Stoichiometric and Oxygen-Deficient VO2 as Versatile Hole Injection

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Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Stoichiometric and Oxygen-Deficient VO2 as Versatile Hole Injection Electrode for Organic Semiconductors Keke Fu,†,∇ Rongbin Wang,†,‡,∇ Takayoshi Katase,§,∥,⊥ Hiromichi Ohta,§ Norbert Koch,†,‡,# and Steffen Duhm*,† †

Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices and Joint International Research Laboratory of Carbon-Based Functional Materials and Devices, Soochow University, Suzhou 215123, China ‡ Institut für Physik & IRIS Adlershof, Humboldt-Universität zu Berlin, Berlin 12489, Germany § Research Institute for Electronic Science, Hokkaido University, N20W10, Kita, Sapporo 001-0020, Japan ∥ Laboratory for Materials and Structures, Institute of Innovative Research, Tokyo Institute of Technology, 4259 Nagatsuta, Midori, Yokohama 226-8503, Japan ⊥ PRESTO, Japan Science and Technology Agency, 7 Gobancho, Chiyoda, Tokyo 102-0076, Japan # Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Berlin 12489, Germany S Supporting Information *

ABSTRACT: Using photoemission spectroscopy, we show that the surface electronic structure of VO2 is determined by the temperature-dependent metal−insulator phase transition and the density of oxygen vacancies, which depends on the temperature and ultrahigh vacuum (UHV) conditions. The atomically clean and stoichiometric VO2 surface is insulating at room temperature and features an ultrahigh work function of up to 6.7 eV. Heating in UHV just above the phase transition temperature induces the expected metallic phase, which goes in hand with the formation of oxygen defects (up to 6% in this study), but a high work function >6 eV is maintained. To demonstrate the suitability of VO2 as hole injection contact for organic semiconductors, we investigated the energy-level alignment with the prototypical organic hole transport material N,N′di(1-naphthyl)-N,N′-diphenyl-(1,1′-biphenyl)-4,4′-diamine (NPB). Evidence for strong Fermi-level pinning and the associated energy-level bending in NPB is found, rendering an Ohmic contact for holes. KEYWORDS: VO2, metal−insulator transition, hole injection electrode, organic semiconductor, Ohmic contact the MI transition temperature can be reduced to 50 K.16 Also, by introducing oxygen vacancies17−19 in VO2, the MI transition temperature can be widely modulated.20,21 Recently, it was shown that the highly oxygen-deficient VO2−δ (with δ = 0.2) is even metallic at 1.8 K,19 with the electronic structure of VO2−δ being slightly different from that of the stoichiometric VO2, see Figure 1a. For some metal oxides, it is known that oxygen vacancies can also influence the work function,22−24 which can reach up to 6.9 eV for materials like MoOx or WOx.25,26 Such high-workfunction materials are thus excellent hole-injecting electrodes in organic electronic devices, as Fermi-level pinning at the highest occupied molecular orbital (HOMO) level manifold of the organic semiconductor can enable Ohmic contacts.27,28 Also, for VO2, the work function is influenced by oxygen vacancies,29 but reported work function values are below 5.7 eV29,30 and

1. INTRODUCTION Metal oxides are suitable for a variety of (opto-)electronic applications as many of their properties, like optical transparency, conductivity, and energy gap, cover a wide range and can be relatively easily tuned via composition.1 Among metal oxides, VO2 is rather unique, as it exhibits a first-order metal− insulator (MI) transition from an insulating phase with a monoclinic structure to a metallic phase with a tetragonal rutile structure at just 40 K above room temperature (RT).2 Details of this transition are still subject of ongoing research.3−8 The conductivity is typically 4−5 orders of magnitude lower below the transition temperature than above it. The energy gap in the insulator phase is 0.6 eV,9,10 and the monoclinic VO2 can thus be regarded as a semiconductor from an application point of view. Moreover, the charge carrier density can be fine-tuned over orders of magnitudes by aliovalent ion doping11 and/or hydrogen incorporation.12,13 VO2 thus has a wide range of potential applications like optical switching, smart windows, and thermal memory.14,15 There are ways to modify the phasetransition temperature of VO2; e.g., by few-percent W doping, © XXXX American Chemical Society

Received: January 1, 2018 Accepted: March 2, 2018

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DOI: 10.1021/acsami.8b00026 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

