Mapping of Domain Orientation and Molecular Order in Polycrystalline Semiconducting Polymer Films...

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Benjamin Watts , Torben Schuettfort , and Christopher R. McNeill *

Mapping of Domain Orientation and Molecular Order in Polycrystalline Semiconducting Polymer Films with Soft X-Ray Microscopy

We utilize scanning transmission X-ray microscopy (STXM) to study the domain structure of polycrystalline fi lms of the semiconducting polymer poly(9,9’-dioctylfl uorene-co-benzothiadiazole) (F8BT). By taking several images at different orientations of the fi lm with respect to the polarization of the X-ray beam, we are able to compute quantitative maps of molecular align-ment/order and molecular orientation, including both the backbone direction and phenyl ring plane orientation, as well as the in-plane and out-of-plane components. We show that polycrystalline F8BT fi lms consist of well-ordered micron-sized domains with the transition from one domain orientation to another characterized either by a smooth transition of orientation or by ∼ 200 nm wide disordered domain boundaries. The morphology of the disordered domain boundaries resemble the electroluminescence patterns observed pre-viously in F8BT light-emitting fi eld-effect transistors suggesting that charge trapping at these disordered domain boundaries facilitates charge recombina-tion in ambipolar operation. A relatively narrow distribution of local average tilt angles is observed that correlates with fi lm structure, with the ordered domains in general showing a higher tilt angle than the disordered domain boundaries. We also use secondary electron detection to image the surface domain structure of polycrystalline F8BT fi lms and demonstrate that the polycrystalline structure extends to the fi lm/air interface. Finally, we calculate ideal NEXAFS spectra corresponding to a perfect F8BT crystal oriented with the 1s – π ∗ transition dipole moment parallel and perpendicular to the electric fi eld vector of a perfectly linearly polarized X-ray beam.

1. Introduction

Organic fi eld-effect transistors (OFETs) based on solution processed conjugated polymers are a commercially attractive technology, allowing for low-cost, large area production on fl exible substrates. [ 1 ] Signifi cant advances have been made in

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DOI: 10.1002/adfm.201001918

T. Schuettfort , Dr. C. R. McNeill Cavendish LaboratoryUniversity of CambridgeJ J Thomson Ave, Cambridge, CB3 0HE, United Kingdom E-mail: [email protected] Dr. B. Watts Swiss Light SourcePaul Scherrer InstitutVilligen-PSI, CH-5232, Switzerland

recent years in the development of new high-performance materials [ 2–5 ] bringing commercialization closer to fruition. The convenience of solution processability, however, typically comes at the expense of structural order with solution-processed conjugated polymers exhibiting higher disorder than small molecules that may be prepared as single crystals. [ 6 ] Nevertheless, the highest polymer fi eld-effect mobilities are typically attained for semicrystalline polymers that show much higher mobili-ties than amorphous polymers. [ 7 ] Solution-processed semicrystalline fi lms are also typically polycrystalline, with charge transport sensitive to grain boundaries, molecular alignment and orientation. [ 8–10 ] The importance of fi lm microstructure on organic fi eld-effect operation is refl ected by the number of extensive structural charac-terization studies that have been carried out on semiconducting polymer fi lms. Key techniques that have been employed to study fi lm microstructure include scanning probe microscopies, [ 8 , 11 ] X-ray scattering, [ 12 , 13 ] electron [ 14 ] and optical microscopies. [ 15 ] Analysis of molecular orientation/alignment can be achieved through the study of optical anisotropy, [ 16 ] such as using ellipsometry [ 17 ] or the use

of cross-polarizers in combination with optical microscopy. [ 15 ] However, optical techniques do not necessarily probe molecular orientation directly, as the optical transition dipole moment is not necessarily aligned with the polymer backbone. [ 18 ] While Raman anisotropy measurements are an alternative, [ 19 ] optical techniques are limited in the spatial resolution that can be achieved, and scattering of light can complicate quantitative analysis. Two recently applied techniques, namely transverse shear microscopy [ 20 ] (a scanning probe technique), and dark-fi eld transmission electron microscopy, [ 21 ] have demonstrated the ability to image domain structure and orientation of poly-crystalline organic thin fi lms with high resolution.

An alternative approach for imaging domain structure and orientation is through the use of scanning transmission X-ray spectromicroscopy (STXM). Soft X-rays with energies close to the carbon 1s ionization energy induce resonant transitions from the carbon 1s orbital to unoccupied π ∗ and σ ∗ antibonding

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Figure 1 . a) Chemical structure of F8BT. b) Geometry of the angle-resolved NEXAFS experi-ment with α the angle between the plane of the sample and the incident X-ray beam. The molecular tilt angle, γ , is defi ned as the angle between the sample surface normal and the 1s to π ∗ resonance direction (yellow arrows normal to the plane of conjugation). The orientation of the C-C 1s to σ ∗ resonances tend to lie within or close to the plane of conjugation and are indicated in magenta (c, d) Angle-resolved TEY NEXAFS spectra of F8BT fi lms as-spun, (c) and annealed at 290 ° C, (d).

