Experimental Relation between Stranski− Krastanov Growth of DIP

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2009, 113, 4234–4239 Published on Web 01/14/2009

Experimental Relation between Stranski-Krastanov Growth of DIP/F16CoPc Heterostructures and the Reconstruction of the Organic Interface Dimas G. de Oteyza,*,† Esther Barrena,‡,§ Yi Zhang,‡ Tobias N. Krauss,‡ Ayse Turak,‡ Alexei Vorobiev,| and Helmut Dosch‡,§ Donostia International Physics Center, Paseo Manuel Lardizabal 4, 20018 San Sebastia´n, Spain, Max-Planck-Institut fu¨r Metallforschung, Heisenbergstrass 3, 70569 Stuttgart, Germany, Institut fu¨r Theoretische and Angewandte Physik, UniVersita¨t Stuttgart, 70550 Stuttgart, Germany, and European Synchrotron Radiation Facility, BP 220, 38043 Grenoble Cedex 9, France ReceiVed: October 28, 2008; ReVised Manuscript ReceiVed: December 05, 2008

By a combined AFM/X-ray study, we unveil a reconstruction at the organic interface accompanying the Stranski-Kranstanov growth of di-indenoperylene (DIP) deposited on fluorinated cobalt-phthalocyanines (F16CoPc). This reconstruction involves an abrupt change in the F16CoPc packing in those areas covered by DIP. After the total completion of the first DIP monolayer, the entire F16CoPc interfacial layer is reconstructed and eventually becomes buried under the growing DIP film. We demonstrate that the morphological transition from smooth to highly textured heterostructures occurring at a threshold temperature of 70 °C is intimately related to the thermal activation of the reconstruction of the underlying F16CoPc layers. This study provides further understanding of the molecular-scale processes that ultimately determine the controlled growth of organic heterojunctions. 1. Introduction The latest advances in the growth and processing of organic semiconducting thin films have led to the introduction of organic-based devices in the electronics market.1,2 The attractiveness of so-called plastic electronics stems from the wide range of novel physical and chemical properties offered by organic materials, as well as the low temperature processing which allows the use of large areas or flexible substrates (such as plastic, cloth, or paper), with a dramatic reduction of production costs. However, as compared to their conventional inorganic counterparts, organic electronic materials and devices still suffer from poor performance, which is one of the most important hurdles to-date for the development of plastic electronics into a well established technology. One route to improve device efficiencies is the optimization of the structure and morphology of the active layers.3-6 Particular attention has to be paid to the regions around the various interfaces, as they have a crucial role in the final device performance.7 While metal-organic interfaces define the charge carrier injection or extraction from the devices,8-10 the organic-dielectric interfaces in thin film transistors strongly influence the charge carrier mobility and threshold voltage.11-13 Furthermore, many devices such as solar cells, organic light emitting diodes, or ambipolar transistors are based on p-n junctions, in which the electronic properties at the additional organic-organic interfaces also play a pivotal role in the device * To whom correspondence should be addressed. E-mail: d_g_oteyza@ ehu.es. Tel.: +34 943015389. Fax: +34 943015600. † Donostia International Physics Center. ‡ Max-Planck-Institut fu¨r Metallforschung. § Universita¨t Stuttgart. | European Synchrotron Radiation Facility.

10.1021/jp809512a CCC: $40.75

performance.14-18 The relevance of structure and morphology at these interfaces has triggered substantial research activities on the growth and structure of organic heterojunctions. The inherent anisotropy of most organic semiconductors is mirrored by way of example in orientation dependent optoelectronic properties of the thin films.19,20 This has motivated a growing number of studies on molecular orientation and potential templating effects in organic-organic heterostructures,19-30 which were eventually beneficially applied to the design of organic-based devices with improved performance.30 Organic heteroepitaxy has also been shown to allow the growth of organic semiconducting films with enhanced azimuthal orientation and/or domain sizes23,31-34 and with a consequent improvement of the electric transport.33,34 However, in spite of the recent progress, it is not yet well understood which factors at molecular scale determine the growth of organic heterostructures, ultimately allowing control and fabrication of desired organic nanostructures. It is clear that accurate experimental data of the growth and structure of the two materials at the interfaces is a necessary key step for advancing the knowledge. We have recently reported that the growth of organic heterostructures based on di-indenoperylene (DIP) deposited on fluorinated copper-phthalocyanines (F16CuPc) can easily be controlled by the growth temperature, ranging from smooth, layered heterostructures, to the formation of highly crystalline DIP islands with tunable size and density.35,36 By in situ X-ray diffraction, we revealed that the formation of three-dimensional(3D)islandsisassociatedwithaStranski-Krastanov type of growth (i.e., a morphology of 3D islands on top of a smooth wetting layer), which involves a structural change of the underlying F16CuPc film in proximity to the organic interface (interfacial reconstruction).36 However, no information is yet  2009 American Chemical Society

