h]
-3
1.6 ꢀ 10 ((1.7 ꢀ 10-3) [h]
-
2
10
4
1.1 ꢀ 10 ((9.8 ꢀ 10-4) [e]
-
2
93
2 ꢀ 10
2 ꢀ 10
4
P(BTimR)-L
P(BTimR)-H
P3HT
3.8 ꢀ 10 ((2.1 ꢀ 10-3) [e]
-
2
68
6
-
2
3
4 ꢀ 10 [h]
-14
75
10
a,b
-2
7
7
P(BTimR)
1.1 ꢀ 10 [e]
2 ꢀ 10
2 ꢀ 10
b
-3
P(BTimR-BT)
8 ꢀ 10 [h]
-12
a
b
c
Measured under vacuum. Data taken from ref 60. h and e indicate hole and electron mobility, respectively. Standard deviation values are shown in
parentheses.
resulting in a higher HOMO energy. On the basis of eq 1, the
calculated HOMO energies are -5.55, -5.45, -5.44, -5.40,
and -5.38 eV for P1, P2a, P2b, P3a, and P3b, respectively. In
comparison to P3HT, which has an oxidation potential onset of
Table 3. OFET Electrical Data for BG/TC Devices
Fabricated from P3b Films Grown on HMDS-Treated
Substrates and Annealed at the Indicated Temperatures
(OFET data measured in ambient)
þ
þ0.33 V vs Fc/Fc and a HOMO energy of -5.13 eV, measured
2
-1 -1 a
annealing temp. (°C)
μ (cm V
s
)
V
t
(V)
Ion/Ioff
under identical conditions, the HOMO energies of all BTI-based
5.5 ꢀ 10 ((1.7 ꢀ 10-4)
-
3
-23
-25
-5
7 ꢀ 10
2 ꢀ 10
4
5
copolymers are substantially lower than that of P3HT by at least
60
-2
0
.25 eV. Thus, incorporating BTI as an electron-withdrawing
120
150
200
250
1.4 ꢀ 10 ((4.6 ꢀ 10-4)
1.6 ꢀ 10 ((1.7 ꢀ 10-3)
-
2
10
4
unit into the polymer backbone should, all other things being
equal, lead to materials with enhanced air stability vs P3HT.
Polymers P1, P2a, and P2b also exhibit two reversible
reductive waves, which indicates that these materials can poten-
tially function as n-type semiconductors. However, note that no
detectable reductive waves are observed for polymer P3. For
polymers P1 and P2 exhibiting reversible reductive waves, the
LUMO energies are calculated from eq 2, whereas for polymer
P3 with no detectable reductive waves, the LUMO energies are
estimated from the HOMO energy and the optical band gap
using the eq 3. Relevant data are collected in Table 1. In com-
parison to the parent homopolymer P(BTimR), which has a
1.5 ꢀ 10 ((8.1 ꢀ 10-4)
-
2
-13
-36
3 ꢀ 10
2 ꢀ 10
4
6.8 ꢀ 10 ((1.2 ꢀ 10-3)
-3
4
a
Standard deviations are shown in parentheses.
The development of mesophases in polymers P3a and P3b may
be associated with the greater microstructural order, which
should also lead to more efficient charge transport. The increas-
ing of molecular weight of homopolymer leads to a more obvious
endothermic transition at higher temperature for P(BTimR)-H
(
Figure S25, Supporting Information).
Field-Effect Transistor Fabrication and Device Character-
60
reduction potential onset of -1.33 eV, P1, P2a, and P2b have
ization. Both bottom-gate/top-contact (BG/TC) and top-
gate/bottom-contact (TG/BC) OFETs were fabricated to in-
vestigate the charge transport properties of the new BTI-based
polymers. For the BG/TC devices, the semiconductor layer was
deposited by spin coating a 5-10 mg/mL polymer solution
þ
reduction potential onsets (vs Fc/Fc ) of -1.47, -1.58, and -
1
.54 V, respectively. The electrochemically derived LUMO
energies are therefore -3.47, -3.33, -3.22, and -3.26 eV for
P(BTimTR), P1, P2a, and P2b, respectively. Therefore, incor-
poration of electron-donating units in the present materials leads
to pronounced destabilization of the LUMO levels.
Polymer Thermal Properties. Thermal analysis of the BTI-
based copolymers was carried out by differential scanning calorim-
etry(DSC) at a temperature ramp rate of 10 °C/min (Figure S24,
Supporting Information). Two heating/cooling cycles were
recorded for each sample to eliminate artifacts arising from residual
(
solvents used: 1,2-dichlorobenzene (DCB) for P1, chloroben-
zene (CB) for P2a, and chloroform for P2b, P3a, and P3b) under
ambient conditions on hexamethyldisilazane (HMDS)-treated,
p-doped Si (001) wafers having a 300 nm thermally grown SiO2
dielectric layer. The capacitance of the 300 nm SiO gate insulator
2
-2
is ∼12 nFcm . Prior to semiconductor deposition, the wafers
were solvent cleaned by immersion in ethanol with sonication
solvent and/or H O. For polymers P1, P2a, and P2b, the DSC
2
and then dried with a filtered stream of N , followed by 5 min
2
thermograms are featureless in the 25-340 °C temperature
range, providing no evidence of mesophase transitions. However,
polymers P3a and P3b exhibit clear thermal transitions during
the heating/cooling cycles located at 287/272 °C for P3a and at
plasma cleaning. Trimethylsilation of the Si/SiO surface was carried
2
out by exposing the silicon wafers to HMDS vapor at room
temperature in a closed air-free container under N . After spin
2
coating, the semiconductor films were annealed under N or in
2
2
56/239 °C for P3b. Interestingly, the polymers functionalized
vacuum at selected temperatures, as summarized in Tables 2 and 3.