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orbitals. [ 22 ] As the transition dipole moment for the 1s to π ∗ transition is perpendicular to the plane of conjugation, near-edge X-ray absorption fi ne structure (NEXAFS) spectroscopy can be used in conjunction with linearly polarized synchro-tron radiation to probe molecular orientation and alignment of the polymer backbone. [ 23 , 24 ] Furthermore, STXM can be used to focus the soft X-ray beam to less than 20 nm in diameter, allowing for the possibility of molecular alignment to be probed with high resolution. [ 25 , 26 ] This approach has been previously used for orientation analysis of number of organic materials systems including Kevlar fi bers, [ 27 ] carbon nanotubes, [ 28 ] liquid crystals [ 29 ] and silk; [ 30 ] and recently employed for the imaging of the grain orientation of polycrystalline pentacene thin fi lms. [ 31 , 32 ] The converse effect has also been used to determine the degree of linear polarization of a soft X-ray beam from the linear dichroism of graphite. [ 33 , 34 ]

Here we utilize X-ray microscopy for the study of fi lms of the polycrystalline conjugated polymer poly(9,9’-dioctylfl u-orene-co-benzothiadiazole) (F8BT). F8BT is a well-studied poly-fl uorene copolymer that has found application in light-emitting diodes, [ 35 ] OFETs, [ 36 ] light-emitting OFETs, [ 37 ] optically pumped lasers [ 38 ] and solar cells. [ 39 ] In addition to mapping the in-plane domain orientation of F8BT fi lms (as has been demonstrated for pentacene fi lms [ 31 ] ), we extend previous analyses by deter-mining the local 3-dimensional resonance orientation, deriving maps of local structural order, and also estimate ideal spectra that demonstrate the full extent of the X-ray optical anisotropy of the material. Furthermore, we demonstrate the ability to image surface as well as bulk domain structure with STXM by taking advantage of the recently developed secondary electron detection mode. [ 40 , 41 ] Our results represent a signifi cant advance

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in the characterization of fi lm microstructure and are used to provide new insights into the interplay between fi lm structure and the operation of light-emitting, fi eld-effect tran-sistors based on polycrystalline F8BT.

2. Results and Discussion

Figure 1 shows the chemical structure of F8BT and angle-resolved Total Electron Yield (TEY) NEXAFS spectra of F8BT fi lms along with the geometry of the angle-resolved TEY NEXAFS experiment. The spot of the X-ray beam in the TEY NEXAFS experiment is of order 1 mm 2 and so probes macroscopic ori-entation. The split peak in the NEXAFS spec-trum of F8BT at 285 eV is due to transitions from the 1s to π ∗ molecular-orbitals and its intensity is sensitive to the orientation of the aromatic structures of the polymer backbone, while the broader peaks at 293 eV and 300 eV are due to 1s to C-C σ ∗ transitions and are additionally sensitive to side chain orientation. The orientation insensitive peak at 287.5 eV corresponds to 1s to C-H σ ∗ transitions. Here, we are chiefl y concerned with the 285 eV π ∗ resonance, which is oriented perpendicular

to the plane of the backbones aromatic rings as shown schemati-cally in Figure 1 b. Note that the orientation of the π ∗ resonance is thus also perpendicular to the local backbone orientation. Note also that while F igure 1 b depicts the F8BT monomer as planar, there is likely to be a torsional angle between the fl uorene and benzothiadiazole units. [ 41 , 42 ] Since it is not possible to deter-mine this torsional angle with the techniques employed in this paper, the orientation of the 1s to π ∗ transition dipole moment will refl ect that of the monomer unit averaging the contribu-tions from the fl uorene and benzothiadiazole units. The macro-scopic NEXAFS spectra of the as-spun, Figure 1 c, and 290 ° C annealed, F igure 1 d show moderate orientation effects, with the π ∗ peak highest for glancing incidence, which could be naively interpreted as indicating that the π ∗ resonance is preferentially directed out of the plane of the fi lm and hence the F8BT aro-matic structures are lying preferentially in the plane of the fi lm as suggested by previous studies. [ 43 , 44 ] However, a quantitative analysis of the tilt angle of the 285 eV π ∗ resonance yields an average tilt angle of 47.5 ° (measured between the resonance vector and the sample surface-normal) – a slightly in-plane ori-entation of the resonance vector that corresponds to a prefer-ence for the F8BT aromatic rings to stand slightly more edge-on than 45 ° . The false conclusion above originates from a failure to appreciate that the sign of the spectral anisotropy is relative to the so-called “magic” tilt angle of 54.7 ° , rather than the intuitive 45 ° angle. [ 22 ] The “magic” tilt angle is described by Stöhr [ 22 ] via arguments based on the geometry of a planar fi lm and the fi eld orientations of a polarized X-ray beam. Since these elements are common to all light-based techniques for the measurement of thin fi lms, Stöhr’s “magic” angle is likely generally relevant and thus our 47.5 ° result is not necessarily in contradiction

