Structure and morphology control in thin films of regioregular poly(3-hexylthiophene)


  • Martin Brinkmann

    Corresponding author
    1. Institut Charles Sadron, CNRS-Université de Strasbourg, 23 rue du Loess, 67034 Strasbourg Cedex, France
    • Institut Charles Sadron, CNRS-Université de Strasbourg, 23 rue du Loess, 67034 Strasbourg Cedex, France
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This review focuses on the structural control in thin films of regioregular poly(3-hexylthiophene) (P3HT), a workhorse among conjugated semiconducting polymers. It highlights the correlation existing between processing conditions and the resulting structures formed in thin films and in solution. Particular emphasis is put on the control of nucleation, crystallinity and orientation. P3HT can generate a large palette of morphologies in thin films including crystalline nanofibrils, spherulites, interconnected semicrystalline morphologies and nanostructured fibers, depending on the elaboration method and on the macromolecular parameters of the polymer. Effective means developed in the recent literature to control orientation of crystalline domains in thin films, especially by using epitaxial crystallization and controlled nucleation conditions are emphasized. © 2011 Wiley Periodicals, Inc. J Polym Sci Part B: Polym Phys 49: 1218–1233, 2011


Plastic electronics has emerged as a major and challenging new research field in the past few decades. This is largely due to the availability of new semiconducting polymers that combine low-cost solution processing, via, for example, spin coating or inkjet printing with interesting optoelectronic properties such as high-charge transport mobilities, electroluminescence or photovoltaic activity.1–7 Charge transport in these polymers is a complex multiscale process. On a local scale, it implies intrachain and interchain transport within crystalline domains.8 Intrachain transport is mainly affected by the conformation of the chain, that is, rupture of conjugation (e.g., by chemical defects altering the regioregularity of the chain), whereas interchain transport depends crucially on the possibility of π-stacking between adjacent chain segments. At a larger scale, the transport is also influenced by the connectivity between crystalline domains. The semicrystalline character of polymers, namely the fact that amorphous and crystalline domains coexist in thin films is another key aspect to be considered in the understanding/modeling of charge transport. In addition, the chains of certain conjugated polymer are sufficiently “flexible“ to allow folding as for “classical“ polyolefins. This folding should also impact charge transport mechanisms. In short, the charge transport is extremely sensitive to the order of the semicrystalline polymers such as regioregular poly(3-alkylthiophene)s (P3AT)s at multiple length scales ranging form molecular to nanoscales and mesoscales. This is the reason why establishing a clear picture of the structure and nanomorphology of conjugated polymers in thin films is a central issue in plastic electronics. This brings forth the stringent need for a better understanding of the crystallization mechanisms of conjugated polymers in general, and more specifically in thin films.

Whereas the structure and crystallization of classical polymers such as polyolefins is now rather well established, the field of conjugated polymer crystallization is largely virgin and should benefit from the past knowledge gathered on semicrystalline polymers.9, 10 In the field of soluble π-conjugated polymers, regioregular (P3AT)s, and especially poly(3-hexylthiophene) (P3HT) have emerged as key materials for the elaboration of organic field effect transistors (OFETs) and organic solar cells (OSCs) due to the combination of facile processability, large charge carrier mobilities (10−2–10−3 cm2 V.s−1) and environmental stability.5 This review focuses on the recent reported literature in the field of structural control in thin films of P3ATs. It provides some guiding principles toward the crystallization control and morphogenesis of conjugated polymers such as regioregular (P3HT). This article will deal with (i) the structure and polymorphism of P3AT, (ii) chain folding, (iii) growth from solution, (iv) nucleation in thin films and (v) epitaxial orientation and nanostructuration.



The P3AT polymer chain is made of a π-conjugated polythiophene backbone and pendent alkyl side groups ensuring the processability of this class of conjugated polymers. As recognized in the early 1990s, the chemical incompatibility between the conjugated polythiophene backbone and alkyl side chains results in the build up of a layered “self-assembled” structure in which π-stacked polythiophene backbones alternate with layers of more or less ordered alkyl side chains11–13 [see Fig. 1(a)]. X-ray diffraction patterns of P3AT films thus exhibit evenly spaced reflections representing the first, second, and higher-order reflections of this characteristic lamellar structure. The corresponding layer period increases linearly with the number of carbon atoms in the side chain n [see Fig. 1(b)].14 For P3HT, the lamellar period is in the range 1.55–1.70 nm, depending on the weight-average molecular weight [Mw; see Fig. 1(c)].15 The second characteristic length scale is the π-stacking of the polythiophene backbones with a stacking period of 0.38–0.39 nm. Finally, the semicrystalline structure of P3HT is also characterized by a secondary “lamellar” structure related to the alternation of crystalline domains and amorphous interlamellar zones as observed for instance in epitaxied films.16 This period is also a function of Mw and depends on the ability of the chains to fold (see Structural Models section).

Figure 1.

