Semiconducting polymers have attracted a great deal of attention as essential materials for polymeric optoelectronic devices such as organic solar cells [1,2,3,4,5], field effect transistors [4,5,6,7,8,9,10,11], and sensors [12, 13]. The molecular features of semiconducting polymers are a rigid backbone and a flexible side chain, leading to a semicrystalline structure [14,15,16,17] and hierarchical chain dynamics [18, 19]. Since the crystallization kinetics depend on the molecular weight (MW) [20, 21] via chain dynamics, the aggregation states, which are generally frozen before reaching equilibrium, are assumed to also be a function of MW. On the other hand, the carrier formation after photoirradiation into an optoelectronic device is a controlling factor of performance because this determines the carrier density. Given that carriers are generated between ordered molecules in a crystallite, the carrier formation process depends on the MW. Thus, the MW effect on the photocarrier formation for semiconducting polymers in thin films has been a central issue for a better understanding and design of semiconducting polymer devices.

The aggregation states of semiconducting polymers in thin films have been examined by atomic force microscopy (AFM) [22,23,24,25,26,27] and grazing-incidence wide-angle X-ray diffraction (GIWAXD) [22,23,24,25, 28,29,30,31,32]. They were generally dependent on MW, as mentioned above. For example, in the case of regioregular poly(3-hexylthiophene) (P3HT), a typical semiconducting polymer, a low-MW sample formed a well-ordered crystalline phase, while a high-MW sample formed a less-ordered crystalline phase [33]. Shorter chains could reach a thermally stable state within a finite annealing time, while longer chains were kinetically trapped in a nonequilibrium state. Although the low-MW sample formed better crystallites in terms of ordering, many grain boundaries existed, which were the sharp interface between/among crystallites. This leads to limited charge transport due to a lack of electronic overlap between neighboring grains. Thus, the films used for the optoelectronic devices were intentionally fabricated to be in a nonequilibrium state even at the expense of crystalline ordering by controlling the preparation process, such as the selection of a solvent and an evaporation rate [34, 35], so that the grain boundary became more diffuse. In some cases, a strategy using a high-MW sample has also been taken. This is because in addition to the less crystalline ordering, longer chains cross the boundary as a bridging pathway for the conductivity [36]. For this reason, much effort has been made to synthesize semiconducting polymers with a higher MW [33].

Meanwhile, the thermal molecular motion of semiconducting polymers has been examined by differential scanning calorimetry (DSC) [14, 37,38,39,40,41], solid-state 13C nuclear magnetic resonance (13C NMR) [38, 42], dynamic mechanical analysis (DMA) [14, 39,40,41, 43] and quasi-elastic neutron scattering (QENS) [44, 45]. For example, P3HT exhibited β-, α1- and α2- relaxation processes corresponding to the molecular motion of alkyl side chains, thiophene backbones and crystalline regions, respectively [43]. The presence of such a hierarchy of molecular motion was also found for other poly(3-alkyl thiophene)s (P3ATs) [18, 46] and poly(benzodithiophene)s [47]. While the relaxation temperature for the β-process was determined by the size scale of the alkyl side chain, that for the α1-process was determined by the plasticizing effect of the alkyl side chain [46]. However, the MW dependence of the relaxation behavior for P3ATs is not yet fully understood.

Femtosecond transient absorption spectroscopy (TAS) enables us to gain access to the carrier formation process in semiconducting polymers after photoirradiation [43, 46, 48,49,50,51,52,53,54,55,56]. In the case of regioregular P3HT, free carriers, so-called polarons (P), are generated from singlet excitons (S) [48] as well as from polaron pairs (PP) [43], which are electron-hole pairs bound by the electrostatic force, under the zero-electric field. Combining the temperature-dependent TAS with DMA revealed that the release of the twisting motion of thiophene rings activated the P formation for the regioregular P3HT film [43]. Additionally, the activation energy for the twisting motion could be correlated with that for P formation [46]. However, the most studied P3HT thus far is polydisperse. In addition, the details of the effect of MW on the P formation process are not yet fully understood.