OFET devices were completed by vapor deposition of the top
gold electrodes through a shadow mask to define devices with
channel lengths of 25-100 μm and widths of 500-2000 μm.
Device characterization was typically performed under ambient
with longer side chains exhibit lower transition temperatures.
In essence, the side chains have already melted at temperatures
approaching the transition temperatures and act as solvents,
34
inducing the polymer backbones to melt at lower temperatures.
1
411
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
Figure 5. OFET response plots for P1- and P3b-based devices. (a) Output plots as a function of gate bias for P1-based devices measured under vacuum,
b) transfer plots (V = 100 V), and (c) output plots of P3b-based devices measured in ambient conditions.
(
D
conditions in a probe station as described in the Experimental
Section (Supporting Information). Table 2 collects the average
OFET performance for the present polymer series using opti-
mized semiconductor film deposition conditions (at least five devices
were measured for each sample), while Figure 5 shows selected
output and transfer plots for polymer P1 and P3b OFETs
copolymer P1, in which only a single thiophene ring is inserted
into the repeat structure, is the only polymer in this class exhibiting
ambipolar transport in vacuum, due to the relatively low-lying
LUMO (-3.33 eV) compared to the other copolymers in this
series having LUMOs from -3.22 to -3.26 eV. However, the P1
LUMO is not sufficiently low-lying to enable electron transport
51,52
(see Table S1 in Supporting Information for additional data).
in ambient;
thus, the ambipolar behavior of the P1-based
The present family of BTI-based copolymers and analysis of
OFETs is only observed in vacuum. This polymer exhibits very
-
4
2
-1 -1
-4
the OFET response provides valuable information on the effects
of two different structural modifications in BTI systems: (i) sequen-
tial π-core expansion of the oligothiophene donor comonomer,
affecting polymer electronic structure, and (ii) modification of
the alkyl side chain length, which should affect film processability
and crystallinity. The OFET trends summarized in Table 2 are in
excellent agreement with the electrochemical measurements
Table 1). In fact, the HOMO energies extracted from the
CV measurements increase in the order P1 (-5.55 eV) < P2
-5.45 eV) < P3 (-5.39 eV), while the hole field-effect
mobilities follow exactly the same trend P1 < P2 < P3, specifically
similar hole (3.6 ꢀ 10 cm V s ) and electron (4.6 ꢀ 10
-1 -1
2
4
cm V s ) mobilities in vacuum with I /I ratios ∼10 and
on off
8-70
6
thus has potential in complementary circuits.
P-type behavior is
detected under ambient conditions with a hole mobility of
-
4
2
-1 -1
∼4.5 ꢀ 10 cm V s . However, the threshold voltages
measured for this polymer, for both electron and hole transport,
are extremely large for practical applications, which may be due
to the large formal energy barriers for efficient charge injection
(1.67 eV for electron injection and 0.55 eV for hole injection) as
well as trapping sites at the interface between the semiconductor
(
(
15
and the SiO dielectric substrate.
2
-
4
-3
-2
2
-1 -1
μ = 3.6 ꢀ 10 , 1.7 ꢀ 10 , and 1.6 ꢀ 10 cm V
s
for P1,
We also investigated the device performance of homopolymer
P2b, and P3b, respectively. These results indicate that by
inserting electron-donating thiophene rings into the copolymer
architecture, the HOMO energies approach the Fermi level of
the gold electrodes (-5.0 eV), thus facilitating charge injection
and increasing the field-effect hole mobility. The best perfor-
mance for the present BG/TC devices is found for polymer P3b,
P(BTimR) for both M batches. The average electron mobility
n
-
2
2
-1 -1
of P(BTimR)-L = 1.1 ꢀ 10 cm V
s
for the optimized
devices, comparable to our previous results. The high M
n
batch P(BTimR)-H exhibits average electron mobility = 3.8 ꢀ
-
2
2
-1 -1
6
10 cm V
s
, I /I = 2 ꢀ 10 , and a threshold voltage V =
on off
t
þ68 V (Figure 6). Since both batches have identical optical and
electrochemical properties, the significant difference in device
performance must originate from the film morphology, which
will be discussed in the following section.
-
2
2
-1 -1
which exhibits a field-effect mobility of 1.6 ꢀ 10 cm V
s , a
4
threshold voltage of ∼-5 V, and a current on/off ratio of ∼10 .
In the marked contrast to the above results, the LUMO levels
are destabilized upon incorporating an increased number of
thiophene subunits into the present polymers, thus rendering
electron injection less favorable. Lacking any donor unit, homo-
polymer P(BTimR) exhibits n-type behavior with an electron
Polymer Film Microstructures and Morphologies. In order
to better understand the effects of the donor comonomer type
and the various alkyl substituents on the electrical performance of
the present BTI-based polymers, film microstructure and surface
morphology were studied by specular X-ray diffraction (XRD)
2
-1 -1
60
mobility = 0.011 cm V
s
under inert atmosphere. However,
1
412
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
Figure 6. OFET response plots for the high molecular weight batch homopolymer P(BTimR)-H measured under vacuum. (Left) Output plots as a
function of gate bias; (right) transfer plots (V = 150 V).
D
Figure 7. (a) Θ-2Θ X-ray diffraction scans of BTI-based polymer films spin cast onto hydrophobic HMDS-treated substrates, (b) cartoon showing
suggested lamellar stacking for P2a in the polymer thin films, and (c) AFM images of thin films of the present polymers deposited on HMDS-treated
Si/SiO substrates and annealed at 150 (P2a, P3a, and P3b) and 200 °C (P2b). Scale bars correspond to 1 μm.
