Single step fabrication of N-doped graphene/Si3N4/SiC heterostructures

Emilio Vélez-Fort1,2, Emiliano Pallecchi1, Mathieu G. Silly3, Mounib Bahri1, Gilles Patriarche1, Abhay Shukla2,

Fausto Sirotti3, and Abdelkarim Ouerghi1 ()

1 CNRS-Laboratoire de Photonique et de Nanostructures (LPN), Route de Nozay, Marcoussis 91460, France 2 Université Pierre et Marie Curie (CNRS – IMPMC), 4 Pl. Jussieu, Paris 75005, France 3 Synchrotron-SOLEIL, Saint-Aubin, BP48, F91192 Gif sur Yvette Cedex, France

Received: 24 November 2013

Revised: 28 January 2014

Accepted: 1 March 2014

© Tsinghua University Press

and Springer-Verlag Berlin

Heidelberg 2014

KEYWORDS

epitaxial graphene,

spectroscopy,

nitrogen-doped graphene,

low-energy electron

microscopy,

electronic properties

ABSTRACT

In-plane heteroatom substitution of graphene is a promising strategy to modify

its properties. The ability to dope graphene with electron-donor nitrogen

heteroatoms is highly important for modulating electrical properties of graphene.

Here we demonstrate a transfer-free method to directly grow large area quasi

free-standing N-doped graphene bilayers on an insulating substrate (Si3N4).

Electron-bombardment heating under nitrogen flux results in simultaneous

growth of N-doped graphene and a Si3N4 layer on the SiC surface. The decoupling

of N-doped graphene from the substrate and the presence of Si3N4 are identified

by X-ray photoemission spectroscopy and low-energy electron diffraction. The

substitution of nitrogen atoms in the graphene planes was confirmed using

high resolution X-ray photoemission spectroscopy which reveals several atomic

configurations for the nitrogen atoms: Graphitic-like, pyridine-like, and pyrrolic-

like. Furthermore, we demonstrated for the first time that N-doped graphene

could be used to efficiently probe oxygen molecules via nitrogen atom defects.

1 Introduction

Graphene is a novel two-dimensional material with

great potential for applications in nanoelectronics,

mainly due to its extremely high electron mobility,

which is a consequence of the high Fermi velocity and

the linear dispersion relation around the Dirac point

[1–3]. It is ambipolar and exhibits ballistic transport

on a micrometer scale up to room temperature. In

addition, graphene is also interesting for spintronic

devices, such as spin qubits and spin diodes, because

of its long spin-coherence time and long spin relaxation

time [4–6].

One of the most promising techniques to synthesize

graphene for electronic devices is the epitaxial growth

on silicon carbide (SiC). The epitaxial growth of

graphene offers the possibility to produce wafer scale

graphene domains and is compatible with silicon

technology [7, 8]. The idea behind this method is to

sublime the Si contained in the SiC by increasing

gradually the temperature of the SiC substrate. At

sufficiently high temperature, the remaining carbon

Nano Research

DOI 10.1007/s12274-014-0444-9

Address correspondence to abdelkarim. ouerghi@lpn.cnrs.fr

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2 Nano Res.

atoms rearrange and form a graphitic honeycomb

lattice on top of the SiC. On the Si face of a silicon

carbide wafer, the first carbon layer is strongly bound

to the substrate which drastically changes its electronic

proprieties. This layer, also referred to as the interface

or buffer layer, displays a large band gap and is

therefore insulating despite its honeycomb lattice

configuration. However, the next graphitic layer

which is grown on top of the interface layer shows the

well-known electronic proprieties of graphene, namely

a zero bandgap and the linear dispersion around the

K and K' points. Moreover, the presence of Si dangling

bonds underneath the buffer layer degrades the charge

carrier mobility due to the high intrinsic doping [9].

