Synthetic principles directing charge transport in low-band-gap dithienosilole-benzothiadiazole copolymers

Article

pubs.acs.org/JACS

Synthetic Principles Directing Charge Transport in Low-Band-Gap

Dithienosilole−Benzothiadiazole Copolymers

†

,⊥

‡

†

∥

Pierre M. Beaujuge, Hoi Nok Tsao, Michael Ryan Hansen, Chad M. Amb, Chad Risko,

Jegadesan Subbiah, Kaushik Roy Choudhury, Alexei Mavrinskiy, Wojciech Pisula, Jean-Luc Bred

́

as,^∥

§

‡

§

‡

,†,∥

Franky So, Klaus Mu

̈

llen, and John R. Reynolds*

†

The George and Josephine Butler Polymer Research Laboratory, Center for Macromolecular Science and Engineering, Department

of Chemistry, University of Florida, Gainesville, Florida 32611, United States

‡

Max Planck Institute for Polymer Research, Ackermannweg 10, D-55128 Mainz, Germany

§

Department of Materials Science and Engineering, University of Florida, Gainesville, Florida 32611, United States

∥

Center for Organic Photonics and Electronics and School of Chemistry and Biochemistry, Georgia Institute of Technology, Atlanta,

Georgia 30332, United States

⊥

King Abdullah University of Science and Technology (KAUST), Thuwal 23955-6900, Saudi Arabia

S Supporting Information

*

ABSTRACT: Given the fundamental diﬀerences in carrier

generation and device operation in organic thin-ﬁlm transistors

(

OTFTs) and organic photovoltaic (OPV) devices, the

material design principles to apply may be expected to diﬀer.

In this respect, designing organic semiconductors that perform

eﬀectively in multiple device conﬁgurations remains a

challenge. Following “donor−acceptor” principles, we de-

signed and synthesized an analogous series of solution-

processable π-conjugated polymers that combine the elec-

tron-rich dithienosilole (DTS) moiety, unsubstituted thiophene spacers, and the electron-deﬁcient core 2,1,3-benzothiadiazole

(

BTD). Insights into backbone geometry and wave function delocalization as a function of molecular structure are provided by

density functional theory (DFT) calculations at the B3LYP/6-31G(d,p) level. Using a combination of X-ray techniques (2D-

WAXS and XRD) supported by solid-state NMR (SS-NMR) and atomic force microscopy (AFM), we demonstrate fundamental

correlations between the polymer repeat-unit structure, molecular weight distribution, nature of the solubilizing side-chains

appended to the backbones, and extent of structural order attainable in p-channel OTFTs. In particular, it is shown that the

degree of microstructural order achievable in the self-assembled organic semiconductors increases largely with (i) increasing

molecular weight and (ii) appropriate solubilizing-group substitution. The corresponding ﬁeld-eﬀect hole mobilities are enhanced

2

−1 −1

by several orders of magnitude, reaching up to 0.1 cm V

s

with the highest molecular weight fraction of the branched alkyl-

substituted polymer derivative in this series. This trend is reﬂected in conventional bulk-heterojunction OPV devices using

−3

2

−1 −1

PC BM, whereby the active layers exhibit space-charge-limited (SCL) hole mobilities approaching 10 cm V s , and yield

71

improved power conversion eﬃciencies on the order of 4.6% under AM1.5G solar illumination. Beyond structure−performance

correlations, we observe a large dependence of the ionization potentials of the polymers estimated by electrochemical methods

on polymer packing, and expect that these empirical results may have important consequences on future material study and

device applications.

INTRODUCTION

While π-conjugated semiconductors are ﬁnding widespread

application in a broad range of operating systems, such as thin-

most favorable morphologies in bulk-heterojunction (BHJ)

solar cell devices. In general, organic semiconductors designed

for a speciﬁc device application tend to not be suﬃciently

eﬀective in other device conﬁgurations.

An excellent illustration of the complexity encountered in

attempting to combine (i) solution processability, (ii) high

charge-carrier mobilities, (iii) wide-ranging optical absorption

proﬁles and large absorption coeﬃcients, (iv) appropriate

energy band structure, and (v) environmental stability with the

■

1

2

ﬁlm transistors (OTFTs), sensors, photovoltaic devices

3

4

(

OPVs), electrochromic displays, or light-emitting diodes

5

OLEDs), material performance and charge-transport require-

ments can vary substantially from one operating system to

another. For example, while most existing organic semi-

conductors with substantial eﬃciencies in photovoltaic devices

do not necessarily fulﬁll the charge-carrier mobility require-

ments for thin-ﬁlm transistor applications, those with packing-

dominated charge-transport characteristics may not induce the

Received: February 26, 2012

Published: May 18, 2012

©

2012 American Chemical Society

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Journal of the American Chemical Society

Article

same material is represented by the thieno[3,2-b]thiophene-

based all-donor polymer analogues developed by McCulloch et

al. and others, namely, poly(2,5-bis(3-alkylthiophen-2-yl)-

beyond 600 nm), thus minimizing the external quantum

eﬃciency (EQE) at longer wavelengths. Along these lines, it is

worth noting that PBTTT possesses a substitution pattern of

solubilizing side-chains, allowing size-controlled intercalation of

6

thieno[3,2-b]thiophene)s (PBTTT) (Chart 1). With ﬁeld-

9

fullerenes. Besides its plausible consequences on polymer

Chart 1. Molecular Structures of PBTTT and PDPP3T

backbone organization, the intercalation of nanometer-sized

molecular objects is suspected to govern the optimized

composition of the photovoltaic blend, generally toward higher

9

concentration of fullerene, hence possibly disfavoring the light-

harvesting capability of the blend. A last critical parameter

limiting solar cell performance could be its small ionization

potential (IP, 5.1 eV relative to vacuum), inherent to most all-

thienylene conjugated polymers, which contributes to minimiz-

ing the device open-circuit voltage (V_O_C) and in turn further

limits solar cell eﬃciency. Overall, as long as a substantial

optical density for photovoltaic blends can be extended to

longer wavelengths and maintained over a broad spectral range,

improving the charge dynamics, including carrier mobilities,

achieved with polymers in thin ﬁlms could enhance both their

OTFT and OPV performance.

eﬀect hole mobilities as high as 0.6 cm V_s−1 and current on/

oﬀ ratios on the order of 10 in bottom-gate/bottom-contact

2

−1

7

transistor conﬁgurations, PBTTTs stand out as p-channel

semiconductors when processed from their mesophase. This

level of performance (carrier mobilities approaching that of

amorphous silicon) can be obtained under controlled thermal

annealing conditions inducing the formation of large crystalline

7c

In recent years, low-band-gap polymers with alternating

electron-rich and electron-deﬁcient building units along the

same backbone (hence following principles introduced earlier

10

6

,7

by Havinga and co-workers) have become especially useful in

both thin-ﬁlm transistors and BHJ photovoltaic devices with

fullerenes. These polymers provide the ability to span the

optical absorption spectrum across the visible and into the near-

IR, and improved stability to ambient electrochemical oxidative

processes as larger IPs (i.e., low-lying “HOMO” energy levels)

can be introduced in conjunction with larger electron aﬃnities

(EAs) (i.e., low-lying “LUMO” energy levels). Among these, a

domains with nanometer-scale substructures (of ca. 10 nm)

that translate into a high level of microstructural organization.

At the same time, in spite of their semicrystalline character,

PBTTTs have been shown to possess limited power conversion

eﬃciencies on the order of 2.3% in BHJ solar cells when

8

blended with PC BM. Diﬀerences between the in-plane and

6

1

out-of-plane charge-carrier mobilities are commonly held

responsible for photovoltaic performance limitations. These

diﬀerences are typically attributed to a large extent of edge-on

molecular alignment expected to reduce the dimensionality of

charge transport across the active layer in highly crystalline

semiconducting polymers. Another limiting factor is repre-

sented by the lack of spectral coverage, as in PBTTT, for which

the onset of absorption lies in the 650−680 nm range while its

spectral absorption peaks at ca. 540 nm. The capture of light is

in turn restricted to the most energetic photons as opposed to

being extended to where the photon ﬂux is maximum (i.e.,

1

poly(diketopyrrolopyrrole-terthiophene) (PDPP3T) with

alternating electron-deﬁcient diketopyrrolopyrrole cores and

electron-rich unsubstituted thiophenes (Chart 1) demonstrates

−

2

near-balanced hole and electron mobilities on the order of 10

2

−1 −1

cm V

s

in bottom-gate/bottom-contact OTFTs. At the

same time, PDPP3T exhibits a photoresponse that extends into

the near-IR wavelengths (absorption onset at ca. 900 nm),

providing for short-circuit current densities on the order of 12

mA cm⁻and power conversion eﬃciencies over 4.5% in BHJ

2

a

Chart 2. Design Principles Used for the Synthesis of P1−P5

a

13

On the left: the all-donor TS6, TS6T1, and TS6T2 as originally developed by Marks and co-workers for OTFT applications. On the right:

derivatization of the TS6 backbones with the electron-deﬁcient 2,1,3-benzothiadiazole (BTD) unit to lead to donor−acceptor repeat unit structures

narrowing the polymer optical gap.

