Electronic functionalisation of graphene via external doping and dosing

Electronic functionalisation of graphene via external doping and dosing Ursel Bangert1 and Recep Zan2 There exist many reports on functionalisation of graphene on a non-spatially resolved scale; this report concentrates on reviewing atomic-scale interactions of functionalising agents, i.e. on the electronic behaviour of single atoms, which are introduced as adatoms or lattice site impurities for the purpose of external doping and dosing to achieve bandgap engineering and electrical contacting of graphene; it also reviews the associated defects. Emphasis is put on visualisation of such interactions by advanced imaging in conjunction with localised spectroscopy techniques. Whereas the existing literature describing the development of such techniques in the application to graphene warrants a review in its own right, here the authors focus on observations, with modelling support, of the interaction phenomena themselves and not on the evaluation of measurements by such techniques. Atomic resolution transmission electron microscopy (TEM) combined with electron energy loss spectroscopy (EELS) in imaging and scanning mode, as well as scanning tunnelling microscopy (STM) are the most frequently applied techniques in aid of revealing topography and defect assisted interactions of graphene with foreign atoms and molecules. Electron-probe based investigations additionally lead to electron beam assisted interactions of foreign species with graphene. The graphene–metal interaction observed in a transmission electron microscope is a prevalent example of how reactions occurring between metals and graphene can be emphasised and thereby assessed: metal-mediated etching of graphene has proven to be a common phenomenon after metal dosing, e.g. to fabricate electrical contacts. The review reports furthermore on investigations revealing atomic position, bonding and dynamics of non-metal p- and n-dopants as well as on revealing the functionalisation of graphene via molecular self-assembly, intercalation and nano-sculpting. Literature till the end of 2013/begin of 2014 is reviewed. Keywords: Graphene, Doping, Adatoms, Metal contacting, Lattice-site impurities, (S)TEM, STM, EELS

Introduction Although the unique electronic structure of graphene results in extremely high carrier mobilities, ballistic transport, quantum Hall effect, high thermal conductivity and mechanical strength, as well as ambipolar behaviour in an electric field with charge carriers being tuneable to be either electrons or holes, the lack of a bandgap is a serious limitation for graphene’s use in electronics; the on/off ratios are three orders of magnitude below that of logic transistors. Thus bandgap ‘engineering’ to increase the on/off ratio is of paramount importance for the use of graphene in electronics. From the above said, doping of graphene is of great interest in

1

Department of Physics and Energy, University of Limerick, Limerick, Ireland Department of Physics, Faculty of Arts and Sciences, Nig˘de, University, 51000 Nig˘de, Turkey

2

*Corresponding author, email [email protected]

ß 2015 Institute of Materials, Minerals and Mining and ASM International Published by Maney for the Institute and ASM International Received 3 June 2014; accepted 17 October 2014 DOI 10.1179/1743280414Y.0000000047

order to shift the Fermi level and introduce a bandgap, not only to make graphene suitable for electronic applications, but also to provide functionality in devices, such as solidstate gas sensors, and even to open the possibility of single molecule detection.1 With regards to electronic applications bandgap opening has mainly been attempted and achieved by applying a gate voltage.2 This has resulted, however, in very small bandgaps, i.e. of several tens of millielectrovolt. Alternative approaches that are expected to lead to similar effects include the reduction of the device dimension; hence nano-sculpting/cutting of graphene to achieve nano-ribbons by lithography and by metalnanoparticle controlled catalytic etching in various atmospheres has been attempted (e.g. in hydrogen and ammonia; see for example Refs. 3–5). Other attempts include introduction of defects, doping by exposing graphene to gases (including hydrogen and fluorine), adatoms and molecules (external doping), applying strain or depositing graphene onto substrates such as h-BN or SiC. Most of the above techniques only open a gap of less

