Spatial Orientation and Order of Structure-Defining Subunits in Thin

Article pubs.acs.org/Macromolecules

Spatial Orientation and Order of Structure-Defining Subunits in Thin Films of a High Mobility n‑Type Copolymer Arthur Markus Anton,*,† Robert Steyrleuthner,‡,§ Wilhelm Kossack,† Dieter Neher,§ and Friedrich Kremer† †

Institut für Experimentelle Physik I, Universität Leipzig, Leipzig, Germany Institut für Physik und Astronomie, Universität Potsdam, Potsdam, Germany

§

S Supporting Information *

ABSTRACT: Orientation and order of distinct molecular subunits in solid layers of the high mobility n-type copolymer poly[N,N′bis(2-octyldodecyl)-1,4,5,8-naphthalenediimide-2,6-diyl]-5,5′-(2,2′bithiophene) P(NDI2OD-T2) are investigated by means of infrared transition moment orientational analysis. This novel spectroscopic technique based on concurrent absorbance measurements of structure-specific bands in dependence on inclination and polarization of the incoming light enables to determine the complete tensor of absorption independently for each transition moment. As a result, for nanometer thin films pronounced in-plane anisotropy arising from self-aggregated order is detected, which, however, is no longer discernible for micrometer thick samples. In contrast, the out-of-plane orientation (inclination of molecular subunits) is retained irrespective of the widely varying layer thicknesses (150 nm vs 1.4 μm). Thus, the conception of the sample morphology occurs as stratification of slightly misaligned layers of oriented polymers; with increasing film thickness the macroscopic in-plane order diminishes, whereas the out-of-plane orientation is preserved.

■

domains with respect to the substrate and π−π stacking proceeding along the out-of-plane direction.18 Studies using near-edge X-ray absorption fine structure (NEXAFS) spectroscopy19,20 or IR spectroscopy21−23 showed (while rotational symmetry with respect to the substrate normal was assumed or measured) that the molecular planes defined through the naphthalenediimide (NDI) and the bithiophene (T2) subunits are inclined relative to the substrate. Although the in-plane order of the molecular structure has been studied by a variety of methods, quantitative analyses of the polymer orientation are seldom.13,15,24 For example, a combination of high-resolution and scanning transmission electron microscopy revealed a polymer backbone long-range correlation approaching length scales of a micrometer,25 whereas a microscopy study with polarized light found that fibrils formed by the polymer can keep parallel to each other over regions as wide as 1−2 cm2.14 By means of scanning transmission X-ray microscopy (STXM)26 or polarized charge modulation microscopy (p-CMM),27 it was possible to extract the local degree of order or degree of alignment, respectively. In order to reveal the molecular organization in thin polymer films, we employ the method of infrared transition moment orientational analysis (IR-TMOA), which allows for determining the complete molecular absorbance tensor.23,28−31 A film

INTRODUCTION Great attention is paid to organic semiconducting materials due to their possible applicability as organic field effect transistors (OFETs), solar cells (OSCs), light-emitting diodes, and biosensors, for instance.1−5 On the basis of the physical properties of soft matter, those polymers are convenient for devices that resist demands on elasticity and flexibility as well as easy customization in shape.6 In addition, they are suitable for low-cost deposition methods on large scale as printing and coating.7−10 Especially poly[N,N′-bis(2-octyldodecyl)-1,4,5,8naphthalenediimide-2,6-diyl]-alt-5,5′-(2,2′-bithiophene) P(NDI2OD-T2), a conjugated donor−acceptor copolymer, has been successfully established as electron conductor in n-type OFETs and electron acceptor in OSCs realizing power conversion efficiencies up to 5.7%.11,12 It is believed that charges can be transported efficiently along the polymer backbone and the π−π stacks, whereas the long nonconjugated side chains represent hopping barriers.13 Thus, varying the microstructure affects the electron mobility in bulk over 2 orders of magnitude.14−16 In order to be able to tailor the macroscopic charge transport on the basis of the molecular organization following the structure−property paradigm,6 the understanding of the structure-forming processes responsible for the final orientation and order is essential.16,17 Early work on thin films of P(NDI2OD-T2) made by dissolving the polymer in chlorobenzene (CB) revealed a face on orientation of crystalline © XXXX American Chemical Society

Received: November 6, 2015 Revised: January 20, 2016

A

DOI: 10.1021/acs.macromol.5b02420 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

Figure 1. (a) Spectra of the P(NDI2OD-T2) sample films at normal incidence (ϑ = 0°). (b, c) None of the structure-specific bands of sample 1 exhibit a polar (φ)-dependent absorbance at ϑ = 0°. The enhanced absorbance of the b band for s-polarized (φ = 0°) and that of the c band for ppolarized light (φ = 90°) indicates a preferred orientation of the according TMs parallel and perpendicular to the substrate, respectively. (d) In the case of sample 2 the polar-dependent absorbance at ϑ = 0° demonstrates anisotropically distributed TMs with respect to the substrate. The diminishing difference between the absorption of TM b for s- and p-polarized light with increasing tilt arises from the superposition of the stronger absorption for p-polarized light at ϑ = 0° and the increasing absorption for s-polarized light with rising tilt similar to sample 1. (e) A monomer segment of P(NDI2OD-T2) with indicated TMs (a, b, and c) employed for the structural elucidation as well as the molecular planes (NDI red; T2 green). For sake of clarity the aliphatic side chains are omitted; the chemical structure can be found in the inset. All spectra are shifted vertically; the data of sample 2 in panel a are multiplied by 3 in addition.

more complex geometry of sample 2. The deduced values from all samples represent global properties (average over detection area), which do not suffer from local attributes and, hence, provide additional and complementary information to those derived from scattering experiments.

