Article pubs.acs.org/JPCC
Interface between Graphene and SrTiO3(001) Investigated by Scanning Tunneling Microscopy and Photoemission Horacio Coy-Diaz, Rafik Addou, and Matthias Batzill* Department of Physics, University of South Florida, Tampa, Florida 33620, United States S Supporting Information *
ABSTRACT: Graphene, grown by chemical vapor deposition, is transferred onto Nb-doped SrTiO3(001) surface and the interface properties are characterized by scanning tunneling microscopy and photoemission spectroscopy. Charge doping of graphene changes from n- to p-type with vacuum annealing and correspondingly opposite space charge regions are observed in SrTiO3 substrate. Formation of an ordered surface reconstruction of the SrTiO3 substrate underneath the graphene is observed. The surface restructuring can be measured in scanning tunneling microscopy because the graphene closely follows to the substrate topography. This causes at the atomic-level a wavy graphene morphology on the SrTiO3 (001)-c(6 × 2) surface reconstruction. Prolonged high temperature (above 700 °C) vacuum annealing causes formation of hexagonal holes with ‘armchair’ edges in the graphene and an eventual disappearance of the graphene. Etching of the graphene is assumed to be caused by reaction with released substrate oxygen.
1. INTRODUCTION Transfer of CVD-grown graphene to arbitrary substrates has become ubiquitous in graphene research. Initially, graphene was transferred to Si/SiO2 as a convenient substrate, where a doped Si substrate could be used as a gate electrode and SiO2 as the gate dielectric.1−5 However, it was recognized that graphene properties of SiO2 supported graphene is strongly affected by the support6−16 and ultraflat substrates like hex-BN17−20 or potentially MoS221 exhibit far fewer carrier scattering sites and thus a higher charge mobility of graphene. Therefore, hex-BN is a very promising material for graphene support although large scale synthesis of good quality hex-BN and growth of hex-BN on a gate electrode material for bottom electrode architectures are challenges that still need to be addressed. Recently, it was reported that by using SrTiO3(001) (STO) wafers as substrates for graphene a better charge carrier mobility can be obtained than for SiO2.22 This good property has been related to the very high dielectric constant of STO, which results in efficient screening of impurity charges.23,24 Furthermore, conducting Nb-doped STO may be used as a gate electrode and thus NbSTO/STO may be used as a gated support for graphene.25,26 In order to understand the materials properties of the graphene/STO heterostructure a detailed description of the interface is needed. Here we investigate the interface between STO and transferred graphene with the use of ultrahigh vacuum (UHV) scanning tunneling microscopy (STM) and photoemission spectroscopy. We find that vacuum annealing is necessary to prepare a well-defined interface and that at very high temperature the graphene may be etched by oxygen from the substrate. Furthermore, surface reconstructions of STO © XXXX American Chemical Society
may persist underneath the graphene and modulate the topography of graphene.
2. EXPERIMENTAL METHODS In order to facilitate characterization of structural and interface band alignment, Nb-doped, n-type STO was used as the substrate. Prior to graphene transfer the STO samples were etched in buffered hydrofluoric acid and annealed in air at 950 °C for 8 h. Graphene was prepared by CVD growth in a tube furnace on high purity copper foil. The graphene was transferred to the STO substrate by spin-coating a layer of poly(methyl methacrylate) (PMMA), followed by chemically etching of the copper in ammonium persulfate. The graphene/ PMMA was then rinsed in DI water and captured with the STO substrate. Subsequently, the PMMA was dissolved in acetone. To remove any organic residue from the surface the graphene/ STO samples were annealed in air at 350 °C for 5 h.27 After graphene transfer, samples were characterized by scanning electron microscopy (SEM) and ambient atomic force microscopy (AFM) in noncontact mode. In addition, the samples were loaded into two ultrahigh vacuum (UHV) surface science characterization chambers with base pressures in the low 10−10 Torr regime. Because of the nature of the graphene sample, substrate preparation was limited to vacuum annealing. The same procedures were used for graphene/STO as for bare STO substrates. One UHV chamber was equipped with an Omicron VT-scanning tunneling microscope, operated at room Received: August 1, 2013
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Figure 1. Top left: SEM image of a partially graphene-covered Nb-STO substrate. Top right: AFM image of bare substrate. Bottom left: AFM image of graphene-covered substrate. Bottom right: line profile of graphene-covered (green) and bare substrate (red).
