(Conjugated) Polymer Thin Films - American Chemical Society

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Fabrication and Mechanical Characterization of Semi-Free-Standing (Conjugated) Polymer Thin Films Jaime Martín,* Miguel Muñoz, Mario Encinar, Montserrat Calleja, and Marisol Martín-González* Instituto de Microelectrónica de Madrid (IMM-CSIC), Calle de Isaac Newton 8, Tres Cantos, 28760 Madrid, Spain S Supporting Information *

ABSTRACT: Polymers undergo severe low-dimensionality effects when they are confined to ultrathin films since most of the structural and dynamical processes involving polymer molecules are correlated to length scales of the order of nanometers. However, the real influence of the size limitation over such processes is often hard to identify as it is masked by interfacial effects. We present the fabrication of a new type of nanostructure consisting of poly[[9-(1-octylnonyl)-9H-carbazole-2,7-diyl]-2,5-thiophenediyl-2,1,3-benzothiadiazole-4,7diyl-2,5-thiophenediyl] (PCDTBT) thin film that is held up exclusively over tips of poly(ether−ether−ketone) (PEEK) nanopillars. The fabrication method exploits the nonwetting behavior of PCDTBT onto an ordered PEEK nanopillar array when the mobility of the PCDTBT molecules is enhanced by a solvent annealing process. We use this new configuration to characterize the mechanical behavior of free-standing thin film regions, thus in the absence of underlaying substrate, by means of an atomic force microscope (AFM) setup. First, we study how the finite thickness and/or the presence of the underlying substrate influences the mechanical modulus of the material in the linear elastic regime. Moreover, we analyze deep indentations up to the rupture of the thin film, which allow for the measurement of important mechanical features of the nanoconfined polymer, such as its yield strain, the rupture strain, the bending rigidity, etc., which are impossible to investigate in thin films deposited on substrates.

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INTRODUCTION Polymers undergo severe low-dimensionality effects when they are confined to ultrathin films since most of the structural and dynamical processes involving polymer molecules (crystallization, molecular motions, chain conformation, phase-separation, etc.) are correlated to length scales of the order of nanometers.1−3 However, the real influence of the size limitation over such processes is often hard to identify and thus has become a matter of controversy among the scientific community.4,5 Pure confinement effects are considered to originate directly from the hindering of certain physical processes. These effects are due to the fact that the associated length scale of the physical property in question comes into conflict with the dimensions of the material itself. Nevertheless, these effects are usually masked by the impact of interfacial interactions between the polymer chains and the confining surface, i.e., substrates, pore walls, etc. Usually such interactions are of attractive nature, so polymer chains located at the interface get perturbed dynamically, conformationally, and structurally. In this way, strong interfacial phenomena have been observed on the glass transition temperature, 5,6 crystallization kinetics,7 orientation of crystals,8,9 fragility,10 mechanical properties,11 and a variety of dynamic processes from segmental12,13 to reptation3 dynamics of polymers. Moreover, from a technological point of view, most of the applications of thin films require the polymer to be in close contact with an attractive surface14,15 (commonly a substrate), © XXXX American Chemical Society

so the properties of such a material are influenced by both pure confinement effects and interfacial effects. Hence, the identification of both kinds of effects is an important issue from fundamental and technological perspectives. Furthermore, the presence of the underlying substrate directly prevents to carry out a range of experiments, such as large-strain mechanical tests or strain−stress measurements, from which important mechanical features of the nanomaterial might be obtained.16,17 The direct strategy to identify pure confinement effects and interfacial effects would be to avoid one of them, for instance, through the study of a free-standing thin film. However, the fabrication of free-standing polymer thin films is far from being a trivial matter since the traditional thin film preparation methods, namely spin-coating, dip-coating, drop-casting, and blade-coating, are in fact based on the presence of a substrate. In this context, Forrest et al. managed to develop the only experimental procedure used to date to achieve free-standing thin films.6,18 Since then, the procedure has been widely used by Napolitano and Wübbenhorst,19 McKenna et al.,16 DalnokiVeres et al.,17 Torkelson et al.,20 and others without significant improvements or variations. The method consists of spin coating a polymer solution onto a freshly cleaved substrate Received: August 21, 2013 Revised: October 8, 2013

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dx.doi.org/10.1021/la4032267 | Langmuir XXXX, XXX, XXX−XXX

