- Email: [email protected]
PCDTBT heterojunctions: Influence on nanoscale charge transport
Polymer 77 (2015) 70e78
Contents lists available at ScienceDirect
Polymer journal homepage: www.elsevier.com/locate/polymer
Competition between phase separation and structure confinement in P3HT/PCDTBT heterojunctions: Influence on nanoscale charge transport Alvaro Rodríguez-Rodríguez a, Michelina Soccio b, Daniel E. Martínez-Tong a, 1, rrez a, * Tiberio A. Ezquerra a, Benjamin Watts c, Mari-Cruz García-Gutie a b c
Instituto de Estructura de la Materia (IEM-CSIC), Serrano 121, 28006 Madrid, Spain di Bologna, via Terracini 28, 40131 Bologna, Italy Dipartimento di Ingegneria Civile, Chimica, Ambientale e dei Materiali, DICAM-Universita Swiss Light Source, Paul Scherrer Institut, 5232 Villigen, Switzerland
a r t i c l e i n f o
a b s t r a c t
Article history: Received 6 July 2015 Received in revised form 25 August 2015 Accepted 4 September 2015 Available online 9 September 2015
A strong impact of crystal morphology on hole mobility is evidenced due to competition between phase separation and structure confinement in P3HT/PCDTBT heterojunctions. We find that domain sizes of both components decrease as film thickness decreases, suggesting an initial advantageous scenario for the efficiency of bulk heterojunction solar cells. However the P3HT/PCDTBT (1:1) films with a thickness of 165 nm and thicker present a dense crystal needle-like morphology while no evidence of needle-like motifs appears in the films 58 nm thick, indicating that confinement inhibiting crystallization takes place due to the very thin P3HT domains of only 40 nm in the thinner sample. The structural studies were correlated with nanoscale charge transport by conductive-AFM. A significant zero-field hole mobility increase is observed for the P3HT/PCDTBT films with increasing thickness. Even more interesting is the fact that a zero-field hole mobility increase of about two orders of magnitude is observed for P3HT domains in a 165 nm blend film compared to a neat P3HT film with similar thickness. This observation is related to the highly conductive network induced in P3HT by the presence of the PCDTBT phase, consisting of needle-like crystals growing from the P3HT domains and acting as bridges through the PCDTBT domains. © 2015 Elsevier Ltd. All rights reserved.
Keywords: Confinement All-polymer heterojunctions Charge transport
1. Introduction Organic photovoltaics (OPVs) are one class of solar-energy conversion devices, offering the advantages of low cost, lightweight, solution-processability and mechanical flexibility over existing photovoltaic technologies [1e3]. Polymer solar cells (PSCs) are attractive due to a number of advantageous features [2], including their thin-film architecture and low material consumption resulting from a high absorption coefficient, their use of organic solution processes and low manufacturing energy requirements. Efficient PSCs typically employ a bulk-heterojunction (BHJ) photoactive layer, where an electron-donating (p-type)
* Corresponding author. rrez). E-mail address: [email protected] (M.-C. García-Gutie 1 partement de Physique, Faculte des Sciences, Universite Present address: De libre de Bruxelles (ULB), Boulevard du Triomphe, 1050 Brussels, Belgium. http://dx.doi.org/10.1016/j.polymer.2015.09.012 0032-3861/© 2015 Elsevier Ltd. All rights reserved.
material and an electron-accepting (n-type) material form a nanosized phase-separated interpenetrating network, which can provide large enough heterointerface areas for efficient exciton dissociation and a continuous pathway for charge transport [4]. Fullerene derivatives are the most effective and commonly used acceptor materials. Polymer/fullerene BHJ solar cells have proven power conversion efficiencies of over 10% now reported [5e7]. Allpolymer solar cells (all-PSCs) have been also developed, where an n-type semiconducting polymer is used as the electron acceptor instead of a fullerene derivative blended with a p-type polymer. Although the efficiency of all-PSCs remains at ~4% [8,9], they offer some potential advantages including superior optical absorption and greater synthetic flexibility in tuning semiconducting properties such as the bandgap and energy levels [10,11]. The bulk heterojunction concept relies on the high purity of donor and acceptor phases within the characteristic exciton diffusion length of ~10 nm, requiring percolating interconnected network morphology. The understanding of phase separation and
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
morphology in all-polymer BHJ is crucial for the optimization of device performance [11,12]. In general, a mixing of donor and acceptor on a length scale smaller than the exciton diffusion length optimizes charge generation, while coarser morphologies optimize the separation of charges away from the interface and their collection at the device electrodes. On the other hand, crystal morphology [13,14] and structure confinement effects [15e17] such as crystallization inhibition in domains of tens of nanometers also play very important roles on charge transport and device performance. In this paper, we present the morphology developed as the result of competition between phase separation and structure confinement in all-polymer bulk-heterojunctions based on the blends of poly-3-hexylthiophene (P3HT) and poly[N-90 -heptadecanyl-2,7-carbazole-alt-5,5-(40 ,70 -di-2-thienyl-20 ,10,30 benzothiadiazole)] (PCDTBT). The influence of morphology on nanoscale charge transport is also reported.
2. Experimental section 2.1. Materials Poly[N-90 -heptadecanyl-2,7-carbazole-alt-5,5-(40 ,70 -di-2 thienyl-20 ,10,30 benzothiadiazole)] (PCDTBT) and poly-3hexylthiophene (P3HT) were purchased from Ossila. PCDTBT (Mw ¼ 35400 g/mol, PDI ¼ 2.4), P3HT (Mw ¼ 34100 g/mol, PDI ¼ 1.7; regioregularity ¼ 94.7%). The chemical structures of P3HT and PCDTBT are shown in Fig. 1. The HOMO and LUMO energy levels of P3HT and PCDTBT with respect to the vacuum level as supplied by Ossila are (HOMO: 4.7 eV; LUMO: 2.7 eV) and (HOMO: 5.35 eV; LUMO: 3.5 eV) respectively.
2.2. Sample preparation P3HT and PCDTBT were dissolved separately in chlorobenzene (CB) and afterwards blended with a weight ratio of 1:1. Thin films of pristine polymers and of the blend were prepared by spin-coating on silicon wafers for investigation with grazing incidence wide angle X-ray scattering techniques (GIWAXS) and on Arsenic ndoped Silicon substrates (resistivity ~ 0.001 Ucm) for morphology and electrical currentevoltage (IeV) characterization by atomic force microscopy (AFM) in tapping mode and by conducting atomic force microscopy (C-AFM) respectively. A fixed amount of 0.2 mL of polymer solution was dropped by a syringe on a square silicon substrate placed in the center of a rotating metallic horizontal plate. A rotation rate of 2400 rpm was kept during 120 s. For scanning transmission X-ray spectro-microscopy (STXM) measurements thin films spin-coated on silicon substrates were then floated off into a very dilute NaOH/water solution (0.25 wt%) and finally picked up with Transmission Electron Microscopy (TEM) grids. Solutions with concentrations of 12, 24 and 36 mg/mL were used in order to obtain films with different thicknesses (Table 1).
71
2.3. Ultravioletevisible spectroscopy Absorption spectra of polymer thin films spin-coated on silicon substrates were recorded in reflection geometry using a UVevisible spectrophotometer (UV-3600, Shimadzu). 2.4. X-ray microscopy Scanning Transmission X-ray Spectro-Microscopy (STXM), with image contrast based on Near-Edge X-ray Absorption Fine Structure (NEXAFS) spectromicroscopy, was performed at the PolLux beamline [18e20] at the Swiss Light Source, Paul Scherrer Institute, Villigen, Switzerland. TEM grids-supported films were mounted in the sample chamber which was evacuated to low vacuum. The transmitted X-ray intensity through the film was recorded using a scintillator and photo-multiplier tube and measured as a function of energy (275.0e345.0 eV with a resolution of 0.1 eV) and position (with X-ray focus of about 25 nm diameter and XeY sample position of 1 nm). Transmitted X-ray intensity was converted to an Xray optical density (defined as OD ¼ ln (I/I0)) by recording the Xray intensity through an empty TEM grid. Analysis of images was performed by the aXis2000 software package [21]. 2.5. Grazing incidence X-ray scattering Information of the inner film structure was obtained by grazing incidence X-ray scattering experiments performed at the synchrotron beamline BW4 at HASYLAB (DORIS, DESY) in Hamburg. The experimental setup has been previously described [22]. An Xray wavelength of l ¼ 0.13808 nm, with a beam size (horizontal vertical) of 40 20 mm2 and an exposure time of 4 min was used in our measurements. Scattered intensity was recorded by a MarCCD detector of 2048 2048 pixels with a resolution of 79.1 mm per pixel. A sample-detector distance of 10.6 cm was used to investigate crystallinity and orientation by GIWAXS. The incident angle (ai) was set to 0.40 , which was above the critical angles of P3HT and PCDTBT and therefore the structure of the full film thickness was detected. 2.6. Scanning probe measurements AFM measurements were done in air under ambient conditions using a commercial scanning probe microscope (MultiMode 8 equipped with a conductive-AFM (C-AFM) module and the Nanoscope V controller, Bruker). The topography AFM images were collected in tapping mode using silicon probes (NSG30 probes by NT-MDT); topographic images were collected from multiple locations to examine the film uniformity. For C-AFM measurements, PtIr-coated silicon tips with a spring constant of 0.2 N/m and a tip radius of ca. 25 nm (SCM-PIC by Bruker) were used. The contact area was estimated using the Hertz model [23,24]. In these measurements, the conducting probe makes contact with the sample (the tip acts as a nanoelectrode) and either measures the current as a function of the applied voltage at certain points on a surface (IeV curve) or maps out a current image at a fixed bias. A scheme with the C-AFM measurement set-up is shown in Fig. 2. The bias was applied to the conducting substrate, and the current was measured by a preamplifier. For each sample, several IeV curves were collected at various locations to obtain an average IeV curve. 3. Results and discussion 3.1. Optical properties
Fig. 1. Chemical structures of P3HT (left) and PCDTBT (right).