copper foil28 has demonstrated, the majority of the graphene consists of monolayer and only some bilayer graphene is present in small regions. Our Auger spectroscopy data of the graphene on copper foil, used in this experiment, are also consistent with mainly monolayer graphene (see Figure S2). The contrast in SEM is consistent with previous assignment of the darker regions to bilayer graphene28 and the remaining surface to monolayer. This assignment is further corroborated by comparison of SEM data on the as-grown graphene on copper with that after transfer to STO (see Figure S3). Also, AFM characterization of graphene transferred from the copper foil to a SiO2/Si substrate shows that the smaller regions are due to bilayer (see Figure S4). Apart from a few wrinkles in the graphene, the graphene follows the terrace structure of the substrate closely, as generally reported.29 It is well-known that preparation of STO by buffered-HF etching can leave fluorine impurities behind at the surface. This may be a concern for transferred graphene as fluorine may act as a charged impurity that can scatter charge carriers in graphene. Our XPS studies show significant fluorine contamination for the as prepared samples, i.e., even after 8 h annealing in air at 950 °C after buffered HF etching. However, our XPS studies, shown in Figure 2, indicate that fluorine can be removed by annealing in vacuum at 300 °C. This indicates that there exist pathways for the fluorine to desorb from the surface or diffuse into the bulk. 3.1. Interface Band Alignment. Photoemission spectroscopy measurements of core-level, valence band, and work function (by means of secondary (inelastic scattered) electron cutoff determination) are a well-established method for determining interface band alignments. Generally, this is done by monitoring photoemission peaks as the interface is being formed by vacuum deposition. Unfortunately, graphene deposition onto STO in vacuum is currently not possible. Therefore, we have to compare photoemission peak position on two separate samples, i.e., a bare and graphene covered
temperature, and the second UHV system was a MuMETAL chamber for photoemission spectroscopy. The system was equipped for both X-ray photoemission spectroscopy (XPS) and ultraviolet photoemission spectroscopy (UPS) measurements. A Mg/Al dual anode X-ray source and He-VUV lamp were used for XPS and UPS, respectively. An Omicron-Sphera II hemispherical energy analyzer has been used for measuring the photo electron spectra. Evaluation of core-level photoemission was done by fitting of Gaussian/Lorentzian line shape to the peak after subtracting a Shirley background for emission from the STO substrate. The C 1s peak was fitted with a Doniach−Sunjic peak shape to accommodate for the typical asymmetric line shape of sp2 carbon. The work function and valence band were studied by UPS using a He II VUV photon source (Omicron, HIS 13). In order to shift the secondary electron background above the work function of the analyzer, the sample was biased at −9.2 V for work function measurements. The work function was determined by the difference between the UV photon energy hω = 21.2 eV and the measured spectral width, which is determined by the onset of the secondary electron emission relative to the Fermi edge, which was independently determined on a metal sample. In order to compare photoemission data from clean and graphene covered STO substrates and thus to construct an interface band diagram, both samples were treated with identical vacuum annealing procedures.
3. RESULTS AND DISCUSSION Figure 1 shows the SEM and ambient-AFM images of a (partially) graphene-covered STO sample. The samples used were almost completely covered by graphene, and only in the corners used for sample handling uncovered areas remain; details are illustrated in Supporting Information Figure S1. In the SEM image of the graphene covered regions two contrasts can be observed. As previous work on CVD grown graphene on B
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photoemission peak positions for a graphene/covered STO sample and a bare STO sample. Table 1 summarizes the Table 1. Summary of All the Measured Binding Energies for a Bare and Graphene-Covered Nb-STO Substrate after Annealing at 800 °C for 30 min O 1s Ti 2p3/2 Sr 3d5/2 C 1s work function VBM
Figure 2. XPS survey spectra for graphene/Nb-STO before and after annealing at 300 °C. Specifically note that the strong fluorine contamination disappears after vacuum annealing.
