Determination of Thermal Transition Depth Profiles in Polymer

Article pubs.acs.org/Macromolecules

Determination of Thermal Transition Depth Profiles in Polymer Semiconductor Films with Ellipsometry Christian Müller,*,†,‡,§ L. Mattias Andersson,‡ Ovidio Peña-Rodríguez,† Miquel Garriga,† Olle Inganas̈ ,‡ and Mariano Campoy-Quiles*,† †

Institut de Ciència de Materials de Barcelona (ICMAB-CSIC), Esfera UAB, Bellaterra 08193, Spain Biomolecular and Organic Electronics, Department of Physics, Chemistry & Biology, Linköping University, 58183 Linköping, Sweden § Department of Chemical and Biological Engineering/Polymer Technology, Chalmers University of Technology, 41296 Göteborg, Sweden ‡

S Supporting Information *

ABSTRACT: Geometric confinement and interface effects can significantly alter the thermodynamic properties of thin polymer films. Phase transition temperatures have been shown to strongly depend on film thickness below a critical thickness threshold. It has been suggested that this behavior is due to an interface-induced continuous variation in phase transition temperatures throughout the depth of the films. Here we employ variable-temperature spectroscopic ellipsometry to demonstrate the existence of these depth profiles. We examine four different polymer semiconductors that are of interest for organic light-emitting diodes, solar cells, and field-effect transistors. In contrast to insulating polymers, these light-absorbing materials provide detailed information about structural changes as a function of depth due to wavelengthdependent attenuation. This concept enables us to investigate a broad range of thermodynamic processes including the glass transition, crystallization as well as crystalline and liquid-crystalline melting. In general, for the here investigated systems, higher transition temperatures are found at the free surface. Finally, the deduced profiles are used to predict the thickness dependence of the mean phase transition temperature.

■

(methyl methacrylate) (PMMA),7,8 and polyester9 but also conjugated polymers such as polyfluorenes,10−12 quinoxalineand carbazole-based copolymers,13,14 and conjugated polymer:fullerene blends.15−17 Typically, to simplify the analysis, variable-temperature ellipsometry is performed at wavelengths at which thin films are transparent, i.e., the extinction coefficient κ ∼ 0. As a result, transition temperatures averaged over the depth of the investigated thin films are obtained, which disregard the likely presence of vertical variations, i.e., thermal transition depth profiles. However, in many cases knowledge of the average thermal behavior of thin semiconductor films is insufficient to fully describe the response of thin-film optoelectronic devices, which often display significant sensitivity to interface effects. Even a partial phase transition that only occurs close to e.g. a substrate interface is likely to strongly influence the device performance. Analysis of interface effects usually relies on (1) inclusion of an optical marker at a precise film depth18 or, more commonly, (2) a systematic variation in film thickness.2−9,12,14,15 Unfortunately, the first method requires demanding processing steps and may alter the nanostructure of the investigated

INTRODUCTION Polymer semiconductors currently attract considerable interest for a range of advanced optoelectronic applications including light-emitting diodes (LEDs), solar cells, and field-effect transistors (FETs). Thin-film devices are the most technologically relevant architectures and typically comprise a 50−200 nm thick active layer that is printed or coated from solution. Precise knowledge of the thermal behavior of the semiconductor material is critical for selecting appropriate postdeposition annealing protocols, which are needed to optimize the nanostructure as well as for assessing the longterm thermal stability. Confinement effects as well as the presence of a free surface and the substrate interface can significantly alter the thermal behavior, rendering bulk measurements insufficient. Instead, it is necessary to directly record the phase transition behavior of thin films, which can be carried out with variable-temperature spectroscopic ellipsometry (SE). Modeling of the measured ellipsometric functions tan Ψ(λ) and cos Δ(λ) provides information regarding film thickness d and complex refractive index (n + iκ). Phase transitions are accompanied by a change in volume expansion coefficient, which can be detected as a change in slope in plots of d(T). This method has been successfully applied to study thin films of commodity polymers such as polystyrene (PS),1−5 poly(α-methylstyrene),6 poly© 2013 American Chemical Society

Received: April 28, 2013 Revised: July 2, 2013 Published: September 9, 2013 7325

dx.doi.org/10.1021/ma400871u | Macromolecules 2013, 46, 7325−7331

Macromolecules

Article

Table 1. Number- and Weight-Average Molecular Weight, Mn and Mw (provided by supplier); Polydispersity Index = PDI

a

Regioregularity ∼99% (provided by supplier).