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available neither on the kinetics of such novel growth process, nor on the general applicability of this growth scenario. In order to answer these open questions, we have now studied heterostructures of DIP deposited on fluorinated cobalt phthalocyanines (F16CoPc). The similarity between F16CoPc and F16CuPc, which differ only in the central metal atom, makes this study especially interesting, because it allows a direct comparison of the observed growth behavior. In contrast to the nonfluorinated “sister” molecule CoPc, F16CoPc has so far been scarcely studied,37,38 and in turn, its crystalline structure in thin films deposited on SiO2 is still unknown. Note that molecular fluorination leads to an increase in the ionization potential, and to a favored n-type semiconducting behavior.39-43 Hence, DIP/ F16CoPc heterostructures form an organic p-n junction of potential interest for organic-based devices. By combination of in situ surface sensitive X-ray diffraction, X-ray reflectivity and atomic force microscopy (AFM) we unravel the first stages of the reconstruction during formation of the organic DIP/F16CoPc interface and provide direct evidence of the intimate relation between the various growth morphologies of the heterostructures and the structure at the interface. 2. Experimental Section A. Sample Preparation. The organic heterostructures have been prepared by organic molecular beam deposition on Si(100) wafers covered by their native oxide. The substrates were cleaned with acetone and ethanol, then annealed under ultra high vacuum (UHV) up to 500 °C, and subsequently cooled to the growth temperature. The molecules were commercially purchased and purified twice by gradient sublimation before use. The deposition was performed by thermal evaporation from home-built Knudsen cells and controlled by means of a calibrated quartz crystal monitor (QCM). The vacuum conditions during evaporation were typically in the low range of 10-8 mbar in the portable UHV chamber used for the X-ray diffraction experiments and in the range of 10-10 mbar for the samples characterized by AFM. B. X-ray Diffraction. In situ X-ray reflectivity and diffraction measurements have been performed in a specially designed portable UHV chamber, equipped with a beryllium window, a QCM, two Knudsen cells, and a heatable sample holder. The measurements have been carried out at the beamline ID10B of the ESRF (Grenoble), with a wavelength of 1.56 Å. The reflectivity profiles correspond to θ-2θ scans. The diffuse background was eliminated subtracting the recorded offset intensity, and the purely specular intensity has then been fitted with the Parratt algorithm,44 employing a multilayer model in which each individual layer is characterized by its thickness, electronic density, absorption coefficient, and the root-meansquare roughness of the associated interfaces. First the data associated with the bare substrate were fitted with a layer for Si and another for SiO2. The substrate parameters were kept fixed for the fitting of the subsequently evaporated films, for which additional layers were added whenever necessary to obtain satisfactory fits. The fit parameters for a certain thickness were used as starting parameters for the subsequent fits of the next growth step. X-ray reflectivity data from laterally heterogeneous samples can be modeled by the addition of the contributions corresponding to each of the laterally differentiated parts.45 For the samples with 2.4 and 4.8 ML DIP, in which the F16CoPc film is covered by a DIP wetting layer and 3D islands, the data have been modeled by the intensity corresponding to a smooth, laterally homogeneous Si/SiO2/F16CoPc/DIP heterostructure plus the

Figure 1. (a) X-ray reflectivity data (symbols) and fit (solid line) of a 5 ML thin F16CoPc film on SiO2 at 120 °C. The inset corresponds to the fitted electron density profile, in which the background colors help with the distinction of the different layers. From right to left: Si (dark gray), SiO2 (light gray), interfacial F16CoPc (red), and upright standing F16CoPc (orange). (b) In-plane GIXD spectrum. The observed reflections are labeled according to the β and βbilayer structure. The inset corresponds to a topography AFM image of a 3 ML thin F16CoPc film on SiO2.

