Axel Wilson
email: axel.wilson@diamond.ac.uk

Aude Bailly

Romain Bernard

Yves Borensztein

Alessandro Coati

Bernard Croset

Hervé Cruguel

Ahmed Naitabdi

Mathieu G Silly

Mathieu SILLY c Marie-Claire Saint-Lager

Alina Vlad

Nadine Witkowski

Yves Garreau

Geoffroy Prevot
email: prevot@insp.jussieu.fr

Gas-induced selective re-orientation of Au-Cu nanoparticles on TiO 2 (110)

come

I) Introduction

Supported metallic nanoparticles (NPs) have a prominent place in the field of catalysis.

Haruta's discovery showing that Au NPs activity exceptionally improves when their size is reduced down to a few nm 1,2,3 triggered an intense and fruitful research activity in catalysis. 4,5,6,7,8 This activity was further stimulated by studies showing that the addition of a second metal could improve the catalytic properties of NPs by either promoting their activity and selectivity or increasing their stability. 9,10,11,12,13 The enhancement of the activity and/or the stability of the NPs is often attributed to synergistic effects where the presence of more than one metal in the catalyst improves its catalytic properties beyond the sum of the properties of the different metals tested separately. In the case of Au-Cu NPs it has been shown that the addition of Cu to Au has a promoting effect on different reactions such as carbon monoxide oxidation reaction, 14,15,16,17,18 PROX reaction (preferential CO oxidation in the presence of hydrogen), 15,19,20,21,22,23,24 propene epoxidation, 25,26 selective hydrogenation of butadiene 27,28 and cinnamaldehyde 29 . More recently Au-Cu NPs have shown a promising efficiency in the challenging fields of fuel production from carbon dioxide 30,31,32,33 and of wastewater treatment, 34 combining the catalytic properties of the NPs and the photo-catalytic properties of a TiO 2 support. However, synergistic effects alone do not explain all the discrepancies between monometallic and bimetallic NPs regarding the catalytic performance and it is often difficult to determine the origin of this difference. 35 Several studies have shown that in addition to the size, the structure of the NPs and the ability of Au and Cu to form an alloy, and/or Cu to form an oxide, have a crucial influence on the activity of the NPs. For Au-Cu NPs deposited on silica gel 15 and SBA-15, 14 X-Ray Diffraction (XRD) spectra measured after reduction under H 2 at 820K for a series of composition ranging from pure Au to pure Cu showed the formation of alloyed particles of about 3 -4 nm in diameter. In the case of TiO 2 supported Au-Cu NPs, the presence of a Cu oxide layer sandwiched between the substrate and the Au NPs has been shown to improve the activity of the NPs. 16 It was thus suggested that the synergistic effect occurs between Au and Cu oxide. 36 Bauer et al. have compared the activity of Au-CuO x /SiO 2 NPs to Pt/Al 2 O 3 NPs for the CO oxidation below 150°C and measured a significant improvement for Au-CuO x /SiO 2 as well as a stabilisation attributed to an 'anchoring' effect of the NPs over the silica support. 37 Delannoy et al. reached a similar conclusion for selective oxidation of butadiene by Au-Cu NPs on TiO 2 . 28 An initial activation phase of the catalyst was observed during the reaction which was tentatively attributed to the segregation of Cu at the surface of the NPs. The influence of catalyst pre-treatments such as reduction and calcination have recently been investigated by Liao et al. for the PROX reaction for Au-Cu on CeO 2 23 and on Al 2 O 3 NPs. 24 In both cases, a reduction pre-treatment leading to alloyed NPs was shown to improve the CO conversion and the stability of the NPs in the presence of reactants. Liquid phase selective hydrogenation of cinnamaldehyde by Au-Cu/CeO 2 NPs was investigated in mild conditions (70°C, 1 atm). 29 The structure of the NPs after the reduction pre-treatment was investigated using HAADF-STEM and was

shown to form alloyed NPs with no visible segregation.

This highlights the growing interest for in-situ experiments where the structure of the NPs is continuously monitored as the system goes through the whole catalytic process from activation up to poisoning. Grazing incidence x-ray diffraction (GIXD) using synchrotron radiation is an excellent technique to carry out this type of investigation as the low incidence angle allows to be extremely sensitive to the surface and the high photon energy used (5 -25 keV) allows the beam to penetrate dense reactive phases (gas or liquid) with limited absorption. In this study, in-situ GIXD was used to fully characterize the structure of Au-Cu NPs deposited onto TiO 2 (110) single crystals. The evolution of the structure of each type of Au-Cu NPs was monitored in-situ during subsequent and simultaneous exposure to oxygen and carbon monoxide. The structural characterization was then complemented with x-ray photoemission spectroscopy (XPS) investigation that corroborates GIXD findings and unveils the diffusion mechanisms leading to the formation of Cu oxide that may explain the enhanced activity of bimetallic Au-Cu NPs.