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Figure 2 . 5 μ m × 5 μ m unprocessed X-ray microscopy images of poly-crystalline F8BT fi lms produced by annealing at a) 150 ° C, b) 200 ° C, c) 240 ° C, and d) 290 ° C.

to previous studies. [ 43 , 44 ] Note that the GIWAXS experiments reported by Donley et al. [ 43 ] are sensitive only to the crystal-line portions of the sample and so cannot be directly compared to photon absorption based measurements, such as NEXAFS and STXM, whose sensitivity is not dependent on crystallinity. Note also that our fi lms here are quenched following annealing rather than slow-cooled as adopted by Donley et al. [ 43 ] Thus our fi lms are likely to be less crystalline, preventing direct compar-ison with these previous GIWAXS measurements. We adopt the quenching protocol here as it is employed in transistor fab-rication [ 37 ] facilitating direct comparison with the light emission profi le in polycrystalline light-emitting F8BT FETs.

Annealing of F8BT fi lms above the glass transition tempera-ture ( ∼ 140 ° C [ 43 ] ) has been observed to produce polycrystalline fi lms with micron-sized semicrystalline domains. [ 15 ] While these domains may be readily imaged by conventional polar-ized optical microscopy for more aggressive annealing, smaller domains are likely to be present in fi lms annealed at lower tem-peratures that are not resolved by optical microscopy. Knowledge of fi ner structure is especially important for photonic applica-tions of F8BT such as lasing where scattering from domains rep-resents a signifi cant loss mechanism. [ 38 ] Figure 2 presents raw transmission polarized X-ray microscopy images of polycrystal-line blend fi lms annealed at temperatures between 150 ° C (just above the glass transition temperature [ 43 ] ) and 290 ° C taken at 285 eV. For the fi lm annealed at 150 ° C, domains can be clearly discerned of size ∼ 100 to 150 nm. Domain size increases sig-nifi cantly with annealing at 200 ° C (domain size ∼ 500 nm) and saturates for higher anneal temperatures with little difference between the domain size of 240 ° C- and 290 ° C- annealed fi lms. Contrast in these images is due to differences between the in-plane orientation of the 1s – π ∗ transition dipole moment (per-pendicular to the polymer backbone) relative to the major elec-tric fi eld vector of the incident polarized X-ray beam. (The X-ray beam is incident normal to the plane of the fi lm in all of the STXM experiments discussed here and hence the electric fi eld vector is always parallel to the plane of the substrate.) Rotation of the sample with respect to the polarization of the incident synchrotron light, Figures 3 and 4, brings about a change in contrast, with the image taken with the sample rotated by 90 ° representing a negative of the original. (The images in Figure 3 map the optical density (OD = ln( I 0 /I )) measured at each point, where I 0 is the incident X-ray fl ux and I is the transmitted X-ray fl ux.) Furthermore, when an image is acquired at a photon energy of 320 eV where absorption is due to ionization of 1s electrons to the continuum alone (due to the absence of mole-cular resonances at this energy), Figure 3 d, little contrast is observed with local variations in intensity refl ecting only vari-ations in fi lm thickness. Due to the strong anisotropy observed in these images, the observed domains are likely to span the thickness of the fi lm.

Overlapping the three images of Figure 3 taken at different rotations enables quantitative analysis of domain orienta-tion and structural order. The variation in optical density with sample rotation can be used to determine the average, local directed resonance intensity of each pixel. Figure 4 defi nes the co-ordinate system adopted for the analysis, with X, Y, Z representing the laboratory co-ordinate system and S1, S2, n the sample co-ordinate system that is rotated by θ . The fi lm is

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oriented with its surface normal n directed along the Z axis, with the incident, elliptically polarized X-ray beam propagating along the Z axis and major and minor electric fi eld vectors oriented along the X and Y axes respectively. The elliptical polarization of the X-ray beam can, for the purposes of this work, be viewed as equivalent to 80% linear polarization. The sample is rotated about the surface normal with θ representing the rotation of the sample frame with respect to the electric fi eld vector of the incident X-ray beam. O is the local transition dipole moment and subtends an angle of γ to n . The aligned and thickness normalized optical density data can be fi t to a cosine squared relation in order to fi nd the local in-plane reso-nance orientation and magnitude using:

OD = M

[P cos2($ +2 ) + (1 − P) sin2($ +2 )

]+ C

(1)

where C is a constant resonance intensity observed at all angles and P is the Stöhr polarization value of the incident X-ray beam ( P = |ΕX |2

|ΕX |2 + | ΕY |2 = ∼0. 9 ). M is defi ned as the magni-tude of the in-plane component of the normalized, directed res-onance intensity, D ( D is oriented along O but with magnitude ∣∣O∣∣2