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Figure 1. (a) Schematic illustration of the procedure to obtain semi-free-standing (SFS) polymer thin films. First nanoporous hard template of anodic aluminum oxide (AAO) is synthesized (b); then, the molten poly(ether−ether−ketone) (PEEK) is infiltrated into the nanopores. After etching the template, a PEEK nanopillar array is obtained (artificially colored image) (c).29 Afterward, the PEEK nanorod array is submerged into a solution of PCDTBT in 1,2-dichlorobenzene. Finally, the solution is dried at ambient conditions (room temperature and 1 atm pressure), and the sample was exposed to 1,2-dichlorobenzene vapor for a week. Thus, a PCDTBT thin film supported exclusively onto the tips of the PEEK nanopillars was obtained. (d) Artificially colored surface SEM image of the semi-free-standing PCDTBT (purple) supported onto the PEEK nanopillars (brownish).

(commonly mica) and then floating the film onto water. Next, one must pick up the film from water and transfer to the sample holder in which the analysis will later be performed. This procedure requires delicate handling and often leads to damage in the thin film or to an incorrect positioning (folding, stretching, etc.) of the film over the sample holder, which could modify its properties. Here, we present a new fabrication method which provides access to semi-free-standing (SFS) thin polymer films, which does not require any difficult manipulation process, as the formation of the nonsupported film takes place spontaneously. The obtained thin film is held up exclusively at some specific points, leaving the rest of the film free-standing. Therefore, on the one hand, it allows studying the thin polymer film in the absence of substrate−polymer interfacial effects and, on the other hand, performing those measurements which are impossible to carry out on films deposited onto a substrate. In this context, we show the characterization of the mechanical behavior of free-standing thin film regions performed by means of an atomic force microscope (AFM) setup. First, we study how the finite thickness and/or the presence of the underlying substrate influences the mechanical modulus of the material in the linear elastic regime. Next, we phenomenologically analyze deep indentations of the thin film up to its rupture, which

should give access to important mechanical features of the nanoconfined polymer, such as its yield strain, the rupture strain, the bending rigidity, etc. It is noteworthy that these features are impossible to investigate in supported thin films. Access to the exploration of the cited mechanical features is of relevance not only from the point of view of the mechanical properties of the material but also because such exploration might offer the possibility of discovering related fundamental aspects of nanoconfined polymers, like the glass transition phenomenon16 or the nature of the chain entanglement network.17 It is worth mentioning that the SFS thin polymer film might also give rise to novel optoelectronic nanodevices and sensors, which might take advantage of this new configuration for enhancing their performance. With this in mind, we selected the poly[[9-(1-octylnonyl)-9H-carbazole-2,7-diyl]-2,5-thiophenediyl-2,1,3-benzothiadiazole-4,7-diyl-2,5-thiophenediyl] (PCDTBT) as a model polymer of technological interest to integrate the SFS thin film. PCDTBT is known to be one of the best performing polymers since it presents a nearly perfect internal quantum efficiency,21 a photovoltaic power conversion efficiency of 7.2%,22,23 and a thermoelectric power factor of 19 μW m−1 K−2.24 B


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Figure 2. (a) SEM surface view of the semi-free-standing PCDTBT thin film. The base of the PEEK nanopillars can be appreciated through the holes of the PCDTBT film. (b, c) Lateral views of the SFS PCDTBT thin film. In (b) a small piece of the PCDTBT film onto the tip of a nanopillar can be seen (indicated with a white arrow), which has been used to measure its thickness (Figure S3).

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were randomly chosen in the supported sample and on the bulklike sample, whereas they were placed specifically on the film between pillars on the suspended film and on the film region located over the PEEK nanopillars. Typically 50 points were analyzed, and for each one, ten consecutive curves have been recorded at a rate of one per second. All the force−displacement experiments were carried out using the same Tap300-G probe with a spring constant of 27.8 N/m, calibrated following the Sader method,27 and 80 nm of tip radius (see Supporting Information section 2). By determining visually the contact points the force−displacement curves were transformed to force−indentation curves. Finally, the Young’s modulus was extracted by fitting the Hertz model for spherical indentors.28 On the other hand, deep indentation experiments at high loading forces (few μN) were carried out using either sharp (lower than 10 nm of tip radius) or blunted probes and only semiquantitatively analyzed.