(a) Schematic model showing the hierarchical organization of the semicrystalline structure of P3HT at three main length scales: (i) the lamellar structure involves layers of π-stacked polythiophene backbones separated by layers of n-hexyl side chains (period of 1.6 nm), (ii) the π-stacking distance between two chains of 0.38 nm and (iii) a periodicity related to the alternation of crystalline and amorphous zones. The disordered P3HT chains in the amorphous zones comprising, chain ends, chain folds, and tie molecules are represented as red segments (Reprinted from ref.16, with permission from Wiley-VCH). (b) Observed lamellar period d1 0 0 corresponding to the vector a of both Forms I and II of various P3ATs as a function of the number of carbon atoms of the side chain n (Reprinted from ref.14, with permission from American Chemical Society). (c) Dependence of unit cell parameters a (black squares) and b (red full circles) as a function of P3HT number-average molecular weight Mn (Reprinted from ref.15, with permission from American Chemical Society).

Considering the family of P3AT with n = 3 to 15, two characteristic layer periods have been observed indicating that P3AT exist in two different crystal structures, called Forms I and II, which differ mainly by the side chain conformation. The polymorphism of P3ATs was recognized early on by Prosa et al.11, 13 The two polymorphs are characterized by different layer periodicities as illustrated in Figure 1(b), which shows the dependence of the layer spacing with the number of carbon atoms in the side chains. The layer period of Form I is always larger than that of Form II.

Form I is usually obtained by casting and spin coating from chlorinated solvents (chloroform, o-dichlorobenzene, etc.). It is the structure encountered in most studies dealing with OFETs and OSCs. Form II is mostly formed only in thin films as it requires specific preparation conditions and/or macromolecular characteristics. As reported by Prosa et al., Form II of poly(3-octylthiophene) (P3OT) and poly(3-dodecyl-thiophene) (P3DDT) is obtained by casting a film from a 1–3% (w/w) solution in xylene at room temperature and slow evaporation of the solvent.11, 13 Alternatively, Form II of poly(3-butylthiophene) (P3BT) can be obtained by exposure of the cast films to CS2 vapors.17, 18 Meille et al.14 suggested that the occurrence of Form II is related to the molecular mass and the regioregularity of the samples. For P3HT, Form II was reported only for low molecular mass, that is, 2–3 kDa.19 The structural conversion of Form II → Form I is possible via thermal annealing at 75 °C for P3OT and 159 °C for P3BT.13 Its irreversible character suggests that Form I is the thermodynamically stable form. It is important to note that the Form II of both P3OT and P3BT were observed in samples prepared from specific solvents, for example, m-xylene and CS2, respectively. Moreover, in the case of P3OT, Form II is obtained from a dried gel in m-xylene, which reminds the situation of a specific polymorph of poly(4-methyl-1-pentene) prepared from a gel in cyclohexane.20 The comparison with the established literature for other semicrystalline polymers suggests that Form II could be a solvent-induced polymorph.1

Table 1. Form I and Form II Unit Cell Parameters for Different P3ATs
inline image

Structural Models

Structural models for both Forms I and II (see Fig. 2 and table 1) have been proposed by several groups since the early work of Prosa et al.13 Tashiro and coworkers constructed four possible structural models of Form I of P3HT21, 22 based on X-ray diffraction data obtained on highly oriented thin films. Meille et al. determined the crystal structure for Form II of P3BT by using the electron diffraction (ED) data of Lu et al.17, 18 and combining it with a Rietveld analysis of powder X-ray diffraction patterns.23 In general, the difficulty in refining the crystal structure of a semicrystalline polymer lies in the limited size of single crystal domains which precludes the use of structural methods based on single crystal X-ray diffraction. As demonstrated in the work by Lotz and coworkers, ED in a transmission electron microscope (TEM) is a very powerful alternative method to unravel the structure of polyolefins24–27 as it can be applied to micron-sized single crystalline domains. However, whereas numerous polyolefins can be grown in the form of lamellar single crystals, single crystal growth of P3ATs is difficult and has been reported only occasionally.28

Figure 2.

Crystal structures of Form I P3HT29 and Form II P3BT23 viewed along two different crystallographic directions (hydrogen atoms have been omitted for clarity).

Brinkmann and coworkers29 recently reinvestigated Form I of P3HT by ED analysis performing rotation-tilt experiments on highly oriented and crystalline P3HT of low Mw (7.9 kDa). A rotation-tilt TEM sample holder is used to acquire characteristic projections in reciprocal space as shown in Figure 3. The ED patterns are analyzed by a trial-and-error method to determine a structural model. Figure 2 depicts the resulting structure of Form I P3HT. It is worth to note that this methodology opens the possibility to determine the reflection rules which are necessary to identify the space group of the crystal structure. In alternative structural refinement methods, the space group of a polymer crystal structure is rather inferred from symmetry considerations.

Figure 3.

Rotation-tilt experiments on an oriented thin film of P3HT (7.9 kDa) grown by slow-rate directional crystallization in TCB and showing a (0 1 0) orientation. For a θc∼40° tilt around the c axis, the electron beam (ke shows the the q-vector direction of the incident electron beam) is approximately parallel to the n-hexyl side chains as seen in the inset showing the relative orientation of the unit cell to the incident electron beam. For θc ∼ −40°, ke is almost perpendicular to the side chains. Note the quasi absence of reflections on the l = ±1 layer lines. (Reprinted from ref.29, with permission from American Chemical Society).

