Complementary Semiconducting Polymer Blends: The Influence of Conjugation-Break Spacer Length in Matrix Polymers

Article abstract of DOI:10.1021/acs.macromol.6b00050

The concept of complementary semiconducting polymer blends (c-SPBs) for efficient charge transport was recently proposed and established by our group. In this study, we aim to reveal the influence of the length of conjugation-break spacers (CBSs) on charge transport properties of the matrix polymers and their corresponding complementary polymer blends. A series of 11 DPP-based semiconducting polymers DPP-Cm (m = 2-12) that incorporate CBSs of 2-12 methylene units along the polymer backbones were prepared and characterized. The UV-vis spectra and the ultraviolet photoelectron spectroscopy (UPS) measurements show that the CBS length has marginal influence on the polymer absorption spectra, energy levels, and band gaps. It also has little impact on polymer decomposition temperatures. However, the CBS length has a profound influence on polymer phase transition and the heat of fusion. As for the melt transitions, an odd-even effect is observed from DPP-C2 to DPP-C7, in which polymers with even-numbered CBSs show higher melting points than their adjacent odd-numbered derivatives. The trend is opposite for heat of fusion. The polymers with odd-numbered CBSs exhibit larger heat of fusion, indicating higher ordering and crystallinity. The odd-even effect is also found in surface morphologies of the polymers by atomic force microscopy (AFM). The polymers with the even CBSs have a more interconnected feature that appear more fibrillar than the polymers with the odd linkages. As far as charge carrier mobility is concerned, the average number drops from 0.023 cm² V^-1 s^-1 to 7.9 × 10^-6 cm² V^-1 s^-1 as the CBS moves from C2 to C12. It is intriguing to observe that even-numbered polymers outperform the adjacent odd-numbered polymers, despite the fact that the latter show higher ordering and crystallinity in thin films. When these polymers are mixed with fully conjugated DPP-C0 (2 wt %, designated as tie chain polymer), the obtained c-SPBs witness a dramatic increase (2-4 orders of magnitude) in charge carrier mobility. Interestingly, the odd-even effect is not found for charge transport in the c-SPBs. This work reveals that the length of CBSs plays a significant role in charge transport properties of the matrix polymers and reconfirms that efficient charge transport properties of the c-SPB result from the interactions between matrix polymers and tie chain polymers. This begins to provide guidelines as to what spacer lengths may be utilized to offer the best balance between processing and charge transport properties.

Full text of DOI:10.1021/acs.macromol.6b00050

Article
pubs.acs.org/Macromolecules
Complementary Semiconducting Polymer Blends: The Influence of
Conjugation-Break Spacer Length in Matrix Polymers
Xikang Zhao,† Yan Zhao,† Qu Ge,§ Kamal Butrouna,∥ Ying Diao,§ Kenneth R. Graham,∥
and Jianguo Mei*,†,‡
‡
†Department of Chemistry and Birck Nanotechnology Center, Purdue University, West Lafayette, Indiana 47907, United States
§Department of Chemical & Biomolecular Engineering, University of Illinois at Urbana−Champaign, Urbana, Illinois 61801, United
States
∥Department of Chemistry, University of Kentucky, Lexington, Kentucky 40506, United States
S
* Supporting Information
ABSTRACT: The concept of complementary semiconducting polymer blends (c-SPBs) for efficient charge transport was
recently proposed and established by our group. In this study, we aim to reveal the influence of the length of conjugation-break
spacers (CBSs) on charge transport properties of the matrix polymers and their corresponding complementary polymer blends.
A series of 11 DPP-based semiconducting polymers DPP-Cm (m = 2−12) that incorporate CBSs of 2−12 methylene units along
the polymer backbones were prepared and characterized. The UV−vis spectra and the ultraviolet photoelectron spectroscopy
(UPS) measurements show that the CBS length has marginal influence on the polymer absorption spectra, energy levels, and
band gaps. It also has little impact on polymer decomposition temperatures. However, the CBS length has a profound influence
on polymer phase transition and the heat of fusion. As for the melt transitions, an odd−even effect is observed from DPP-C2 to
DPP-C7, in which polymers with even-numbered CBSs show higher melting points than their adjacent odd-numbered
derivatives. The trend is opposite for heat of fusion. The polymers with odd-numbered CBSs exhibit larger heat of fusion,
indicating higher ordering and crystallinity. The odd−even effect is also found in surface morphologies of the polymers by atomic
force microscopy (AFM). The polymers with the even CBSs have a more interconnected feature that appear more fibrillar than
the polymers with the odd linkages. As far as charge carrier mobility is concerned, the average number drops from 0.023 cm2 V−1
s−1 to 7.9 × 10−6 cm2 V−1 s−1 as the CBS moves from C2 to C12. It is intriguing to observe that even-numbered polymers
outperform the adjacent odd-numbered polymers, despite the fact that the latter show higher ordering and crystallinity in thin
films. When these polymers are mixed with fully conjugated DPP-C0 (2 wt %, designated as tie chain polymer), the obtained c-
SPBs witness a dramatic increase (2−4 orders of magnitude) in charge carrier mobility. Interestingly, the odd−even effect is not
found for charge transport in the c-SPBs. This work reveals that the length of CBSs plays a significant role in charge transport
properties of the matrix polymers and reconfirms that efficient charge transport properties of the c-SPB result from the
interactions between matrix polymers and tie chain polymers. This begins to provide guidelines as to what spacer lengths may be
utilized to offer the best balance between processing and charge transport properties.