To succeed in developing high performance graphene-

based electronic devices, we must be able to dispense

with the interface layer without losing the exceptional

intrinsic carrier mobility of graphene [10–14]. Further,

doped graphene structures have attracted considerable

interest among the scientific community [15–20]. Like

other doped carbon forms, N-doped graphene is

expected, for example, to have enhanced properties for

Li-battery applications [21]. Reports of the synthesis of

N-doped graphene are scarce and the development of a

simple method to synthesize N-doped graphene is

needed. Nitrogen doping can enhance the conduc-

tivity of graphene and induce n-type semiconductor

behavior [22]. More importantly, incorporating nitrogen

atoms in different bonding configurations into the

carbon network would provide the N-doped graphene

with more functional groups for property design.

Ammonia-assisted doping methods have been reported

for graphene grown with different techniques such

as chemical vapor deposition, arc discharge of carbon

electrodes and thermal annealing of graphene oxide

[23–25].

Recent discoveries demonstrate that a widely

tunable bandgap can be realized in bilayer graphene

and bilayer graphene–boron nitride heterostructures

(h-BN), which opens a new door for applications of

graphene in electronic and optical devices [26, 27].

Here, we demonstrate an approach to grow quasi

free-standing N-doped graphene on a Si3N4 insulating

substrate. This result is obtained by exposing the

substrate to nitrogen gas during the graphene growth.

This method offers a number of advantages: (i) It is a

single step process; (ii) the tuning of the nitrogen flux

intensity can induce quasi free-standing growth; (iii)

the graphene is doped in situ, during the growth.

More precisely, we report on the formation of N-doped

bilayer graphene on SiC and its electronic properties.

The graphene layers were investigated by low-energy

electron diffraction (LEED), atomic force microscopy

(AFM), X-ray photoelectron spectroscopy (XPS), Raman

scattering and transmission electron microscopy

(TEM). The XPS spectra show three types of nitrogen

doping sites: Graphitic, pyridinic and pyrrolic nitrogen.

Decoupling of the buffer layer from the substrate

and the associated physical and electronic structural

changes are observed via LEED. The LEED and XPS

results demonstrate that a decoupled graphene bilayer

is formed on top of the SiC substrate. Moreover, we

show that a Si3N4 interfacial layer is formed by dilute

nitrogen situated between the graphene and SiC

substrate.

In the present work, the reaction scheme for direct

growth of quasi free-standing bilayer graphene on

insulating substrates is shown in Fig. 1(a). The N-doped

graphene and Si3N4 layers were formed on the SiC

substrate at high temperature, where different nitrogen

species were incorporated into the graphene and

the SiC substrate spontaneously. During the vacuum

annealing process, both carbon-trapped nitrogen

and silicon-trapped nitrogen atoms contribute to the

Figure 1 (a) Schematic illustration of our concurrent method for growing N-doped graphene/SiC. The nitrogen flux was used to reduce the desorption of Si atoms at high temperature. N-doped graphene was formed on the SiC surface after semi-vacuum (2 × 10–5 Torr) annealing at high temperature. LEED patterns at 124 eV for (b) pristine graphene and (c) N-doped graphene. (d) A typical AFM image of the doped sample after graphitization at 1,300 °C under N2 flux.

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3 Nano Res.

nitridation of the SiC surface and convert the epitaxial

graphene into N-doped graphene. Pristine graphene

(PG) was synthesized similarly by annealing a SiC

substrate under argon atmosphere. After growth, N-

doped graphene (NG) and pristine graphene (PG) can

be transferred ex situ for further characterization.