/a where a is the optical density of the fi lm at 320 eV) with

β the direction of the in-plane component of D in the sample frame. Figure 5 plots the results of this fi tting, with Figure 5 a depicting the fi tted value of M at each point, Figure 5 b the fi tted value of C for each point, and Figure 5 c the fi tted value of β for each point. Since the fi lm is thin relative to the length of the polymer chains (the fully extended length of an individual chain with mass M p ∼ 163 kg mol − 1 is ∼ 400 nm, while the fi lm imaged here is ∼ 100 nm thick), we can assume that the polymer chains prefer to lie within the plane of the fi lm, an assumption that is supported by ellipsometric and grazing incidence X-ray diffraction studies. [ 43 ] Therefore we can deduce the preferred

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Figure 3 . a–c) OD transmission images with different rotations of the sample about the normal to the plane of the fi lm. The relative orientation of the electric fi eld of the probing X-ray beam is indicated by an arrow in each image. d) Optical density image at 320 eV that shows no change with sample rotation and is used to normalize images for small variations in fi lm thickness.

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Figure 4. Geometry of the co-ordinate systems adopted for the analysis. X,Y,Z represents the laboratory co-ordinate system, S1,S2,n the sample co-ordinate system, and n is the fi lm surface normal. The X-ray beam propagates along the Z axis with its major electric fi eld vector oriented along the X axis. The sample is rotated about the surface normal with θ representing the rotation of the sample frame with respect to the major electric fi eld vector of the incident X-ray beam. O

_ is the local transition

dipole moment, which subtends an angle of γ to n and β to S1.

polymer backbone orientation as perpendicular to the in-plane resonance direction, β , and overplot Figure 5 c with black lines representing the orientation and fully extended length of the polymer backbones. (Note that the polymer chains may not be fully extended, especially in disordered regions; we use this length for the black lines for illustrative purposes only.)

We now extend the analysis to consider the unobserved out-of-plane resonance intensity by considering two facts. Firstly, the sample fi lm is made of a single material and the data has been normalized for thickness/density variations. Therefore, the sample’s spectrum can only vary according to the alignment of the molecules with respect to the electric fi eld of the incident X-ray beam. Further, if one were to sum such spectra measured over all electric fi eld orientations, then the resulting total spec-trum would be independent of the molecular alignment and orientation. The implication of this is that the total π ∗ resonance magnitude, which involves M , C and the unobserved out-of-plane resonance magnitude (i.e. all orientations), should sum to a constant value at each point in our data set.

The second fact our analysis relies on is what the C term from our fi tting results represents in a three-dimensional sense. The initial description of C given above was the degree to which the π ∗ resonance was observed equally in all in-plane directions. In three dimensions however, C could be interpreted in two

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Figure 5 . Results of the fi tted of the transmission images to equation 1 , showing the positional dependence of M , (a), C , (b), and β , (c). While the colors in (c) indicate β , corresponding to the in-plane orientation of the 285 eV π ∗ resonance, black lines have been overplotted to indicate the local orientation and estimated (fully extended) length of the polymer backbone.

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ways. In one case, the polymer backbones could be aligned in one direction, but with the aromatic ring structures oriented to equally face the two dimensions perpendicular to the backbone axis, possibly via twisting of the polymer backbone. However, this would require that a signifi cant proportion of the ∼ 400 nm long molecular chains be oriented in the out-of-plane direc-tion, which is physically unreasonable for a ∼ 100 nm thick fi lm. The second, and only physically reasonable case is that C must indicate the degree to which the fi lm is amorphous in a three dimensional sense also. Thus C can be thought of as arising from an isotropic transition dipole moment of magnitude


√

C , giving rise to a total isotropic resonance intensity of 3 C for all three-dimensionally possible X-ray polarization directions.

Now, with knowledge of the total summed resonance inten-sity (equal to the sum of the directed and isotropic contribu-tions) it is possible to deduce the out-of-plane component of the directed resonance intensity and so have full knowledge of the π ∗ resonance in each pixel of the imaged area. Since the total summed resonance intensity can be expressed as

T =∣∣D ∣∣ + 3C (2)

we make an estimate of the total summed resonance intensity by assuming that the out-of-plane component of the directed resonance intensity is zero in at least one pixel of our meas-ured data set (this assumption will be subsequently verifi ed). Since a lack of out-of-plane directed resonance intensity must be refl ected in a greater observed intensity of the in-plane directed resonance, M , and/or the isotropic component, C , then T must be equal to the maximum observed value of M + 3 C . In other words, in pixels where

∣∣D∣∣ = ∣∣D‖

∣∣ = M then Equation 2 becomes

T = M + 3C (3)

The histogram in Figure 6 shows the range of summed reso-nance intensity values observed (i.e. T = M + 3C ), with the inset showing further detail in the region of the highest observed values. Discarding a few outliers, we are left with a maximum observation of the summed, normalized resonance intensity of T = 3.4. We now have suffi cient information to calculate the full three dimensional orientation and degree of alignment within the fi lm as follows.