even more than 10 cm2 V−1 s−1 in many cases, are now being
INTRODUCTION
■
reported for polymeric semiconductors.10−12 Considering that
With their technological relevance in flexible, soft, stretchable,
implantable, and printed electronic devices, semiconducting
the first solution-processed conjugated polymers, regiorandom
polythiophenes, only gave charge carrier mobilities of 10−5 cm2
polymers have attracted tremendous interest from a wide range
V−1 s−1,13 these latest achievements are truly encouraging and
of disciplines.1−5 One of the major efforts across disciplinary
boundaries is to understand and improve charge transport
properties of organic semiconductors, as efficient charge
transport is crucial to all electronic devices.6−9 Resulting from
these efforts, charge carrier mobilities over 1 cm2 V−1 s−1, and
Received: January 9, 2016
Revised: March 10, 2016
© XXXX American Chemical Society
A
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Scheme 1. Synthetic Routes for the Monomers and Polymers (DPP-Cm, m = 2−12)
exciting. The puzzling fact is that these high-mobility donor−
acceptor (D−A) type polymers exhibit relatively poor
crystallinity, as revealed by X-ray diffraction, while most
theories suggest that high crystallinity and long-range ordering
are required for efficient charge transport in polymer thin
films.14−16 These theories drove the development of a series of
highly crystalline polythiophenes with higher charge-carrier
mobilities, such as the most widely known regioregular poly(3-
hexylthiophene) (rr-P3HT)1 and poly(2,5-bis(3-hexadecyl-
thiophen-2-yl)thieno[3,2-b]thiophene) (PBTTT).17 A rising
effort is now devoted to the reconciliation of this conflict
between solid-state order and charge-carrier mobility. For
instance, it has been recently proposed that charge transport in
high-mobility semiconducting polymers is quasi-one-dimen-
sional (1D),18 predominantly occurring along the backbone.
This requires only occasional intermolecular hopping through
short π-stacking bridges. Hence, a planar, torsion-free polymer
backbone with a low degree of energetic disorder is crucial for
high charge carrier mobilities in highly disordered D−A
polymers.19 In addition, Salleo and co-workers proposed a
unified model (tie-chain model) of how charge carriers travel in
polymer thin films.15 They argued that the limiting step for
charge transport is trapping, caused by lattice disorder, and that
short-range intermolecular aggregation is sufficient for efficient
long-range charge transport. In this model, it is intrachain
charge transport between aggregated regions that enables these
high mobilities. This tie chain model reconciles well with
experimental results that polymers having weak and broad X-
ray diffraction peaks, but nonetheless still possessing local
order, can have relatively high mobilities. With these prior
investigations, there are a few questions that remain
unanswered. For example, is there a platform that can decouple
intrachain and interchain charge transport? Is there a way to
control the order and density of π-aggregates in the polymer
thin films? Can one achieve both local and long-range order for
D−A polymer thin films, and how would it impact charge
transport? Along with the development of high mobility
polymers, an accompanied problem has also emerged, that is,
the relatively poor solubilities of these polymers in common
organic solvents, which is considered to be a leading factor for
batch-to-batch variations in material synthesis, purification, and
device fabrication.20 In addition, the processing media is often
limited to high boiling point, carcinogenic and chlorinated
aromatic solvents. This fact not only imposes an environmental
concern but also diminishes the value of solution processing
and increases the manufacturing cost.
With these questions in mind, our group has recently
introduced complementary semiconducting polymer blends.21
The demonstrated blend is composed of a matrix polymer
(DPP-C3) and a tie chain polymer (DPP-C0). The DPP-C3
contains a propylene conjugation-break spacer (CBS) along the
polymer backbone. As a result, intrachain charge transport is
halted. On the other hand, the matrix polymer DPP-C3
presents higher crystallinity and much enhanced solubility,
together with a much lower melting transition that potentially
endows melt-processability. The tie chain polymer DPP-C0 is a
fully conjugated polymer, which is responsible for connecting
the π-aggregates of the matrix polymer in the blend. We
recently have showed that the incorporation of only 1 wt % of
tie chain polymer DPP-C0 in a DPP-C3 matrix led to 2 orders
of magnitude improvement in charge carrier mobility. Through
a combination of various morphological and electrical
characterizations, we established the concept of complementary
semiconducting polymer blends (c-SPB) for efficient charge
transport.
In this study, we aim to uncover the influence of conjugation-
break spacer length. In this aspect, we prepared and
characterized a set of 11 matrix polymers (DPP-Cm, m = 2−
12) with the CBS varying from 2 to 12 methylene groups and
studied the charge transport properties of their c-SPBs with a
fully conjugated tie chain polymer (DPP-C0, 2 wt %). This
work reveals that the length of the CBS plays a significant role
in determining both morphological and electrical properties of
the polymers. An odd−even effect is observed, which affects the
morphologies and charge transport properties of the polymer
thin films. This systematic investigation reconfirms that efficient
charge transport properties of complementary blends indeed
result from the interactions between matrix polymers and tie
chain polymers and begins to provide guidelines as to which
spacer lengths may be utilized to offer the best balance between
processing and charge transport properties.