EXPERIMENTAL SECTION

Fabrication of AAO Templates. Self-ordered AAO templates were prepared by a two-step electrochemical anodization of aluminum.25 First, 99.999% pure aluminum foils (Advent Research Materials, England) were cleaned and degreased by sonication in acetone, water, isopropanol, and ethanol. The foils were then electropolished in a perchloric acid/ethanol (1/3) solution under a constant voltage of 20 V. Afterward, the first anodization was carried out using an aqueous solution of phosphoric acid (1 wt %) and aluminum oxalate (0.01 M) as electrolyte. The potential, temperature, and length of the reaction were 205 V, 4.5 °C, and 6 h, respectively. The grown anodic layer was chemically etched with a mixture of phosphoric acid (7 wt %) and chromic oxide (1.8 wt %). Finally, the second anodization was performed in the same conditions for 5 min. Fabrication of PEEK Nanopillar Array. For the PEEK (Goodfellow, Ltd.) infiltration, the hydroxylated AAO pore walls were cleaned with different polarity solvents (water, ethanol, and acetone) and heat. Next, a piece of commercial PEEK was placed onto the surface of the AAO template at 390 °C for 25 min. The sample was cooled at 3 °C/min down to 325 °C and then at 1 °C/min down to 250 °C, in order to improve crystallinity, as the mechanical properties of the PEEK nanostructure are enhanced in direct correlation with increased crystallinity. All the temperatures were selected according to the melting (∼350 °C) and crystallization (∼300 °C) temperatures obtained from DSC (Supporting Information Figure S1a). To release the PEEK nanopillar array, AAO templates were selectively etched with NaOH 10 wt %. Fabrication of the Semi-Free-Standing PCDTBT Thin Film. The PCDTBT, supplied by Solaris Chem Inc., had a weight-average molecular weight (Mw) of 53 000 g/mol and a polydispersity index of 1.5. A 4 g/L solution of PCDTBT in 1,2-dichlorobenzene was prepared and filtered through 0.45 μm pore diameter PTFE filters. The PEEK nanopillar array was submerged into such solution for a day, and then the solution was dried at ambient conditions (room temperature and 1 atm pressure) on the surface of the PEEK nanopillar array. Once the solution was completely dried, the sample was exposed to 1,2-dichlorobenzene vapor for a week (vapor annealing). Fabrication of the Supported PCDTBT Thin Film. A drop of the same PCDTBT solution was casted onto a glass slide and dried at ambient conditions to form the PCDTBT thin film. Then, the supported films were annealed in 1,2-dichlorobenzene vapor for a week. Bulklike PCDTBT Sample. A 200 μm thick PCDTBT planar piece was used as received. Characterization. Scanning electron microscopy (SEM) analysis was carried out using a Hitachi SU8000 microscope, operating at an accelerating voltage of 1 kV. Atomic force microscopy experiments were conducted using a Nanotec Cervantes AFM system.26 The topography images were taken in dynamic mode using commercial sharpened silicon cantilevers (Tap300-G from BudgetSensors) with resonant frequencies of 300 kHz. A set of force−displacement measurements at low loading forces (20 nN) have been performed leading to shallow indentations (4 nm) of the sample. The locations

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RESULTS The SFS thin film fabrication method is schematically presented in Figure 1a. It begins with the synthesis of a nanoporous anodic aluminum oxide (AAO) template (Figure 1b). The AAO templates were synthesized by a two-step anodization method25,30 and exhibited a polydomain hexagonal array of cylindrical nanopores. The pore diameter was 120 nm, and the pore depth amounted to approximately 400 nm. Then, the nanopores were infiltrated with melted poly(ether−ether− ketone) (PEEK) to obtain an ordered array of nanopillars (Figure 1c).29 Figure 1c shows the obtained PEEK nanopillar array after it has been released from the AAO template (see also Figure S1b,c). The diameter of the PEEK nanopillars was 120 nm, while their length was approximately 400 nm, as confirmed from the height profile obtained by atomic force microscopy (AFM) (Figure S2a,b). The PEEK nanopillar array was then submerged into a 4 g/L solution of PCDTBT in 1,2dichlorobenzene. After the removal of the PEEK nanostructure from the solution and the evaporation of the solvent, the sample was exposed to 1,2-dichlorobenzene vapor for a week. Thus, a PCDTBT thin film was formed, which was exclusively supported onto the tips of the PEEK nanopillars (Figure 1d). Note that the configuration of this material is absolutely new. Figure 2 shows SEM surface view (a) and lateral views (b, c) of the obtained SFS PCDTBT thin film. As can be observed, the PEEK nanopillars hold up the PCDTBT film, leaving the film free-standing in the region between the anchoring points. The PCDTBT slightly falls around the pillars describing reliefs of approximately 150 nm in height (Figures 2b,c and Figure S2c). The thickness of the film was 40 nm, as measured from the small film piece shown in Figure 2b (marked with a white arrow) (see with higher detail in Figure S3). The formation of the SFS thin film lays on the fact that PCDTBT does not wet the PEEK surface. Therefore, when the mobility of the PCDTBT molecules is enhanced during the vapor annealing, C