II) Experiment

All experiments were performed on UHV set-ups with a base pressure of a few 1 × 10



. The NPs lattice parameters were determined for each peak using the centre of a Gaussian fit. The size of as-prepared NPs were determined after the synthesis using the Scherrer equation 38 and a height/diameter parameter for the NPs of 0.89, obtained from previous investigations. 42 However, due to the experimental set-up utilised (point detector and detector slits) and the fact that the aspect ratio of the crystalline domain of the NPs was evolving during the experiment, no adequate parameter was found to input in the Scherrer formula to satisfactorily represent the evolution of the size of the NPs.

XPS experiments were performed at the TEMPO beamline of SOLEIL synchrotron. Au 4f and Cu 3p core levels were recorded in the same kinetic energy window using normal emission of the photoelectrons and photon energies of 250 eV, 700 eV and 1050 eV with an overall resolution E/E of about 5000. Intensities were normalized to the incoming photon flux recorded on a gold mesh located prior to the analysis chamber. Shirley background has been removed from all XPS spectra.

The binding energy (BE) has been calibrated with respect to the Fermi energy measured on the Ta stripes holding the sample.

III) GIXD results

a) NPs structure under UHV

Hereafter, we present the results obtained under UHV on a reference sample composed of Au 52 Cu 48 NPs of about 2 nm in size grown on a rutile TiO 2 (110) substrate. Those NPs have been obtained by the successive evaporations of 9.5 × 10 14 at.cm -2 of Au at 570K and 8.8 × 10 14 at.cm -2 of Cu at room temperature. This procedure has previously been shown to lead to the formation of bimetallic particles. 42 Figure 1(a) shows a (h, k, l=0.05) map of the intensity scattered by the sample in the reciprocal space. Intense sharp peaks (red/black) matching integer values of h and k, correspond to Bragg reflections of the TiO 2 substrate. The broad spots at non-integer values of h and k and the ring centred at the origin are due to the scattering by the NPs. The broad spots are attributed to crystalline NPs that have grown on the TiO 2 (110) substrate with specific epitaxial relationships detailed in figure 1(b). The ring corresponds to randomly oriented NPs. In the following, the first type of NPs will be referred as epitaxial NPs and the second type as randomly oriented NPs. Figure 1(b) is a schematic representation of the different features present in figure 1(a). The black spots correspond to the substrate Bragg peaks, blue, red and green spots to diffraction peaks of oriented NPs and the purple arc to the signal obtained from randomly oriented NPs. at.cm -2 of Au and 1.4 × 10 15 at.cm -2 of Cu on a TiO 2 substrate at 415 K have been exposed, at room temperature, to CO and to a CO+O 2 mixture. In order to follow the evolution of the structure of each type of NPs the h-, k-and hk-reference scans corresponding to the red lines shown in figure 1(b) have been acquired repeatedly before and during gas exposure.

Figure 2 shows the comparison between the same h-, k-and hk-scans acquired in UHV after growth (green triangles in figure 2(a),(b) and (c) respectively), after a first CO exposure (P CO = 1 × 10 - 6 mbar; red squares) and after a subsequent exposure to a CO + O 2 (1:1) mixture (P CO+O2 = 2 × 10 -6 mbar; blue circles). The structure of the as-prepared Au-Cu NPs is determined from the reference scans realized under UHV condition (green triangles). The peak around 28.7 nm -1 , which is present on all the scans, corresponds to the (111) ring feature previously ascribed to randomly oriented NPs on the substrate. In the case of UHV, it is associated with a lattice parameter of 0.380 nm indicating that NPs are alloyed. The presence of a small (220) common peak on the h-scan around 45.6 nm -1 shows the presence of epitaxial NPs. However, the absence of the (220) 001 peak in the k-scan and the weak intensity of the (220) 111 peak in the hk-scan as compared to the ring intensity, suggest that most of the NPs are randomly oriented. A lattice parameter of 0.390 nm was extracted from the (220) common peak showing that epitaxial NPs are also alloyed. From the width of the diffraction spots, we estimate that the randomly-oriented particles have a mean size of 1.8 nm, whereas (112) epitaxial nanoparticles have a mean size of 2.0 nm. The difference in lattice parameter between epitaxial and randomly oriented NPs could either be attributed to a slight difference in composition or to a different of strain between the two types of NPs. We have previously shown that for small Au particles, a decrease of the lattice constant is observed when the size decays. 42,43 In the present case, we do not observe significant differences between the value of the position of the (111) peak found for randomly-oriented particles and for (112) oriented particles. The difference with the position of the (220) common peak could thus more probably reflect a lattice distortion of the cubic lattice. However the proximity of the (220) common peak and the TiO 2 Bragg peak may influence the value of the fit. showing that a reversible process is occurring between oxygen introduction and annealing.