International Materials Reviews

2015

VOL

60

NO

3

133

Bangert and Zan

Electronic functionalisation of graphene

than 0?36 eV, and the most promising method, therefore, remains chemical modification or doping. A further objective of prime importance, which involves controlled incorporation or addition of adatoms on graphene, is the optimisation of electrical contacts: contact resistances to date are higher6 and carrier mobilities are more limited than expected. Graphene is expected to exhibit ballistic transport. Xia et al.7 found that this is approximately the case at low temperatures in gated graphene devices (e.g. for Pd contacts), where the only charge carrier scattering process is induced by the graphene–metal coupling, whereas at higher temperatures diffusive transport prevails because of changes in the scattering mean free path and the graphene–metal coupling length. Many respective papers describe effects of metals on graphene and resulting contacting issues from a ‘macroscopic’ point of view in that they deal with general solid state and electronic band structure aspects. There is, for example, a considerable number of papers dealing with surface enhanced Raman scattering (SERS) induced by a variety of metal deposits (often Ag and Au) in form of films and nano-crystals on graphene (e.g. Refs. 8 and 9). From the G-band shifting and the relative direction of the 2D-band shift in Raman spectra p- and n-doping are furthermore inferred, e.g. there is agreement that Ag deposition induces n-doping of graphene, whereas Au deposition induces p-doping, and that the substantial SERS and doping effects decrease with increasing number of graphene layers.10,11 However, the effects that metal atoms sitting in specific sites invoke on Raman spectra (e.g. whether they arise from charge transfer or strain) are still under dispute.12 So, relatively little information has so far been obtained of the metal–graphene interaction on the atomic-scale. Later in this review, the authors will address more local, atomic structure-related explanations of the effects introduced by metal deposition and contacting down to the single-atom level. More recently, doping has been attempted by molecular adsorbates.1,13 In first experiments to induce charge carriers in graphene via adsorption of various gases including NH3, H2O, and NO2,13 Hall effect measurements proved that NH3 induces electrons, whereas the latter two types of adsorbates result in holes as charge carriers. The above gases have been detected at remarkably low concentrations, and NO2 has even been detected in the extreme limit of single molecules.1 Wehling et al.14 tried to obtain a clearer picture of the physics behind the scattering mechanisms inferred by molecular absorbates of the NO2 system, and conducted a study in which the authors combined theoretical modelling based on density function theory (DFT), using the local density approximation (LDA) as well as the general gradient approximation (GGA), with electric field effect and Hall measurements. A general relation between the doping strength and whether adsorbates are open- or closed-shell systems was found, which showed that single, open shell NO2 molecules are strong acceptors, whereas the closed shell dimer N2O4 causes only weak p-doping. Experimental studies via scanning tunnelling microscopy (STM) and scanning tunnelling spectroscopy (STS) by Castellanos-Gomex et al.,15 following hydrogenation of the graphene surface with an Ar/H2 plasma, showed that this can open a bandgap of 0?4 eV. The authors introduced a method to statistically analyse thousands of

134


2015

VOL

60

NO

3

STS spectra, demonstrating furthermore that moderate annealing of the crystals is enough to close this bandgap and the samples can be hydrogenated again to yield a similar semiconducting behaviour. Further results of modifying graphene by adsorbates/ adatoms in order to achieve electronic doping are reported by Liu et al.16 The authors carried out O2 oxidation and infer creation of strong hole doping in single layer graphene on grounds of the observed Gpeak shifts in Raman spectra. Several authors have demonstrated doping (e.g. with N) of graphene directly during chemical vapour deposition (CVD) growth. The atomic configuration of the dopant atoms was primarily assessed by X-ray photoelectron spectroscopy (XPS) in these studies, revealing the fraction of dopant atoms by the emission peak intensity and the bonding type by its energy. C5H5N,17 NH3,18 N3H319 and solid precursors20 have been used as nitrogen sources. The authors of these studies report mainly pyridinic or pyrrolic N-sites,17,19,20 although in the case of N3H3 post-growth annealing at 500uC was found to convert the pyridinic nitrogen into graphitic nitrogen.19 One of the latest attempts to controllably incorporate functional atomic species directly into graphene is via low-energy ion implantation in a similar way to largescale integrated semiconductor technologies, in order to be able to precisely tailor dose and position of, e.g. electronic dopants.21,22 Whereas the authors in Ref. 21 have used low energy nitrogen ions to dope graphene on a Ni(111) surface, and show, using XPS, that N incorporates in both graphitic and pyridinic nitrogen positions, the authors of Ref. 22 carry out low-energy ion-implantation (at 20–30 eV) into suspended graphene and observe via atomically resolved imaging methods that the N-sites are mostly graphitic. In support, DFT calculations of low energy ion implants in graphene predict that at irradiation energies of 50 eV and below, the probability for substitutional incorporation of the ions is high. While the macroscopic properties of graphene devices are readily measured and characterised, understanding the interaction of impurities with the graphene lattice necessitates investigations and direct ‘visualisation’ at the atomic level. In the following, the authors concentrate on reviewing studies of effects of impurity atoms/ molecules/nano-clusters in/on grapheme, which employ an element of local, down to atomic-scale, imaging and spectroscopy in the (scanning) transmission electron microscope (S)TEM and the STM.

