Impact of interface geometric structure on organic–metal interface energetics and subsequent films electronic structure

Journal of Electron Spectroscopy and Related Phenomena 174 (2009) 28–34

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Impact of interface geometric structure on organic–metal interface energetics and subsequent films electronic structure H. Yamane a,∗ , K. Kanai b,1 , Y. Ouchi a , N. Ueno c , K. Seki a a b c

Department of Chemistry, Nagoya University, Nagoya 464-8602, Japan Research Center for Materials Science, Nagoya University, Nagoya 464-8602, Japan Graduate School of Advanced Integration Science, Chiba University, Chiba 263-8522, Japan

a r t i c l e

i n f o

Article history: Available online 14 March 2009 PACS: 72.80.Le 68.43.−h 73.20.At Keywords: Organic semiconductor Organic–metal interface Electronic structure Geometric structure Angle-resolved ultraviolet photoemission spectroscopy

a b s t r a c t We present the electronic structure of various pentacene thin films grown on Au(1 1 1), Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0) surfaces studied by angle-resolved ultraviolet photoemission spectroscopy using synchrotron radiation. A systematic variation of the metal surface such as the substrate metal and its surface symmetry allows a comprehensive discussion on the correlation between the electronic structure and the interface geometric structure. In the monolayer regime, we observed the evidence of the formation of the organic–metal interface state depending on the metal surface, i.e., the interface geometric structure. This evidence is explained by the different organic–metal and intermolecular interactions, which originate from the hybridization of the molecular orbitals with the metal wavefunction. These interface geometric and electronic phenomena can be a seed for the subsequent film growth and resultant films electronic structure. © 2009 Elsevier B.V. All rights reserved.

1. Introduction Recently, extensive research has been carried out for applying organic semiconductors to optoelectronic devices such as lightemitting diodes, field-effect transistors, and solar cells. In all of these devices, organic interfaces such as organic–metal (O–M) and organic–organic (O–O) can be regarded as a field in which electrons and holes play various roles in functionalities of organic devices, e.g., charge injection/extraction, charge separation, and charge transport. Among these, the O–M interface energetics of the molecular energy levels near the Fermi level (EF ) such as the highest-occupied and lowest-unoccupied molecular orbital (HOMO and LUMO) as well as the vacuum level (VL), i.e., energy level alignment, is of fundamental importance in discussing barrier heights at the interface [1,2]. However, it is generally not easy to estimate the O–M interfacial electronic structure precisely just by comparing the molecular ionization energy, the molecular electron affinity, and the work function (˚m ) of metals since there is a possibility of

∗ Corresponding author. Present address: Institute for Molecular Science, Okazaki 444-8585, Japan. Tel.: +81 564 55 7394;fax: +81 564 55 7391. E-mail address: [email protected] (H. Yamane). 1 Present address: Research Core for Interdisciplinary Sciences, Okayama University, Okayama 700-8530, Japan. 0368-2048/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.elspec.2009.03.002

the presence of interface-specific electronic state, namely interface state, as well as the interface dipole layer () [1,2]. These interface phenomena seriously dominate and modify the interface energetics and hence device performance. The O–M interface state has (i) intrinsic nature due to the formation of organic interfaces or (ii) extrinsic one due to the presence of defects and impurities in the film and/or the interface. Thereby, we may be able to manipulate the interface energetics by controlling the interface state. In the case of O–M interfaces, the interface states observed in ultraviolet photoemission spectroscopy (UPS) spectra, although diverse, can be categorized into following two main structures: (A) tailing into the HOMO–LUMO gap and (B) new-peak structure. The interface state with tailing structure (type A) is formed by the molecular level broadening due to O–O [3] and/or O–M [4] interaction as depicted in Fig. 1 (a) and (b). In this scenario, the HOMO and LUMO are simply broadened to form continuous levels in the original HOMO–LUMO gap of the molecule, and the two electrons in the original HOMO occupy these states up to the topmost occupied level in the induced continuous levels, which is called by the chargeneutrality level (CNL). The initial energy difference between the EF of the metal and the CNL of the organic semiconductor is canceled to achieve thermodynamic and electrical equilibrium with the alignment of the top occupied states at both sides of the interface. On the other hand, the interface state with new-peak structure (type B) has been observed as a result of the charge transfer (CT) [5–8]

H. Yamane et al. / Journal of Electron Spectroscopy and Related Phenomena 174 (2009) 28–34

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were studied by other groups in detail [15–20]. We observed the evidences of the formation of the interface state, depending on the substrate metal and its surface symmetry. The observed difference in the formation of the Pn–metal interface state is caused by the different O–M interaction and the resultant substrate-mediated intermolecular interaction. Furthermore, we found that the geometric and electronic structure of the first monolayer film can be a seed for the subsequent film growth and the resultant films electronic structure. 2. Experimental

Fig. 1. Possible origins of the O–M interface state: molecular level broadening due to (a) O–O interaction and (b) O–M interaction, and new-peak formation due to (c) O–M charge transfer and (d) molecular level splitting by orbital hybridization.

