Acene Ring Size Optimization in Fused Lactam Polymers Enabling High n-Type Organic Thermoelectric Performance

Article abstract of DOI:10.1021/jacs.0c10365

Three n-type fused lactam semiconducting polymers were synthesized for thermoelectric and transistor applications via a cheap, highly atom-efficient, and nontoxic transition-metal free aldol polycondensation. Energy level analysis of the three polymers demonstrated that reducing the central acene core size from two anthracenes (A-A), to mixed naphthalene-anthracene (A-N), and two naphthalene cores (N-N) resulted in progressively larger electron affinities, thereby suggesting an increasingly more favorable and efficient solution doping process when employing 4-(2,3-dihydro-1,3-dimethyl-1H-benzimidazol-2-yl)-N,N-dimethylbenzenamine (N-DMBI) as the dopant. Meanwhile, organic field effect transistor (OFET) mobility data showed the N-N and A-N polymers to feature the highest charge carrier mobilities, further highlighting the benefits of aryl core contraction to the electronic performance of the materials. Ultimately, the combination of these two factors resulted in N-N, A-N, and A-A to display power factors (PFs) of 3.2 μW m-1 K-2, 1.6 μW m-1 K-2, and 0.3 μW m-1 K-2, respectively, when doped with N-DMBI, whereby the PFs recorded for N-N and A-N are among the highest reported in the literature for n-type polymers. Importantly, the results reported in this study highlight that modulating the size of the central acene ring is a highly effective molecular design strategy to optimize the thermoelectric performance of conjugated polymers, thus also providing new insights into the molecular design guidelines for the next generation of high-performance n-type materials for thermoelectric applications.

Full text of DOI:10.1021/jacs.0c10365

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Article
Acene Ring Size Optimization in Fused Lactam Polymers Enabling
High n‑Type Organic Thermoelectric Performance
Hu Chen,*,∞ Maximilian Moser,*,∞ Suhao Wang, Cameron Jellett, Karl Thorley, George T. Harrison,
Xuechen Jiao, Mingfei Xiao, Balaji Purushothaman, Maryam Alsufyani, Helen Bristow, Stefaan De Wolf,
Nicola Gasparini, Andrew Wadsworth, Christopher R. McNeill, Henning Sirringhaus, Simone Fabiano,
and Iain McCulloch
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ABSTRACT: Three n-type fused lactam semiconducting poly-
mers were synthesized for thermoelectric and transistor applica-
tions via a cheap, highly atom-efficient, and nontoxic transition-
metal free aldol polycondensation. Energy level analysis of the
three polymers demonstrated that reducing the central acene core
size from two anthracenes (A-A), to mixed naphthalene−
anthracene (A-N), and two naphthalene cores (N-N) resulted in
progressively larger electron affinities, thereby suggesting an
increasingly more favorable and efficient solution doping process
when employing 4-(2,3-dihydro-1,3-dimethyl-1H-benzimidazol-2-
yl)-N,N-dimethylbenzenamine (N-DMBI) as the dopant. Mean-
while, organic field effect transistor (OFET) mobility data showed
the N-N and A-N polymers to feature the highest charge carrier
mobilities, further highlighting the benefits of aryl core contraction to the electronic performance of the materials. Ultimately, the
combination of these two factors resulted in N-N, A-N, and A-A to display power factors (PFs) of 3.2 μW m−1 K−2, 1.6 μW m−1 K−2,
and 0.3 μW m−1 K−2, respectively, when doped with N-DMBI, whereby the PFs recorded for N-N and A-N are among the highest
reported in the literature for n-type polymers. Importantly, the results reported in this study highlight that modulating the size of the
central acene ring is a highly effective molecular design strategy to optimize the thermoelectric performance of conjugated polymers,
thus also providing new insights into the molecular design guidelines for the next generation of high-performance n-type materials
for thermoelectric applications.
their lightweight and flexible nature; (iv) the significantly
higher abundance and lower cost of their constituent elements;
(v) their lower thermal conductivity.1,2,6,7 Nonetheless, organic
materials also display various drawbacks compared to their
inorganic counterparts, including reduced Seebeck coefficients
(S) and electrical conductivities (σ), ultimately resulting in
lower thermoelectric performance on average. The develop-
ment of higher performing organic thermoelectric (OTE)
materials is thus imperative to promote their wider adoption.