Characterisation methods There exists a number of comprehensive reviews on transmission electron microscopy (TEM) studies of carbon materials including those of graphene via high-resolution imaging and spectroscopy in aberration-corrected machines (e.g. Ref. 23). Such high resolution transmission electron microscopes (HRTEM) employ spherical and chromatic aberration correctors of their magnetic lenses24–26 ˚ at 80 keV and achieve image resolutions as high as 0?8 A ˚ and somewhat less at lower accelerating voltages (e.g. y1 A at 60 keV; the knock-on damage threshold in graphene is y80 keV), and are operated in scanning (STEM) and stationary (HREM) beam mode.

Bangert and Zan


1 Electron energy loss spectra acquired by accumulating 1 s exposures while scanning repeatedly the area around a Siatom for 200 s. The inset presents a 50 frame average from the stack of HAADF images created during the acquisition, showing (left) a trivalently and (right) a tetravalently bonded Si atom in the graphene lattice. The filled curve overlays of the Si-edge EEL spectrum are CASTEP calculations using the shown model insets. Images are taken in a Nion UltraSTEM at 60 keV with 31 pA beam current, 86–190 mrad HAADF detector angle and 35 mrad spectrometer semiangle (adapted from Ref. 35)

These instruments enable analysis, simultaneously and with single atom resolution and sensitivity, of local atomic configurations, their chemical identity and spectral response in monolayer graphene. HRTEM has played a vital role in atomic structure determination right from the time when first claims of successful extraction of monolayer graphene had been made in 2004: it provided direct evidence of the graphene lattice as well as of atomic defects in graphene layers,27 and, importantly, of the structure, stability and dynamics of graphene edges28,29 and of reconstructions of these edges beyond the commonly assumed zigzag and armchair configurations.30 By varying the accelerating voltage, HRTEM has furthermore given useful insight into defect dynamics under irradiation (Refs. 31–33, see next paragraph). Studies of atomic-scale elemental distributions in particular employ annular dark field (ADF) imaging and electron energy loss spectroscopy (EELS). The former technique reveals the chemical nature and site of atomic species by the fact that the signal strength of incoherently, elastically scattered electrons follows a yZ2 law (Z is the atomic number). These electrons are scattered into angles deviating significantly from the original beam direction, hence the name ‘high angle annular dark field (HAADF) imaging’ is given to this technique. Electron energy loss spectroscopy is an absorption spectroscopic technique, and in a STEM, spectra can be obtained from defined, sub-atomic areas. Carried out in combination these techniques provide measurements that are directly interpretable and give indisputable results: the most successful studies up to date of local bonding characteristics (visualisation of the density of states) of individual impurities have been achieved with Si-atoms incorporated in graphene.34,35 When carried out via sequential fast-scanning ADF imaging one is able to directly observe point defect diffusion within the graphene lattice with atoms resolved and identified via quantitative image analysis. Noise and distortions of atomic positions induced, e.g. by scanning, can be reduced by summing multiple ADF frames obtained of stationary defects. Electron energy-loss spectrum imaging (EEL SI) of single atoms allows the delocalisation of