2
and tapping mode atomic force microscopy (AFM). Figure 7
shows Θ-2Θ XRD scans and AFM images for the present
polymer films under the conditions yielding optimum BG/TC
OFET performance. A single family of Bragg reflections is found
for all the polymers. It is also seen that introducing thiophene
rings into the main chain structure enhances thin film crys-
or a greater degree of alkyl group interdigitation. From the XRD
data in Figure 7a, it is also clear that replacing n-decyl by
n-dodecyl as the side chain enhances the film crystallinity in all
cases, probably because of improved self-assembly driven by side
chain crystallization, which is also in agreement with the d spacings
found for P2a and P2b. The greatest crystallinity is found for
polymer P3b, with Bragg reflections up to third order. Further-
more, AFM images (Figure 7c) also indicate that the greatest crys-
tallinity is for polymer P3b, films of which are characterized by
better-defined grains. These results together with the highest
HOMO energy are in accord with the greater OFET mobility for
P3b versus the other BTI-based polymers.
For the most ordered, highest mobility polymer of this series
P3b, we also investigated how film microstructure and BG/TC
OFET performance evolve as a function of annealing at increas-
ing temperatures. Figure 8 clearly shows that P3b crystallinity is
enhanced when the annealing temperature is increased from 60 to
150 °C and then becomes less ordered as the annealing temperature
33
7
1
tallinity, presumably due to a more ordered π-π stacking of
72
the conjugated cores.
From the Bragg reflections shown in Figure 7, the out-of-plane
d spacings are estimated to be 23.9, 25.2, 25.2, 23.3, and 27.60 Å
for polymers P1, P2a, P2b, P3a, and P3b, respectively. To better
understand this trend, Figure 7b shows a sketch indicating the
possible lamellar stacking of the polymers under study. From this
diagram note that the backbone tilt angle should strongly depend
on the substituent chain length and the extent of alkyl chain
interdigitation. Interestingly, in the case of P2a and P2b, the
d spacings are identical, while the data for P2b having a different
substituent alkyl chain lengths indicate either a different tilt angle
1
413
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
and P(BTimR)-H thin films annealed at the temperatures that
yield the highest electrical performance, 300 and 270 °C,
respectively (see Supporting Information for the AFM images
of the polymer films annealed at other temperatures). It is evident
from both the XRD and the AFM data that the higher M batch
n
polymer, P(BTimR)-H, yields far more crystalline films, exhibit-
ing more intense XRD peaks and better-defined grains in the
AFM images. This trend is in good agreement with electrical
performance where the FET mobility is increased from 1.1 ꢀ
-1 -1
-
2
2
-1 -1
-2
2
1
0
cm V
s
for P(BTimR)-L to 3.8 ꢀ 10 cm V
s
for P(BTimR)-H. These results clearly indicate the importance
of controlling the molecular weight of these polymer semicon-
25,26
ductors to achieve optimum device performance.
Top-Gate/Bottom-Contact Transistors. The semiconduct-
ing properties of polymers P(BTimR)-H, P3a, and P3b were
also characterized in top-gate/bottom-contact (TG/BC) OFET
architectures. These devices were fabricated on glass substrates with
gold source and drain electrodes (30 nm) deposited by thermal
evaporation (L = 50 μm, W = 500 μm). The semiconductor films
were deposited by spin coating 5 mg/mL DCB solutions and
Figure 8. (a) Θ-2Θ X-ray diffraction scans and (b) AFM images of
2
polymer P3b deposited on HMDS-treated Si/SiO substrates and annealed
annealed in a vacuum oven at 110 °C. Either PMMA (poly-
at the indicated temperatures. AFM scale bars correspond to 1 μm.
2
(
methylmethacrylate), capacitance =6.2 nF/cm ) or Polyera Acti-
2
vInk D2200 (polyolefin-polyacrylate, capacitance =5.4 nF/cm )
was spin coated on top of the semiconductor films as the gate
dielectric. Gold was then deposited by thermal evaporation as the
gate electrode. Homopolymer P(BTimR)-H-based devices were
evaluated under vacuum, and transfer plots were recorded at
-
150 V (V ) and V from -10 to 150 V. Copolymer P3a- and
D
G
P3b-based devices were evaluated under ambient, and transfer
plots were recorded at -60 V (V ) and V from 20 to -60 V.
D
G
Transfer plots for P(BTimR)-H, P3a, and P3b using the two
different gate dielectrics are presented in Figure 10, and extracted
OFET data are summarized in Table 4.
The data collected in Table 4 clearly demonstrate the en-
hanced OFET performance of these TG/BC devices vs the
BG/TC OFETs discussed earlier. P(BTimR)-H-based devices
2
-1 -1
Figure 9. (a) Θ-2Θ X-ray diffraction scans of homopolymers
P(BTimR)-L and P(BTimR)-H spin cast onto HMDS-treated sub-
strates and annealed at 300 and 270 °C, respectively (temperatures yielding
the highest electrical performance). (b) AFM image of a P(BTimR)-H film
annealed at 270 °C, and (c) AFM image of a P(BTimR)-L film annealed at
exhibit a maximum electron mobility = 0.19 cm V
on off
s
and
5
I /I = 10 using D2000 as dielectric under vacuum. However,
these OFETs no longer function when exposed to air due to the
high polymer LUMO energies. P3a-based OFETs exhibit similar
2
-1 -1
hole mobilities for both PMMA (μ ≈ 0.061 cm V
s ) and
300 °C. AFM scale bars = 1 μm.