The normalized ultravioletevisible (UVevisible) absorption
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
72
Table 1 Description of the investigated samples: Solution concentration, average film thickness measured by AFM (for the P3HT/PCDTBT blends the domain thickness of each component is indicated) and zero-field hole mobility. Sample
Concentration (mg/mL)
Thickness (nm)
m0 (cm2 V1 s1)
P3HT PCDTBT P3HT/PCDTBT_1 P3HT/PCDTBT_2 P3HT/PCDTBT_3
24 24 12 24 36
145 190 58 (40P3HT-75PCDTBT) 165 (140P3HT-190PCDTBT) 295 (270P3HT-320PCDTBT)
2.3 102 e 1.2 105 0.61 1.21
A
V
P3HT:PCDTBT n Si Metallic support Fig. 2. C-AFM experimental setup.
spectra of pristine P3HT, PCDTBT and the blend P3HT/PCDTBT_2 films spin-cast from chlorobenzene (see Table 1), are compared in Fig. 3. P3HT features a main absorption band (1P3HT) with the maximum at about 524 nm, related to the electronic pep* transition [25]. PCDTBT exhibits two pronounced absorption bands: the 1PCDTBT at approximately 385 nm attributable to the pep* transition of their conjugated main chains and the 2PCDTBT at about 563 nm attributable to the intramolecular charge transfer (ICT) interaction between carbazole moieties and benzothiadiazole moieties [26]. Both, P3HT and PCDTBT show vibronic structure, suggesting good light-harvesting ability. Fig. 3 also shows internal order in the blend, as evidenced by the red shift of its absorption spectrum with respect to the neat absorption spectra and appearance of vibronic structure [27].
3.2. Surface morphology AFM in tapping mode was used to characterize the surface topography at the nanoscale and the thickness of P3HT/PCDTBT (1:1) films prepared from different solution concentrations. The surface topography images of as-cast films are shown in Fig. 4
Absorbance (a.u.)
1,0
1PCDTBT
1P3HT
P3HT PCDTBT P3HT/PCDTBT
0,8
0,6
0,4
2PCDTBT
0,2
0,0 300
400
500
600
700
Wavelength (nm) Fig. 3. Normalized UVevis absorption spectra of thin films of P3HT (blue), PCDTBT (red) and the blend P3HT/PCDTBT_2 (green). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
(except for the P3HT/PCDTBT_3 sample that is shown in Fig. S1 of the Supplementary Information). Neat P3HT and PCDTBT thin films are included for comparison. The topography of the P3HT (Fig. 4a) and PCDTBT (Fig. 4b) films is smooth and very small fiberlike crystallites can be observed for the P3HT (inset in Fig. 4a). P3HT/PCDTBT (1:1) films show the characteristic morphology of lateral phase separation, indicating immiscibility of the two polymers [28,29]. P3HT/PCDTBT_2 and P3HT/PCDTBT_3 thin films present a 50 nm difference in height between phases for an average film thickness of 165 and 295 nm respectively, the P3HT/PCDTBT_1 thin film shows 35 nm difference in height between phases for an average film thickness of 58 nm (see Table 1). It is shown that domain size as well as film thickness increase as increasing polymer concentration (Figs. 4c, d and S1a). It is also interesting to point out the presence of a dense needle-like morphology in films with 165 and 295 nm average thicknesses (insets in Figs. 4d and S1a respectively) while no evidence of needle-like motifs appears in the thin film with 58 nm average thickness (inset in Fig. 4c). The samples were then annealed at 140 C for 20 min (above the glass transition temperature of both polymers [30]) and cooled to room temperature aiding the thermodynamic equilibration of the samples, but no change in morphology was observed. 3.3. Composition mapping and internal structure Since there is not a single technique able to fully characterize the morphology of a thin-film polymer blend, the use of complementary techniques is required. Although AFM enables imaging of surface topography with high spatial resolution, to distinguish regions with different chemical composition can be cumbersome. Scanning Transmission X-ray Spectro-Microscopy (STXM) is a technique that provides an appropriate combination of spectroscopy and imaging capabilities at low radiation doses. As a powerful synchrotron-based technique, STXM can quantitatively map the chemical composition of organic films with nanoscale resolution [31,32]. The STXM images can provide information on the domain sizes, shapes, and purities of thin film polymer blends with a spatial resolution of ~20 nm [33]. In order to obtain quantitative chemical composition of the blend thin films we have used the singular value decomposition (SVD) mathematical procedure [31,34]. SVD requires previous knowledge of the mass absorption coefficients of the blend components. Fig. 5 shows the mass absorption coefficients as a function of energy for P3HT and PCDTBT pristine films of known thickness (Table 1). The transmitted X-ray intensity measured was converted to an X-ray optical density (defined as OD ¼ ln(I/I0)) where I is the transmitted intensity and I0 is the incident intensity. Since OD ¼ mrt, where m is the mass absorption coefficient, r is the density and t the sample thickness, we have obtained the mass absorption coefficient of each homopolymer considering density values of 1.13 and 1.33 g/cm3 for P3HT and PCDTBT respectively [35]. The SVD procedure also requires a series of images acquired at a number of energies that equal or exceed the number of compositional components. STXM images (Fig. 5 top) of the same area of
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
73
Fig. 4. Tapping mode AFM height images of (a) P3HT, (b) PCDTBT, (c) P3HT/PCDTBT_1 and (d) P3HT/PCDTBT_2.
the P3HT/PCDTBT_2 thin film were taken at energies indicated by arrows in the NEXAFS spectra which correspond to 280 eV (preedge), 284.2 eV (PCDTBT resonance), 287.8 eV (P3HT resonance) and 320 eV (chemically insensitive). By using the aXis2000 software package [21] where the SVD procedure is integrated, the information contained within the set of STXM images taken at the
284.2 eV
320 eV
287.8 eV
4
5x10
2
mass absorption coefficient (cm /g)
280 eV
selected energies can be transformed into maps quantifying the composition and thickness of the sample in each pixel [31,34]. The corresponding composition and thickness maps of the P3HT/ PCDTBT_2 thin film are shown in Fig. 6. On the one hand, these images reveal the chemical nature of the different observed
4
4x10
4
3x10
283
286
289
Energy (eV)
4
2x10
4
1x10
275
285
295
305
315
325
335
345
Energy (eV)
Fig. 5. Reference NEXAFS spectra for pristine films of P3HT (red line) and PCDTBT (blue line). Vertical arrows indicate the photon energies used to obtain the raw STXM images (6 mm 6 mm) of the same sample area of a P3HT/PCDTBT_2 thin film, presented at the top. The inset highlights differences in the NEXAFS spectra of P3HT and PCDTBT in the C-1s / p* region. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 6. 6 mm 6 mm images of quantitative composition (a) and thickness (b) maps of the P3HT/PCDTBT_2 thin film, calculated from the set of raw images in Fig. 5 top.
74
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
domains. On the other hand they reveal a structure of thinner P3HT-rich domains enclosed by a thicker PCDTBT-rich matrix. The thicknesses of domains are in agreement with the values measured by AFM. In order to complement our STXM studies, we have used GIWAXS in order to probe changes in molecular and nanoscale structure in P3HT, PCDTBT and in the blend P3HT/PCDTBT_2 thin films, before and after annealing at 140 C for 20 min. Fig. 7a shows the 2D scattering pattern for the P3HT as-cast film. Four main peaks can be seen, three out of plane reflections h00 having q-values of 3.8, 7.59 and 11.55 nm1, and one in plane peak superposition of the 020 and the weak 002 reflections [36], with a q-value of 16.39 nm1. In agreement with previous reports [36,37], the pattern reveals the characteristic crystalline reflections corresponding to the crystal sheets formed by the pep stacking of the thiophene rings h00 and the interchain distances 0k0. In addition, the P3HT as-cast film is uniaxially oriented with mainly an edge-on configuration which corresponds to the usual conformation adopted by P3HT thin films, consisting of polymer chains parallel to the substrate. The 2D scattering pattern of PCDTBT as-cast film, shown in Fig. 7b, reveals a lower degree of crystallinity compared to P3HT. The broad peak at q ¼ 14.55 nm1 in the out-of-plane direction, indexed as 010, corresponds to pep stacking between PCDTBT back-bones while the diffraction peak at q ¼ 2.95 nm1 also in the out of plane direction, indexed as 100, corresponds to the alkyl side-chain packing [14,30,35]. The blend P3HT/PCDTBT_2 as-cast thin film presents a 2D scattering pattern (Fig. 7c) very similar to P3HT, where only the P3HT reflections are detected. These
reflections appear at similar q values as those of P3HT. No evidence of PCDTBT reflections is observed. These data reveal that P3HT and PCDTBT do not co-crystallize. After annealing the samples at 140 C for 20 min, the P3HT thin film exhibits more intense reflections (Fig. 7d) suggesting an increase in crystallinity and a similar trend is observed for the blend (Fig. 7f). However, for PCDTBT Fig. 7e shows that while the 100 peak is better defined, the 010 loses intensity after annealing. This effect is even more evident for a thinner PCDTBT sample, as can be seen in Fig. S2. This trend can be explained by considering that annealing improves side-chain order while reducing the pep stacking, as has been previously proposed by Wang et al. [30]. 3.4. Nanoscale charge transport properties Conductive-AFM (C-AFM) was used to characterize the electrical properties at the nanoscale of P3HT/PCDTBT (1:1) blend thin films with different thicknesses. Neat P3HT films were included in the study for the sake of comparison. A scheme with the C-AFM measurement set-up is shown in Fig. 2. Hole current imaging was used to examine conductivity variations and to map the P3HT conducting network in the P3HT/PCDTBT (1:1) thin films. In these measurements, PtIr-coated silicon probes were used to ensure that holes are the major injected carriers [23]. Fig. 8 shows contact mode AFM topography and current images recorded simultaneously by applying a constant voltage of 5 V on the conducting substrate. The current image of the P3HT thin film (Fig. 8d) shows a little current variation throughout the film. However, the P3HT/
Fig. 7. 2D GIWAXS images of as-cast thin films: a) P3HT, b) PCDTBT, c) P3HT/PCDTBT_2. And after annealing at 140 C for 20 min: d) P3HT, e) PCDTBT, f) P3HT/PCDTBT_2.