IV) XPS results

Two other samples have been prepared for XPS experiments with compositions of Au 27 Cu 73 and Au 53 Cu 47 . The XPS spectra of Au 4f and Cu 3p core-levels, recorded on these two samples for a photon energy of 250 eV, are displayed in the figure 6. The top spectra corresponds to the Au 27 Cu 73 sample, whose composition has been obtained from Rutherford Backscattering Spectroscopy (RBS) measurement after the removal of the sample from the XPS chamber. The Au 4f and Cu 3p intensities have been extracted by the deconvolution of each component using Voigt functions.

Taking into account the cross section at 250 eV of photon energy [START_REF] Yeh | Atomic Data and Nuclear Data Tables[END_REF] and the mean free path of electrons in copper and gold, [START_REF] Powell | NIST Electron Inelastic-Mean-Free-Path Database -Version 1.2[END_REF] the quantitative analysis of Au 4f and Cu 3p leads to a composition of Au 24 Cu 76 , close to the one obtained using RBS. This validates the procedure used to determine the atomic fraction of Cu and Au by The bottom spectra of the figure 6 have been recorded on a second sample, with Au 53 Cu 47 composition, as given by the flux of evaporation. This composition is confirmed by XPS analysis at 250 eV photon energy, giving Au 51 Cu 49 . From XPS, we can exclude the formation of a large amount of a CuO oxide, which would have led to a chemical shift larger than 1 eV on the Cu 2p and Cu 3p peaks. The small chemical shifts observed with XPS may be associated with the formation of a shell of Cu 2 O or to chemisorbed oxygen. However, the absence of significant shift on the Auger peak indicate that most Cu atoms remain in a metallic state. This high stability towards oxidation could be due to the low O 2 partial pressure used in the present experiments. It has been shown using Scanning Tunneling Microscopy (STM) that the morphology of Au-Cu NPs of about 1 nm in size remains stable during the introduction of 10 -6 mbar of oxygen if they incorporate more than 20% of Au. [START_REF] Wilson | [END_REF] The fact that the individual volume of the NPs remains the same means that the Cu does not diffuse away from the particles but forms a core-shell structure with a Au-rich core and a Cu 2 O shell too poorly crystallised to scatter significant intensity. Such a phenomenon was studied using XPS and density functional theory (DFT) for Au-rich polycrystalline surfaces by Völker and co-workers 47 and for Cu 3 Au(111) by Tsuda and co-workers. 48,49 In both cases the XPS/DFT interplay shows that the alloyed surfaces under UHV are slightly enriched in Au and that a clear inversion occurs during the introduction of oxygen with Cu diffusing to the outermost layers of the surface. More recently, Dhifallah and co-workers have suggested using firstprinciples calculation that the Cu segregation may also depend on the surface orientation. 50 b) Thermal stability of the core/shell structure In each of the GIXD experiments presented in figure 4 and 5 the gas has been pumped down and the system maintained under UHV at room temperature to investigate a possible diffusion of the Cu atoms back into the Au enriched core of the NP. Indeed, it has been shown that Cu atoms deposited onto ≈ 2 nm Au NPs at room temperature diffuse into them to form an alloy. 42 The fact that such diffusion could not be observed reinforce the interpretation made previously that the copper forms an oxide shell when exposed to oxygen. This shell is stable at room temperature, under UHV or when exposed to 10 -6 mbar of CO (see figure 5). This interpretation is in agreement with the conclusion of Sandoval and co-workers 18 regarding the formation of Cu oxide but in the present case, XPS demonstrate that the amorphous Cu oxide cannot be 'sandwiched' between the substrate and Au clusters. Moreover, in such a configuration, the epitaxial relationship between the NPs and the substrate would have been lost, which is not the case here.

The Cu oxide shell can be reduced by annealing the sample as shown in figure 4 and 5 at a temperature above 430 K. The decrease of the lattice parameter is therefore interpreted as Cu atoms being progressively released from the shell and forming a Au-Cu alloy within the NP. As expected from this scenario, the re-introduction of oxygen at room temperature after annealing regenerates the Cu oxide shell through a segregation phenomenon as seen in figure 5 Two interpretations can be made to explain the total disappearance of randomly oriented NPs in the presence of the gas mixture. i), a NP to NP diffusion mechanism occurs where randomly oriented NPs, which may also be the smallest, sinter. Au and Cu atoms diffuse from these small NPs toward larger, epitaxial and more stable NPs, resulting in an increase of the corresponding diffracted intensity. ii) an intra-atomic mobility triggers the NP re-crystallisation, favouring a (111) NPs //(110) TiO2 epitaxy. The second hypothesis is more likely because a ripening mechanism cannot explain that a specific epitaxial relationship is favoured. The irreversibility of the disappearance of the ring feature, even after annealing, is in agreement with both previous interpretations. The disappearance of randomly oriented NPs is probably triggered by the fact that the chemical potential is higher for these NPs than for epitaxial ones. This could be due to a weaker interfacial bonding or to a smaller size (1.8 nm instead of 2.0 nm). On the one hand, the adhesion energy of a Au particle on TiO 2 (110) has been shown to depend on its geometry. 51 On the other hand, from the surface energy of Au and Cu 52 , a difference in chemical potential of 0.05 eV/at is expected between 1.8 nm and 2.0 nm size Au 50 Cu 50 particles.