RESULTS AND DISCUSSION
■
Monomer and Polymer Synthesis. The synthetic routes
for the CBS monomers and the diketopyrrolopyrrole (DPP)-
based polymers are summarized in Scheme 1. Similar to
previous studies,22−24 the linear alkyl chains were selected as
the conjugation-break spacers. For the monomers with three or
B
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Table 1. Physical Properties of the CBS Polymers
abs
max
λ
(nm)
energy levels (eV)
a
c
d
e
f
g
h
i
j
k
Mn (kDa)/PDI
Td (°C)
Tm (°C)
Tc (°C)
ΔHfus (J/g)
solution
film
Eogpt (eV)
EHOMO
ELUMO
b
DPP-C0
DPP-C2
DPP-C3
DPP-C4
DPP-C5
DPP-C6
DPP-C7
DPP-C8
DPP-C9
DPP-C10
DPP-C11
DPP-C12
30.7/3.61
400
404
394
394
393
401
396
397
404
390
406
405
740, 811
591, 633
595, 639
597, 630
594, 631
591, 632
590, 630
594, 627
590, 629
589, 629
589, 628
589, 628
725, 789
596, 647
601, 650
606, 645
605, 654
606, 655
603, 652
603, 654
601, 652
603, 649
599, 651
1.32
1.40
1.40
1.39
1.41
1.40
1.40
1.42
1.39
1.39
1.42
1.40
−4.97
−5.19
−5.17
−5.19
−5.14
−5.13
−5.17
−5.16
−5.19
−5.12
−5.11
−5.15
−3.65
−3.79
−3.77
−3.78
−3.75
−3.73
−3.77
−3.74
−3.78
−3.73
−3.69
−3.75
b
6.6/1.2 (8.74/2.1)
9.4/1.3 (13.3/1.8)
221
162
166
138
153
114
114
110
105
98
189
117
123
93
3.4
5.1
4.8
7.3
4.9
8.6
6.4
9.6
7.4
9.1
9.7
b
b
b
7.60/1.2 (12.2/2.2)
12.9/1.3 (14.9/2.0)
8.2/1.3
117
89
12.3/1.4
10.3/1.3
71
13.9/1.4
86
15.4/1.3
75
17.2/1.5
77
14.5/1.4
94
71
600, 651
a
b
c
d
Tetrahydrofuran as the eluent at room temperature. Trichlorobenzene as the eluent at 180 °C. Decomposition temperature. Melting
e
f
g
h
temperature. Crystallization temperature. Enthalpy of fusion. In chloroform solution. Spin-coated films on glass substrates, annealed at 120 °C.
i
j
k
Calculated from the onset absorption Eogpt = 1240/λ
(nm). Obtained from UPS, uncertainty 0.05 eV. Calculated using the equation ELUMO =
abs
onset
EHOMO + Eogpt
.
more methylene units, we followed the same synthetic protocol
as previously revealed.25 Thiophene end-capped alkyl com-
pounds (ThCmTh, m = 3−12) were prepared by the lithiation
of thiophene at cryogenic temperature, followed by the addition
of dibromoalkanes. The bis-stannylation ThCmTh gave the
resulting monomers Me3SnThCmThSnMe3, which appear to be
white solids or colorless liquids. For the C2 derivative, bis(5-
trimethylstannylthiophen-2-yl)ethane, however, the same pro-
tocol failed to provide 1,2-dithienylethane (ThC2Th) during
the first step because of the undesired elimination reaction. As
an alternative, the Huang−Minlon modification of the Wolff−
Kishner reduction of 2,2′-thenil was performed, yielding 1,2-
dithienylethane (ThC2Th) quantitatively.26 Subsequently,
ThC2Th was stannylated with butyllithium and trimethyl-
stannyl chloride to give 1,2-bis(5-(trimethylstannyl)thiophen-2-
yl)ethane (Me3SnThC2ThSnMe3) in an excellent yield. For the
monomers (Me3SnThCmThSnMe3) bearing even-numbered
methylene units, colorless crystals were obtained as pure
products. For the odd-numbered compounds, white solids were
obtained with short (m = 3) or long (m = 11) spacers. The
monomers with moderate length (m = 5, 7, and 9) were
obtained as colorless liquids. These different melting behaviors
of the monomers are a result of competition between π−π
interactions of the thiophene rings and the van der Waals
interactions between the alkylene chains. The structure of the
monomers was confirmed by 1H NMR, 13C NMR, and HR-MS.
Matrix polymers (DPP-Cm) were prepared by Stille
polymerization with the aid of microwave irradiation. The
synthetic details can be found in the Experimental Section. The
obtained polymers were purified by Soxhlet extraction with
acetone, hexane, and chloroform. The chloroform fraction was
precipitated into methanol. The solids were collected by
vacuum filtration and dried at 60 °C under high vacuum. The
polymers were characterized by 1H NMR, gel permeation
chromatography (GPC), thermal gravimetric analysis (TGA),
differential scanning calorimetry (DSC), UV−vis spectroscopy,
and ultraviolet photoelectron spectroscopy (UPS). The
molecular weights and polydispersities were determined by
GPC with tetrahydrofuran as eluent under room temperature as
well as trichlorobenzene at 180 °C for those relatively poorly
soluble polymers. The results are summarized in Table 1. The
molecular weights of the polymers are between 6.6 and 17.2
kDa with small polydispersities of 1.2−1.5 when THF is used as
the eluent under room temperature. We observed that the
molecular weights increase with the length of the CBSs. We
suspected that this trend could be a result of solubility
difference because polymers with short CBSs clearly exhibit
lower solubilities in THF solutions. In other words, the GPC
data only reflect the low molecular weights fractions of the
relatively poorly soluble polymers, and high molecular weight
portions are filtered out.27 To confirm, we managed to carry
out GPC measurements with trichlorobenzene at 180 °C at
which all the polymers are fully dissolved and no polymers are
filtered out. The results are highlighted in parentheses as shown
in Table 1. For polymers measured under high temperature, an
increase in both molecular weight and polydispersity is
observed. From the high temperature GPC measurements, it
confirms that all prepared polymers have comparable molecular
weights and polydispersities.