Figure 1. (a) Band structure sketch of stoichiometric VO2 (insulating and metallic state) and VO2−δ drawn after ref 19. In the insulating states, the V 3d orbitals are split into three bands usually termed d||, d||*, and π*. The d|| and π* bands are partially filled in the metallic states and the Fermi level is shifted up in oxygen-deficient VO2−δ. (b, c) Valence region ultraviolet photoelectron spectroscopy (UPS) spectra of VO2 after “process A” and “process B”, measured at the indicated temperature. (d) Secondary electron cutoff (SECO) spectra measured at room temperature immediately after process A/B. resistivity measurements show the MI transition at 335 K. The corresponding X-ray diffraction patterns and resistivity traces are shown in the Supporting Information, as well as the surface morphology obtained by scanning force microscopy (Figure S1). The photoemission experiments were carried out in an ultrahigh vacuum (UHV) in a customized SPECS photoelectron spectroscopy system34 consisting of interconnected load-lock chamber, preparation and evaporation chambers (base pressure: 3 × 10−10 mbar), and analysis chamber (base pressure: 2 × 10−10 mbar). The XPS spectra were recorded using a monochromatized Al Kα source (1486.7 eV) and UPS spectra were recorded using the monochromatized He Iα excitation line (21.22 eV). Despite the insulating sapphire substrate, no signs of charging were detected during the XPS and UPS measurements because the charge carrier density of VO2 films is sufficiently high, also in the insulating state. For measurements at elevated temperatures, the margin between the actual sample temperature and that measured by a thermocouple placed next to the sample is estimated to be ±15 K. The UPS secondary electron cutoff (SECO) spectra were measured with a sample bias of −3 V to clear the analyzer work function. The energy difference between the SECO onset and the Fermi level (EF) is the sample work function.

thus little attention was paid to VO2 as a hole-injecting contact for organic semiconductors.31,32 In this work, we show that the work function of clean, stoichiometric VO2 can be as high as 6.7 eV, rendering this metal oxide a prime hole-injecting contact, as demonstrated here with the prototypical organic hole transport material, N,N′-di(1-naphthyl)-N,N′-diphenyl-(1,1′-biphenyl)4,4′-diamine (NPB). Moreover, our temperature-dependent X-ray photoelectron spectroscopy (XPS) and ultraviolet photoelectron spectroscopy (UPS) measurements disentangle the impact of oxygen vacancies on the one hand and the MI phase transition on the other hand on the surface electronic structure of VO2, and we disclose how both contributions depend on the details and sequence of surface treatment.

2. EXPERIMENTAL DETAILS An epitaxial VO2 film of 115 nm thickness was grown on a A-plane sapphire (α-Al2O3) substrate by pulsed laser deposition.33 The epitaxial relationship between the VO2 film and the α-Al2O3 substrate is (100)[010]VO2||(112̅0)[0001] α-Al2O3. Temperature-dependent B


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Figure 2. Impact of heating in UHV on a VO2 sample (initially prepared by process B): (a−c) the O 1s and V 2p core levels measured at room temperature (295 K), 403 K, and again at 295 K after cooling down in UHV. (d) A comparison of the O 1s core level spectra. (e) The UPS secondary electron cutoff and (f) the valence band region of the same sample.

O 1s and V 2p core levels evidence the presence of V5+, in addition to V4+ expected for stoichiometric VO2 (Figure S2b,c). As indicated in Figure 1a, also oxygen vacancies can lead to a metallic sample. Moreover, the V 3d-derived peak5−8,10 (d||) centered around 1 eV binding energy (BE) is not very well resolved after process A, which also supports a nonuniform and nonstoichiometric surface. In contrast, after process B, i.e., sputtering with tenfold higher ion current and the final annealing step done in oxygen (6 × 10−4 mbar), barely any carbon contamination can be detected (Figure S2a). The V 2p3/2 core level (Figure 2a) is centered at 516.2 eV BE, in accordance with the reported values for VO2,37−39 and the spectrum shows only V 4+ contributions; VO 2 is thus stoichiometric after process B. This is further supported by valence spectra, as the V 3d-derived peak at ∼1 eV BE is well developed over the entire temperature range with an MI transition (Figure 1c). Figure 1d shows the SECO spectra (measured at 295 K) of VO2 films after the two treatments. Notably, although the work function is only 5.1 eV after process A, it increases to 6.7 eV after process B, which is substantially higher than hitherto reported VO2 work function values, which range from 4.80 to 5.65 eV.29,40,41 The likely cause for lower work function values obtained earlier is the residual surface contamination, most commonly carbon-based, the proper

The VO2 surfaces were subjected to Ar-ion sputtering (15 min at a bias of 600 V) to remove surface contamination and subsequent annealing in UHV (20 min at 425 K) cycles to heal the structural surface defects. For process A, the discharge current was 0.5 mA, whereas it was 5.4 mA for process B, to investigate whether higher sputter rates (and thus shorter overall sputter times to completely remove contamination) have an influence on the final surface properties. Because it is known that annealing of metal oxide in UHV can result in an oxygen-deficient surface, for process B, the final annealing step was done in oxygen atmosphere with a pressure of 6 × 10−4 mbar, which is known to help in recovering the intended stoichiometry.35,36 NPB was sublimated in situ onto VO2 from resistively heated Knudsen cells with a deposition rate of about 2 Å/min (as controlled by a quartz-crystal microbalance). All evaporation steps and all NPB/VO2 measurements were done at room temperature (295 K). The NPB thin films could be removed completely by process B and the VO2 sample could be used repeatedly.

