Charge Transport in Highly Face-On Poly(3-hexylthiophene) Films

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Charge Transport in Highly Face-On Poly(3-hexylthiophene) Films Deepak Gargi,† R. Joseph Kline,‡ Dean M. DeLongchamp,‡ Daniel A. Fischer,‡ Michael F. Toney,§ and Brendan T. O’Connor*,† †

North Carolina State University, Raleigh, North Carolina 27695, United States National Institute of Standards and Technology, Material Measurement Laboratory, Gaithersburg, Maryland 20899, United States § Stanford Synchrotron Radiation Lightsource, Menlo Park, California 94025, United States ‡

S Supporting Information *

ABSTRACT: We report that the π-stacking direction in poly(3-hexylthiophene) (P3HT) films can be made to orient strongly out-of plane by uniaxially straining films in orthogonal directions, providing a valuable opportunity to evaluate charge transport in a very unusual microstructure for this material. The structure of the films was characterized using UV− visible spectroscopy, X-ray diffraction, and near-edge X-ray absorption fine structure spectroscopy, showing that unstrained films have a weakly edge-on stacking character with a large orientation distribution, whereas films strained biaxially by 100% in orthogonal directions have highly face-on stacking. In the biaxially strained films the face-on packing occurs while the P3HT long axis orientation is found to be only weakly anisotropic inplane. Charge transport is characterized in an organic thin-film transistor (OTFT) configuration, showing that the saturated field effect mobility in the biaxially strained films is greater than that for unstrained films for channel lengths ≤10 μm. The mobilities are found to have different channel-length dependence, attributed primarily to differences in the field-dependent charge-transport behavior, resulting in the mobility being comparable for channel lengths of 20 μm. The results suggest that edge-on packing is not a prerequisite for relatively high-field-effect mobility in P3HT-based OTFTs.

1. INTRODUCTION In polymer semiconductors, molecular organization is known to play a critical role in determining electrical properties.1−4 Charge transport depends on structural features ranging from local polymer chain order to larger scale crystalline size and quality. In particular, correlating structure and charge transport in poly(3-hexylthiophene) (P3HT) has been widely studied due to its early recognition as a high charge carrier mobility material. In regioregular P3HT, several early studies showed evidence that edge-on stacking (with the conjugation ring plane perpendicular to the substrate plane, as illustrated in Figure 1) improves in-plane charge transport in an organic thin film transistor (OTFT).1,5,6 This is logical given that the intermolecular π−π electronic orbital coupling is in the plane of the film for this packing configuration. However, recently several polymer semiconductors including {[N,N-9-bis(2octyldodecyl)naphthalene-1,4,5,8-bis(dicarboximide)-2,6-diyl]alt-5,59-(2,29-bithiophene)} [P(NDI2OD-T2)],7,8 diketo pyrrolo−pyrrole−bithiophene (DPPT) copolymers,9 and indacenodithiophene−benzothiadiazole (IDT-BT) copolymers,10 have been shown to have both face-on packing and high inplane field-effect mobility. These findings suggest that edge-on packing is not a prerequisite for high in-plane mobility in semicrystalline polymer semiconductors. Previous work on the role of out-of-plane packing on field-effect charge mobility in P3HT used processing methods such as spin-coating speed and tailored substrate surface energy that successfully altered the © XXXX American Chemical Society

Figure 1. Illustration of the P3HT backbone stacking orientation relative to a substrate for an edge-on (a) and a face-on (b) stacking configuration.

Received: May 22, 2013 Revised: July 29, 2013

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Figure 2. Schematic of the biaxial strain process. (a) PDMS stamp is placed on a P3HT film that is on an OTS-treated silicon substrate. The PDMS is then removed taking the P3HT film with it. (b) P3HT-PDMS stack is uniaxially strained to varying extents and then (c) laminated onto a second OTS-treated silicon substrate. (d,e) Process is repeated with the P3HT film being strained in the direction transverse to the initial strain direction. (f) P3HT film is then printed onto a third substrate for film characterization.