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Table 1. GPC Estimated Molecular Weights of the DTS-BTD Copolymers P1−P5 (from TCB), TGA-Estimated Decomposition

Temperatures (°C), Local Absorption Maxima for the Polymers (in toluene solution and in thin ﬁlms), and π-Stacking (π, nm)

and Chain-to-Chain Lamellar Distances (d, nm) by 2D-WAXS

M (kDa) (PDI)

λ_a_b_s(nm) in toluene

λ_a_b_s(nm) thin ﬁlm

2D-WAXS

d (nm)

n

polymer

TCB

T_d(°C)

422

1

2

1

2

π (nm)

P1.1

P1.2

P2.1

P2.2

P3

10.3 (3.6)

12.7 (6.3)

18.3 (6.1)

36.1 (6.0)

16.3 (8.2)

16.0 (8.0)

18.3 (3.3)

26.4 (6.5)

392

408

577

404

425

459

462

464

483

427

583

n/a

1.95

1.80

a

653 (770)

627

672 (770)

643

0.35

0.36

n/a

b

c

429 (312)

458

(1.95/1.87 )

1.75

469

680

653

424

426

451

602

611

1.94

P4

∼480

429

617

645

0.36

1.84

a

P5.1

P5.2

703 (764)

697 (764)

1.70

429

a

b

c

Shoulder. Unexpected signature for early degradation. Fiber was annealed at 200 °C for 40 min.

1

12

solar cells. PDPP3T and analogues that combine substantial

transistor mobilities and solar cell eﬃciencies suggest that

polymer thin-ﬁlm device performance could be improved by

building a more fundamental understanding of the complex

interplay between molecular structure and intrinsic charge-

transport properties in donor−acceptor systems.

means of developing fundamental understanding of the

correlation between charge transport and OPV eﬃciency.

Knowing the optimized polymer/PCBM blend ratio for each

polymer derivative was instrumental to measuring the SCL hole

mobilities in hole-only devices composed of a BHJ of the

corresponding donors and acceptor. The control solar cells

showed PCEs on the order of 4.6% in the absence of small

molecule additive-assisted device processing (i.e., no morphol-

ogy control); this value is in excellent agreement with that

previously reported for the same polymer (i.e., 4.7% in

With the goal to identify synthetic principles that impart

particularly large carrier mobilities in narrow band gap

semiconducting polymers, we designed a series of π-conjugated

systems (P1−P5, Chart 2) that incorporate distinct ratios of

electron-rich to electron-deﬁcient moieties in the monomer

repeat unit, along with variations (i.e., concentration and

bulkiness) of the solubilizing side-chains along the backbones.

P1−P5 contain a combination of electron-rich substituents

14a,15

average),

and approaching the NREL-certiﬁed eﬃciency

of ca. 5.2% recently reported using an inverted OPV device

3

d,14b,16

conﬁguration.

13

RESULTS AND DISCUSSION

Polymer Synthesis. The sequence of Pd-mediated Stille

couplings used to access the polymers P1−P5 is shown in

Scheme S1, along with a detailed description of our synthetic

protocols (see the Supporting Information).

examined earlier by Marks and co-workers in OTFTs

■

composed of dithienosilole (DTS) and unsubstituted thio-

phene spacers. This backbone of electron-rich units was

modiﬁed by incorporating the electron-deﬁcient core 2,1,3-

benzothiadiazole (BTD) in order to (i) study the eﬀect of the

donor−acceptor motif on intrinsic charge transport, and (ii)

narrow signiﬁcantly the optical gap of the corresponding

polymers, in turn shifting their absorption spectra toward

longer wavelengths.

Table 1 shows the number-average molecular weights as

determined by high temperature GPC (using trichlorobenzene

(TCB) as eluent), the temperatures where the onset of thermal

decomposition occurs as determined by TGA (under N ), and

2

In this contribution, we correlate structural variations

imparted to the polymer backbones with their charge-transport

properties in thin ﬁlms, using OTFT and OPV devices as

platforms for the study. In particular, using a conjunction of X-

ray techniques (2D-WAXS and XRD) supported by solid-state

NMR (SS-NMR) and atomic force microscopy (AFM), we

emphasize the key role of molecular weight distribution, and

the nature of the solubilizing substituents appended to the

backbones on the degree of microstructural organization that

can be attained with each polymer. While these variables add a

higher level of complexity to the equation of design principles

governing the performance of polymers in devices, it is

increasingly apparent that small structural changes can have a

tremendous impact on both charge transport and photovoltaic

the local absorption maxima in thin ﬁlm and solution (toluene)

for various batches of P1−P5. These polymers exhibit excellent

thermal stability as evidenced by decomposition temperatures

on the order of 400 °C (except for an early sign of degradation

at ca. 300 °C in the case of P2). Noticeably, two substantial

polymer fractions could be isolated by Soxhlet extraction for all

of the polymers produced (except for P3 and P4, inherently

limited by solubility): a low molecular weight fraction soluble in

chloroform (PX.1) and a high molecular weight fraction only

soluble in hot chlorobenzene (PX.2). These fractions yield

various degrees of structural order and distinct device

performance, results that will be discussed in the subsequent

sections.

Fiber 2D X-ray Scattering Analysis. The bulk micro-

structural organization of the polymers was examined by 2D

wide-angle X-ray scattering analysis (2D-WAXS) from extruded

1

4

properties in thin-ﬁlm devices. Of all the DTS-BTD

backbones explored, the highest molecular weight fraction

17

(

P5.2) of the strictly alternating donor−acceptor derivative P5

ﬁbers obtained following previously reported methods. In

14c

functionalized with 2-ethylhexyl branched substituents stands

out by demonstrating up to 0.1 cm V

earlier work,

we showed how small structural changes

2

−1 −1

s

ﬁeld-eﬀect (hole)

imparted to the polymer repeat unit on going from P1.1, to

P2.1, to P3, and to P4 aﬀect the degree of microstructural

order achievable in the bulk and in thin ﬁlms of the respective

semiconducting polymers. The extent of “crystallinity” of each

polymer analogue was further correlated to their performance

in thin-ﬁlm transistors where a trend involving variables such as

mobility in bottom-gate/bottom-contact transistors, and by

approaching 10⁻cm V

3

2

−1 −1

s

space-charge-limited (SCL) hole

mobility in hole-only devices composed of a BHJ of the

polymer with the electron-acceptor PC BM. It is essential to

7

1

note that, in this study, control solar cell devices were used as a

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Journal of the American Chemical Society

Article

the relative concentration of electron-donating and -accepting

units along the backbone, or the extent of appended

solubilizing side-chains, arose across the series. Here, we

further underscore the key role of molecular weight

distribution, along with that of the nature of the solubilizing

substituents (linear vs branched), on the extent of achievable

microstructural organization in DTS-BTD donor−acceptor

systems.

by the equatorial nature of the Bragg reﬂections observed along

the Debye rings in the most ordered patterns (e.g., Figure 1b

and d). For the same reason, examining the wide angle

scattering regions, the lamellae show a strong propensity to π-

stack in a plane normal to the extrusion direction, rather than

along the ﬁber as is typically seen with small molecular weight

discotics forming columnar assemblies (and scattering in the

meridional direction).

Figures 1 and 2 (along with Figure S1) overview the 2D-

WAXS patterns obtained for P1−P5 as a function of their

Figure 1a and b shows the diﬀerences in ordering behavior

between two molecular weight fractions of the strictly

alternating copolymer P1, namely, P1.1 (M = 10.3 kDa in

n

TCB) and P1.2 (M = 12.7 kDa in TCB). Here, the ﬁrst

n

striking indication of a pronounced eﬀect of molecular weight

on the degree of observable structural order in the bulk is the

clear apparition of wide-angle reﬂections characteristic of

interchain π-stacking interactions (ca. 3.5 Å) on going from

P1.1 to P1.2. These reﬂections are not as pronounced in

intensity as those seen in the patterns of P2.1 and P4 (see

Figure S1), which could be a simple consequence of the

molecular weight limitation. In addition, the disruption of

isotropy along the smaller-angle Debye rings translates well into

a net increase in orientation between lamellae, and is

accompanied by a reduction of the d-spacing by 1.5 Å (varying

from 1.95 to 1.80 nm).

The same decrease in lamellar distance is also illustrated in

Figure 1c and d comparing two distinct molecular weight

fractions of polymer P2, namely, P2.1 (M = 18.3 kDa in TCB)

n

and P2.2 (M = 36.1 kDa in TCB). In this case, while the short

n

π-stacking distance of 3.6 Å is retained from one batch to the

other (and regardless of the extrusion conditions and thermal

annealing subjected to the ﬁber), the lamellar spacing dropped

from 1.87 to 1.75 nm. Visible in Figure 1d, the presence of a

higher-order anisotropic reﬂection at small angles is additional

evidence for the even more pronounced molecular-weight-

promoted ordering eﬀect in P2 in comparison to P1. Moreover,

a signiﬁcant distinction between P2 and P1 is represented by

the higher intensity π-stacking reﬂections of P2, which further

translates into a larger degree of relative chain-to-chain order.

Variations in microstructural organization with molecular

weight are expected given that the degree of molecular entropy

in conjugated polymers (e.g., entanglement) varies with average

chain lengths. As empirically observed in the case of P1 and P2,

shorter lamellar distances between backbones can be obtained

with higher average molecular weight fractions (see Table 1).

This could be explained by considering polymer chain-ends as

defects disrupting the self-assembling polymer lattices.