and the molecular level splitting due to the strong O–M interaction [9–11] as depicted in Fig. 1(c) and (d). The formation of the CT state is qualitatively expected for the combinations of strong acceptor (or donor) molecules and low (or high) ˚m metals [1,2,5–8]. In this case, the interfacial electronic structure is explained by the electronic structure of molecules in the anion (or cation) state. Furthermore, it has recently been proposed that the magnitude of CT is seriously affected by the O–M bonding distance and the distortion of the molecule upon adsorption [7,8]. When the O–M interaction becomes much stronger, origin of the interface state shifts to the molecular level splitting. In this case, the interfacial electronic structure is completely different from the electronic structure of molecules in the neutral, anion, and cation states such as the positions and symmetries of molecular energy levels. These new-peak structures are formed by the occupation/inoccupation and the splitting of former unoccupied/occupied levels due to CT and the orbital hybridization at the interface. Such interfacial charge phenomena induce a shift of the chemical potential in organic semiconductors. Other possible origins of the interface state are, for example, the formation of chemical reaction products [12], polaron/bipolaron [13], and metal-induced gap states by the penetration of metal wavefunction into organic layer [14]. In order to obtain the general picture of organic interfaces, due to complicated nature and mixed origins of the above-mentioned interface states, the systematic and quantitative accumulation of the reliable experimental data for well-defined interfaces is necessary. In this work, by angle-resolved UPS (ARUPS) using synchrotron radiation, we studied systematically the valence electronic structure of various thin films of pentacene (Pn), one of important organic semiconductors, grown on Au(1 1 1), Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0) surfaces, in which the interface geometric structures

The experiments were performed at the BL8B2 of the synchrotron radiation facility UVSOR at the Institute for Molecular Science, Okazaki, Japan. The system consists of an organic-preparation chamber, a metal-cleaning chamber, and a measurement chamber, with base pressures of 9 × 10−8 , 1 × 10−8 , and 2 × 10−8 Pa, respectively. The single crystal substrates of Au(1 1 1), Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0) (purity 5N, MaTecK GmbH) were cleaned by repeated cycles of Ar+ ion sputtering (primary ion-beam energy of 2 keV and current density of 30 mA/cm2 under Ar pressure of 8 × 10−3 Pa) and annealing up to about 870 K in the metal-cleaning chamber. An inert surface of highly oriented pyrolytic graphite (HOPG) was also used in this study. The ZYA-grade HOPG substrate was cleaved in air just before loading into the metal-cleaning chamber and cleaned in UHV by heating at 620 K for 15 h. The cleanliness of the substrate surface was confirmed by the surface state and ˚m in the ARUPS spectra, and the order of atomic arrangements of the substrate surface was confirmed by low-energy electron diffraction (LEED). The Pn sample (purity 98%: Tokyo Chemical Industry Co. Ltd.) was purified by three cycles of sublimation in Ar gas stream of 13 Pa. The purified material was evaporated onto the clean surfaces at elevated temperature, which realizes the formation of flat-lying Pn monolayer films: 380 K for HOPG [21], 450 K for Au(1 1 1) and Cu(1 1 1), 470 K for Cu(1 0 0), and 500 K for Cu(1 1 0) [10]. Furthermore, by depositing the molecules onto these Pn monolayer films kept at about 350 K, we prepared Pn multilayer films. The film thickness and the deposition rate (less than 0.1 nm/min) were measured with a calibrated quartz crystal microbalance. The ARUPS measurements were performed by using a VGARUPS10 system with a hemispherical electron energy analyzer and a multichannel detector in the measurement chamber. The synchrotron radiation was monochromatized by a plane-grating monochromator [22]. In the present experiments, we set the total-energy and angular resolutions at about 100 meV and 0.8◦ , respectively. All ARUPS measurements were performed at the surface temperature of 300 K at the photon energy of h= 20 eV and the photon incidence angle of ˛= 60◦ . In order to estimate ˚m , the ARUPS spectra in the region of the secondary-electron cutoff were measured with a −5 V bias applied to the sample to detect the electrons with a kinetic energy close to 0 eV. 3. Results and discussion 3.1. Overview of the interface geometric structure Prior to showing the experimental results, we briefly summarize the previous works on the interface geometric structure of the Pn monolayer films on Au(1 1 1), Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0), as shown in Table 1. According to the previous structural experiments such as LEED and scanning tunneling microscopy (STM) [15–20], the Pn molecules on Au(1 1 1), Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0) orient with their molecular plane parallel to the substrate surface, while the lateral molecular ordering at these interfaces depends on the substrate metal and its surface symmetry. By depositing the

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Table 1 Summary of the interface geometric structure of the flat-lying Pn monolayer films on Au(1 1 1), Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0): degree of ordering, number of domain, and bonding distance for the carbon atom (dC ).