different materials systems of interest for a broad range of optoelectronic applications, including light-emitting diodes, solar cells, and organic field-effect transistors. Moreover, these five materials exhibit four types of phase transitions, which indicates that the here proposed method is general.

material. The second approach, on the other hand, is limited because it is necessary to change critical processing parameters such as the initial polymer concentration in solution or the spin-coating speed to alter the film thickness, which can strongly influence the degree of molecular ordering and orientation.19,20 As a result, it can be challenging to distinguish confinement, interfacial, and processing effects. Therefore, we set out to establish a method based on variable-temperature SE that permits to determine the entire thermal transition depth profile for a single thin-film specimen. This goal is achieved by harnessing the optical response at wavelengths at which the material absorbs light, i.e., κ > 0. Interestingly, conjugated polymers often exhibit both an absorption band and a transparency region within the visible spectrum, where most ellipsometers operate. In contrast, most commodity polymers are optically transparent across the visible spectrum, which may explain why our approach has not been considered previously. Monitoring changes in κ is better suited for the investigation of phase transitions than monitoring changes in d. This is because each phase is characterized by a specific absorption spectrum, and therefore the presence of a given phase is more readily detectable. Therefore, analysis of the full variable-temperature SE spectrum, i.e. both the transparent and the absorbing region, can reveal information about thermal processes as a function of depth. In this paper we demonstrate that our approach permits to record vertical phase transition profiles across thin polymer semiconductor films. We illustrate the potential of our approach by investigating five

■

EXPERIMENTAL SECTION

Materials. All materials were used as received; see Table 1 for chemical structure and molecular weights. Regioregular poly(3hexylthiophene) (P3HT) and poly(9,9-dioctylfluorene) (PFO) were purchased from Sigma-Aldrich. Poly[2,7-(9,9-dioctylfluorene)-alt-5,5(4,7-di-2-thienyl-2,1,3-benzothiadiazole)] (APFO3) and poly[2,7-(9,9dioctylfluorene)-alt-5,5-(5,10-di-2-thienyl-2,3,7,8-tetraphenyl-pyrazino[2,3-g]quinoxaline)] (APFO-Green9) were provided by Prof. Mats Andersson, Chalmers University of Technology (cf. refs 21 and 22 for details on synthesis; both polymers are phenyl end-capped). [6,6]Phenyl-C61-butyric acid methyl ester (PCBM) was obtained from Solenne BV. Sample Preparation. Thin films for ellipsometry, optical microscopy, and FET device studies were spin-coated from odichlorobenzene (APFO3, 9:1 APFO3:PCBM, APFO-Green9 and P3HT; 20−30 g L−1) or mixed xylenes (PFO; 20 g L−1), followed by drying for several hours at ambient temperature and atmosphere in order to permit evaporation of residual solvent. Ellipsometry. Variable-temperature SE was carried out with a RC2 instrument from J.A. Woollam Co., Inc., at a fixed angle of incidence (70°), equipped with a nitrogen-flushed hot stage. Spin-coated thin films on silicon substrates with a 2.0 nm native silicon oxide top layer were heated by 10 °C min−1. The integration time for each real time SE scan during heating was 2 s. Temperature calibration was 7326


Macromolecules

Article

copy.29 In another study, we have employed variable-temperature SE to study the phase behavior of spin-coated APFO3 films.15 Sufficiently thick films with d ∼ 60−190 nm featured a thickness-independent liquid-crystalline melting temperature 30 T̅ SE lc ∼ 179 °C, which however increased for d ≤ 50 nm. Here, we employ variable-temperature SE on a 70 nm thick APFO3 film that had been spin-coated on a silicon substrate (see Experimental Section for more details). Initially, we extracted Tlc from the final change in slope of Ψλ(T) as displayed in Figure 1, which according to our previous work