Optical and Thermal Properties. The optical properties
of these polymers were evaluated from UV−vis measurements
and are summarized in Table 1. The solution and solid-state
UV−vis spectra are exhibited in Figures S3 and S4. In
chloroform solutions, the polymers exhibit a structured
absorbing band with absorption maximum near 630 and 590
nm, which is attributed to the well-defined chromophore along
the polymer backbone. This is different from their fully
conjugated polymer (DPP-C0), which presents a broad and
featureless absorption band. In thin films, a bathochromic shift
about 20 nm is observed for all polymers. A broad absorption at
around 800 nm implies the formation of π-aggregates in the
thin films.25 The absorption spectra of c-SPB films (2 wt %
DPP-C0 blended with the CBS polymers) are also measured
and presented in Figure S5. Compared to the pure polymer
films, no significant change is observed. The energy levels of the
polymers, as measured by ultraviolet photoelectron spectros-
copy (UPS), are summarized in Table 1, and representative
spectra are shown in Figure S6. All polymers with a CBS have
HOMO levels of −5.15 0.04 eV, which is within the 0.05
eV uncertainty associated with the measurement.
Thermal stability of the polymers was studied by thermal
gravimetric analysis (TGA), and the results are summarized in
Table 1 and Figure S7. All polymers show good thermal
stability, and no decomposition is observed until around 400
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°C. Phase transition behaviors of the polymers were studied by
differential scanning calorimetry (DSC). The results are
summarized in Table 1 and provided in Figure 1 and Figure
beyond dodecane.28 Heat of fusion, surprisingly, presents an
opposite trend. An ascending trend is observed as the length of
CBS increases, as is shown in Figure 1. For instance, the
enthalpy of fusion for DPP-C2 was 3.4 J/g, which was only
about one-third of DPP-C12 (9.7 J/g). An odd−even effect is
also observed for the heat of fusion. Counterintuitively, the
odd-numbered polymers exhibit larger heat of fusion than the
adjacent even-numbered polymers. For instance, DPP-C9 has
the heat of fusion of 9.6 J/g, while the numbers for DPP-C8
and DPP-C10 are 6.4 and 7.4 J/g, respectively. Heat of fusion is
directly correlated to the energy required to break the
intermolecular forces, suggesting that the odd-numbered
polymers have higher ordering and crystallinity, with the
structural similarities among all polymers.29 This observation is
in a good agreement with atomic force microscopy (AFM) and
grazing incidence X-ray diffraction (GIXRD) measurements.
Charge Transport Properties of Matrix Polymers and
the Complementary Polymer Blends. The charge transport
properties of DPP-Cm (m = 2−12) and their c-SPBs with the
tie chain polymer DPP-C0 were investigated using bottom-gate,
bottom-contact device configurations. The silicon wafer was
modified with octadecyltrichlorosilane (OTS-18). The gold
(Au) source-drain electrodes were prepatterned by photo-
lithography. The polymer thin films were formed by spin-
casting of 1,2-dichlorobenzene solutions (5 mg/mL for DPP-
C2 and 10 mg/mL for the rest). For c-SPB thin films, matrix
polymers were premixed with the tie chain DPP-C0 (2 wt %).
After annealing at 80 or 120 °C under the nitrogen atmosphere
in the glovebox for the desired duration, the devices were tested
in air. Average mobilities were calculated from more than five
batches of devices with more than four devices in each batch for
statistical meaningful results.
Figure 1. Melting point and heat of fusion as a function of the number
of methylene units in the CBSs.
S8. Melting transitions are observed in the range of 90−220 °C
for all polymers, while the corresponding crystallizations occur
between 190 and 70 °C. It is observed that melting transitions
correlate with the length of the CBSs. The longer the CBSs, the
lower the melt transition temperatures are observed for the
polymers. For instance, DPP-C2 exhibits a broad melting peak
at 221 °C, while the melting point of DPP-C12 is as low as 71
°C. This phenomenon is a result of increasing the flexibility of
the polymer backbone, resulting in lower melting points.
Interestingly, a distinct morphology difference is observed for
polymers with shorter CBSs (i.e., DPP-C2 to DPP-C8). The
even-numbered polymers exhibit higher melting points than the
adjacent odd-numbered polymers. For example, DPP-C6
presents a melting point at 153 °C, while DPP-C5 and DPP-
C7 exhibit melting transitions at 138 and 114 °C, respectively.
This trend fades for the polymers with CBSs beyond eight
methylene groups. A monotonous trend is observed for the
melting points of DPP-C8 (114 °C) to DPP-C12 (94 °C). This
overall trend is similar to the melting points of n-alkanes, in
which an obvious odd−even effect is observed between
methane and undecane and a monotonous trend is observed
For all matrix polymers, typical p-channel charge transport
behaviors are observed. The OFET results show a general
descending trend in mobility as the length of CBS increases.