3. RESULTS AND DISCUSSION The temperature-dependent valence spectra of VO2 after process A, i.e., Ar-ion sputtering at low ion current and annealing only in UHV (details in Section 2), reveal the MI transition by the increased density of states (DOS) close to EF and the evolution of a characteristic Fermi-edge with increasing temperature (Figure 1b). However, after process A, we still find carbon contamination on the VO2 surface (Figure S2a), and the C


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Figure 3. (a) UPS spectra of SECO and HOMO region during the stepwise evaporation of NPB on VO2 substrate (prepared by process B). (b) Energy-level diagram of NPB/VO2/Al2O3 based on these UPS spectra and XPS spectra of the same interface (Figure S4).

S3) derived from O 2p bands were also apparent in these earlier works, although different photon energies cause intensity variations. In contrast, for V2O5 at room temperature, the valence band maximum (VBM) is expected to be at 2.5 eV BE and there should be no features in 0−2 eV BE range in the UPS spectrum;42,50 V2O3 at room temperature is metallic and a clear Fermi edge should be visible.51,52 However, for our samples at room temperature, the VBM does clearly not cross the Fermi level and shows a defined peak between 0 and 2 eV BE, thus ruling out the presence of V2O5 or V2O3. To further investigate the role of oxygen vacancies, stoichiometric VO2 prepared by process B was subjected to heating without oxygen atmosphere, i.e., in UHV. The V 2p core levels measured at T = 403 K (Figure 2b) show clear contributions of V3+, and the area ratio of V4+ and V3+ yields that the sample becomes VO2−δ with δ = 0.06. The V 2p levels remain virtually constant after cooling the sample to room temperature (Figure 2c). Likewise, the O 1s signal (Figure 2d) changes after initial heating in UHV but does not further change upon cooling to RT. This implies that heating stoichiometric VO2 in UHV just mildly above the MI transition temperature leads to the formation of oxygen vacancies, as evidenced by the reduction of some V ions from V4+ to V3+. After cooling down to room temperature, the oxygen loss cannot be recovered due to the lack of an oxygen reservoir in

removal of which was not commented on in these reports. It is known that adsorbed molecules lower the work function of many surfaces, particularly metals and metal oxides, due to the “push-back” effect, i.e., Pauli-repulsion of electron density between the two materials in contact.42−45 In addition, Kelvinprobe based work function measurements29,41 have an uncertainty due to the fact that only the contact potential difference between tip and sample is measured,46 and proper calibration in air is demanding. For comparison, Greiner et al.22 have reported the work function for V metal (0), V2O3 (+3), and V2O5 (+5) to be 4.0, 5.0, and 7.0 eV, respectively. There appears to be a correlation between the metal oxidation state and the work function, and our work function values of 6.7 eV for VO2 (+4) is in line with this. One should, however, not expect a simple or linear relationship. As seen from Figure 2a,b,e, when heating VO2 from 295 to 403 K, the work function decreases with the presence of small amount of V3+ (or the presence of O vacancies), consistent with the above tendency. To further substantiate the presence of stoichiometric VO2 after process B (in addition to the fact that only V4+ was identified in the core levels), we compare the valence band of VO2 in the insulating state with the reported VO2 single crystal10,39,47 and thin film data.48,49 A well-defined V 3dderived peak at ∼1 eV BE without an apparent Fermi edge (Figure 1c) is fully consistent with refs 10, 39, and 47−49. Furthermore, the two broader peaks (labeled I and II in Figure D