At this point, one has to determine whether the disappearance of randomly oriented NPs is induced by the mixture of gas or if such effect could be seen for pure oxygen exposure. The segregation mechanism has been shown to occur in the presence of oxygen only (see figure 5) and if, as stated above, the re-orientation mechanism was a consequence of Au and Cu atoms segregating it would happen in these conditions as well. However, if this phenomenon has been observed under a CO+O 2 mixture on the Au 38 Cu 62 sample, it has not been measured on the Au 49 Cu 51 sample under pure O 2 . Nevertheless, in that case, a higher Au content may enhance the NP stability. We have indeed previously shown by STM that Au x Cu 1-x NPs were stable under 10 -5 mbar O 2 for x>0.2. [START_REF] Wilson | [END_REF] It is worth to note that the atomic mobility needed for the reorientation of the NPs is greater than the one required for Cu segregation to the surface, since it implies that all atoms, including those at the interface with the substrate, are involved. Au-Cu/TiO 2 NPs have also been shown to catalyse the carbon monoxide oxidation at room temperature and above. 18 Thus, considering that the reaction is exothermic (2.93 eV per CO 2 molecules), one can suggest that the energy released by the reaction, occurring here at room temperature, leads to an enhanced atomic mobility, and then to a gas-induced selective re-orientation of the NPs at the pressure at which this study has been performed.

VI) Conclusion

The combined use of GIXD and XPS allows us to comprehend the atomic diffusion and oxidation mechanisms at play in the case of 2 nm Au-Cu NPs exposed to oxygen and carbon monoxide. When exposed to 10 -6 mbar of oxygen the Cu atoms, initially alloyed with Au, segregate to the surface of the NPs to form a Cu oxide shell stable in UHV at room temperature and reducible by annealing the sample over 430K. In the presence of a CO + O 2 mixture, this segregation phenomenon triggers a second phenomenon referred here as gas-induced selective re-orientation where randomly oriented NPs use the mobility induced by the segregation to re-orient their Au-rich core following the specific epitaxial relationship (111) NPs //(110) TiO2 , [1-10] NPs //[001] TiO2 . It remains unsure whether the presence of carbon monoxide is mandatory to trigger the re-orientation mechanism. After reduction, the alloyed NPs keep their epitaxial relationship making this procedure an efficient way to selectively orient Au-Cu NPs on a rutile TiO 2 (110) substrate.

  mbar. Rutile TiO 2 (110) single crystals (from MaTecK GmbH) were cleaned by several cycles of Ar sputtering (P Ar = 7 × 10 -5 mbar, 20 minutes, E = 1 keV) and annealing under UHV (10 minutes at 1000 K), leading to slightly reduced samples. Au and Cu were evaporated in situ under UHV either from ebeam evaporators or from effusion cells, using high purity Au and Cu (Alfa Aesar). The amounts of Au and Cu evaporated were measured during the experiments with a flux monitor in the case of the ebeam evaporators and with a quartz microbalance in the case of the effusion cells. During each experiment a calibration sample was realized by depositing a large amount of Au and Cu (typically 2 ML of each) on a TiO 2 sample. The quantity of Au or Cu present on the calibration samples was precisely determined using Rutherford backscattering spectroscopy (RBS) performed with a van de Graaff accelerator at INSP. An excellent agreement was found between doses estimated using the flux monitor or the quartz microbalance and the values obtained using RBS. Ultrapure O 2 and CO were introduced in the chambers using precision leak valves. GIXD experiments were performed at the SixS beamline of SOLEIL synchrotron and at the BM32 beamline of ESRF synchrotron. GIXD measurements were performed at an energy of 15 keV and a grazing incidence angle of 0.2°. The x-ray diffraction data presented in this study are indexed according to the surface unit cell of TiO 2 (110) described by the following parameters: a = 0.296 nm, b = c = 0.648 nm, and α = β = γ = 90°. h, k, and l Miller indices are used for indexing reflections in the reciprocal space. As a matter of clarity some of the data have been plotted against the modulus of the momentum transfer corresponding to

FIG

  FIG 1: (a Au 52 Cu 48 Bragg r scattere the direc

  FIG 2. Di direction h=0, l=0 mbar CO rectangl vertical

  Figure 4 (respect increase

Figure 5 :

 5 Figure 5: Time of evolution of: a) lattice parameter of Au 51 -Cu 49 NPs measured from the position of the (220) common peak, (b) O 2 (blue) and CO (purple) pressures and (c), temperature.