The data from the OFET devices annealed at 120 °C are
summarized in Table 2 and Figure 2. Their output and transfer
curves are shown in Figure S9. DPP-C2 exhibits the highest
mobility among all matrix polymers. When annealed at 120 °C,
−1
the maximum mobility of 0.032 cm2 V s−1 and the average
mobility of 0.022 cm2 V−1 s−1 are obtained. The polymer with
the longest CBS in this series, namely DPP-C12, gives an
average mobility of 7.9 × 10−6 cm2 V−1 s−1, which is about 3
orders of magnitude lower than DPP-C2. Exponential fit is used
to calculate the decreasing rate of the mobility. The result
Table 2. Charge Transport Characteristics of DPP-Cm and c-SPB with 2 wt % of Tie Chain Polymer DPP-C0 Annealed at 120
°C
matrix polymers DPP-Cm
c-SPBs with DPP-C0 (2 wt %)
m (length of CBS)
μavg (cm2/(V s))
μmax (cm2/(V s))
Vth (V)
Ion/Ioff
μavg (cm2/(V s))
μmax (cm2/(V s))
Vth (V)
Ion/Ioff
2
0.023
0.032
−9.3
−9.8
−8.4
−6.9
−5.3
−6.6
−3.4
−7.1
−12.5
−14.1
−14.4
107−108
0.78
1.29
1.23
1.00
0.77
0.54
0.31
0.15
0.15
0.11
0.098
0.029
−10.2
−12.1
−11.7
−8.4
−8.9
−9.1
−7.4
−8.9
−10.2
−8.2
−9.8
∼107
3
0.012
0.025
∼107
0.81
106−107
106−107
106−107
106−107
106−107
∼106
4
0.013
0.025
∼107
0.60
5
0.0023
0.0055
106−107
∼107
0.53
6
0.0023
0.0052
0.26
7
8.2 × 10−4
5.4 × 10−4
1.2 × 10−4
1.6 × 10−4
4.9 × 10−5
7.9 × 10−6
0.0019
105−106
105−106
105−106
105−106
∼105
0.099
0.081
0.081
0.066
0.052
0.012
8
0.0011
9
1.7 × 10−4
2.5 × 10−4
1.0 × 10−4
1.3 × 10−5
105−106
105−106
105−106
∼105
10
11
12
∼104
D
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we have shown that the addition of small percentage of DPP-
C0 (1 wt %) into DPP-C3 leads to nearly 2 orders of
magnitude improvement in charge mobility.21 Similar behaviors
are observed for the entire matrix polymers studied in this
work, as plotted in Figure 2 and presented in Figure S10. In
other words, 2 or 3 orders of magnitude improvement in charge
mobility is generally observed. For instance, c-SPBs with DPP-
C2 or DPP-C3 as the matrix polymer show the highest
mobilities of ∼0.8 cm2 V−1 s−1, which are 1 or 2 orders of
magnitude higher than the pure matrix polymers. When using
DPP-C11 or DPP-C12 as the matrix polymer, c-SPBs provide
average mobilities of 0.052 and 0.029 cm2 V−1 s−1, which are
over 1000 times higher than their matrix polymers. By contrast
to the pure matrix polymers, charge transport in the c-SPBs is
significantly less sensitive to the length of CBSs. For instance,
the c-SPBs containing DPP-C8, -C9, -C10, and -C11 present
nearly the same charge mobilities. Similarly, c-SPBs with DPP-
C2 or DPP-C3 as the matrix polymer show nearly identical
charge mobilities, despite that the charge mobility of DPP-C2 is
on average 2 times higher than that of DPP-C3. In addition, the
odd−even effect is absent for the c-SPBs. In our previous work,
we have confirmed the applicability of the tie chain model for
describing charge transport behaviors in the c-SPBs. That is the
tie chain polymer connects crystalline π-aggregates of the
matrix polymer, leading to enhanced charge transport proper-
ties. The phenomena observed here can be explained by this
model. For the matrix polymers, DPP-Cm, interchain transport
is the primary pathway for charge carriers because of the
absence of π-conjugation and delocalization along the back-
bone. Interchain transport is largely dependent on the density
of π-aggregates, which are further correlated to the length of
CBS. The longer the CBS, the lower the π-aggregate density
and the smaller probability to form effective charge transport
network. This relationship essentially explains the observation
that charge mobility drops roughly 6 times with adding two
methylene units to the CBS for matrix polymers. In the c-SPBs,
however, efficient intrachain charge transport originating from
the addition of tie chain polymer, DPP-C0, changes the
transport landscape. An interconnected network is formed
between crystalline π-aggregates and tie chains, which makes
charge transport behaviors no longer solely dependent on the
density of π-aggregates in matrix polymers. To summarize, we
are able to show how the density of π-aggregates can affect the
charge transport properties. The present study also reconfirms
that efficient charge transport properties of the complementary
blends are resulted from the interactions between matrix
polymers and tie chain polymers.
Figure 2. Charge carrier mobility as a function of number of
methylene groups in the CBSs in both pure DPP-Cm (red) and c-SPB
with 2 wt % of DPP-C0 (blue) OFETs annealed at 120 °C. Error bar
represents standard deviation.
reveals that for every two methylene units in the CBS, the
mobilities of the polymers would be 6 times lower on average.