2 Methods

Graphene samples were grown on a n-type Si ter-

minated 6H-SiC(0001) wafer. The substrate was first

etched in a hydrogen flux at 1,500 °C at 200 mbar

for 15 min in order to clean the surface and form a

step-ordered structure. After etching, the substrate was

transferred into an ultrahigh vacuum (UHV) chamber

with a base pressure of 1 × 10–9 Torr. The surface of

the substrate was first degassed for several hours at

600 °C to eliminate all the impurities and then annealed

under a Si flux (2 monolayers per minute (ML/min)) at

900 °C to remove the native oxide. The graphitization

process was performed under a nitrogen flux and the

base pressure of the system was maintained around

2 × 10–5–2 × 10–4 Torr. The growth was performed by

annealing the surface with an electron bombardment

technique. The sample was heated gradually from

800 °C to 1,300 °C. After the growth of the graphene

layer, the sample was quenched to 500 °C under

nitrogen flux (1 ML/min) [28]. In order to compare our

results, pristine graphene was also synthesized with a

similar method but using an argon flux instead of the

nitrogen flux. The chemical, electronic and structural

properties of the resulting films were investigated

by means of XPS, LEED, AFM, TEM and Raman

spectroscopy. The micro Raman spectroscopy was

performed at room temperature with a Renishaw

spectrometer using 532 nm laser light focused on the

sample by a DMLM Leica microscope with a 100 ×

(NA = 0.75) objective and a power of 5 mW with a spot

size of about 1 μm. High-resolution core level X-ray

synchrotron radiation experiments were carried out at

the third generation SOLEIL storage ring (Saint-Aubin,

France) operated in the top-up mode at 430 mA, using

the photoemission end station of the TEMPO beamline

which covers a photon energy range of 50–1500 eV.

The experiments were performed at room temperature

with a base pressure of 2 × 10–10 mbar. Photoelectrons

were analyzed at a take-off angle of 90° with respect

to the sample surface using a Scienta SES 2002 hemi-

spherical analyzer and a delay-line detector [29].

3 Results and discussion

First, we characterized the graphene crystallographic

structure by low-energy electron diffraction. Figure 1(b)

shows the LEED image of pristine graphene. A single

phase of graphene layer is formed as deduced from

the very well ordered (6 3 6 3) R30° LEED pattern

that corresponds to a carbon-rich termination. The

smallest hexagons around the integer spots are due

to double diffraction as a result of a (6 3 6 3) R30°

reconstruction of the interfacial layer. These LEED

findings match previously reported results [8, 30].

Figure 1(c) shows the typical LEED pattern of the

doped sample after graphitization at 1,300 °C under

N2 flux. Two contributions can be clearly distinguished

in the LEED diagram. The first one arising from the

(1 × 1) graphene layer confirms the presence of this

material at the surface of the sample (blue arrows).

The second one, which is characterized by isotropic

(1 × 1) spots, is the 6H-SiC substrate (white arrows).

This finding indicates the presence of oriented single-

domain graphene patches forming a layer, and the

absence of rotational disorder and twinning domains.

Importantly, the absence of the six satellites spots in

the pattern of N-doped graphene clearly demonstrates

that the (6 3 6 3) R30° reconstruction is not present

in our sample. This structure is well known and

refers to the buffer layer (interface layer) [31]. This

result indicates that the buffer layer is not present for

epitaxial graphene grown under N2 flux in our growth

conditions. A careful analysis of the XPS (discussed

below) shows that N-doped graphene is a bilayer

while PG a monolayer. In order to assess the mor-

phology of the graphene layer, AFM was performed in

tapping mode. Figure 1(d) shows typical AFM image

of the sample. It is clearly observed that the terraces

on the SiC surface are atomically flat and separated

by straight step edges. The average terrace width is

600 nm. We also see white spots which are possibly

due to physisorbed nitrogen on the surface. We then

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measured the Raman spectra on the same sample at