  Figure 6 at a pho

  during the second introduction of oxygen. c) Gas-induced selective re-orientation/crystallisation of the NPs The disappearance of the ring feature upon introduction of a mix of CO + O 2 implies a deep and global reorganisation of individual NPs. The sample presented in section b) for which the reorientation was observed, was synthesized under UHV on a substrate at 415 K to improve the diffusion of individual Au and Cu atoms and obtain larger NPs. Despite this procedure, the signal scattered by randomly oriented NPs was the majority. In other experiments, we observed that sample annealing up to 670 K in UHV does not change the intensity of the ring and does not improve the epitaxial relationships of the NPs (not shown here).

Acknowledgements

Axel Wilson has received funding from the European Union's Horizon 2020 research and innovation programme under the Marie Skłodowska-Curie grant agreement (GA) No 665593 awarded to the Science and Technology Facilities Council. Experiments were performed on the SixS and TEMPO beamlines at SOLEIL Synchrotron, France and on the BM32 beamline at the ESRF, France.


Fethullah Günes

Hakim Arezki

Debora Pierucci

David Alamarguy

José Alvarez

Jean-Paul Kleider

Yannick J Dappe

Abdelkarim Ouerghi

Mohamed Boutchich
email: mohamed.boutchich@geeps.centralesupelec.fr

TUNING THE WORK FUNCTION OF MONOLAYER GRAPHENE ON 4H-SIC (0001) WITH NITRIC ACID

Keywords: Graphene, Work function, doping, Raman spectroscopy, UPS spectroscopy, DFT calculations

Chemical doping of graphene is a key process for the modulation of its electronic properties and the design and fabrication of graphene based nanoelectronic devices. Here, we study the adsorption of diluted concentrations of nitric acid (HNO 3 ) onto monolayer graphene/4H-SiC (0001) to induce a variation of the graphene work function (WF). Raman spectroscopy indicates an increase in the defect density subsequent to the doping. Moreover, ultraviolet photoemission spectroscopy (UPS) was utilized to quantify the WF shift. UPS data show that the WF of the graphene layer decreased from 4.3 eV (pristine) down to 3.8 eV (30% HNO 3 ) and then increased to 4.4 eV at 100% HNO 3 concentration. These observations were confirmed using density functional theory (DFT) calculations. This straightforward process allows a large WF modulation rendering the molecularly modified graphene/4H-SiC(0001) a highly suitable electron or hole injection electrode.

Introduction

Graphene is a 2D-material combining various outstanding properties in a single material, like quantum Hall effect or massless Fermions among others 1,2 . Its high mobility, optical transparency, along with flexibility, high mechanical strength and environmental stability make it a very attractive material for a broad range of applications [3][4] . Graphene can also be exploited for precise metrological applications, as it has been demonstrated recently with exceptionally accurate measurements of the quantum Hall resistance quantization 5,6 .

In order to implement the potential of such material into viable applications, large-scale wafers of high quality graphene grown on insulating substrates are required for device integration 7,8 . Moreover, to achieve performant devices, the electronic properties of the material must be modulated 9,10,11 . In that respect, the work function (WF) is a remarkable property that can be tuned for various optoelectronic applications like transparent electrodes for optoelectronic devices, or multiple tunnel heterostructures matching the energy offsets of common transport materials, to achieve low and/or large charge injection barriers 12 . To reach this goal, the modification of the work function of graphene has to be obtained without damaging the overall characteristics of the material. A possible approach to vary the WF is to modify the surface by depositing a monolayer of a molecular donor or acceptor. Recent studies have demonstrated possible WF variations using as acceptor materials tetrafluoro-tetracyanoquinodimethane (F4-TCNQ), gold colloids or metals 13,14,15 for example.

Although these methods have demonstrated WF modulation, they present major drawbacks. On one hand, the transparency of graphene is reduced, which represents a strong limitation for optoelectronic applications where graphene could play a role. On the other hand, these adsorbed elements create scattering interfaces that may severely affect the transport properties of graphene. At last, deposition and temperature dependence are potential limitations for the device fabrication process. Among the dominant graphene synthesis methods, chemical vapor deposition (CVD) and epitaxy on SiC substrate are the leading technologies in terms of scale and properties. For electronic applications, the growth of graphene by CVD on metals requires transferring the graphene onto an insulating substrate. This last operation involves chemicals such as PMMA and copper etchant known to induce a p-type doping of the material 16 . Although this approach allows the transfer of graphene on any functional substrate, the electronic properties of the graphene sheet are altered by these manipulations. Furthermore, epitaxial graphene on hexagonal silicon carbide (SiC) wafers presents homogeneous films with large areas and high electron mobility 17 . The SiC substrate is suitable for opto-electronic applications since it is transparent over a very broad frequency spectrum. It has a large thermal conductivity, and is compatible with highfrequency devices where losses due to residual conductivity of the substrate have to be minimized by using insulating materials. We have recently demonstrated the WF modulation on trilayer graphene on SiC (0001) using in-situ nitrogen flux 18 . Furthermore, Das et al have demonstrated that Nitric Acid (HNO 3 /NA) on CVD graphene is an efficient doping agent inducing p-type variation of the WF 19,20 .