performed with an external thin-film temperature sensor (Omega Engineering, UK) that had been glued onto a silicon wafer; error of quoted temperatures is estimated to be ±3 °C. Analysis of Ellipsometric Data. Variable-temperature SE measurements were analyzed using the WinElli II piece of software by SOPRALAB. First, heating scans of spin-coated films were analyzed in order to relate our study to the thermal properties and the stability on heating of as-prepared films, which are most relevant for applications. A detailed description of the different simulations and tests is presented throughout the text and complemented in the Supporting Information. In brief, the optical constants of films before onset and after completion of thermal transitions were deduced using the standard critical point model of the dielectric function.23 At intermediate temperatures different effective medium approximations were applied together with discrete sublayering of the film. Several models were considered, and in all cases the best results in terms of convergence to a unique solution and the smallest standard deviation were obtained with a bilayer structure, which comprised a pristine sublayer and one that had undergone the thermal transition (see Supporting Information for more details). Film expansion was accounted for by blending the optical constants through the Bruggeman model24 with air. Within this approach, we model the optical constants as a function of temperature for homogeneous phases as a mixture of the optical constants at a given temperature and void fraction. Besides film expansion and the concomitant reduction in density, raising the temperature may broaden the absorption peaks. This effect is not included in our simple description of the temperature-dependent index of refraction. We minimized this effect by taking as reference the refractive index values of the homogeneous film at temperatures close to the beginning and the end of the phase transition. To summarize, we used four fitting parameters: the thickness and void fraction of each of the two sublayers. The temperature dependence of the optical properties of the silicon substrate was also considered. Thermal Analysis. Differential scanning calorimetry (DSC) was carried out under nitrogen from 20 to 300 °C at a scan rate of 10 °C min−1 with a PerkinElmer Pyris 1 DSC instrument. Optical microscopy of spin-coated films on glass slides was carried at a scan rate of 10 °C min−1 with an Olympus BH2 polarizing microscope equipped with a nitrogen flushed Mettler FP82HT hot stage. FET Charge Transport Measurements. FETs were fabricated in bottom-gate bottom-contact configuration on a highly doped silicon wafer with a thermally grown 100 nm silicon dioxide layer; the two layers served as the gate electrode and gate insulator, respectively. Au source and drain electrodes with a Cr adhesion layer were defined by standard photolithography. The channel length L and width W varied from 9 to 37 μm and from 2 to 16 mm, respectively. Variabletemperature electronic characterization of FETs was carried out with a Keithley 4200 parameter analyzer in high vacuum. Saturated hole mobilities μh were extracted from the transfer characteristics using the expression

⎛∂ I sd μ h = ⎜⎜ ⎝ ∂Vg

⎞2 2L ⎟⎟ ⎠ WC i

Figure 1. (a) Extinction coefficient κ at 50 °C as a function of wavelength λ for an APFO3 film (room temperature thickness dRT ∼ 70 nm). (b) Ψλ(T) at λ = 550 and 800 nm. Liquid-crystalline melting temperatures Tlc and their difference ΔTlc are indicated.

corresponds to liquid-crystalline melting of APFO3.15 An average T̅ SE lc ∼ 185 °C can be estimated from the transparent region, e.g. at λ = 800 nm, where light probes the whole depth of the film and as a result Tlc is wavelength-independent. In contrast, at the absorption maximum around λ ∼ 550 nm, only the uppermost free surface of the film is sampled, as the penetration depth 1/α = λ/4πκ ∼ 60 nm, which yields a surface maximum Tfree ∼ 215 °C that is ΔTlc ∼ 30 °C above the lc average T̅ SE . Certainly, this observation implies that for some lc other section of the film Tlc < T̅ SE lc , which results in a depth profile of this transition temperature. A mathematically rigorous approach for deduction of the full depth profile Tlc(z) can be obtained by dividing the film into a large number of sublayers6 and introducing an expression of the form

(1)

where Isd is the source-drain current (saturation regime), Vg the gate voltage, and Ci the insulator capacitance.

■

RESULTS AND DISCUSSION In a first set of experiments we examined the liquid-crystalline melting of the fluorene copolymer APFO3 that has been widely investigated for organic photovoltaics.15,21,25−29 Previously, we have studied this phase transition of APFO3 in detail. Differential scanning calorimetry (DSC) second heating thermograms of the here used molecular weight of APFO3 revealed a distinct endotherm with a peak temperature T̅ DSC ∼ lc 196 °C, which corresponds to bulk liquid-crystalline melting as confirmed with variable-temperature polarized optical micros-

d

Tlc(λ) =

7327

∫0 Tlc(z)I(z , λ) dz d

∫0 I(z , λ) dz

(2)