Pn/Au(1 1 1) Pn/Cu(1 1 1) Pn/Cu(1 0 0) Pn/Cu(1 1 0)

Ordering

Domain

dC /nm

Short range Short range Long range Long range

Multi Multi Multi Single

N/A 0.234 0.233 N/A

molecules onto the Cu(1 1 0) surface at 400–600 K, the molecules form a highly ordered monolayer with the stripes of Pn along the [0 0 1] substrate direction, retaining the orientation of the molecule ¯ substrate direction [15,16]. The with its long axis along the [110] LEED pattern of the Pn/Cu(1 1 0) interface reveals the presence of a single-domain film structure [16]. When the surface symmetry increases from twofold Cu(1 1 0) to fourfold Cu(1 0 0), the LEED pattern shows a multi-domain film structure [17]. In the following, we call these interfaces as long-range ordered interfaces, which show clear LEED pattern as well as STM. On the other hand, on the Cu(1 1 1) and Au(1 1 1) surfaces with sixfold symmetry, the LEED pattern shows the absence of any diffraction spots, although the STM local image reveals the presence of the ordered structure [18–20]. This evidence may be explained by its incommensurate structure appearing in various domains. We call these interfaces as short-range ordered interfaces in the following. It has been recognized that the O–M bonding distance at the interface is an important factor in discussing the formation of the interface state [4,7–9] and the interface dipole [23,24]. However, it is in principle difficult to estimate the O–M bonding distance just by using the conventional density-functional theory, which cannot adequately treat van der Waals interactions at the interface. The van der Waals interaction as well as the exchange–correlation interaction is proposed as a key to understand the interface energetics and the adsorption geometry [25–27]. For the accurate theoretical determination of the bonding distance, one has to compare with the bonding distance obtained by an experimental determination method, X-ray standing wave (XSW) [28]. The XSW experiment, which gives site-specific information on the bonding distance between adsorbates and substrate surfaces, has

recently been applied to the O–M interface [7,8,28–31]. Among these, Koch et al. reported the bonding distance for the carbon atom (dC ) at the Pn/Cu(1 1 1) interface to be 0.234 nm [29] with fair agreement on the Pn–Cu(1 0 0) bonding distance of 0.233 nm calculated by Ferretti et al. [9]. The Pn–Cu(1 1 1) bonding distance of 0.234 nm is shorter than those in the case of other organic monolayer films on Cu(1 1 1) studied by XSW such as perfluoro Pn (dC = 0.298 nm) [29], hexadecafluoro copper phthalocyanine (F16 –CuPc) (dC = 0.261 nm) [30], and perylene-3,4,9,10-tetracarboxylic acid dianhydride (PTCDA) (dC = 0.266 nm) [31]. On the other hand, from the thermal desorption experiments by the group of Witte and Wöll [16,20,32], the Pn sublimation enthalpy amounts to 157 kJ/mol leading to an onset of evaporation at temperatures of about 390 K under UHV, while Pn monolayer films chemisorbed on metal surfaces remain stable up to much higher temperatures: 500 K for Au(1 1 1) [20], 520 K for Ag(1 1 1) [32], 750 K for Cu(1 1 1) [33], and 770 K for Cu(1 1 0) [16]. The difference in the desorption temperature for various Pn monolayer films suggests that the Pn–Au(1 1 1) bonding distance is probably longer than the Pn–Cu(1 1 1) bonding distance of 0.234 nm. The different O–M bonding distance may lead the different evidence in the formation of the interface state as in the case of the PTCDA/metal interface [7]. 3.2. Formation of the metal–surface-dependent interface state Fig. 2 shows the ARUPS spectra for the various Pn/metal interfaces as a function of the photoelectron emission angle () or the azimuthal angle (), where the surface Brillouin zone for the (1 1 1), (1 0 0), and (1 1 0) surfaces with the definition of  are also shown: (a and b) the  dependence of the Pn/Au(1 1 1) and the ¯ direction), (c and d) Pn/Cu(1 1 1) interfaces fixed at  = 0◦ (¯ – M the  dependence with a step of 3◦ for the Pn/Cu(1 0 0) and the Pn/Cu(1 1 0) interfaces fixed at  = 58◦ , and (e) the Pn/HOPG interface fixed at  = 38◦ ( = integrated by the azimuthal disorder of the single-crystal domains in the HOPG surface). The abscissa is the binding energy (Eb ) relative to the substrate EF . As a reference, in Fig. 2(a)–(e), we also show the UPS spectrum of a gas-phase Pn molecule [34], which is shifted to align the lowest Eb peak to that of the Pn monolayer films. For both the Pn/Au(1 1 1) and Pn/Cu(1 1 1) interfaces (Fig. 2(a) and (b)), one can see the sharp

Fig. 2. The ARUPS spectra for the various Pn/metal interfaces as a function of the photoelectron emission angle () or the azimuthal angle (): (a and b) the  dependence ¯ direction), (c and d) the  dependence with a step of 3◦ for the Pn/Cu(1 0 0) and the Pn/Cu(1 1 0) of the Pn/Au(1 1 1) and the Pn/Cu(1 1 1) interfaces fixed at  = 0◦ (¯ – M interfaces fixed at  = 58◦ , and (e) the Pn/HOPG interface fixed at  = 38◦ . The surface Brillouin zone for the (1 1 1), (1 0 0), and (1 1 0) surfaces with the definition of  are also shown.