This downward trend is in good agreement with the fact that
polymers with longer CBSs have a lower density of π−π
aggregates. It is known that π−π aggregates are crucial for
interchain charge transport. Previously, we have shown that
DPP-C20 failed to show any OFET performance.21 According
to the observed relationship here, the mobility of DPP-C20
−1
would be 1.0 × 10−8 cm2 V
properly measured.
s−1, which is too low to be
Annealing of the devices at 80 °C was also performed
because some of the polymers have melt transitions below the
previous annealing temperature of 120 °C. Similar results were
obtained (see Figure S11). Comparing the OFETs perform-
ances at two different annealing temperatures, we find that for
those polymers with high melting points (>130 °C, DPP-C2 to
DPP-C6), higher annealing temperature (120 °C) leads to a
better performance. For those polymers with low melting
points (<120 °C, DPP-C7 to DPP-C12), no additional
improvement is observed with the annealing temperature of
120 °C. The additional thermal energy provided during
annealing allows the polymers to reorganize, which results in
more ordered film morphologies, larger crystal grains, and less
grain boundaries. For DPP-C7 to DPP-C12, the melting points
are relatively low, and the annealing temperature at 80 °C is
sufficient enough to achieve the desired morphologies. In
contrast, for DPP-C2 to DPP-C6, a higher annealing
temperature of 120 °C is required to achieve optimum
morphologies.
Besides the observed descending trend, the odd−even effect
is also observed for this set of matrix polymers. Ratios between
measured charge carrier mobility vs exponential fit values are
shown in Figure S12. Polymers with even-numbered CBSs
generally show higher mobilities than the adjacent odd-
numbered counterparts. For instance, DPP-C10 exhibits the
highest mobility of 2.5 × 10−4 cm2 V−1 s−1 and an average
mobility of 1.6 × 10−4 cm2 V−1 s−1. The adjacent odd-
numbered derivatives, DPP-C9 and DPP-C11, present highest
mobilities of 1.7 × 10−4 and 1.0 × 10−4 cm2 V−1 s−1 and the
average mobilities of 1.2 × 10−4 and 4.9 × 10−5 cm2 V−1 s−1,
respectively. The only exception is DPP-C12 which has
mobility lower than the adjacent odd-numbered DPP-C11.
To investigate the influence of the length of CBS on the
charge transport properties of the c-SPBs, the fully conjugated
tie chain polymer DPP-C0 (2 wt %) was blended with the
matrix polymers DPP-Cm (m = 2−12). In our concept paper,
Thin Film Morphology Characterization. To correlate
charge transport properties with thin film morphologies,
tapping-mode atomic force microscopy (AFM) and grazing
incidence X-ray diffraction (GIXRD) were used to probe the
thin films of matrix polymers and their c-SPBs (with 2 wt % of
the tie chain polymer DPP-C0). The AFM images are shown in
Figure 3 and Figures S13−S16. Compared with the films
without annealing (Figure S13), most of the films annealed at
120 °C (Figure 3 and Figure S15) show larger domain sizes.
Figure 3 highlights the AFM images of DPP-C2 to DPP-C7
films annealed at 120 °C. The polymers with the even-
numbered linkages have a more interconnected feature that
appear more fibrillar than the polymers with odd-numbered
CBSs.30 It is also noted that the AFM image of DPP-C7 thin
film annealed at 120 °C is much smoother. Similar features are
also observed for DPP-C9 to DPP-C12 thin films (Figure S15).
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polymers are smaller than the adjacent even-numbered
polymers, while the d−d spacing distances present an opposite
trend. The only exception is the d−d spacing distances of DPP-
C6. For example, DPP-C4 has a π−π stacking distance of 3.65
Å, which is larger than both DPP-C3 (3.61 Å) and DPP-C5
(3.64 Å). In addition, DPP-C4 has the d−d spacing distance of
21.6 Å, which is smaller than both DPP-C3 (22.7 Å) and DPP-
C5 (22.3 Å). The molecular packing distances observed here
are similar to DPP-based fully conjugated polymers. 1D regular
XRD measurement was also carried out to obtain more
information on the bulk material. The results agree very well
with the GIXRD measurements, as is shown in Figure S21. The
opposite trend in odd−even effects between the π−π and the
lamellar stacking distances suggests that the even-numbered
polymers have higher tilting angles with respect to the surface
normal than the odd-numbered counterparts. We note that this
effect has been observed in small molecule organic semi-
conductors recently.31 The origin of such an effect shall be
linked to the tilting angles of the CBSs between the conjugated
segments. For the c-SPB films, a similar odd−even effect is also
present in the π−π stacking direction. Such a rich odd−even
effect in semiconducting polymers resulted from varying the
length of conjugation-break spacers along the polymer
backbones has not been well-known. It opens a new dimension
in molecular design to control the polymer properties by using
proper conjugation break spacers.
Figure 3. AFM images and cross-section figures of polymer thin films
from DPP-C2 to DPP-C7. The polymer thin films were fabricated by
spin-coating on OTS-modified Si/SiO2 substrates and annealed at 120
°C.
Odd−Even Effect, Physical Property, Thin Film
Morphology, and Charge Transport. The odd−even effect
revealed in this study offers a new dimension of controlling
properties of semiconducting polymers through the rational
molecular design, namely introducing conjugation-break
spacers along the polymer backbone. In earlier work, Pei et
al. prepared a set of small molecules by linking two conjugated
moieties with alkyl spacers and studied their influence on
microwire growth and crystal packing. From the crystal
structures, the odd-numbered aromatic moieties tend to exhibit
bent or “V-shaped” configuration, whereas zigzag or “Z-shaped”
configuration is favored in even-numbered derivatives. It was
also observed that the odd-numbered derivatives exhibit a
strong tendency of one-dimensional growth. In contrast, the
even-numbered molecules do not show any preferred crystal
growing tendency.32 In the current study, we confirmed that
the odd−even effect observed in small molecules can be
extended into polymeric materials. A summary of observations
from this work are as follows: (1) Polymers with odd-
numbered CBSs show better solubilities, lower melting
transitions, smaller π−π stacking distances, two-dimensional
lamellar-like morphologies, and higher crystallinity. (2) When
the length of CBS exceeds eight methylene units, the odd−even
effect in melting temperature becomes more subtle. (3) Charge
mobilities of this set of matrix polymers inversely increase as
the CBS length increases. Generally speaking, the longer the
CBS, the lower the π-aggregate density and the less likely to
form an efficient charge transport network for matrix polymers.