Following this reasoning, a lower-molecular-weight polymer

fraction contains a higher concentration of distinct chain-ends

and necessarily meets inherent limitations in terms of attainable

structural order. In contrast, the possibility that the degree of

entanglement promotes shorter molecular distances by

Figure 1. Fiber 2D wide-angle X-ray scattering (2D-WAXS) of (a)

P1.1 (extruded at 170 °C), (b) P1.2 (extruded at 170 °C), (c) P2.1

extruded at 170 °C and annealed at 200 °C), and (d) P2.2 (extruded

(

at 190 °C).

Figure 2. Fiber 2D wide-angle X-ray scattering (2D-WAXS) of P5.1

“

tightening” the polymer lattice is not consistent with the

(

extruded at 210 °C).

higher degree of chain orientation seen with the large molecular

weight fractions in general (as clearly evidenced by a

diminution of the scattering accounting for the lamellar

organization).

repeat unit structure, their molecular weight distribution, or as a

function of the postextrusion thermal annealing condition

subjected to the ﬁber (Figure S1), and highlight the diﬀerences

in attainable structural order across the polymer series. The X-

ray data, including π-stacking distances (when observable) and

interlamellar spacings, are summarized in Table 1.

In general, looking at the small angle scattering regions, P1−

P5 are prone to assemble in lamellar superstructures

preferentially oriented in the extrusion direction, as evidenced

Based on solution processability considerations, copolymers

P3 (M = 16.3 kDa) and P4 (M = 16.0 kDa) were examined

n

without further attempt to achieve higher molecular-weight

fractions. As illustrated in Figure S1e, and in contrast with the

higher molecular-weight fractions of P1 and P2, P3 revealed

relatively isotropic diﬀractions at small angles (d = 1.94 nm)

with the absence of observable π-stacking at wider angles. In

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1

comparison, and as shown in Figure S1f, the pattern of P4

couplings as in the case of H NMR in solids. In fact, the

observed line width diﬀerences for the isotropic chemical shifts

of a number of diﬀerent NMR active nuclei have been exploited

for a large range of compounds. This includes organic and

inorganic materials and has led to a better understanding of the

short and intermediate range order present in both semi-

(

bithiophene-substituted analogue of P3) reveals π-stacking

distances as low as 3.6 Å and fewer scattered small-angle

reﬂections, thus supporting a more pronounced relative

orientation of the backbones, and narrower lamellar distances

on the order of 1.84 nm.

As illustrated in Figure S1a−d, postextrusion thermal

annealing of the ﬁbers of P1.2 and P2.2 does induce some

additional ordering characteristics in the bulk polymers, which

is evidenced by the more pronounced π-stacking reﬂections,

though the eﬀect remained relatively modest.

21

crystalline and amorphous materials. Thus, combining the

information available from observing two diﬀerent nuclei ( H

and C), using MAS (magic angle spinning) NMR methods,

it is possible to get a more detailed picture of both local packing

1

13

22

arrangements and molecular order.

In parallel, comparing the highest molecular weight of P1

attained, namely, P1.2, and its branched 2-ethylhexyl-

substituted counterpart P5.1, whose X-ray scattering pattern

is shown as Figure 2, it appears that side-chain dynamics can

further promote the extent of achievable microstructural

organization both by imparting proper interchain spacing in

the solid state, and by introducing appropriate material

solubility. On one hand, in the case of P5.1, the bulkier nature

of the 2-ethylhexyl substituents (versus the linear octyls) allows

a higher degree of polymerization to be attained by preventing

premature precipitation during polymerization. As discussed

above, longer backbones can induce higher ordering

capabilities. On the other hand, and as reported in Table 1,

the shorter nature of the hexyl segment in 2-ethylhexyl versus

that of a linear octyl (calculated to be shorter by 1.3 Å at the

Figure 3 shows the solid-state NMR results obtained for the

high molecular-weight fraction of the DTS-BTD derivative

18

DFT-ωB97XD/6-31G(d,p) level) induces an interlamellar

distance of only 1.70 nm, which is signiﬁcantly shorter than the

1.80 nm spacing measured for the higher molecular weight

fraction of P1 (P1.2), and is also the shortest lamellar spacing

observed across this polymer series. Nonetheless, the 0.1 Å

increase in π-stacking distance on going from P1.2 to P5.1,

which could be explained by the presence of the 2-ethyl

substituent branching out of the core solubilizing chain, may

also be “loosening” the polymeric network and promoting the

favorable self-assembly of P5.1. In light of these empirical

results, 3.5 Å could therefore be interpreted as being a lower

threshold for material processing (i.e., closer stacking largely

reducing processability) and relatively unhindered nano-

structural organization in stackable π-conjugated polymers.

Overall, the distinct and particularly favorable ordering behavior

of P5.1 when compared to its linearly substituted counterparts

P1−P4 is expected to impact device performance (vide infra).

Solid-State NMR. Insights into the molecular order and

local packing arrangement of the DTS-BTD donor−acceptor

polymers can tentatively be produced from solid-state NMR by

Figure 3. Morphology changes in P5.2 probed by 2D solid-state

H− H DQ-SQ correlation and C{ H} MAS NMR after thermal

annealing (a, d) and before annealing at 200 °C (15 min) (b, c). A

cross-polarization time of 1.0 ms was used for the 1D spectra (a, b),

and a back-to-back recoupling and reconversion period of 40 μs was

employed for the 2D correlation spectra (c, d). All spectra were

recorded at 16.4 T using a fast spinning speed of 50.0 kHz.

1

13

1

probing the local molecular environments of proton ( H) and

13

19

carbon ( C) sites reﬂected in their chemical shift values. The

distinction relies on the fact that the packing of self-assembled

aromatic-based systems induced by π−π interactions is mainly

observed as a nucleus independent shift to higher or lower

chemical shift on the order of 1−10 ppm as compared to the

P5.2 (26.4 kDa) substituted with 2-ethylhexyl branched

solubilizing side chains, before and after sample annealing at

2

0

13

1

isotropic chemical shift (δ ) observed in solution NMR. This

200 °C. Before annealing, the C{ H} CP (cross-polar-

ization)/MAS spectrum (Figure 3b and c) exhibits a good

resolution in the region characteristic for the side chains (5−45

ppm). After annealing (Figure 3a and d), these resonances are

even more resolved (decrease in linewidths of about 40−50

Hz), a result suggesting that improved order of the side chains

has been achieved. Upon side-chain reorganization at elevated

temperatures, the π-conjugated backbones are expected to

iso

1

makes H NMR the obvious choice for probing this type of

interaction, since this nucleus has a much lower chemical shift

1

3

19

dispersion (∼15 ppm) as compared to C (∼220 ppm). The

structural order or disorder is, on the other hand, mainly

reﬂected in the line width of the observed isotropic chemical

shifts for the diﬀerent chemical moieties of the probed

1

9

13

molecule. For this reason, C is a more appropriate method

for probing local order since this nucleus has a larger chemical

shift range leading to a better chemical shift resolution.

Moreover, it is not inﬂuenced by homonuclear dipole−dipole

1

rearrange. The 2D H− H DQ-SQ (double quantum−single

quantum) correlation spectra shown in Figure 3c and d give

some indications on the evolution of the packing of the

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polymer backbone with thermal annealing. Here, improved

resolution with thermal annealing is apparent from the oﬀ-

diagonal cross peaks. These signals originate from the close

spatial proximity between the proton sites at the polymer

backbone and those of the attached 2-ethylhexyl side chains.

(toward lower frequencies) for the BTD protons in comparison

to that observed for P5.2 (Figure 3b); this result suggests

increased nonuniformity of packing of the BTD accepting units

in P2.1 in comparison to P5.2. The elongation to higher

frequency is most likely due to a DQ-SQ correlation arising

from the packing of acceptor-acceptor groups on top of one

another.

1

The improved H resolution is consistent with the better local

1

3

1

order observed in the C{ H} CP/MAS NMR spectra. Neither

of the 2D spectra in Figure 3c and d shows any signiﬁcant DQ-

SQ correlations between DTS and BTD, and DTS and DTS

groups. However, the donor−acceptor correlation (DTS-BTD)

is observable at longer recoupling time (Figure S2), in

agreement with our recent results obtained with an analogous

polymer backbone for which the donor units are bridged by a

Polymer Electrochemical Properties. Taking into

consideration the range of diﬀerences in intrinsic material

structural organization properties emphasized above, the thin-

ﬁlm electrochemical properties of P1−P5 were next to be

examined with speciﬁc attention to the possible interplay

between propensity to order and subtle/unexpected variations

in redox onsets. The redox properties of P1−P5 were

investigated via cyclic voltammetry (CV) and diﬀerential

pulse voltammetry (DPV) (see Figures S4−S10) with the

purpose of correlating the polymer oxidation and reduction

potentials and respective electrochemical energy gaps induced

by the successive structural modiﬁcations to the repeat unit

with the theoretical trends predicted at the DFT level. A three-

electrode electrochemical cell composed of a platinum disk

23

carbon center substituted with linear dodecyl side chains. In

Figure 3c and d, the resonance close to the diagonal from the

two diﬀerent BTD protons displays an elongated and slightly

split line shape, extending toward lower frequency. These

features suggest that the DTS-BTD polymer backbones are

indeed π-stacked, but also show that the BTD acceptor units

are heterogeneously packed throughout the sample; that is, the

acceptor groups are not always located at the exact same

position with respect to the neighboring polymer chains. These

results suggest that the molecular structure changes induced on

the donor, including bridging atom substitution and side-chain

substituent variations, are not a critical parameter inﬂuencing

polymer packing according to SS-NMR. We note that it has

been suggested in a recent study that the incorporation of

silicon bridging centers on the donor leads to an increase of the

angle between donor and acceptor groups, going from 10° in

carbon-bridged donor-containing systems to 19° in silicon-

2

working electrode (0.02 cm ), a platinum ﬂag counter

+

electrode, and a Ag/Ag reference electrode (10 mM AgNO ,

3

0.1 M Bu PF in MeCN) was employed, and all results were

4

6

+

subsequently calibrated to Fc/Fc (for consistency with values

reported in the literature). The estimated polymer energy levels

(IP and EA ) are presented in Table S1 (see the Supporting

Information) and provided relative to the vacuum level,

SS

24

considering that the SCE is 4.7 eV versus vacuum and Fc/

+

25

Fc is 0.38 eV vs SCE, that is, ∼5.1 eV relative to vacuum.