The structural models of Forms I and II displayed in Figure 2 show characteristic features. In both structures, the polythiophene backbone adopts a trans-planar conformation. Nevertheless, the unit cells of Forms I and II are markedly different and show specific packing schemes of the polymer backbone and n-alkyl side chains. The shorter layer period of Form II P3ATs has been attributed to side chain interdigitation.5 Interdigitation of the n-butyl side chains is indeed observed in the structure of Form II P3BT proposed by Meille and coworkers,8 whereas in the structural model of Form I P3HT, n-hexyl side chains are not interdigitated. Forms I and II differ also by the stacking of the polythiophene backbones. In form I P3HT, the polythiophene backbones are separated by 0.38 nm along the b axis when compared with 0.47 nm along the c axis in Form II P3BT. Moreover, the inclination of the conjugated backbones with respect to the stacking directions (b axis for P3HT and c axis for P3BT) is significantly different. Nevertheless, in both crystal structures, relatively short interatomic contacts are observed between successive polythiophene backbones. In Form I of P3HT, short 0.34-nm contacts are observed, in agreement with the distances obtained from density functional theory calculations.30, 31 An additional difference between Forms I and II concerns the shift of successive polymer chains along the c axis. In Form I P3HT, two successive chains along the π-stacking direction (b axis) are shifted by 0.15 nm parallel to the chain axis direction. In Form II, the thiophene cycles almost overlap in the projection along the c axis (see Fig. 2). Altogether, these structural characteristics should result in very different π-orbital overlaps, that is, different transfer integrals. Significant differences in charge transport properties between Forms I and II of P3ATs are therefore expected.

Chain Folding

Following the initial work of Mena-Osteriz et al.,32–36 several groups used scanning tunneling microscopy (STM) to study the 2D growth of P3ATs on substrates of highly oriented pyrolytic graphite (HOPG). Interestingly, these monolayer-thick epitaxied films provided a direct image of chain folding, a major characteristic of a semicrystalline polymer.9 On P3ATs bearing rather long alkyl side chains (decyl and dodecyl), STM studies provided clear evidence for the interdigitation of side chains in the 2D crystalline domains. Semiempirical calculations on P3ATs were conducted to model the folding of the chains. In P3AT, tight hairpin folds can be explained by a sequence of 7–10 alkylthiophene monomers all in cis conformation as opposed to the all-trans conformation observed for the extended P3AT chains (see Fig. 4).

Figure 4.

(a and b) STM showing the short-range ordering of regioregular P3AT. (c and d) 2D crystalline packing as modeled by semiempirical calculations. The “hairpin” intramolecular folds are composed of several 3-alkylthiophene monomers in all-cis conformation, whereas folds of larger radius may include both cis and trans conformations (Reprinted from ref.32, with permission from Wiley-VCH).

Indirect evidence for the folding of P3AT chains can be inferred from the evolution of the lamellar period with increasing molecular weight Mn. As pointed out in various studies, the width of nanofibrils as well as the lamellar period in epitaxied layers tends first to increase with Mn and reach a plateau when Mn > 10 kDa.37, 38 This is a typical fingerprint of the onset of chain folding as observed for instance in the case of oligo(urethane)s or polyethylene/parafins.39 A second hint for the possibility of P3AT chains to fold lies in the variation of the a-axis parameter (layer period) with increasing Mn [see Fig. 1(c)]. This variation of the unit cell parameters a is a manifestation of the higher proportion of re-entrant folded chains in a crystalline domains for higher Mn. The re-entrant folds are expected to exert some stress on the crystalline packing. This result is in direct line with earlier reports on chain folding in aromatic stiff-chain polymers such as poly(ether-ether-ketone) showing that the interchain distance along the b axis is increased in the polymer with respect to oligomers.40

Morphology and Structure in Thin Films

In flexible polymers, chain folding is at the root of the formation of thin lamellar crystals. These lamellar crystals further organize on a larger length scale into spherulites.9, 27 However, in spin-coated films of P3HT, no such lamellar crystals or spherulites are formed. Rather, the P3HT films are made of a network of crystalline nanofibrils embedded in a matrix of amorphous material.

However, different groups have shown that a strong correlation exists between charge transport and nanomorphology/crystallinity of the P3HT films41–49 which depend crucially on the thin-film processing conditions. The morphology and crystallinity of P3HT thin films is affected by (i) the thin-film preparation method (drop-casting, spin coating, inkjet printing), (ii) the nature of the substrates and (iii) the macromolecular parameters, for example, average molecular weight, regioregularity and polydispersity (PDI) of the polymer. In most reports, P3HT films are prepared by spin coating from solution in solvents such as chloroform. However, the nanomorphology and the crystallinity of the films can be significantly modified by using high-boiling solvents,49 solvent mixtures, or “additives,” that is, poor solvents such as 1,8-diiodooctane.50–52 For instance, charge mobility is improved by a factor of 10 (>0.1 cm2 V.s−1) when P3HT OFETs are prepared from 1,2,4-trichlorobenzene (TCB) rather than chloroform.49 Thermal annealing further enhances charge transport by modifying the nanomorphology and crystallinity of the film.43

The orientation of the crystalline domains is a key parameter in the optimization of the transport properties. For P3HT, charge transport is favorable along both the chain axis direction (c axis) and the π-stacking direction (b axis), whereas low transport is expected along the alkyl side chains (a axis). For an OFET, high in-plane charge transport between source and drain electrodes is desired, whereas in an OSC, the transport should be high along the normal to the film surface to evacuate photogenerated charges. These conditions require two distinct orientations of the crystalline domains on the substrate: “edge-on” orientation of the polymer chains with the π-stacking in-plane in OFETs and “face-on” orientation in OSCs (see Fig. 5). Edge-on and face-on orientations of P3HT chains can be favored by the choice of the proper processing parameters. Edge-on orientation of P3HT is favored for: (i) spin coating from solutions in high-boiling solvents, (ii) functionalization of the SiO2/Si substrates with self-assembled monolayers of octadecyltrichlorosilane, hexamethyldisilazane, or 3-aminopropyltriethoxysilane.46, 53 Face-on orientation is observed for spin coating from solutions in CHCl3 or deposition of P3HT films by a friction transfer method.54

Figure 5.