However, the odd−even effect is notably present in charge
transport, with even-numbered polymers outperforming their
adjacent odd-numbered counterparts. Considering that odd-
numbered polymers exhibit smaller π−π stacking distances and
higher crystallinity, this observation is unexpected. The possible
explanation is that Z-shaped configuration in even-numbered
polymers is more beneficial for the formation of interconnected
charge transport network than V-shaped configuration in odd-
numbered polymers, hinted by the crystal structures in small
With longer and more flexible CBSs, the melting temperatures
of polymers (from DPP-C7 to DPP-C12) are lower than the
annealing temperature of 120 °C. This renders less-featured
AFM images for these polymers. Subsequently, we carried out
AFM on these polymer thin films annealed at 80 °C and were
able to obtain clear images for DPP-C7 and DPP-C9 to C12, as
is revealed in Figure S14. In addition, AFM measurements
suggest that polymer with longer CBSs films become “softer”,
which leads to difficulties in getting clear AFM images even
using low force-constant AFM probe tips (k = 7.8 N/m). From
grazing incidence X-ray diffraction (GIXRD) experiments in
the next section, however, we can see the morphological and
packing difference between odd-numbered polymers and even-
numbered polymers, including the polymers with longer spacer
lengths (DPP-C8 to DPP-C12). Furthermore, we performed
AFM on c-SPB thin films. Little difference is observed between
the matrix polymer thin films and their corresponding blend
thin films, as exhibited in Figure S16. It implies that the
addition of 2 wt % of tie chain polymer DPP-C0 does not alter
the thin film morphology macroscopically.
Grazing incidence X-ray diffraction (GIXRD) was performed
to illustrate the structure−property relationship for matrix
polymer and their c-SPB thin films. The results are summarized
in Figures S17−S20. From the 2D-GIXRD patterns, we observe
that the π−π stacking peaks (010) of the matrix polymer and c-
SPB thin films appeared in the in-plane direction, and the layer-
by-layer (d−d spacing) peaks (h00) appear in the out-of-plane
direction. This indicates that the “edge-on” orientation is
dominant in all thin films. The corresponding 1D-GIXRD data
and peak information are shown in Figures S18 and S19. The
packing distances are calculated and plotted in Figure S20. The
results show clearly that the π−π stacking distance increase
with the length of CBSs. In a good agreement with the AFM
results, the odd−even effects are confirmed by X-ray diffraction
experiments. The π−π stacking distances of the odd-numbered
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spectral shifts, and the photon flux was adjusted to minimize sample
damage.
molecules. (4) Charge transport in the c-SPBs is improved
dramatically (3−4 orders of magnitude improvement on
average). The odd−even effect is absent in the c-SPBs. This
is because the interconnected network form by π-aggregates in
matrix polymer is no longer the only deciding factor of charge
transport. Instead, the interactions between π-aggregates and tie
chains and their newly formed interconnected network play a
dominant role in charge transport.
General Method for Polymerization. 3,6-Bis(5-bromothiophen-
2-yl)-2,5-bis(4-decyltetradecyl)-2,5-dihydropyrrolo[3,4-c]pyrrole-1,4-
dione (88 μmol) and ditin monomers (88 μmol) were added into a 35
mL microwave vessel. Upon stirring, the monomers were dissolved in
15 mL of anhydrous toluene and degassed with nitrogen for 15 min.
Pd2(dba)3 (2.5 mol %) and tris(o-tolyl) phosphine (4 mol %) were
added under nitrogen. The vessel was capped with a snap cap and
transferred to the CEM Discover automatic microwave reactor. The
reaction conditions were listed as follows: power cycling mode; power,
200 W; power cycles, 100; temperature, 150−180 °C; heating, 120 s;
cooling, 30 s; pressure, 150 psi; stirring, high. After the polymerization
was complete, the mixture was taken up and precipitated into
methanol. The solids were collected by a high quality glass thimble,
which was purified by Soxhlet extraction with acetone, hexane, and
chloroform. The chloroform fraction was precipitated into methanol.
The collected polymer was dried at 60 °C under vacuum with the
yields in the range of 61%−78%.
Fabrication of FET Devices. A heavily n-doped Si wafer with a
300 nm SiO2 surface layer (capacitance of 11 nF/cm2) was employed
as the substrate with Si wafer serving as the gate electrode and SiO2 as
the dielectric. The gold S/D electrodes were sputtered and patterned
by photolithography technique. The device channel width was 1400
μm; the channel length was 5 μm for pure DPP-Cm OFETs and 50
μm for blend OFETs. For the octadecyltrichlorosilane (OTS)
modification, the silicon wafer (with Au bottom contact) was first
cleaned with hot piranha solution (H2SO4 (98%):H2O2 (30% water
solution) = 7:3). It was then further subjected to sonication
sequentially in water and acetone for 6 min each. After drying at an
oven, the silicon wafer was then put in a Petri dish with a small drop of
OTS in the center. The dish was then covered and heated in a vacuum
oven at 120 °C for 3 h, resulting in the formation of an OTS self-
assembled monolayer on the surface. The OTS-modified substrates
were rinsed successively with hexane, ethanol, and chloroform and
dried by nitrogen. The semiconductor layer was deposited on the
OTS-treated Si/SiO2 substrates by spin-coating with spin speed of
2000 rpm for 2 min. The concentrations of the solutions used for spin-
coating were 3 mg/mL for polymer DPP-0, 5 mg/mL for DPP-C2,
and 10 mg/mL for other DPP-Cm polymers. Before spin-coating, the
solutions were heated up to 50 °C. The devices were annealed in a N2-
purged glovebox at 80 or 120 °C and then tested in open air.