The objective of this study is to gain a better understanding of how the carrier generation process is related to the aggregation states, in addition to the chain dynamics, using well-defined P3HT in films. To do so, we first prepared monodisperse P3HT with various MWs from a commercially available polydisperse P3HT using preparative gel permeation chromatography (GPC). Then, the monodisperse P3HT thin films were characterized structurally and dynamically by GIWAXD and DMA, respectively. The photocarrier formation process in the films was examined by TAS measurements. Finally, integrating all the information obtained improves our current understanding of the carrier formation process for semiconducting polymers.

Experimental section

Materials and sample preparation

Regioregular P3HT was purchased from Sigma Aldrich Inc. (St. Louis, MI, USA). The number-average molecular weight (Mn) and polydispersity index (PDI) of P3HT were 26k and 2.4, respectively, as measured by GPC (LC-2000 Plus, Jasco Co., Tokyo, Japan) with polystyrene standards. The head-to-tail regioregularity of P3HT was confirmed by 1H nuclear magnetic resonance measurement (1H NMR, Unity Inova 500, Varian, CA, USA) operating at 500 MHz. Monodisperse P3HT (PDI < 1.1) with Mn values of 15k, 35k, 63k, 133k, and 171k was obtained by preparative GPC (PLC761, GL Sciences Inc., Tokyo, Japan) and used in this study.

The melting point (Tm) was determined by differential scanning calorimetry (DSC, an EXTAR DSC 6220, Hitachi High-Tech Science Co., Tokyo, Japan). The temperature ranged from 303 to 513 K, and the heating and cooling rates were 5 K min−1. Films were prepared from 2 wt% chloroform solutions on polyimide films and silicon, BaF2, and quartz substrates. The P3HT films were melted at 513 K under vacuum for 30 min, subsequently cooled to 474 K, and kept for 48 h to promote crystallization. The film thickness evaluated on the basis of the step height, which was created with a blade, using atomic force microscopy (AFM, E-sweep with an SPI 3800 controller, Hitachi-High Tech Science Co., Tokyo, Japan) was approximately 200 nm.


The crystalline structure of P3HT in the thin films was examined by GIWAXD measurements, which were conducted at BL03XU [57, 58] of SPring-8, Hyogo, Japan. The wavelength (λ) and incident angle of the X-rays and the camera length were 0.10 nm and 0.125° and 450 mm, respectively. Since the total reflection angle (θc) for the P3HT film was calculated to be 0.11°, X-rays should pass through the whole region of the sample along the direction normal to the surface. The scattering X-rays were collected as a function of angle (θ) with a high-throughput imaging Plate X-ray area detector, R-AXIS VII from Rigaku.

The loss modulus (E′′) of the P3HT film supported by a polyimide film was examined as a function of frequency and temperature by DMA (Rheovibron DDV-01FP, A&D Co., Ltd. Tokyo, Japan). Applying Takayanagi’s parallel mode, the contribution of the P3HT component to E′′ was extracted. The measurements were carried out at a heating rate of 1 K min−1 under a dry N2 atmosphere.

The photocarrier formation process in the P3HT thin films was examined by TAS measurements as a function of temperature. Our system consists of a transient absorption spectrometer (Helios, Ultrafast Systems, FL, USA) and a regenerative amplified Ti:sapphire laser (Solstice, Spectra-Physics, CA, USA). The laser provided fundamental pulses with λ = 800 nm and a width of 100 fs (fwhm) at a repetition rate of 1 kHz, which were split into two beams with a beam splitter to generate pump and probe pulses. One of them was converted into pump pulses, which were mechanically modulated with a repetition rate of 500 Hz, with λ = 400 nm with a second harmonic generator. Another was converted into probe pulses with the λ regions from 450 to 750 nm (visible (VIS)) or 850 to 1550 nm (near infrared (NIR)). The film was set in a cryostat (Oxford Instrument, Optistat CF-V, Oxfordshire, UK) and excited by pump pulses with a laser power of 200 μW. The transient absorption spectra, and thereby decays, were followed over the time range from −5 to 100 ps as a function of temperature, from 120 to 420 K, under vacuum.