Macromolecules

Article

where z is the depth, Tlc(λ) the change in Tlc with wavelength, and I(z,λ) the contribution of the light intensity profile across the film to the reflected beam. To retrieve Tlc(z), a numerical inversion of eq 2 is necessary, which is unfeasible because I(z,λ) is unknown. In other words, the exact contribution of different sublayers to the reflected beam and to Tlc(λ) cannot be explicitly formulated. Attempts to approximate I(z,λ) by the Beer−Lambert law or even to calculate the full electric field distribution within the film did not result in good fits of the experimental data for a consistent number of samples and measurements. In order to avoid the physical and mathematical limitations associated with eq 2, we chose to employ an alternative analysis method of the SE data that is based on a combination of two phases with distinct optical constants that represent the liquidcrystalline phase and isotropic liquid (Figure 2). At temper-

bottom, and (4) a triple layer with either phase at the center. We applied each of these models to the recorded variabletemperature SE data. A bilayer structure of a liquid-crystalline layer on top of an isotropic liquid layer consistently yields the lowest standard deviation when fitting the SE spectra across the whole temperature range (Supporting Information, Figure S2). Inverting the order of the phases led to very bad fits of the SE data, and triple layers resulted in one of the layers with thickness equal to zero (within the fitting error), which gives confidence to our model. The thickness evolution of the two sublayers provides an easy measure for the depth up to which the APFO3 film has undergone liquid-crystalline melting. We visualize this trend by plotting the thickness of the liquid-crystalline sublayer against temperature (Figure 3). Conceptually, our data are indicating

Figure 2. Schematic representation of the modeling approach. The refractive indices n and extinction coefficients κ are determined for the homogeneous liquid-crystalline phase and the melt at T ≪ T̅ SE lc and T ≫ T̅ SE lc , respectively. Then, the makeup of a bilayer profile is calculated for intermediate temperatures T ∼ T̅ SE lc .

Figure 3. (a) Temperature evolution of the glassy (open diamonds) liquid-crystalline (red circles) and liquid sublayer (open diamonds) of an APFO3 film. Note that the total film thickness expands from 70 to 83 nm during heating. (b) Depth profile Tlc(z) deduced from the thickness of the liquid-crystalline sublayer of APFO3 films with dRT ∼ 57 nm (open triangles), 70 nm (red circles; cf. part a) and 100 nm (open diamonds). The average T̅ SE lc is indicated with a solid black line.

atures T ≪ T̅ SE lc no sublayer of the APFO3 film has undergone liquid-crystalline melting. Likewise, the complete film is an isotropic liquid for T ≫ T̅ SE lc . On the basis of this assumption, it is straightforward to extract the optical constants of both phases, which we performed at 150 and 250 °C (Supporting Information, Figure S2) by analyzing the corresponding SE data at those temperatures. At intermediate temperatures T ∼ T̅ SE lc the film may be composed of a mixture of the initial liquid-crystalline phase and the final isotropic liquid, for which the optical constants have been deduced. Mathematically, this intermediate film configuration can be described as (1) a homogeneous mixture (blend) of the two phases, (2) a gradually vertically segregated mixture of the two phases forming a vertical profile, (3) a bilayer with the isotropic liquid phase either on the top or

that at a temperature of ∼165 °C the half of the film closer to the substrate has undergone the phase transition and the other half has not. To melt the next 5 nm of film, an additional temperature increase is needed. In other words, if we think of the film as a composition of infinitesimal sublayers, we have to raise the temperature up to 165 °C in order to melt a sublayer with infinitesimal layer width situated at a depth of 40 nm. This argument allows deducing the phase transition profile by assigning Tlc(z) to the temperature at which the liquidcrystalline sublayer had a thickness of z. Mathematically, this means replotting the points in Figure 3a for which there exist 7328


Macromolecules

Article

systematic error arising from the experiment, such as a thermal gradient across the thin film induced by substrate heating, convection effects, etc. In order to confirm that liquid-crystalline melting of APFO3 films commences at the substrate interface, we performed variable-temperature FET measurements. We used a bottomgate, bottom-contact configuration, which predominately probes the substrate interface (Figure 4). Labram et al. as