inelastic scattering events to be quantified, and quantum mechanical calculations employing DFT are able to describe the delocalisation effect with good accuracy. These capabilities have opened new opportunities to probe the defect structure, defect dynamics, and even local optical properties in 2D materials with single atom sensitivity. However, there has been less focus so far on observations of the interaction of single foreign atom species with graphene. Similar to defects, atoms on graphene, as well as their movements and atomic-scale restructuring of the graphene lattice, on time scales .1 s (for image statistics purposes) can directly be monitored via the above mentioned advanced instruments and techniques. Figure 1 demonstrates that not only the atomic site of the Si-atom in graphene can be directly seen in HAADF STEM images, but it can also be verified via atomic resolution EELS that the atom is indeed Si, and what is more, DFT calculations reveal that the Si atom in the substitutional case sits slightly above the graphene sheet (side-view model in the left-hand panel in Fig. 1). Furthermore, the bonding state can be determined from the fine structure in the EEL spectra; the left-hand panel in Fig. 1 for trivalently bonded Si shows a significantly different onset of the Si–L absorption edge than the righthand panel for tetravalently bonded Si. The other most used technique for studying surface morphology as well as electronic structure on the atomic scale is STM: STM studies of graphene have been carried out on a number of substrates such as SiO2, SiC, Cu, Ir, Ru and others (e.g. Refs. 36–39), where the graphene layers were obtained by, e.g. exfoliation, CVD or graphitisation of SiC. More recently, STM has also been carried out on suspended graphene.40 Characterisation of such films has also typically exploited Z-imaging STS and low energy diffraction pattern analysis, all in situ and in ultra-high vacuum (UHV).

Electron-beam and topography assisted interactions in graphene Effects of accelerating voltage on defect dynamics Atomic interactions are occurring inadvertently on graphene during imaging in a TEM, although they


2015

VOL

60

NO

3

135

Bangert and Zan


˚ nominal thickness: a Before etching; 2 HAADF images of graphene etching in the presence of an aluminium layer of 2 A b after the start of the hole formation; c after hole enlargement in subsequent scans; d after continued etching as a result of a sustained supply of Al atoms to the hole’s edge (some Al atoms are indicated by red arrows in b–d); e after the etching process has almost stopped because the Al atom supply has ceased; f A lower magnification overview of the Al distribution and hole evolution. Black contrast in the images signifies holes, dark grey contrast graphene, lighter grey contrast hydro-carbon contamination on graphene and white contrast Al-clusters and -atoms. The scale bar is the same in a–e, 1 nm. Images were taken in a Nion UltraSTEM at 60 keV with a 86–190 mrad HAADF detector angle (taken from Ref. 45)

depend on the above mentioned beam energy displacement threshold. The latter is y80 keV in pristine graphene, but varies at defects and edges. Employing accelerating voltages around and below this threshold can therefore give insight into changes in the atomic structure of carbon nano-systems under electron irradiation. This concerns, e.g. electronic properties of defects associated with Stones–Wales transformations; these defects can be formed by single electron impacts and, remarkably, at electron energies below the threshold for atomic displacements through mechanisms of irradiation driven bond rotations.31 Kotakoski et al. furthermore find that regions with the defects withstand electron irradiation at ,100 keV, i.e. at higher energies than the knock-on threshold; rather than being perforated, graphene tends to amorphise. Warner et al.32 find that at accelerating voltages lower than the knock-on threshold electron beam irradiation induces variations in the distance between dislocations moving along a zigzag lattice direction by single bond rotation or through the loss of two carbon atoms. Chuvilin et al.33 report that indeed, the dominant effect of the electron irradiation at 80 kV in graphene is not atom removal, but atom rearrangement (e.g. the Stone–Wales-type bond rotation), which can lead to carbon ribbon or chain formation. Chains form efficiently by self-organisation during continuous removal of atoms from a graphene bridge. The authors point out that the high robustness under irradiation of these linear carbon chains combined with the ease of electron beam fabrication at the nanometers scale could provide a route to a synthesis of electronic device components with ultimate confinement.