Color online 2D Grazing Incidence X-ray diffraction (GIXRD) patterns and AFM topographs of drop-cast (a and b) and spin-cast (c and d) P3HT films on the SiO2/Si substrates held at room temperature from CHCl3 and warm CH2Cl2. The insets in (a) and (b) illustrate schematic diagrams for edge-on and face-on orientations of RR P3HT in the films, respectively. The inset in (d) represents 1D out-of-plane X-ray profiles extracted from 2D patterns of the spincast CHCl3 and CH2Cl2 films (Reprinted from ref.48(b), with permission from American Insitute of Physics).

Several studies of the P3HT structure by grazing incidence X-ray diffraction in spin-coated films have however demonstrated that crystalline perfection and the preferred edge-on orientation are not uniform in the bulk of the films and depend on the film thickness.46, 53, 55, 56 For instance, enhanced crystallinity is observed at the air/P3HT film surface. Moreover, the level of preferred orientation is a function of film thickness. Higher preferred edge-on orientation is observed in thin films (thickness below 20–25 nm), whereas the orientational distribution of the crystallites becomes progressively random for thicker films (∼200 nm). This thickness effect suggests that the preferred orientation of nanocrystals is enforced by the substrate surface which is responsible for some specific pinning of the nanocrystal's orientation.55 Pietsch and coworkers have also shown that thermal annealing affects differently films of thickness below and above a threshold value of 20 nm.55

If scattering methods can give some insight into the average nanocrystal size and orientation in the films, the connectivity between crystalline domains cannot be directly observed by these methods. To address this point, Brinkmann and Rannou used high-resolution transmission electron microscopy (HR-TEM) in the low-dose mode to preserve the beam-sensitive P3HT films. The study was performed on oriented epitaxied films of P3HT with different molecular weights Mw.57 Increasing Mw from 7.3 to 69.6 kDa eq. PS results in several important structural changes. First, the structural coherence and extension of crystalline lamellae tends to decrease with increasing Mw. In particular, Brinkmann and Rannou observed an increase of the fluctuations of the lamellar thickness whose average value does not scale with Mw. The increase of the total lamellar periodicity is mainly related to the increase of the width of the amorphous interlamellar zones: the higher the Mw of the polymer, the lower the average crystallinity in the thin films. For samples with Mw = 69.6 kDa eq. PS, the crystalline lamellae tend to be strongly interconnected by small tie-crystallites consisting of small ordered bridging areas comprising a few π-stacked P3HT polymer chains (see Fig. 6). This enhanced connectivity may account, at least in part, for the increase in the charge carrier mobility in OFETs when the molecular weight of P3HT increases. For Mw ≥18.8 kDa eq. PS, preferential tilt angles of the polymer chains in the crystalline lamellae are observed. The increasing disorder in the crystalline packing of P3HT chains in the lamellar domains, the occurrence of certain fold planes corresponding to relatively high tilt angles of the polymer chains in the crystalline lamellae and the increase of the interchain distance (i.e., the a axis) with increasing Mw are attributed to the impact of chain folding. The folding of comb-like and stiff π-conjugated Rr-P3HT chains is expected to induce stress on the chain packing and is believed to account for the increased disorder in the structure of the crystalline lamellae.

Figure 6.

(a) Underfocused BF TEM image of a thin area of an oriented thin film of P3HT with Mw = 18.8 kDa eq. PS grown by directional epitaxial crystallization. (b) Low-dose HR-TEM image of an equivalent area showing the crystalline lamellae of P3HT with edge-on orientation. The edge of the crystalline lamellae is indicated by a line and the average chain orientation in the lamellae by an arrow. The average in-plane orientation of the chain axis is also indicated by an arrow. (c). Schematic representation of the microstructure in semicrystalline P3HT as a function of molecular weight. The crystalline domains are shown in projection along the π-stacking direction. The chains in the crystalline lamellae are drawn in green while their sections located in the amorphous interlamellar zones are red. (Reprinted from ref.57, with permission from the American Chemical Society).

The impact of the semicrystalline character of the polymer also affects strongly the optical properties of the P3HT films.58–60 As seen in Figure 7, the absorption spectrum of P3HT is composed of two contributions: (i) a lower energy component which arises from crystalline regions formed by weakly interacting H-aggregates and (ii) a higher energy component associated with more disordered chains presumably located in majority in the amorphous zones of the film. A quantitative analysis of the absorption spectrum was proposed by Spano and coworkers showing that the overall spectrum, especially the ratio between the 0-0 and 0-1 components, is very much related to the degree of excitonic coupling within the crystalline domains W. The intensity ratio between 0-0 an 0-1 components was defined as:

equation image(1)

where Ep is the phonon energy of the main oscillator mode coupled to the electronic transition and W is the free exciton bandwidth.