While electrical conductivities of >1000 S cm−1 and power
factors (PFs) of >100 μW m−1 K−2 have been achieved for
several p-type organic semiconductors,8−11 even the highest
INTRODUCTION
■
With the continuously growing demand for energy and the
need to reduce greenhouse gas emissions it is imperative to
develop sustainable and environmentally friendly energy
solutions to effectively combat global warming. Given the
ubiquity of waste heat produced from industrial processes,
thermoelectric generators capable of converting thermal
gradients into electrical energy are a promising option to
contribute solving the energy issues faced in the 21st century.1
Thermoelectric generators are particularly attractive since they
do not contain any moving mechanical components, thus
reducing wear and improving the durability of devices.2
Moreover, their facile scalability should enable their applic-
ability to a range of systems, including sensors, electronic
components for the Internet of Things (IoT) and wearable
electronics.1,3,4 Although the vast majority of thermoelectric
research has been focused on inorganic materials,5 their
organic counterparts display several advantages, including (i)
their compatibility with solution-based processing and low-cost
roll-to-roll printing techniques; (ii) their reduced toxicity; (iii)
Received: September 28, 2020
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Figure 1. Chemical structures of the rigid polymers investigated within this study.
performing n-type polymers lag behind these values.12−15 This
performance disparity presents a significant issue, since
maximizing the energy conversion process in a thermoelectric
generator necessitates a balanced performance of both p- and
n-type components. Developing new n-type polymers with
improved thermoelectric performance and formulating struc-
ture−property relationships are therefore of utmost impor-
tance.
studies evaluating the effects of rational chemical modification
in terms of device performance.15,24,26 To provide further
insights into the design of high-performance n-type OTE
semiconductors, we have investigated the modulation of the
aromatic core size of a fused lactam polymer backbone to tune
the thermoelectric performance. This involved synthesizing
three fused acene containing polymers formed by the aldol
condensation of bis-isatin and bis-oxindole monomers,
specifically N-N, A-N, and A-A, with their structures illustrated
in Figure 1. Tuning the acene core size was deemed as a
particularly interesting molecular engineering strategy, given
that electron polaron binding energies and hence the charge
carrier transport abilities of organic semiconductors are
dependent on the size and rigidity of the constituent acene
core, and the length of the acene itself can tune the LUMO by
varying the spacing along the backbone of the electron
withdrawing lactam units.27
Currently, the most widely explored n-type polymers for
OTE are naphthalenetetracarboxylic diimide (NDI) based
polymers, specifically N2200. In thermoelectric devices, N2200
affords σ in the range of 10−3 S cm−1 and PF in the order of
10−2 μW m−1 K−2.16−20 The modest performance of N2200 is
typically ascribed to (i) its poor miscibility with the employed
dopant19,21 and (ii) the strong localization of its lowest
unoccupied molecular orbital (LUMO) and hence the polaron
on the NDI fragment, which arises from a combination of the
large dihedral angle between the NDI and bithiophene (T2)
moieties and the poor-electron accepting abilities of the T2
unit.16,18 The potential unsuitability of employing donor−
acceptor (D−A) polymer backbones for OTEs has also been
demonstrated in a related set of studies on p-type OTE
materials, in which a D−A polymer backbone was compared to
its all-donor analogue.22 The results obtained indicate that the
D−A structure poses an obstacle to efficient doping, with the
acceptor moiety being inactive for p-type doping. In an
analogous fashion, one can envisage that this concept is likely
to also apply when D−A polymers are employed as n-type
OTE materials, where the donor units will be inactive for n-
type doping. With this in consideration, conjugated ladder type
polymers that exhibit an all-acceptor polymer backbone seem
to be particularly attractive for n-type OTE, especially when
considering (i) their excellent thermal stability rendering them
ideal candidates for thermoelectric applications, where thermal
stability is crucial.15,18,23,24 (ii) The presence of electron-
withdrawing carbonyl groups in the conjugated polymer
backbone is beneficial as they contribute to reducing the
LUMO energy, thus rendering the n-type doping of these
materials thermodynamically more favorable.15,24,25 (iii) Their
highly rigid conjugated polymer backbones facilitate polaron
delocalization, consequently aiding charge carrier transport.18
Although the performance of ladder type polymers for n-
type thermoelectric applications has improved considerably in
recent times and PF values on the order of 100−101 μW m−1
K−2 have been reported, the elucidation of structure−property
relationships of conjugated ladder type polymers for n-type
OTE has remained scarce in the literature, with only few
POLYMER SYNTHESIS
■
The syntheses of the bis-oxindole and bis-isatin monomers
followed a general strategy reported in the literature.25 For the
bis-oxindole monomers, this involved acylation of the
corresponding aryl diamine with lauroyl chloride, followed
by reduction of the resulting amide with lithium aluminum
hydride to the analogous amine. The amine containing aryl
species was subsequently reacted with chloroacetyl chloride
followed by an intramolecular Heck-type cyclization to yield
the desired bis-oxindole. A more detailed explanation of the
synthetic procedures can be found in the Supporting
Information.