In this work, we have monitored the structural and electronic modifications of epitaxial graphene for different concentrations of diluted Nitric Acid (NA) 15 vol% (NA15), 30 vol% (NA30), 70 vol% (NA70) and 100% (NA100) of nitric acid (NA) solutions in deionized water (DI water) for 2min. The electronic properties of the doped graphene samples were characterized by Raman spectroscopy. The changes in work function after NA doping were measured using ultraviolet photoemission spectroscopy (UPS). Based on these observations, the influence of NA doping on graphene work function is discussed. To highlight the effect of nitric acid doping, these measurements were compared with the obtained values for pristine graphene. We demonstrate in the present work that the NA doping permits the WF modulation with respect to the dilution. Indeed, the WF first increases the electronic doping between 15 and 30 vol% (NA15, NA30) and then decreases the electronic doping between 70 and 100 vol% (NA70, NA100). This study on epitaxial graphene presents striking differences compared to the work of Das et al on CVD graphene where only p-type work function shifts is observed. We clearly demonstrate that both systems react differently to an identical chemical functionalization.

These observations are supported by density functional theory calculations (DFT).

Experimental section

The graphene was grown on semi-insulating SiC(0001) substrates 17,21 . The SiC sample was then heated to 1000 °C in a UHV chamber and further heated to 1525 °C in an Ar atmosphere to favor the formation of large and homogeneous graphene layer. The substrate was then cooled down from 1525°C to room temperature 18,22 . The samples were subsequently transferred ex situ from the growth chamber to the AFM microscopy, Raman spectroscopy and UPS measurement systems.

The morphology of the nitrogen doped epitaxial graphene was studied after growth by atomic force microscopy (AFM). The Raman measurements were carried out with a frequency-doubled NdYag laser operating at 532 nm as excitation source with up to 20 mW continuous wave output power. The spectra were recorded through a 300 mm imaging spectrometer equipped with both a 600 lines/mm and 1800 lines/mm grating, and a backilluminated CCD. A 100X objective with respective numerical aperture of 0.9 and 1.4, and a 50-µm core diameter multimode fiber acting as a pinhole, were used to collect the signal. The expected lateral spatial resolution in the x-y plane can reach 150 nm (num app=1.4) and 250 nm (num app=0.9). A particular attention was paid to the incident laser power density to prevent local heating effects that may induce a shift of the Raman peaks. UPS measurements for the study of the work function have been performed using a PHI 5000 Versaprobe spectrometer (Physical Electronics) operating at a base pressure of 10 -9 mbar, using a He I discharge lamp (hν = 21.2 eV).

Results and discussion

The typical morphology of the graphene sample is displayed in the AFM images of Figure 1. The step direction and terrace width were determined by the incidental misorientation of the substrate surface with respect to the crystallographic (0001) plane. The surface is highly uniform over a large scale with atomically flat, 3 µm large terraces and 1.8 nm high steps. On defect-free areas of the sample, the terraces typically extend undisturbed over 50 µm in length. In AFM phase mode (Figure1 (b)), the homogenous contrast indicates the thickness homogeneity, that corresponds to a graphene monolayer.

Figure 2 displays the typical spectra over the range 1200-3000 cm -1 of the pristine graphene and NA doped graphene samples 23 . For each sample, to reduce the possible errors affecting the peaks positions and the intensity due to local inhomogeneity, we collected the average Raman spectra over an area of 1µm 2 , before (pristine) and after the doping. We observed the main features of graphene as well as the second-order Raman signal that originates from the SiC substrate visible between 1200 and 3000 cm -1 . Graphene contributes to Raman spectra with three main peaks: i) D band at ~1350 cm -1 (defect induced mode), ii) G band at ~1590 cm -1 (in-plane vibrational mode) and, iii) the 2D band (two-phonon mode) at ~2700 cm - 1 24,25 . The pristine sample does not present any noticeable D band indicating a low density of defects and a high crystalline phase i.e. I D /I G ratio ~0.13. On the other hand, all NA doped samples exhibit a higher intensity D band. The I D /I G ratio increases from ~ 0.13 for the pristine sample to an average of 2.5 for doped samples confirming the alteration of the structure after NA dip. In addition, we note the appearance of a (G+D) band ~2950 cm -1 caused by the disorderinduced feature in sp 2 carbon associated with defects. The 2D band full width half-maximum (FWHM) increases from 35 cm -1 for the pristine sample to an average of 42 cm -1 for NA15, NA30 and NA70, and raised to 50 cm -1 for the maximum doping for NA100. These Raman shifts clearly indicate a doping effect induced by the NA dip at various concentrations. The bending of the graphene sheet at the kinks of the terraces induces local physical stress on the carbon atoms of the graphene layer that results in an increase in the chemical reactivity of these atoms, and then we can suppose that the incorporation of dopants occurs at the edges states of the epitaxial graphene 26,27 .