down to the measurable thickness of 150 nm is prepared (sample 2; cf. Experimental Section and ref 23), which results in a macroscopic in-plane anisotropy (within the 10 mm in diameter detection area), even though the polymer is spincoated from solution. Thus, we are able to characterize quantitatively the inclination of the molecular subunits and the alignment order of the polymer backbone simultaneously. For reference reasons a 1.4 μm thick film is prepared as well (sample 1), for which it is known to show rotational symmetry with respect to the substrate normal.23 For the sake of convenience, we will first discuss sample 1 and proceed with the

■

RESULTS In the IR spectra (Figure 1a) peaks are identified at ν = 1707 cm−1 (a band) and ν = 1667 cm−1 (b band), corresponding respectively to the symmetric (ν s) and antisymmetric CO stretching vibration (νas) located at the B


Article

Macromolecules NDI unit (Figure 1e).21 In addition, we assess the C−H out-ofplane bending vibration (δ) at the T2 unit (c band; ν = 792 cm−1).21 Since each IR-active vibration is associated with a specific transition moment (TM) that is localized at one of those two structural units (either NDI or T2) and exhibits a distinctive orientation with respect to the molecule (Figure 1e), we assign normalized vectors to them: vector a refers to the TM of νs(CO), b to that of νas(CO), and c to that of δ(C−H). These vectors are described in the sample fixed reference frame through spherical coordinates: an azimuthal angle Θ and a polar angle Φ.23 Because of the specificity of the IR spectral range, we determine the orientation of the NDI unit as well as that of the T2 part independently of each other as follows: (1) The NDI plane’s normal vector n is given by the vector cross product of a and b (n = a × b), which are both located within the NDI plane. (2) The normal vector of the T2 plane is represented through c. Since every individual thiophene ring exhibits one TM of δ(C−H) that runs perpendicular to the ring, c represents the average over both rings from the T2 unit. The distribution of orientation of the described TMs (and the corresponding molecular moieties) is deduced by IRTMOA.23,28−31 For that a sample film is inclined (at an angle ϑ = −60° to +60°; Figure 2) relative to the optical axis, while

order parameter Sii along the direction of consideration (here on the example of the principal axis i) reads as Sii =

⎞ 1 ⎛ 3 μ̲ ′ ii ⎜⎜ − 1⎟⎟ 2 ⎝ ∑m μ̲ ′ mm ⎠

(1)

where μ′ is a diagonal matrix with the absorbance eigenvalues on the main diagonal μ′ii and m ∈ [i, j, k]. Sjj and Skk are defined analogously. It can be shown that Sii = 0 in the case that all TMs are randomly distributed, Sii = 1 for perfectly aligned TMs in parallel to i, and Sii = −1/2 when all TMs are oriented perpendicular to i. During the curse of determining the orientation of the molecular subunits the absorbance eigenvalues are employed. Afterward, the results are compared with the literature, for which it is easier to use the molecular order parameter. Molecular Orientation of Sample 1. The spectrum of sample 1 (Figure 1a−c) reveals no anisotropy at normal incidence (ϑ = 0°), proving equally distributed TMs in all directions parallel to the film plane. With rising tilt the b band absorbs stronger for s-polarized (Figures 1b and 3a) and the c band for p-polarized light (Figures 1c and 3a). This indicates a preferred alignment of b parallel and of c perpendicular to the substrate (xy) plane.23 Because the absorbance of all three bands (a, b, and c) of sample 1 is symmetric with respect to polarization and inclination (Figure 3a), the principal axes (i, j, k) coincide with the sample coordinate system (x, y, z). In detail, we find for all three bands (a, b, and c) the Euler angle β < 3°. That means the principal axis system is tilted less than 3° away from the sample-fixed coordinate system, and hence, we ascribe the principal axis k to the sample fixed axis z representing the direction perpendicular to the film plane. Furthermore, the absorbance eigenvalues of the remaining two directions are equal (Ai ≈ Aj), indicating no preferred alignment in any direction parallel to the substrate. Thus, the Euler angle α, which has the meaning of a polar angle, is arbitrary and we identify the principal axes system (i, j, k) with the sample coordinate system (x, y, z). In summary, we find uniaxial order with respect to the sample normal (z axis; Figure 4a,b). In order to deduce the molecular structure from the absorbance pattern, we consider laterally extended planar aggregates composed of extended polymer chains as proposed by Rivnay et al.18 and discussed recently.23 The azimuthal orientation (Θ) of the particular TMs can then be determined from the eigenvalues Ai, Aj, and Ak using eq S5 (cf. Supporting Information).23 We find that TM b runs parallel to the xy plane (Θb = 90°; Table 1), whereas a is inclined by 39° relative to it (Θa = 51°). We set the orientation of TM b as parallel to the x axis (Φb = 0°) and consider a as perpendicular to b (Φa = 90°) in accordance with the literature21 and later on experimentally confirmed through the findings for sample 2. Consequently, the NDI plane is tilted 39° relative to the substrate as well. The T2 plane, instead, shows an angle of 34° relative to the substrate (Θc = 34°, Figure 4b). The planes’ cutting angle χ is deduced on the condition that the polymer backbone is laterally extended and the inclination of the molecular planes arises from a rotation around the direction of TM b. Thus, χ is solely based on experimental data. Note that the cutting angle can be different from the dihedral angle often deduced from simulations. While the former is defined through the alignment of the NDI unit relative to the T2 plane (which is the mean of the two thiophene ring planes

Figure 2. Schematic of the measurement geometry for IR-TMOA. The incoming IR beam (k0) is propagating along the z′ direction and its electric field E is polarized (φ) within the x′y′ plane. During the sample-fixed coordinate system (x,y,z) is inclined at an angle ϑ relative to the laboratory frame (x′,y′,z′), the polarization-dependent absorbance is recorded. Thus, the projection of the transition moment (TM) onto the x′y′ plane is systematically varied, and information on the tensor of absorption with respect to all three axes (x,y,z) can be obtained.