H. Yamane et al. / Journal of Electron Spectroscopy and Related Phenomena 174 (2009) 28–34

peak at about Eb of 0.4 eV in the normal-emission ( = 0◦ ) spectra of the bare surface (gray line). This sharp peak is ascribed to the Shockley-type surface state of metal surfaces [35,36], and is quenched upon adsorption of the Pn monolayer film. Recently, using ARUPS, scanning tunneling spectroscopy, or two-photon photoemission spectroscopy, surviving of the Shockley-type surface state upon adsorption of monolayer films has been reported for the interfaces of rare-gases/fcc(1 1 1) (fcc = Au, Ag, and Cu) [36], nalkane/Au(1 1 1) [37], C60 /Cu(1 1 1) [38], PTCDA/Au(1 1 1) [39], and PTCDA/Ag(1 1 1) [40,41]. Interestingly these interfaces show the modification of the Shockley-type surface state in the peak position at the ¯ point and its dispersion parabola. In the case of the weakly interacting interfaces, the magnitude of the shift at the ¯ point is in the range of 200 meV [36–39], while the strongly chemisorbed PTCDA/Ag(1 1 1) interface shows large upshift of about 660 meV [40,41]. Although origin of the modification of the Shockley-type surface state upon adsorption is still in debate, such a phenomenon may be a probe for the magnitude of the O–M interaction at the interface [37]. Hence, these results suggest that the O–M interaction at both the Pn/Au(1 1 1) and Pn/Cu(1 1 1) interfaces may be strong since the Shockley-type surface state is quenched from the valence level region upon adsorption of the Pn monolayer film as in the case of the PTCDA/Ag(1 1 1) interface [40,41]. We now turn to discussion on the appearance of the Pn-induced peaks at the Pn/metal interfaces. The first example is the case of the laterally short-range ordered interfaces of Pn/HOPG, Pn/Au(1 1 1), and Pn/Cu(1 1 1). In the case of the Pn/HOPG interface displayed in Fig. 2(e), one can see a sharp HOMO-derived peak at Eb of 1.34 eV as reported in the earlier work [21]. The observed peak shows an asymmetric lineshape caused by a vibrational progression towards the higher Eb side as indicated by arrows, and the full-width at half-maximum (FWHM) including the vibration satellite is about 200 meV. The relative energy of the Pn-induced peaks such as HOMO and HOMO-1 agree well with the UPS spectrum of the gas-phase Pn molecule [21,34]. This is the reflection of the persistence of molecular characteristic in the film, and is the case of the weak physisorption. At the Pn/Au(1 1 1) interface (Fig. 2(a)), the Pn-induced peak appears at Eb of 0.96 eV. Although it is difficult to distinguish the other Pn-induced peaks due to the presence of the large Au5d bands in the ARUPS spectra, deeper lying Pn-induced peaks appear by changing  and h (not shown). Thereby, we found that the relative energy of the Pn-induced peaks for the Pn/Au(1 1 1) interface agrees with that for the gas-phase Pn molecule. Thus, we can assign the first Pn-induced peak at Eb of 0.96 eV to the HOMO level of the neutral Pn. However, the FWHM of the HOMO-derived peak is 450 meV, which is much broader than that for the Pn/HOPG interface. This broadening is ascribed to the molecular level broadening due to the weak O–M interaction as depicted in Fig. 1(b). The change in ˚m , i.e., VL shift (), upon adsorption is −0.85 eV, in which negative  indicates the lowering of VL. Since there are no evidences of the strong CT between Pn and Au(1 1 1), the magnitude of  may be ascribed to the charge redistribution due to the image effect (polarization of the electron cloud of molecules) [42] and the push-back effect (polarization of the electron cloud of metals) [1]. On the other hand, the Pn-induced electronic structure on Cu(1 1 1) shown in Fig. 2(b) is rather different from that on Au(1 1 1). The Pn-induced peaks on Cu(1 1 1) appear at Eb s of 0.70 and 1.32 eV with the energy separation of 0.62 eV, furthermore, the steplike structure appears at EF upon adsorption. These structures cannot be explained by the UPS spectra of the substrate and the gas-phase Pn molecule. As in the following, these structures are ascribed to the molecular level splitting due to the hybridization of the MOs with the metal wavefunction (Fig. 1(d)), which seriously modify the energies and symmetries of molecular energy levels. As the origin of the apparent difference in the interfacial electronic structure between the Pn/Au(1 1 1) and Pn/Cu(1 1 1) interfaces, we consider