room temperature shown in Fig. 2(a). The measure-

ments were performed for the N-doped bilayer

graphene (NG, red curve) on 6H-SiC(0001) and for the

pristine graphene (PG, black curve). The graphene

layer has several overtone peaks in the 1,000–2,000 cm–1

regime. The peak at 1,522 cm–1 is considered to be an

overtone of the L-point optical phonon from the

SiC. The Raman signals from the graphene show

prominent characteristic peaks at 1,620 cm–1 (G) and

2,732 cm–1 (2D), which give evidence of carbon sp2

reorganization [32]. The G peak value of 1,620 cm–1

could indicate strong N-doping. Also the blue-shifted

position of the 2D peak at 2,732 cm–1 indicates

compressive strain of the graphene layer during the

post-growth cooling down procedure or charge

transfer doping from the substrate. The blue shift of

the G and 2D band has also been reported in epitaxial

graphene on 6H-SiC wafers [33]. The 2D peak is

however much broader than in isolated graphene

with a full width at half maximum (FWHM) of around

70 cm–1. The broadening can be attributed both to

defect scattering and two- or more layer graphene

formation. The strong intensity of the D peak (in

comparison to the G peak) is a clear indication of

zone-boundary phonons scattering on defects, not

allowed in ideal graphene due to the large wave-

vector of the K-point phonon involved. Its presence is

an indication of disorder, such as finite size domains.

The D bands of N-doped graphene can be significantly

high with respect to PG because formation of pyridinic

and pyrrolic nitrogen sites is accompanied by defects

inside the graphene network and functional edges of

the graphene layer. The ratio of the integrated inten-

sities of D to G bands (ID/IG) can be used to estimate

the crystallite size (La) of graphene. We calculated La

of graphene by the following equation [34],

1

10 4 Da laser

G

nm 2, 4.10I

LI

where λlaser is the Raman excitation wavelength (λlaser =

532 nm). The values of ID/IG for PG and N-doped gra-

phene are 0.14 and 0.6, which correspond to crystallite

sizes of 137 nm and 32 nm, respectively.

Cross-sectional TEM samples were prepared using

a focused ion beam (FIB). Protective layers of amor-

phous carbon were deposited on the sample, initially

with the electron beam to avoid surface damage.

Images were obtained in STEM mode using a Cs-probe

aberration corrected JEOL 2200 FS STEM. Figure 2(b)

shows a cross-sectional TEM image for the N-doped

graphene on SiC(0001). The high resolution image

in Fig. 2(b) shows that Si3N4 was grown epitaxially on

SiC(0001) at this temperature. The TEM diffraction pat-

tern indicates and confirms that (112—

0)Si3N4//(0001)SiC,

namely the c-axis of Si3N4, is parallel to the SiC(0001)

surface. The line profile in Fig. 2(b) of the SiC planes

confirms that their spacing is indeed constant (2.43 Å

and 2.66 Å) and attests to the abruptness of the

formation of the Si3N4 and SiC interface under the

growth conditions reported here. No detectable sign

Figure 2 (a) Micro-Raman spectra of the D, G and 2D bands, for the N-doped (red) and pristine (black) graphene. (b) High-resolution STEM image obtained of the sample annealed under N2 (inset STEM image showing the interface layer).

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5 Nano Res.

of interdiffusion and intermixing has been observed.

We measured a distance of 4 Å between each graphene

layer; this value is very close to the bulk interplanar

distance (3.376 Å).

To investigate the elemental composition as well

as the chemical bonding environment of N atoms in

N-doped graphene, XPS measurements were carried

out for N-doped and pristine graphene, as shown in

Fig. 3. The experimental C 1s spectra and their fitting

curves are reported in Fig. 3(a) for the pristine and

N-doped graphene layers. Raw data for PG represented

by red disks are shown together with the results of

the curve fitting of the C 1s spectra on the top panel

(black curve). Given the asymmetrical shape of the

graphene component, three contributions have to be

considered to obtain consistent fits using a least-square

fitting procedure: The small component (IL) located

at 285.4 eV (FWHM = 1 eV), the stronger one (G) at

284.6 eV (FWHM = 0.53 eV), and the bulk component

(SiC) at 283.8 eV (FWHM = 0.95 eV). These three peaks

(IL, G and SiC) can be, respectively, ascribed to interface

layer, graphene layer and bulk contribution. No other

chemically shifted peak is needed to fit the spectrum.