The polymer ﬁlms were cycled until they reached a stable and

reproducible redox response prior to characterization. The

respective polymer optical gaps as determined from the onset

of their low-energy transition (thin-ﬁlm value) are also

presented in Table S1.

14b

bridged donor-containing ones. This variation in backbone

geometry could explain why the appearance of the donor−

acceptor correlations are only observed at longer recoupling

times (see Figure S2).

In order to identify a possible molecular-weight dependence

of π-stacking and lamellar organization, two samples of P1 with

distinct molecular weight distributions were characterized by

solid-state NMR (P1.1, 10.3 kDa; and P1.2, 12.7 kDa). The

results from these investigations are summarized in Figure S3

It is worth noting that, while cyclic voltammetry prevails as

the most common electrochemical method presented through-

out the π-conjugated polymer literature, DPV is now clearly

established to allow for higher sensitivity, yielding sharper redox

26

(

see the Supporting Information). Important similarities exist

onsets as a result of a reduction of charging currents. This, in

turn, increases the accuracy of the energy gaps estimated

electrochemically; as such, DPV values will be the primary

values referred to throughout this work, while CV values are

only reported for comparison with other important contribu-

tions on the topic. The beneﬁts of using DPV over CV in this

study are illustrated in Figure 4, which overlays the voltammo-

grams obtained via the two approaches. Here, our interpreta-

tion of the CV of P2.2 reveals a diﬀerence of 150 mV in the

onset of oxidation between the CV and DPV of the same

polymer ﬁlm, with the CV onset being somewhat ambiguous.

As an added level of complexity, the various methods of

1

3

1

among the chemical shifts seen in all C{ H} CP/MAS NMR

spectra (Figure S3a−d), and only minor improvements in line

widths are observed in the region of the side chains after the

samples have been annealed. Following the conclusions drawn

earlier for P5, this provides evidence that thermal annealing

aﬀects the reorganization of appended octyl side chains, albeit

to a relatively modest extent for the two molecular weight

fractions of P1 shown here. Interestingly, the eﬀect of annealing

at 200 °C appears to be less pronounced for P1 than for P5. A

possible explanation for this diﬀerence could be that the longer

octyl side chains introduce more conformational disorder in

comparison to the shorter branched 2-ethylhexyl chains of P5.

In theory, linear octyl side chains can adopt both “linear” (all-

trans) and “random coil” (predominant gauche) conformations

in the condensed state of the copolymer. In contrast, the

shorter branched 2-ethylhexyl pendant chains would be

somewhat less ﬂexible, while allowing more interlamellar

reorganization (in consistency with the larger π-stacking

distance of 3.6 Å estimated for P5 by 2D-WAXS). In other

words, the structural disorder introduced by the linear octyl

side chains cannot be easily removed by annealing the sample.

27

determining electrochemical onsets (e.g., intersecting tan-

gents) can lead to diﬀerent interpretations of where the onset

actually occurs within the range 0.1−0.4 V, whereas DPV

provides a sharp onset at which the oxidation is observed at

0.10 V.

In addition to using DPV to determine redox onsets, we

found that particularly thin ﬁlms of polymer drop-cast from

−1

dilute solutions (0.1−0.25 mg mL ) onto platinum disk

electrodes from chlorobenzene solutions yielded more electro-

chemically reversible voltammograms, and narrower electro-

chemical energy gaps when compared to thicker ﬁlms deposited

from more concentrated casting solutions (see Figures S11−

1

Importantly, the 2D H− H DQ-SQ correlation spectrum of

P1.2 (Figure S3e) shows an even more elongated line shape

8

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Figure 5. DPV (step size 2 mV, step time 38 ms, pulse amplitude 100

mV) of thin ﬁlms of P1.1, P1.2, and P5.1 drop-cast on a platinum disk

electrode immersed in 0.1 M Bu NPF /PC solution. Arrows indicate

Figure 4. CV (blue curve, 50 mV/s scan rate) and DPV (black curve,

step size 2 mV, step time 38 ms, pulse amplitude 100 mV) of a thin

ﬁlm of P2.2 drop-cast from chlorobenzene on a platinum disk

immersed in 0.1 M Bu NPF /PC solution. Arrows indicate direction

4

6

scan direction.

4

6

of scans.

polymer need to be analyzed to reveal much lower oxidation

potentials by CV (onset around −0.1 V). This result further

underscores the greater sensitivity attained in using DPV vs CV

when determining electrochemical onsets.

In relating oxidation potentials and thin-ﬁlm structural

ordering capabilities, it is also worth noting that the oxidation

potential of P2.1 is 120 mV higher than that of its higher

molecular weight counterpart P2.2 (0.22 V versus 0.10 V,

respectively) while P2.2 exhibits more “crystallinity” in 2D-

WAXS measurements.

S15). This behavior is attributed to the neutral state polymer

network being essentially insulating in character, inhibiting ion

transport, and increasing uncompensated resistance. The use of

thick polymer layers, in essence more polymer on the electrode,

is expected to increase resistance and the total amount of

charge passing. This induces larger ohmic polarization and thus

creates larger overpotentials through the relationship shown in

eq 1, whereby the potential of the working electrode E at a

certain current ﬂow is equal to the sum of the null potential

In contrast, the reduction potentials of all the polymers were

nearly identical as measured by CV and DPV, yielding onset

(

E ), ohmic polarization (or uncompensated resistance, IR ),

n

u

+

activation polarization (η ), and concentration polarization

values on the order of −1.6 and −1.5 V versus Fc/Fc ,

act

2

8

(

η_c_o_n_c). Since activation polarization is intrinsic to the analyte

suggestive of a polymer reduction potential governed by the

reduction of the electron-deﬁcient BTD units. In addition, for

all polymers except P1.2, the peak reductive current (as

measured by CV) either remained the same or decreased upon

increasing ﬁlm thickness (see Table S2), suggesting that the

negative charge injected into the polymers was not mobile

throughout the bulk material. This assertion is supported by the

FET charge transport study (vide infra), whereby no evidence

for n-type transport was observed in any of the polymers in

combination with the dielectrics used in our laboratories.

Looking more closely at the DPV results, it seems that the

reduction onsets of P1.1−P4 approach −1.5 V, with the

exception of that observed for the more ordered system P5.1

showing a reduction at −1.39 V. Here, it is reasonable to

associate the measurement of a lower reduction potential to the

more pronounced “crystallinity” in thin ﬁlms of P5.1.

at a given interface, and since concentration polarization can be

neglected for a ﬁlm suﬃciently thin, employing relatively thin

polymer ﬁlms can be expected to limit polarization eﬀects in

general. As a result, the overpotential required for current ﬂow

can be minimal, and the same thin ﬁlms can be used for the

measurement of electrochemical onsets close in value to the

true thermodynamic potentials.

E = E + IR + η ⁺^ηconc

n

u

act

(1)

Importantly, as can be seen from Table S1, signiﬁcant

diﬀerences in redox potential values arise on going from one

conjugated backbone to another, on changing the solubilizing

substitution pattern, and depending on the molecular weight

fraction considered. In particular, P1.1, P1.2, and P5.1 all

contain the same conjugated backbone, but the onsets of

oxidation diﬀer by 300 mV as evidenced by their respective

DPVs shown in Figure 5. In this case, it is interesting to note

that decreasing the oxidation potential correlates with

increasing the thin-ﬁlm structural ordering capabilities. In

particular, while the lower molecular weight fraction of P1,

namely, P1.1, demonstrated less “crystallinity” than P1.2 by

Overall, these results support the idea that redox potentials in

π-conjugated polymers are not only governed by the chemical/

molecular structure of the repeat units, but also by the distinct

microstructural ordering propensities of each system; that is,

the ﬁnal measurable properties are emergent characteristics of

the material. This is highlighted by the comparable ionization

potentials and electron aﬃnities determined for similarly sized

oligomers across the polymer series at the B3LYP/6-31G(d,p)

level of theory. Accordingly, in addition to the repeat unit

structure parameter, molecular-weight distribution, side-chain

concentration, and side-chain structure represent a set of

important factors that inﬂuence the magnitude of the ionization

potentials in the solid-state polymer materials. Of course, the

same parameters should also be expected to impact the overall

2D-WAXS analysis, the oxidation potential of P1.2 is found to

be 140 mV lower than that of P1.1 (0.22 and 0.08 V,

respectively). In comparison, the 2-ethylhexyl-substituted

analogue of P1, namely, P5.1, showing particularly pronounced

“crystallinity” from 2D-WAXS with signatures of long-range

order, yields the lowest potential of oxidation at −0.12 V versus

+

Fc/Fc . It should be noted that this potential is strikingly

diﬀerent from independently reported values estimated by CV

1

5

6,11,29

(+0.25 V). Here we found that particularly thin ﬁlms of this

performance of thin-ﬁlm devices.