Figure 7.

(a) Normalized absorbance spectra taken on films spun from different solvents. A theoretical spectrum for mesitylene is also shown.58 (b) Exciton bandwidth W as deduced from eq 1 and percentage of film made up of crystalline aggregates (Reprinted from ref.60, with permission from American Institue of Physics).

Figure 7 depicts the typical exciton bandwidth obtained from eq 1 for P3HT films spin coated from solution in different solvents as well as the percentage of the film composed of crystalline aggregates. As noted earlier in this report, higher crystallinity of the films is indeed achieved in the case of high-boiling point solvents such as TCB, and this can be further demonstrated by a quantitative analysis of the absorption spectrum as demonstrated by Clark et al.60


Crystallization from Solution

Usually, crystallization of semicrystalline polymers can be achieved either from the melt or from a dilute solution. Moderate crystallinity is usually observed when P3HT films are subjected to isothermal crystallization after melting, presumably because of the limited mobility of polymer chains in the rigid amorphous phase.61

A number of groups have investigated the growth and crystallization of P3HT from solution. The motivation of these studies was to decouple the crystallization mechanism, which should be slow for improved crystal perfection, from the rather fast thin film processing as in the case of spin coating. In spin-coated films, the crystallization kinetics of P3HT is dictated by the kinetics of solvent evaporation. Thus, crystal growth occurs in far-from-equilibrium conditions when a solvent such as chloroform is used. A first means to slow down the aggregation kinetics is to use a high-boiling point solvents such as TCB which allows the formation of large nanofibrils with a preferred (1 0 0) contact plane.49 However, several groups have demonstrated the possibility to produce crystalline nanofibrils or whiskers of P3HT from solution in a controlled manner with a high yield.62–66 Starting from a hot solution of P3HT in a “marginal” solvent such as xylene, slow cooling triggers aggregation and crystallization of nanofibrils as evidenced by a characteristic 610 nm peak in the absorption spectra.64 By successive cycles of centrifugation and crystallization, it was possible to isolate the Mw fractions prone to crystallize and obtain pure suspensions of crystalline P3HT fibrils that were used for the elaboration of organic solar cells (see Fig. 8). In a similar way, Samitsu et al.63 demonstrated that high-Mw P3ATs mainly form whiskers whereas low-Mw fractions remain in the filtrate. This observation, in common with other reports, suggests that only the P3HT chains that can fold are incorporated into the crystalline nanofibrils.

Figure 8.

(a) Absorption spectra of a 1 wt % P3HT solution in p-xylene with different proportions of nanofibers and well-solubilized P3HT in the range 97–0% nanofibers. AFM images of P3HT nanofibers obtained from a 0.05 (b) and a 0.5 wt % (c) P3HT solution in cyclohexanone. The films were formed by dipping the SiO2 substrates for 2 min in the corresponding solutions (Reprinted from ref.64, with permission from Wiley VCH).

The crucial role of molecular weight distribution on the crystallization of P3HT in solution was also demonstrated by Liu et al.67 By precipitation of P3HT at different temperatures in anisole or dimethylformamide (DMF)/anisole mixtures, Liu et al. managed to grow either long nanofibrils made of folded chains or nanoribbon-like crystals made of extended P3HT chains. Nanoribbons were only obtained for P3HT with Mn ≤ 10.2 kDa (see Fig. 9). This observation is again direct evidence for the existence of a transition between extended-chain and folded-chain crystallization of P3HT with increasing Mn. A similar finding was made by Brinkmann and Rannou on oriented P3HT films grown from TCB: extended chain lamellar crystals were observed for the hexane fraction of P3HT with Mw = 7.9 kDa.38

Figure 9.

(a and b) TEM image of P3HT nanoribbons (6.0 and 10.2 kDa, respectively). (c and d) Corresponding AFM topographic images with corresponding section profiles. (e) Selected area electron diffraction images corresponding to a P3HT nanoribbon. (f) Schematic illustration of the chain stacking orientation in a nanoribbon. (Reprinted from ref.67, with permission from The American Chemical Society).

One of the problems frequently encountered when polymers crystallize from an evaporating solution is the formation of a gel. This is a crucial issue in the case of inkjet printing. In the case of P3HT, Ihn et al. have shown that gelation occurs in drop-cast films upon slow evaporation of solvents like xylene.62 Several rheological investigations have shown that gelation of P3ATs depends on the type of solvent used, the length of the alkyl side chains and the macromolecular parameters of the P3ATs.68–70 The mechanism of thermoreversible gelation has been extensively studied by Nandi and coworkers.68, 69 They suggested a two-step process involving successively (i) the coil-to-rod transformation of the P3HT chain and (ii) the crystallization of P3HT nanofibrils. A similar result was obtained by Koppe et al.,70 who proposed that aggregation of P3HT chains is followed by the formation of a network of interconnected nanofibrils which is responsible for the increased viscosity of the solution (see Fig. 10).

Figure 10.

(a) Rotational rheometer viscosity measurements of various Mw compositions dissolved in o-xylene (1 wt %) at 25 °C. A-Mw (Mw = 72,800 g mol−1), L-Mw (Mw = 26,200 g mol−1). Full lines are linear fits to the experimental viscosity data. The inset shows the slopes of the linear fits over Mw with increasing contents of L-Mw mixed into A-Mw (0, 50, 80, and 90 wt %). (b) Photos and schematic demonstration of the two-step gelation process of a 1 wt % o-xylene-P3HT solution of A-Mw batch (Mw = 72,800 g mol−1). (Reprinted from ref.70, with permission from the American Chemical Society, 2010).