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
75
Fig. 8. Contact mode AFM topography (left) and current (right) images of P3HT (a, d), P3HT/PCDTBT_1 (b, e) and P3HT/PCDTBT_2 (c, f). Blue and red crosses indicate the positions where IeV curves presented in Fig. 9a were measured. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
PCDTBT (1:1) films show domains with high (dark region) and low (bright region) current (Fig. 8e and f). By comparison with the compositional maps (Fig. 6), we can assign the high current regions to the P3HT-rich domains and the low current ones to PCDTBT. It is known that the thickness of thin films prepared by spin-coating is mainly related to the solution concentration. Table 1 shows the thicknesses of neat P3HT and PCDTBT thin films and the average thickness as well as the thicknesses of P3HT-rich and PCDTBT-rich domains of P3HT/PCDTBT (1:1) thin films prepared from solutions with different concentrations. We observe in Fig. 8 that the current in the P3HT-rich domains increases enormously with increasing thickness, from hundreds of pA for the blend with a P3HT-rich domain thickness of 40 nm (Fig. 8e) to several nA for blends with P3HT-rich domains thicker than 100 nm (Figs. 8f and S1c). In addition, we pointed out previously the presence of a dense needlelike morphology in P3HT/PCDTBT_2 and P3HT/PCDTBT_3 films with 165 and 295 nm thickness respectively (Figs. 4d and S1a) while no evidence of needle-like motives appears in the 58 nm thick P3HT/PCDTBT_1 film (Fig. 4c). The current images of P3HT/
PCDTBT_2 and P3HT/PCDTBT_3 films (Figs. 8f and S1c), clearly show a fibrous network where the strongest current is measured. These results evidence that P3HT needle-like crystals grow from the P3HT-rich domains acting as bridges through the PCDTBT-rich domains. However, the current image of the P3HT/PCDTBT_1 thin film (Fig. 8e) does not show any fibrous network nor evidence of needle-like motifs (Fig. 4c). A fact that can be explained by confinement inhibiting crystallization due to the very thin domains of P3HT of only 40 nm as it is supported by GIWAXS measurements (Fig. S3). In addition to current imaging with C-AFM, local IeV curves can also be collected. This technique of measuring IeV characteristics by C-AFM has the advantage that only a very small region of the film, comparable to the tip contact area, is probed [23,38,39]. Thus, several areas can be examined aiming to provide information on the film electrical heterogeneity. Several IeV curves were collected from different locations within a P3HT-rich domain (red crosses in Fig. 8f) and from different locations within a PCDTBT-rich domain (blue crosses in Fig. 8f). IeV curves measured in both regions
76
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
including the average curves (red and blue for the P3HT rich domain and the PCDTBT rich domain respectively) are shown in Fig. 9a. As one can see, while there is a negligible electrical current in the PCDTBT-rich domains, a significant electrical current is measured in the P3HT-rich domains. The same procedure was used for all the samples investigated. Average IeV curves for the neat P3HT, neat PCDTBT and P3HT-rich domains in P3HT/PCDTBT (1:1)
films with different thicknesses are shown in Fig. 9b. The highest current for a similarly applied bias is measured for the P3HT/ PCDTBT_2 thin film. Since the samples have different thicknesses, Fig. 9c presents the I-E curves where E ¼ V⁄L is the electric field and L is the P3HT domain thickness. The logelog plot of the jJj jVj data, where J is the current density, is presented in the inset of Fig. 9b. It shows that for the investigated samples the electrical response exhibits two regimes: initially a short Ohmic range (JfV) and a more extended regime (JfV2) denominated space-charge limited current (SCLC) regime. In this case (SCLC), the nanoscale charge mobility can be extracted by using the following expression [40,41]:
J¼
9 V2 εr ε0 m 3 8 L
(1)
where εr is the dielectric constant of the polymer for which we assumed a value of 3 as proposed in the literature [42], ε0 ¼ 8.85 1012 F/m is the vacuum permittivity, m is the charge mobility, V is the applied voltage, and L is the P3HT domain thickness. In order to verify whether the calculated carrier mobilities are field dependent we have used the approach for a field-dependent carrier mobility described by the Poole-Frenkel equation [40,41,43]:
m ¼ m0 eðE=E0 Þ
1=2
(2)
where E is the electric field, m0 is the zero-field mobility, and E0 is the field coefficient. The Poole-Frenkel plots (lnm vs. E1/2) are shown in Fig. 10. All the samples exhibit similar behavior showing that the mobility depends weakly on the electric field. A significant zero-field hole mobility increase is observed for the P3HT/PCDTBT (1:1) films with increasing thickness (Table 1): from 1.2 105 cm2 V1 s1 for the P3HT/PCDTBT_1 thin film to 0.61 cm2 V1 s1 and 1.21 cm2 V1 s1 for the P3HT/PCDTBT_2 and P3HT/PCDTBT_3 thin films respectively. It is also interesting to note that a zero-field hole mobility increase of about two orders of magnitude is observed for the P3HT rich domains in the P3HT/ PCDTBT_2 thin film compared to the neat P3HT sample with similar thickness. It indicates a P3HT and PCDTBT synergy in the blend, probably related to the highly conductive needle like morphology induced in P3HT domains by the presence of the PCDTBT phase, as it is shown in Figs. 4d and 8f.
Fig. 9. a) IeV curves measured in different locations into a rich P3HT domain (red crosses in Fig. 8f) and in different locations into a rich PCDTBT domain (blue crosses in Fig. 8f). Average IeV curves, red and blue for the P3HT rich domain and the PCDTBT rich domain are also included. b) Average IeV curves for pristine P3HT and PCDTBT thin films and P3HT rich domains in P3HT/PCDTBT (1:1) thin films with different thicknesses. The inset shows the logelog plot of the jJjejVj data where two different regimes of the electrical response can be observed. c) Average I-E curves for pristine P3HT and PCDTBT thin films and P3HT rich domains in P3HT/PCDTBT (1:1) thin films with different thicknesses. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 10. Hole mobilities for pristine P3HT (black) and P3HT rich domains in P3HT/ PCDTBT_1 (green), P3HT/PCDTBT_2 (red) and P3HT/PCDTBT_3 (blue). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
4. Conclusions In summary, using several complementary techniques we have investigated the optical, structural and hole transport properties of the P3HT/PCDTBT (1:1) polymer blend films as a function of film thickness. It is seen that the optical absorption of the blend is improved compared to the neat P3HT and PCDTBT UVevis absorption spectra, covering a broader range of the solar spectrum. We find that the P3HT/PCDTBT is an immiscible blend with lateral phase separation when spin-coated. The domain sizes of both components decrease as film thickness decreases. This suggests an initial advantageous scenario for the efficiency of bulk heterojunction solar cells, but a competition between the favorable decrease of domain size and the resulting structure confinement takes place. The P3HT/PCDTBT (1:1) thin films with a thickness of 165 nm and greater present a dense crystal needle-like morphology while no evidence of needle-like motifs appears in the 58 nm thick film. In addition, conductive-AFM was used to characterize the electrical properties at the nanoscale. It clearly shows that blend films 165 and 295 nm thick present a fibrous network where the strongest current is measured. These results evidence that P3HT needle-like crystals grow from the P3HT-rich domains, acting as bridges through the PCDTBT-rich domains. However, the current image of the 58 nm thick blend film neither shows fibrous network, nor evidence of needle-like motifs. This fact can be explained by confinement inhibiting crystallization due to the very thin domains of P3HT of only 40 nm thickness. A strong impact of the crystal morphology on hole mobility is evidenced. A significant zero-field hole mobility increase is observed for the P3HT/PCDTBT (1:1) thin films with increasing thickness: from 1.2 105 cm2 V1 s1 for the 58 nm thick film to 0.61 cm2 V1 s1 and 1.21 cm2 V1 s1 for the 165 nm and a 295 nm thick films respectively. Even more interesting is the fact that a zero-field hole mobility increase of about two orders of magnitude is observed for the P3HT rich domains in the 165 nm P3HT/PCDTBT (1:1) thin film compared to the neat P3HT thin film with similar thickness, probably related to the highly conductive needle like network induced in P3HT by the presence of the PCDTBT phase. Acknowledgments The authors gratefully acknowledge the financial support of the Spanish Ministry of Economy and Competitiveness (MINECO) through the project MAT 2011-23455. A. R-R is indebted to MINECO for the tenure of a FPI BES-2013-062620 contract associated to the project MAT2012-33517. We thank the Swiss Light Source for beamtime at PolLux. The PolLux end station was financed by the German Minister für Bildung und Forschung (BMBF) through contracts 05KS4WE1/6 and 05KS7WE1. Parts of this research were carried out at the light source DORIS III at DESY, a member of the Helmholtz Association (HGF). We would like to thank J. Perlich for assistance in using beamline BW4. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.polymer.2015.09.012. References [1] S.R. Forrest, The path to ubiquitous and low-cost organic electronic appliances on plastic, Nature 428 (2004) 911e918. [2] G. Dennler, M.C. Scharber, C.J. Brabec, Polymer-fullerene bulkheterojunction solar cells, Adv. Mater. 21 (2009) 1323e1338. [3] Y. Tang, C.R. McNeill, All-polymer solar cells utilizing low band gap polymers as donor and acceptor, J. Polym. Sci. Pt. B Polym. Phys. 51 (2013) 403e409.