8

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and shifts from ca. 410 nm to ca. 480 nm. Noticeably, the dual-

band absorption of P4 is the most coalesced and balanced of all,

while maintaining high absorption coeﬃcients across the visible

spectrum.

The empirical trends for the ﬁrst transition energies are well

reproduced by time-dependent density functional theory

(

TDDFT) calculations at the B3LYP/6-31G(d,p) level; see

Tables S1 and S7 (even if the eﬀects of electron self-interaction

from the general DFT formalism result in a red shift of the

transition energies). The S →S transitions for P1−P4 are

0

1

principally described as HOMO→LUMO transitions in the

one-electron picture, though other close-lying frontier molec-

ular orbitals (in particular HOMO-1 and LUMO+1) contribute

to the description of the excitation.

Overall, P1−P4 absorb more eﬀectively where the solar

Figure 6. Absorption coeﬃcients as a function of wavelength for P1−

P4 (along with those for P3HT). The absorption coeﬃcients (in

photon ﬂux is the most intense (that is in the 600−800 nm

1

3

range) in comparison to their all-donor parents, a result that

could promote their eﬃciency in solar cell devices.

−1

−3

cm ) are derived from the mass absorption coeﬃcients (in L g

−1

cm ) measured in toluene solution by assuming a density of 1 g cm

Transistor Fabrication and Characterization. In order

to probe the eﬀect of the synthetic modiﬁcations made on

charge transport, bottom-gate/bottom-contact FET devices

were constructed on heavily doped Si wafers covered with 200

nm of thermally grown dielectric SiO₂passivated with

hexamethyldisilazane (HMDS). The gold contact electrodes

were patterned using conventional photolithographic methods

to yield 50 μm long and 3500 μm wide channels. The

semiconducting polymer layers were drop-cast from o-

for all polymers. (Figure adapted with permission from ref 4b.

Polymer Optical Properties. The absorption spectra of

P1−P4 were recorded in toluene solutions and for thin ﬁlms

(

see Table 1). In Figure 6, the mass absorption coeﬃcients of

−1

P1−P4 (in L g cm ) have been converted to “thin-ﬁlm”

absorption coeﬃcient values (in cm ) by assuming a density of

1

g cm ). The UV−visible spectrum of regioregular poly(3-

hexylthiophene) (P3HT) is superimposed for comparison.

As expected from previous work investigating the spectral

distribution of π-conjugated polymers with intramolecular

−1

−

3

30

g cm for all polymers (e.g., poly(3-octylthiophene) ∼ 1.05

−1

−3

4b

dichlorobenzene (2 mg mL ) onto a hot-plate preheated to

00 °C, or spin-cast from the same solvent, and were

subsequently annealed for 15 min at various temperatures

vide infra). The solution-processing of the active layer and

1

(

4b,31

device testing were carried out under N atmosphere in a

donor−acceptor interactions,

P1−P4 possess a “dual-

2

glovebox. Table 2 gives an overview of the ﬁeld-eﬀect transistor

performance obtained across the polymer series, including hole

mobility and on/oﬀ current ratios, comparing drop-cast and

spin-cast processing methods. The molecular weight and 2D-

WAXS results described in Table 1 are also reported in Table 2

in order to facilitate drawing direct correlations with the FET

mobility results.

The ﬁrst striking dependence of the electronic device

performance on the polymer molecular weight is represented

by the higher molecular weight fraction of P1, namely P1.2 (M_n

= 12.7 kDa in TCB), raising the FET hole mobility from 2 ×

band” of absorption evolving as a function of the ratio of

electron-rich to electron-deﬁcient building units incorporated

along the backbone. In comparison with their all-donor π-

13

conjugated parents, the onsets of absorption of P1−P4 are

red-shifted by ca. 150 nm toward longer wavelengths, with the

smallest optical gap obtained for P1.2 (1.4 eV), in which the

concentration of electron-deﬁcient units along the backbone is

the largest across the polymer series. While their maximum of

absorption peaks in the range 600−650 nm, the second local

maximum located at shorter wavelengths is more sensitive to

the concentration of electron-rich units present in the backbone

a

Table 2. Overview of Field-Eﬀect Transistor Performance Obtained across Polymer Series

M (PDI)

n

2D-WAXS

FET drop-cast

XRD^d^r^o^p

FET spin-cast

polymer

TCB

π

d

I :I

μ_S_A_T

anneal.

d

I :I

μ_S_A_T

anneal.

on oﬀ

1 × 10²

5 × 10³

on oﬀ

c

−6

−2

−4

−2

−3

−2

−1

c

P1.1

P1.2

P2.1

P2.2

P3

10.3 (3.6)

12.7 (6.3)

18.3 (6.1)

36.1 (6.0)

16.3 (8.2)

n/a

1.95

1.80

2 × 10

1 × 10

3 × 10

4 × 10

3 × 10

1 × 10

2 × 10

5 × 10

1 × 10

100

150

100

225

100

175

100

250

150

0.35

0.36

n/a

1.85

3 × 10²

1 × 10⁻⁴

200

b

3

(1.95/1.87 )

1.75

6 × 10

2 × 10⁴

1.79

1.99

1 × 10⁴

2 × 10³

7 × 10

9 × 10

−3

200

−4

1.94

3

× 10⁵

3

2 × 10³

1 × 10⁴

7 × 10⁻⁴

5 × 10⁻³

P4

16.0 (8.0)

18.3 (3.3)

0.36

1.84

1.70

1 × 10

× 10⁴

1 × 10

1.83

1.74

200

6

5

P5.1

a

Number average molecular weight in TCB (M , kDa). π-Stacking (π, nm) and chain-to-chain lamellar distances (d, nm) by 2D-WAXS. Charge-

n

2

−1 −1

carrier mobilities at saturation (μ , cm V s ) and current on/oﬀ ratios (I :I ) in drop-cast FETs for copolymers P1.1, P1.2, P2.1, P2.2, P3, P4,

sat

on oﬀ

and P5.1 on HMDS-passivated SiO . Annealing temperature (°C) for the drop-cast FETs. XRD chain-to-chain distances (d, nm) as measured from

2

the drop-cast FET devices. Charge-carrier mobilities at saturation (μ ) and current on/oﬀ ratios (I :I ) in spin-cast FETs for copolymers P1.2,

sat

on oﬀ

b

c

P2.2, P3, P4, and P5.1 on HMDS-passivated SiO . Fiber was annealed at 200 °C for 40 min. FET made on PTES-passivated SiO .

2

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1

0⁻to 1 × 10⁻²cm V s⁻¹in a drop-cast transistor processed

6

2

−1

commercial viability aspects, and is in excellent agreement with

following the same conditions as those used with P1.1 (M =

the ability of this polymer to “crystallize” from the processing

n

14c

1

0.3 kDa in TCB). The same device showed only a slight

solvent (see Figure S1f and previously reported work ),

without being subjected to more elevated temperatures as

necessary with its counterpart P2.2.

−

2

−1 −1

increase to 1 × 10 cm V

s

on annealing at a temperature

identiﬁed as providing the best device performance for the

drop-cast polymer (150 °C here), a result consistent with the

modest increase in microstructural order observed in the ﬁber

annealed at elevated temperatures (see Figure S1a and b). This

increase by ca. 4 orders of magnitude obtained with P1.2 is all

the more surprising that an average molecular weight diﬀerence

of less than 3 kDa exists on going from P1.1 to P1.2, but

correlates well with the diﬀerences observed in the respective

X-ray patterns of the two fractions (see Figure 1a and b), and is

attributed to the combination of a net diﬀerence in polymer

processing, inﬂuencing the thin-ﬁlm morphology, and to the

absence of π-stacking interactions in ﬁlms processed with P1.1.

These important results further reinforce the necessity to attain

number-average molecular weights that are suﬃciently high to

be able to remove the largest extent of this dependence from

the polymer performance data.

Interestingly, in comparing the linear octyl-substituted P1.2,

whose highest achievable molecular weight was found to not

exceed 12.7 kDa, to its branched 2-ethylhexyl-substituted

counterpart P5.1, a clear FET performance enhancement was

2

−1

observed, thus reaching mobility values as high as 0.1 cm V

−1

s

in drop-cast devices (see Table 2). Figure S17 shows the

FET transfer curve and the corresponding output character-

istics at various applied gate-voltages for P5.1. Similarly here,

the device properties of P5.1 are consistent with the high

degree of microstructural order observed by 2D-WAXS (see

Figure 2). Following those considerations, it is further worth

noting that the molecular weight of P5.1 remains relatively

modest when compared to that attained with P2.2, and the

synthesis of a longer backbone distribution is expected to lead

to even higher mobilities as no clear evidence that a limit in

terms of FET performance as been reached so far in our work

with P5.