Results and discussion

Aggregation structures

Panels (a)–(e) in Fig. 1 show GIWAXD patterns for monodisperse P3HT thin films with a thickness of 200 nm with various Mn values. As a general trend, an intense diffraction was observed at q = 3.9 nm−1 along the direction normal to the surface. Additionally, diffraction arcs or spots were also observed at q = 7.8 and 11.7 nm−1, which were twice and three times q for the most intense arc, at 3.9 nm−1. They could be assigned to (100), (200), and (300) of the monoclinic unit cell for P3HT, originating from the interlayer distance [32, 59, 60, 59, 61,62,63]. Additionally, the in-plane diffraction observed at q = 16.5 nm−1 was assignable to (020), arising from the π-π stacking of thiophene rings. Based on these assignments, it can be deduced that P3HT tends to form crystals with a monoclinic unit cell, in which the interlayer backbone (100) d-spacing distance is 1.6 nm and the aromatic π-π stacking distance is ~0.39 nm. Since the peaks were more clearly observed for samples with a smaller Mn, it seems reasonable to claim that the P3HT crystalline structure became more ordered with decreasing Mn.

Fig. 1
figure 1

GIWAXD patterns for P3HT thin films with Mn values of a 15k, b 35k, c 63k, d 133k, and e 171k. f Mn dependence of 1/g obtained from paracrystal analysis

To examine how the chain length affects the ordering of the crystalline structure of P3HT in the thin film geometry, the paracrystal analysis proposed by Hosemann [64, 65] was applied to sector-averaged intensity profiles along the out-of-plane and in-plane directions. In the paracrystalline lattice model, the individual vectors between adjacent unit cells vary in magnitude and direction due to the large displacement of the lattice points from their ideal positions, resulting in a loss of the long-range crystallographic order. Assuming that the paracrystal coordination statistics are in the form of a Gaussian distribution, the paracrystalline lattice factor of the hth-order reflection (Z(h)) is defined as follows [64, 65]:

$$Z\left( h \right) = \frac{{1 - \exp \left( { - 4\pi ^2g^2h^2} \right)}}{{\left( {1 - \exp \left( { - 2\pi ^2g^2h^2} \right)} \right)^2 \, + \, \left( {4\sin ^2\left( {2\pi h} \right)} \right)\exp \left( { - 2\pi ^2g^2h^2} \right)}}$$

where g is the standard deviation of the Gaussian distribution divided by the average lattice vector a. The integral breadth of a reflection (δβ) is expressed by

$$\left( {\delta \beta } \right)^2 = \left( {\frac{1}{{a^2}}} \right)\left( {\frac{1}{{N^2}} + \pi ^4g^4h^4} \right)$$

Here, N is the number of scattering units. Figure S1 shows 1D GIWAXD profiles for the P3HT thin films along the out-of-plane and in-plane directions. The (δβ)2 values were obtained from the 1D profiles, leading to the g value via Eq. (2). Panel (f) of Fig. 1 shows the 1/g value for the P3HT thin films as a function of Mn. The 1/g value increased with decreasing Mn, indicating that our hypothesis mentioned above is correct. The N value also increased with decreasing Mn (not shown). That is, the paracrystalline lattice was more ordered with decreasing Mn.