Nb-STO (eV)
G/Nb-STO (eV)
530.16 458.90 133.38
530.29 459.04 133.58 284.37 4.40 3.18
5.05 2.93
photoemission measurements, and Figure 4 shows the pertinent photoemission data and the deduced interface band alignment. Shifts in the STO core-level positions as well as a shift of the valence band maximum (VBM) of the STO (note that graphene exhibits low density of states in the region of the STO valence band and thus does not contribute significantly to the measured valence band) indicate a downward band bending of the STO of ∼0.2 eV for the graphene covered surface compared to the bare STO surface. The C-1s core level is shifted by 0.04 eV compared to the peak position in HOPG. The C-1s peak shape for graphene on STO maintains the characteristic line-shape of sp2 carbon as a comparison with HOPG or graphene on copper (see Figure S5) illustrates. Consequently, there is no covalent interaction between graphene and STO. Only weak van der Waals interaction is also the expected adsorption for the inert graphene with a widebandgap oxide. In the absence of any covalent bonding between graphene and the substrate, chemical shifts are negligible, and therefore a shift in XPS can be assigned to a shift of the Fermi level in graphene30−32 due to an interface charge transfer. The
sample. Although the non-vacuum transfer of graphene compromises the interface quality, it is the relevant condition used for example in transport measurements of STO supported graphene. Fortunately, the STO surface is fairly inert, and simple annealing in vacuum results in well-defined surface structure as we will show by STM, below. The presence of interface impurities (e.g., water) for low vacuum annealing temperatures is apparent from the evolution of the core-level binding energies of the substrate and the graphene. Figure 3 shows the peak position of the Sr 3d5/2 and C 1s peak as a function of annealing temperature. Both peaks shift to lower binding energy with increasing annealing temperature. These shifts are more easily discussed after we determine the band alignment of the clean graphene/STO interface, which we discuss next. For samples vacuum-annealed to 700 °C the core-level position stabilizes, which we interpret as the formation of a “clean” interface. Under these conditions we can compare
Figure 3. Peak position of the (a) C 1s and (c) Sr 3d5/2 as a function of annealing temperature. XPS core level spectra of (b) C 1s and (d) Sr 3d5/2 for two different annealing temperatures. Annealing at each temperature was done for 30 min. C
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Figure 4. Photoemission studies of band alignment of the graphene/Nb-STO interface. (a) XPS core level spectra of C 1s peak. The peak is shifted by ∼0.04 eV to lower binding energy, indicating a down ward shift of the Fermi level below the Dirac point. The relationship between Fermi level position, C 1s core level, and work function is schematically illustrated in the inset of (a). (b) Core level spectra of Sr 3d5/2; the band bending in NbSTO is evaluated using the Sr 3d5/2 core level shift for Nb-STO and G/Nb-STO. (c) He II UPS spectra for G/Nb-STO and Nb-STO. The work function was calculated by the difference between the UV photon energy hω = 21.2 eV and the measured spectral width, which is determined by the onset of the secondary emission peak relative to the Fermi edge. Additionally, the shift in the valence band maximum is shown. (d) Interface electronic structure of graphene on Nb-STO as derived from photoemission studies.