77
[4] G. Yu, J. Gao, J.C. Hummelen, F. Wudl, A.J. Heeger, Polymer photovoltaic cells [mdash] enhanced efficiencies via a network of internal donor-acceptor heterojunctions, Science 270 (1995) 1789e1791. [5] Z. He, C. Zhong, S. Su, M. Xu, H. Wu, Y. Cao, Enhanced power-conversion efficiency in polymer solar cells using an inverted device structure, Nat. Phot. 6 (2012) 591e595. [6] C.M. Amb, S. Chen, K.R. Graham, J. Subbiah, C.E. Small, F. So, J.R. Reynolds, Dithienogermole As a Fused Electron Donor in Bulk Heterojunction Solar Cells, J. Am. Chem. Soc. 133 (2011) 10062e10065. [7] C.E. Small, S. Chen, J. Subbiah, C.M. Amb, S.-W. Tsang, T.-H. Lai, J.R. Reynolds, F. So, High-efficiency inverted dithienogermole-thienopyrrolodione-based polymer solar cells, Nat. Phot. 6 (2012) 115e120. [8] E. Zhou, J. Cong, K. Hashimoto, K. Tajima, Control of Miscibility and Aggregation Via the Material Design and Coating Process for High-Performance Polymer Blend Solar Cells, Adv. Mater. 25 (2013) 6991e6996. [9] D. Mori, H. Benten, I. Okada, H. Ohkita, S. Ito, Low-bandgap donor/acceptor polymer blend solar cells with efficiency exceeding 4%, Adv. Energy Mater. 4 (2014) 1301006. [10] C.R. McNeill, N.C. Greenham, Conjugated-Polymer Blends for Optoelectronics, Adv. Mater. 21 (2009) 3840e3850. [11] A. Sepe, Z. Rong, M. Sommer, Y. Vaynzof, X. Sheng, P. Muller-Buschbaum, D.M. Smilgies, Z.-K. Tan, L. Yang, R.H. Friend, U. Steiner, S. Huttner, Structure formation in P3HT/F8TBT blends, Energy Environ. Sci. 7 (2014) 1725e1736. [12] C.R. McNeill, A. Abrusci, I. Hwang, M.A. Ruderer, P. Mueller-Buschbaum, N.C. Greenham, Photophysics and photocurrent generation in polythiophene/ polyfluorene copolymer blends, Adv. Funct. Mater. 19 (2009) 3103e3111. [13] F.S. Kim, G. Ren, S.A. Jenekhe, One-dimensional nanostructures of p-conjugated molecular systems: assembly, properties, and applications from photovoltaics, sensors, and nanophotonics to nanoelectronics, Chem. Mat. 23 (2010) 682e732. [14] H.-W. Wang, E. Pentzer, T. Emrick, T.P. Russell, Preparation of low band gap fibrillar structures by solvent-induced crystallization, ACS Macro Lett. 3 (2013) 30e34. [15] J.J. Hernandez, D.R. Rueda, M.C. Garcia-Gutierrez, A. Nogales, T.A. Ezquerra, M. Soccio, N. Lotti, A. Munari, Structure and morphology of thin films of linear aliphatic polyesters prepared by spin-coating, Langmuir 26 (2010) 10731e10737. [16] M.-C. Garcia-Gutierrez, A. Linares, I. Martin-Fabiani, J.J. Hernandez, M. Soccio, D.R. Rueda, T.A. Ezquerra, M. Reynolds, Understanding crystallization features of P(VDF-TrFE) copolymers under confinement to optimize ferroelectricity in nanostructures, Nanoscale 5 (2013) 6006e6012. rrez, R.J. Davies, [17] W. Pisula, M. Kastler, B. El Hamaoui, M.-C. García-Gutie C. Riekel, K. Müllen, Dendritic morphology in homeotropically aligned discotic films, ChemPhysChem 8 (2007) 1025e1028. € ge, A. Jaggi, B. Sarafimov, [18] J. Raabe, G. Tzvetkov, U. Flechsig, M. Bo M.G.C. Vernooij, T. Huthwelker, H. Ade, D. Kilcoyne, T. Tyliszczak, R.H. Fink, C. Quitmann, PolLux: a new facility for soft x-ray spectromicroscopy at the Swiss Light Source, Rev. Sci. Instrum. 79 (2008) 113704. € ge, R. Fink, H. Ade, The PolLux [19] U. Flechsig, C. Quitmann, J. Raabe, M. Bo microspectroscopy beam line at the Swiss light source, AIP Conf. Proc. 879 (2007) 505e508. [20] U. Frommherz, J. Raabe, B. Watts, R. Stefani, U. Ellenberger, Higher Order Suppressor (HOS) for the PolLux microspectroscope beamline at the Swiss light source SLS, AIP Conf. Proc. 1234 (2010) 429e432. [21] A.P. Hitchcock, aXis 2000danalysis of X-ray Images and Spectra, 2000. http:// unicorn.mcmaster.ca/aXis2000.html. [22] J. Perlich, J. Rubeck, S. Botta, R. Gehrke, S.V. Roth, M.A. Ruderer, S.M. Prams, € rstgens, P. Müller-Buschbaum, Grazing incidence M. Rawolle, Q. Zhong, V. Ko wide angle x-ray scattering at the wiggler beamline BW4 of HASYLAB, Rev. Sci. Instrum. 81 (2010) 105105. [23] M. Dante, J. Peet, T.-Q. Nguyen, Nanoscale charge transport and internal structure of bulk heterojunction conjugated polymer/fullerene solar cells by scanning probe microscopy, J. Phys. Chem. C 112 (2008) 7241e7249. [24] O.G. Reid, K. Munechika, D.S. Ginger, Space charge limited current measurements on conjugated polymer films using conductive atomic force microscopy, Nano Lett. 8 (2008) 1602e1609. €hler, J.S. Wilson, J.-S. Kim, C.M. Ramsdale, [25] P.J. Brown, D.S. Thomas, A. Ko H. Sirringhaus, R.H. Friend, Effect of interchain interactions on the absorption and emission of poly(3-hexylthiophene), Phys. Rev. B 67 (2003) 064203. [26] C. Duan, W. Cai, B.B.Y. Hsu, C. Zhong, K. Zhang, C. Liu, Z. Hu, F. Huang, G.C. Bazan, A.J. Heeger, Y. Cao, Toward green solvent processable photovoltaic materials for polymer solar cells: the role of highly polar pendant groups in charge carrier transport and photovoltaic behavior, Energy Environ. Sci. 6 (2013) 3022e3034. [27] P. Vanlaeke, A. Swinnen, I. Haeldermans, G. Vanhoyland, T. Aernouts, D. Cheyns, C. Deibel, J. D'Haen, P. Heremans, J. Poortmans, J.V. Manca, P3HT/ PCBM bulk heterojunction solar cells: relation between morphology and electro-optical characteristics, Sol. Energy Mater. Sol. Cells 90 (2006) 2150e2158. [28] M. Boltau, S. Walheim, J. Mlynek, G. Krausch, U. Steiner, Surface-induced structure formation of polymer blends on patterned substrates, Nature 391 (1998) 877e879. [29] S.Y. Heriot, R.A.L. Jones, An interfacial instability in a transient wetting layer leads to lateral phase separation in thin spin-cast polymer-blend films, Nat. Mater 4 (2005) 782e786.
78
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
[30] T. Wang, A.J. Pearson, A.D.F. Dunbar, P.A. Staniec, D.C. Watters, H. Yi, A.J. Ryan, R.A.L. Jones, A. Iraqi, D.G. Lidzey, Correlating structure with function in thermally annealed PCDTBT: PC70BM photovoltaic blends, Adv. Funct. Mater. 22 (2012) 1399e1408. [31] H. Ade, A.P. Hitchcock, NEXAFS microscopy and resonant scattering: composition and orientation probed in real and reciprocal space, Polymer 49 (2008) 643e675. [32] H. Ade, H. Stoll, Near-edge X-ray absorption fine-structure microscopy of organic and magnetic materials, Nat. Mater. 8 (2009) 281e290. [33] B. Watts, C.R. McNeill, J. Raabe, Imaging nanostructures in organic semiconductor films with scanning transmission X-ray spectro-microscopy, Synth. Met. 161 (2012) 2516e2520. [34] I.N. Koprinarov, A.P. Hitchcock, C.T. McCrory, R.F. Childs, Quantitative mapping of structured polymeric systems using singular value decomposition analysis of soft X-ray images, J. Phys. Chem. B 106 (2002) 5358e5364. [35] X. Lu, H. Hlaing, D.S. Germack, J. Peet, W.H. Jo, D. Andrienko, K. Kremer, B.M. Ocko, Bilayer order in a polycarbazole-conjugated polymer, Nat. Commun. 3 (2012) 795. [36] Z. Wu, A. Petzold, T. Henze, T. Thurn-Albrecht, R.H. Lohwasser, M. Sommer, M. Thelakkat, Temperature and molecular weight dependent hierarchical equilibrium structures in semiconducting poly(3-hexylthiophene), Macromolecules 43 (2010) 4646e4653.
[37] P. Kohn, Z. Rong, K.H. Scherer, A. Sepe, M. Sommer, P. Müller-Buschbaum, R.H. Friend, U. Steiner, S. Hüttner, Crystallization-induced 10-nm structure formation in P3HT/PCBM blends, Macromolecules 46 (2013) 4002e4013. [38] D.A. Kamkar, M. Wang, F. Wudl, T.-Q. Nguyen, Single nanowire OPV properties of a fullerene-capped P3HT Dyad investigated using conductive and photoconductive AFM, ACS Nano 6 (2012) 1149e1157. [39] J.M. Mativetsky, Y.-L. Loo, P. Samori, Elucidating the nanoscale origins of organic electronic function by conductive atomic force microscopy, J. Mater. Chem. C 2 (2014) 3118e3128. [40] P.W.M. Blom, M.J.M. de Jong, M.G. van Munster, Electric-field and temperature dependence of the hole mobility in poly(p-phenylene vinylene), Phys. Rev. B 55 (1997) R656eR659. [41] L. Bozano, S.A. Carter, J.C. Scott, G.G. Malliaras, P.J. Brock, Temperature- and field-dependent electron and hole mobilities in polymer light-emitting diodes, Appl. Phys. Lett. 74 (1999) 1132e1134. chet, Molecular[42] C. Goh, R.J. Kline, M.D. McGehee, E.N. Kadnikova, J.M.J. Fre weight-dependent mobilities in regioregular poly(3-hexyl-thiophene) diodes, Appl. Phys. Lett. 86 (2005) 122110. [43] I.H. Campbell, D.L. Smith, C.J. Neef, J.P. Ferraris, Consistent time-of-flight mobility measurements and polymer light-emitting diode currentevoltage characteristics, Appl. Phys. Lett. 74 (1999) 2809e2811.