the electric field of the collimated beam is polarized (φ = 0° to 180°) prior to passing through the specimen. Thus, the projections of the TMs onto the plane of polarization are altered stepwise. This modification is measured by the dependence of the integrated absorbance (area under the curve) of the vibrational bands on ϑ and φ. On the basis of this dependence a structural model is fitted to the experimental absorbance pattern (Figure 3) and, as a result, a diagonal matrix with the eigenvalues of the absorption tensor (Ai, Aj, Ak; Tables 1 and 2) as well as the relative orientation of the principal axes (i, j, k) to the sample fixed coordinate system (x, y, z) described by three Euler angles (α, β, γ) in ZXZ convention (Tables S1 and S2 in the Supporting Information) is obtained.23,28 Subsequently, the set of eigenvalues (Ai, Aj, Ak) can optionally be converted into a set of order parameters (Sii, Sjj, Skk), which is more common than the characterization using the eigenvalues but does not provide additional information. The C


Article

Macromolecules

Figure 3. Integrated absorbance (area under the curve) of the structure-related bands of (a) sample 1 and (b) sample 2 along with the fitted structural model. (a) When sample 1 is inclined (ϑ ≠ 0°), the absorption of νas(CO) is increased for s-polarized light (φ = 0° and 180°), whereas the absorption of δ(C−H) is enhanced for p-polarized light (φ = 90°), which indicates the orientation of TM b as parallel and that of TM c as perpendicular to the substrate. (b) The pronounced polarization-dependent absorption at normal incidence (ϑ = 0°, compare with panel a) demonstrates anisotropically orientated TMs along one direction parallel to the substrate. Because of the reduced film thickness, IR-TMOA could not be performed properly for δ(C−O) (TM c) of sample 2. For sake of clarity and comparability all values are normalized to the mean absorbance at normal incidence.

Molecular Orientation of Sample 2. In the case of sample 2 a pronounced in plane anisotropy of the absorption is evident for the a band as well as the b band (Figures 1d and 3b). In particular, the b band shows a strong absorption (Ai = 1.29 ± 0.04, Table 2) along one direction of the xy plane (principal axis i) and a weaker absorption (Aj = 0.55 ± 0.01) along the direction perpendicular to the former (principal axis j) but still part of the xy plane. Normal to the sample plane (principal axis k, which coincides with z) essentially no absorption is found (Ak = 0.00 ± 0.06). Consequently, vector b is distributed only within the xy plane (Θb = 90°; Table 2) and we consider i as the direction of the preferred orientation (Φb = 0°) and absorption along j as arising through deviations from the alignment. In order to characterize the in-plane anisotropy, we enhance the density distribution function from sample 1 and model the molecular distribution assuming the polymer backbones as dispersed obeying a Gaussian function centered at i (Φb = 0°). Its distribution width ωb then is determined through the aspect ratio of the particular absorbance values Ai/Aj = 2.35 and results in ωb = 38.6° (eqs S1, S2, and S6 of the Supporting Information, Figure 4c).

in the T2 unit), the latter is based on the orientation of atomic bonds between the NDI part and the first adjacent thiophene ring. Torsion between the two rings in the T2 unit gives rise to a deviation of the normal vector of the adjacent ring from that of the T2 plane. Consequently, the cutting angle and the dihedral angle differ by half of the torsion angle. Two cases of χ have to be considered (Figure 4b):23 (a) Both planes, the NDI and the T2 one, rotate in opposite directions around the polymer backbone (considered as parallel to b). Then χ is given by the sum of the individual inclination angles: χ = 73° (39° + 34°). (b) The two planes are tilted in the same direction around the backbone. Consequently, the cutting angle is given by the difference of the two inclinations: χ = 5° (39°−34°). These values of χ are deduced under the presumption that both planes rotate around a better: common axis parallel to TM b (which is here identical to i and x), which is appropriate for the rigid planar structure of the aromatic subunits. We consider a deviation of the relative polar angle of the planes’ normal vectors (Φn − Φc = 180° for case a and 0° for case b) of more than 30° as not reasonable. In such extreme cases χ would differ by −3° (case a) and +13° (case b) from the values derived above (Figure S2). D


Article

Macromolecules

Table 1. Relevant Parameters of the Structural Model Resulting from the Experimental Absorbance Pattern of Sample 1a quantity

a

b

Ai Aj Ak Sii Sjj Skk Θ (deg) Φ (deg)

2.40 ± 0.03 2.50 ± 0.04 3.18 ± 0.24 −0.05 ± 0.01 −0.04 ± 0.01 −0.09 ± 0.05 51 90

8.77 ± 0.09 8.92 ± 0.09 0.00 ± 0.07 0.24 ± 0.01 0.26 ± 0.01 −0.50 ± 0.04 90 0

n(NDI)

χ

c(T2) 0.59 ± 0.06 0.59 ± 0.06 2.59 ± 0.04 −0.27 ± 0.01 −0.27 ± 0.01 0.53 ± 0.02 34

39 −90

73/5

a

The values Ai, Aj, and Ak denote the absorbance eigenvalues according to the principal axes (i, j, k) and Sii, Sjj, and Skk the corresponding order parameters. Θ gives the azimuthal angle (eq S5) and Φ the polar angle of the TM, respective the plane’s normal vector, whereas χ represents the planes’ cutting angle as sum/difference of the inclinations in accord to case a or case b. The angles are accurate up to ±5°. The full parameter set including the Euler angles (α, β, γ) and the refractive index of the sample material n are given in Table S1.