136


2015

VOL

60

NO

3

It is worth noting that graphene is notoriously unclean, and that this, not least, restricts studies of large-area samples in the electron microscope even in UHV conditions. The authors have explored ways to clean graphene by leaving samples for prolonged times under UHV conditions and applying in situ electron-beam showering; however, even after short exposure to ambient conditions thereafter and re-insertion into UHV, contamination returns (Bangert et al., unpublished results). Algara-Siller et al.41 present a dry-cleaning process under ambient atmosphere using activated carbon, which results in mm-size clean areas. The effect of cleanliness on device applications of graphene remains a huge unknown and requires further investigation.

Interactions of graphene with foreign elements and molecules Electron beam assisted interactions are also being taken advantage of to study, e.g. the dynamics of light atoms and molecules42 and of nano-crystals on graphene under the ebeam,43 as well as interactions of the latter directly with graphene. Metal atoms, for example, have been found to exhibit catalytic activity in the etching of graphene.43–45 Figure 2 (from Ref. 45) shows the expansion of a hole in graphene through catalytic reaction with Al-atoms, moving along its edge. The authors of Refs. 43–45 suggest that in the presence of O-molecules (abundant in contaminants) metal atoms that are mobilised by the e-beam become oxidised and then reduced via CO or CO2 formation, whereby C-atoms leave the graphene lattice. In fact, because of its thinness and the resulting wealth of atomicscale observations graphene is increasingly being considered as the ultimate transparent support film for TEM

Bangert and Zan


3 a–c Atomic structures of reconstructed double vacancy defects in graphene as obtained from DFT calculations; d–f experimental TEM images of the same structures (courtesy of J. Meyer, J. Kotakoski et al., submitted for publication); a, d Double vacancy V2(5–8–5); b,e V2(555–777) transformed from the V2(5–8–5) defect by rotating a bond (marked in panel a); c,f V2(5555–6–7777) defect formed from V2(555–777) by another bond rotation (bond marked in pane b) (taken from Ref. 53)

samples. Along the above lines, Erni et al.46 use HREM to image small molecules trapped on graphene. On the basis of their experimental observation, following restoration of the exit-plane wave (for phase images the restorable image frequencies lie between 50 and 320 pm47) the authors conclude, in combination with DFT calculations, that hydrogen adatoms can trap molecules to specific sites on graphene that remain stable at room temperature. So adatoms and molecules mutually trap each other in specific energetically favourable sites. What likely occurs in the process of forming these configurations is that due to their higher adhesion energy, hydrogen atoms first attach to graphene and then trigger molecules to find stable positions on top of adjacent C-sites. Provided that there are ways to attach hydrogen to graphene in a controlled manner and amount, this mechanism could be employed to functionalise graphene with specific molecular groups. Furthermore corrugation-enhanced covalent chemical binding of graphene with hydrogen atoms has been predicted48,49 through controlled introduction of ripples (1-D features, by thermal cycling), whose orientations, wavelengths and amplitudes can be further engineered.50 Theoretical investigation of the chemical reactivity of different binding sites on graphene ripples with H2 and F2 molecules as a function of the amplitude–wavelength ratio (A/l) show no region-selectivity for fluorination (as the electro-negativity of F is dominant over the topography influence on the binding energy), but there are regional differences for hydrogenation; also, a y0?15 eV bandgap is opened. Although HREM observations show that single noble metal atoms show strong interaction with graphene edges, which might indicate the possibility of functionalisation of graphene nano-ribbons with metal atoms,51 the latter are mobile under the e-beam and presumably also under other conditions, e.g. electric fields and currents, requiring more investigations into a better control of graphene metallisation. The same applies to light-atom doping or attachments of organic molecules/groups: observations in many publications (e.g. Ref. 52) show that adatoms or groups thereof are extremely mobile on graphene and adhere

to adsorbates/contamination on graphene, or cluster together to form large atom-number patches.