The synthesis of the bis-isatin monomers on the other hand
was based on a Martinet isatin approach. Specifically, this
involved reacting the same aryl diamine starting material
employed for the bis-oxindole synthesis with diethyl
ketomalonate followed by alkylation with the desired alkyl
iodide. The polymers N-N, A-N, and A-A were ultimately
synthesized by acid catalyzed aldol polycondensation between
the enolizable carbonyl unit of the bis-oxindole monomers and
the electrophilic carbonyl unit of the corresponding bis-isatin
comonomers. This polymerization approach offers numerous
advantages compared to the traditional transition-metal
catalyzed one, as it does not necessitate the use of expensive
transition metal catalysts such as palladium and it does not
require the formation and use of highly toxic organostannane
coupling partners, such as those required for Stille cross-
couplings.23,25 Ultimately, the resulting polymers were
obtained in high molecular weights, with number-average
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molecular weights (Mn) > 50 kDa, which is often a prerequisite
for high-performance materials in organic electronic applica-
tions.28,29
to their naphthalene counterparts. The IP extracted from the
HOMO region of the UPS spectra for the three semi-
conductors displayed the same trend as their EA, with the all
naphthalene-based polymer, N-N, showing the largest IP and
the all anthracene-based polymer, A-A, possessing the smallest
IP, indicating that the orbital electron densities are strongly
determined by the electron withdrawing properties of the
lactam groups.
UV−vis−NIR absorption spectroscopy was conducted on
both pristine and doped thin-film samples of the three
semiconductors. The corresponding UV−vis−NIR absorption
spectra can be found in Figure 2, where in the pristine state, all
polymers exhibited two main absorption peaks, one with an
absorption onset around 1200 nm and one with an absorption
onset around 590 nm. The UV−vis absorption profiles of the
three semiconductors also showed a significant difference in
terms of vibronic features. In fact, the absorption profile of N-
N displayed a maximum absorption wavelength (λmax,film) at
915 nm and no well-resolved fine structure. On the other hand,
both anthracene containing polymers, A-N and A-A,
demonstrated more resolved fine structures with A-N and A-
A displaying low energy shoulders, which in turn resulted in a
red-shift of their thin film absorption maxima to 1039 and
1064 nm, respectively.
Addition of 25 mol % dopant, specifically 4-(2,3-dihydro-
1,3-dimethyl-1H-benzimidazol-2-yl)-N,N-dimethylbenzen-
amine (N-DMBI), was used to gain insights into the doping
process of the polymers. N-DMBI was chosen to solution-dope
the polymers, thus rendering the doping process compatible
with low cost roll-to-roll printing techniques.33 As highlighted
in Figure 2, introduction of N-DMBI in each polymer film
resulted in the appearance of a new absorption feature with an
absorption maximum at approximately 1900 nm, which we
ascribed to a charged species of the polymers. Given that the
maximum absorption wavelength of this optical transition
remained unchanged at the various doping levels, see Figure
POLYMER CHARACTERIZATION
■
The energy levels of the polymers were investigated through a
combination of low energy inverse photoemission spectrosco-
py (LE-IPES), ultraviolet photoelectron spectroscopy (UPS),
and UV−vis−NIR spectroscopy and are summarized in Table
1. LE-IPES, UPS, and UV−vis−NIR spectroscopy were
Table 1. Summary of the Energy Levels and Optical
Properties of the Polymers under Investigation
a
b
c
polymer
IP (eV)
EA (eV)
λmax,film (nm)
Eg (eV)
N-N
A-N
A-A
5.34
5.29
5.18
3.94
3.83
3.72
915
1039
1064
1.04
1.03
1.03
a
Extracted from the HOMO region of the UPS spectrum.