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substrate. For the latter, we measure the SFS PCDTBT thin film region directly lying on the tip of PEEK nanopillars (see Figure 3a). PEEK is a rigid substrate for PCDTBT (EPEEK ≈ 4000 MPa), but it presents a much lower surface energy than glass (γPEEK = 42 mN/m, γglass > 1000 mN/m), and therefore, it can be considered as a nonattractive rigid surface for PCDTBT. The fact that the PCDTBT does not wet the entire PEEK surface during the solvent annealing process further suggests this assumption. Finally, a 200 μm thick bulk PCDTBT piece was also measured. The Young’s moduli of the different points and samples were derived analyzing the force curves by means of a Hertz model28 (eq 1), which is a reliable description of the results given the good fitting quality (see Figure S6).

the contact area of the PCDTBT−PEEK interfaces tend to be minimized. Thus, the film tends to move away from the PEEK, so it rises up until the system reaches a state of equilibrium, which corresponds to the SFS configuration. The distance from the PCDTBT film to the midpoint between nanopillars amounted to approximately 300 nm. Therefore, the mechanical response of the thin polymer film in the absence of any substrate could be studied in that region (Figure 3a). To do this, AFM nanomechanical measurements

F=

16E 1/2 3/2 R δ 9

(1)

where F is the force, E the Young’s modulus, R the tip radius, and δ the sample indentation. Figure 3a shows a 3D AFM image showing the SFS film in the region on the left and a region without SFS film on the right. The blue and red arrows indicate representative points where the mechanical measurements were performed. The blue arrow indicates a free-standing point and the red arrow points out an area in which the SFS film is over the tip of a nanopillar. The set of histograms of elastic moduli measured for the different points and samples is presented in Figure 3b, and their values are collected in Table 1. The bulk PCDTBT showed a Table 1. Statistical Analysis of the Young Modulus Data for the Different Samples/Regions Young modulus (MPa) sample/indentation region

av

std dev

skewness

kurtosis excess

free-standing on pillar-tip on glass bulk

190 200 325 270

70 75 60 30

0.26 0.51 0.26 −0.41

−0.61 −1.1 −0.91 0.50

narrow distribution with an average value and standard deviation of 270 ± 40 MPa, in good agreement with the values recently obtained by Leclerc et al.31 The kurtosis excess of bulk PCDTBT is the lowest among the analyzed samples, and the peak appears slightly skewed to the left. The distribution of the film supported by the glass (denoted as “on glass” in the plot) was wider than the latter (average and standard deviation of 325 ± 60 MPa). It showed a peak at value near the bulk (mode: 300 MPa) and was skewed to higher values of the elastic modulus. Moreover, the free-standing film (denoted as “free-standing”) showed a wide bimodal distribution (modes: 130 and 220 MPa) with values mostly placed to the left of the bulk peak (average and standard deviation: 190 ± 70 MPa). Finally, the Young’s modulus distribution for the film on the nanopillars (denoted as “on pillar-tip”) was the widest (average and standard deviation of 200 ± 75 MPa), showing the main peak at similar values as in the free-standing film region (mode: 160 MPa) and a long tail at higher values. Several issues can be extracted from the above analysis. First, the fact that the PCDTBT thin film supported onto glass presented a higher Young modulus than its bulk counterpart suggests the typical substrate-induced effects on the mechanical properties of thin polymer films. This stiffening may be consequence of two causes: The first would be that the