two effective layers by exchanging the axes and its names (T by Tlc(z) and thickness by depth). The deduced Tlc(z) profile is monotonous but highly nonlinear. Close to the substrate liquid-crystalline melting ∼ 150 °C. A gradual increase to T ∼ already occurs at Tsubstrate lc T̅ SE lc up to a depth z ∼ 15 nm is followed by a rapid rise to surface ∼ 215 °C at the free top surface. Tfree lc The deduced Tlc(z) depth profile can be rationalized with the presence of a more densely packed free surface layer as compared to the bulk, which Anselmo et al. have recently reported for APFO3.31 This surface layer may arise because of a higher degree of in-plane orientation or, potentially, vertical separation of the polymer according to its molecular weight distribution. Indeed, the liquid-crystalline melting of APFO3 ∼ 100 strongly depends on the chain length, ranging from T̅ DSC lc °C for a number-average molecular weight Mn ∼ 2 kg mol−1 to ∼ 290 °C for Mn ∼ 38 kg mol−1.29 Our analysis indicates T̅ DSC lc that the free surface layer is about 15−20 nm thick and, in addition, does not vary with film thickness for a range d ∼ 57− 100 nm (Figure 3b). However, molecular weight gradients are expected to scale with d in the sense that the percentage of a given molecular weight in the mixture does not depend on the thickness and thus thicker films should have the same relative Mn distribution leading to a surface skin with equal relative thickness, not absolute thickness. Therefore, we propose that variations in molecular packing, and not molecular weight, give rise to the observed Tlc(z) depth profile. The Tg of the here studied spin-coated samples can be easily determined with variable-temperature SE because solutionprocessing results in considerable disorder, which is frozen in at room temperature. When heated above Tg the extent of the liquid-crystalline phase develops, as evidenced by optical microscopy as well as the sudden decrease in film thickness (cf. refs 15 and 29), which results in a distinct change in the optical constants. We used the same modeling approach to examine the glass transition temperature Tg of APFO3 films. Again, a bilayer model yields the lowest standard deviation. We observe a distinct change between 115 and 125 °C, which agrees with DSC data (cf. Figure 3a and Table 2). Evidently, the sharp Tg(z) profile has a markedly different shape than the Tlc(z) profile, which occurs over a broad temperature range ΔTlc ∼ 60 °C. Therefore, we consider it unlikely that the here deduced thermal transition profiles are due to any possible

Figure 4. Hole mobility of an APFO3 film during heating, extracted from a representative FET measurement (L ∼ 25 μm, W ∼ 10 mm).

well as Andersson have demonstrated that changes in the saturated hole charge-carrier mobility μh can be related to phase transitions that occur in the semiconductor active layer.32,33 We observe distinct changes in μh around 110 and 160 °C, which are consistent with T̅ substrate and Tsubstrate , respectively (cf. Table g lc 2). Evidently, our FET and SE measurements are in good agreement. In a further set of experiments we explored if our modeling approach can be used to record thermal transition profiles in other materials. We chose to study (1) the liquid-crystalline fluorene copolymer APFO-Green9,22 (2) the semicrystalline polyfluorene PFO,10,12 (3) the semicrystalline polythiophene P3HT,34 and (4) a 9:1 blend of APFO3 and the fullerene derivative PCBM, which forms a homogeneous phase.15 For each system, a bilayer model gave the best results, which are shown in Figure 5 (see also Supporting Information, Figures S8−S10). We were able to confirm that glass transitions generally are characterized by a sharp depth profile (Table 2). In contrast, liquid-crystalline melting of APFO-Green9 and 9:1 APFO3:PCBM, crystallization of PFO, and crystalline melting of P3HT yield pronounced profiles that occur over a much broader temperature range and in each case feature the highest transition temperature at the free surface (Figure 5). Interestingly, it appears that the surface skin is in all cases about 10−20 nm thick. The Tlc(z) profiles of APFO-Green9, APFO3, and 9:1 APFO3:PCBM appear over vastly different temperature ranges but are similar in shape, which may be due to the fact that a comparable thermal transition is monitored. The crystallization temperature Tc(z) profile of PFO is consistent with a recent surface report by Roigé et al., which established that Tfree > c PL/Raman T̅ c by comparing variable-temperature atomic force microscopy (AFM) and Raman measurements.35 In that study, nonresonant Raman (in the transparent region) was employed to obtain an average value of the phase transition throughout the film depth and roughness values deduced from temperature-dependent AFM images provided the phase

Table 2. Phase Transition Temperature Values Obtained from Second Heating DSC Thermograms and SE material APFO3 APFO3:PCBM APFO-Green9 PFO P3HT

Ttransition (°C)

T̅ DSC (°C)

T̅ SE (°C)

Tf ree surface (°C)

Tsubstrate (°C)

ΔT (°C)

Tg Tlc Tg Tlc Tg Tlc Tc Tm

109 196

109 185 110 157 177 242 90d 170

125 215 120 185 200 260 115d 230

115a 155a 100 125 170 210 75 150

10 60 20 60 30 50 40 80

192b 287c 92 224

cf. FET measurements in Figure 4: Tsubstrate ∼ 110 °C and Tsubstrate ∼ g lc b 160 °C on the substrate. Taken from ref 41 for APFO-Green9 with Mn ∼ 11 kg mol−1 and Mw ∼ 25 kg mol−1. cClearing temperature from variable-temperature optical microscopy. dcf. ref 35: photolumines∼ 82 °C and cence and Raman spectroscopy yield an average T̅ PL/Raman c surface ∼ 92 °C. atomic force microscopy yields Tfree c a