Figure 5. C 1s/Ti 2p3/2, C 1s/Sr 3d5/2, and C 1s/O 1s XPS signal ratios. (a) XPS signal ratios as a function of the annealing temperature. Annealing at each temperature was done for 30 min. (b) Time sequence for C 1s/Ti 2p3/2, C 1s/Sr 3d5/2, and C 1s/O 1s ratios at an annealing temperature of 825 °C. (b) XPS core level spectra of C 1s for different annealing times at 825 °C.
imply any chemical bonding between the graphene and STO. The dipole is related to the well-known “pillow” effect of adsorbed molecules on surfaces.34,35 The “pillow” effect can be basically described by Pauli repulsion between electrons form the substrate and from graphene that causes charge redistribution at the interface compared to the free surfaces. Following this analysis of the band alignment of the “clean” interface, we can now better understand the changes in the core level binding energies shown in Figure 3 as a function of vacuum annealing. The shifts in the STO core levels and the graphene C 1s peak are consistent with a transition in space charge region of the substrate from an upward band bending to a downward band bending and a transition of the charge doping in graphene from an n-type doped graphene for a “dirty” interface to a p-type doped graphene for a “clean” interface. Realizing the dependence of the charge transfer doping in graphene on the annealing temperature and interface
relationship between C 1s binding energy shift and charge doping (Fermi level shift) is schematically illustrated in the inset of Figure 4a. Using the C 1s peak of HOPG as a reference line for charge neutral graphene, a shift of 0.04 eV to lower binding energies indicates a p-type doping of graphene, consistent with the observed formation of a negative space charge region (downward band bending) in STO. Finally, formation of an interface dipole is determined from the change in the work function. We measure a work function of clean STO of 5.05 eV and after graphene transfer of 4.4 eV. Our measured work function for graphene on STO is in good agreement with previously reported work functions measured by Kelvin probe microscopy.33 Taking the band bending in the STO substrate into account, the difference in work function of the surface with and without graphene allows us to determine the interface dipole between graphene and STO of 0.45 eV. This is schematically indicated in Figure 4d. This dipole is a consequence of interface electron redistribution but does not D
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cleanliness is important for interpreting transport measurements in STO supported graphene. 3.2. Thermal Stability of Graphene/STO in Vacuum. In this section we address the thermal stability of graphene on STO in vacuum. Monitoring the C 1s signal with annealing temperature and normalizing it to a substrate signal shows that the graphene is stable up to annealing temperatures of 700 °C, and then the carbon signal starts to decrease. This is shown in Figure 5a. A time sequence for prolonged annealing at 825 °C is shown in Figure 5b, and Figure 5c shows XPS spectra for a few selected time points of the C 1s and Sr 3p region. With the reasonable assumption that the substrate signal does not change the C 1s/Sr 3d (or O 1s or Ti 2p) ratio is proportional to the area covered by graphene. Thus, the decrease in the ratio indicates a consumption of graphene carbon, and the change in the peak ratio is the rate at which graphene is consumed. From STM (see below) we know that hexagonal holes are formed in the graphene at high annealing temperatures, and thus it is reasonable to assume that graphene is etched from the edges of these holes. The rate at which graphene is etched at a constant temperature may give some information on the mechanism of this process. For instance, if carbon was detached from graphene edges at a constant rate, as one might expect if there was a mechanism for carbon to simply escape into the gas phase or the bulk of STO, then the rate at which the C 1s signal decreases should increase with the length of the perimeter of the holes in the graphene sheet. Since the length increases with time (at least initially until holes start to coalesce), an increase in the carbon etching rate may be anticipated. However, contrary to this expectation, the rate at which carbon is consumed with time decreases and in fact can be fitted fairly well by a simple exponential decay with a decay constant λ = 240 ± 5 min. Such a decay implies that the consumption of graphene is proportional to the area covered by graphene. This is possible if an etchant (for instance oxygen) is captured by the graphene covered area and transported to the graphene edge where it may react with edge atoms and subsequently desorb as CO or CO2. STO reduces by annealing in vacuum.36 This is also evident from the color change during annealing of undoped STO in vacuum from transparent to a blue color. The blue color is due to oxygen vacancy-related color centers in the material.. Thus, we speculate that at high annealing temperatures some oxygen from the bulk diffuses to the surface where it is initially trapped underneath the graphene. Those trapped oxygen species may emerge at holes in the graphene where it reacts with the undercoordinated graphene edge atoms. 