2
-1 -1
D2000 (μ ≈ 0.068 cm V
s ) dielectric materials and I /I
on off
5
ratios ≈ 10 . These mobilities are ∼10 times larger than those of
the corresponding BG/TC OFETs. However, the enhancement
in OFET performance for the P3b-based devices is not as
is increased further. AFM images of the polymer films
Figure 8b) reveal similar evolution of morphologies for anneal-
(
2
-1 -1
ing temperatures ranging from 60 to 200 °C, which are char-
acterized by the presence of small grains. In contrast, the AFM
image of a film annealed at 250 °C shows no grains, in agreement
with the lower crystallinity found in the XRD experiments
significant, with average mobilities = 0.026 cm V
s
using
2
-1 -1
PMMA as the dielectric and 0.064 cm V
s
using D2200 as
the dielectric. The 3-fold lower mobility in P3b/PMMA devices
versus P3b/D2200 devices can be explained by the poor film
uniformity observable in the latter films under the optical
microscope. The performance enhancement in TG/BC OFET
devices has also been demonstrated for other polymeric
semiconductors and offers several advantages over the tradi-
tional bottom-gate/bottom-contact (BG/BT) OFET configura-
tion: (i) lower contact resistance due to reduction of current
crowding effects in staggered contact structures; (ii) wider
selection of gate dielectric materials to minimize charge trapping
at the dielectric/semiconductor interfaces; (iii) better device
air stability afforded by encapsulation effects provided by the
(Figure 8a). Table 3 summarizes OFET performance for P3b-
based devices as a function of the annealing temperature. Note
that the OFET data mirror the XRD/AFM variations and micro-
structural trends. Thus, intermediate annealing temperatures
14
-
2
(
120-200 °C) result in optimum performance with μ > 10
2 -1 -1
cm V
s , while lower performance is observed for P3b films
7
3
annealed at both 60 and 250 °C. For P(BTim-R)-H homo-
polymer films, annealing at 270 °C yields optimum OFET
response, in agreement with the morphology evolution indicated
by XRD and AFM (Figure S26, Supporting Information).
In order to better understand the different device performances of
74
74
overlying gate dielectrics and gate electrodes; (iv) favorable
architecture for printed electronics in which every component
the P(BTimR) homopolymers having different M s, Figure 9
n
18
compares the XRD spectra and AFM images for P(BTimR)-L
can be applied by printing technology.
1
414
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
G
Figure 10. OFET transfer plots of current vs V for representative top-gate/bottom-contact P(BTimR)-H, P3a, and P3b devices with PMMA and
D2200 as the indicated polymeric dielectrics.
Table 4. OFET Electrical Data for Polymers P(BTimR)-H, P3a, and P3b TG/BC Devices Using PMMA and ActivInk D2200 as
the Gate Dielectrics
PMMA dielectric
D2200 dielectric
2
-1 -1 a
2
-1 -1 a
polymer
μ (cm V
s
)
I
on/Ioff
V
t
(V)
μ (cm V
s
)
Ion/Ioff
t
V (V)
4
5
P(BTimR)-H
P3a
0.013 (0.017)
0.061 (0.067)
0.026 (0.031)
10
10
61
0.141 (0.189)
0.068 (0.077)
0.064 (0.067)
10
76
5
5
-18
-21
10
-19
-17
4
5
P3b
5 ꢀ 10
10
a
Data in parentheses are the highest measured mobilities.
It is also instructive to compare the device performance of P3
with that of P3HT in TG/BC structures. P3HT exhibits hole
2
-1 -1
2
3
mobilities ≈ 0.02-0.1 cm V
s
and I /I = 10 -10 in
on off
TG/BC devices using Au/PMMA and Au/D2200 as contact/
dielectric materials. The hole mobilities of P3a and P3b are
comparable to that of P3HT under the same conditions. How-
14
ever, the I /I ratio for devices fabricated from P3HT is ca.
on off
2
-3 orders of magnitude lower than that of devices fabricated
from P3, which reflects the lower-lying P3 HOMO.
Transistor Performance Stability. For n-type polymers,
device air stability is mainly governed by electron affinities.
Clearly, the BTI subunit is not sufficiently electron withdrawing
to stabilize electron transport for both the n-type and the ambipolar
polymers reported here. For semiconducting p-type polymers,
75
OFET air stability is largely controlled by the ionization potentials.
Low ionization potentials usually result in oxidative doping by
O , which erodes OFET performance by decreasing the I /I
2
on off
ratio due to increased off currents and threshold voltage shifts to
4
9,76
positive values.
As seen in Figure 4 and Table 1, introduction
of the BTI subunit in conjugated polymers increases the oxida-
tion potentials of the copolymers as measured by cyclic voltam-
metry with respect to that of P3HT by at least 0.25 eV. This
translates into lower-lying HOMO levels and consequently higher
ionization potentials. This important characteristic makes the
copolymers studied in this contribution far more resistant to O2
doping than P3HT, which is demonstrated by analysis of OFET
performance as a function of time.
Figure 11. (a) Transfer plot (V = 100 V) for a fresh P3b-based device
D
(solid line) and after 9 months storage in air (dashed line). Temporal
evolution of OFET performance in air: (b) carrier mobility, (c) thresh-
old voltage, and (d) Ion/Ioff ratio for the present BTI-based polymers.
Average values are shown.
4
4
ratio, which slightly increases from 10 to 2 ꢀ 10 due to the
Figure 11 shows transfer plots for P3b-based BG/TC OFETs
measured immediately after device fabrication and after storage
for 9 months under ambient conditions and also temporal evolution
of P3b-based OFET ambient characteristics. From these plots, an
decreased off current, which may originate from degradation of
7
7,78
unintentionally p-doped polymer radical cations in air.