2. RESULTS AND DISCUSSION 2.1. Structure Characterization. The in-plane orientation of the polymer long axis was characterized using ultraviolet− visible (UV−vis) optical spectroscopy with polarized light. With the primary optical transition dipole moment (π−π*) in P3HT parallel to the polymer backbone axis,15 the average orientation of the polymer backbone can be determined by considering the dichroic ratio (R = A∥/A⊥) of the film, where A is the absorbance of light that is polarized parallel (∥) and perpendicular (⊥) to the strain direction. Straining the polymer film in one direction results in a significant increase in the dichroic ratio, as shown in Figure 3. By straining the film in the transverse direction, the dichroic ratio decreases, approaching ∼1 for films strained equally in two orthogonal directions. The dichroic ratio of ∼1 for biaxially strained films may indicate either an isotropic backbone or a biaxially oriented backbone; the specific orientation distribution cannot be determined because only its first moment is measured.16 To assist in determining if the polymer is biaxially oriented or isotropic inplane, X-ray diffraction measurements are performed on the films, providing information on the crystalline orientation distribution, given below. In addition to the polymer long-axis alignment, the absorbance characteristics provide information on the P3HT local order.17,18 A comparison of the absorbance of an unstrained P3HT film and a 100% × 100% biaxially strained film show clear differences, as given in Figure 3. To quantitatively compare film order, the absorbance of the aggregate P3HT is modeled using a weakly coupled Haggregate model.18−20 The model fitting parameters include the free exciton bandwidth (W) and Gaussian line width (σ). A decrease in W corresponds to an increase in the average conjugation length, and a decrease in σ is related to an increase in local aggregate order.18 Further information on the model is provided in the Supporting Information. The best-fit values for the unstrained films, and biaxially strained films (100% × 100%) are provided in Table 1. The biaxially strained films have a significantly smaller W than the unstrained film and a larger σ, corresponding to longer conjugation lengths but aggregates with greater disorder. From the absorption model, the percentage of aggregate P3HT is also estimated to increase from 46% for an unstrained film to 57% when the film is biaxially strained. While the percentage of aggregate P3HT is likely an overestimate, as discussed in the Supporting

local stacking orientation but also altered other important structural features (particularly order), making a direct comparison difficult.1,5,6,11 Here we use a biaxial strain method to investigate the difference in field-effect mobility for P3HT films with a variation in out-of-plane packing behavior. As we describe below, the biaxial strain process allows us to vary the out-of-plane molecular packing in P3HT films while largely maintaining order and general film quality. To control the out-of-plane orientation of P3HT, we employ a biaxial strain approach, as illustrated in Figure 2. This approach begins by spin-casting a P3HT solution onto an octyltrichlorosilane (OTS)-treated silicon substrate to promote edge-on packing and improved aggregate order. A polydimethylsiloxane (PDMS) slab is then laminated onto the P3HT film and quickly removed, resulting in the P3HT adhering to the PDMS and complete removal of the P3HT film from the initial substrate.12 The P3HT−PDMS stack is then uniaxially strained to varying extents, and while the elastomer is held under the specified strain, the film is printed onto a second OTS-treated silicon surface. To obtain biaxially strained films, we repeated the process with the uniaxially strained P3HT film delaminated a second time by the PDMS and strained in the transverse direction. The biaxially strained films are then printed onto a number of substrates including silicon and glass for structural and electrical characterization. The strained films are compared with films that are processed similarly without the application of strain. These unstrained films are first cast onto OTS-treated silicon, picked up by the PDMS elastomer, and then printed onto the appropriate secondary substrate without the application of strain. Using PDMS as a stamp to pick up and print P3HT films without applying strain has previously been shown to retain film quality and chargetransport characteristics.13,14 In addition, uniaxially straining P3HT films have been shown to align the polymer backbone in the direction of strain, resulting in significant charge transport anisotropy.14 Below, the structure of the P3HT films is characterized in detail, followed by charge-transport analysis of films in an OTFT configuration. To investigate the difference in charge transport between films with different out-of-plane πstacking orientation, we focus on printed P3HT films that are unstrained, and films biaxially strained by 100% × 100% in orthogonal directions. B

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Figure 4. Two-dimensional X-ray diffraction image plate results for (a) an unstrained film and (b,c) a 100% × 100% biaxially strained film. For the biaxially strained film, the X-ray beam is parallel to the first strain direction in panel b and parallel to the second strain direction in panel c.