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that the Pn–Au(1 1 1) bonding distance is probably larger than the Pn–Cu(1 1 1) bonding distance of 0.234 nm as we described in Section 3.1 in this paper. The  of −1.02 eV at the Pn/Cu(1 1 1) interface may originate from the effects of the orbital hybridization and corresponding CT as well as the charge redistribution due to the image effect and the push-back effect. Next, we discuss the second example: the laterally longrange ordered interfaces of Pn/Cu(1 0 0) and Pn/Cu(1 1 0). For the Pn/Cu(1 0 0) interface with the multi-domain film structure (Fig. 2(c)), the Pn-induced structures appear with nearly the same energy separation as for the Pn/Cu(1 1 1) interface. Dissimilarity of the Pn-induced structures on Cu(1 0 0) to that on Cu(1 1 1) is the increase in the peak intensity and its clear dependence, which is explained by the improvement on the symmetry selection rules in ARUPS due to the formation of the long-range ordered film structure. The single-domain Pn/Cu(1 1 0) interface in Fig. 2(d) also forms the interface state originating from the molecular level splitting as for the Pn/Cu(1 1 1) and Pn/Cu(1 0 0) interfaces. It is interesting to note that the resultant split levels at the Pn/Cu(1 1 0) interface show dispersive behavior with . Using the single-domain Pn/Cu(1 1 0) interface, we discuss the origin of the interface state at the Pn/Cu(1 1 1), Pn/Cu(1 0 0), and Pn/Cu(1 1 0) interfaces as well as the origin of the dispersive behavior observed only at the Pn/Cu(1 1 0) interface. When we consider the symmetry selection rules in ARUPS for detection of the photoelectrons in the plane of incident-light polarization under the present experimental setup, the initial state should be symmetric with respect to the incident plane. For example, the HOMO of the neutral Pn should show weak emission at  = 0◦ and strong emission at  = 90◦ ; however, the lowest Eb peak of the Pn/Cu(1 1 0) shows strong emission at  = 0◦ and no emission at  = 90◦ . Since the simple consideration mentioned above does not work, which is discussed in detail in Ref. [10], we propose that the Pn-induced structures at the Pn/Cu(1 1 1), Pn/Cu(1 0 0), and Pn/Cu(1 1 0) interfaces originate from the molecular level splitting due to the hybridization between the MOs and the metal wavefunction, which seriously modifies the energy and symmetry of molecular energy levels. This scenario is supported by the theoretical calculation done by Ferretti et al. [9]. Further discussion on the origin of the resultant split levels at the Pn/Cu interfaces is given in Section 3.3. Since the dispersive behavior in the interface states was observed not at the multi-domain Pn/Cu(1 0 0) interface but at the single-domain Pn/Cu(1 1 0) interface, we can straightforwardly correlate the dispersive behavior with the lateral molecular ordering. In fact, from the energy-versus-momentum, E(k), relation in the topmost Pn-induced peak along the ¯ – X¯ direction [10], the lateral lattice constant can be estimated to 1.6 nm, which agrees well with the lattice constant obtained from the geometrical studies using LEED and STM [15,16]. Furthermore, from the observed E(k) rela¯ tion along the ¯ – Xdirection, we estimated the effective mass of the hole (m∗h ) for the upper branch to be 0.24m0 at 300 K [10]. Such a light m∗h of 0.24m0 with the lateral lattice constant of 1.6 nm suggests the presence of the strong lateral intermolecular interaction. In the case of thin films of large ␲-conjugated planar molecules with flat-lying orientation, however, the lateral intermolecular interaction is in general dominated by the rather weak intermolecular ␴– ␴ interaction of less than ∼ 0.1 eV. For the discussion on the origin of the strong lateral intermolecular interaction, a key parameter may be the observed m∗h along the ¯ – X¯ direction of 0.24m0 , which is nearly the same to the effective mass of the Shockley-type surface state along the ¯ – Y¯ direction of 0.26 m0 at the clean Cu(1 1 0) surface [43]. Although there is no known surface state along the ¯ – X¯ direction [43], we would like to propose that the observed E(k) dispersion of the Pn/Cu(1 1 0) interface state originates from the substrate-mediated intermolecular interaction, which is caused by the hybridization between the MOs and the (modi-