Figure 3 (a) C 1s XPS spectra for epitaxial pristine graphene annealed at 1,350 °C at hν = 340 eV (top panel) and for N-doped graphene at hν = 340 eV (middle panel) and hν = 600 eV (bottom panel). The spectra were fitted using a Doniach–Sunjic line-shape analysis. (b) Si 2p XPS spectra for pristine and N-doped graphene at hν = 340 eV (surface sensitive) and hν = 600 eV (bulk sensitive). XPS measurements were performed at = 45° emerging angle with respect to the sample normal.

The XPS C 1s spectra for the N-doped graphene are

presented in Fig. 3(a) (middle and bottom panels).

Again, we found the two peaks at 283.2 eV (FWHM =

0.85 eV) and 284.6 eV (FWHM = 1.1 eV) which are

attributed to bulk Si–C bonds and C–C bonds

(graphene) and a new one at 286 eV (FWHM = 1.3 eV)

arising from the C–N bonds [35]. The main difference

for N-doped graphene compared to PG is the absence

of the interface layer peak at 285.4 eV. This observation,

which is consistent with our LEED results, confirms

the absence of a carbon-rich buffer layer. We notice an

important difference between the FWHM of graphene

peaks in PG and N-doped graphene: This can be

explained by a disorder induced on the graphene

layer due to the presence of nitrogen (seen as well in

the Si 2p core level spectrum). Also, the SiC bulk

component in the C 1s and Si 2p (Fig. 3) core level

spectra shifts 0.6 eV towards lower binding energy.

This reveals that the NG/SiC interfacial region is

modified for growth under N2 flux, because of a

different band bending [10]. It is worth noting the lack

of C–O components in the C 1s core-level spectrum.

This shows that the oxygen does not interact covalently

with the carbon layers in our material. The number

of layers present in each of our samples was also

deduced from our photoemission results. Assuming

homogenous two-dimensional graphene growth,

we can estimate the thickness of graphene using a

simple attenuation model. The estimated thickness

for our samples was 2 ML for the decoupled N-doped

graphene and monolayer for the PG.

In order to investigate the decoupling of the doped

graphene layer and the SiC substrate, we turn now to

the investigation of the interfacial region. Here the

Si 2p XPS spectra allow information about the com-

position of this region to be extracted. Figure 3(b)

shows the Si 2p spectra for N-doped (middle and

bottom panels) and pristine epitaxial graphene (top

panel). Both spectra consist of a dominant Si 2p

spin–orbit split doublet at 101.5 eV for PG and 100.9 eV

for N-doped graphene. The small shoulder at 102.1 eV

in pristine graphene is attributed to the interface

layer area of the SiC(0001) substrate. For the N-doped

graphene spectrum, we found additional Si 2p

components at higher binding energies: At 101.3 eV,

101.6 eV and 103.1 eV, due to Si–N–C units, Si3N4 and

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6 Nano Res.

Si–O bonds respectively. The peak located at 101.6 eV

is a first important signature of silicon nitride for-

mation [36]. At lower photon energy (340 eV, middle

panel) where the surface sensitivity is higher, this

signal from Si3N4 is strong, whereas it becomes weaker

at higher photon energy or less surface-sensitive con-

ditions (600 eV, bottom panel). This implies that the

graphene/SiC interfacial region can also be modified

under a N2 flux and results in silicon nitride formation

between the N-doped graphene layer and the SiC

substrate.