Contents lists available at ScienceDirect
Polymer journal homepage: www.elsevier.com/locate/polymer
Competition between phase separation and structure confinement in P3HT/PCDTBT heterojunctions: Influence on nanoscale charge transport Alvaro Rodríguez-Rodríguez a, Michelina Soccio b, Daniel E. Martínez-Tong a, 1, rrez a, * Tiberio A. Ezquerra a, Benjamin Watts c, Mari-Cruz García-Gutie a b c
Instituto de Estructura de la Materia (IEM-CSIC), Serrano 121, 28006 Madrid, Spain di Bologna, via Terracini 28, 40131 Bologna, Italy Dipartimento di Ingegneria Civile, Chimica, Ambientale e dei Materiali, DICAM-Universita Swiss Light Source, Paul Scherrer Institut, 5232 Villigen, Switzerland
a r t i c l e i n f o
a b s t r a c t
Article history: Received 6 July 2015 Received in revised form 25 August 2015 Accepted 4 September 2015 Available online 9 September 2015
A strong impact of crystal morphology on hole mobility is evidenced due to competition between phase separation and structure confinement in P3HT/PCDTBT heterojunctions. We find that domain sizes of both components decrease as film thickness decreases, suggesting an initial advantageous scenario for the efficiency of bulk heterojunction solar cells. However the P3HT/PCDTBT (1:1) films with a thickness of 165 nm and thicker present a dense crystal needle-like morphology while no evidence of needle-like motifs appears in the films 58 nm thick, indicating that confinement inhibiting crystallization takes place due to the very thin P3HT domains of only 40 nm in the thinner sample. The structural studies were correlated with nanoscale charge transport by conductive-AFM. A significant zero-field hole mobility increase is observed for the P3HT/PCDTBT films with increasing thickness. Even more interesting is the fact that a zero-field hole mobility increase of about two orders of magnitude is observed for P3HT domains in a 165 nm blend film compared to a neat P3HT film with similar thickness. This observation is related to the highly conductive network induced in P3HT by the presence of the PCDTBT phase, consisting of needle-like crystals growing from the P3HT domains and acting as bridges through the PCDTBT domains. © 2015 Elsevier Ltd. All rights reserved.
Keywords: Confinement All-polymer heterojunctions Charge transport
1. Introduction Organic photovoltaics (OPVs) are one class of solar-energy conversion devices, offering the advantages of low cost, lightweight, solution-processability and mechanical flexibility over existing photovoltaic technologies [1e3]. Polymer solar cells (PSCs) are attractive due to a number of advantageous features [2], including their thin-film architecture and low material consumption resulting from a high absorption coefficient, their use of organic solution processes and low manufacturing energy requirements. Efficient PSCs typically employ a bulk-heterojunction (BHJ) photoactive layer, where an electron-donating (p-type)
* Corresponding author. rrez). E-mail address: [email protected] (M.-C. García-Gutie 1 partement de Physique, Faculte des Sciences, Universite Present address: De libre de Bruxelles (ULB), Boulevard du Triomphe, 1050 Brussels, Belgium. http://dx.doi.org/10.1016/j.polymer.2015.09.012 0032-3861/© 2015 Elsevier Ltd. All rights reserved.
material and an electron-accepting (n-type) material form a nanosized phase-separated interpenetrating network, which can provide large enough heterointerface areas for efficient exciton dissociation and a continuous pathway for charge transport [4]. Fullerene derivatives are the most effective and commonly used acceptor materials. Polymer/fullerene BHJ solar cells have proven power conversion efficiencies of over 10% now reported [5e7]. Allpolymer solar cells (all-PSCs) have been also developed, where an n-type semiconducting polymer is used as the electron acceptor instead of a fullerene derivative blended with a p-type polymer. Although the efficiency of all-PSCs remains at ~4% [8,9], they offer some potential advantages including superior optical absorption and greater synthetic flexibility in tuning semiconducting properties such as the bandgap and energy levels [10,11]. The bulk heterojunction concept relies on the high purity of donor and acceptor phases within the characteristic exciton diffusion length of ~10 nm, requiring percolating interconnected network morphology. The understanding of phase separation and
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
morphology in all-polymer BHJ is crucial for the optimization of device performance [11,12]. In general, a mixing of donor and acceptor on a length scale smaller than the exciton diffusion length optimizes charge generation, while coarser morphologies optimize the separation of charges away from the interface and their collection at the device electrodes. On the other hand, crystal morphology [13,14] and structure confinement effects [15e17] such as crystallization inhibition in domains of tens of nanometers also play very important roles on charge transport and device performance. In this paper, we present the morphology developed as the result of competition between phase separation and structure confinement in all-polymer bulk-heterojunctions based on the blends of poly-3-hexylthiophene (P3HT) and poly[N-90 -heptadecanyl-2,7-carbazole-alt-5,5-(40 ,70 -di-2-thienyl-20 ,10,30 benzothiadiazole)] (PCDTBT). The influence of morphology on nanoscale charge transport is also reported.
2. Experimental section 2.1. Materials Poly[N-90 -heptadecanyl-2,7-carbazole-alt-5,5-(40 ,70 -di-2 thienyl-20 ,10,30 benzothiadiazole)] (PCDTBT) and poly-3hexylthiophene (P3HT) were purchased from Ossila. PCDTBT (Mw ¼ 35400 g/mol, PDI ¼ 2.4), P3HT (Mw ¼ 34100 g/mol, PDI ¼ 1.7; regioregularity ¼ 94.7%). The chemical structures of P3HT and PCDTBT are shown in Fig. 1. The HOMO and LUMO energy levels of P3HT and PCDTBT with respect to the vacuum level as supplied by Ossila are (HOMO: 4.7 eV; LUMO: 2.7 eV) and (HOMO: 5.35 eV; LUMO: 3.5 eV) respectively.
2.2. Sample preparation P3HT and PCDTBT were dissolved separately in chlorobenzene (CB) and afterwards blended with a weight ratio of 1:1. Thin films of pristine polymers and of the blend were prepared by spin-coating on silicon wafers for investigation with grazing incidence wide angle X-ray scattering techniques (GIWAXS) and on Arsenic ndoped Silicon substrates (resistivity ~ 0.001 Ucm) for morphology and electrical currentevoltage (IeV) characterization by atomic force microscopy (AFM) in tapping mode and by conducting atomic force microscopy (C-AFM) respectively. A fixed amount of 0.2 mL of polymer solution was dropped by a syringe on a square silicon substrate placed in the center of a rotating metallic horizontal plate. A rotation rate of 2400 rpm was kept during 120 s. For scanning transmission X-ray spectro-microscopy (STXM) measurements thin films spin-coated on silicon substrates were then floated off into a very dilute NaOH/water solution (0.25 wt%) and finally picked up with Transmission Electron Microscopy (TEM) grids. Solutions with concentrations of 12, 24 and 36 mg/mL were used in order to obtain films with different thicknesses (Table 1).
71
2.3. Ultravioletevisible spectroscopy Absorption spectra of polymer thin films spin-coated on silicon substrates were recorded in reflection geometry using a UVevisible spectrophotometer (UV-3600, Shimadzu). 2.4. X-ray microscopy Scanning Transmission X-ray Spectro-Microscopy (STXM), with image contrast based on Near-Edge X-ray Absorption Fine Structure (NEXAFS) spectromicroscopy, was performed at the PolLux beamline [18e20] at the Swiss Light Source, Paul Scherrer Institute, Villigen, Switzerland. TEM grids-supported films were mounted in the sample chamber which was evacuated to low vacuum. The transmitted X-ray intensity through the film was recorded using a scintillator and photo-multiplier tube and measured as a function of energy (275.0e345.0 eV with a resolution of 0.1 eV) and position (with X-ray focus of about 25 nm diameter and XeY sample position of 1 nm). Transmitted X-ray intensity was converted to an Xray optical density (defined as OD ¼ ln (I/I0)) by recording the Xray intensity through an empty TEM grid. Analysis of images was performed by the aXis2000 software package [21]. 2.5. Grazing incidence X-ray scattering Information of the inner film structure was obtained by grazing incidence X-ray scattering experiments performed at the synchrotron beamline BW4 at HASYLAB (DORIS, DESY) in Hamburg. The experimental setup has been previously described [22]. An Xray wavelength of l ¼ 0.13808 nm, with a beam size (horizontal vertical) of 40 20 mm2 and an exposure time of 4 min was used in our measurements. Scattered intensity was recorded by a MarCCD detector of 2048 2048 pixels with a resolution of 79.1 mm per pixel. A sample-detector distance of 10.6 cm was used to investigate crystallinity and orientation by GIWAXS. The incident angle (ai) was set to 0.40 , which was above the critical angles of P3HT and PCDTBT and therefore the structure of the full film thickness was detected. 2.6. Scanning probe measurements AFM measurements were done in air under ambient conditions using a commercial scanning probe microscope (MultiMode 8 equipped with a conductive-AFM (C-AFM) module and the Nanoscope V controller, Bruker). The topography AFM images were collected in tapping mode using silicon probes (NSG30 probes by NT-MDT); topographic images were collected from multiple locations to examine the film uniformity. For C-AFM measurements, PtIr-coated silicon tips with a spring constant of 0.2 N/m and a tip radius of ca. 25 nm (SCM-PIC by Bruker) were used. The contact area was estimated using the Hertz model [23,24]. In these measurements, the conducting probe makes contact with the sample (the tip acts as a nanoelectrode) and either measures the current as a function of the applied voltage at certain points on a surface (IeV curve) or maps out a current image at a fixed bias. A scheme with the C-AFM measurement set-up is shown in Fig. 2. The bias was applied to the conducting substrate, and the current was measured by a preamplifier. For each sample, several IeV curves were collected at various locations to obtain an average IeV curve. 3. Results and discussion 3.1. Optical properties
Fig. 1. Chemical structures of P3HT (left) and PCDTBT (right).