On
storing the P3b-based FETs in ambient up to 9 months, the hole
-
2
2
-1 -1
mobility is slightly decreased to 1.1 ꢀ 10 cm V
s
with an
-
2
2
-1 -1
identical field-effect mobility of 1.6 ꢀ 10 cm V
s
and a
identical Ion/Ioff ratio and similar V
t
≈ -9 V. These results
threshold voltage of ∼-5 V are extracted for P3b-based devices
underscore the impressive stability of P3b since in the case of air
instability, a decreased Ion/Ioff ratio is expected due to oxidative
after 2 months ambient exposure. The only difference is the I /I
on off
1
415
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
Figure 12. (a) Schematic layout of the complementary inverters fabricated from polymers P3b and P(BTimR)-H and optical image of the inkjet-
printed devices, (b) static switching characteristics of an inverter, and (c) gains of the corresponding inkjet-printed inverter.
14,69,70
O doping, which would increase the off current. Figure 11 also
2
inverters.
The inverting voltage (V ) is slightly shifted in
IN
shows the temporal evolution of the OFET ambient character-
istics for P2a-, P2b-, and P3a-based devices. All BG/TC devices
fabricated from these materials are robust and air stable, with only
the positive direction with respect to 1/2 VDD due to slightly
different V s and charge carrier mobilities between the p- and the
t
n-channel materials.
slightly lower hole mobility, the same order I /I ratio, and
on off
similar V after 200 days storage. The only exception is those
t
’
CONCLUSIONS
based on P2b, which exhibit a V shift of þ30 V. The exact reason
t
for the large V shift is unknown and cannot be only explained by
A new series of BTI-oligothiophene copolymers was synthe-
t
its HOMO energy, since the most structurally similar polymer
sized and characterized. These systems contain electron-deficient
BTI subunits and oligothiophenes having sequentially varied
conjugation lengths as electron-rich cosubunits. The resulting
copolymers exhibit similar band gaps (∼1.80 eV); however, the
energies of the frontier molecular levels (FMOs) vary based
on the donor/acceptor interactions, which greatly affect OFET
performance. As the conjugation length of donor blocks increases
from monothiophene (P1) to bithiophene (P2) to tetrathio-
phene (P3), the p-type behavior becomes more pronounced and
the holes become more mobile, resulting in greater hole mobi-
lities. Compared to n-type homopolymer P(BTimR), insertion
of monothiophene subunits results in ambipolar P1 with ba-
P2a-based devices exhibit essentially no V shift over the same
t
storage period. For the present TG/BC devices, comparable air
stabilities are again observed due to the low HOMO energies of
these materials as well as the encapsulation effects of the TG/BC
7
4
architecture. In fact, field-effect mobilities remain basically
constant over 60 days of air exposure. As an example, TG/BC2
P3a-based devices exhibit a field-effect mobility of ∼0.06 cm
-
1
-1
V
s
, a threshold voltage of ∼-24 V, and an I /I ratio
on off
5
of ∼10 when D2200 is used as the gate dielectric. When storage
2
-1 -1
was extended to 9 months, the hole mobility (0.06 cm V
s )
and threshold voltage (-24 V) also remain unchanged; however,
-
4
2
-1 -1
the I /I ratio is decreased by 10ꢀ (see the transfer plots in
on off
lanced hole and electron mobilities of ∼10 cm V
s .
Figure S28, Supporting Information).
Further increasing the oligothiophene donor conjugation length
-
3
Inkjet-Patterned Polymeric CMOS Inverters. Complemen-
tary inverters were fabricated by inkjet printing/patterning the
present semiconducting polymers (see Supporting Information for
detailed procedure). For monolithic integration of the p- and
n-channel OFETs, P3b and P(BTimR)-H were chosen as the
p-type and n-type semiconductors, since they show the highest
hole and electron mobilities in this series, respectively. Their
solutions were sequentially inkjet-printed onto photolithography
patterned Au bottom-contact electrodes. The polymer gate
dielectric PMMA was then spin coated on the conjugated poly-
mers for a top-gated geometry. Figure 12 shows the voltage
transfer characteristics (VTCs) of the resulting complementary
inverter (p-channel, W/L = 5.0 mm/10 μm; n-channel, W/L =
in P2 and P3 yields p-type response with hole mobilities of 10
2 2
-
-1 -1
and 10 cm V
s , respectively. We also investigated the
influence of P(BTimR) homopolymer molecular weight on the
film morphology and device performance. The high molecular
weight batch, P(BTimR)-H, exhibits far greater performance
than the low molecular weight one. Through device optimization
using a TG/BC OFET architecture, P(BTimR)-H exhibits a
2
-1 -1
high electron mobility of ∼0.2 cm V
s
in vacuum. Device
optimization by employing TG/BC OFET architectures leads to
hole mobilities approaching 0.1 cm V
2
-1 -1
s
for P3a and P3b,
comparable to that of P3HT devices fabricated under identical
conditions. Furthermore, both the present BG/TC and TG/BC
devices fabricated from P3a and P3b are remarkably air stable,
1
.0 mm/10 μm) at various supply voltages (VDD). The static
and key OFET performance parameters (hole mobility, I /I
on off
inverter characteristics show negligible bias hysteresis and high
ratio and V ) exhibit negligible changes after storing the devices
t
voltage gains (∼ 40) at V = -100 V; the voltage gain reported
DD
under ambient conditions for 200 m days. Finally, complemen-
tary inverters fabricated from P3b and P(BTimR)-H OFETs
here is comparable to that of current state-of-the-art organic
1
416
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
achieve a high voltage gain of ∼40. The present results demon-
strate enhanced charge carrier mobility by optimizing polymer
semiconductor molecular weight, tunability of charge carrier
polarity, and achieving high hole mobility with good device air
stability by adjusting frontier molecular orbital energies and
optimizing film microstructures.