Upon comparing P2.1 (M = 18.3 kDa in TCB) to P2.2 (M

n

=

36.1 kDa in TCB), the transistor mobility gains about 2

−4

−2

2

−1

orders of magnitude, raising from 3 × 10 to 4 × 10 cm V

For comparison of diﬀerent processing methods, P1−P5

were integrated in spin-cast (as opposed to the above-discussed

−

1

s , in a drop-cast transistor subsequently annealed at a

temperature identiﬁed as providing the best device performance

for the respective drop-cast polymers (Figure S16). This result

is consistent with the pronounced increase in lamellar

orientation and the decrease in d-spacing observed in the

bulk of the polymer by 2D-WAXS for the larger molecular

weight fraction (see Figure 1c and d), as well as with the slight

increase in π-stacking intensity seen on thermal annealing in

general (see Figure S1c and d).

drop-cast) thin-ﬁlm transistors (HMDS-passivated SiO ); the

2

devices were all subjected to a thermal annealing step at 200 °C

(

15 min), since all semiconducting copolymers had previously

shown close to peak performance after exposure at 200 °C in

this device conﬁguration. Field-eﬀect mobilities and current on/

oﬀ ratios are reported in Table 2. In the present device

conﬁguration, the charge-carrier mobilities for the polymers are

less than 2 orders of magnitude lower than those determined

with the drop-cast architecture, which also points toward a

certain degree of structural order for the spin-coated polymers.

Here, P2.2 and P5.1 show the highest hole mobilities −7 ×

Similarly, the device thermal annealing conditions were

found to aﬀect the performance of P3 and P4 in FETs (though

to a lesser extent), with hole mobilities increasing to values as

high as 5 × 10⁻cm V

2

−1 −1

s

in the case of P4 after annealing,

−3

2

−1 −1

1

0

and 5 × 10 cm V s , respectively, and the mobility

−2

2

−1 −1

and a peak mobility of 1 × 10 cm V

s

for P3 after

values obtained for P3 and P4 are relatively close to these, in

agreement with the processing method employed.

annealing. Figure 7 illustrates the trend in ﬁeld-eﬀect hole

mobilities observed following the successive thermal annealing

steps subjected to the drop-cast transistor devices made with

P2.2, P3, and P4. In the case of P4, the high as-cast mobility

Thin-Film X-ray Diﬀraction Analysis. X-ray diﬀraction

XRD) was performed on polymer thin ﬁlms (for P1.2, P2.2,

(

P3, P4, and P5.1) obtained following the same solution-

processing and temperature-annealing conditions as those

employed during the transistor fabrication. As reported in

Table 2, the d-spacings estimated by XRD closely followed the

trend set by the 2D wide-angle X-ray scattering method, with

values only slightly higher than those estimated from the

polymer ﬁbers; this reinforces the relevance of the measure-

ments performed from the bulk semiconducting materials (vide

supra). For example, P2.2, P4, and P5.1 also show the smallest

interlamellar distances in drop-cast ﬁlms, estimated to be as low

as 1.79, 1.83, and 1.74 nm, respectively. The diﬀractograms for

the drop-cast polymer ﬁlms are shown in Figure S18.

Interestingly, the Bragg reﬂection seen at the small angles for

each polymer is relatively pronounced and nearly identical in

intensity across the analogous series. In this sense, looking at

mechanically oriented ﬁbers by 2D-WAXS provided much

deeper insight into the varying degrees of microstructural order

attained by each polymer upon thermal annealing. Other XRD

experiments (low-temperature annealing) revealed that only P4

shows the ability to ‘crystallize’ from the processing solvent

without being subjected to more elevated temperatures as

necessary with its counterparts P2.2 and P5.1.

value of 2 × 10⁻cm V

2

−1 −1

s

is worth emphasizing considering

Figure 7. Evolution of ﬁeld-eﬀect hole mobilities at saturation (μ ,

sat

2

−1 −1

cm V s ) for P2.2, P3, and P4 drop-cast on HMDS-passivated

SiO as a function of annealing temperature (annealing maintained 15

2

min each time).

8

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Journal of the American Chemical Society

Article

Figure 8. AFM tapping-mode images (heights) of (a) P1.2 and (b) P2.2, drop-cast on HMDS-treated SiO (the ﬁlm surfaces are represented), along

2

with that of (c) P3 spin-cast on HMDS-treated SiO for comparison. Each device received a postpolymer processing thermal treatment of 15 min at

2

a temperature identiﬁed as providing the best device performance (vide supra).

Table 3. Solar Cell Device Performance Observed for P1.2, P2.2, P3, P4, P5.1, and P5.2 at Optimized Polymer:PC BM blend

7

1

−2

a

Composition (w/w) (AM1.5G illumination, 100 mW cm )

solar cell devices

hole-only devices

polymer

PX:PC BM

thick.

J_S_C

V_O_C

FF

PCE

μ(PX)

μ(PX:PC BM)

7

1

71

1.0 × 10⁻

3.0 × 10

2.0 × 10⁻

1.0 × 10⁻

8.0 × 10⁻

2.4 × 10⁻

5

4

5

3

5

4

1.0 × 10

−5

P1.2

P2.2

P3

1:2

110

120

105

110

8.52

9.80

0.63

0.51

0.70

0.59

0.61

0.62

37

49

34

43

51

55

2.01

2.44

1.58

2.02

3.80

4.59

−

−4

1.9 × 10

−5

6.65

1.0 × 10

−4

P4

7.89

4.8 × 10

−4

P5.1

P5.2

1:1.5

12.29

13.44

2.2 × 10

−4

6.8 × 10

a

−2

Thickness of the active layer (nm), short-circuit current density (mA cm ), open-circuit voltage (V), device ﬁll-factor (%), power conversion

eﬃciency (%). The devices were subjected to a post-polymer processing thermal treatment at 200 °C (for 1−5 min), under inert atmosphere. The

solar cell device structure is ITO/PEDOT/PX:PC BM/LiF/Al. All the solar cell results were collected under atmospheric conditions. Zero-ﬁeld

70

2

−1 −1

hole mobility (cm V s ) in the pristine copolymers P1.2, P2.2, P3, P4, P5.1, and P5.2, and in the polymer phase of the optimized blends

PX:PC BM). Values derived from hole-only devices (ITO/PEDOT-PSS/PX or PX:PCBM/Au), and obtained by ﬁtting the J−V data to trap-free

(

71

single-carrier SCLC model. Devices with post-polymer processing thermal treatment at 150 °C (5-6 min).

AFM Topography. In addition to the X-ray analyses

performed for the bulk and thin-ﬁlm semiconducting polymers,

atomic force microscopy (AFM) was employed in the tapping

mode to image the topography of the transistor device thin-ﬁlm

active layers obtained from the two distinct deposition methods

applied in this study (drop-cast versus spin-cast deposition).

The analysis was carried out to probe the ﬁlm-forming

propensities across the polymer series, and to tentatively

correlate microstructural order-directed morphology with

device performance. While the morphology may arguably diﬀer

at the polymer−dielectric interface, it is expected that a certain

degree of macroscopic and microstructural order is retained, for

the drop-cast ﬁlms in particular.

2D-WAXS (see Figure 2). In parallel, P5.1 shows an increase in

hole mobility (1 × 10⁻cm V s ) of about half an order of

magnitude in comparison with that of P2.2 and P4. In

comparison, the morphologies obtained from spin-casting the

1

2

−1 −1

semiconducting polymers onto HMDS-passivated SiO were

2

found to be signiﬁcantly more homogeneous as exempliﬁed

with P3 in Figure 8c and further detailed in Figure S20. Beyond

the apparent ﬁlm similarities, it is worth noting that the image

relative to P5.1 does show coarser features reminiscent of those

observed in the drop-cast ﬁlms (Figure S19e), a feature induced

by the more pronounced propensity of the polymer to self-

assemble into structured domains (as suggested by the 2D-

WAXS analysis).

Figure 8 illustrates the various morphologies obtained for the

drop-cast (Figure 8a,b and Figure S19) and spin-cast (Figure 8c

and Figure S20) depositions of the semiconducting polymers

Solar Cell Device Fabrication and Testing. The

photovoltaic properties of P1−P5 were examined in donor−

acceptor bulk heterojunction solar cells (BHJs) employing

(6,6)-phenyl-C -butyric acid methyl ester (PC BM) as the

on HMDS-passivated SiO . While the topographic patterns for

2

71

P1.2 and P3 appear to possess a similar nodular character when

electron acceptor. All solar cell results were collected under

atmospheric conditions and simulated AM 1.5G solar

illumination (at an irradiation intensity of 100 mW cm ).

−2

2

compared to others, the p-type mobility of P1.2 (1 × 10 cm

−1

−2

2

−1 −1

−2

V

s ) and that of P3 (1 × 10 cm V s ) at peak

performance are also about half an order of magnitude less than

The copolymer-PC BM blends were spin-cast from chlor-

71

−2

2

−1 −1

those of their counterparts P2.2 (4 × 10 cm V s ) and P4

obenzene in devices using PEDOT-PSS-coated ITO glass

substrates. The blend compositions, as well as the thickness of

the active layer to be solution-processed, were optimized. The

results are summarized in Table 3. The semiconducting active

layers were subjected to appropriate thermal annealing before

top electrode deposition (see Table 3). Figure 9 superimposes

the J−V curves for devices made with P1.2, P2.2, P3, P4, and

P5.2 at optimized thickness and polymer:PC BM composi-

2

−1 −1

(

5 × 10⁻cm V s ) in drop-cast devices. The improved

charge-carrier mobilities of P2.2 and P4 may stem from their

much more ﬁbrilar morphology in thin ﬁlms, suggesting the

presence of domains with more pronounced alignment in

comparison with the nodular case. Following the same

argument, the presence of more extended ordered domains is

suggested in the ﬁlms of P5.1 (Figure S19e). The more

71

“crystalline” character of P5.1 has previously been evidenced by

tion.