The normalized ultravioletevisible (UVevisible) absorption
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
72
Table 1 Description of the investigated samples: Solution concentration, average film thickness measured by AFM (for the P3HT/PCDTBT blends the domain thickness of each component is indicated) and zero-field hole mobility. Sample
Concentration (mg/mL)
Thickness (nm)
m0 (cm2 V1 s1)
P3HT PCDTBT P3HT/PCDTBT_1 P3HT/PCDTBT_2 P3HT/PCDTBT_3
24 24 12 24 36
145 190 58 (40P3HT-75PCDTBT) 165 (140P3HT-190PCDTBT) 295 (270P3HT-320PCDTBT)
2.3 102 e 1.2 105 0.61 1.21
A
V
P3HT:PCDTBT n Si Metallic support Fig. 2. C-AFM experimental setup.
spectra of pristine P3HT, PCDTBT and the blend P3HT/PCDTBT_2 films spin-cast from chlorobenzene (see Table 1), are compared in Fig. 3. P3HT features a main absorption band (1P3HT) with the maximum at about 524 nm, related to the electronic pep* transition [25]. PCDTBT exhibits two pronounced absorption bands: the 1PCDTBT at approximately 385 nm attributable to the pep* transition of their conjugated main chains and the 2PCDTBT at about 563 nm attributable to the intramolecular charge transfer (ICT) interaction between carbazole moieties and benzothiadiazole moieties [26]. Both, P3HT and PCDTBT show vibronic structure, suggesting good light-harvesting ability. Fig. 3 also shows internal order in the blend, as evidenced by the red shift of its absorption spectrum with respect to the neat absorption spectra and appearance of vibronic structure [27].
3.2. Surface morphology AFM in tapping mode was used to characterize the surface topography at the nanoscale and the thickness of P3HT/PCDTBT (1:1) films prepared from different solution concentrations. The surface topography images of as-cast films are shown in Fig. 4
Absorbance (a.u.)
1,0
1PCDTBT
1P3HT
P3HT PCDTBT P3HT/PCDTBT
0,8
0,6
0,4
2PCDTBT
0,2
0,0 300
400
500
600
700
Wavelength (nm) Fig. 3. Normalized UVevis absorption spectra of thin films of P3HT (blue), PCDTBT (red) and the blend P3HT/PCDTBT_2 (green). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
(except for the P3HT/PCDTBT_3 sample that is shown in Fig. S1 of the Supplementary Information). Neat P3HT and PCDTBT thin films are included for comparison. The topography of the P3HT (Fig. 4a) and PCDTBT (Fig. 4b) films is smooth and very small fiberlike crystallites can be observed for the P3HT (inset in Fig. 4a). P3HT/PCDTBT (1:1) films show the characteristic morphology of lateral phase separation, indicating immiscibility of the two polymers [28,29]. P3HT/PCDTBT_2 and P3HT/PCDTBT_3 thin films present a 50 nm difference in height between phases for an average film thickness of 165 and 295 nm respectively, the P3HT/PCDTBT_1 thin film shows 35 nm difference in height between phases for an average film thickness of 58 nm (see Table 1). It is shown that domain size as well as film thickness increase as increasing polymer concentration (Figs. 4c, d and S1a). It is also interesting to point out the presence of a dense needle-like morphology in films with 165 and 295 nm average thicknesses (insets in Figs. 4d and S1a respectively) while no evidence of needle-like motifs appears in the thin film with 58 nm average thickness (inset in Fig. 4c). The samples were then annealed at 140 C for 20 min (above the glass transition temperature of both polymers [30]) and cooled to room temperature aiding the thermodynamic equilibration of the samples, but no change in morphology was observed. 3.3. Composition mapping and internal structure Since there is not a single technique able to fully characterize the morphology of a thin-film polymer blend, the use of complementary techniques is required. Although AFM enables imaging of surface topography with high spatial resolution, to distinguish regions with different chemical composition can be cumbersome. Scanning Transmission X-ray Spectro-Microscopy (STXM) is a technique that provides an appropriate combination of spectroscopy and imaging capabilities at low radiation doses. As a powerful synchrotron-based technique, STXM can quantitatively map the chemical composition of organic films with nanoscale resolution [31,32]. The STXM images can provide information on the domain sizes, shapes, and purities of thin film polymer blends with a spatial resolution of ~20 nm [33]. In order to obtain quantitative chemical composition of the blend thin films we have used the singular value decomposition (SVD) mathematical procedure [31,34]. SVD requires previous knowledge of the mass absorption coefficients of the blend components. Fig. 5 shows the mass absorption coefficients as a function of energy for P3HT and PCDTBT pristine films of known thickness (Table 1). The transmitted X-ray intensity measured was converted to an X-ray optical density (defined as OD ¼ ln(I/I0)) where I is the transmitted intensity and I0 is the incident intensity. Since OD ¼ mrt, where m is the mass absorption coefficient, r is the density and t the sample thickness, we have obtained the mass absorption coefficient of each homopolymer considering density values of 1.13 and 1.33 g/cm3 for P3HT and PCDTBT respectively [35]. The SVD procedure also requires a series of images acquired at a number of energies that equal or exceed the number of compositional components. STXM images (Fig. 5 top) of the same area of
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
73
Fig. 4. Tapping mode AFM height images of (a) P3HT, (b) PCDTBT, (c) P3HT/PCDTBT_1 and (d) P3HT/PCDTBT_2.
the P3HT/PCDTBT_2 thin film were taken at energies indicated by arrows in the NEXAFS spectra which correspond to 280 eV (preedge), 284.2 eV (PCDTBT resonance), 287.8 eV (P3HT resonance) and 320 eV (chemically insensitive). By using the aXis2000 software package [21] where the SVD procedure is integrated, the information contained within the set of STXM images taken at the
284.2 eV
320 eV
287.8 eV
4
5x10
2
mass absorption coefficient (cm /g)
280 eV
selected energies can be transformed into maps quantifying the composition and thickness of the sample in each pixel [31,34]. The corresponding composition and thickness maps of the P3HT/ PCDTBT_2 thin film are shown in Fig. 6. On the one hand, these images reveal the chemical nature of the different observed
4
4x10
4
3x10
283
286
289
Energy (eV)
4
2x10
4
1x10
275
285
295
305
315
325
335
345
Energy (eV)
Fig. 5. Reference NEXAFS spectra for pristine films of P3HT (red line) and PCDTBT (blue line). Vertical arrows indicate the photon energies used to obtain the raw STXM images (6 mm 6 mm) of the same sample area of a P3HT/PCDTBT_2 thin film, presented at the top. The inset highlights differences in the NEXAFS spectra of P3HT and PCDTBT in the C-1s / p* region. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 6. 6 mm 6 mm images of quantitative composition (a) and thickness (b) maps of the P3HT/PCDTBT_2 thin film, calculated from the set of raw images in Fig. 5 top.
74
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
domains. On the other hand they reveal a structure of thinner P3HT-rich domains enclosed by a thicker PCDTBT-rich matrix. The thicknesses of domains are in agreement with the values measured by AFM. In order to complement our STXM studies, we have used GIWAXS in order to probe changes in molecular and nanoscale structure in P3HT, PCDTBT and in the blend P3HT/PCDTBT_2 thin films, before and after annealing at 140 C for 20 min. Fig. 7a shows the 2D scattering pattern for the P3HT as-cast film. Four main peaks can be seen, three out of plane reflections h00 having q-values of 3.8, 7.59 and 11.55 nm1, and one in plane peak superposition of the 020 and the weak 002 reflections [36], with a q-value of 16.39 nm1. In agreement with previous reports [36,37], the pattern reveals the characteristic crystalline reflections corresponding to the crystal sheets formed by the pep stacking of the thiophene rings h00 and the interchain distances 0k0. In addition, the P3HT as-cast film is uniaxially oriented with mainly an edge-on configuration which corresponds to the usual conformation adopted by P3HT thin films, consisting of polymer chains parallel to the substrate. The 2D scattering pattern of PCDTBT as-cast film, shown in Fig. 7b, reveals a lower degree of crystallinity compared to P3HT. The broad peak at q ¼ 14.55 nm1 in the out-of-plane direction, indexed as 010, corresponds to pep stacking between PCDTBT back-bones while the diffraction peak at q ¼ 2.95 nm1 also in the out of plane direction, indexed as 100, corresponds to the alkyl side-chain packing [14,30,35]. The blend P3HT/PCDTBT_2 as-cast thin film presents a 2D scattering pattern (Fig. 7c) very similar to P3HT, where only the P3HT reflections are detected. These
reflections appear at similar q values as those of P3HT. No evidence of PCDTBT reflections is observed. These data reveal that P3HT and PCDTBT do not co-crystallize. After annealing the samples at 140 C for 20 min, the P3HT thin film exhibits more intense reflections (Fig. 7d) suggesting an increase in crystallinity and a similar trend is observed for the blend (Fig. 7f). However, for PCDTBT Fig. 7e shows that while the 100 peak is better defined, the 010 loses intensity after annealing. This effect is even more evident for a thinner PCDTBT sample, as can be seen in Fig. S2. This trend can be explained by considering that annealing improves side-chain order while reducing the pep stacking, as has been previously proposed by Wang et al. [30]. 3.4. Nanoscale charge transport properties Conductive-AFM (C-AFM) was used to characterize the electrical properties at the nanoscale of P3HT/PCDTBT (1:1) blend thin films with different thicknesses. Neat P3HT films were included in the study for the sake of comparison. A scheme with the C-AFM measurement set-up is shown in Fig. 2. Hole current imaging was used to examine conductivity variations and to map the P3HT conducting network in the P3HT/PCDTBT (1:1) thin films. In these measurements, PtIr-coated silicon probes were used to ensure that holes are the major injected carriers [23]. Fig. 8 shows contact mode AFM topography and current images recorded simultaneously by applying a constant voltage of 5 V on the conducting substrate. The current image of the P3HT thin film (Fig. 8d) shows a little current variation throughout the film. However, the P3HT/
Fig. 7. 2D GIWAXS images of as-cast thin films: a) P3HT, b) PCDTBT, c) P3HT/PCDTBT_2. And after annealing at 140 C for 20 min: d) P3HT, e) PCDTBT, f) P3HT/PCDTBT_2.