3.3. Scanning Tunneling Microscopy of Graphene/ STO. STM studies of the evolution of monolayer-graphene/ Nb-STO(001) surface with increasing annealing temperature are shown in Figure 6. Samples, vacuum annealed to ∼300 °C for 1 h, exhibit two-dimensional clusters on the STO terraces. Two structures may be differentiated: an open structure with random cluster shapes and sizes and terraces on which the clusters are denser and organized in cross-hatched structure with preferential cluster orientation in the [110] and [−110] crystallographic direction. In the open structure a uniform cluster height of ∼4 Å is measured. These 4 Å tall clusters disappear after annealing to 600 °C for 1 h, and the “crosshatched” clusters become better defined. Now the surface consists of flat terraces and regions that exhibit elongated clusters in the [110] and [−110] direction. In Figure 6b, we can see that these clusters are ordered and a substructure consisting
Figure 6. STM of monolayer-graphene covered STO at different vacuum annealing temperatures: (a) 300 °C (Vbias = 1.5 V, It = 0.5 nA) and (b) 600 °C (Vbias = 1.3 V, It = 0.5 nA). With increasing temperatures the substrate is getting better ordered, indicating material rearrangement underneath of the graphene.
of rows along the long axis of the clusters with a separation of 2.2 ± 0.3 nm, i.e., 4 × √2aSTO = 2.2 nm. Since we cannot resolve the structure along the rows, this kind of reconstruction may be labeled as (n√2 × 4√2)R45°, where n√2 corresponds to the periodicity along the rows, which could not be determined from the STM image (note that the entire surface is graphene covered, which makes a determination of the atomic structure of the underlying STO substrate more challenging). A large number of possible reconstructions for the STO(001) surface are reported,3637,38 many after ion sputtering and annealing procedures.39−44 Although there is no specific report for a 4√2-R45° structure, some of the various “centered” unit cell surface reconstructions may account for this island morphology. Even after prolonged annealing at 600−700 °C these two substrate terminations remain, as is apparent in large scale STM images as shown in Figure 7a. One structure exhibits the strongly corrugated partially reconstructed units of the 4√2-R45° kind. The other structure appears as flat terraces. Zoom-in on the flat areas allows imaging the atomic structure of monolayer graphene, as shown in Figure 7c. Careful investigation of the graphene structure reveals a periodic wave structure of the graphene. Comparison of this wavy graphene with a bare STO surface suggests that the waviness in the graphene is a consequence of a STO c(6 × 2) reconstruction underneath the graphene. The c(6 × 2) reconstruction is known to exist for air and vacuum-annealed STO samples.45−47 Figure 7d shows an overlay of a highresolution image of the graphene layer on top of a bare STO(001)-c(6 × 2) surface structure. It should be mentioned that the imaging condition for the bare STO (Vbias = 1.4 V, It = 0.4 nA) is vastly different from that of the graphene (Vbias = 10 mV, It = 5 nA). Nevertheless, the corrugation of the c(6 × 2) reconstruction of the substrate is clearly reproduced in the graphene. The c(6 × 2) reconstruction also exhibits some defects, specifically some brighter protrusions are imaged along the row structure. These have been previously reported and have been attributed to additional TiO2 units.47 While the convolution between topography and electronic effects can make atomicE
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Figure 7. Monolayer-graphene on STO after annealing to 700 °C. In the large scale image (450 nm × 450 nm, Vbias = 1.5 V, It = 0.5 nA) shown in (a) two surface structures of the graphene covered STO are observed. In addition, hexagonal holes are imaged that correspond to holes in the graphene layer. The apparent step height between graphene-covered surface and the bare STO substrate is measured in (b) to 3−4 Å, which corresponds well to usual van der Waals between graphene and weakly interacting substrates. From comparison with atomic resolution images of the graphene (e.g., shown in (c)) indicate that the edge orientation of the holes is of the “armchair” kind. A zoom-into the black square indicated in (a) on the seemingly flat terrace is shown in (c) (Vbias= 10 mV, It = 5 nA). In the high-resolution image the honeycomb structure of the graphene is clearly observed. In addition, a periodic wavy modulation of the graphene is observed with ridges along the [100] direction of the STO substrate. This modulation is a consequence of a c(6 × 2) STO reconstruction. This is clearly illustrated in (d), where two STM images are partially superimposed. The STM image of the graphene covered surface was taken in the same area as (c) and the image of the bare STO-c(6 × 2) substrate was taken on the same sample after prolonged vacuum annealing that removed most of the graphene. The faint waviness of the graphene is of the same periodicity as that of the STO-c(6 × 2) reconstruction illustrating that graphene follows the substrate modulation of the surface reconstruction. A small misalignment between the modulation of the two images is due to unavoidable differences in thermal drift of the STM. The inset in the lower right corner shows the c(6 × 2) unit cell (white rectangle) and the primitive unit cell (black, dashed line).