contribution of those areas covered with rough DIP islands.45 The incoherent addition of the intensities is further justified by the large separation between DIP islands. The GIXD data were measured performing detector scans in the surface plane, with the grazing angle fixed at the critical angle value of the F16CoPc (Rc ) 0.19°). The data were first appropriately transformed from angular to reciprocal space coordinates, which were in turn used for the determination of the real space unit cell parameters.46,47 C. Atomic Force Microscopy. Measurements of the resulting morphology of the heterostructures were performed in tapping and contact mode under ambient conditions with a commercial AFM (Nanotec Electronica S.L.) and further analyzed with the freeware WSxM.48 3. Results and Discussion A. Growth of F16CoPc Thin Films. Figure 1a shows the X-ray reflectivity data obtained from a 5 monolayer (ML) thin film of F16CuPc grown on SiO2 at a low deposition rate (∼1 Å/min) and for a substrate temperature of 120 °C. From the fit of the reflected X-ray intensity (for details see the experimental part), we determine the electronic density profile projected along the surface normal, F(z), thereby obtaining accurate information on the structure and morphology of the multilayer.44 The leastsquares fit (solid line) corresponds to the density profile depicted in the inset. It reveals a layered structure for the film, with an interlayer spacing of 14.70 Å. This spacing is in the range of the lateral molecular dimensions and thus implies an upright standing configuration of the molecules within the layers. The electron density profile shows a rather smooth F16CoPc film, as only the two topmost layers (5th and 6th) are incomplete with

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TABLE 1: Unit Cell Parameters of the Two Crystalline Structures of F16CoPc (and Those Previously Determined for F16CuPc) and That of the DIP Film on Top F16CoPc βbilayera a [Å] b [Å] c [Å] R [deg.] β [deg.] γ [deg.] layer height [Å] V/unit cell [Å3] molec./unit cell density [gr/cm3]

F16CoPc β

F16CuPc βbilayerb

F16CuPc βb

DIP 8.55 ( 0.07 7.13 ( 0.07 16.52 ( 0.1 92.5 ( 0.2 90 90 16.5 ( 0.1 1006 ( 14 2 1.32 ( 0.02

14.61 ( 0.05 3.19 ( 0.05

20.43 ( 0.05 4.84 ( 0.09

14.61 3.31

20.26 ( 0.06 4.87 ( 0.04

∼90 14.70 ( 0.03 685 ( 12 1 2.08 ( 0.04

83.4 ( 0.4 14.70 ( 0.03 1444 ( 27 2 1.98 ( 0.04

∼90 14.1 ( 2.6 682 ( 126 1 2.10 ( 0.39

84.1 ( 0.2 14.3 ( 0.05 1403 ( 14 2 2.04 ( 0.02

a

A rectangular unit cell is assumed because the exact angle cannot be definitely concluded from only two reflections. The parameters “a” and “b” (and their calculated error bars) are based on this assumption, but would vary for changes in “γ”. b Values as taken from ref 51.

a partial coverage of ∼80% and ∼20%, respectively. In addition, the reflectivity data give unambiguous evidence for the formation of a low density interfacial layer of 6.8 Å thickness between the SiO2 and the first layer of upright standing molecules (marked with a red background color in the density profile). A similar layer was found in the growth of F16CuPc on SiO2 and was assigned to a disordered, poorly packed layer composed, in average, of two layers of lying down molecules at the interface.49,50 Figure 1b shows the complementary grazing incidence X-ray diffraction (GIXD) data, which accesses the crystalline structure in the surface plane of the F16CoPc film. In perfect analogy to F16CuPc films on SiO2,51 the data indicate the presence of two different structures, denoted βbilayer and β in what follows. Their unit cell parameters are summarized in Table 1 together with those previously obtained for F16CuPc. Systematic X-ray measurements for different film thicknesses indicate that the first two layers of upright standing molecules arrange in the βbilayer structure. Upon subsequent growth, polymorphism sets in and both the βbilayer and β structures grow simultaneously, favoring the latter as the coverage increases. The resulting film structure at room temperature and high temperature growth conditions are the same. The inset in Figure 1b shows a topographic atomic force microscopy (AFM) image of a 3 ML thin F16CoPc film on SiO2 grown at 120 °C evidencing elongated crystallites (of few hundreds of nm length) randomly oriented in the surface plane, as for F16CuPc films.49,52 B. Growth of DIP/F16CoPc Heterostructures. The heterostructure growth has been performed by stepwise deposition of DIP onto previously grown F16CoPc films and characterized in situ by X-ray reflectivity and GIXD after each of the growth steps. Figure 2a shows the X-ray reflectivity data associated with a heterostructure grown at 120 °C. The underlying F16CoPc film is 2.5 ML thick, and DIP has been deposited in steps corresponding to 0.4, 0.8, 1.2, 1.6, 2.4, and 4.8 ML of nominal thickness. The electron density profiles (obtained from fitted X-ray reflectivity data) reveal the evolution of the film surface and organic-organic interface morphology during the heterostructure growth (for sake of clarity, only the density profiles corresponding to selected DIP thicknesses are depicted in Figure 2b). The profile for the pristine F16CoPc film discloses the same low density interfacial layer, followed by two complete layers of upright standing molecules and an incomplete layer with ∼50% of coverage (Figure 2b). Upon deposition of 0.4 ML DIP, there is an increase in the electron density of the incomplete upmost layer, caused by the nucleation and subsequent growth of DIP from the edges of the F16CoPc islands, which is facilitated by the comparable layer heights. Upon complete filling of this uppermost layer, the growth of the next DIP layer