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
75
Fig. 8. Contact mode AFM topography (left) and current (right) images of P3HT (a, d), P3HT/PCDTBT_1 (b, e) and P3HT/PCDTBT_2 (c, f). Blue and red crosses indicate the positions where IeV curves presented in Fig. 9a were measured. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
PCDTBT (1:1) films show domains with high (dark region) and low (bright region) current (Fig. 8e and f). By comparison with the compositional maps (Fig. 6), we can assign the high current regions to the P3HT-rich domains and the low current ones to PCDTBT. It is known that the thickness of thin films prepared by spin-coating is mainly related to the solution concentration. Table 1 shows the thicknesses of neat P3HT and PCDTBT thin films and the average thickness as well as the thicknesses of P3HT-rich and PCDTBT-rich domains of P3HT/PCDTBT (1:1) thin films prepared from solutions with different concentrations. We observe in Fig. 8 that the current in the P3HT-rich domains increases enormously with increasing thickness, from hundreds of pA for the blend with a P3HT-rich domain thickness of 40 nm (Fig. 8e) to several nA for blends with P3HT-rich domains thicker than 100 nm (Figs. 8f and S1c). In addition, we pointed out previously the presence of a dense needlelike morphology in P3HT/PCDTBT_2 and P3HT/PCDTBT_3 films with 165 and 295 nm thickness respectively (Figs. 4d and S1a) while no evidence of needle-like motives appears in the 58 nm thick P3HT/PCDTBT_1 film (Fig. 4c). The current images of P3HT/
PCDTBT_2 and P3HT/PCDTBT_3 films (Figs. 8f and S1c), clearly show a fibrous network where the strongest current is measured. These results evidence that P3HT needle-like crystals grow from the P3HT-rich domains acting as bridges through the PCDTBT-rich domains. However, the current image of the P3HT/PCDTBT_1 thin film (Fig. 8e) does not show any fibrous network nor evidence of needle-like motifs (Fig. 4c). A fact that can be explained by confinement inhibiting crystallization due to the very thin domains of P3HT of only 40 nm as it is supported by GIWAXS measurements (Fig. S3). In addition to current imaging with C-AFM, local IeV curves can also be collected. This technique of measuring IeV characteristics by C-AFM has the advantage that only a very small region of the film, comparable to the tip contact area, is probed [23,38,39]. Thus, several areas can be examined aiming to provide information on the film electrical heterogeneity. Several IeV curves were collected from different locations within a P3HT-rich domain (red crosses in Fig. 8f) and from different locations within a PCDTBT-rich domain (blue crosses in Fig. 8f). IeV curves measured in both regions
76
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
including the average curves (red and blue for the P3HT rich domain and the PCDTBT rich domain respectively) are shown in Fig. 9a. As one can see, while there is a negligible electrical current in the PCDTBT-rich domains, a significant electrical current is measured in the P3HT-rich domains. The same procedure was used for all the samples investigated. Average IeV curves for the neat P3HT, neat PCDTBT and P3HT-rich domains in P3HT/PCDTBT (1:1)
films with different thicknesses are shown in Fig. 9b. The highest current for a similarly applied bias is measured for the P3HT/ PCDTBT_2 thin film. Since the samples have different thicknesses, Fig. 9c presents the I-E curves where E ¼ V⁄L is the electric field and L is the P3HT domain thickness. The logelog plot of the jJj jVj data, where J is the current density, is presented in the inset of Fig. 9b. It shows that for the investigated samples the electrical response exhibits two regimes: initially a short Ohmic range (JfV) and a more extended regime (JfV2) denominated space-charge limited current (SCLC) regime. In this case (SCLC), the nanoscale charge mobility can be extracted by using the following expression [40,41]:
J¼
9 V2 εr ε0 m 3 8 L
(1)
where εr is the dielectric constant of the polymer for which we assumed a value of 3 as proposed in the literature [42], ε0 ¼ 8.85 1012 F/m is the vacuum permittivity, m is the charge mobility, V is the applied voltage, and L is the P3HT domain thickness. In order to verify whether the calculated carrier mobilities are field dependent we have used the approach for a field-dependent carrier mobility described by the Poole-Frenkel equation [40,41,43]:
m ¼ m0 eðE=E0 Þ
1=2
(2)
where E is the electric field, m0 is the zero-field mobility, and E0 is the field coefficient. The Poole-Frenkel plots (lnm vs. E1/2) are shown in Fig. 10. All the samples exhibit similar behavior showing that the mobility depends weakly on the electric field. A significant zero-field hole mobility increase is observed for the P3HT/PCDTBT (1:1) films with increasing thickness (Table 1): from 1.2 105 cm2 V1 s1 for the P3HT/PCDTBT_1 thin film to 0.61 cm2 V1 s1 and 1.21 cm2 V1 s1 for the P3HT/PCDTBT_2 and P3HT/PCDTBT_3 thin films respectively. It is also interesting to note that a zero-field hole mobility increase of about two orders of magnitude is observed for the P3HT rich domains in the P3HT/ PCDTBT_2 thin film compared to the neat P3HT sample with similar thickness. It indicates a P3HT and PCDTBT synergy in the blend, probably related to the highly conductive needle like morphology induced in P3HT domains by the presence of the PCDTBT phase, as it is shown in Figs. 4d and 8f.
Fig. 9. a) IeV curves measured in different locations into a rich P3HT domain (red crosses in Fig. 8f) and in different locations into a rich PCDTBT domain (blue crosses in Fig. 8f). Average IeV curves, red and blue for the P3HT rich domain and the PCDTBT rich domain are also included. b) Average IeV curves for pristine P3HT and PCDTBT thin films and P3HT rich domains in P3HT/PCDTBT (1:1) thin films with different thicknesses. The inset shows the logelog plot of the jJjejVj data where two different regimes of the electrical response can be observed. c) Average I-E curves for pristine P3HT and PCDTBT thin films and P3HT rich domains in P3HT/PCDTBT (1:1) thin films with different thicknesses. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 10. Hole mobilities for pristine P3HT (black) and P3HT rich domains in P3HT/ PCDTBT_1 (green), P3HT/PCDTBT_2 (red) and P3HT/PCDTBT_3 (blue). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
4. Conclusions In summary, using several complementary techniques we have investigated the optical, structural and hole transport properties of the P3HT/PCDTBT (1:1) polymer blend films as a function of film thickness. It is seen that the optical absorption of the blend is improved compared to the neat P3HT and PCDTBT UVevis absorption spectra, covering a broader range of the solar spectrum. We find that the P3HT/PCDTBT is an immiscible blend with lateral phase separation when spin-coated. The domain sizes of both components decrease as film thickness decreases. This suggests an initial advantageous scenario for the efficiency of bulk heterojunction solar cells, but a competition between the favorable decrease of domain size and the resulting structure confinement takes place. The P3HT/PCDTBT (1:1) thin films with a thickness of 165 nm and greater present a dense crystal needle-like morphology while no evidence of needle-like motifs appears in the 58 nm thick film. In addition, conductive-AFM was used to characterize the electrical properties at the nanoscale. It clearly shows that blend films 165 and 295 nm thick present a fibrous network where the strongest current is measured. These results evidence that P3HT needle-like crystals grow from the P3HT-rich domains, acting as bridges through the PCDTBT-rich domains. However, the current image of the 58 nm thick blend film neither shows fibrous network, nor evidence of needle-like motifs. This fact can be explained by confinement inhibiting crystallization due to the very thin domains of P3HT of only 40 nm thickness. A strong impact of the crystal morphology on hole mobility is evidenced. A significant zero-field hole mobility increase is observed for the P3HT/PCDTBT (1:1) thin films with increasing thickness: from 1.2 105 cm2 V1 s1 for the 58 nm thick film to 0.61 cm2 V1 s1 and 1.21 cm2 V1 s1 for the 165 nm and a 295 nm thick films respectively. Even more interesting is the fact that a zero-field hole mobility increase of about two orders of magnitude is observed for the P3HT rich domains in the 165 nm P3HT/PCDTBT (1:1) thin film compared to the neat P3HT thin film with similar thickness, probably related to the highly conductive needle like network induced in P3HT by the presence of the PCDTBT phase. Acknowledgments The authors gratefully acknowledge the financial support of the Spanish Ministry of Economy and Competitiveness (MINECO) through the project MAT 2011-23455. A. R-R is indebted to MINECO for the tenure of a FPI BES-2013-062620 contract associated to the project MAT2012-33517. We thank the Swiss Light Source for beamtime at PolLux. The PolLux end station was financed by the German Minister für Bildung und Forschung (BMBF) through contracts 05KS4WE1/6 and 05KS7WE1. Parts of this research were carried out at the light source DORIS III at DESY, a member of the Helmholtz Association (HGF). We would like to thank J. Perlich for assistance in using beamline BW4. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.polymer.2015.09.012. References [1] S.R. Forrest, The path to ubiquitous and low-cost organic electronic appliances on plastic, Nature 428 (2004) 911e918. [2] G. Dennler, M.C. Scharber, C.J. Brabec, Polymer-fullerene bulkheterojunction solar cells, Adv. Mater. 21 (2009) 1323e1338. [3] Y. Tang, C.R. McNeill, All-polymer solar cells utilizing low band gap polymers as donor and acceptor, J. Polym. Sci. Pt. B Polym. Phys. 51 (2013) 403e409.