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fied) metal wavefunction as well as the geometrical film structure/ ordering. 3.3. Evolution of the electronic structure with the film thickness It has been reported that the Pn molecules on Cu(1 1 0) exhibit various structural phases with the film thickness [16]. At the monolayer coverage, the Pn molecules form a long-range ordered single-domain structure with planar adsorption geometry. For thin multilayers (less than 2 nm), the molecules give a second structural phase with a laterally rotated of 21◦ and a vertically tilted of 28◦ structure with respect to the first monolayer structure. Finally, when exceeding a thickness of about 2 nm, crystalline films with an upright-standing molecular orientation can be obtained. The upright-standing Pn multilayer film on Cu(1 1 0) is the model of the thin-film phase crystalline structure of Pn [16]. Fig. 3 shows the film thickness dependence of the ARUPS spectra on the Pn/Cu(1 1 0) system. In the first monolayer (1 ML) film, as discussed already, we observed the distinctive electronic structure with five interface-specific peaks in the Eb range of 0–4 eV as labeled by asterisk in Fig. 3. This electronic structure corresponds well with that obtained from the theoretical calculation by Ferretti et al. [9]. According to their theoretical calculation [9], the first and second interface states are ascribed to the resultant split levels of the former LUMO level, and other three interface states are ascribed to the resultant split levels of the former HOMO level as indicated in the energy diagram in Fig. 3. These interface states rapidly get weaker when the second monolayer (2 ML ≈0.5 nm) film is grown on the 1-ML film. Such a decrease in the peak intensity is explained by the different electronic structure and the adsorption geometry between the 1-ML and 2-ML films that changes the symmetry selection rule in ARUPS. The electronic structure of the 1-nm film (≈3–4 ML equivalence) is slightly different from that of the 2-ML film. This is the reflection of the coexistence of the second structural phase in the 2-ML film and the subsequent overlayers structure with different molecular orientation. When the film thickness reaches to 2 nm, which is the branch point from slightly tilted to upright-standing molecular orientation [16], the contribution from the Cu 3d bands completely disappears. After the 2-nm film thickness, the lineshape of the HOMO-derived peak at around Eb of 1 eV changes gradually with increasing the film thickness. Finally at

the 11-nm film thickness, the HOMO-derived peak shows an asymmetric lineshape, which consists of two prominent components with the energy separation of 440 meV as indicated by downed arrows in Fig. 3. We found for the 11-nm film that the observed components in the HOMO-derived peak show continuous and periodic shift with  for both ¯ − X¯ and ¯ − Y¯ azimuths, which is discussed in detail in Ref. [11]. The lateral lattice constant obtained from this dispersive behavior agrees well with that obtained from the LEED of the upright-standing Pn crystalline film on Cu(1 1 0) [11]. These results lead a conclusion that the observed HOMO lineshape and its dispersive behavior of the 11-nm film originate from the energy-band dispersion, E(k), which is attributable to the strong intermolecular ␲– ␲ interaction in the upright-standing Pn crystalline film [11]. From the observed E(k)relation, we estimated m∗h ¯ and 1.86m0 (¯ – Y¯ ) at 300 K within the tightto be 3.02m0 (¯ – X) binding approximation [11]. Such a HOMO lineshape due to the strong intermolecular interaction has been observed also for various Pn crystalline films [44–46]. On the other hand, the growth of Pn multilayer film on clean Cu(1 1 0) substrates at room temperature (∼ 300 K) leads to rather rough film since the multilayers start to grow before completion of the monolayer [16]. In this case, the sharp LEED pattern and the dispersive behavior in the ARUPS spectra cannot be observed. The important point of this evidence is that the quality of the subsequent multilayer film could be improved significantly by first growing a well-ordered monolayer film. In order to examine the effect of the degree of ordering of first monolayer films on the subsequent film growth and its electronic structure, in Fig. 3, we show a normal-emission spectrum of the Pn multilayer film prepared on Cu(1 1 1). The HOMO-derived peak of the Pn multilayer film on Cu(1 1 1) is less intense and broader than that on Cu(1 1 0). Furthermore, the ARUPS spectrum shows the trace of the substrate Fermi edge, indicating a pronounced and distinct Stranski–Krastanov growth mode, which is also suggested by the thickness dependence of C 1s and Cu 2p core-level photoemission spectra [47]. In this case, the broad HOMO-derived peak involves the contribution from both the interface and the bulk with different polarization energies. Since the HOMO-derived peak of the Pn multilayer film on Cu(1 1 1) does not show the clear dispersive behavior, although broad, the energy-band dispersion may be negligibly small due to the weak intermolecular ␲– ␲interaction. The

Fig. 3. Thickness dependence of the ARUPS spectra, schematics of the thin-film growth, the LEED pattern for the 11-nm film with the incident-electron energy of 32 eV, and the energy diagram on the Pn/Cu(1 1 0) system. The ARUPS spectra were recorded at normal emission ( = 0◦ , black/thin curve) and at off-normal emission ( = 58◦ , red/thick curve). The normal-emission spectrum of the Pn multilayer film on Cu(1 1 1) is shown for comparison. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)

H. Yamane et al. / Journal of Electron Spectroscopy and Related Phenomena 174 (2009) 28–34 Table 2 Summary of the substrate work function ˚m , vacuum level shift (vs. substrate surface) , ionization energy Is , and hole injection barrier εh of monolayer and multilayer films of Pn on Cu(1 1 0), Cu(1 0 0), Cu(1 1 1), Au(1 1 1), HOPG [21,44], CuPc/HOPG [44], and SiO2 [44]. The Is and εh of the Pn multilayer films on Cu(1 1 1) and Au(1 1 1) were estimated from the bulk component by the peak fitting due to the co-existence of the surface and bulk contributions in the HOMO-derived peak (see Fig. 3). Substrate