unstrained films show features of weak edge-on stacking crystals with alkyl-stacking (h00) peaks out-of-plane and π−π stacking (0l0) peak detected both in-plane and out-of-plane with a broad orientation distribution. For the biaxially strained film the alkyl stacking direction (h00) is seen in-plane and π−π stacking (0l0) is primarily out-of-plane, indicative of highly face-on packing. For both the unstrained and biaxially strained films the scattering patterns from the two-orthogonal sample orientations are similar. While the unstrained film is expected to have a preferred out-of-plane packing orientation with an isotropic inplane angular distribution, this is not known a priori for the biaxially strained films. Two aspects of the biaxially strained films must be considered: (1) the consistency of face-on packing with in-plane azimuthal angle (ϕ)21 and (2) whether there is preferential in-plane alignment of the crystalline P3HT. To determine if the face-on packing is consistently observed with azimuthal angle, in-plane line scans from 2D images were taken at 30° intervals with the scattering vector going from parallel to perpendicular to the first strain direction. As given in Figure 5, it is found that face-on stacking is observed at each scattering vector angle. Whereas this does not probe all orientations, the consistent scattering behavior suggests that face-on stacking exists for all ϕ. Note that the difference in scattering intensity between scans does not imply preferential in-plane alignment due to the noncircular film shape resulting

Figure 3. (a) Normalized absorbance for the unstrained film and a 100% × 100% biaxially strained film using polarized light aligned along the two strain directions. (b) Optical dichroic ratio and mobility anisotropy (μfirst/μsecond) with increasing uniaxial strain and films strained by 100% along the first direction and then strained to varying extents in the second direction. The top value along the abscissa refers to strain in the first direction, and the bottom value refers to strain in the second direction. The mobility anisotropy and absorbance are shown to have a similar trend (slightly offset for clarity). Inset: illustration of a film uniaxially strained in the first direction, then a transverse strain applied in the second direction.

Information, the model suggests that strain-induced crystallization is occurring in the biaxially strained films. Note that the strained films are plastically deforming and the thickness of the unstrained film is ∼80 nm and the 100% × 100% biaxially strained film is ∼30 nm, as measured by spectroscopic ellipsometry. Grazing incidence X-ray diffraction (GIXD) was used to probe the crystalline P3HT packing characteristics. Twodimensional image plate results for an unstrained and a 100% × 100% biaxially strained film printed onto silicon substrates are given in Figure 4. For the biaxially strained films, the X-ray beam was incident along the primary strain directions. The

Table 1. Film Structure Characterization Parameters from UV-Visible and NEXAFS Spectroscopya film

W (meV)

σ (meV)

E00 (eV)

Rπ* top

Rπ* bottom

unstrained biaxial strain first direction biaxial strain second direction

184 126 132

79 88 92

2.03 2.06 2.06

0.08 ± 0.01 −0.47 ± 0.01 −0.38 ± 0.01

−0.21 ± 0.01 −0.52 ± 0.01 −0.38 ± 0.01

a Parameters include best-fit values for the weakly coupled H-aggregate absorption model applied to an unstrained P3HT film and a 100% × 100% biaxially strained film. These parameters include the exciton bandwidth (W), the Gaussian line width (σ) and the 0-0 transition energy (E00). From NEXAFS spectroscopy, the π* orbital dichroic ratio (Rπ*) values are given as measured for both the top and bottom of the films.

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Figure 5. (a) GIXD line scans for a 100% × 100% biaxially strained film taken at different orientations of the scattering vector with respect to the first strain direction. (b) ϕ scans of a 112% uniaxially strained film and a 100% × 100% biaxially strained film. The ϕ scan of the uniaxially strained film is for the in-plane (200) peak, while the ϕ scan of the biaxially strained film is for the (100) in-plane peak.