77
[4] G. Yu, J. Gao, J.C. Hummelen, F. Wudl, A.J. Heeger, Polymer photovoltaic cells [mdash] enhanced efficiencies via a network of internal donor-acceptor heterojunctions, Science 270 (1995) 1789e1791. [5] Z. He, C. Zhong, S. Su, M. Xu, H. Wu, Y. Cao, Enhanced power-conversion efficiency in polymer solar cells using an inverted device structure, Nat. Phot. 6 (2012) 591e595. [6] C.M. Amb, S. Chen, K.R. Graham, J. Subbiah, C.E. Small, F. So, J.R. Reynolds, Dithienogermole As a Fused Electron Donor in Bulk Heterojunction Solar Cells, J. Am. Chem. Soc. 133 (2011) 10062e10065. [7] C.E. Small, S. Chen, J. Subbiah, C.M. Amb, S.-W. Tsang, T.-H. Lai, J.R. Reynolds, F. So, High-efficiency inverted dithienogermole-thienopyrrolodione-based polymer solar cells, Nat. Phot. 6 (2012) 115e120. [8] E. Zhou, J. Cong, K. Hashimoto, K. Tajima, Control of Miscibility and Aggregation Via the Material Design and Coating Process for High-Performance Polymer Blend Solar Cells, Adv. Mater. 25 (2013) 6991e6996. [9] D. Mori, H. Benten, I. Okada, H. Ohkita, S. Ito, Low-bandgap donor/acceptor polymer blend solar cells with efficiency exceeding 4%, Adv. Energy Mater. 4 (2014) 1301006. [10] C.R. McNeill, N.C. Greenham, Conjugated-Polymer Blends for Optoelectronics, Adv. Mater. 21 (2009) 3840e3850. [11] A. Sepe, Z. Rong, M. Sommer, Y. Vaynzof, X. Sheng, P. Muller-Buschbaum, D.M. Smilgies, Z.-K. Tan, L. Yang, R.H. Friend, U. Steiner, S. Huttner, Structure formation in P3HT/F8TBT blends, Energy Environ. Sci. 7 (2014) 1725e1736. [12] C.R. McNeill, A. Abrusci, I. Hwang, M.A. Ruderer, P. Mueller-Buschbaum, N.C. Greenham, Photophysics and photocurrent generation in polythiophene/ polyfluorene copolymer blends, Adv. Funct. Mater. 19 (2009) 3103e3111. [13] F.S. Kim, G. Ren, S.A. Jenekhe, One-dimensional nanostructures of p-conjugated molecular systems: assembly, properties, and applications from photovoltaics, sensors, and nanophotonics to nanoelectronics, Chem. Mat. 23 (2010) 682e732. [14] H.-W. Wang, E. Pentzer, T. Emrick, T.P. Russell, Preparation of low band gap fibrillar structures by solvent-induced crystallization, ACS Macro Lett. 3 (2013) 30e34. [15] J.J. Hernandez, D.R. Rueda, M.C. Garcia-Gutierrez, A. Nogales, T.A. Ezquerra, M. Soccio, N. Lotti, A. Munari, Structure and morphology of thin films of linear aliphatic polyesters prepared by spin-coating, Langmuir 26 (2010) 10731e10737. [16] M.-C. Garcia-Gutierrez, A. Linares, I. Martin-Fabiani, J.J. Hernandez, M. Soccio, D.R. Rueda, T.A. Ezquerra, M. Reynolds, Understanding crystallization features of P(VDF-TrFE) copolymers under confinement to optimize ferroelectricity in nanostructures, Nanoscale 5 (2013) 6006e6012. rrez, R.J. Davies, [17] W. Pisula, M. Kastler, B. El Hamaoui, M.-C. García-Gutie C. Riekel, K. Müllen, Dendritic morphology in homeotropically aligned discotic films, ChemPhysChem 8 (2007) 1025e1028. € ge, A. Jaggi, B. Sarafimov, [18] J. Raabe, G. Tzvetkov, U. Flechsig, M. Bo M.G.C. Vernooij, T. Huthwelker, H. Ade, D. Kilcoyne, T. Tyliszczak, R.H. Fink, C. Quitmann, PolLux: a new facility for soft x-ray spectromicroscopy at the Swiss Light Source, Rev. Sci. Instrum. 79 (2008) 113704. € ge, R. Fink, H. Ade, The PolLux [19] U. Flechsig, C. Quitmann, J. Raabe, M. Bo microspectroscopy beam line at the Swiss light source, AIP Conf. Proc. 879 (2007) 505e508. [20] U. Frommherz, J. Raabe, B. Watts, R. Stefani, U. Ellenberger, Higher Order Suppressor (HOS) for the PolLux microspectroscope beamline at the Swiss light source SLS, AIP Conf. Proc. 1234 (2010) 429e432. [21] A.P. Hitchcock, aXis 2000danalysis of X-ray Images and Spectra, 2000. http:// unicorn.mcmaster.ca/aXis2000.html. [22] J. Perlich, J. Rubeck, S. Botta, R. Gehrke, S.V. Roth, M.A. Ruderer, S.M. Prams, € rstgens, P. Müller-Buschbaum, Grazing incidence M. Rawolle, Q. Zhong, V. Ko wide angle x-ray scattering at the wiggler beamline BW4 of HASYLAB, Rev. Sci. Instrum. 81 (2010) 105105. [23] M. Dante, J. Peet, T.-Q. Nguyen, Nanoscale charge transport and internal structure of bulk heterojunction conjugated polymer/fullerene solar cells by scanning probe microscopy, J. Phys. Chem. C 112 (2008) 7241e7249. [24] O.G. Reid, K. Munechika, D.S. Ginger, Space charge limited current measurements on conjugated polymer films using conductive atomic force microscopy, Nano Lett. 8 (2008) 1602e1609. €hler, J.S. Wilson, J.-S. Kim, C.M. Ramsdale, [25] P.J. Brown, D.S. Thomas, A. Ko H. Sirringhaus, R.H. Friend, Effect of interchain interactions on the absorption and emission of poly(3-hexylthiophene), Phys. Rev. B 67 (2003) 064203. [26] C. Duan, W. Cai, B.B.Y. Hsu, C. Zhong, K. Zhang, C. Liu, Z. Hu, F. Huang, G.C. Bazan, A.J. Heeger, Y. Cao, Toward green solvent processable photovoltaic materials for polymer solar cells: the role of highly polar pendant groups in charge carrier transport and photovoltaic behavior, Energy Environ. Sci. 6 (2013) 3022e3034. [27] P. Vanlaeke, A. Swinnen, I. Haeldermans, G. Vanhoyland, T. Aernouts, D. Cheyns, C. Deibel, J. D'Haen, P. Heremans, J. Poortmans, J.V. Manca, P3HT/ PCBM bulk heterojunction solar cells: relation between morphology and electro-optical characteristics, Sol. Energy Mater. Sol. Cells 90 (2006) 2150e2158. [28] M. Boltau, S. Walheim, J. Mlynek, G. Krausch, U. Steiner, Surface-induced structure formation of polymer blends on patterned substrates, Nature 391 (1998) 877e879. [29] S.Y. Heriot, R.A.L. Jones, An interfacial instability in a transient wetting layer leads to lateral phase separation in thin spin-cast polymer-blend films, Nat. Mater 4 (2005) 782e786.
78
Rodríguez-Rodríguez et al. / Polymer 77 (2015) 70e78 A.
[30] T. Wang, A.J. Pearson, A.D.F. Dunbar, P.A. Staniec, D.C. Watters, H. Yi, A.J. Ryan, R.A.L. Jones, A. Iraqi, D.G. Lidzey, Correlating structure with function in thermally annealed PCDTBT: PC70BM photovoltaic blends, Adv. Funct. Mater. 22 (2012) 1399e1408. [31] H. Ade, A.P. Hitchcock, NEXAFS microscopy and resonant scattering: composition and orientation probed in real and reciprocal space, Polymer 49 (2008) 643e675. [32] H. Ade, H. Stoll, Near-edge X-ray absorption fine-structure microscopy of organic and magnetic materials, Nat. Mater. 8 (2009) 281e290. [33] B. Watts, C.R. McNeill, J. Raabe, Imaging nanostructures in organic semiconductor films with scanning transmission X-ray spectro-microscopy, Synth. Met. 161 (2012) 2516e2520. [34] I.N. Koprinarov, A.P. Hitchcock, C.T. McCrory, R.F. Childs, Quantitative mapping of structured polymeric systems using singular value decomposition analysis of soft X-ray images, J. Phys. Chem. B 106 (2002) 5358e5364. [35] X. Lu, H. Hlaing, D.S. Germack, J. Peet, W.H. Jo, D. Andrienko, K. Kremer, B.M. Ocko, Bilayer order in a polycarbazole-conjugated polymer, Nat. Commun. 3 (2012) 795. [36] Z. Wu, A. Petzold, T. Henze, T. Thurn-Albrecht, R.H. Lohwasser, M. Sommer, M. Thelakkat, Temperature and molecular weight dependent hierarchical equilibrium structures in semiconducting poly(3-hexylthiophene), Macromolecules 43 (2010) 4646e4653.
[37] P. Kohn, Z. Rong, K.H. Scherer, A. Sepe, M. Sommer, P. Müller-Buschbaum, R.H. Friend, U. Steiner, S. Hüttner, Crystallization-induced 10-nm structure formation in P3HT/PCBM blends, Macromolecules 46 (2013) 4002e4013. [38] D.A. Kamkar, M. Wang, F. Wudl, T.-Q. Nguyen, Single nanowire OPV properties of a fullerene-capped P3HT Dyad investigated using conductive and photoconductive AFM, ACS Nano 6 (2012) 1149e1157. [39] J.M. Mativetsky, Y.-L. Loo, P. Samori, Elucidating the nanoscale origins of organic electronic function by conductive atomic force microscopy, J. Mater. Chem. C 2 (2014) 3118e3128. [40] P.W.M. Blom, M.J.M. de Jong, M.G. van Munster, Electric-field and temperature dependence of the hole mobility in poly(p-phenylene vinylene), Phys. Rev. B 55 (1997) R656eR659. [41] L. Bozano, S.A. Carter, J.C. Scott, G.G. Malliaras, P.J. Brock, Temperature- and field-dependent electron and hole mobilities in polymer light-emitting diodes, Appl. Phys. Lett. 74 (1999) 1132e1134. chet, Molecular[42] C. Goh, R.J. Kline, M.D. McGehee, E.N. Kadnikova, J.M.J. Fre weight-dependent mobilities in regioregular poly(3-hexyl-thiophene) diodes, Appl. Phys. Lett. 86 (2005) 122110. [43] I.H. Campbell, D.L. Smith, C.J. Neef, J.P. Ferraris, Consistent time-of-flight mobility measurements and polymer light-emitting diode currentevoltage characteristics, Appl. Phys. Lett. 74 (1999) 2809e2811.