b
c
Determined from LE-IPES. Calculated from the onset of absorption
in thin film samples.
employed to determine the electron affinity (EA), ionization
potential (IP), and optical gaps (Eg) of the polymers,
respectively; see Figure S1 and Figure 2. The UPS and LE-
IPES spectra onsets were assigned as 2σ from the leading-edge
Gaussian fitted peak.30−32 As shown in Table 1, all three
polymers exhibited a similarly small Eg, while the EA of the
polymers became progressively shallower, moving from 3.94 to
3.72 eV, upon increasing the size of the aromatic core in the
polymer backbone, i.e., upon going from an all naphthalene to
a mixed naphthalene−anthracene to an all anthracene polymer
backbone. This was attributed to the effective dilution of the
electron withdrawing properties of the four lactam moieties per
polymer repeat unit, consequently rendering the larger
anthracene containing systems less electron deficient compared
Figure 2. UV−vis−NIR thin film absorption spectra of the pristine and doped (a) N-N, (b) A-N, and (c) A-A and normalized doped absorption
spectra of (d) N-N, (e) A-N, and (f) A-A.
C
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S2, we tentatively attributed this transition to the polaronic
species of the various semiconductors.
(100) peak of N-N appeared largely in-plane (along Qxy), while
the (100) peak of A-A appeared predominantly out-of-plane
(along Qz). The A-N polymer shows a somewhat intermediate
character suggesting a transition from a predominantly face-on
stacking for N-N to a predominantly edge-on stacking for A-A.
These changes in texture are more obvious from the horizontal
and vertical line profiles; see Figure S7 in the Supporting
Information. While this could be attributed to an intrinsic
material property, an alternative possibility, corroborating
previous literature reports, is that the progressively lower
solubility of the polymers in the processing solvent upon acene
core expansion leads to preaggregation in solution, forming
lamella sheets, which deposit in the plane of the substrate, with
a preferential edge-on deposition of the backbone, orthogonal
to the lamella.34,35
The π−π stacking peak for the three materials occurred at
17.4 nm−1 for N-N, 17.3 nm−1 for A-N, and 17.6 nm−1 for A-A,
corresponding to similar π−π stacking distances of ∼3.6 Å for
all polymers. It is noted that the A-N polymer exhibits broader
lamellar stacking peaks indicating a smaller crystallite size and/
or increased lattice disorder. In particular, the full width at half-
maximum (fwhm) values of the in-plane (100) peak were
0.018 Å−1, 0.033 Å−1, and 0.021 Å−1 for N-N, A-N, and A-A,
respectively.
The thermal properties of the polymers were investigated by
both thermogravimetric analysis (TGA) and differential
scanning calorimetry (DSC). The TGA and DSC traces of
N-N, A-N, and A-A can be found in Figures S3−S6 in the
Supporting Information. All of the polymers displayed
excellent thermal stabilities showing no signs of decomposition
up to 360 °C, thus suggesting their suitability for thermo-
electric applications. Moreover, the DSC traces recorded for all
of the polymers were featureless, highlighting the absence of
any significant morphological changes upon heating or cooling
of the samples.
To gain further insights into the morphological properties
and the relative degree of order of the various polymers,
grazing-incidence wide-angle X-ray scattering (GIWAXS) was
performed on polymer films cast on silicon substrates. The
two-dimensional scattering patterns and one-dimensional line
profiles are illustrated in Figure 3.
To complement the experimentally recorded data, density
functional theory (DFT) calculations were carried out on
tetramers of the various polymers employing either an
optimally tuned ωB97XD functional and 6-31G* as the basis
set (geometry optimizations) or a B3LYP-D3 functional and 6-
31G* as the basis set (torsional potentials). The DFT
optimized structures of the three tetramers are shown in
Figures S11, S13, and S15 in the Supporting Information. As is
common in the field, the long solubilizing alkyl chains were
replaced by the shorter methyl groups to simplify the
simulations.