Table 2. Relevant Parameters of the Structural Model Resulting from the Experimental Absorbance Pattern of Sample 2a quantity

a

b

Ai Aj Ak Sii Sjj Skk Θ (deg) Φ (deg)

0.17 ± 0.01 0.26 ± 0.01 0.17 ± 0.04 −0.07 ± 0.03 0.14 ± 0.04 −0.07 ± 0.11 47 90

1.29 ± 0.04 0.55 ± 0.01 0.00 ± 0.06 0.55 ± 0.02 −0.05 ± 0.01 −0.50 ± 0.07 90 0

n(NDI)

χ

c(T2) sample 1 (Table 1)

sample 1 (Table 1) 43 −90

34 77/9

a

The values Ai, Aj, and Ak denote the absorbance eigenvalues according to the principal axes (i, j, k) and Sii, Sjj, and Skk the corresponding order parameters. Θ gives the azimuthal angle (eqs S6 and S7) and Φ the polar angle of the TM, respective the plane’s normal vector, whereas χ represents the planes’ cutting angle as sum/difference of the inclinations in accord to case a or case b. The angles are accurate up to ±5°. The full parameter set including the Euler angles (α, β, γ) and the refractive index of the sample material n are given in Table S2.

For TM a we find that the principal axes coincide (within experimental accuracy) with the principal axes of b with the strongest absorption (Aj = 0.26 ± 0.01; Table 2) along the direction j being perpendicular to TM b. In contrast to TM b, the eigenvalue Ak does not vanish (Ak = 0.17 ± 0.04) which, in combination with the alignment of j parallel to the xy plane, explicitly proves that TM a is symmetrically inclined with respect to the xy plane. Thus, the molecular organization of TM a is modeled assuming that the mean vector is preferentially aligned at a particular inclination (Θa) and runs perpendicular to TM b (Φa = 90°).21 The distribution function is designed to account for the inclination in the direction of j (Φa = 90°) but to preserve the alignment parallel to the xy plane for the perpendicular course (direction of i, Φ = 0°). Agreement between the experimentally determined eigenvalues (respectively their ratios Ai/Aj, Ai/Ak, and Aj/Ak) and the corresponding measures calculated on the bases of the model structure (eqs S1−S3 and S7) is found for an inclination of 43° relative to the substrate (Θa = 47°) and a distribution width of ωa = 39.8°. These values corroborate our findings for TM b of sample 2 (ωa ≈ ωb) and are in agreement with the literature (Θ = 45°−52°).20−23 Furthermore, they indicate only a slightly smaller azimuthal angle (Θa = 47°) compared to the thick film (sample 1: Θa = 51°). This increase is presumably related to a higher disorder in sample 1 (a perfect uniaxial distribution with an azimuthal angle of 54.7°, the “magic angle”, cannot be distinguished from an isotropic sample). Consequently, the NDI unit shows an inclination of 43° with respect to the substrate (Table 2). A distribution function taking the case into account in which the inclination of TM a is preserved for all polar angles (Θ0 = Θa in eq S7, rotation around z axis) has been tested as

well. In that case the derived values (ωa = 52.3°, Θa = 57.6°) appear to be unreasonable in comparison to the results for TM b of sample 2 (ωb = 38.6°) and TM a of sample 1 (Θa = 51°). Note that the dichroic ratios are caused by a deviation of TM a from its mean alignment, which originates from a rotation around an axis perpendicular to the inclined plane. The same holds for TM b with the difference that the axis of rotation is distinct to that of TM a. As a result, the particular TMs of type a are not entirely inclined at Θa; they are rather distributed around the mean inclination allowing for different rotation axes. This is consistent with the derived values and the measurement uncertainty. In case of TM c the small film thickness of sample 2 and the accompanying less absorbance of the c band do not allow for determining the orientation of this vector properly. On the basis of a comparable inclination of the TMs a and b between samples 1 and 2 (Tables 1 and 2), we therefore assume the inclination of c of sample 2 as similar to the corresponding part of sample 1. The planes’ cutting angle χ then results either in 77° (case a) or in 9° (case b). A 30° deviation in the relative polar angles of the planes’ normal vectors would give rise to a difference of −3° (case a) and +11° (case b) (Figure S2).

■

DISCUSSION Inclination of the Molecular Subunits. It has been reported that on aluminum or barium fluoride substrates TM b exhibits a distinctive inclination to that,22,23 whereas a strict alignment is found for zinc selenide (ZnSe). This implicates that the structural organization is affected by the employed substrate.10 When comparing the inclination of the TMs and molecular subunits with identically prepared (except for the substrate) samples,23 it appears that TMs a and c orient at E


Article

Macromolecules

Figure 4. Orientation of the transition moments as deduced from IR-TMOA. (a) As evident by the symmetric absorption TM a (red) is equally distributed at the lateral surface of a cone (opening angle Θa = 51°, Table 1); TM b (black) is found within the xy plane (Θb = 90°). Their vector cross product gives the NDI plane’s normal vector n (red; Θn = 39°), which is consequently rotationally symmetric distributed. The same holds true for TM c (T2 unit, green). (b) In case both planes are tilted in opposite directions around the polymer backbone ((assumed as parallel to TM b which is parallel to i) the cutting angle χ results from the sum of the particular inclinations (73°; case a). For the opposite case (rotation in the same direction) χ is given by the difference (5°; case b). (c) At a reduced sample thickness macroscopic anisotropy in absorption arises. TM b is distributed only within the xy plane (Θb = 90°, ωb = 38.6°, Table 2). TM a is symmetrically tilted with respect to the xy plane (Θa = 47°, ωa = 39.8°, only one direction is shown) which gives rise to absorption perpendicular to the substrate, despite the eigenvector j runs parallel to it. For convenience we set i as parallel to x and j parallel to y. For sake of clarity only selected examples are shown; note that TMs a and c may be tilted in both directions (upward or downward).