8

953

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Journal of the American Chemical Society

Article

EQE is assigned to the long wavelength absorption band of the

polymer. The short wavelength maximum can be attributed to

the absorption by PC BM combined with that of the short

71

wavelength absorption band of the polymer (see Table 3), and

the onset of photocurrent at about 900 nm is in agreement with

the polymer thin-ﬁlm optical absorption (also shown in Figure

S21). Considering that no small molecule processing additive

was used during device preparation (in order to avoid

interfering with the self-assembling properties of the materials

described in this study), the excellent results obtained with

P5.2 versus P1.2 could be the consequence of (i) the higher

molecular weight, (ii) its greater ability to order (as suggested

by 2D-WAXS), or more likely (iii) a combination of these two

factors. As suggested earlier in this report, it is expected that the

presence of the 2-ethyl substituent appended to the hexyl main

chain of the solubilizing side groups promotes favorable self-

assembly of the polymer backbone by imparting proper

interchain spacing, less hindered torsional motions between

repeat units, and adequate processing ability. In any case,

diﬀerences in carrier mobility should be expected to result from

these intrinsic structural disparities (vide infra).

Figure 9. I−V curves of P1.2, P2.2, P3, P4, and P5.2 based PSCs (at

best polymer:PCBM composition) under AM 1.5 solar illumination,

−

2

1

00 mW cm . Devices with postpolymer processing thermal

treatment at 70 °C (30 min). The device structure is ITO/

PEDOT/PX:PC BM/LiF/Al with PX = P1.2, P2.2, P3, P4, or P5.2.

70

Inspecting Figure 9 further, it is essential to note that the

open-circuit voltages (V_O_C) observed for the devices made are

system-dependent and vary considerably across the polymer

series: being as low as 0.51 eV for P1.2 and as high as 0.70 eV

for P3. Given the close oxidation potentials estimated

electrochemically for the polymers and the comparable IPs

As shown in Table 3, a polymer:PC BM ratio of 1:2 (by

7

1

weight) was found to be optimum across the entire linearly

substituted polymer series P1−P4. Here, the results obtained

throughout the structural X-ray analyses of the polymers, and

the charge-transport study in transistors, justiﬁed using the

largest molecular weight fractions of P1 and P2, respectively

P1.2 and P2.2, for the solar cell device fabrication. The

polymer that provided the most limited PV device performance

was P3, which showed a relatively modest power conversion

eﬃciency (PCE) of 1.58%, correlated to the low intensity of

^c^a^l^c^u^l^a^t^e^d^b^y^D^F^T^,ⁱ^t^a^p^p^e^a^r^s^t^h^a^t^t^h^e^v^a^rⁱ^a^tⁱ^oⁿ^sⁱⁿ^VOC

observed across the polymer series could depend less on the

intrinsic energetics of the individual polymer chains, than on

the degree of microstructural order in the polymer-rich

32

domains and the ﬁlm morphology for example.

Insight into the morphology obtained on blending P5.2 with

the electron-acceptor PC71BM at optimized polymer:fullerene

composition is provided in Figure S22. AFM was used in the

tapping-mode to image device regions where the top contacts

were not present. The 2D image of the P5 based device surface

shown in Figure S22 reveals a relatively nodular morphology

(sharp features of 10 nm in height, rms roughness = 1.567 nm)

where large domains of higher heights are dispersed within the

−2

photogenerated current (ca. 7 mA cm ). Interestingly, P3

shows the least pronounced degree of achievable micro-

structural organization (as evidenced by 2D-WAXS), with no

observable π-stacking, relatively long interlamellar distances (d

=

1.94 nm), and a clear lack of lamellar orientation (see Figure

S1e). In comparison, P1.2 and P4 show almost exactly the

same device performance, with PCE values in the order of 2%

and ﬁll factors (FF) exceeding 35%. In the linearly substituted

polymer series P1-P4, P2.2 represents the most eﬀective

system in BHJs with PC BM, demonstrating the highest

33

active layer. Consistent with other contributions on the topic,

the larger domains with taller heights are tentatively assigned to

the PCBM-rich phase, with no indication of “overgrown”

PC71BM-cluster formation. In fact, it is now well established

that a certain extent of phase separation between donor and

acceptor in the BHJ remains a necessary condition in the

formation of the bicontinuous interpenetrating network

7

1

−

2

photogenerated current densities (ca. 10 mA cm ), a

substantial FF of 49%, and up to 2.44% of PCE. Here again,

when looking at the results obtained by 2D-WAXS (see Figure

1d), P2.2 shows the more pronounced degree of achievable

3

4

structural organization, with a narrow d-spacings of 1.75 nm

and higher-order small-angle reﬂections suggesting better

packing properties.

However, it is only by varying the side-chain substitution

pattern of P1 to yield the polymer analogue P5, here

conﬁguration desired to obtain eﬃcient solar cell devices.

In contrast, and as previously proposed,

34,35

a pronounced

phase segregation (demixing) between PC71BM and the

polymer (not observed here) would likely impede the BHJ

device performance by conﬁning the dissociated charges to the

segregated domains.

represented by its molecular weight fraction P5.2 (M = 26.4

n

kDa in TCB), that an outstanding performance enhancement is

observed. In this case, the measured photogenerated current

Space Charge-Limited Current−Voltage Measure-

ments. Considering the distinct solar cell responses observed

across the polymer series P1−P5 in BHJs with PC BM,

inspection of the space charge limited current (SCLC) in the

composite active layers (at optimized composition) was

envisaged to correlate carrier transport and device performance.

With the photogenerated current in donor−acceptor excitonic

solar cells dominated by a combination of exciton generation

and dissociation rates, and by the recombination processes, the

charge transport in the p-type and n-type semiconductors

−2

density exceeds 13 mA cm , and the FF improves by some

0% when compared to that of the linearly substituted analogue

7

1

5

P1.2 (FF going from 37% to 55%). Overall, the measured PCE

was found to average 4.59%, in good agreement with other

recent contributions emphasizing the photovoltaic properties of

3

d,14a,b,15

this particular system.

The EQE data of the

corresponding P5.2-based BHJ device is shown as Figure

S21, in which the long wavelength maximum peaking at 51% of

8

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Journal of the American Chemical Society

Article

should be balanced to avoid the buildup of charges commonly

observed in materials with low intrinsic drift carrier mobilities

polymer P1.2, to P3, to P4, and to P2.2, likely as a result of (i)

increasingly prominent ordering abilities and (ii) increasingly

higher polymer molecular weights.

36

and short carrier lifetimes. In this regard, the SCL transport in

the pristine polymers was examined to gain insight into the

inﬂuence of microstructural order on the charge-carrier

properties of P1−P5 in vertically stacked devices (i.e., factoring

out the ﬁeld eﬀect dependence as in thin-ﬁlm transistors). The

electron mobility in the PCBM-rich phase of a polymer:PCBM

blend has previously been determined to be on the order of

It is worth noting that neither P5.1 nor P5.2 reveal

drastically superior hole mobilities by the SCLC approach,

such that the clear photovoltaic performance enhancement

observed in Figure 9 could result from a simple morphology

eﬀect (in absence of further insight into the polymer charge-

transport characteristics in blends with PC BM). In fact, the

71

−3

2

−1 −1 37

1

0

cm V s .

In systems where the dark current is space-charge limited

i.e., in media with low mobility and low intrinsic carrier

zero-ﬁeld hole mobilities in blends of P5 with PC BM (as

71

optimized to best photoresponse) are found to be more than

one (P5.1) and nearly two (P5.2) orders of magnitude larger

than that for P1.2 (see Table 3), hence strongly supporting the

considerably improved photovoltaic response seen in the

branched alkyl-substituted analogue P5.2 based devices.

Importantly, the hole transport in the polymer−fullerene

blends is almost systematically lower than that in the pristine

polymers, except in the case of P5.2, thus pointing to a

favorable intimate polymer−fullerene self-assembly achieved on

blending the donor and the acceptor in the latter case.

Favorable changes in polymer−fullerene self-assembly have

previously been proposed on mixing PPV-based donor

polymers and PCBM in order to account for hole mobility

(

concentration), the mobility can be simply extracted from the

current−density−voltage response, as illustrated by the Mott-

Gurney equation for trap-free SCL currents:

2

_J₌₉_/₈_μ_εV

_d3

(2)

where ε is the dielectric constant, μ is the charge-carrier

mobility, and d is the sample thickness. Hole-only devices with

gold electron-blocking counter electrodes were fabricated with

the following conﬁguration: ITO/PEDOT-PSS/polymer or

polymer:PC BM/Au. The ﬁlms were annealed at 150 °C for

36b,39

enhancements in BHJs.

This phenomenon has been

71

attributed to an increase in polymer-backbone ordering in the

presence of PCBM.

5

−6 min prior to Au deposition.