First, geometry optimizations of the tetramers were
performed, revealing that the anti configuration, i.e., the one
in which the carbonyl moieties of neighboring bis-isoindigo
units are oriented in opposite directions, is the thermodynami-
cally most favorable configuration; see Figure S8 in the
Supporting Information. Closer evaluation of the 120−180°
region of the torsional potential energy surface (see Figure 4)
showed an energetic minimum at 165°, thus indicating the
almost planar configuration assumed by the conjugated
polymer backbone for the three polymers. Also note the
relatively shallow torsional potential energy surface (PES)
incurred at 165° by each of the three polymers, which enables
dihedral angle fluctuations of 15° at room temperature. This
feature, in combination with intermolecular packing inter-
actions between neighboring polymer chains, is likely to enable
the polymers to assume an even more planar configuration in
the solid state, thus benefiting orbital overlap and in turn also
polaron delocalization and intermolecular charge transport of
these materials.
Single point DFT calculations on the optimized geometries
were performed to determine the MO distribution of the
neutral polymer LUMO and the singly occupied molecular
orbital (SOMO) of the anionic polymer over the eight acene
fragments comprising the tetramers of the various polymers. As
shown in Figure 4, the neutral polymer LUMOs were primarily
distributed on the central four acene units of all the polymers.
The slight difference in the shape of the MO contribution trace
of A-N compared to the ones of N-N and A-A was due to the
Figure 3. GIWAXS data of the three polymers: two-dimensional
scattering patterns of (a) N-N, (b) A-N, and (c) A-A. (d)
Comparison of the azimuthally integrated one-dimensional scattering
profiles recorded for the three polymers. Profiles are offset vertically
for clarity.
As shown in Figure 3, all polymers exhibited a semicrystal-
line microstructure with both lamellar stacking and π−π
stacking peaks evident. Despite having identical side chains, the
lamellar stacking peaks of the three polymers appeared at
slightly different values of Q, corresponding to different
lamellar stacking distances. For N-N the (100) peak appeared
at 2.29 nm−1, corresponding to a d-spacing of 27.5 Å. The
(100) peak of A-N in comparison occurred at 2.25 nm−1
(corresponding to a d-spacing of 28.0 Å), while the (100) peak
of A-A was recorded at 2.44 nm−1 (corresponding to a d-
spacing of 25.8 Å). This indicates that the A-A polymer had
the shortest lamellar stacking distance which is likely afforded
by the longer backbone repeat unit distance, which in turn may
allow for some degree of interdigitation or space-filling of the
side chains. The films also exhibited differences in texture: the
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Figure 4. (a) Torsional potentials around the centermost CC bond in the neutral tetramers of N-N, A-N, and A-A focusing on the 120−180°
region. Computed MO distributions of the (b) neutral LUMO and (c) anionic polymer SOMO for the tetramers of N-N, A-N, and A-A. The MO
distribution traces of the neutral LUMO of N-N and A-A are overlapping.
Figure 5. Organic field effect transistor output curves for (a) N-N, (b) A-N, and (c) A-A and transfer curves for (d) N-N, (e) A-N, and (f) A-A,
respectively.
mixed naphthalene and anthracene cores present within the A-
N polymer. As seen in Figure 4, the introduction of two
different acene cores resulted in an asymmetric LUMO
distribution of the A-N polymer, which in turn was due to
the stronger contributions from the lactam groups in the
central two naphthalene cores to the LUMO in comparison to
those in the anthracene cores. These findings therefore indicate
that the naphthalene bis-oxindole unit had a more electron
deficient character compared to its anthracene counterpart,
thus corroborating the findings from UPS and LE-IPES.
We envisaged the SOMO of the anionic polymer
distribution to be of particular interest given its analogy to
the electron polaron of the corresponding polymers. The MO
distribution of the anionic polymer SOMO showed more
localized character than the neutral LUMO, but significant
differences were observed across the three polymers (see
Figure 4), with the N-N polymer showing the broadest
distribution and the A-A polymer the narrowest one. To better
quantify the width of the MO distribution curve and use it as
an indicator of the delocalization of the anionic polymer
SOMO and hence polaron, we calculated the full width at half-
maximum (fwhm) of the MO distribution curves. The fwhm
values computed for the N-N, A-N, and A-A polymers were 2.4
fragments, 2.3 fragments, and 2.0 fragments, respectively, thus
revealing the anionic polymer SOMO of N-N to be the most
delocalized and that of the A-A polymer to be the least
delocalized.