We observe that, for the NA15 sample the position of the G-band before and after doping is within the resolution of the Raman system i.e. ~ 1 cm -1 . However, there is a red shift of 14 cm -1 in the position of the 2Dband assigned to an n-type doping. The defect induced D band also increases with respect to the pristine sample as expected for doped samples. For this sample, the D band intensity was not changed, mainly because the spot size of the confocal Raman was smaller than the size of the terraces affected by the edges of epitaxial graphene, thus inducing a D-band intensity even before doping. In the case of NA30 sample, we note a significant red shift for both the G-band and the 2D-band positions, of 8 and 19 cm -1 respectively, showing a strong n-type doping behavior 28 . The D-band also shows a prominent intensity increment after doping with NA30, which is due to the sp 3 -activation by nitrogen intercalation in the conjugated sp 2 network of graphene. Note that the epitaxial graphene is intrinsically n-doped due the charge transfer occurring from the SiC substrate through the interface layer 29 . This point will be further discussed in the UPS section. NA70 sample does not present a relevant shift in the G-band position. However, the red shift of the 2D-band is reduced to 6 cm -1 , denoting a reduced n-type doping effect as opposed to NA15 and NA30. NA100 shows a different behavior as the G-band position remains stable but the 2D band exhibits a blue shift of 11 cm -1 indicating p-type like doping 30 . We believe that this effect at high NA concentrations results from a balance between two different simultaneous effects, i.e. p and ntype doping from physisorption of the molecules and nitrogenation reactions near the terraces, respectively.

To evaluate the electronic quality of our graphene layers, we determined the carrier mobility of the pristine and NA100% doped graphene using van der Pauw measurements at room temperature. The electron density in several areas of these samples was typically 9 x 10 12 cm -2 and 6 x 10 12 cm -2 respectively, with a macroscopically averaged electron mobility inferred from the Hall voltage estimated at about 2500 cm 2 /(Vs) and 1700 cm 2 /(Vs) respectively. We have measured in previous works the electronic properties on such a monolayer graphene using the same growth mechanism with hydrogen, nitrogen introduced in-situ and ammonia (liquid or vapor) as doping media 18 , 22 . The measurements showed outstanding properties including quantum Hall effect and the high mobility reached up to 11000 cm -².V -1 .s -1 at 1.7K for carrier density ~10 12 cm -². The lowest mobility achieved was 1400 cm -².V -1 .s -1 at 300K 17 . Note that Günes et al 14 achieved a sheet resistance of 100Ω/o using HNO 3 doping on CVD graphene. These data clearly indicate that although doping alters to some extent the lattice of graphene, it does not render the material electronically poor. The orders of magnitude are sufficient for many mainstream applications where contacts are required 31 , 32 .

For that purpose and to complete the picture of the electronic properties of doped monolayer epitaxial graphene, we performed UPS measurements in order to monitor the WF as function of the NA concentration.