Figure S23 shows the empirical dark current densities passing

across hole-only devices made with the pristine copolymers

P1.2, P2.2, P3, P4, and P5.1. The applied voltage V was

corrected for the built-in voltage V_b_iresulting from the

diﬀerence in work function between electrodes, and appears

as the eﬀective electric ﬁeld E_e_ﬀto facilitate comparison of

devices with diﬀerent thicknesses. Here, it should be noted that,

in all cases, the current densities were found to scale

quadratically with the applied voltage (see Figure S23, inset),

hence reﬂecting space-charge-limited transport.

CONCLUSIONS

■

Following “donor−acceptor” principles, we have designed and

synthesized a series of low-band-gap π-conjugated polymer

analogues that combine the electron-rich dithienosilole (DTS)

moiety, unsubstituted thiophene spacers, and the electron-

deﬁcient core 2,1,3-benzothiadiazole (BTD). These polymers

(P1−P5 ) were designed to incorporate (i) distinct ratios of

electron-rich to electron-deﬁcient moieties in the monomer

repeat unit, and (ii) variations (i.e., concentration and

bulkiness) of the solubilizing side-chains along the backbones,

with the goal of identifying synthetic principles that impart clear

improvements in carrier mobilities in low-band-gap polymers.

In this contribution, important correlations were drawn

between the structural variations imparted to the polymer

backbones and their charge-transport properties in thin ﬁlms,

using OTFT and OPV devices as platforms for the study. Based

on a combination of X-ray techniques (2D-WAXS and XRD)

supported by SS-NMR and AFM, we emphasize the key role of

(i) molecular weight distribution and (ii) the nature of the

solubilizing substituents appended to the backbones on the

degree of microstructural organization that can be attained with

each polymer. Overall, small structural changes induced along

the DTS-BTD backbones resulted in notable variations in

charge transport and photovoltaic properties in thin-ﬁlm

devices. Of all the DTS-BTD backbones inspected, the highest

molecular weight fraction (P5.2) of the strictly alternating

donor−acceptor derivative P5 functionalized with 2-ethylhexyl

branched substituents stands out by demonstrating ﬁeld-eﬀect

From the results presented in Table 3, it should be noted

that the pristine copolymers P2.2 and P4 demonstrate hole

mobilities higher than those of P1.2 and P3 by more than 1

order of magnitude, following the SCL transport model

governed by eq 1. Further, it is interesting to note that these

results can be correlated to the trend set by the 2D-WAXS

polymer ﬁber analysis (vide supra), whereby P2.2 and P4

showed particularly small π-stacking distances (3.6 Å) and the

shortest lamellar distances (1.75 and 1.84 nm, respectively),

while essentially no π-stacking and a large d-spacing were

observed in the case of P3 for instance. With the relatively low

molecular weight of P1.2 (in comparison with P2.2 and P4)

and its higher concentration of chain-ends possibly hindering

the formation of well-ordered nanoscale domains, more modest

charge-carrier mobilities had been anticipated for this polymer.

In parallel, the especially high zero-ﬁeld mobility value

−

3

2

−1 −1

estimated with P4 (1.0 × 10 cm V s ) is nearly identical

to that observed in spin-cast bottom-gate/bottom-contact FET

devices (7 × 10⁻cm V s ), hence possibly suggesting a

4

2

−1 −1

certain isotropy in the hole transport properties of the polymer

2

−1 −1

(

although direct correlations are not systematic in considering

mobilities as high as 0.1 cm V

contact transistors. The same material approaches a space-

s

in bottom-gate/bottom-

38

SCL versus FET charge-carrier measurements ). This result

correlates with the introduction of unsubstituted conjugated

spacers along the backbone of P4 that further diminish the

concentration of solubilizing side-chains along the π-conjugated

main-chain. Overall, and as illustrated in Figure S23, the

device current densities at low device biases (actual device bias

range in solar cells) increase on going from the pristine

charge-limited (SCL) hole mobility of 10⁻cm V

3

2

−1 −1

s

in hole-

only devices composed of a BHJ of the polymer with the

electron-acceptor PC BM; the control solar cells show PCEs

7

1

4c

on the order of 4.6% in the absence of small molecule additive-

assisted device processing (i.e., no morphology control),

conﬁrming the importance of charge transport in DTS-BTD−

8

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Journal of the American Chemical Society

Article

based OPV devices. These results clearly reinforce the

argument that maximizing charge transport in conjugated

polymers implies that the appropriate combination of (i) repeat

unit structure, (ii) molecular weight distribution, and (iii)

solubilizing side-chain substituents be found for a given set of

aromatic building units, as each of these parameters impacts

polymer self-assembling properties. Remarkably, we found a

large dependence of the ionization potentials of the polymers

estimated by electrochemical methods on polymer packing.

These empirical results may have important consequences on

future material study and device applications.

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llen, K. J. Mater.

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MacDonald, I.; Shkunov, M.; Sparrowe, D.; Tierney, S.; Wagner, R.;

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̈

ASSOCIATED CONTENT

■

*

S

Supporting Information

Sparrowe, D.; Tierney, S.; Zhang, W. Adv. Mater. 2009, 21, 1091−

Experimental details, synthesis of the monomers and polymers,

additional 2D-WAXS, SS-NMR, CV and DPV data, OTFT and

OPV device fabrication details, XRD, AFM, EQE, SCLC data,

DFT calculations. This material is available free of charge via

the Internet at http://pubs.acs.org.

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D. J.; Heeney, M.; McCulloch, I. Adv. Funct. Mater. 2010, 20, 4098−

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AUTHOR INFORMATION

■

Heeney, M.; McCulloch, I.; McGehee, M. D. Nano Lett. 2009, 9,

4153−4157.

(10) (a) Havinga, E. E.; Hoeve, W.; Wynberg, H. Polym. Bull. 1992,

Corresponding Author

reynolds@chemistry.gatech.edu

2

1

(

9, 119−126. (b) Havinga, E. E.; Hoeve, W.; Wynberg, H. Synth. Met.

Notes

993, 55, 299−306.

The authors declare no competing ﬁnancial interest.

11) 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.

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ACKNOWLEDGMENTS

2

(

2

■

We acknowledge the funding of this work as follows. J.R.R.: the

Air Oﬃce of Scientiﬁc Research (FA9550-09-1-0320) for new

materials development. J.R.R. and F.S.: the Oﬃce of Naval

Research (N00014-11-1-0245) for transport measurements and

device construction/testing. J.-L.B.: the Oﬃce of Naval

Research (N00014-11-1-0211) for computational studies.

K.M.: the German Science Foundation (Korean-German IR

TG), the European Community‘s Seventh Framework

Programme ONE-P (grant agreement no. 212311), DFG

Priority Program SPP 1355, DFG MU 334/32-1, DFG Priority

Program SPP 1459, and ESF Project GOSPEL (ref no.: 09-

EuroGRAPHENE-FP-001) for characterization studies. M.R.H.

acknowledges Dr. Robert Graf for helpful discussions and Prof.

Hans Wolfgang Spiess for his continued support. P.M.B.

(

Mater. 2010, 22, 5409−5413. (d) Bijleveld, J. C.; Gevaerts, V. S.; Di

Nuzzo, D.; Turbiez, M.; Mathijssen, S. G. J.; de Leeuw, D. M.; Wienk,

M. M.; Janssen, R. A. J. Adv. Mater. 2010, 22, E242−E246. (e) Li, Y.;

Singh, S. P.; Sonar, P. Adv. Mater. 2010, 22, 4862−4866. (f) Bronstein,

H.; Chen, Z.; Ashraf, R. S.; Zhang, W.; Du, J.; Durrant, J. R.; Shakya

Tuladhar, P.; Song, K.; Watkins, S. E.; Geerts, Y.; Wienk, M. M.;

Janssen, R. A. J.; Anthopoulos, T.; Sirringhaus, H.; Heeney, M.;

McCulloch, I. J. Am. Chem. Soc. 2011, 133, 3272−3275.

(

13) (a) Usta, H.; Lu, G.; Facchetti, A.; Marks, T. J. J. Am. Chem. Soc.

006, 128, 9034−9035. (b) Lu, G.; Usta, H.; Risko, C.; Wang, L.;

Facchetti, A.; Ratner, M. A.; Marks, T. J. J. Am. Chem. Soc. 2008, 130,

670−7685.

14) (a) Hsiang-Yu, C.; Jianhui, H.; Amy, E. H.; Hoichang, Y.; Houk,

2

7

(

̈

acknowledges Dr. Uwe Rietzler and Dr. Rudiger Berger for

their support at the AFM facilities of MPIP-Mainz.

K. N.; Yang, Y. Adv. Mater. 2010, 22, 371−375. (b) Scharber, M. C.;

Koppe, M.; Gao, J.; Cordella, F.; Loi, M. A.; Denk, P.; Morana, M.;

Egelhaaf, H.-J.; Forberich, K.; Dennler, G.; Gaudiana, R.; Waller, D.;

Zhu, Z.; Shi, X.; Brabec, C. J. Adv. Mater. 2010, 22, 367−370.

(c) Beaujuge, P. M.; Pisula, W.; Tsao, H. N.; Ellinger, S.; Mullen, K.;

Reynolds, J. R. J. Am. Chem. Soc. 2009, 131, 7514−7515. (d) Piliego,

C.; Holcombe, T. W.; Douglas, J. D.; Woo, C. H.; Beaujuge, P. M.;

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Synthetic principles directing charge transport in low-band-gap dithienosilole-benzothiadiazole copolymers

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