Graphene defects and their interaction with foreign species Hence, the capability to create stable adatoms/molecules, or, more ideally, lattice integrated species, so as to gain control over electronic properties via the carrier type and density to electronically dope graphene, presents a bottleneck. One possibility would be to anchor such dopant atoms/molecules at defects. This section points out issues surrounding defect-trapping of foreign atoms. Defects in graphene have been studied in great detail, these concern mainly faulted rings [e.g. 55–77 (Stone– Wales), 555–777, 5555–6–7777] and vacancy defects.28,53–57 Figure 3, taken from Ref. 53, for example, shows reconstructed di-vacancy defects. It should be noted that most observed defects, e.g. mono- and di-vacancies, are created during imaging around the knock-on threshold (y80 keV), and that graphene itself is remarkably perfect (as witnessed in low voltage ,60 keV experiments) and has a preponderance to ‘heal’ itself.58 Wang et al.59 demonstrate that using defects is an efficient way to dope graphene: in a two-step process they create defects (i.e. vacancies) by high energy atom/ion bombardment and fill these vacancies with desired dopants. In this way different elements (Pt, Co, and In) have been successfully introduced in single-atom form. Figure 4 shows in situ observation (taken from Ref. 59) in an aberration-corrected and monochromated TEM operated at 60 keV, of Pt incorporated at defects in graphene. The high binding energy of the metal–vacancy complex ensures its stability (compare Fig. 5b and c, taken from Ref. 59). Wang et al. find using DFT calculations that the migration barrier for, e.g. Pt adatoms is sufficiently high to ensure that vacancy–Pt complexes are stable enough to survive prolonged electron beam irradiation. Later in this article, the authors show that this situation changes completely in the presence of O-ions (and disconcertingly the latter are always present because


2015

VOL

60

NO

3

137

Bangert and Zan


4 HRTEM images of a Pt atom: Trapped in a a bi-vacancy and b a trivacancy. c and d atomic models and e and f simulated HRTEM images for the Pt–vacancy complexes in a,b. g Binding energies for different configurations. For the monovacancy, the Pt atom resides out of the lattice plane in order to minimise the energy of lattice distortion because of atom size mismatch. h–j Video clips show the evolution of the Pt–vacancy complexes (indicated by arrows) under 60 keV electron irradiation. The electron beam current density is estimated to be 76106 e s21 nm22, i.e. 100 A cm22. The graphene sheet finally collapsed after an observation time span of y16 min. Scale bar: 1 nm (taken from Ref. 59)

of atmospheric and contamination conditions). Edgetrapped Pt atoms, on the other hand, were found to migrate easily along the edge. The occurrence of edgetrapped Pt atoms points to the possibility of edge-doping, which is critical for creating graphene nano-ribbon devices. The authors’ DFT results show furthermore that the Pt– vacancy complexes are non-magnetic for different vacancies, while all Co–Vn complexes are magnetic, and that graphene can be either p-doped with Pt and Co dopants or n-doped with In-dopants. Figure 5c and d shows the spin polarisation in the Co–V2 complex, reflected by a splitting into spin-up and spin-down bands. The local density of states at the Fermi level is mainly attributed to the localised states around the Co atom, as shown in Fig. 5f. The charge difference in Fig. 5h indicates that Co-integration results in a slight p-doping. Robertson et al.60 were able, by manipulating a focused electron beam in an AC-TEM, to control the formation of mono- and di-vacancies in graphene to subsequently trap single Fe atoms to form covalently bonded metal–defect complexes. These Fe-containing defect structures are stable in comparison to Fe atoms incorporated into the graphene edge but undergo limited e-beam mediated migration to adjacent lattice sites; this manoeuvre could be promising for selecting magnetic and non-magnetic states in doped graphene. It should be noted also that trapping of even larger foreign atom ensembles, namely Si6 clusters, in graphene nano-pores has been directly evidenced in atomic resolution HAADF-STEM observations, by Lee et al.61 Coming back to edge states – these represent defects that attract adatoms. Warner et al.62 probe the energy