Firstly, we already know that the total amorphous intensity at each pixel is given by 3 C . Therefore, the total aligned intensity magnitude,

∣∣D∣∣ for any point (where D may now have an out-

of-plane component) is given by:

∣∣D

∣∣ = 3.4 − 3C (4)

Secondly, the local tilt angle of the resonance orientation direction can be calculated via:

( = sin−1

(√M∣∣D∣∣

) (5)

where γ is the angle between the aligned resonance direction and the sample surface-normal. The square root operation is required in equation 5 because we are dealing with a cosine squared relationship, as can also be seen in Equation 1 . Figure 7 a shows a map of the local out-of-plane tilt angle, γ . Here, one can see that a narrow range of tilt angles dominates the imaged area, but lower tilt angles tend to occur in the more amorphous regions (i.e. higher values of C , as shown in Figure 5 b). (Note that the tilt angles presented in Figure 7 still represent an average tilt angle of each pixel and there will still likely be a dis-tribution of tilt angles within each pixel.) A histogram of these tilt angles, grouped by degree of local order (defi ned subse-quently) is presented in Figure 7 b, with numerical details listed in Table 1 . As expected from Figure 7 a, the most well-ordered fi lm areas (purple) show a narrow distribution (FWHM of 7.4 ° )


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Figure 6 . Histogram showing the distribution of summed resonance intensities. The inset highlights the upper tail, with a value of 3.4 chosen for the maximum observed total summed resonance intensity.

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Figure 7 . a) Map of local molecular tilt angle. b) Histogram of molecular tilt angle grouped according to local degree of order (red: < 25%; blue: 25% – 50%; green: 50% – 75%; purple: > 75%). The vertical dashed line represents the “magic” tilt angle of 54.7 ° .

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Table 1. Details of molecular tilt angle distributions shown in F igure 7 . The overall mean tilt angle is 47.2 ° .

Local Order Proportion [%] Distribution Center [ ° ] FWHM [ ° ]

< 25% 1.4 38.5 36.3

25%–50% 16.7 44.0 19.1

50%–75% 42.8 47.5 11.4

> 75% 39.0 49.4 7.4

of tilt angles, centered at 49.4 ° and thus indicating a slightly more out-of-plane orientation of the 285 eV π ∗ resonance than the 54.7 ° “magic” tilt angle that NEXAFS spectra are normally interpreted with respect to, but a slightly in-plane orientation relative to the intuitive 45 ° angle. In terms of molecular orienta-tion, this result corresponds to an average 49.4 ° angle between the plane of the aromatic ring units in a highly ordered region of fi lm and the surface plane. The distributions representing less well-ordered regions of fi lm in Figure 7 b display a clear trend of lower average tilt angle and broader distribution as the order decreases, in agreement with the tilt angles mapped in Figure 7 a (referenced with Figure 5 b). The accuracy of the measured tilt angles can be expected to depend on the appro-priateness of the choice of T, the maximum observed summed resonance intensity. Repetition of the above calculations show that a variation in T of ± 0.1 ( ∼ 3%) causes a mean tilt angle vari-ation of ± 1.4 ° (Further details can be found in the Supporting Information). For our choice of T = 3.4, the overall mean tilt angle is 47.2 ° , in close agreement with the 47.5 ° measured by macroscopic, surface sensitive TEY NEXAFS (Figure 1 ). An exact match between the microscopic and macroscopic data is not necessarily to be expected however, given that the STXM measurements are performed in transmission (and are there-fore representative of the bulk of the sample), while the TEY NEXAFS experiments measure the sample surface (depth sen-sitivity of ∼ 3 nm [ 45 ] ). This relationship between surface and bulk structure is discussed further below.

With our extensive knowledge of the system, it is now pos-sible to quantitatively parameterize and map the local degree of alignment of polymer chains (i.e. the local degree of order). The directed resonance intensity, D , closely follows the local degree of molecular alignment and so the local degree of order is well represented by the percentage of the total resonance intensity that lies within D . In other words:

Or der = D/3.4 × 100%

(6)


In this way, 100% order would correspond to all neighboring chains aligned in the same direction, and 0% order would refl ect an amorphous region with no correlation between the orientation of neighboring chains. Figure 8 a maps the percentage order observed in the imaged area, with well-ordered domains showing order approaching 100%, surrounded by amorphous regions reaching below 20% order. Furthermore, Figure 8 a clearly shows that the interfaces between the well-ordered,


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micron-sized domains are not sharp, but involve either smooth transitions or ∼ 200 nm wide, highly disordered boundary regions. In order to visualize the disordered domain bounda-ries more clearly, we also plot the rate of change of backbone orientation, Figure 8 b, that highlight these disordered regions. Interestingly, the morphology of these inter-domain regions closely resemble the electroluminescence pattern observed in light-emitting fi eld effect transistors based on polycrystalline F8BT fi lms. In a previous publication comparing electrolumi-nescence (EL) and cross-polarized photoluminescence (PL) images, a poor correlation between the structure of EL and PL was noted, with the regions of higher EL were tentatively attrib-uted to areas of high current density. [ 16 ] However our imaging of domain structure shows a strong correlation between the EL image and domain boundaries suggesting that recombination is occurring predominantly in these disordered inter-domain regions. As the charge carrier mobility of the well-ordered domains is likely to be signifi cantly higher than in the disor-dered inter-domain regions, electrons and holes will spend less time in the inter-domain regions resulting in a lower charge density there. The disordered inter-domain regions will have

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Figure 8 . Plots of a) local molecular order and b) local rate of change of resonance, and hence also the in-plane curvature of the polymer back-bone. Here, ‘local’ refers to the 50 nm pixel size of the measured data.