The UPS spectra around the secondary electron threshold region of the graphene layer treated by NA solutions are shown in Figure 3(a). The secondary electron threshold was determined by extrapolation between the background and straight solid lines in the secondary electron threshold region of the UPS spectra. The work function was determined from the secondary electron threshold using the equation Φ = hν -E th , where hν and E th are the photon energy of excitation light (He I discharge lamp, 21.2 eV) and the secondary electron threshold energy, respectively. Figure 3(b) shows the work function of graphene pristine and HNO 3 doped graphene. Decomposition of HNO 3 molecules takes place upon the interaction with the graphene layer where the active edge-state electrons play a catalytic role. While the HNO 3 concentration increases, the amount of species physisorbed onto the graphene layer increases. The work function of pristine graphene layer is 4.3 eV as reported by different groups 33,34 . Subsequently, the work function decreases from 4.3 eV to 4.1 eV after doping with 15% NA. The WF further decreases to 3.8 eV with 30% NA. Note that this reduction implies an n-type doping that is consistent with the variation observed on the Raman spectra on Figure 2, suggesting an increase of the nitrogen content between pristine and 30% NA. Beyond 30% NA, the WF increases first to 4.05eV at 70% NA and reaches 4.41 eV for pure NA. For the latter we also observed a blue shift on the Raman spectrum of NA100 confirming a p-type doping that compensates the initial n-type nature of the sample. The changes of the Fermi level position with the NA concentration is illustrated in Figure 3 (b). Our results clearly indicate that for low NA concentration, the nitrogenation reduces the WF towards the n-type doping as the nitrogen content (donor) increases in the sample 35 . This is in contrast to the observation of Das et al where the WF presents a monotonic increase as a function of NA%. Note that the graphene utilized here is initially n-type due to a residual doping from the SiC substrate. We believe that the first HNO 3 molecules dissociate and saturate all available defects sites or steps edges. This induces a nitrogen doped graphene layer (NA15% and NA30%) and therefore decreases the WF 35 . On the other hand, beyond NA30%, the HNO 3 can physisorb on the surface inducing a p-type doping, lowering the Fermi level and as a result increasing the WF. This was confirmed by ab initio calculations describing the electronic structure of graphene on SiC as well as the interaction with the HNO 3 molecules were performed using a very efficient DFT localized orbital molecular dynamic technique (FIREBALL) 36 . Numerical localized orbital basis sets of sp 3 (Si) and s = 4.1 (H) 37 . In this study we have considered supercells of 5 ML SiC(0001), with a zero-layer graphene and a AB stacked graphene plane on top. The lateral size roughly corresponds to a 4x4 unit cell of graphene. The bottom layer is saturated with hydrogen atoms. On top of the supercell, we have set from 1 to 4 HNO 3 molecules per unit cell to observe the graphene doping evolution with the molecular concentration, as represented in Figures 4 (a), (b) and(c). The geometry of those systems has then been relaxed at 0 K, using a sample of 32 k-points in the surface Brillouin zone, as previously used in similar calculations 38 , maintaining the last three bottom layers in bulk positions. The equilibrium configuration is obtained when forces are less than 0.01 eV/Å. The final distances between the molecules and the graphene plane have been determined using the well-known LCAO-S 2 + vdW formalism 39 , 40 , 41 which specifically takes into account van der Waals interaction in the frame of DFT. Finally, a set of 300 special k points along the Γ-K-M path has been used for the band structure calculations on the optimized geometries. These calculations aim at reproducing the physisorbed process leading to a p-doping, as we did not model the graphene terraces. Results of the band structure calculations are represented in Figures 4 (d),(e) and (f), corresponding to one, two or four deposited HNO 3 molecules. We can observe a variation of the graphene doping from -350 meV for one molecule (which represents already a p-type doping with respect to the pristine situation at -400 meV), to -180 meV for two and until +190 meV for four molecules. Consequently, the calculations also confirm a p-type doping for the physisorption process in good agreement with the WF variation recorded experimentally.

Conclusions

We have investigated the doping strength of nitric acid on monolayer graphene grown on SiC (0001). The Raman measurements confirm the doping induced by various concentration of nitric acid and indicate the incorporation of atoms into the graphene lattice. The UPS measurements showed that n-type and p-type doping are competitive processes and the work function of graphene layer can be continuously tuned from 4.3 eV to 4.4 eV with a minimum of 3.8 eV upon depositing the molecular acceptor HNO 3 . This observation is supported by DFT calculations considering the physisorption of HNO 3 molecules. This simple approach allows the modulation of WF and could permit the design of tuneable electrodes to improve the carrier injection as well as band gap-engineered heterostructures.

  for C, N, O and Si, and s for H were used with spatial cutoff radii (in atomic units) s = 4.5, p = 4.5 (C), s = 4.2, p = 4.2 (N), s = 3.3, p = 3.8 (O), s = 4.8, p = 5.4

  This work is supported by a public grant overseen by the French National Research Agency (ANR) as part of the "Investissements d'Avenir" program (Labex NanoSaclay, reference: ANR-10-LABX-0035), as well as the the French Ministère des affaires étrangères et européennes (MAEE), the Centre National de la Recherche Scientifique (CNRS) through the STIC ASIA programme grant 3226/DGM/ATT/RECH. The XPS/UPS equipment was supported by a grant from Region Ile-de-France.

Figure 1 :

 1 Figures captions:

Figure 2 :

 2 Figure 2: Raman spectra of pristine, NA15, NA30, NA70 and NA100 samples. D peak magnitude appears with NA concentration indicating the creation of defects due to nitrogen incorporation. (G+D) band indicate a similar effect. Raman shifts indicated by arrows evidence the doping trends. The presence of the G and 2D bands across the overall area indicates homogenous growth of graphene.

Figure 3 :

 3 Figure 3: (a) UPS data for pristine, NA30, NA70 and NA100 samples indicating the WF shift as a function of nitric acid concentration. (b) Illustration of the Fermi level shift in the graphene Dirac cones.

Figure 4 : 1 Figure 1 Figure 2 Figure 4

 41124 Figure 4: (a), (b), and (c) atomic structures used for the calculations, considering one graphene layer on top of a SiC(0001) surface with graphene buffer layer, and one, two and four HNO 3 molecules respectively, (d), (e) and (f), corresponding band structures showing the evolution of the graphene doping from n to p-type. The graphene Dirac cone is highlighted in red for the three configurations.