similar angles (TM a: 51°/47° vs 51°; TM c: 34° vs 29°) within the experimental uncertainty (±5°). However, in the case of TM b (90° vs 71°) the difference has to be taken into account, but the reason why only that TM is selectively influenced is not clear so far. The inclination of the NDI plane is reduced (39° vs 45°), whereas that of the T2 part is increased (34° vs 29°). Thus, the cutting angle for case 1 (χ = 73°) is similar to previous values (χ = 74°),23 for case 2, instead, χ differs (5° vs 16°). Molecular Order. In order to compare our findings concerning the in-plane anisotropy with previous results derived from STXM and p-CMM measurements,26,27,32 we express the degree of order (as discussed in refs 26 and 27) and the molecular order parameter in terms of the ratio of absorption values arising from crystalline and amorphous parts of the sample (M/C; cf. Supporting Information). Sciascia et al. derived an averaged degree of order of O(3) = 17% while assuming a possible orientation of the absorbing moieties in all three directions of space.26 For that value one finds a ratio M/C = 0.61. Despite the TMs for X-ray and optical transitions are extended over the distinct subunits, whereas in the case of IR the TMs are well located at one of these two moieties, it is possible to obtain comparable information. Since the possible orientations of the aromatic structures forming the polymer backbone is limited by steric hindering and two NDI subunits

are separated by only one T2 part, the in-plane order of the T2 subunits is expected to be similar to that of the NDI counterparts, which are a priori labeled with TM b. Thus, the in-plane order of TM b reflects the in-plane order of the polymer backbone. The ratio M/C = 0.61 gives then an in= 0.43. This value is in fair plane order parameter of S(2) i agreement with the experimental result of Si = 0.55 (Table 2). Martino et al. obtained a degree of alignment of O(2) = 11% and O(2) = 22% from the optical density analysis and p-CMM measurements, respectively, while considering arrangements only parallel to the substrate.27 For those values a ratio of M/C = 0.25 and 0.56 can be found which corresponds to an = 0.39 and S(2) = 0.42. order parameter of S(2) i i In addition, the separation between absorption arising from crystalline and amorphous parts allows for estimating the degree of aggregation (considering aggregates are synonymous with crystallites and all of them oriented in the same direction). For the in-plane order parameter of Si = 0.55 (Table 2) one finds a degree of aggregation of 35% (eq S12) and 44% (eq S14) assuming an alignment of the aggregates in three and in two directions of space, respectively. Similar values have been found recently for UV/vis measurements.15 Structural Organization. On the basis of the similar inclination of the TMs a and b, and hence of the NDI plane for the nanometer and micrometer thick films, we inquire for a F


Article

Macromolecules

Figure 5. Schematic of the revised notion and interpretation of the molecular orientation and order. (a) The extended polymer chains with inclined NDI and T2 units aggregate resulting in (b) lamellar-like structures. Because of the pronounced correlation length parallel to the substrate, these lamellae are oriented and (c) build up flat plate-like sheets. Those oriented polymer assemblies then induce the macroscopic anisotropy evident in the case of sample 2. Since the coherence length is orders of magnitude smaller than the film thickness, the absorption, as average over oriented (Θ) but slightly misaligned (Φ) layers, appears rotationally symmetric as for sample 1. Note that the stacking can occur in a segregating or an alternating fashion as indicated in part b.

coupling of the in-plane orientation of the aggregates is quantified by the coherence length in thickness direction (6.5 nm),15 which is much smaller than in the lateral direction (34 nm).15 Thus, the order parallel to the substrate is maintained, resulting in a macroscopic anisotropy as found for sample 2 (Figure 4c and Figure S1d). Growing film thickness then gives rise to a structural superposition of several layers of aggregates, which are slightly misaligned with respect to the ones underneath, and hence, cause the diminishing of the macroscopically observable in-plane anisotropy (Figure 4a and Figure S1f). Please note that the two forms of π−π stacking, either segregated (donor on donor and acceptor on acceptor) or alternating (donor on acceptor and vice versa) as depicted in Figure 5b, have been shown to be possible by Brinkmann et al.,42 but only the former was observed in spin-coated films from a chlorobenzene solution similar to those investigated here.15 However, the particular stacking is not in the focus of this work. Thus, both forms are considered as possible. The preparation procedure can be excluded as possible origin of the apparent anisotropy because spin-coating, on the one hand, would result in a radial structure, whereas, on the other hand, the movement of a drying front (from the edge of the circular-shaped IR window into the middle or vice versa) causes concentric rings. Both of them would result in a rotationally symmetric distribution of the TMs within the macroscopic detection area (10 mm in diameter) at the middle of the sample, which is evidently not the case at least for sample 2. Moreover, the notion of extended aggregates comprising lamellae is further supported by considering P(NDI2OD-T2) as a block copolymer composed of a rigid (NDI-T2) core surrounded by two flexible blocks of aliphatic (OD) side chains. Then, the aggregation behavior depends in general on the following points: (i) The relative volume fraction of the respective blocks (OD:NDI-T2:OD) results from the molecular weights of the different parts and the corresponding densities. For the NDI-T2 core one finds 376.4 g mol−1 and we employ 1.27 g cm−3 adopted from pyrene;43 for the OD side chains one finds 281.5 g mol−1 and we use 0.79 g cm−3 adopted