(13) Hoth, C. N.; Choulis, S. A.; Schilinsky, P.; Brab ꢀe c, C. J. Adv.
Mater. 2007, 19, 3973–3978.
(14) Yan, H.; Chen, Z.; Zheng, Y.; Newman, C.; Quinn, J. R.; Dotz,
F.; Kastler, M.; Facchetti, A. Nature 2009, 457, 679–686.
(
(
15) Zaumseil, J.; Sirringhaus, H. Chem. Rev. 2007, 107, 1296–1323.
16) Sirringhaus, H.; Kawase, T.; Friend, R. H.; Shimoda, T.;
Inbasekaran, M.; Wu, W.; Woo, E. P. Science 2000, 290, 2123–2126.
(
17) McCulloch, I.; Heeney, M.; Chabinyc, M. L.; DeLongchamp,
’
ASSOCIATED CONTENT
D.; Kline, R. J.; Coelle, M.; Duffy, W.; Fischer, D.; Gundlach, D.;
Hamadani, B.; Hamilton, R.; Richter, L.; Salleo, A.; Shkunov, M.;
Sporrowe, D.; Tierney, S.; Zhong, W. Adv. Mater. 2009, 21, 1091–1109.
(18) Allard, S.; Forster, M.; Souharce, B.; Thiem, H.; Scherf, U.
Angew. Chem., Int. Ed. 2008, 47, 4070–4098.
(19) Murphy, A. R.; Fr ꢀe chet, J. M. J. Chem. Rev. 2007, 107, 1066–
096.
(
(
S
Supporting Information. Monomer/polymer synthesis
b
and characterization; detailed procedure for fabrication and
1
13
characterization of complementary inverters; H and C NMR
spectra of all monomers; H NMR spectra and DSC curves of all
polymers; XRD and AFM images of polymer films; overlapped
transfer plots of P3a-based TG/BC OFETs after different
storage time; OFET device data for different deposition condi-
tions of polymers. This material is available free of charge via the
Internet at http://pubs.acs.org.
1
1
20) Anthony, J. E. Chem. Rev. 2006, 106, 5028–5048.
21) Osaka, I.; McCullough, R. D. Acc. Chem. Res. 2008, 41, 1202–
1
214.
22) Liu, S.; Wang, W.; Briseno, A. L.; Mannsfeld, S. C. B.; Bao, Z.
Adv. Mater. 2009, 21, 1217–1232.
23) Payne, M. M.; Parkin, S. R.; Anthony, J. E.; Kuo, C. C.; Jackson,
T. N. J. Am. Chem. Soc. 2005, 127, 4986–4987.
(
(
’
AUTHOR INFORMATION
(
(
24) Anthony, J. E. Angew. Chem., Int. Ed. 2008, 47, 452–483.
25) Kline, R. J.; McGehee, M. D.; Kadnikova, E. N.; Liu, J.; Fr ꢀe chet,
Corresponding Author
a-facchetti@northwestern.edu; t-marks@northwestern.edu
J. M. J. Adv. Mater. 2003, 15, 1519–1522.
26) Kline, R. J.; McGehee, M. D.; Kadnikova, E. N.; Liu, J.; Fr ꢀe chet,
J. M. J.; Toney, M. F. Macromolecules 2005, 38, 3312–3319.
27) Ahmed, E.; Kim, F. S.; Xin, H.; Jenekhe, S. A. Macromolecules
(
’
ACKNOWLEDGMENT
(
We thank ONR (N00014-05-1-0766), AFOSR (FA9550-08-
2009, 42, 8615–8618.
1
-0331), and Polyera Corp. for support of this research and the
(28) McCullough, R. D.; Tristramnagle, S.; Williams, S. P.; Lowe,
R. D.; Jayaraman, M. J. Am. Chem. Soc. 1993, 115, 4910–4911.
NSF-MRSEC program through the Northwestern University
Materials Research Science and Engineering Center for char-
acterization facilities (DMR-0520513). The research leading
to these results has also received funding from the European
Community's Seventh Framework Programme under Grant
Agreement 234808. Research fulfilled at Hanbat National Uni-
versity and ETRI was supported by Development of Next Generation
RFID Technology for item-level applications (2008-F052-01)
funded by the Ministry of Knowledge Economy (MKE) of Korea
and Korea Basic Science Research Program through the National
Research Foundation of Korea (NRF) funded by the Ministry of
Education, Science and Technology (2010-0023180). We thank
Dr. S. Lu of Polyera Corp. for GPC measurements.
(
29) McCullough, R. D.; Williams, S. P. J. Am. Chem. Soc. 1993, 115,
1608–11609.
30) Bao, Z.; Dodabalapur, A.; Lovinger, A. J. Appl. Phys. Lett. 1996,
9, 4108–4110.
31) Sirringhaus, H.; Tessler, N.; Friend, R. H. Science 1998, 280,
741–1744.
32) Cho, J. H.; Lee, J.; Xia, Y.; Kim, B.; He, Y.; Renn, M. J.; Lodge,
1
6
1
(
(
(
T. P.; Frisbie, C. D. Nat. Mater. 2008, 7, 900–906.
(33) Sauve, G.; Javier, A. E.; Zhang, R.; Liu, J.; Sydlik, S. A.;
Kowalewski, T.; McCullough, R. D. J. Mater. Chem. 2010, 20, 3195–
3
201.
34) Osaka, I.; Zhang, R.; Sauve, G.; Smilgies, D. M.; Kowalewski, T.;
McCullough, R. D. J. Am. Chem. Soc. 2009, 131, 2521–2529.