determined by collecting carbon K-edge spectra with multiple incident angles of polarized soft X-rays. The average orientation of the π* orbitals is given by the dichroic ratio (Rπ*), which is the difference between the extrapolated intensities at 90 and 0° incidence divided by the sum.3,6,30 Rπ* can vary from 0.7 for perfectly edge-on to −1 for perfectly face-on conjugated ring orientation.3,30 The values of Rπ* for the top surface and bottom surface of the unstrained and biaxially strained films is given in Table 1, as calculated from NEXAFS spectra provided in Figure S2 in the Supporting Information. For the unstrained film, the Rπ* for the top surface is indicative of a modest edgeon orientation, while the bottom surface is mildly face-on. The biaxially strained film has Rπ* values consistent with highly face-on stacking P3HT with similar Rπ* for both top and bottom surfaces. These values are consistent with the crystalline P3HT orientation in the bulk of the film, as given by the GIXD images, suggesting that the packing character is consistent throughout the film for both the unstrained and biaxially strained films. Combining the results of the UV−vis spectroscopy, NEXAFS spectroscopy, and X-ray diffraction, a number of structural characteristics can be described. The unstrained films have weakly edge-on packing with a large orientation distribution, whereas the biaxially strained films have a predominantly faceon packing motif. The crystalline material of the biaxially strained film is largely isotropic in the plane of the film, with weak preferential alignment of the polymer backbone along the first strain direction. Because the crystalline material is weakly anisotropic and the absorbance dichroic ratio of the biaxially strained film is ∼1, the amorphous material is likely isotropically distributed in-plane. Although the biaxial strain process involves significant mechanical manipulation of the film, the quality, as measured by optical spectroscopy and diffraction, appears to stay largely intact. In the biaxially strained films, the absorbance features suggest an increase in the average P3HT conjugation length and a slight decrease in aggregate quality. 2.2. Charge Transport. The in-plane charge transport in the films was studied in a bottom-gate, bottom-contact, thinfilm transistor configuration with a focus on the saturated field effect mobility. For uniaxially strained films, mobility anisotropy has been reported with the saturated field-effect mobility increasing parallel to the strain direction and decreasing perpendicular to it.14 As a transverse strain is applied to a uniaxially strained film, the mobility decreases along the first

in different X-ray scattering volumes at different azimuthal angles. To determine the in-plane distribution, a ϕ-scan of the (100) peak (the lamellar stacking peak, which is in-plane for face-on molecular orientations) was measured on relatively circular films. The background scattering was estimated and subtracted by performing a ϕ scan at a q value adjacent to the scattering peak (q = 0.45 Å−1) to minimize variation in diffraction intensity due to nonuniform scattering volume and to account for air scattering. The ϕ scan, given in Figure 5, shows that there is a slight preference of the face-on P3HT crystals to orient with the polymer backbone parallel to the first strain direction but with much less anisotropy than the uniaxially strained film. To quantitatively determine the degree of alignment of the crystalline material in the biaxially strained films, we applied an orientation factor using the ϕ-scan data. Order parameters have been developed for both uniaxial and biaxially aligned polymers.22,23 Applying a biaxial order parameter along the primary stain directions (as developed by White and Spruiell)23 results in no discernible preferential alignment along the second strain direction. Applying a uniaxial 2D order parameter, S = 2⟨cos2(ϕ)⟩ − 1,14,22 where S = 1 for complete uniaxial alignment and S = 0 for an isotropic distribution, results in S = 0.15 for alignment along the first strain direction. Comparing the order parameter to a uniaxially strained film by 112% [(200) peak], where S = 0.84,14 conveys the weak preferential alignment of the crystalline material in the biaxially strained film. While the ϕ scan only probes crystals with (100) in the plane of the film, the consistent face-on packing suggests that this orientation distribution will hold for all crystalline P3HT. Near-edge X-ray absorption fine structure (NEXAFS) spectroscopy measurements were also performed on the unstrained and 100% × 100% biaxially strained films to characterize the molecular orientation near the film surface.3,6,24 The partial electron yield (PEY) measurement provides surface-weighted information from ∼6 nm into the sample surface.3,25 NEXAFS spectroscopy also measures both crystalline and amorphous regions. The film tops were measured directly, and the film bottoms were accessed using mechanical delamination with the PDMS stamp.9 For the delaminated film bottoms, the NEXAFS measurement volume corresponds well to the region where the majority of charge transport is believed to occur in bottom-gate OTFTs.26−29 The average orientation of the conjugated plane (the 1s→π* transition, which is orthogonal to the thiophene ring plane) can be quantitatively D