The measurements of the solution viscosity demonstrated that gelation is determined by the Mw distribution of P3HT: the higher the Mw, the faster and stronger the gelation process (see Fig. 10). By blending two P3HT samples with molecular weights A-Mw = 72.8 kDa and L-Mw = 26.6 kDa, Koppe et al. showed that the viscosity η* can be gradually lowered as the content of L-Mw increases. The higher viscosity of the high-Mw samples has been attributed to the higher density of entanglement points between the polymer chains. This study of the gelation mechanism of P3ATs illustrates vividly the importance of understanding the physicochemical processes occurring in P3ATs solutions for specific processing conditions in the perspective of improving OSCs performances.

Homogeneous Nucleation of P3HT

Although the generation of spherulitic morphologies has not been reported in spin-coated films used for the fabrication of OFETs or OSCs, recent studies by Reiter and coworkers and Yang and coworkers have shown that such morphologies can indeed be generated under specific preparation conditions.18, 71 As a matter of fact, when P3BT films are subjected to solvent annealing in a vapor of CS2, spherulites were generated.18 Going one step further, Reiter and coworkers have demonstrated the possibility to effectively control the nucleation of P3HT spherulites using the so-called self-seeding methods. Self-seeding exploits a fundamental property of semicrystalline polymers, that is, the simultaneous presence of kinetically trapped metastable states, resulting in a distribution of melting temperatures.72, 73 The most stable crystals can be selected by choosing the appropriate melting temperature and are subsequently used as seeds under controlled crystallization conditions. Reiter and coworkers have developed a simple methodology to control homogeneous nucleation in thin films by using a controlled swelling/deswelling of the films.74 For P3HT, spherulitic domains of 10–100 mm diameter were obtained reproducibly by controlling the saturated vapor pressure of CS2 which determines the nucleation density of P3HT spherulites in a thin film (see Fig. 11).

Figure 11.

(a) Spherulitic morphology of P3HT films (40 nm) grown from CS2 solvent after an initial swelling at 91.0% (Pmath image) and 30 h at Pmath image. The sequence of polarized optical microscope images shows the dependence of the nucleation density of spherulites versus Pmath image. (b) Graph showing the variation of homogeneous nucleation density of spherulites versus Pmath image. (Reprinted from ref.71, with permission from Wiley-VCH).

Heterogeneous Nucleation of P3HT

Following the approach used by Chris Li to nucleate crystalline lamellae of semicrystalline polymers on carbon nanotubes,75 Liu et al. demonstrated the possibility to nucleate heterogeneously crystalline nanofibrils of P3HT on multiwall carbon nanotubes (MWCNTs).76 The Bright Field image in Figure 12(a) shows the preferential nucleation of crystalline P3HT nanofibrils on the surface of MWCNTs. The long axis of the nanofibrils corresponding to the π-stacking direction is oriented perpendicular to the Carbon nanotube's axis. The nucleating ability of MWCNTs on P3HT was further demonstrated by following the intensity of the 610 nm peak (characteristic of P3HT crystallization) for a sample with and without CNTs. As seen in Figure 12(b,c), the kinetics of P3HT aggregation is significantly enhanced in presence of MWCNTs, demonstrating the nucleating ability of MWCNTs for P3HT. The nucleating ability of CNTs is presumably related to a preferential (epitaxial) interaction between the conjugated polythiophene backbone and the surface of the CNT, similarly to the epitaxial orientation of P3ATs on the flat HOPG surface.32,36

Figure 12.

(a) Transmission electron microscopy (TEM) images of P3HT supramolecular structures on MWCNTs (P3HT/MWCNT mass ratio = 7). The width of the nanowires is about 12–15 nm. (Scale bar:100 nm). (b) The crystallization process of P3HT on MWCNTs monitored by in situ UV–vis spectroscopy at room temperature. [[P3HT] = 0.05 mg mL−1, P3HT/MWCNT mass ratio = 7, collecting time of the curves (from bottom to top at 600 nm): 0, 5, 15, and 30 min, 1, 1.5, 2, and 3 h] (c) UV–vis absorbance change at 600 nm of the P3HT suspension with (squares) and without (circles) CNTs during the crystallization process. Solid lines are the fitted first-order kinetics curves. (Reprinted from ref.76, with permission from Wiley-VCH).


The understanding of charge transport is a complex multiscale process that must take into account the structure of P3HT at different length scales. Moreover, the charge transport is expected to be highly anisotropic with predominant intrachain transport along the conjugated backbone of the polymer. For this reason, uniaxial orientation of a polymer film can be a way to generate geometrically and structurally simpler systems to measure the anisotropy of charge transport. There are different methods to orient the polymer chain direction along a given direction in the plane of a substrate, for example, (i) by mechanical rubbing of the polymer, for example, with a velvet cloth, (ii) friction-transfer, (iii) by deposition of a solution using directional casting or (iv) by the use of an orienting substrate. Several of these methods have been tested on P3ATs.