Hirshfeld atomic charges of each atom of the neutral
polymer were subtracted from the equivalent atomic charges of
the anionic polymer, leading to a visualization of where the
extra charge upon reduction was found along the polymer
chain. Within the fragment analysis, there is strong agreement
with the MO contributions. Plotting on an atom-by-atom basis
(see Figure S18 in the Supporting Information) shows that
most of the charge is located on only a few atoms, although the
same trends in charge distribution can be observed.
Interestingly, the mixed A-N polymer shows two similar
environments bearing similar amounts of charge. Most of this
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charge is stabilized by the bridging CC double bond carbon
atoms between fragments, and the carbonyl groups; see Figure
S19 in the Supporting Information. The comparison between
N-N and A-A is also interesting in that the fragment approach
suggests that the N-N polymer has a greater degree of
delocalization. Turning to a distance comparison, A-A shows
similar delocalization lengths to N-N, but N-N spreads the
charge over more repeat units and over more of the more
stabilizing CC and CO groups. Thus, in addition to
changes in energy levels relevant to electronic doping, the
change from naphthalene to anthracene also spaces out the
atoms capable of stabilizing the excess charge, leading to
localization on only a few groups. Consequently, this suggested
the polaron in the N-N polymer to be the most delocalized and
hence N-N to be most suited toward thermoelectric
applications, given that previous studies have indicated a
strong relationship between polaron delocalization and
electrical conductivity.16
Lastly, the energy levels and optical properties of the
semiconductors were computed. The IP values calculated for
N-N, A-N, and A-A were 5.34 eV, 5.14 eV, and 5.11 eV, thus
reflecting the trend recorded in the UPS measurements. A
good fit between the experimental and computational trends
obtained for the EA of the polymers was also observed with the
simulated EA for N-N being the largest (3.21 eV), for A-N
being in the middle (3.16 eV), and for A-A being the smallest
(3.11 eV).
Table 2. Organic Field Effect Transistor and Thermoelectric
Performance Summary of the Polymers under Investigation
a
a
μe,lin
polymer (cm2 V−1 s−1
μe,sat
σmax
PFmax
)
(cm2 V−1 s−1
)
(S cm‑1)
(μW m−1 K−2
)
N-N
A-N
A-A
0.20
0.15
0.05
0.33
0.23
0.07
0.65
0.26
0.018
3.2
1.6
0.25
a
Obtained from TGBC OFET measurements.
delocalization led to polymers with a higher σmax. This in turn
further supported previous literature evidence of a correlation
between the degree of delocalization of the electron polaron
and the electrical conductivity of materials.16
Next, the thermopower of the three semiconductors was
measured. For all three members of the series negative
thermopower values were obtained, hence indicating that
electrons were the majority carrier within each of the polymers.
The Seebeck coefficient for N-N, A-N, and A-A was highest at
a dopant ratio of 3 mol % and decreased progressively upon
increasing the dopant ratio. At each N-DMBI ratio, N-N and
A-A displayed the lowest and highest S, respectively.
As anticipated and highlighted in Figure 6, the evolution of σ
and S as a function of dopant ratio followed an opposite trend,
thus signifying that the maximum power factor for the
polymers was obtained at a dopant concentration that would
best balance the two quantities. The maximum PF values of the
three semiconductors were recorded at different dopant
concentrations with N-N affording the highest PF of 3.2 μW
m−1 K−2 at 9 mol % N-DMBI concentration, A-N a PF of 1.6
μW m−1 K−2 at 15 mol % N-DMBI content, and A-A a PF of
0.25 μW m−1 K−2 at 25% dopant ratio. The PF values recorded
for N-N and A-N are among the highest recorded for n-type
OTE materials to date, while the PF recorded for A-A is still
higher compared to the thoroughly studied and optimized
N2200 polymer.14,15,17,19,20,24
With the establishment of the suitability of the series of
ladder type semiconducting polymers, their electrical con-
ductivity, σ, which mathematically is given in eq 1, was
investigated next.