Figure 2. (a) X-ray reflectivity data (symbols) and fits (solid lines) measured in situ at different steps of the heterostructure growth at 120 °C. The inset represents the diffraction geometry. (b) Electronic density profiles obtained from fits to the reflectivity measurements. For the sake of clarity, only the profiles corresponding to selected coverages (following the color code of panel (a)) are shown. The electron density of the DIP islands at higher coverages is not represented for the sake of clarity. The background colors help with the distinction of the different layers. From right to left: Si (dark gray), SiO2 (light gray), interfacial F16CoPc (red), F16CoPc (orange), DIP (green). (c) AFM topography image (2 × 2 µm2) and schematic representation of a DIP/ F16CoPc heterostructure grown at 120 °C, corresponding to 6 ML DIP on 3 ML F16CoPc.

starts. After all the initially exposed F16CoPc is covered by a DIP wetting layer, the growth of 3D DIP islands sets in, monitored by the emergence of a sharp DIP(001) Bragg peak accompanied by only minor changes in the Kiessig fringes at lower qz values (Figure 2a). From the width of the Bragg peak (∆qz) we estimate (according to disl ≈ 2π/∆qz) that the DIP islands have average heights of 15 and 30 ML after deposition of 2.4 and 4.8 ML DIP, respectively. The Stranski-Krastanov growth behavior of DIP, resulting in textured organic heterostructures has been corroborated by AFM measurements. Figure

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Figure 3. (a) GIXD data at different stages of the heterostructure growth at 120 °C. The DIP reflections are appropriately labeled. The shift of the two main F16CoPc reflections (βbilayer Bragg peaks) is marked by the vertical lines and the arrows. The inset represents the diffraction geometry. (b) Intensity vs DIP coverage of the “as grown” (black squares) and reconstructed (red circles) F16CoPc (01) βbilayer Bragg peak, normalized to the initial intensity and to the final saturation intensity values, respectively. The green line represents the evolution of the percentage of surface covered by DIP as obtained from the reflectivity data. The schemes at the right depict the suggested reconstruction scenario, which takes place only in those areas covered by DIP.

2c shows by way of example the topography of a sample with 6 ML DIP deposited onto 3 ML F16CoPc at 120 °C. DIP islands are observed, with average heights of ∼23 nm, corresponding to roughly 14 ML of DIP.53 The in-plane structure corresponding to each of the previous growth steps has been further monitored by GIXD (Figure 3a). Upon DIP growth, a structural reconstruction of the underlying F16CoPc film is uncovered by the disappearance of the βbilayer reflections and appearance of new peaks at q| ) 0.53 Å-1 and q| ) 1.90 Å-1(associated with real space lattice spacings of d| ) 11.85 Å and d| ) 3.31 Å, respectively). This structural change has been examined in detail during the first stages of the DIP/ F16CoPc interface formation. The well defined position of the Bragg peaks (of the original and new structure) remains fixed during the reconstruction process. Therefore, a gradual transition among the different packing structures (which would imply a continuously shifting peak position) can be ruled out, and an abrupt change in the F16CoPc packing is confirmed. The intensity of the in-plane Bragg peaks of the original and new structure change gradually with opposing trends. Because the intensity of the Bragg peaks is directly related to the amount of material ordered with a specific structure, we can quantify the amount of reconstructed F16CoPc as a function of DIP coverage. Figure 3b depicts the normalized intensity of the βbilayer F16CoPc (01) reflection, for the original and the reconstructed structures vs DIP coverage. The peak of the original structure decreases at the same rate as the new peak increases, confirming the transformation of the βbilayer into the reconstructed structure. In addition, the percentage of surface covered by DIP, as obtained from the previously discussed electron density profiles, presents a striking agreement with the evolution of the new peak