Molecular motion

To reveal the thermal molecular motion of monodisperse P3HT in the film, DMA measurements were conducted. Panel (a) of Fig. 2 shows the temperature dependence of E′′ for films of P3HT with various Mn values at a frequency (f) of 20 Hz. Two relaxation peaks were observed at ca. 200 and 310 K, which can be assigned to the β and α1 relaxation processes, corresponding to the side-chain motion and twisting motion of thiophene rings, respectively [43, 46]. In our previous study using a P3HT thin film, which was prepared by spin-coating and annealed at 333 K, above Tg and below Tm, the crystalline relaxation process, or the α2 process, was observed at approximately 390 K. However, it was not observed in these melt-crystallized samples.

Fig. 2
figure 2

a Temperature dependence of the dynamic loss modulus (E′′) for films of monodisperse P3HT with various Mn values at a frequency of 20 Hz. Mn dependence of b Tα1 and Tβ and c ΔH*α1 and ΔH*β for films of monodisperse P3HT

Figure 2b shows the Mn dependence of the β and α1 relaxation temperatures (Tβ and Tα1) for the monodisperse P3HT thin films. The Tβ and Tα1 values were determined by curve fitting the data in Panel (a) using two Gaussian functions. The Tβ was almost constant ca. 200 K independent of Mn. Tα1 also remained unchanged at 310 K from 171k down to 35k. However, this was not the case for 15k. The apparent activation energy (ΔH*) for the β and α1 relaxation processes was extracted on the basis of the relation between the frequency and the reciprocal number of the temperature, at which a peak was observed, as follows:

$${{{{{{{\mathrm{ln}}}}}}}}\;f\sim - {\Delta H}^ {\ast} /RT$$

where R is the gas constant. Panel (c) of Fig. 2 shows the Mn dependence of ΔH*β and ΔH*α1. The ΔH*β value was constant at 55 kJ mol−1, irrespective of Mn. On the other hand, ΔH*α1 was constant at 170 kJ mol−1 in the Mn range from 63k to 171k and decreased with Mn. These trends are in accordance with the Mn dependence of Tβ and Tα1 shown in Fig. 2b.

Since the α1 relaxation strongly affects the carrier formation rather than β relaxation [43, 46], the α1 relaxation process is the focus of this study. In general, ΔH* for the α process is proportional to the dynamic fragility index, which corresponds to ΔH* at Tg [66]. The value reflects the extent to which the cooperativity changes as T decreases from the liquid state through the glass transition [66,67,68]. Dargent et al. examined the fragility index for a uniaxially drawn film of poly(ethylene terephthalate) (PET), which is a typical semicrystalline polymer, to discuss the glass transition of the amorphous region by thermally stimulated depolarized current measurements [69]. When the drawing ratio increased, the crystallinity increased, and the fragility index decreased. This means that when the amorphous phase is confined in the space surrounded by crystallites, the extent of the segmental cooperativity is reduced. In this study, the crystalline ordering increased with decreasing Mn, as shown in Fig. 1f. Nevertheless, Fig. 2c shows that ΔH*α1 remained almost unchanged down to an Mn of 35k. Thus, it seems reasonable to claim that the α1 relaxation process for P3HT in the film arises from the twisting motion not in the amorphous phase but in the crystalline phase. This molecular picture was consistent with the twist glass transition proposed by Yazawa and Asakawa based on solid-state 13C NMR [38].

Then, the deviation of the ΔH*α1 value at the lowest Mn employed is discussed. The fragility index is also related to the dynamic heterogeneity [70]. Alternatively, more simply, the twisting motion can be easily released by the presence of chain ends that possess greater mobility. Either way, taking the above into account, it can be claimed that the α1 relaxation process is the cooperative twisting motion of several conjugated thiophene rings in the crystalline phase. Although we have previously identified the origin of the α1 process [43], it is a new finding that the relaxation process is in the crystalline region.