Based on the peak area ratio of the silicon nitride

component and the SiC component in the fitted Si 2p

spectra, we can estimate the thicknesses of the silicon

nitride passivation layers in our sample to be about

1.0 nm. The passivation of SiC suggests that higher

substrate temperature under N2 flux leads to a higher

nitridation rate when other variables (nitrogen partial

pressure and time of graphitization) are kept the

same. These values are close to the N 1s binding energy

of hexagonal silicon nitride (h-Si3N4). Furthermore,

the ratio of peak area for these N 1s component peaks

to the Si 2p component peaks at 101.6 eV were cal-

culated to be around 1.3 confirming that Si3N4 is the

main product in the passivation layer on the SiC

surface. Figure 4(a) shows the decomposition of the

N 1s spectrum for the nitrogen sample. The N 1s peak

can be resolved into five components centered at

397.7 eV (N1), 398.9 eV (N3), 398 eV (N2), 400.9 eV

(N4), and 401.4 eV (N5) representing Si–N, N–O,

pyridinic, pyrrolic, and graphitic type of N atoms

present in the sample, respectively. The differences

between N2, N4 and N5 are shown schematically in

Fig. 4(b), a top view of a graphene layer with nitrogen

defects. The graphitic-like configuration corresponds

to a direct substitution of a carbon atom by a nitrogen

atom, while the pyridinic and pyrrolic structures

correspond to a nitrogen atom bonded to two carbon

atoms in a six-atom ring and a five-atom ring

respectively. Our experimental results suggest that

nitrogen atoms present in the preparation chamber

interact with the Si atoms of the substrate to form a

Si3N4 interfacial layer, thus preventing the formation

of the carbon-rich buffer layer. For pristine graphene,

as expected, the buffer layer is present.

Figure 4 (a) High-resolution N 1s spectra of the N-doped graphene at hν = 480 eV (surface sensitive, top panel) and N 1s at hν = 600 eV (bulk sensitive, bottom panel). (b) Sketches of N-doped graphene with the four suggested doping sites, i.e., graphitic, pyrrolic and pyridinic nitrogen and Si–N bond formation.

The Si3N4 thin films are about 1 nm thick, suggesting

a good quality of the Si3N4 layer. It is found that the

interface between silicon nitride films and graphene

is sharp. By comparing the difference in spacing

between graphene and SiC, the thickness of the grown

graphene is estimated to be a bilayer, consistent with

the Raman spectroscopy results. SiC has a certain

roughness and contains charge traps which induce

Coulomb scattering centers and residual charge carrier

inhomogeneities in the graphene lattice, ultimately

limiting the graphene mobility. Si3N4 is a promising

substrate candidate for graphene since it offers several

benefits: Its interaction with graphene is weak and

the Si3N4/graphene interface is thermally stable up to

800 °C under ultrahigh vacuum (UHV) [37]. Because

of its high dielectric constant, large band gap (5 eV),

and excellent insulating and mechanical properties,

Si3N4 has been extensively used in current Si-based

technology as an insulator or chemical barrier. More

recently, it was predicted by first-principles calculations

that the plane of graphene remains flat, and the linear

dispersion relationship is well preserved, when

graphene is placed on β-Si3N4 substrate, and the

intrinsically high mobility of graphene is not affected

significantly by the weak interaction between Si3N4

and graphene [38]. In view of the fact that nitrogen

atoms play a critical role in the passivation of the SiC

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interface, as revealed by both LEED and XPS, it is

expected that a nitrogen flux may be used to assist

the growth of a high quality passivation layer on SiC

surfaces because of its lower kinetic energy and

higher activity. Besides the formation of these simple

bonds, our experiments suggest that much more

complicated bonds may also be created during such a

nitridation processes, making it difficult to control

the structural and electronic properties at the interface.

Recently, we have reported elsewhere that the growth

of N-doped graphene prepared by subliming Si from

SiC heated to high temperatures under N2 gas and

silicon flux [22]. In our earlier work, the presence of

silicon atoms hinders the nitridation of the SiC and

leads to the formation of the ( 6 6) R30° interface

layer.