σ = neμ
(1)
where n is the number density of charge carriers, e the
elementary charge, and μ the charge carrier mobility.
As highlighted in Figure 6 the electrical conductivity for each
polymer is lowest in the absence of any N-DMBI. However,
increasing the N-DMBI concentration resulted in an increase
in the conductivity with σ increasing over 4 orders of
magnitude for each polymer. The maximum electrical
conductivity (σmax) recorded for N-N, A-N, and A-A were
0.65 S cm−1, 0.26 S cm−1, and 0.018 S cm−1 respectively. The
recorded trend in σmax therefore is consistent with the trend in
the electron mobility (μe) for the three polymers, measured in
organic field effect transistor (OFET) devices that featured a
top-gate bottom-contact (TGBC) architecture; see Figure 5.
The electron mobilities recorded in the linear (μe,lin) and
saturation (μe,sat) regimes for the three polymers were 0.20 cm2
V−1 s−1 and 0.33 cm2 V−1 s−1 for N-N, 0.15 cm2 V−1 s−1 and
0.23 cm2 V−1 s−1 for A-N, and 0.05 cm2 V−1 s−1 and 0.07 cm2
V−1 s−1 for A-A, respectively, thus highlighting that N-N and
A-N had comparable mobilities while further increasing the
acene core size was detrimental for charge carrier transport; see
Table 2.
Comparing the differences recorded in the σmax with the
differences recorded in the μe for the polymers led to the
conclusion that the difference in the polymers’ μe is not
sufficient to fully explain the difference in σmax. In fact, the
larger difference in σmax suggests that increasing the acene core
size also results in lower charge carrier densities, n, generated
during the doping process. Finally, the recoded σmax trend also
followed the trend in the simulated anionic polymer SOMO
delocalization, whereby increasing the anionic polymer SOMO
CONCLUSION
■
In summary, three new n-type fused lactam semiconducting
polymers only differing in the relative size of their constituent
acene cores were synthesized for thermoelectric and transistor
applications. Compared to traditional transition-metal cata-
lyzed polymerization reactions for conjugated polymers, the
materials developed herein were obtained via a convenient
transition-metal free aldol polycondensation reaction, thus
avoiding the use of expensive palladium-based catalysts and
any highly toxic organometallic synthetic intermediates. TGA
measurements indicated excellent thermal stability of each
polymer, with none of the materials showing any sign of
decomposition up to 360 °C. Experimental and computational
energy level analysis of the three polymers on the other hand
demonstrated that reducing the central acene core size from
two anthracene (A-A) to mixed naphthalene−anthracene (A-
N) and two naphthalene cores (N-N) progressively resulted in
larger EA, thereby suggesting an increasingly more favorable
and efficient doping process when employing N-DMBI as the
dopant. In parallel, both OFET mobility data and DFT
simulations also suggested the benefits of acene core
contraction on the thermoelectric performance, with the N-
N polymer affording the highest electron mobility and the
most delocalized electron polaron and the A-A polymer
incurring the lowest values. Ultimately, these concepts were
clearly reflected in the thermoelectric performance data
F
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Figure 6. (a) Electrical conductivities, (b) Seebeck coefficients, and (c) power factors recorded for N-N, A-N, and A-A as a function of N-DMBI
ratio employed.
recorded for the polymers with N-N, A-N, and A-A affording
maximum electrical conductivities and power factors of 0.65 S
cm−1 and 3.2 μW m−1 K−2, 0.26 S cm−1 and 1.6 μW m−1 K−2,
and 0.018 S cm−1 and 0.25 μW m−1 K−2 when doped with N-
DMBI respectively. The PF values recorded for N-N and A-N
are thus among the highest reported in the literature for n-type
polymers. Importantly, the results reported in this study
highlight that modulating the size of the central acene ring is a
highly effective molecular design strategy to optimize the
thermoelectric performance, thus also providing new insights
into the molecular design guidelines for the next generation of
high-performance n-type materials for thermoelectric applica-
tions.