Figure 3. (a) Artificially colored 3D AFM image of the SFS thin film (in purple color). Uncovered PEEK nanopillars are colored in ochre. The blue and red arrows indicate representative points where the mechanical analyses were performed. Red arrow: indentation on the SFS film is over the tip of a nanopillar; blue arrow: indentation on the free-standing region. (b) Young’s moduli histograms of the freestanding region of the SFS thin film (blue, denoted as “free-standing”); of the SFS thin film over the tip of PEEK nanopillars (red, denoted as “on pillar-tip”); of the thin film on glass (green, denoted as “on glass”); and of bulk PCDTBT (gray, denoted as “bulk”). The analysis was done by fitting a Hertz model to the force indentation curves. The lines are nonparametric distribution fittings to the histograms (i.e., kernel density estimations).

(specifically shallow indentations up to 10% of the film thickness) were performed using a probe with a blunted tip in order to ensure linear elastic deformations (see Figure S5). For comparative reasons, equivalent experiments were also done on a PCDTBT thin film of the same thickness deposited onto an attractive (glass substrate, Figure S4) and nonattractive D


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Figure 4. Deep indentation experiments. (a, b) AFM images of the same area of the PCDTBT SFS thin film before and after indentations, respectively. (c) Representative force indentation curve for the deep indentation experiments. The different membrane deformation mechanisms have been indicated. The inset shows the derivative of the force indentation curve where different relevant points/regimes can be identified.

mechanical properties of thin polymer film are convoluted by the presence of a stiffer glass surface.32 The second effect would be due to the formation of layers of reduced molecular mobility in the polymer film near attractive surfaces like glass,3,33 which eventually would yield a more rigid material. In order to figure out which of these effects might be at the origin of such stiffening, we turned our attention to the PCDTBT thin film over the nonattractive surface (over the pillar tip). The Young modulus of the PCDTBT thin film over the tip of the PEEK nanopillar was much lower than that of the thin film supported on glass. Therefore, the stiffening observed in the latter is likely due to the restricted mobility of PCDTBT molecules as a consequence of the attractive nature of the polymer−glass interactions. This result may evidence the strong influence of the surface energy of substrates on the mechanical properties of thin polymer films. Likewise, the Young modulus of the PCDTBT free-standing thin film was lower than both that of the bulk and the thin film on glass. This may be explained in terms of the presence of soft layers near the polymer−air interfaces35 (free surfaces) and also using the packing frustration criteria addressed by Napolitano et al.36 The enhancement of the segmental and chain mobility at polymer−air interfaces is a well-documented phenomenon18,37,38 and usually leads to a significant stress relaxation at such interfaces39 and, eventually, to a softening of the material. However, this might not be a fully universal behavior for polymers. In fact, O’Connell and Mckenna found a sever stiffening of a free-standing thin polymer film upon reduction of the thickness.16 On the other side, it is known that PCDTBT develops a kind of short-range ordering, probably of crystalline nature,24,40 as evidenced by differential scanning calorimetry and X-ray diffraction (see Figure S10). Hence, mechanical properties of the SFS film may be additionally/further affected by structural changes taking place in the nanoconfined polymer, typically a reduction of crystallinity,12 which would decrease further the modulus of the SFS thin film. It is also worth noting that the absence of substrate and the presence of a nonattractive substrate are the same scenario for the polymer thin film in terms of the Young modulus. Furthermore, if one would consider that the two modes observed in the histogram