mechanism that allows for maintaining the inclination (Θ) of the molecular moieties even when the film thickness is drastically increased, whereas the polar dependence (Φ) gets lost. This condition is conceivably satisfied by strongly oblate assemblies, which align parallel to the substrate during spin coating and resist disarrangements (stay preferentially parallel) because of their shape. In case there were not such assemblies, the molecules would orient randomly with rising film thickness resulting in isotropic absorbance patterns for all TMs and molecular order parameters of 0. Such oblate structures are in agreement with previously proposed planar aggregates comprising extended polymer chains18 and preaggregates already present in the polymer solution.33 Furthermore, from the spectral anisotropy of the a and b band for sample 2 we conclude that the polymer backbones incorporated into the planar aggregates, and hence the assemblies itself, have a preferential orientation within the substrate plane, which is disguised in the case of the thicker sample 1. Thus, we modify the concept of the structural organization on the basis of the model proposed by Rivnay et al.18 and integrate recent findings from the literature as follows (Figure 5): Extended polymer chains with the backbone oriented preferentially parallel to the substrate18 (Θb = 90° from the present work; Θb = 71° from ref 23) form lamellae, as shown by Steyrleuthner et al., Schuettfort et al., and Lemaur et al.15,19,34 for instance. Inside such lamellae the molecular planes of the NDI and T2 subunits are inclined (resolved by Schuettfort et al., Giussani et al., and Anton et al.20−23 as well as evident in the present work). Additional indication of molecular plains tilted with respect to the stacking direction is provided by the increased π−π stacking distance (3.93 Å)18 as compared to thiophene-based (3.3−3.5 Å) 35−39 or naphthalenediimide-based derivatives (3.44−3.55 Å).40,41 Furthermore, these lamellae are primarily aligned (apparent in the micrographs in ref 15 and the anisotropy in absorption for sample 2 in the present work) and build up flat plate-like aggregates. The alignment appears only in thin films, where the G


Article

Macromolecules

only as well as for thin film interference arising from the ZnSe substrate by Fourier-filtering and a straight baseline was subtracted. In case of infinitely thin films the integrated absorbance is proportional to the molecular ensemble average (denoted as ⟨...⟩) over the scalar product (μ·E)2

from nonadecane.43 Thus, the volume fraction amounts about 35:29:35. (ii) The degree of polymerization is considered as 20, because one of the two the aliphatic side chains is represented through the sum formula C20H41 and one monomeric unit through CH2. Thus, one OD part is composed of 20 monomeric units. (iii) The Flory−Huggins parameter depends on the interaction energies between segments from different blocks and from the same.44 In case those energies are equal, the parameter would result in 0. However, since the OD side chains are hydrophobic but the aldehyde groups on the NDI part are partially charged, the blocks are hardly miscible. When employing the solubility parameter for octane45 (or decane) and bromonaphthalene45 an assumption for the Flory−Huggins parameter can be made resulting in ≈2.45 Taking remarks i to iii into account, one obtains a set of parameters which corresponds to self-aggregated lamellae.44,46 In conjunction with the interdigitated structure of the side chains,34,47 the observed long-range order appears natural, even though the polymer is spin-coated on the substrate.48,49 Consequently, by tuning the chemical structure of the sample molecule and choice of the adequate supporting material, one may take advantage of the self-assembly propensity and obtain functionalized films directly by spin-coating without the need of extensive pre- or post-treatment.50,51 Detailed studies facing the (mathematical) description of the orientational distribution of the molecular moieties in dependence on the film thickness as well as the preparation conditions including the employed solvents and the substrate material are intended for future prospects.

A ∝ ⟨(μ · E)2 ⟩

Thus, the current polarization and inclination directly influence the amount of absorbed light. Furthermore, eq 2a can be transformed into

A ∝ ET· μ̲ ·E

(2b)

where μ represents a matrix defined through μab = ⟨μaμb⟩ with a, b ∈ [x, y, z]. Because μ is a symmetric matrix in the reference frame of the sample coordinate system, it can be diagonalized resulting in a new matrix μ′ with Ai, Aj, and Ak (the eigenvalues of μ) on the main diagonal. This diagonalization represents a transformation from the sample coordinate (x, y, z) into the principal axes system (i, j, k) of the absorption tensor described by three Euler angles α, β, and γ in ZXZ convention. The reader should be aware of the fact that the values obtained by IR-TMOA are approximations. The exact solution can only be determined by solving the equation of the propagation of the electromagnetic wave in an anisotropic medium.52,53 However, the orientation determined by IR-TMOA deviates not more than 3° from the exact solution and is covered by the measurement accuracy of ±5°.28 Further details concerning the sample preparation and the measurement procedure can be found in ref 23.

■

ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b02420. The complete sets of the parameters of the absorbance tensor (Euler angles and sample material’s refractive index) for samples 1 and 2, polarized micrographs of the sample films showing birefringent properties for sample 1 as a result of long-range order, a detailed discussion of the mathematical description of the molecular distribution functions, and the connection between the molecular order parameter and the degree of order or alignment are provided. (PDF)

■

CONCLUSION In summary, we provide quantitative refinements of the molecular organization of the high mobility n-type copolymer P(NDI2OD-T2) as proposed by Rivnay et al.18 The films consist of platelets of oriented lamellae, which comprise inclined molecular subunits. Their tilt (azimuthal orientation) is hardly affected by the film thickness (Θa = 47° vs 51°; Θb = 0°). The in-plane orientation of the platelets and hence of the subunits is biased, indicating self-aggregated long-range order, even though the samples are spin-coated from solution. Through stratification of slightly misaligned layers the macroscopic order is gradually lost with increasing film thickness, resulting in rotational symmetry. Thus, we go one step further in understanding the assembly process on the molecular scale as a prerequisite for tailoring devices’ functionality.

■

(2a)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (A.M.M.). Present Address ‡

Fachbereich Physik, Freie Universität Berlin, Berlin, Germany.

Notes

EXPERIMENTAL SECTION

The authors declare no competing financial interest.