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a lower charge carrier mobility with the high degree of struc-tural disorder likely to facilitate charge trapping. Therefore we expect a relatively high charge density in these disordered regions that will facilitate charge recombination. Thus at least for polycrystalline F8BT light-emitting fi eld-effect transistors, the electroluminescence maps do not provide a visualization of lateral current fl ow but rather refl ect regions of high charge density and hence high recombination effi ciency.

The images above have all been acquired through (or derived from) transmission imaging and therefore represent the bulk structure of the fi lm. The strong anisotropy observed demon-strates that these features extend vertically through the majority of the fi lm; however it is not clear from the transmission data whether this bulk structure extends to the surface. Since charge transport in OFETs is an interfacial process, with the accumu-lation layer localized within the fi rst few nm of the fi lm/die-lectric interface, sensitivity to surface structure is paramount when correlating fi lm structure and device processes. Recently, we have demonstrated the ability to simultaneously acquire images of surface structure with scanning transmission X-ray microscopy [ 41 ] through the use of a channeltron detector that detects photoelectrons emitted from the surface of the fi lm. [ 40 ] Figure 9 presents simultaneously acquired bulk and surface images acquired on a polycrystalline F8BT fi lm. The surface imaged here is the “top” of the fi lm or the fi lm/air interface (as opposed to the substrate/fi lm interface). Light-emitting fi eld-effect transistors are typically fabricated with a top gate geom-etry, meaning that the fi lm/air interface will interface with the subsequently deposited dielectric layer and thus represents the active, accumulation layer of the fi lm. The morphology of the surface image in Figure 9 b shows very similar structure to that of the bulk image shown in Figure 9 a, indicating that the bulk polycrystalline structure is also present at the surface. There are some subtle differences between bulk and surface images, with some structures present on the surface that are not refl ected in the bulk. However in general there is a strong correlation between surface and bulk images (see Figure 9 c) demonstrating that charge transport in the accumulation layer is directly affected by the polycrystalline domain structure.

Since we have shown, both through quantitative tilt angle anal-ysis and imaging, that the surface and bulk fi lm structures are very similar, we can check that our earlier assumptions are correct by comparing the STXM data to the TEY NEXAFS data shown in Figure 1 c. Such a comparison is possible by using the detailed, STXM-derived molecular alignment and orientation data to simu-late macroscopic, polar-angle dependent NEXAFS experiments. Since the imaged fi lm area is small relative to a statistically signifi -cant number of domains, we integrate about θ in order to smooth out any bias in β , the in-plane resonance orientation (relative to the sample coordinate system). This can be calculated via: [ 22 ]

I = P 13 M

[1 + 1

2 3 cos2 (") − 1 3 cos2 (() − 1)]

+ (1 − P )[

12 M sin2 (()

]+ C

(7)

where I is the observed resonance intensity and α is the polar angle (measured as depicted in Figure 1 b). Figure 10 a compares the calculated polar dependence of resonance inten-sity compared to the macroscopic measurement (Figure 1 c). The shaded area represents the spread of values observed,

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Figure 10 . a) Comparison of the macroscopic NEXAFS measurements made at different polar angles of X-ray from Figure 1, circles, to a simulation of the same experiment calculated from the STXM-derived information shown in Figures 5 and 8 (with rotational averaging to remove any in-plane direction bias). The shaded area shows the range of values observed, while the solid line shows the average over the imaged area. The horizontal dashed line represents the resonance intensity that would be observed for a completely amorphous sample. The data sets all converge at the “magic” angle of X-ray incidence for 80% linear polarization of 52.3 ° . b) Estimate of ideal F8BT NEXAFS spectra that might be measured in an experiment involving a perfectly linearly polar-ized X-ray beam incident on a perfect F8BT crystal oriented exactly to align the electric fi eld vector of the X-ray beam parallel (0 ° ), perpen-dicular (90 ° ) or at the “magic” incidence angle (54.7 ° ) to the 1s – π ∗ resonance direction (which is normal to the plane of the aromatic ring structures).

280 300 320

0

1

2

3

Photon Energy (eV)

Res

onan

ce In

tens

ity (

a.u.

)

90°54.7°0°

290 310

Polar Angle (degrees)

Res

onan

ce In

tens

ity (

a.u.