J. Phys. Chem. C, Vol. 113, No. 11, 2009 4237 corresponding to the reconstructed F16CoPc. These findings provide evidence that F16CoPc undergoes a reconstruction in those areas covered by DIP, which increase as more DIP is deposited. Once the F16CoPc surface is covered by a complete DIP wetting layer, the whole underlying F16CoPc film is reconstructed. Thus, additional DIP deposition does not cause further change in the amount of reconstructed F16CoPc and the intensity of the corresponding Bragg peaks remains constant (except for the expected decrease due to absorption). This growth scenario is schematically presented in Figure 3b. In contrast to the changes observed for F16CoPc, the analysis of the peak positions of DIP shows no evidence for strain or structural changes. The DIP wetting layer and DIP islands present the same structure reported for thin-films.54-56 Strain relief is the most usual reason for Stranski-Krastanov type of growth in inorganic heteroepitaxy. In this case, in the absence of strain, the transition to 3D island growth is expected to arise from a change in the interfacial energy landscape as a result of the F16CoPc reconstruction, driven by the strong C-F · · · H-C interactions at the organic-organic interface.57,58 The similarities observed between the growth of heterostructures of DIP on F16CoPc and on F16CuPc,36 two molecules that present similar in-plane film structures before and after the reconstruction, suggest an epitaxial relationship between the underlying phthalocyanine layers and the DIP films as a driving force for the reconstruction. However, since the samples present a twodimensional powder, this can not be determined by the present experiment. In order to determine how many layers of F16CoPc are involved in the reconstruction, we have carried out measurements on heterostructures with a systematic variation of the underlying F16CoPc film thickness. Figure 4a depicts GIXD data obtained from heterostructures consisting of DIP deposited atop a 2.5 ML and a 5 ML thin film of F16CoPc. We find that the reconstruction of the underlying F16CoPc also takes place for thicker F16CoPc films. However, while a complete reconstruction is observed for the thinnest films, contributions of both the original and the reconstructed structures are observed for thicker films. Because of the direct proportionality between Bragg peak intensity and the amount of material with the corresponding crystalline structure, we conclude from the intensity ratio of the original βbilayer Bragg peak before and after the reconstruction that only 2-3 ML adjacent to the organic-organic interface are involved in the reconstruction. The temperature dependence of the growth of the heterojunctions has also been examined. For this, we have grown heterostructures of ∼2 ML DIP on top of 2.5 ML F16CoPc at different substrate temperatures ranging from 30 to 120 °C. The GIXD data measured in situ at the different temperatures are summarized in Figure 4b.59 Whereas at 80 °C and higher temperatures, there is a complete reconstruction of the underlying F16CoPc film upon DIP growth, at 30 °C, the F16CoPc film remains unchanged. At 70 °C, the majority of the F16CoPc film still exhibits the pristine structure, with only a small contribution of the reconstructed one. Hence, the critical temperature to activate the reconstruction process is placed within a temperature window between 70 and 80 °C. AFM measurements on heterostructures grown at different temperatures show smooth morphologies at low growth temperatures and pronounced islands at high temperatures (Figure 4b), thereby proving the direct correlation between the Stranski-Krastanov growth of DIP 3D islands and the F16CoPc reconstruction process. The full structure and crystalline coherence of the DIP islands have been determined by GIXD measurements employing a