Cu(1 1 0) Cu(1 0 0) Cu(1 1 1) Au(1 1 1) HOPG CuPc/HOPG SiO2

˚m /eV

4.60 4.61 4.92 5.50 4.44 4.52 4.32

Monolayer

Multilayer

/eV

Is /eV

εh /eV

/eV

Is /eV

εh /eV

−0.90 −0.92 −1.02 −0.85 0.01 −0.02 0.03

3.70 3.69 3.90 5.19 5.54 4.77 4.72

∼0 ∼0 ∼0 0.54 1.09 0.27 0.37

−0.90 −0.89 −1.07 −0.95 0.01 −0.05 0.05

4.50 4.79 5.00 5.45 5.15 4.78 4.80

0.80 1.07 1.15 0.90 0.70 0.31 0.43

observed broad lineshape may originate from the inhomogeneous film structure with different polarization energies, as observed in the CuPc/HOPG system [48]. Similar evidence is also reported for the case of Au(1 1 1) [20,49]. Such a difference in the crystalline structure of Pn multilayer films is also reflected in the ionization energy (Is ) since it is very sensitive to the packing structure, i.e., crystallinity, of the film: Is = 5.00 eV for Pn(thick)/Cu(1 1 1) and Is = 4.50 eV for Pn(thick)/Cu(1 1 0). These evidences again demonstrate the importance of the geometric structure of first monolayer films for the subsequent film growth and the resultant films electronic structure. In Table 2, we compile ˚m , , Is , and hole injection barrier (εh ) of monolayer and thick multilayer films of Pn on various substrate surfaces. For the monolayer films on Cu(1 1 1), Cu(1 0 0), and Cu(1 1 0), the εh are almost zero, indicating a metallic character of the film due to the formation of the interface state just at EF by the orbital hybridization and the corresponding CT at the interface. Other monolayer films have a semiconducting character. For the multilayer films, it seems that the εh is independent to ˚m , , and Is ; however, when we use wide energy-gap substrate surfaces such as SiO2 and CuPc-coated HOPG, the εh is obviously decreased [44]. Judging from this comparison, we consider that the energetics of thick multilayer films is dominated by the existence of the density of states near EF at the interface/surface, which may penetrate into subsequent organic layers, even if it rather small. When the organic film becomes thick enough to avoid from the interface/surface wavefunction penetration [50], the EF alignment may be achieved with the possible modification of the film structure for the total-energy relaxation in the system. 4. Conclusion In this work, we performed detailed ARUPS experiments on the electronic structure of the Pn thin films on various metal surfaces systematically. We have observed the metal–surface-dependent electronic structure at the Pn/metal interfaces with the following findings: (i) molecular level broadening at the Pn/Au(1 1 1) interface, (ii) molecular level splitting with possible modification of the orbital symmetry as well as the energy position at the Pn/Cu(1 1 1), Pn/Cu(1 0 0), and Pn/Cu(1 1 0) interfaces, and (iii) two-dimensional intermolecular band dispersion of the resultant split levels only at the highly ordered Pn/Cu(1 1 0) interface. These interface states can be deduced to originate from the different O–M interaction; among these, the strong O–M interaction can cause the molecular level splitting due to the hybridization of the MOs with the metal wavefunction, which may also lead to the band dispersion by the substrate-mediated intermolecular interaction. We have demonstrated that the interface geometric structure acts like a seed for not only formation of the interfacial electronic structure but also