7329


Macromolecules

Article

Figure 6. Estimated film thickness dependence of the average phase transition temperature T̅ SE normalized by the bulk value T̅ DSC for APFO3 (red circles), APFO-Green9 (green triangles), P3HT (blue diamonds), and PFO (black inverted triangles).

assumes that the changes in nanostructure are limited to the decrease in the thickness of the central part of the film, which could be a great simplification in some cases.19,20 We find that T̅ lc(d) of APFO3 and APFO-Green9 substantially increases with decreasing film thickness, which is in agreement with our previous study.15 Likewise, for T̅ m(d) of P3HT we obtain a reduction for thinner films, which has also been observed for other semicrystalline polymers such as polyethylene.38−40 Prediction of T̅ c(d) of PFO is more challenging because the contributions from the substrate interface and free surface are similar in magnitude. A previous study has revealed a nonmonotonic trend with a maximum T̅ SE c at d ∼ 80−100 nm.12 We were able to reproduce this behavior with the recorded Tc profile by assuming that the substrate interface contribution dominates over free surface effects.

Figure 5. Tlc(z) depth profiles of APFO-Green9 (green triangles), APFO3 (red circles), and a 9:1 APFO3:PCBM blend film (open circles); melting depth profile Tm(z) of a P3HT film (blue diamonds); and recrystallization depth profile Tc(z) of a PFO film (black inverted triangles). The depth z is normalized with respect to the total film thickness (dAPFO‑Green9 ∼ 70 nm, dAPFO3 ∼ 70 nm, d9:1APFO3:PCBM ∼ 65 nm, dP3HT ∼ 115 nm, dPFO ∼ 120 nm).

transition temperature at the free surface. (Note that crystallization increases surface roughness.) Interestingly, this report showed a more than 10 °C difference in the deduced crystallization temperatures. The same study also reported even larger differences between average and surface values for films consisting of blends of P3HT and PCBM. Moreover, the melting temperature profile Tm(z) of P3HT is consistent with grazing-incidence wide-angle X-ray scattering (GIWAXS) measurements of P3HT thin films on silicon substrates, which indicate that the polymer crystallinity decreases for a thickness of less than 40 nm.36,37 Our SE analysis reveals a substrate layer of comparable thickness that is characterized by Tsubstrate as low as 150 °C. This observation suggests the m presence of smaller P3HT crystallites and, most likely, more amorphous material. It is worth mentioning that the accuracy of the presented method relies on having two distinct reference refractive indices at temperatures below and above the phase transition. This could include very different absorption spectra, densities, or anisotropies of the two phases. In this respect, first-order transitions will be more readily detectable using this technique. Moreover, there may also be restrictions in terms of thicknesses, since access to the bottom of the layers will be compromised for very thick films (>250 nm). (See Supporting Information for additional details on the model.) We complete our analysis with an attempt to predict the thickness dependence of the here investigated thermal transitions based on the recorded depth profiles. The profiles obtained for APFO3 films (Figure 3b) suggest that a reduction in film thickness predominately affects the central part of the film. Provided that this behavior is general, we can estimate the thickness dependence of the mean phase transition temperatures for the investigated materials, which should be a continuous function (see Figure 6). We note that this analysis

■

CONCLUSIONS We have demonstrated that variable-temperature spectroscopic ellipsometry can be effectively used to record vertical thermal transitions profiles in thin polymer semiconductor films. The optical constants of the material before onset and after completion of the thermal transition were combined to model a two-phase film configuration at intermediate temperatures. We analyzed a wide range of materials with this method, including the liquid-crystalline fluorene copolymers APFO3 and APFO-Green9, the semicrystalline polymers PFO and P3HT, and a 9:1 APFO3:PCBM blend. In all cases, a simple bilayer model best described the experimental data, which permitted us to extract thermal transition profiles by plotting the sublayer thickness as a function of temperature. Generally, we observe sharp glass transition profiles whereas crystallization, crystalline melting, and liquid-crystalline melting profiles occur over a broad temperature range. Material close to the substrate interface displays depressed transition temperatures that gradually increase toward the free surface to reach values larger than the corresponding bulk transition temperatures. Our results are consistent with additional probes, such as variable-temperature AFM and FET mobility measurements. Finally, we have shown that the recorded profiles are in agreement with previous reports that have explored the influence of film thickness on thermal transitions. Clearly, our results help to understand the often complex phase behavior of thin polymer semiconductor films and may provide guidance for novel preparation protocols that exploit these phenomena. 7330