scale topographical assignments in STM of oxide surfaces challenging, the fact that similar bright protrusions are imaged after covering the surface with graphene suggests that these protrusions imaged in STM are indeed of topographical origin. Regardless of the extent the graphene follows the substrate corrugation, it appears that the substrate induces some local topographical bending of the graphene at the nanometer scale, and this may induce modification of the electronic properties in graphene. Further studies of the effect of periodic substrate reconstruction on the properties of supported graphene are planned for the future. For samples annealed at 700 °C we also observe the formation of hexagonal holes in the graphene sheet as highlighted in Figure 7a. These holes are tens of nanometers wide. Cross sections, as the one shown in Figure 7b, across the edge between graphene-covered and the bare STO substrate indicate an apparent height of 3−4.5 Å. While in STM absolute height measurements between two dissimilar materials may be influenced by differences in tunneling probabilities, these values are consistent with monolayer graphene. From atomic
resolution images of the graphene nearby the hexagonal holes (e.g., shown in Figure 7c) we can infer that the edges of the holes are of the armchair type. This is in agreement with the notion that armchair edges are more stable than zigzag edges that are known to be only metastable.48 The formation of these holes is considered the onset of the etching of the graphene discussed above in the XPS measurements.
4. CONCLUSIONS In summary, we used surface science methods to investigate a graphene/metal oxide interface. From XPS and STM measurements it is apparent that vacuum annealing at ∼700 °C results in a controlled interface without adsorbates, which results in stable C 1s and substrate core-level peak positions. Under these conditions charge transfer between the STO and graphene is observed, making the graphene p-type. STM demonstrates that graphene closely adheres to the substrate in agreement to previous observations on SiO2 and other substrates. This close adhesion of graphene to the substrate extends to atomic-scale substrate reconstructions. On STO we show that a quasi-oneF
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dimensional c(6 × 2) substrate reconstruction results in a wavelike modulation of the graphene morphology, suggesting the possibility of modifying graphene properties by suitably corrugated substrates. Prolonged annealing in vacuum at high temperatures (∼800 °C) can result in an etching of the graphene most likely by oxygen from the STO bulk reacting with free edges of the graphene.
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ASSOCIATED CONTENT
S Supporting Information *
(i) Description of sample dimensions and graphene coverage; (ii) characterization of CVD grown graphene on Cu foil and transferred to SiO2/Si substrates; (iii) comparison of C 1s peak shape for graphene on copper and after transfer to STO (annealed to 700 °C in vacuum). This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Author
*E-mail
[email protected] (M.B.). Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS Financial support from the National Science Foundation (NSF) under Award # DMR-1204924 and CHE-0840547 and from the Department of Energy Basic Energy Sciences (DOE-BES) under Grant DE-FG02-09ER16082 is acknowledged.
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