Photo-carrier formation

We now turn to the photocarrier formation process. Figure 3 shows typical TAS spectra in the visible and NIR regions acquired at room temperature for the film of monodisperse P3HT with an Mn of 35k. Time 0 was defined as the moment at which the bleaching intensity from the 0-2 transition at approximately 510 nm was maximized [46]. In this definition, the optical density (ΔOD) for PP in the visible region and S in the NIR region was also maximized at t = 0. This means that the conversion from hot excitons to S and/or PP was completed within an infinitely short time. The peak positions of PP and S for monodisperse Mn of 35k P3HT were 650 and 1170 nm, respectively, and were almost comparable to those for polydisperse P3HT [43]. Thus, it is conceivable that the energy states of PP and S were not sensitive to the crystalline ordering for P3HT, if any. On the other hand, the peak position for P, which appeared approximately 1 ps after excitation, was 910 nm. In our previous work, the peak position of P in the polydisperse P3HT thin film with a lower crystalline ordering was 1050 nm [43]. Therefore, it is clear that the peak position was dependent on the crystalline ordering of P3HT.

Fig. 3
figure 3

TAS spectra for monodisperse P3HT with Mn of 35k in a visible and b near infrared (NIR) regions, respectively, at 298 K. Transient absorption signals for P3HT at approximately 1170 nm, 650 nm and 910 nm correspond to singlet excitons (S), polaron pairs (PP) and polarons (P), respectively. The dotted and broken curves in Panel b denote the extracted spectra for S and P after 1 ps, respectively

ΔOD for S, PP, and P was expressed on the basis of the peak-top value. The detailed procedure of the spectral analysis is described elsewhere [46]. Figure 4 shows the fraction of the excited species just after the photo irradiation (t = 0) as a function of Mn at room temperature. In our previous study using a polydisperse P3HT film, P was not directly generated from hot excitons under this excited condition [43]. Basically, P was not directly generated from hot excitons either in the current study using monodisperse samples with Mn larger than 35k. However, monodisperse P3HT with an Mn of 15k, which exhibited the best crystalline ordering among the samples, was not the case. Thus, it can be claimed that the aggregation states of P3HT affect P formation from hot excitons at room temperature.

Fig. 4
figure 4

Mn dependence of fraction among S, PP and P at t = 0

The time dependence of ΔOD for S, PP and P was analyzed by coupled differential equations expressed by Eqs. (4)–(6) [46].

$$\frac{{d\left[ {{{{{{{\mathbf{S}}}}}}}} \right]}}{{dt}} = -\! \left( {k_{{{{{{{{\mathbf{S}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}0} + k_{{{{{{{{\mathbf{S}}}}}}}} \to {{{{{{{\mathbf{P}}}}}}}}} + k_{{{{{{{{\mathbf{S}}}}}}}} \to {{{{{{{\mathbf{PP}}}}}}}}} + k_{{{{{{{{\mathbf{S}}}}}}}} + {{{{{{{\mathbf{S}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}0}\left[ {{{{{{{\mathbf{S}}}}}}}} \right]} \right)\left[ {{{{{{{\mathbf{S}}}}}}}} \right] \\ + k_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}}\left[ {{{{{{{{\mathbf{PP}}}}}}}}} \right] + k_{{{{{{{{\mathbf{P}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}}\left[ {{{{{{{\mathbf{P}}}}}}}} \right]$$
$$\frac{{d\left[ {{{{{{{{\mathbf{PP}}}}}}}}} \right]}}{{dt}} = -\! \left( {k_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}0} + k_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{P}}}}}}}}} + k_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}} + k_{{{{{{{{\mathbf{PP}}}}}}}} + {{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}0}\left[ {{{{{{{{\mathbf{PP}}}}}}}}} \right]} \right)\left[ {{{{{{{{\mathbf{PP}}}}}}}}} \right] \\ + k_{{{{{{{{\mathbf{P}}}}}}}} \to {{{{{{{\mathbf{PP}}}}}}}}}\left[ {{{{{{{\mathbf{P}}}}}}}} \right] + k_{{{{{{{{\mathbf{S}}}}}}}} \to {{{{{{{\mathbf{PP}}}}}}}}}\left[ {{{{{{{\mathbf{S}}}}}}}} \right]$$
$$\frac{{d\left[ {{{{{{{\mathbf{P}}}}}}}} \right]}}{{dt}} = -\! \left( {k_{{{{{{{{\mathbf{P}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}0} + k_{{{{{{{{\mathbf{P}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}} + k_{{{{{{{{\mathbf{P}}}}}}}} \to {{{{{{{\mathbf{PP}}}}}}}}} + k_{{{{{{{{\mathbf{P}}}}}}}} + {{{{{{{\mathbf{P}}}}}}}} \to {{{{{{{\mathbf{S}}}}}}}}0}\left[ {{{{{{{\mathbf{P}}}}}}}} \right]} \right)\left[ {{{{{{{\mathbf{P}}}}}}}} \right] \\ + k_{{{{{{{{\mathbf{S}}}}}}}} \to {{{{{{{\mathbf{P}}}}}}}}}\left[ {{{{{{{\mathbf{S}}}}}}}} \right] + k_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{P}}}}}}}}}\left[ {{{{{{{{\mathbf{PP}}}}}}}}} \right]$$