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constant, i.e., no adsorption-induced band bending occurs in VO2. In contrast, the NPB-related N and C core levels exhibit a gradual shift of 0.4 eV to higher BE with increase in the NPB coverage from a nominal 4−160 Å. The magnitude of this shift is the same as observed for the HOMO and HOMO−1 features in the valence region. Apparently, all molecular electronic levels shift in parallel, indicative of an electrostatic origin. In line with the present understanding of the interfacial-level alignment of organic semiconductors on electrodes, our observations can be explained by a strong Fermi-level pinning at the HOMO-level manifold of NPB. As the work function of VO2 (6.7 eV) is substantially higher than the ionization energy of NPB (ca. 5.4 eV55), the electron transfer from NPB to VO2 must occur to establish an electronic equilibrium at the interface. The positive charge density thus induced within the NPB layer leads to a diffusion-driven density profile away from the interface, resulting in energy-level bending,56,57 i.e., the discretized analogue of band bending in covalent semiconductors. Because this energy-level bending amounts to 0.4 eV in the present case (inferred from UPS and XPS above), 1.8 eV of the total work function change (2.2 eV) upon interface formation remain to be explained. Because no band bending within VO2 occurs, the 1.8 eV value is due to an interface dipole, i.e., localized right at the interface between the two materials. This interface dipole has most probably two contributions, one due to the charge transfer from VO2 into the first layer of the organic semiconductor58,59 and a second due to the push-back effect55,60 due to molecule-induced Pauli repulsion at the VO2 surface. The presence of positively charged NPB molecules at the interface, inferred from the observation energy-level bending within the NPB layer, can be evidenced from an analysis of the C 1s and N 1s core levels. These levels are indicative of the charge state of the molecules, i.e., those of the cations are shifted toward a higher binding energy compared to those of neutral NPB. The thick-film C 1s spectrum (160 Å NPB coverage and topmost curve in Figure 4) is from the neutral NPB only (energy-level bending is saturated by the thickness and the density of charged NPB can reasonably be assumed to be zero or negligible) and exhibits a double peak, the larger peak from carbon atoms bound to other carbons (CC), and the smaller one from carbons bound to nitrogen atoms (CN). Assuming that only the BE of both species changes rigidly upon NPB cation formation, the spectra of thinner NPB films require a second double peak for adequate fitting, as shown for 2, 4, 8, and 16 Å NPB coverage in Figure 4. From the relative area of the two double peaks, we can directly infer the fraction of positively charged NPB molecules in the monolayer coverage range (diameter of one molecule ca. 10 Å), which amounts to little over 30%. An analogous fit made for the N 1s spectra results in a slightly smaller fraction of 26% (see Figure S6), but due to the lower overall signal from N 1s, the evaluation is less robust than that for the C 1s spectra. We refrain from also further evaluating the intermediate film thicknesses, as due to the superposition of a few molecular layers in the signal, additional broadening due to energy-level bending would yield less reliable evaluations. Both the substantial energy-level bending within NPB and the estimation of positive NPB fraction close to the VO2 interface evidence a hole concentration at the interface that is markedly beyond the intrinsic free carrier concentration of the wide-gap organic semiconductor at RT, i.e., the interface is Ohmic and thus well suited for hole injection and extraction, as required in numerous electronic and optoelectronic devices. In practical

UHV, and also no diffusion of oxygen from the bulk to the surface occurs. Figure 2f shows the impact of oxygen vacancies and the metal−insulator transition on the DOS close to EF. For stoichiometric VO2 analyzed below the transition temperature (black curve), the V 3d-derived DOS is pinned at EF, but no Fermi edge is visible. As heating the sample in UHV above the transition temperature induces oxygen vacancies and the UPS measurements under O2 atmosphere are not possible in our setup, no spectra of stoichiometric VO2 above the transition temperature could be acquired. For VO2−δ measured above the transition temperature (at 403 K, red curve), the Fermi edge is well pronounced. For the nonstoichiometric VO2−δ measured below the transition temperature (after heating in UHV) of VO2 (blue curve), the oxygen vacancies lead to the persistence of the Fermi edge. However, there is less spectral weight close to EF compared to the spectrum of VO2−δ measured at 403 K. Overall, these UPS results agree with the model suggested in ref 19 and schematically shown in Figure 1a. However, the pronounced differences in the DOS near EF of VO2−δ measured at 403 and 295 K shows that the MI transition still contributes substantially to the DOS also for nonstoichiometric VO2−δ. Moreover, in the RT spectra, the V 3d-derived peak is centered at ∼1 eV BE for both VO2 and VO2−δ. Consequently, it can be speculated that for our samples, in contrast to the proposed model,19 the additional DOS at EF is not due to an upward shift of the occupied d-band but due to oxygen-vacancy-derived gap states. This proposition is supported by the shift of the maximum of the d-derived photoemission intensity toward EF upon measuring VO2−δ at elevated temperature (403 K, red spectrum in Figure 2f), but detailed theoretical modeling should help clarify this issue in subsequent work. Within this experimental sequence just discussed above, the sample work function (Figure 2e) changes hardly from stoichiometric insulating (6.50 eV) to nonstoichiometric VO2−δ measured at 403 K (6.28 eV), but it decreases notably to 5.75 eV for VO2−δ measured at RT. This substantial work function decrease upon cooling in UHV to room temperature is ascribed to surface contaminations, as the C 1s signal becomes rather notable (Figure S3). Yet, after process B, we find consistently that the work function of stoichiometric VO2 is between 6.5 and 6.7 eV (Figures 1d, 2e, and 3a). To substantiate that the ultrahigh work function of stoichiometric VO2 makes it indeed a suitable hole injection/ extraction contact in electronic devices, the organic hole transport material NPB was deposited stepwise on VO2 conditioned by process B, but kept at room temperature thereafter throughout. After the deposition of submonolayer NPB (nominal coverage: 2 Å), the SECO shifts to a considerably lower kinetic energy (Figure 3a) and continues to shift with increasing NPB coverage. In total, the work function decreases from the bare VO2 value of 6.7 eV to 4.5 eV for a multilayer NPB coverage (nominal coverage: 160 Å). With increasing coverage, the NPB valence features, most notably the HOMO and HOMO−1-related peaks centered at ∼1.0 and ∼1.5 eV BE, become apparent and shift to a higher BE. The broadening of the NPB HOMO and HOMO−1 features beyond 96 Å coverage is indicative of an increased energetic disorder, most likely due to increasing conformational variation of the molecules in multilayers, e.g., as the intramolecular twist angle can sample more values.53,54 Upon NPB deposition, the V and O core levels (Figure S4) become attenuated by the molecular overlayer, but their BE stays E