of the SFS thin film might be correlated to different spatial regions of the sample as exposed in the Supporting Information, we obtain that the distance from the center of the pillar where the SFS thin film could be considered as freestanding is 205 nm. Finally, deep indentations were performed on the SFS thin film up to its rupture in order to determine the complete mechanical behavior of the nanoconfined polymer. The force curve behavior can only be described here by considering global membrane deformation mechanisms, instead of local deformation contact problems (like Hertzian models). The AFM images of the PCDTBT SFS thin film before and after the deep indentations are shown in Figures 4a and 4b, respectively. Such indentations caused severe plastic deformations and eventually the rupture of the SFS thin film, as can be observed in Figure 4b. A representative force indentation curve is presented in Figure 4c, which has been phenomenologically analyzed. As it can be seen in Figure 4c, at indentation depths of a few tens of nanometers the force curve became linear with a slope of k = 9 N m−1 (region 1 in Figure 4c). For clarity, the derivative of the force has also been plotted against the distance, δ. This linear dependence can be related to the bending deflection as a whole (see Figure S6).41−43 Thus, assuming, as a first approximation, a circular membrane inscribed in the triangle defined by three nanopillars (radius Rc = 150 nm) and considering the mathematical expression deduced by refs 41, 43, and 44, we obtained that the value of the bending rigidity, κ, of our freestanding thin PCDTBT film was κ = 10−15 J. Whereas this kind of analysis can only be considered as approximate, this value for the bending rigidity is the same as what is calculated from the Young’s modulus, extracted from shallow indentations experiments (see Supporting Information section 2). At around 30 nm the slope decreased, requiring less force to be deformed as the indentation progressed. This fact can be associated with the beginning of plastic deformation, and it is typical for thermoplastics. Hence, the yield strain of the thin film could be identified to be of 75% of the film thickness in this geometry. At larger depths the force reached a plateau with zero slope, in an analogous way to that of thermoplastics presenting necking and drawing (region 2 in Figure 4c). Then, the force increased E


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again up to the point of breaking, where the curve shows a decrease of the force (observed better in the derivative of the force, inset in Figure 4c). The rupture of the film took place at around 220 nm, so the strain at rupture of the free-standing thin film was measured to be 550% of the film thickness. Beyond this point the mechanical behavior of the nanoconfined material changed and became hard to analyze since resulted from a complex set of processes (region 3 in Figure 4c): the stretching of the ruptured film, friction forces between the tip walls, and the broken film; additional bending as more and more tip crosses the film, etc. Finally, the tip reached the PEEK substrate at values between 400 and 500 nm (region 4 in Figure 4c), which agrees well with the height of the nanopillars, and the slope of the curve in this region is related to the PEEK stiffness. Note that important mechanical features of the polymer confined in thin films, such as bending rigidity, yield elongation, rupture elongation, etc., cannot be characterized in a typical supported polymer thin film. In conclusion, a new procedure for achieving semi-freestanding thin polymer films has been developed. Thus, a PCDTBT semiconducting polymer thin film (40 nm thick) has been fabricated, which is held up solely on the tips of PEEK nanopillars. Thus, a new kind of nanostructure has been developed in this work. The fabrication method exploits the nonwetting behavior of PCDTBT onto an ordered PEEK nanopillar array when the mobility of the PCDTBT molecules is enhanced by a solvent annealing process. Moreover, the mechanical behavior of the free-standing PCDTBT thin film was analyzed and compared to reference samples. First, we observed that the PCDTBT polymer confined to the thin film in the absence of any substrate (free-standing region of the SFS thin film) showed a Young modulus of 190 ± 70 MPa, which means a 30% lower than bulk and 45% lower than that of a thin film of the same thickness (40 nm) deposited onto glass. Thus, this softening observed here can be considered to be a real lowdimensionality effect taking place in the nanoconfined polymer material, as interfacial effects can be ruled out in the SFS thin film. Both dynamical (enhanced dynamics) and structural (a reduction of the crystallinity) changes experimented by nanoconfined polymer might be at the origin of the softening of the semifree standing film. Furthermore, deep indentation experiments enabled determining important mechanical features of the nanostructure, such as the bending rigidity (κ = 10−15 J), yield strain (75%), rupture strain (550%), etc. These kinds of experiments are only plausible due to the free-standing nature of the thin film. Direct measurements of mechanical properties of nanoscopic polymer materials are important for predicting the collapse of the nanostructures as a consequence of capillary forces caused by surface tension of rinsing liquids, resistance of a potential diffusion barrier (nanomembranes) for nanoscopic packaging, and so on.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (J.M.). *E-mail [email protected] (M.M-G.) Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the ERC Starting Grants “NanoTEC” (ERC-StG-2008-240497) and “NANOFORCELLS” (ERC-StG-2011-278860). The authors thank Dr. Á lvaro San Paulo for fruitful discussions and Adam Wilson for technical help during the manuscript writing.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

Details of the fabrication of materials, thickness measurements, a simple model correlating the mechanical response of the SFS thin film with different spatial regions, details of shallow and deep indentations measurements, and characterization of the internal structure of bulk PCDTBT. This material is available free of charge via the Internet at http://pubs.acs.org. F


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(Conjugated) Polymer Thin Films - American Chemical Society

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