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expect the angular dependence of the measured mobility to be accurate. As seen in Figure 7, the mobility in the uniaxially strained film is highly anisotropic, while the biaxially strained film is largely isotropic with a slightly higher mobility measured in the primary strain directions. The angular dependence of mobility is fit to an effective mobility, μeff = liμijlj, where μij is a 2D tensor and li and lj are unit vectors in perpendicular directions.31,32 The tensor values of the uniaxially strained film are found to be μ11 = 0.004 ± 0.001 cm2 V−1 s−1 and μ22 = 0.031 ± 0.001 cm2 V−1 s−1, and the tensor values of the biaxially strained film are μ11 = 0.0086 ± 0.0003 cm2 V−1 s−1 and μ22 = 0.0087 ± 0.0004 cm2 V−1 s−1. The off-diagonal components of the tensors are zero. While we find that the saturated field-effect mobility for the biaxially strained film is greater than the unstrained film for L ≤ 10 μm, the films have different channel-length dependencies resulting in roughly equivalent mobility for L = 20 μm, as shown in Figure 6. The channel-length dependence of the charge mobility is often attributed to the electric-fielddependent mobility or contact-resistance effects. A gated transmission line model was applied to the unstrained and biaxially strained films to determine if the difference in the channel-length dependence is due to contact resistances effects.33,34 In this model, the total resistance, RTOT = ∂VDS/ ∂ID is calculated from the output characteristics in the linear regime and plotted as a function of channel length for a given gate voltage, as shown in Figure S3 in the Supporting Information. With RTOT = RCH(L) + RC, the slope for a given gate voltage (VG) gives the channel resistance (RCH), and the intercept extrapolated to L = 0 gives the contact resistance (RC). The resistance analysis is considered for gate voltages of −30, −50, and −70 V, the range over which the field-effect mobility is determined. A negative contact resistance is determined for the biaxially strained films at all gate biases and unstrained films at low gate bias. The negative contact resistance values do not imply a negative resistance but rather that the channel conductance increases faster than the channel length decreases, suggesting a strong field-dependent mobility.34 The contact resistance is found to be VG-dependent, similar to previous reports,35 and the contact resistance is positive (0.82 MΩ) for the unstrained film at VG = −70 V. At this gate voltage, the channel resistance is 0.42 MΩ/μm, which results in the channel resistance being greater than the contact

strain direction and increases along the second strain direction, as shown in Figures 3b and 6 (inset). This results in biaxially

Figure 6. Saturated field-effect mobility as a function of channel length for the 100% × 100% biaxially strained films and unstrained films. (Inset) Saturated field-effect mobility of films strained by 100% in the first direction and strained to varying extents in the second strain direction for 5 μm channel length devices.

strained films with equal strain in two orthogonal directions having a mobility anisotropy approaching one. As given in Figure 6, for a 100% × 100% biaxially strained film, the mobilities in the two primary strain directions are similar, slightly favoring the second strain direction. Not only is the mobility similar along the primary strain directions, for OTFTs with channel length L ≤ 10 μm the mobility along the biaxial strain directions is greater than the mobility found in the unstrained film. To probe the angular distribution of the saturated field effect mobility, we patterned a circular array of source drain electrodes by photolithography, with a 15° angular resolution and L = 10 μm. Two P3HT films are compared, a uniaxially strained film (strained by 100%) and a biaxially strained film (100% × 100%), with results shown in Figure 7. We note that the test structures had measured field-effect mobilities slightly lower than the conventional test structures. While the cause of this lower performance is not completely understood, we

Figure 7. Field-effect mobility measurements for films printed onto a circular array of source-drain electrodes spread over 360° at an angular resolution of 15° and with 10 μm channel lengths. Left: the saturated field-effect mobility for a 100% uniaxially strained film. Right: the saturated field-effect mobility for a 100% × 100% biaxially strained film. Squares are the experimental data, and the dashed line is the effective mobility taken as a 2D mobility tensor. E