■

Thin films of P(NDI2OD-T2) (ActiveInk N2200, Polyera Corp., USA) are prepared on an IR-transparent zinc selenide (ZnSe) substrate (Crystal GmbH, Germany) by spin-coating from a chlorobenzene solution. Through the use of two different concentrations, 35 g/L (sample 1) or 10 g/L (sample 2), we obtain distinctive samples in thickness: 1400 ± 100 nm (sample 1) or 150 ± 20 nm (sample 2). Subsequently, the samples are dried and annealed at 200 °C under a nitrogen atmosphere well below the melting temperature (300−320 °C).24 Infrared measurements are accomplished on a Bio-Rad FTS 6000 FTIR spectrometer equipped with a liquid nitrogen-cooled mercury cadmium telluride (MCT) detector (Kolmar Technologies Inc., USA) using a spectral resolution of 2 cm−1. The integrated absorbance A is obtained through fitting a pseudo-Voigt function to each peak after the spectra were corrected for atmospheric water by subtracting absorption with comparable strength arising from gaseous water

ACKNOWLEDGMENTS Thanks is given to Z. Chen and Prof. A. Facchetti (Polyera Corp., USA) for supplying the sample polymer as well as Prof. M. Beiner (Martin-Luther Universität Halle Wittenberg and Fraunhofer-Institut für Mikrostruktur von Werkstoffen und Systemen IMWS, Halle (Saale)) and W. K. Kipnusu (University of Leipzig) for fruitful discussions. Financial support by the Deutsche Forschungsgemeinschaft (DFG, Projects B05 and B08 within SFB/TRR 102 “Polymers under multiple constraints: restricted and controlled molecular order and mobility”), through the Leipzig School of Natural Sciences “Building with Molecules and Nano-Objects” (BuildMoNa), by the Sächsische Forschergruppe FOR 877 “From Local H


Article

Macromolecules

(30) Kipnusu, W. K.; Kossack, W.; Iacob, C.; Jasiurkowska, M.; Sangoro, J. R.; Kremer, F. Z. Phys. Chem. 2012, 226, 797−805. (31) Mahmood, N.; Anton, A. M.; Gupta, G.; Babur, T.; Knoll, K.; Thurn-Albrecht, T.; Kremer, F.; Beiner, M.; Weidisch, R. Polymer 2014, 55, 3782−3791. (32) Watts, B.; Schuettfort, T.; McNeill, C. R. Adv. Funct. Mater. 2011, 21, 1122−1131. (33) Steyrleuthner, R.; Schubert, M.; Howard, I.; Klaumünzer, B.; Schilling, K.; Chen, Z.; Saalfrank, P.; Laquai, F.; Facchetti, A.; Neher, D. J. Am. Chem. Soc. 2012, 134, 18303−18317. (34) Lemaur, V.; Muccioli, L.; Zannoni, C.; Beljonne, D.; Lazzaroni, R.; Cornil, J.; Olivier, Y. Macromolecules 2013, 46, 8171−8178. (35) Janzen, D. E.; Burand, M. W.; Ewbank, P. C.; Pappenfus, T. M.; Higuchi, H.; da Silva Filho, D. A.; Young, V. G.; Brédas, J.-L.; Mann, K. R. J. Am. Chem. Soc. 2004, 126, 15295−15308. (36) Scherlis, D. A.; Marzari, N. J. Am. Chem. Soc. 2005, 127, 3207− 3212. (37) Graf, D. D.; Duan, R. G.; Campbell, J. P.; Miller, L. L.; Mann, K. R. J. Am. Chem. Soc. 1997, 119, 5888−5899. (38) Facchetti, A.; Yoon, M.-H.; Stern, C. L.; Katz, H. E.; Marks, T. J. Angew. Chem., Int. Ed. 2003, 42, 3900−3903. (39) Dag, S.; Wang, L.-W. J. Phys. Chem. B 2010, 114, 5997−6000. (40) Zhan, X.; Facchetti, A.; Barlow, S.; Marks, T. J.; Ratner, M. A.; Wasielewski, M. R.; Marder, S. R. Adv. Mater. 2011, 23, 268−284. (41) Katz, H. E.; Siegrist, T.; Schön, J. H.; Kloc, C.; Batlogg, B.; Lovinger, A. J.; Johnson, J. ChemPhysChem 2001, 2, 167−172. (42) Brinkmann, M.; Gonthier, E.; Bogen, S.; Tremel, K.; Ludwigs, S.; Hufnagel, M.; Sommer, M. ACS Nano 2012, 6, 10319−10326. (43) CRC Handbook of Chemistry and Physics, 90th ed.; Lide, D. R., Ed.; CRC Press (Taylor and Francis Group): Boca Raton, FL, 2009. (44) Lee, M.; Cho, B.-K.; Zin, W.-C. Chem. Rev. 2001, 101, 3869− 3892. (45) Koenhen, D. M.; Smolders, C. A. J. Appl. Polym. Sci. 1975, 19, 1163−1179. (46) Bates, F. S. Science 1991, 251, 898−905. (47) Olivier, Y.; Niedzialek, D.; Lemaur, V.; Pisula, W.; Müllen, K.; Koldemir, U.; Reynolds, J. R.; Lazzaroni, R.; Cornil, J.; Beljonne, D. Adv. Mater. 2014, 26, 2119−2136. (48) Schacher, F. H.; Rupar, P. A.; Manners, I. Angew. Chem., Int. Ed. 2012, 51, 7898−7921. (49) Black, C. T.; Ruiz, R.; Breyta, G.; Cheng, J. Y.; Colburn, M. E.; Guarini, K. W.; Kim, H.-C.; Zhang, Y. IBM J. Res. Dev. 2007, 51, 605− 633. (50) Segalman, R. A.; McCulloch, B.; Kirmayer, S.; Urban, J. J. Macromolecules 2009, 42, 9205−9216. (51) Lazzari, M.; López-Quintela, M. Adv. Mater. 2003, 15, 1583− 1594. (52) Yeh, P. Optical Waves in Layered Media; Goodman, J. W., Ed.; John Wiley & Sons, Inc.: New York, 1988. (53) Stenzel, O. The Physics of Thin Film Optical Spectra: An Introduction; Ertl, G., Lüth, H., Mill, D. L., Eds.; Springer: Berlin, 2005.