35) Lim, E.; Jung, B. J.; Lee, J.; Shim, H. K.; Lee, J. I.; Yang, Y. S.; Do,
L. M. Macromolecules 2005, 38, 4531–4535.
36) Lu, G.; Usta, H.; Risko, C.; Wang, L.; Faccehtti, A.; Ratner,
(
(
’
REFERENCES
(
(
(
(
1) Forrest, S. R. Nature 2004, 428, 911–918.
2) Facchetti, A. Mater. Today 2007, 10, 28–37.
3) G €u nes, S.; Neugebauer, H.; Sariciftci, N. S. Chem. Rev. 2007, 107,
M. A.; Marks, T. J. J. Am. Chem. Soc. 2008, 130, 7670–7685.
(37) Zhang, M.; Tsao, H. N.; Pisula, W.; Yang, C.; Mishra, A. K.;
M €u llen, K. J. Am. Chem. Soc. 2007, 129, 3472–3473.
(38) McCulloch, I.; Heeney, M.; Bailey, C.;Genevicius, K.; MacDonald,
I.; Shkunov, M.; Sparrowe, D.; Tierney, S.; Wagner, R.; Zhang, W.;
Chabinyc, M. L.; Kline, R. J.; McGehee, M. D.; Toney, M. F. Nat. Mater.
2006, 5, 328–333.
1
324–1338.
4) Arias, A. C.; MacKenzie, J. D.; McCulloch, I.; Rivnay, J.; Salleo, A.
Chem. Rev. 2010, 110, 3–24.
5) Friend, R. H.; Gymer, R. W.; Holmes, A. B.; Burroughes, J. H.;
(
(
Marks, R. N.; Taliani, C.; Bradley, D. D. C.; Dos Santos, D. A.; Br ꢀe das,
J. L.; Logdlund, M.; Salaneck, W. R. Nature 1999, 397, 121–128.
(39) Ong, B. S.; Wu, Y.; Liu, P.; Gardner, S. J. Am. Chem. Soc. 2004,
126, 3378–3379.
(
(
6) Lo, S. C.; Burn, P. L. Chem. Rev. 2007, 107, 1097–1116.
7) Dimitrakopoulos, C. D.; Malenfant, P. R. L. Adv. Mater. 2002, 14,
(40) Liu, J.; Zhang, R.; Sauve, G.; Kowalewski, T.; McCullough,
R. D. J. Am. Chem. Soc. 2008, 130, 13167–13176.
9
9–117.
(41) Kim, D. H.; Lee, B.-L.; Moon, H.; Kang, H. M.; Jeong, E. J.;
Park, J.-I.; Han, K.-M.; Lee, S.; Yoo, B. W.; Koo, B. W.; Kim, J. Y.; Lee,
W. H.; Cho, K.; Becerril, H. A.; Bao, Z. J. Am. Chem. Soc. 2009, 131,
6124–6132.
(42) Guo, X.; Kim, F. S.; Jenekhe, S. A.; Watson, M. D. J. Am. Chem.
Soc. 2009, 131, 7206–7207.
(43) Wang, Y.; Watson, M. D. Macromolecules 2008, 41, 8643–8647.
(44) Irvin, J. A.; Schwendeman, I.; Lee, Y.; Abboud, K. A.; Reynolds,
J. R. J. Polym. Sci., Part A: Polym. Chem. 2001, 39, 2164–2178.
(45) Guo, X.; Watson, M. D. Org. Lett. 2008, 10, 5333–5336.
(
(
8) Wen, Y.; Liu, Y. Adv. Mater. 2010, 22, 1331–1345.
9) Cheng, Y.-J.; Yang, S.-H.; Hsu, C.-S. Chem. Rev. 2009, 109,
5
868–5923.
10) Thompson, B. C.; Fr ꢀe chet, J. M. J. Angew. Chem., Int. Ed. 2008,
7, 58–77.
11) Argun, A. A.; Aubert, P. H.; Thompson, B. C.; Schwendeman,
(
4
(
I.; Gaupp, C. L.; Hwang, J.; Pinto, N. J.; Tanner, D. B.; MacDiarmid,
A. G.; Reynolds, J. R. Chem. Mater. 2004, 16, 4401–4412.
(12) Beaujuge, P. M.; Reynolds, J. R. Chem. Rev. 2010, 110, 268–320.
1
417
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418

Journal of the American Chemical Society
ARTICLE
(
46) Babel, A.; Jenekhe, S. A. J. Am. Chem. Soc. 2003, 125, 13656–
3657.
47) Tsao, H. N.; Cho, D.; Andreasen, J. W.; Rouhanipour, A.;
Breiby, D. W.; Pisula, W.; M €u llen, K. Adv. Mater. 2009, 21, 209–212.
48) Murphy, A. R.; Liu, J.; Luscombe, C.; Kavulak, D.; Fr ꢀe chet,
J. M. J.; Kline, R. J.; McGehee, M. D. Chem. Mater. 2005, 17, 4892–4899.
49) Meijer, E. J.; Detcheverry, C.; Baesjou, P. J.; van Veenendaal, E.;
de Leeuw, D. M.; Klapwijk, T. M. J. Appl. Phys. 2003, 93, 4831–4835.
50) Anthopoulos, T. D.; Anyfantis, G. C.; Papavassiliou, G. C.; de
Leeuw, D. M. Appl. Phys. Lett. 2007, 90, 122105.
51) Jones, B. A.; Facchetti, A.; Wasielewski, M. R.; Marks, T. J.
J. Am. Chem. Soc. 2007, 129, 15259–15278.
52) Usta, H.; Risko, C.; Wang, Z.; Huang, H.; Deliomeroglu, M. K.;
Zhukhovitskiy, A.; Facchetti, A.; Marks, T. J. J. Am. Chem. Soc. 2009, 131,
586–5608.