Macromolecules

■

Article

(20) DeLongchamp, D. M.; Vogel, B. M.; Jung, Y.; Gurau, M. C.; Richter, C. A.; Kirillov, O. A.; Obrzut, J.; Fischer, D. A.; Sambasivan, S.; Richter, L. J.; Lin, E. K. Chem. Mater. 2005, 17, 5610−5612. (21) Svensson, M.; Zhang, F. L.; Veenstra, S. C.; Verhees, W. J. H.; Hummelen, J. C.; Kroon, J. M.; Inganäs, O.; Andersson, M. R. Adv. Mater. 2003, 15, 988−991. (22) Zhang, F. L.; Bijleveld, J.; Perzon, E.; Tvingstedt, K.; Barrau, S.; Inganäs, O.; Andersson, M. R. J. Mater. Chem. 2008, 18, 5468−5474. (23) Campoy-Quiles, M.; Nelson, J.; Bradley, D. D. C.; Etchegoin, P. G. Phys. Rev. B 2007, 76, 235206. (24) Tompkins, H. G.; Irene, E. A. Handbook of Ellipsometry; William Andrew Publishing: Norwich, NY, 2005. (25) Inganäs, O.; Svensson, M.; Zhang, F.; Gadisa, A.; Persson, N. K.; Wang, X.; Andersson, M. R. Appl. Phys. A: Mater. Sci. Process. 2004, 79, 31−35. (26) Zhang, F. L.; Jespersen, K. G.; Björström, C.; Svensson, M.; Andersson, M. R.; Sundström, V.; Magnusson, K.; Moons, E.; Yartsev, A.; Inganäs, O. Adv. Funct. Mater. 2006, 16, 667−674. (27) Andersson, L. M.; Zhang, F. L.; Inganäs, O. Appl. Phys. Lett. 2007, 91, 071108. (28) Svanström, C. M. B.; Rysz, J.; Bernasik, A.; Budkowski, A.; Zhang, F.; Inganäs, O.; Andersson, M. R.; Magnusson, K. O.; BensonSmith, J. J.; Nelson, J.; Moons, E. Adv. Mater. 2009, 21, 4398−4403. (29) Müller, C.; Wang, E. G.; Andersson, L. M.; Tvingstedt, K.; Zhou, Y.; Andersson, M. R.; Inganäs, O. Adv. Funct. Mater. 2010, 20, 2124−2131. (30) Note the difference in DSC and SE liquid-crystalline melting temperatures, which we ascribe to the thermal history of APFO3, i.e., cooling from the isotropic liquid state and spin-coating from solution. (31) Anselmo, A. S.; Dzwilewski, A.; Svensson, K.; Moons, E. J. Polym. Sci., Part B: Polym. Phys. 2013, 51, 176−182. (32) Labram, J. G.; Domingo, E. B.; Stingelin, N.; Bradley, D. D. C.; Anthopoulos, T. D. Adv. Funct. Mater. 2011, 21, 356−363. (33) Andersson, L. M. Org. Electron. 2011, 12, 300−305. (34) Brinkmann, M.; Wittmann, J. C. Adv. Mater. 2006, 18, 860−863. (35) Roigé, A.; Campoy-Quiles, M.; Ossó, J. O.; Alonso, M. I.; Vega, L. F.; Garriga, M. Synth. Met. 2012, 161, 2570−2574. (36) Joshi, S.; Grigorian, S.; Pietsch, U.; Pingel, P.; Zen, A.; Neher, D.; Scherf, U. Macromolecules 2008, 41, 6800−6808. (37) Jimison, L. H.; Himmelberger, S.; Duong, D. T.; Rivnay, J.; Toney, M. F.; Salleo, A. J. Polym. Sci., Part B: Polym. Phys. 2013, 51, 611−620. (38) Wang, Y.; Ge, S.; Rafailovich, M.; Sokolov, J.; Zou, Y.; Ade, H.; Luning, J.; Lustiger, A.; Marom, G. Macromolecules 2004, 37, 3319− 3327. (39) Wang, Y.; Rafailovich, M.; Sokolov, J.; Gersappe, D.; Araki, T.; Zou, Y.; Kilcoyne, A. D. L.; Ade, H.; Marom, G.; Lustiger, A. Phys. Rev. Lett. 2006, 96, 028303. (40) Bernazzani, P.; Sanchez, R. F. Eur. Phys. J. E 2008, 26, 427−434. (41) Kroon, R. Synthesis and properties of pi-conjugated polymers for organic photovoltaics. PhD Thesis, Chalmers University of Technology, 2013.