Constraints to Macroscopic Transport”, and by the Helmholtz Association (Helmholtz-Energie-Allianz “Hybrid-Photovoltaik”) is highly acknowledged

■

REFERENCES

(1) Mas-Torrent, M.; Rovira, C. Chem. Soc. Rev. 2008, 37, 827−838. (2) Clarke, T. M.; Durrant, J. R. Chem. Rev. 2010, 110, 6736−6767. (3) Brütting, W.; Frischeisen, J.; Schmidt, T. D.; Scholz, B. J.; Mayr, C. Phys. Status Solidi A 2013, 210, 44−65. (4) Torsi, L.; Magliulo, M.; Manoli, K.; Palazzo, G. Chem. Soc. Rev. 2013, 42, 8612−8628. (5) Jesper, M.; Alt, M.; Schinke, J.; Hillebrandt, S.; Angelova, I.; Rohnacher, V.; Pucci, A.; Lemmer, U.; Jaegermann, W.; Kowalsky, W.; et al. Langmuir 2015, 31, 10303−10309. (6) Fazzi, D.; Caironi, M. Phys. Chem. Chem. Phys. 2015, 17, 8573− 8590. (7) Baeg, K.-J.; Caironi, M.; Noh, Y.-Y. Adv. Mater. 2013, 25, 4210− 4244. (8) Arias, A. C.; MacKenzie, J. D.; McCulloch, I.; Rivnay, J.; Salleo, A. Chem. Rev. 2010, 110, 3−24. (9) Krebs, F. C. Sol. Energy Mater. Sol. Cells 2009, 93, 394−412. (10) Alt, M.; Schinke, J.; Hillebrandt, S.; Hänsel, M.; HernandezSosa, G.; Mechau, N.; Glaser, T.; Mankel, E.; Hamburger, M.; Deing, K.; et al. ACS Appl. Mater. Interfaces 2014, 6, 20234−20241. (11) Moore, J. R.; Albert-Seifried, S.; Rao, A.; Massip, S.; Watts, B.; Morgan, D. J.; Friend, R. H.; McNeill, C. R.; Sirringhaus, H. Adv. Energy Mater. 2011, 1, 230−240. (12) Mori, D.; Benten, H.; Okada, I.; Ohkita, H.; Ito, S. Energy Environ. Sci. 2014, 7, 2939−2943. (13) Rivnay, J.; Mannsfeld, S. C. B.; Miller, C. E.; Salleo, A.; Toney, M. F. Chem. Rev. 2012, 112, 5488−5519. (14) Luzio, A.; Criante, L.; D’Innocenzo, V.; Caironi, M. Sci. Rep. 2013, 3, 3425. (15) Steyrleuthner, R.; Di Pietro, R.; Collins, B. A.; Polzer, F.; Himmelberger, S.; Schubert, M.; Chen, Z.; Zhang, S.; Salleo, A.; Ade, H.; et al. J. Am. Chem. Soc. 2014, 136, 4245−4256. (16) Osaka, I.; Takimiya, K. Polymer 2015, 59, A1−A15. (17) Fabiano, S.; Himmelberger, S.; Drees, M.; Chen, Z.; Altamimi, R. M.; Salleo, A.; Loi, M. A.; Facchetti, A. Adv. Energy Mater. 2014, 4, 1301409. (18) Rivnay, J.; Toney, M. F.; Zheng, Y.; Kauvar, I. V.; Chen, Z.; Wagner, V.; Facchetti, A.; Salleo, A. Adv. Mater. 2010, 22, 4359−4363. (19) Schuettfort, T.; Huettner, S.; Lilliu, S.; Macdonald, J. E.; Thomsen, L.; McNeill, C. R. Macromolecules 2011, 44, 1530−1539. (20) Schuettfort, T.; Thomsen, L.; McNeill, C. R. J. Am. Chem. Soc. 2013, 135, 1092−1101. (21) Giussani, E.; Fazzi, D.; Brambilla, L.; Caironi, M.; Castiglioni, C. Macromolecules 2013, 46, 2658−2670. (22) Giussani, E.; Brambilla, L.; Fazzi, D.; Sommer, M.; Kayunkid, N.; Brinkmann, M.; Castiglioni, C. J. Phys. Chem. B 2015, 119, 2062− 2073. (23) Anton, A. M.; Steyrleuthner, R.; Kossack, W.; Neher, D.; Kremer, F. J. Am. Chem. Soc. 2015, 137, 6034−6043. (24) Rivnay, J.; Steyrleuthner, R.; Jimison, L. H.; Casadei, A.; Chen, Z.; Toney, M. F.; Facchetti, A.; Neher, D.; Salleo, A. Macromolecules 2011, 44, 5246−5255. (25) Takacs, C. J.; Treat, N. D.; Krämer, S.; Chen, Z.; Facchetti, A.; Chabinyc, M. L.; Heeger, A. J. Nano Lett. 2013, 13, 2522−2527. (26) Sciascia, C.; Martino, N.; Schuettfort, T.; Watts, B.; Grancini, G.; Antognazza, M. R.; Zavelani-Rossi, M.; McNeill, C. R.; Caironi, M. Adv. Mater. 2011, 23, 5086−5090. (27) Martino, N.; Fazzi, D.; Sciascia, C.; Luzio, A.; Antognazza, M. R.; Caironi, M. ACS Nano 2014, 8, 5968−5978. (28) Kossack, W.; Papadopoulos, P.; Heinze, P.; Finkelmann, H.; Kremer, F. Macromolecules 2010, 43, 7532−7539. (29) Jasiurkowska, M.; Kossack, W.; Ene, R.; Iacob, C.; Kipnusu, W. K.; Papadopoulos, P.; Sangoro, J. R.; Massalska-Arodz, M.; Kremer, F. Soft Matter 2012, 8, 5194−5200. I


Spatial Orientation and Order of Structure-Defining Subunits in Thin

Recommend Documents