Taking into account all possible transitions from one exciton to another (ij), the time dependence of ΔOD for S, PP and P was fitted with rate constants ki→j as fitting parameters. Figure S2 shows the analysis results for the thin film of P3HT with an Mn of 35k at various temperatures. If a ki→j value so obtained is finite, the corresponding transition process exists. Otherwise, the process does not. For all monodisperse samples employed, there were only four processes: PP→P, P→PP, PP+PP→S0 and S+S→S0. This is qualitatively the same as the result for polydisperse P3HT [46].

Figure 5 shows the temperature dependence of kPP→P for monodisperse P3HT with an Mn of 35k. The data for the different Mn values are shown in Figure S3. As a general trend, the kPPP value was almost constant in a temperature region lower than a threshold and then started to increase with increasing temperature. This result is in good agreement with that for polydisperse P3HT [46]. Here, the threshold temperature for kPPP was defined as TPPP, at which point two linear tangents obtained by linear-least squares fitting [43, 46] were crossed, as marked by an arrow in Fig. 5.

Fig. 5
figure 5

Temperature dependence of kPPP for a thin film of monodisperse P3HT with an Mn of 35k

Panel (a) of Fig. 6 shows TPPP and Tα1 for the monodisperse P3HT system as a function of Mn. TPPP basically coincided with Tα1. This means that the PP→P process was activated by the release of the twisting motion for thiophene rings in the crystalline phase at Tα1. That is, at a higher temperature beyond Tα1, the effect of the thermal molecular motion on the carrier formation was more dominant than that of the aggregation states. Here, it should be discussed that the timescales of the two processes, exciton formation, and α1 relaxation, are not matched. Once the α1 relaxation is released, the conjugation length of π electrons starts to fluctuate on a time scale of seconds. This means that even in a short time domain, the conjugation length may change, depending on the probability associated with the relaxation time of the α1 process. Thus, the two processes with different time scales can be correlated. In the case of P3HT with an Mn of 15k, TPPP seems to be lower than Tα1, although the error of the determination for TPPP was larger than that for P3HT with other Mn values. A plausible explanation for this might be based on the effect of chain end groups, which increased with decreasing Mn and were localized at the surface of crystallites [71].