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3. CONCLUSIONS Previous research on the field-effect mobility in P3HT films with variation in out-of-plane packing typically used variations in materials or casting conditions to achieve the desired contrast. These differences included the use of P3HT with different molecular weight,1,39 regioregularity,1 ink formulation,5,11 or casting conditions.6 However, these processing methods not only alter film texture but also substantially modify aggregate quality, conjugation length, and tie-chain characteristics.2,18 This makes any correlation between out-ofplane stacking and field-effect mobility difficult to isolate. We employ a novel biaxial strain approach to modify P3HT film structure, resulting in an unusual out-of-plane π-stacking character. We find that the unstrained films have a broad πstacking orientation with both edge-on and face-on crystals, whereas the biaxially strained films have highly face-on packing. The biaxially strained film was shown to maintain general film quality, where optical absorbance measurements indicate a slight decrease in aggregate quality and also an increase in average conjugation length. The saturated field-effect mobility of the biaxially strained P3HT films was shown to be higher than the unstrained counterpart for L ≤ 10 μm, and similar for L = 20 μm. The channel-length dependence of the films is attributed primarily to a difference in field-dependent mobility. Surprisingly, out-of-plane π-stacking in P3HT films does not appear to have a major deleterious effect on charge transport. This suggests that a number of other structural features play more significant roles in the measured field effect mobility.

resistance for all channel lengths under consideration and suggests that the devices are not contact-limited. The difference in contact resistance between the unstrained and biaxially strained films may be associated with the molecular orientation of the P3HT in contact with the Au electrodes. Previous work has shown a difference in the energy level alignment of the P3HT−Au interface depending on the polymer orientation.36 The face-on stacking orientation was shown to have better band alignment between the Au work function and P3HT highest occupied molecular orbital, which may lead to a lower contact resistance. As stated above, the observed channel-length dependence may also be attributed to differences in the electric-field dependence of mobility.34,37 To consider the field dependence of the mobility we apply a Poole−Frenkel model to the saturated mobility given by,34 μ = μ0 exp(γ√E) where μ0 is the zero-field mobility, and E is the source-drain electric field (estimated to be VDS/L).34 We find that the mobility of the unstrained and biaxially strained films show a clear difference in field-dependence, as shown in Figure 8, and fit well to the

4. EXPERIMENTAL SECTION 4.1. Film Preparation. P3HT was purchased from Plextronics with a number-average molecular mass Mn = 50 kg/mol, a regioregularity of 99%, and a polydispersity of 2.1.40 The P3HT solution was dissolved in chloroform at 8 mg/mL and heated for a minimum of 4 h at 75 °C. The films were spin coated at 1000 rpm onto OTS-treated Si for 45 s at room temperature in a nitrogen-filled glovebox. Strained films were prepared by first laminating PDMS onto the cast P3HT films and removing the PDMS quickly, resulting in the P3HT film adhering to the PDMS slab and being fully removed from the OTS-silicon substrate. The P3HT−PDMS stack was strained in the direction parallel to the thin film, plastically deforming the P3HT. These strained P3HT films were then laminated onto a secondary substrate and the PDMS was removed. Unstrained films were prepared in the same manner without straining the P3HT−PDMS stack. For the biaxially strained films, the secondary substrate was OTS-treated silicon. The strain process was repeated with the uniaxially strained P3HT film being strained in the transverse direction. All transfer printing was done in a nitrogen-filled glovebox. Final substrates included glass and OTS−silicon with a native or thermal oxide layer. All substrates were cleaned with sonication in acetone, isopropanol, and DI water, followed by UV ozone treatment. The OTStreated silicon consisted of immersing the cleaned silicon substrates in a 0.002 mol/L solution of OTS in anhydrous hexadecane for 12 to 16 h, followed by sonication in chloroform, isopropanol, and deionized (DI) water. The substrate was then heated on a hot plate at 150 °C for 10 min, followed by DI water rinse. PDMS was prepared by mixing the elastomer with the cross-linker at a ratio of 15:1 and heating for 14 h at 60 °C while under low vacuum. The elastomer base-

Figure 8. Saturated field-effect mobility with applied electric field. The data is fit to a Poole−Frenkel transport model with the zero field mobility μ0 and prefactor (γ) given in the inset. In the biaxially strained films the 1 and 2 refer to charge mobility along the first and second strain directions, respectively.