Heil et al. investigated the influence of mechanical rubbing on the field-effect mobility of P3HT OFETs. Oriented P3HT films with dichroic ratio of 5.1 were prepared but surprisingly, the charge mobility was improved when rubbing in the direction perpendicular to the source and drain contacts.77 Using the principle of the friction-transfer method developed by Smith and Wittmann to prepare oriented PTFE layers,78 Nagamatsu et al. fabricated highly oriented P3HT films with typical dichroic ratio 10–100 indicating a very high level of in-plane alignment of the polymer chains with the c-axis parallel to the direction of friction transfer.54 GIXD indicated that the π-stacking direction (b-axis) was perpendicular to the substrate plane, that is, the conjugated backbone adopts a “flat-on” orientation.

Epitaxy of semicrystalline polymers on aromatic organic crystals is an original and elegant method to grow highly crystalline and oriented polymer thin films with a controlled and regular organization of crystalline domains on a surface.79–81 Epitaxy of semicrystalline polymers can be achieved on a large variety of substrates including, inorganic substrates such as NaCl or KCl,82, 83 aromatic molecular crystals,79–81 polymeric alignment layers,78 and infusible aromatic salts.84 These substrates tend usually to direct the chain orientation along certain crystallographic directions of the substrate and act also as nucleating surfaces. For instance, para-terphenyl and anthracene lead to textured and oriented polymer films of polyethylene (PE) when the growth is performed on the surface of uniform single crystals of these aromatic molecules.79 Nucleation of numerous polymers and molecular materials is strongly enhanced on the surface of oriented PTFE substrates via so-called ledge-directed nucleation.85

One of the first examples of efficient epitaxial orientation of P3HT was obtained by using the directional epitaxial crystallization method in 1,3,5-trichlorobenzene (TCB) as proposed by Brinkmann and Wittmann.16, 27 The originality of this approach lies in the use of a crystallizable aromatic solvent, in this case TCB, which can successively play the role of solvent for the polymer and, once crystallized, the role of substrate for epitaxy. After orientation, TCB is readily removed by evaporation in primary vacuum leaving large areas of highly oriented P3HT. The morphology of the P3HT films was investigated by transmission electron microscopy, revealing an alternation of crystalline lamellae and amorphous interlamellar zones in both the bright field and the dark field modes (see Fig. 13). Orientation of P3HT on TCB has been explained in terms of 1D-epitaxy as the stacking periodicity of TCB molecules matches almost perfectly the repeat period of 3-hexylthiophene monomers in the P3HT chain.27 Using the epitaxial orientation of P3HT, Salleo and coworkers investigated the anisotropy of charge transport in oriented films.86 Higher mobility was observed in the direction of the polymer chains. The origin of this anisotropy was attributed to the existence of a high density of grain boundaries between fiber-like domains, limiting charge transport in the direction perpendicular to the chain axis.

Figure 13.

(a) Morphology of a thin P3HT film oriented by directional epitaxial crystallization, after selective removal of the 1,3,5-trichlorobenzene, as observed by optical microscopy with polarized incident light parallel to the P3HT chains (polarization direction indicated by the double arrow). (b) AFM phase-mode image showing the periodic lamellar structure of an oriented P3HT thin film. (c) Electron diffraction pattern of an oriented P3HT film. (d) Dark field image of a P3HT thin film obtained by using the 0 2 0 reflection. The crystalline lamellae appear in bright. (Reprinted from ref.16, with permission from Wiley-VCH).

The directional epitaxial crystallization method was also used to orient polyfluorenes, for example, poly(9,9′-di-n-octyl-2,7-fluorene) (PFO).87, 88 Dialkylfluorenes tend to form crystalline lamellae made of extended chains due to the higher persistence length of the polymer chain. No significant amorphous interlamellar zones were observed in PFO but rather only narrow grain boundaries.

Other substrates are also available for epitaxial orientation, for example, oriented polymeric substrates of poly(tetrafluoroethylene) (PTFE) or some aromatic salts such as potassium 4-bromobenzoate (KBrBz).89 These substrates offer the possibility to melt and anneal the thin films at high temperatures (up to 280 °C). This is particularly relevant as the melting temperature of conjugated polymers is usually quite high (∼240 °C for P3HT). Whereas epitaxial orientation of P3HT was not successful on PTFE, highly oriented and crystalline films of poly(9,9-bis(2-ethylhexyl)fluorene-2,7-diyl) (PF2/6) could be grown with a unique contact plane and a high in-plane orientation of the polymer chains forming extended-chain crystalline lamellae, the thickness of which is scales with Mw.85

As illustrated in parts (a) and (b) of Figure 14, epitaxial growth of P3HT on the surface of an aromatic salt (KBrBz) leads to highly crystalline, oriented and nanotextured P3HT films which consist of a regular network of interconnected semicrystalline domains oriented along two preferential in-plane directions (see schematic illustration in part d of Fig. 14).89 The overall crystallinity and the level of in-plane orientation of the P3HT films are controlled by the temperature of isothermal crystallization (Tiso). Well-defined ED patterns with sharp reflections obtained for Tiso = 180 °C [see Fig. 14(c)] indicate that the crystalline domains grow with a unique (1 0 0) P3HT contact plane on the K-BrBz substrate (so-called edge-on orientation). The P3HT chains are oriented along two preferred in-plane direction of the K-BrBz substrates namely the [0 ±2 1] directions. During annealing of the polymer film, the surface of the aromatic salt undergoes a topographic reconstruction resulting in regular nanostructured “hill and valley” topography, which templates and orients the growth of P3HT. Preferred orientation of P3HT crystalline domains occurs at step edges of the substrate and is favored by the matching between the layer period of P3HT and the terrace height of the K-BrBz substrate. This result opens new perspectives in terms of epitaxial crystallization of conjugated polymers as numerous organic (as opposed to the inorganic KBr) salts such as potassium acid phthalates can be used to generate new textures and nanomorphologies of P3HT films.