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Ex situ characterization of VO2 substrates, XPS and UPS in situ characterization of VO2 substrates; details of UPS and XPS spectra for NPB stepwise deposited on VO2; fitting of N 1s levels (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Hiromichi Ohta: 0000-0001-7013-0343 Norbert Koch: 0000-0002-6042-6447 Steffen Duhm: 0000-0002-5099-5929 Author Contributions ∇

K.F. and R.W. contributed equally to the manuscript. The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Financial support from the Major State Basic Research Development Program of China (973 Program, No. 2014CB932600), an NSFC Research Fund for International Young Scientists (No. 11550110176), the National Key R&D Program of China (No. 2017YFA0205002), the 111 Project of the Chinese State Administration of Foreign Experts Affairs, the Collaborative Innovation Center of Suzhou Nano Science & Technology (NANO-CIC), and the SFB951 (DFG) is gratefully acknowledged. T.K. was supported by PRESTO, JST (JPMJPR16R1), Grant-in-Aid for Young Scientists A (15H05543) and Grant-in-Aid for Challenging Exploratory Research (16K14377) from JSPS. H.O. was supported by Grant-in-Aid for Scientific Research A (17H01314) and Grantin-Aid for Scientific Research on Innovative Areas (25106007) from JSPS.

Figure 4. XPS C 1s core level spectra and their fits of NPB up to ca. monolayer coverage (2−16 Å) and thick film (160 Å), as indicated next to each spectrum. Fitting procedure with two double peaks due to C−C and C−N intramolecular bonds for neutral and positively charged NPB as explained in the text. The CC peaks (higher blue peak) were aligned to zero to remove binding energy shifts due to energy-level bending as a function of film thickness. CC and CN with superscript “+” are from positively charged NPB molecules.

applications, VO2 would probably be employed as a thin film on a highly conductive substrate, such as (polycrystalline) metals or transparent conductive oxides, likely resulting in structurally less ordered vanadium oxide films as studied here. However, it can be expected that annealing under oxygen (maybe even higher pressure than applied here) will as well restore the VO2 composition after sputter cleaning, and thus the high work function. For instance, the MoO3 thin films grown on indium tin oxide61 or Au62,63 exhibited a work function (or slightly higher) comparable to that of a MoO3 single crystal.64

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4. CONCLUSIONS In conclusion, the work function of VO2 has been underestimated and we show that the work function of stoichiometric VO2 can reach up to 6.7 eV by a simple annealing step at only 425 K under moderate O pressure. Moreover, even oxygen deficiencies of up to 6% are not detrimental to achieve high work function values. The Fermi-level pinning of the organic semiconductor NPB shows that VO2 is probably suitable to achieve Ohmic contacts with organic semiconductors that have ionization energies as high as 6 eV, provided that the push-back is not exceedingly strong. Furthermore, even in the insulating state at room temperature, the charge carrier density is relatively high and allows for charging-free UPS measurement, making VO2 a superior hole-injecting electrode in organic electronics devices.

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The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b00026. F


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