Poole−Frenkel model, with prefactor (γ) values of (1.4 and 3.0) × 10−4 (V/m)−1/2, respectively. The field dependence of mobility in organic semiconductors is generally attributed to film disorder. While the biaxially strained films have an increase in mobility at short channel lengths, the strain process may increase interaggregate polymer disorder, resulting in a larger field dependence. These results are similar to a previous study on poly(2,5-bis(3-tetradecylthiophene-2-yl)thieno[3,2-b]thiophene) (pBTTT), where annealed pBTTT films showed an increase in saturated mobility and γ as compared with the ascast film.34 Annealing pBTTT films increases the film grain size and charge mobility,38 and it was speculated that the increase in γ is associated with an increase in grain boundary disorder in the annealed pBTTT films.34 Interestingly, in both the annealed pBTTT film and the biaxially strained P3HT film, the polymer backbone is likely aligned parallel to the substrate plane to a greater extent. This improved backbone alignment may also play a role in the improved mobility and electric-field dependence. F

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are also given. This material is available free of charge via the Internet at http://pubs.acs.org.

to-cross-linker ratio was chosen to maximize the fracture strain of the PDMS. 4.2. Structure Characterization. The absorbance measurements were performed using an Ocean Optics, Jazz spectrometer. The absorbance was measured for films on glass substrates and references to a clean glass slide. X-ray diffraction was performed at the Stanford Synchrotron Radiation Lightsource (SSRL) on beamlines 7−2 and 11−3. On beamline 7−2, the films were illuminated with an 8 keV beam at a constant incidence angle of ∼0.2° using a pointdetector geometry with soller slits to limit the angular acceptance in the plane of the sample film. For ϕ scans, a point detector was used and the sample tilt was adjusted to make the sample normal collinear with the rotation axis of the ϕ motor and maintain a constant incidence angle as the sample was rotated. On 11−3, the films were illuminated with 12.735 keV beam at an incident angle of 0.12° using a 2D geometry with an image plate (MAR2300) ∼33 cm from the sample. All sample chambers were purged with helium during the scattering experiments to reduce beam damage and background scattering. NEXAFS spectroscopy was performed at NIST beamline U7A of the National Synchrotron Light Source at Brookhaven National Laboratory. Carbon K-edge collection was performed in PEY mode with a grid bias of −50 V. The spectra collected were normalized with respect to carbon concentration by their intensity at 330 eV. The thickness of the films was determined using a Woollam variable-angle spectroscopic ellipsometer (VASE). A four-phase (ambient, P3HT, oxide, silicon) model was used and modeled using an isotropic Cauchy fit over the optical range of (750 to 1100) nm with both thickness and Cauchy parameters fit. 4.3. Charge Transport Characterization. The bottomcontact transistor test structures were fabricated by depositing gold/titanium (35 nm/4 nm) electrodes on highly doped (100) silicon wafers with 200 nm dry thermal oxide layer. The electrodes were deposited using e-beam deposition and patterned by photolithography. The channel length varied from 5 to 50 μm and the channel width was constant at 1000 μm. The test beds were also OTS-treated, as described above. Prior to device testing, the transferred films are annealed at 80 °C for 10 min, followed by being slowly cooled to room temperature. Annealing did not have a large impact on measured mobility but improved the on/off ratio of the device. The electrical properties were probed in a nitrogen-purged environment and measured using an HP 4156B semiconductor parameter analyzer. For transfer characteristics, the gate voltage was swept from 30 to −70 V for a drain source voltage of −2 (linear regime) and −70 V (saturation regime). Output characteristics were measured with gate voltage at 10 to −70 V in steps of −20 V and sweeping the drain source voltage from 10 to −70 V. The saturated field-effect hole mobility was calculated from a linear fit of ID1/2 versus VG, fitting a slope over a minimum 5 V range. The error bars are given as one-standard deviation of a minimum of four measured devices. In the gated transmission line model, the total resistance, RTOT = ∂VDS/∂ID, is calculated by taking the slope of the output characteristics in the linear regime using a VDS = 0 to −2 V with a step size of 0.5 V.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: (919) 515-5282. Fax: (919) 515-7968. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This material is based on work supported by the National Science Foundation under Grant No. CMMI-1200340. Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource, a Directorate of SLAC National Accelerator Laboratory and an Office of Science User Facility operated for the U.S. Department of Energy Office of Science by Stanford University.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

Additional details are provided on the H-aggregate absorption model and OTFT contact resistance analysis. NEXAFS spectra G

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The Journal of Physical Chemistry C

Article

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