Figure 14.

(a) Bright Field TEM image of an oriented P3HT film grown on the 1 0 0 crystal surface of K-BrBz at 180 °C. Note the alternation of dark and brighter stripes in the oriented domains. The two preferential in-plane orientations of the P3HT chains (cP3HT) are indicated by arrows. (b) Scheme showing the organization of the nanostructured P3HT films grown on K-BrBz substrate. Crystalline zones are shown in red and amorphous interlamellar zones are in blue. (c) Electron diffraction pattern. The arrows indicate the two preferential in-plane directions of the π-stacking. (d) Schematic illustration of the preferential nucleation of crystalline P3HT domains at step edges of the KBrBz substrate. The lamellar period matches the height of the K-BrBz terraces (Reprinted from ref.89, with permission from American Chemical Society).

A further example of original polymer morphogenesis is illustrated in Figure 15 showing the so-called shish-kebab fibers of P3HT. Highly oriented fibers of regioregular (P3ATs) with a shish-kebab morphology have been prepared by oriented epitaxial crystallization in a mixture of 1,3,5-trichlorobenzene (TCB) and pyridine.90 The superstructure of the P3AT fibers consists of an oriented thread-like core several hundreds of micrometers long (the “shish”) onto which lateral crystalline fibrils made of folded polymer chains (the “kebabs”) are connected with a periodicity in the 18–30 nm range (see Fig. 15). The P3AT chain axis is oriented parallel to the fiber axis, whereas the π-stacking direction is perpendicular to it. The oriented character of the shish-kebab fibers results in polarized optical absorption and photoluminescence. The formation of oriented precursors by epitaxial orientation of polymer chains onto long needles of a molecular crystal, here TCB, appears to be an original alternative to the shish-kebab crystallization usually performed under external flow conditions.

Figure 15.

Mesoscale structure of oriented poly(3-hexylthiophene) (P3HT) shish-kebab fibers grown by directional epitaxial crystallization in a mixture of TCB and pyridine (30 mg of TCB and 20 μL pyridine). (a) Optical micrograph taken with crossed polarizers, showing highly birefringent TCB needles decorated with highly oriented P3HT fibers (the orientation of the polarizer and analyzer are indicated on the top right). Note the extinction of the birefringence of the fibers in the directions parallel to the polarizer and the analyzer. (b) TEM bright field image of a P3HT fiber with a shish-kebab superstructure after removal of both pyridine and TCB. (c) Topographic image (4 × 4 μm2) of P3HT shish-kebab fibers as observed by atomic force microscopy (AFM) operated in tapping mode. (d) Schematic structure of a shish-kebab fiber. The axis directions for the unit cells are indicated. (Reprinted from ref.90, with permission from Wiley-VCH).


The emerging field of conjugated polymer crystallization will certainly witness a gain in activity over the future decades mainly driven by the availability and synthesis of new generations of conjugated polymers, for example, low bandgap polymers combining different monomeric units within a larger “macromonomeric” unit91 or more complex block-copolymer architectures.92–95 Correlations between the molecular structure of the monomeric units made of several chemical moities and the crystalline packing in the unit cells (especially the π–π-stacking) need still to be established. In this perspective, P3HT is certainly a case system to test the applicability of experimental methods and insights gained when investigating conventional polymer crystallization, for example, using epitaxial crystallization methods, self-seeding or nucleating agents. Fundamental research in nucleation/growth control of conjugated polymers is promising. However conjugated polymers have intrinsic characteristic that make them significantly different from more flexible polymers, especially in terms of anisotropy of interactions in the unit cell and rigidity of the conjugated backbone that impacts directly the folding ability of the chains and correspondingly the crystallization mode (folded vs. extended chain crystallization). From the charge transport viewpoint, one important issue remains the possibility to crystallize high-Mw P3ATs in the form of extended-chain crystalline domains, trying to reduce the level of amorphous material and promoting long range crystalline perfection. Some preliminary studies in that direction have been reported recently95 suggesting that high-pressure crystallization of P3HT may lead to higher charge carrier mobilities. Further improvements can be foreseen with mastering of chemical defects in the chains of P3Ats.96 Crystalline perfection will no doubt benefit from the synthesis of P3ATs with low PDI, high regioregularity, and a low level of chemical impurities.


B. Lotz is acknowledged for enlightening discussions and numerous suggestions as well as his critical review of this manuscript. This work is dedicated to my mother who passed away, for her longstanding support and love all along these past years.

Biographical Information

original image

Martin Brinkmann obtained his PhD in physics from the University of Strasbourg in 1997. He spent several years as postdoctoral fellow at CNR in Bologna and M.I.T. in Cambridge (U.S.A) before moving to the Institut Charles Sadron (CNRS) in Strasbourg in 2000 as a senior scientist. His research interests include fundamental aspects of organic thin film growth using transmission electron microscopy, especially epitaxy of π-conjugated materials.