Maximilian Moser − Department of Chemistry and Centre for
Plastic Electronics, Imperial College London, London W12
0BZ, United Kingdom; Department of Chemistry, University
of Oxford, Oxford OX1 3TA, United Kingdom;
orcid.org/0000-0002-3293-9309;
Email: maximilian.moser@chem.ox.ac.uk
Authors
Suhao Wang − Laboratory of Organic Electronics, Department
of Science and Technology, Linköping University, Norrköping
SE-60174, Sweden; orcid.org/0000-0002-6295-7639
Cameron Jellett − Department of Chemistry and Centre for
Plastic Electronics, Imperial College London, London W12
0BZ, United Kingdom; orcid.org/0000-0003-2811-3683
Karl Thorley − Department of Chemistry, University of
Kentucky, Lexington, Kentucky 40506-0055, United States;
orcid.org/0000-0003-0665-3363
ASSOCIATED CONTENT
■
sı
* Supporting Information
The Supporting Information is available free of charge at
George T. Harrison − Physical Science and Engineering
Division, King Abdullah University of Science and Technology
(KAUST), Thuwal 23955-6900, Saudi Arabia
Xuechen Jiao − Department of Materials Science and
Engineering, Monash University, Clayton, Victoria 3800,
Australia; orcid.org/0000-0001-7387-0275
Mingfei Xiao − Department of Physics, University of
Cambridge, Cambridge CB2 1TN, United Kingdom
Balaji Purushothaman − Physical Science and Engineering
Division, King Abdullah University of Science and Technology
(KAUST), Thuwal 23955-6900, Saudi Arabia
Maryam Alsufyani − Physical Science and Engineering
Division, King Abdullah University of Science and Technology
(KAUST), Thuwal 23955-6900, Saudi Arabia
https://pubs.acs.org/doi/10.1021/jacs.0c10365.
General procedures, synthetic details, UV−vis−NIR
spectroscopy, low energy inverse photoemission spec-
troscopy and ultraviolet photoelectron spectroscopy,
thermogravimetric analysis, differential scanning calo-
rimetry, grazing-incidence wide-angle X-ray scattering,
and computational simulation data (PDF)
AUTHOR INFORMATION
■
Corresponding Authors
Hu Chen − Physical Science and Engineering Division, King
Abdullah University of Science and Technology (KAUST),
Thuwal 23955-6900, Saudi Arabia; Email: hu.chen@
kaust.edu.sa
Helen Bristow − Department of Chemistry and Centre for
Plastic Electronics, Imperial College London, London W12
G
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Andrew Wadsworth − Department of Chemistry and Centre
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Henning Sirringhaus − Department of Physics, University of
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Simone Fabiano − Laboratory of Organic Electronics,
Department of Science and Technology, Linköping University,
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7016-6514
Iain McCulloch − Physical Science and Engineering Division,
King Abdullah University of Science and Technology
(KAUST), Thuwal 23955-6900, Saudi Arabia; Department
of Chemistry, University of Oxford, Oxford OX1 3TA,
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Complete contact information is available at:
https://pubs.acs.org/10.1021/jacs.0c10365
Author Contributions
∞H.C. and M.M. contributed equally to this work.
Notes
The authors declare no competing financial interest.
ACKNOWLEDGMENTS
■
The authors acknowledge generous funding from KAUST for
financial support. The research reported in this publication was
sponsored by funding from King Abdullah University of
Science and Technology Office of Sponsored Research (OSR)
under Awards OSR-2018-CARF/CCF-3079, OSR-2015-
CRG4-2572, and OSR-4106 CPF2019. We acknowledge EC
FP7 Project SC2 (610115), EC H2020 (643791), and EPSRC
Projects EP/G037515/1, EP/M005143/1, and EP/L016702/
1. This work was performed in part at the SAXS/WAXS
beamline at the Australian Synchrotron, part of ANSTO.36 S.F.
acknowledges financial support from the Swedish Research
Council (Grant 2016-03979), ÅForsk (Grants 18-313, 19-
310), Olle Engkvists Stiftelse (Grant 204-0256), and the
̈
Advanced Functional Materials Center at Linkoping University
(Grant 2009-00971).
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Andreasen, J. W.; Thiel, W.; Berggren, M.; Crispin, X.; Fazzi, D.;
Fabiano, S. Thermoelectric Properties of Solution-Processed n-Doped
Ladder-Type Conducting Polymers. Adv. Mater. 2016, 28, 10764−
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