53) de Leeuw, D. M.; Simenon, M. M. J.; Brown, A. R.; Einerhand,
R. E. F. Synth. Met. 1997, 87, 53–59.
54) Thompson, B. C.; Kim, Y. G.; McCarley, T. D.; Reynolds, J. R.
J. Am. Chem. Soc. 2006, 128, 12714–12725.
55) Gao, X.; Di, C.; Hu, Y.; Yang, X.; Fan, H.; Zhang, F.; Liu, Y.; Li,
H.; Zhu, D. J. Am. Chem. Soc. 2010, 132, 3697–3699.
56) Gs €a nger, M.; Oh, J.; K €o nemann, M.; H €o ffken, H.; Krause,
A.-M.; Bao, Z.; W €u rthner, F. Angew. Chem., Int. Ed. 2010, 49, 740–743.
57) Zhu, Y.; Champion, R. D.; Jenekhe, S. A. Macromolecules 2006,
9, 8712–8719.
58) Crouch, D. J.; Skabara, P. J.; Lohr, J. E.; McDouall, J. J. W.;
(78) Schmidt, R.; Oh, J. H.; Sun, Y.; Deppisch, M.; Krause, A. M.;
Radacki, K.; Braunschweig, H.; Konemann, M.; Erk, P.; Bao, Z.;
W €u rthner, F. J. Am. Chem. Soc. 2009, 131, 6215–6228.
1
(
(
(
(
(
(
5
(
(
(
(
(
3
(
Heeney, M.; McCulloch, I.; Sparrowe, D.; Shkunov, M.; Coles, S. J.;
Horton, P. N.; Hursthouse, M. B. Chem. Mater. 2005, 17, 6567–6578.
(
59) McCulloch, I.; Bailey, C.; Giles, M.; Heeney, M.; Love, I.; Shkunov,
M.; Sparrowe, D.; Tierney, S. Chem. Mater. 2005, 17, 1381–1385.
60) Letizia, J. A.; Salata, M. R.; Tribout, C. M.; Facchetti, A.; Ratner,
M. A.; Marks, T. J. J. Am. Chem. Soc. 2008, 130, 9679–9694.
61) Fong, H. H.; Pozdin, V. A.; Amassian, A.; Malliaras, G. G.;
(
(
Smilgies, D. M.; He, M.; Gasper, S.; Zhang, F.; Sorensen, M. J. Am. Chem.
Soc. 2008, 130, 13202–13203.
(62) Osaka, I.; Abe, T.; Shinamura, S.; Miyazaki, E.; Takimiya, K.
J. Am. Chem. Soc. 2010, 132, 5000–5001.
(
(
63) Zagorska, M.; Krische, B. Polymer 1990, 31, 1379–1383.
64) Koster, L. J. A.; Mihailetchi, V. D.; Blom, P. W. M. Appl. Phys.
Lett. 2006, 88, 243502.
65) Wu, P. -T.; Xin, H.; Kim, F. S.; Ren, G.; Jenekhe, S. A.
Macromolecules 2009, 42, 8817–8826.
66) Pommerehne, J.; Vestweber, H.; Guss, W.; Mahrt, R. F.;
B €a ssler, H.; Porsch, M.; Daub, J. Adv. Mater. 1995, 7, 551–554.
(
(
(
(
67) Meerholz, K.; Heinze, J. Electrochim. Acta 1996, 41, 1839–1854.
68) Meijer, E. J.; de Leeuw, D. M.; Setayesh, S.; van Veenendaal, E.;
Huisman, B. H.; Blom, P. W. M.; Hummelen, J. C.; Scherf, U.; Kadam, J.;
Klapwijk, T. M. Nat. Mater. 2003, 2, 834–834.
(69) Bijleveld, J. C.; Zoombelt, A. P.; Mathijssen, S. G. J.; Wienk,
M. M.; Turbiez, M.; de Leeuw, D. M.; Janssen, R. A. J. J. Am. Chem. Soc.
2
009, 131, 16616–16617.
70) Kim, F. S.; Guo, X.; Watson, M. D.; Jenekhe, S. A. Adv. Mater.
010, 22, 478–482.
71) Durban, M. M.; Kazarinoff, P. D.; Luscombe, C. K. Macromolecules
010, 43, 6348–6352.
72) Osaka, I.; Zhang, R.; Liu, J.; Smilgies, D.-M.; Kowalewski, T.;
McCullough, R. D. Chem. Mater. 2010, 22, 4191–4196.
(
2
(
2
(
(73) Street, R. A.; Salleo, A. Appl. Phys. Lett. 2002, 81, 2887–2889.
(74) Chen, Z.; Lemke, H.; Albert-Seifried, S.; Caironi, M.; Nielsen,
M. M.; Heeney, M.; Zhang, W.; McCulloch, I.; Sirringhaus, H. Adv.
Mater. 2010, 22, 2371–2375.
(75) Zhang, W.; Smith, J.; Watkins, S. E.; Gysel, R.; McGehee, M.;
Salleo, A.; Kirkpatrick, J.; Ashraf, S.; Anthopoulos, T.; Heeney, M.;
McCulloch, I. J. Am. Chem. Soc. 2010, 132, 11437–11439.
(
76) Bao, Z.; Lovinger, A. J.; Dodabalapur, A. Appl. Phys. Lett. 1996,
9, 3066–3068.
77) Oh, J. H.; Liu, S.; Bao, Z.; Schmidt, R.; W €u rthner, F. Appl. Phys.
Lett. 2007, 91, 212107.
6
(
1
418
dx.doi.org/10.1021/ja107678m |J. Am. Chem. Soc. 2011, 133, 1405–1418