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Figure 4. (a) Dependence of the DIP/F16CoPc heterostructure growth on the F16CoPc thickness. The black vertical arrows mark the remaining peaks corresponding to the original unreconstructed F16CoPc structure. A schematic representation of the resulting heterostructure is shown at the right for the 5 ML F16CoPc example, evidencing the reconstruction of only the two upper F16CoPc layers in proximity to the DIP. (b) GIXD spectra of heterostructures corresponding to ∼2 ML DIP on 2.5 ML F16CoPc at different growth temperatures. The peak positions of the original and the reconstructed F16CoPc structures are marked by black and red vertical lines, respectively. AFM topography images (3 × 3 µm2) and a schematic heterostructure representation are included at the right for low (30 °C) and high (120 °C) temperature examples.

position sensitive detector (PSD) to access a larger portion of the reciprocal space (as schematically represented in the inset of Figure 5a). For the different in-plane peak positions q| (horizontal axis), the intensity distribution along the rods in qz is also obtained (vertical axis), which contains information about the three-dimensional crystallographic structure of the films.46,60-65 Such data are shown in Figure 5a, measured on the heterostructure with 4.8 ML DIP on 2.5 ML F16CoPc grown at 120 °C. Because of the reduced thickness of the F16CoPc film, comprising only about three layers, the F16CoPc rods present streaks at low qz values (this a signature of the two-dimensional crystallinity of the film).66 In contrast, the DIP rods exhibit higher order Bragg diffraction peaks as result of the crystalline island growth.66 From the positions of the peaks, the unit cell parameters of the three-dimensional thin-film structure of DIP have been determined for the first time (summarized in Table 1). The thin-film structure of DIP differs considerably from the structure of DIP bulk crystallites,67 evidencing the crucial role of the substrate. The width of the Bragg peaks (∆q) along a specific direction is related to the crystalline domain sizes d along that direction by d ≈ 2π/∆q. Since X-ray reflectivity data contain no information about the in-plane structure, the width of the observed Bragg peaks reflects the thickness of the whole DIP islands (disl.), as previously discussed in Figure 2a. In contrast, the width along qz of the Bragg diffraction peaks observed in in-plane rods (see the rod profiles in Figure 5a) gives insight

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Figure 5. (a) Two-dimensional GIXD data corresponding to 4.8 ML DIP on 2.5 ML F16CoPc grown at 120 °C. The upper spectrum corresponds to the integrated intensity along qz, in which the various rods corresponding to F16CoPc and DIP are labeled accordingly. The spectra at the right side correspond to the intensity along each of the labeled DIP rods. The positions for the observed DIP Bragg reflections as calculated from the optimized unit cell parameters are overlaid on the graph, shifted to higher q| values by 0.7 Å-1. The measurement geometry is schematically depicted in the inset. (b) Schematic representation of the heterostructure, evidencing the presence of multiple coherently ordered crystalline domains within each of the islands.

into the thickness of the DIP islands that are coherently ordered (dcoh). Our data reveal that DIP islands with average height of disl. ≈ 30 ML (Figure 2a) exhibit coherently ordered crystalline domains of only dcoh ≈ 8 ML. Thus, as schematically outlined in Figure 5b, the crystal structure propagates coherently, on average, only to eight subsequent layers, and the resulting DIP islands are not single crystalline. 4. Conclusions We have disclosed an intimate relationship between the structure at the organic-organic interface and the growth and morphology of p-n heterojunctions composed of DIP on F16CoPc. Growth at temperatures above 70 °C results in a reconstruction of the underlying F16CoPc layers closest to the organic interface, concomitant with the formation of highly crystalline DIP 3D islands via Stranski-Krastanov growth. In situ X-ray reflectivity and GIXD investigations of the initial stages of the DIP/F16CoPc interface formation reveal an abrupt change in the F16CoPc packing in those areas covered by DIP. The amount of reconstructed F16CoPc at the organic interface increases during DIP deposition until total completion of the DIP wetting monolayer. With further DIP deposition, island growth sets in and F16CoPc becomes buried under the growing DIP islands. For growth temperatures below 70 °C, smooth heterostructures form with unmodified F16CoPc structure at the organic-organic interface. The similarities observed with the growth of DIP on F16CuPc highlight the generality of this novel growth scenario to different fluorinated phthalocyanines and the important role of C-F · · · H-C interactions at the organic-organic interface. This investigation not only provides a controllable route (via growth temperature) to finely tune the structure and morphology of organic p-n heterostructures, but also gives new insight into the mechanisms of organic heterojunction formation and their

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