Device Characterization. Device characterization of the
fabricated OFETs was carried out using Keithley 4200 in ambient
air. The field-effect mobility was calculated in the saturation regime by
using the equation IDS = (μWCi/2L)(VG − VT)2, where IDS is the
drain−source current, μ is the field-effect mobility, W is the channel
width, L is the channel length, Ci is the capacitance per unit area of the
gate dielectric layer, VG is the gate voltage, and VT is the threshold
voltage.
CONCLUSION
■
In summary, we have prepared and characterized a set of 11
matrix polymers containing conjugation-break spacers with
various lengths. Our findings reveal a profound odd−even
effect on the properties of this class of polymers, ranging from
solubility, melting transition, thin-film morphology, and charge
transport. This work also reconfirms that efficient charge
transport properties of the complementary blends result from
interactions between π-aggregates in matrix polymers and tie
chains of polymers. At present, we plan to take advantage of the
fact that charge transport in c-SPBs is not overly sensitive to the
length of conjugation-break spacer in matrix polymers and
explore the potentials of melt processing of c-SPBs.
EXPERIMENTAL SECTION
Materials and Characterizations. All reagents and starting
■
materials were purchased from Sigma-Aldrich and used without further
1
purification unless otherwise noted. H and 13C NMR spectra were
recorded using Varian Inova 300 and Bruker ARX 400 at 293 K with
deuterated chloroform as solvent. Room temperature gel permeation
chromatography (GPC) was performed in tetrahydrofuran under
room temperature with a Polymer Laboratories PL-GPC20. High
temperature GPC was carried out in trichlorobenzene under 180 °C.
The molecular weights were calculated using a calibration curve based
on polystyrene standards. Thermal gravimetric analyses (TGA) were
performed using a TA Instruments Q50 with samples heated from 40
to 800 °C at a rate of 10 °C/min under a nitrogen flow (60 mL/min).
Differential scanning calorimetry (DSC) measurements were carried
out using a TA Q5000 calorimeter with an indium standard to
calibrate the instrument and nitrogen as the purge gas (50 mL/min).
Samples were sealed in hermetic aluminum pans. Each measurement
included two cycles with heating and cooling rates at 10 °C/min. UV−
vis−NIR spectra were recorded on an Cary 50 spectrophotometer
(300−1100 nm). High-resolution mass data were measured with
Micromass 70-VSE. Atomic force microscopy images were obtained
on a Veeco Dimension 3100 AFM in tapping mode. Grazing incidence
X-ray diffraction was performed at the Argonne National Laboratory
on beamline 8-ID-E. Data were collected with a two-dimensional
detector (Pilatus 1M) to obtain molecular packing information. The
beam energy was 7.35 keV. Experiments were carried under ambient
condition. One scan was carried for each sample with incident angle
set at 0.2°. For UPS analysis, indium tin oxide (ITO) coated glass
slides (RS = 15 Ω/□) were used as the substrates. The substrates were
cleaned through sequential sonication in sodium dodecyl sulfate
aqueous solution, deionized water, acetone, and isopropanol, followed
by UV-ozone cleaning for 10 min. To encourage polymer wetting, for
films of DPP-C10, -C11, and -C12, the substrates were spun-cast with
a layer of PEDOT:PSS (Clevios P VP Al 4083) at 5000 rpm that was
annealed at 130 °C for 15 min in air. Polymer solutions were prepared
at 10 mg/mL in 1,2-dichlorobenzene, with the exceptions of DPP-C0
and DPP-C2 which were prepared at 3 and 5 mg/mL, respectively, and
stirred overnight at 50 °C. The films were spun-cast at 2500 rpm in a
nitrogen-filled glovebox (H2O and O2 < 0.3 ppm) and transferred
without air exposure into a PHI 5600 ultrahigh-vacuum system (UHV)
for UPS measurements. UPS measurements were performed with a
10.20 eV photon source33 (E-lux, Excitech GmbH) and an 11 in.
diameter hemispherical electron energy analyzer with a multichannel
plate detector. The samples were monitored for damage-induced
ASSOCIATED CONTENT
■
S
* Supporting Information
The Supporting Information is available free of charge on the
ACS Publications website at DOI: 10.1021/acs.macro-
mol.6b00050.
NMR, TGA, UV−vis spectra, DSC curves, additional
AFM figures, GIXD data, OFET characteristics (PDF)
AUTHOR INFORMATION
■
Corresponding Author
*E-mail jgmei@purdue.edu (J.M.).
Author Contributions
X.Z. and Y.Z. made equal contributions to the work.
Notes
The authors declare no competing financial interest.
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F. Liquid-crystalline semiconducting polymers with high charge-carrier
mobility. Nat. Mater. 2006, 5, 328−333.
ACKNOWLEDGMENTS
■
The authors acknowledge the financial support from Purdue
University. GIXRD measurements were carried out at the
Advanced Photon Source−Argonne National Laboratory,
operated by the University of Chicago for the Department of
Energy (DOE), Office of Basic Energy Sciences.
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