)

0 15 30 45 60 75 90

0.8

1

1.2

1.4

1.6

1.8

2Imaged areaAmorphousMacroscopic (90°)Macroscopic (50°)Macroscopic (20°)

(a)

(b)

Figure 9 . Comparison of bulk (a) and surface (b) images of the polycrystalline domain structure of an F8BT fi lm annealed at 290 ° C. c) Highlights differences between the bulk (red) and surface (blue) images.

while the solid line shows the total intensity averaged over all pixels of the imaged fi lm area. Excellent agreement is seen in Figure 10 a between the calculated angular dependence of resonance intensity from the microscopy data and the macro-scopic TEY NEXAFS measurements, confi rming the validity of the assumptions made to deduce the out-of-plane resonance components.

Note that the resonance intensity value of 3.4 that was assumed earlier represents the largest possible observable value, does not appear in the shaded area of Figure 10 a because the required conditions of perfect molecular alignment and alignment with the electric fi eld vector of the perfectly linearly polarized incident X-ray beam were not met in any pixel case. Integrating about θ to remove in-plane directional bias further reduced the possibility of perfectly aligning the molecular orien-tation with the electric fi eld vector of the incident X-ray beam.

An interesting property of angle-dependent NEXAFS spectra is that angular-dependence of the spectra due to the

© 2011 WILEY-VCH Verlag G© 2011 WILEY-VCH Verlag GAdv. Funct. Mater. 2011, 21, 1122–1131

polarization of the X-ray beam and due to the orientation and alignment of the sample molecules is the same. This can be seen in the “magic” angles of X-ray incidence (that depends on the degree of linear polarization; " = sin−1

√1 − 1/3P ) and

molecular tilt ( γ = 54.7 ° ), either of which produce spectra iden-tical to that of a completely amorphous sample, irrespective of other sample parameters. [ 22 ] This property can be exploited by extending the observed angular-dependence of the spectra in Figure 1 c to match the maximum intensity of the π ∗ resonance

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deduced from the STXM data and thus calculate the expected ideal NEXAFS spectrum corresponding to a perfect F8BT crystal oriented with the 1s – π ∗ transition dipole moment aligned parallel to the electric fi eld vector of the incident, per-fectly linearly polarized X-ray beam. Further, by preserving the cosine-squared angular dependence, a complementary NEXAFS spectrum, involving a perfect F8BT crystal oriented with the 1s – π ∗ transition dipole moment aligned perpendicular to the electric fi eld vector of the incident X-ray beam, can also be calcu-lated without assuming that the 1s – π ∗ transition dipole moment vanishes at perpendicular angles. Note, however, that perpendic-ularity to a vector in three dimensional space corresponds to a plane (i.e. two directions) and therefore the calculated “perpen-dicular” spectrum is not quite ideal. Thus, the “perpendicular” spectrum can be interpreted as arising from either an F8BT crystal with rotational symmetry, or an unpolarized incident X-ray beam. The calculated spectra are presented in Figure 10 b, along with the amorphous, or “magic” angle spectrum. Inter-estingly, the 285 eV π ∗ resonance goes to approximately zero, demonstrating that the 1s – π ∗ transition dipole moment does indeed vanish at perpendicular angles and further confi rming our choice of 3.4 as the maximum possible resonance intensity. While the 285 eV π ∗ resonance is observed to dip slightly below zero, which is not physically possible, this is likely a conse-quence of the errors in the measured spectra being magnifi ed and can be ignored. The parallel orientation (0 ° ) ideal spectrum in Figure 10 b displays some interesting additional structure. Specifi cally, a split peak emerges at ∼ 287 eV with further peaks at ∼ 289 eV – 291 eV that are otherwise dominated by broad σ ∗ structure. Further discussion of these ideal spectra, such as peak assignment, will be made in a later publication. However, we note here the potential of this approach for revealing addi-tional electronic structure without the requirement of perfectly ordered samples. This additional information may be of use for comparing to theoretical modeling of the unoccupied electronic structure of conjugated polymers.

3. Conclusions

We have used STXM to provide an in-depth, quantitative analysis of the molecular order and orientation present in the domain structure of polycrystalline F8BT fi lms. In addition to con-fi rming the qualitative information provided by cross-polarized optical microscopy, the STXM data has been used to derive maps of polymer backbone orientation, molecular alignment/order and out-of-plane tilt angle. The observed polycrystalline F8BT fi lms consist of well-ordered domains partly surrounded by ∼ 200 nm wide disordered domain boundaries and otherwise showing smooth transitions between domains. We have also imaged the interfacial domain structure of polycrystalline F8BT fi lms, demonstrating that the polycrystalline structure extends to the fi lm/air interface. The disordered domain boundaries are likely to result in local charge trapping and thus facilitate charge recombination in ambipolar operation, explaining the EL pat-terns seen in polycrystalline F8BT light-emitting fi eld-effect transistors. We fi nd a good agreement between the macroscopic tilt angle measured by TEY NEXAFS and the area-averaged tilt angle from STXM. The STXM data shows that there is a


relatively narrow distribution of tilt angles that correlates with fi lm structure, with the ordered domains in general showing a higher tilt angle than the disordered domain boundaries. Finally, we have combined knowledge of the F8BT spectrum and its angular dependence from both macroscopic and microscopic measurements to calculate ideal NEXAFS spectra corresponding to a perfect F8BT crystal oriented with the 1s – π ∗ transition dipole moment parallel and perpendicular to the electric fi eld vector of a perfectly linearly polarized X-ray beam.