ASSOCIATED CONTENT

S Supporting Information *

Modeling of ellipsometry data. This material is available free of charge via the Internet at http://pubs.acs.org.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (C.M.); mcampoy@ icmab.es (M.C.-Q.). Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS We thank Prof. Mats Andersson (Chalmers University of Technology) for generously providing APFO3 and APFOGreen9 as well as Prof. Hans Arwin (Linköping University) for access to the RC2 ellipsometer and Dr. Isabel Alonso (ICMAB) for useful discussions. The Spanish Ministerio de Economı ́a y Competitividad is acknowledged for financial support through projects MAT2009-10642 and PLE2009-0086. C.M. gratefully acknowledges funding from the CSIC through the JAE-Doc program (European Social Fund) as well as from Formas and the Chalmers Area of Advance Energy.

■

REFERENCES

(1) Beaucage, G.; Composto, R.; Stein, R. S. J. Polym. Sci., Part B: Polym. Phys. 1993, 31, 319−326. (2) Keddie, J. L.; Jones, R. A. L.; Cory, R. A. Europhys. Lett. 1994, 27, 59−64. (3) Dalnoki-Veress, K.; Forrest, J. A.; Murray, C.; Gigault, C.; Dutcher, J. R. Phys. Rev. E 2001, 63, 031801. (4) Singh, L.; Ludovice, P. J.; Henderson, C. L. Thin Solid Films 2004, 449, 231−241. (5) Kawana, S.; Jones, R. A. L. Phys. Rev. E 2001, 63, 021501. (6) Kim, J. H.; Jang, J.; Zin, W. C. Langmuir 2000, 16, 4064−4067. (7) Sharp, J. S.; Forrest, J. A. Phys. Rev. E 2003, 67, 031805. (8) Erber, M.; Tress, M.; Mapesa, E. U.; Serghei, A.; Eichhorn, K. J.; Voit, B.; Kremer, F. Macromolecules 2010, 43, 7729−7733. (9) Erber, M.; Khalyavina, A.; Eichhorn, K. J.; Voit, B. I. Polymer 2010, 51, 129−135. (10) Campoy-Quiles, M.; Etchegoin, P. G.; Bradley, D. D. C. Synth. Met. 2005, 155, 279−282. (11) Sims, M.; K. Z; Campoy Quiles, M.; Xia, R.; Stavrinou, P. N.; Bradley, D. D. C.; Etchegoin, P. J. Phys.: Condens. Matter 2005, 17, 1− 12. (12) Campoy-Quiles, M.; Sims, M.; Etchegoin, P. G.; Bradley, D. D. C. Macromolecules 2006, 39, 7673−7680. (13) Kroon, R.; Gehlhaar, R.; Steckler, T. T.; Henriksson, P.; Müller, C.; Bergqvist, J.; Hadipour, A.; Heremans, P.; Andersson, M. R. Sol. Energy Mater. Sol. Cells 2012, 105, 280−286. (14) Wang, T.; Pearson, A. J.; Dunbar, A. D. F.; Staniec, P. A.; Watters, D. C.; Coles, D.; Yi, H.; Iraqi, A.; Lidzey, D. G.; Jones, R. A. L. Eur. Phys. J. E 2012, 35, 129. (15) Müller, C.; Bergqvist, J.; Vandewal, K.; Tvingstedt, K.; Anselmo, A. S.; Magnusson, R.; Alonso, M. I.; Moons, E.; Arwin, H.; CampoyQuiles, M.; Inganäs, O. J. Mater. Chem. 2011, 21, 10676−10684. (16) Campoy-Quiles, M.; Ferenczi, T.; Agostinelli, T.; Etchegoin, P. G.; Kim, Y.; Anthopoulos, T. D.; Stavrinou, P. N.; Bradley, D. D. C.; Nelson, J. Nat. Mater. 2008, 7, 158−164. (17) Pearson, A. J.; Wang, T.; Jones, R. A. L.; Lidzey, D. G.; Staniec, P. A.; Hopkinson, P. E.; Donald, A. M. Macromolecules 2012, 45, 1499−1508. (18) Ellison, C. J.; Torkelson, J. M. Nat. Mater. 2003, 2, 695−700. (19) Zhokhavets, U.; Gobsch, G.; Hoppe, H.; Sariciftci, N. S. Thin Solid Films 2004, 451, 69−73. 7331


Determination of Thermal Transition Depth Profiles in Polymer

Recommend Documents