Quasi free-standing N-doped graphene can be also

interesting as a metal-free catalyst because of its high

electron conductivity and high content of nitrogen.

N-doped carbon nanotubes (CNTs) from pyrolysis

of iron(II) phthalocyanine and chemical vapor de-

position (CVD)-derived N-graphene have been found

to facilitate the fuel cell cathode oxygen reduction

reaction (ORR) [39]. Recently, we have shown that

nitrogen atoms incorporated into a graphene lattice

contribute additional charge carriers and lead to a

modification of graphene’s electronic properties [22];

this may be useful in molecular sensing applications,

especially the graphene-enhanced Raman scattering

of organic molecules [35]. The changes in the chemical

and electronic properties of NG are expected to modify

the adsorption of acceptor molecules and molecular

oxygen via nitrogen atom defects (pyrrolic and

pyridinic) [40]. In order to understand whether the

incorporation of N atoms in the graphene structure is

sufficient to activate oxygen, we carried the surface

reaction of N-doped graphene and pristine graphene

layer by exposure to air for 7 days. XPS experiments

were performed under UHV in order to study the

oxidation of the two samples. In both cases the XPS

experiments were performed at room temperature, and

physically adsorbed species are desorb during surface

treatment (600 °C for 30 min) in the UHV chamber.

We demonstrate below that the reaction resulting from

exposure to air appears to be different for nitrogen-

doped and pristine graphene samples. Figure 5 shows

Figure 5 XPS overview spectra of the pristine graphene (black line) and the N-doped graphene (red line) after 7 days exposure to air.

XPS core orbital spectra of different elements of

N-doped graphene and PG samples exposed to air.

The binding energies of Si 2p, Si 2s, C 1s, N 1s, and

O 1s were observed at 100 eV, 150 eV, 285 eV, 400 eV,

and 532 eV, respectively. The N-doped graphene

shows an obvious N 1s peak whereas there are no

detectable N peaks in the spectrum of PG. The O 1s

peak is observed only in N-doped graphene samples,

which is possibly due to chemisorbed oxygen on the

graphene surface.

The results show that oxygen can be activated in

the presence of graphitic, pyrrolic and pyridinic N

species. For pyridinic and pyrrolic N located at the

graphene edge/defect, oxygen can be activated via

direct bonding with the lone pair electrons of N. For

graphitic N, because of the higher electronegativity

of N with respect to C, electrons transfer from the

adjacent C to N atoms and N back donates electrons

to adjacent C pz orbitals. The donation and back

donation processes not only facilitate O2 dissociation

on the adjacent C atoms but also help to form a

strong chemical bond between O and C. The higher

peak intensity ratio of O 1s to C 1s in N-doped

graphene shows the N-doped graphene surface has a

stronger oxygen adsorption ability, which may suggest

a potential application in fuel cells.

4 Conclusions

We have developed a general route for the direct

synthesis of N-doped bilayer graphene on Si3N4/SiC

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8 Nano Res.

substrates. This method is a new controllable transfer-

free route that opens the pathway for scalable bilayer

graphene growth. N-doped epitaxial graphene was

investigated for the first time using LEED/AFM

measurements. In addition to the nature of the quasi

free-standing N-doped graphene layer, the presence

of an underlying Si3N4 structure was also confirmed

by XPS. This might be an appealing way to achieve

heterostructure-based graphene/Si3N4/SiC for future

technology. Clear identification of the doped nitrogen

groups provides a solid basis for the rational modi-

fication of graphene properties. Graphitic, pyrrolic

and pyridinic type nitrogen, for example, renders

graphene sensitive to ambient conditions and opens

a way to oxygen sensing.

Acknowledgements

We are grateful to B. Etienne and M. Ridene for fruitful

discussions and N. Vélez for the 3D sketch. This work

was supported by the French Contracts ANR-2010-

MIGRAQUEL and ANR-2011-SUPERTRAMP, and the

RTRA Triangle de la Physique.

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