4. Experimental Section Sample Preparation : F8BT was received from Cambridge Display

Technology Ltd. with a peak molecular weight of 163 kg mol − 1 and used as received. Films were prepared by dissolving F8BT in xylene (mixed isomers, anhydrous) at 12 mg mL − 1 and spin-coating at 1500 rpm to give fi lms ∼ 100 nm thick. Films for X-ray microscopy were prepared on glass substrates pre-coated with a water soluble sodium polystyrene sulfonate sacrifi cial layer (spin-coated from a 50 mg mL − 1 aqueous solution at 5000 rpm and baked at 125 ° C for 10 min) to aid fl oat-off of fi lms onto TEM grids. Films for TEY NEXAFS spectroscopy were prepared on highly doped silicon wafers. Annealing was performed in a nitrogen glove box at 150 ° C, 200 ° C, 240 ° C or 290 ° C for 20 min with quenching to room temperature.

Total Electron Yield (TEY) NEXAFS: TEY NEXAFS measurements were performed at beamline 10–1 of the Stanford Synchrotron Radiation Lightsource, California. Carbon K-edge NEXAFS spectra were obtained with a spectral resolution of about 50 meV at normal (90 ° ), so-called “magic” (50 ° ) and grazing (20 ° ) angles of X-ray incidence. The spectra were measured in total electron yield (TEY) mode by recording the drain current from the sample; the incident photon fl ux was simultaneously recorded via the drain current from a gold mesh grid located immediately upstream of the sample chamber. The NEXAFS spectra were normalized as detailed by Watts et al. [ 46 ]

X-ray Microscopy: NEXAFS spectromicroscopy was performed at beamline 5.3.2 at the Advanced Light Source, Berkeley, California [ 47 , 48 ] and at the PolLux beamline at the Swiss Light Source, Paul Scherrer Institut, Villigen, Switzerland. [ 49 , 50 ] The Advanced Light Source and the Swiss Light Source are open-access facilities, available to all researchers and free of charge for non-proprietary work (proprietary projects incur fees on a cost-recovery basis). [ 51 ] TEM grid-supported fi lms were mounted in the sample chamber which was either evacuated to low vacuum (SLS) or evacuated to low vacuum and subsequently refi lled with 1/3 atm of helium (ALS). The transmitted X-ray intensity through the fi lm was recorded using a scintillator and photo-multiplier tube and measured as a function of energy (280.0 to 320.0 eV with a resolution of 0.1 eV) and position (with X-ray focus better than 25 nm diameter and X–Y sample position resolution better than 1 nm). Transmitted X-ray intensity was converted to an X-ray optical density (defi ned as OD = ln (I 0 /I) ) by recording the X-ray intensity through a region of the TEM grid with no fi lm. For computation of domain orientation and molecular order a rotation stage was used to rotate the sample with respect to the fi xed linear polarization of the synchrotron. Three images were acquired of the same region of the sample at orientations of 0 ° , 45 ° and 90 ° with respect to the X-ray polarization. Detection of surface photoelectrons was achieved using a channeltron biased at 2.5 kV (Model No. KBL10RS, Dr. Sjuts Optotechnik GmbH) with the chamber pumped down to ∼ 10 − 6 mbar using a turbo pump and cold fi nger. A bias of + 400 V was applied to the metal order sorting aperture (OSA) to suppress photoemission of electrons from the OSA that otherwise swamps the signal from the sample. Due to the geometry of the microscope, the channeltron was mounted adjacent to the photomultiplier tube. The channeltron signal thus images the downstream-facing surface of the sample and is normalized to the incident photon fl ux by dividing by the measured transmitted photon counts. Further details of the surface imaging technique and analysis can be found in previous publications. [ 40 , 52 ]


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Supporting Information Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements The authors thank the ALS and SLS for beamtime and Cambridge Display Technology Ltd. for the supply of F8BT. This work was supported by the Engineering and Physical Sciences Research Council, U.K. (Advanced Research Fellowship EP/E051804/1). Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource, a national user facility operated by Stanford University on behalf of the U.S. Department of Energy, Offi ce of Basic Energy Sciences. The ALS is supported by the Director, Offi ce of Science, Offi ce of Basic Energy Sciences, of the U.S. Department of Energy under Contract DE-AC02–05CH11231. PolLux is funded by the BMBF (Project No. 05KS7WE1).

Received: September 13, 2010 Revised: November 29, 2010

Published online: February 15, 2011

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Mapping of Domain Orientation and Molecular Order in Polycrystalline Semiconducting Polymer Films...

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