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subsequent film growth and resultant films electronic structure. Furthermore, the existence of the density of states near EF at the interface/surface may dominate the energetics of thick multilayer films. The type of metal, its surface symmetry, and the specimen temperature during the film growth play crucial role in geometric and electronic characteristics of not only monolayer films but also subsequent multilayer films. Acknowledgements The authors and all members of Solid-State Chemistry Laboratory of Nagoya University are deeply grateful to the group leader, the late Professor Kazuhiko Seki. He was all the time kind and respectful to the scientific research and the education. We will miss him dearly. The present work was partly supported by the Grant-in-Aid for Scientific Research (S) from JSPS (19105005), the JSPS Core-to-Core Program (19002), 21st century Center-of-Excellence Program from MEXT (B11 and G04), and Global Center-of-Excellence Program from MEXT (B08 and G03). References [1] H. Ishii, K. Sugiyama, E. Ito, K. Seki, Adv. Mater. 11 (1999) 605. [2] W.R. Salaneck, K. Seki, A. Kahn, J.-J. Pireaux (Eds.), Conjugated Polymers and Molecular Interfaces: Science and Technology for Photonic and Optoelectronic Applications, Dekker, New York, 2002. [3] H. Fukagawa, S. Kera, T. Kataoka, S. Hosoumi, Y. Watanabe, K. Kudo, N. Ueno, Adv. Mater. 19 (2007) 665. [4] H. Vázquez, Y.J. Dappe, J. Ortega, F. Flores, J. Chem. Phys. 126 (2007) 144703, and references therein. [5] N. Koch, S. Duhm, J.P. Rabe, A. Vollmer, R.L. Johnson, Phys. Rev. Lett. 95 (2005) 237601. [6] Y. Zou, L. Kilian, A. Schöll, Th. Schmidt, R. Fink, E. Umbach, Surf. Sci. 600 (2006) 1240. [7] S. Duhm, A. Gerlach, I. Salzmann, B. Broker, R.L. Johnson, F. Schreiber, N. Koch, Org. Electron. 9 (2008) 111. [8] L. Kilian, A. Hauschild, R. Temirov, S. Soubatch, A. Schöll, A. Bendounan, F. Reinert, T.-L. Lee, F.S. Tautz, M. Sokolowski, E. Umbach, Phys. Rev. Lett. 100 (2008) 136103. [9] A. Ferretti, C. Baldacchini, A. Calzolari, R. Di Felice, A. Ruini, E. Molinari, M.G. Betti, Phys. Rev. Lett. 99 (2007) 046802. [10] H. Yamane, D. Yoshimura, E. Kawabe, R. Sumii, K. Kanai, Y. Ouchi, N. Ueno, K. Seki, Phys. Rev. B 76 (2007) 165436. [11] H. Yamane, E. Kawabe, D. Yoshimura, R. Sumii, K. Kanai, Y. Ouchi, N. Ueno, K. Seki, Phys. Stat. Sol. (b) 245 (2008) 793. [12] S. Kera, H. Setoyama, M. Onoue, K.K. Okudaira, Y. Harada, N. Ueno, Phys. Rev. B 63 (2001) 115204. [13] N. Koch, H. Oji, E. Ito, E. Zojer, H. Ishii, G. Leisinga, K. Seki, Appl. Surf. Sci. 175/176 (2001) 764. [14] J. Tersoff, Phys. Rev. Lett. 52 (1984) 465. [15] S. Lukas, G. Witte, Ch. Wöll, Phys. Rev. Lett. 88 (2001) 028301. [16] S. Söhnchen, S. Lukas, G. Witte, J. Chem. Phys. 121 (2004) 525. [17] C. Baldacchini, M.G. Betti, V. Corradini, C. Mariani, Surf. Sci. 566–568 (2004) 613. [18] J. Lagoute, K. Kanisawa, S. Fölsch, Phys. Rev. B 70 (2004) 245415. [19] C.B. France, P.G. Schroeder, J.C. Forsythe, B.A. Parkinson, Langmuir 19 (2003) 1274. [20] G. Beernink, T. Strunskus, G. Witte, Ch. Wöll, Appl. Phys. Lett. 85 (2004) 398. [21] H. Yamane, S. Nagamatsu, H. Fukagawa, S. Kera, R. Friedlein, K.K. Okudaira, N. Ueno, Phys. Rev. B 72 (2005) 153412. [22] K. Seki, H. Nakagawa, K. Fukui, E. Ishiguro, R. Kato, T. Mori, K. Sakai, M. Watanabe, Nucl. Instrum. Meth. Phys. Res. A 246 (1986) 264. [23] Y. Morikawa, H. Ishii, K. Seki, Phys. Rev. B 69 (2004) 041403 (R). [24] P.S. Bagus, K. Hermann, Ch. Wöll, J. Chem. Phys. 123 (2005) 184109. [25] M. Dion, H. Rydberg, E. Schröder, D.C. Langreth, B.I. Lundqvist, Phys. Rev. Lett. 92 (2004) 246401. [26] K. Lee, J. Yu, Y. Morikawa, Phys. Rev. B 75 (2007) 045402. [27] P. Sony, P. Puschnig, D. Nabok, C. Ambrosch-Draxl, Phys. Rev. Lett. 99 (2007) 176401. [28] A. Hauschild, K. Karki, B.C.C. Cowie, M. Rohlfing, F.S. Tautz, M. Sokolowski, Phys. Rev. Lett. 94 (2005) 036106; R. Rurali, N. Lorente, P. Ordejón, Phys. Rev. Lett. 95 (2005) 209601; A. Hauschild, K. Karki, B.C.C. Cowie, M. Rohlfing, F.S. Tautz, M. Sokolowski, Phys. Rev. Lett. 95 (2005) 209602. [29] N. Koch, A. Gerlach, S. Duhm, H. Glowatzki, G. Heimel, A. Vollmer, Y. Sakamoto, T. Suzuki, J. Zegenhagen, J.P. Rabe, F. Schreiber, J. Am. Chem. Soc. 100 (2008) 7300. [30] A. Gerlach, F. Schreiber, S. Sellner, H. Dosch, I.A. Vartanyants, B.C.C. Cowie, T.-L. Lee, J. Zegenhagen, Phys. Rev. B 71 (2005) 205425.

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