Fig. 6
figure 6

Mn dependence of a TPPP and b ΔEPPP for monodisperse P3HT thin films. For comparison, Tα1 is also plotted in Panel a

The activation energy for the PP→P process (ΔEPPP) was extracted from the temperature dependence of kPPP above TPPP using an Arrhenius-type functional form as follows:

$$k_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{P}}}}}}}}} = A \cdot \exp \left( {\frac{{ - \Delta E_{{{{{{{{\mathbf{PP}}}}}}}} \to {{{{{{{\mathbf{P}}}}}}}}}}}{{RT}}} \right)$$

where A and R are a frequency factor and the gas constant, respectively. Panel (b) of Fig. 6 shows the Mn dependence of ΔEPPP for monodisperse P3HT. The ΔEPPP value was constant at 9.6 ± 0.1 kJ mol−1 above an Mn of 65k and decreased with decreasing Mn. This trend corresponds well to the Mn dependence of ΔH*α1 shown in Fig. 2c, implying that the carrier formation correlated with the chain dynamics at a temperature above TPPP. Taking into account that the α1 relaxation process observed by DMA arose from the twisting motion of thiophene rings in the crystalline phase, it is conceivable that P formation from PP was an event that occurred in the crystalline phase. Figure 7 shows ΔEPPP vs. ΔH*α1 for monodisperse P3HT with various Mn values. Data points were concentrated in a region around ΔEPPP = 9–10 kJ mol−1 and ΔH*α1 = 160–180 kJ mol−1. However, two data points for the samples with lower Mn values of 35k and 15k deviated from the region, although they were superimposed on a straight line through the origin.

Fig. 7
figure 7

Relationship between ΔH*α1 and ΔEPP→P for monodisperse P3HT. The dotted line is a visual guide

We finally come to the effect of Mn on the carrier formation in the film of P3HT via the aggregation states and the thermal molecular motion. At room temperature, which was below Tα1, the paracrystalline lattice was more ordered with decreasing Mn. P3HT with the lowest Mn employed had an advantage for P formation from hot excitons. However, once the temperature exceeded Tα1, the twisting motion for P3HT chains in the crystalline region started to occur. Then, the P formation process from PP was activated, as evidenced by the finding that Tα1 coincided with TPPP. Mn did not significantly impact the α1 relaxation, so P formation from PP was also insensitive to Mn. The Mn dependence of this process is in stark contrast to P formation from hot excitons. Thus, it can be claimed that the P formation process in the thin films of P3HT was regulated by Mn and temperature via the crystalline structure and the thermal molecular motion. That is, below Tα1, the Mn dependence on the P formation process appeared via the difference in the aggregation states in the crystalline phase. On the other hand, above Tα1, since the contribution of the extent of the twisting motion to the P formation became dominant over the one based on the crystalline structure, the Mn dependence was supposed to be less important.


Using monodisperse P3HT with various Mn values, which were prepared from polydisperse P3HT by preparative GPC, the effects of the crystalline structure and thermal molecular motion on photocarrier formation for monodisperse P3HT in melt-crystalized thin films were discussed on the basis of GIWAXD and DMA measurements and TAS spectroscopy. P3HT in thin films supported on silicon wafers formed the crystalline phase with the monoclinic unit cell. The 1/g value, which reflected the paracrystal order, gradually increased with decreasing Mn. Such P3HT crystallites with better ordering facilitated direct P formation from hot excitons. This was the case for neither monodisperse P3HT with larger Mn nor the polydisperse sample. Once the temperature exceeded Tα1 at approximately 310 K, the twisting motion of the thiophene rings of P3HT, named the α1 relaxation process, was released. Although the crystalline ordering increased with decreasing Mn, Mn affected neither the Tα1 nor ΔH*α1 values unless Mn was quite small. Considering that the α1 relaxation process was assigned to the twisting motion in not the amorphous phase but the crystalline one, in which chain ends were segregated out, the result can be understood. TPPP, at which P formation from PP started to be accelerated, was in good accordance with Tα1, indicating that P formation was correlated with the twisting motion of thiophene rings in the crystalline phase. Thus, it can be claimed that the P formation in the P3HT film was regulated by the thermal molecular motion in addition to the crystalline structure. We believe that this fundamental knowledge on the relation among crystalline structure, thermal molecular motion, and photocarrier formation in a typical semiconducting polymer of P3HT will be useful for designing and constructing polymeric optoelectronic devices.