138


2015

VOL

60

NO

3

states at graphene edges with single atom sensitivity and revealed that these edges easily and quickly attract surface atoms, leading to functionalisation by adatoms (e.g. C and N). Furthermore, graphene ribbons of the order of 1?5 nm width have low-energy edge states compared to a single graphene edge; hence restricting the size of graphene to a ribbon leads to different density of states compared to bulk graphene, but these states also quickly attract adsorbates and become functionalised. Rodrı´guez-Manzo et al.63 create lattice defects in carbon nanotubes (CNTs) and in graphene by focusing an electron beam in an STEM onto a 0?1 nm spot on places in pristine graphene at elevated temperatures, upon which metal atoms (Fe, Co, Mo) migrating on the graphenic surfaces, again at elevated temperatures, are trapped by these defects (e.g. at 400uC). Depending on the size of the defect, single metal atoms or clusters of several atoms can be localised in or on nanotubes or graphene layers. The aim is to place metal atoms with almost atomic precision, following e-beam patterning, in graphenic structures and to create a predefined pattern of foreign atoms in graphene or CNTs. Metal atoms can either be located on top of a coherent (non-defective) graphenic plane in a weakly bonded adatom configuration or form covalent bonds with carbon atoms at the edge of a graphene layer or at defective sites such as a lattice vacancy or a reconstructed area.64,65 While adatoms on a perfect surface are highly mobile, as proven experimentally as well as theoretically (see next section) metal atoms occupying defective sites in CNTs or graphene can be trapped in more or less stable positions.66 The fast migration of metal atoms on top of

Bangert and Zan


5 In situ observation of the Co atom trapping process: a HRTEM image decomposed from a video of a vacancy (a white spot indicated by the arrow) before trapping a Co atom; b A Co atom was trapped to form a Co–vacancy complex; c The complex was stable under electron irradiation. Scale bar: 1 nm. d Spin up and e spin down band structures and density of states of the Co–bivacancy complex. f Local density of state isosurface (0?1), g charge density isosurface (0?1), h charge difference isosurface (¡0?01 for red/blue colour) and i spin density isosurface (0?001e Bohr23) of the Co–bivacancy complex (taken from Ref. 59)

or within graphenic lattices has already been studied both experimentally67 and theoretically.66,68–71 Figure 6 from Ref. 71 depicts scenarios of Fe- and Cr- atoms on graphene: Fe-atoms sit on bridge sites between two Catoms, Cr atoms etch graphene, starting from bridge occupations (Fig. 6b), leaving holes behind after their removal (Fig. 6c), a process that progresses after repeated steps (Fig. 6d). Diffusion in the lattice via substitutional sites occurs at a much lower rate because the breaking of covalent bonds in the host lattice requires higher activation energies. The binding energies of metal atoms in vacancies in the graphene lattice66 and at graphene edges are too high to allow de-trapping, neither thermally nor by irradiation. The de-trapping of single metal atoms from defects that we observe points, however, to a lower than anticipated binding energy of the atom to the defect. It also appears that the trapping of single metal atoms occurs in graphenic systems with more than one layer. Intershell bonds between carbon atoms may exist and stabilise these defects against annealing.72 Even if no dangling bonds exist at such defects, the strongly distorted p-electron system may form bonds of sufficient strength to the metal atoms. However, the observation of de-trapping shows that strong covalent bonding to a single metal atom in a vacancy does not exist. Again, the role of the e-beam in other than simply knock-on effects, e.g. as mediator in electro-chemically induced reactions, because of the high mobility and large range of e-beam associated secondary electrons, has to be considered. The same effects could then be caused by currents flowing in graphene devices with metal contacts. This could have detrimental effects and requires further investigations. Similar negative effects with electrical contacts will arise from the high mobility of metal atoms on graphene; TEM

inspection after evaporation of metal contacts shows that all metal clusters sit in the contamination that is abundant on graphene.70,71 Cretu et al.73 confirm, via HREM imaging, attractive interaction between reconstructed defects in graphene and metal atoms, caused by an interplay of local strain around defects and electronic adsorption effects. The authors derive this from the observation that W atoms do not execute a random walk; instead there seems to be an attraction between a W atom and a trapping site over a surprisingly long distance. Possible trapping centres are single- and di-vacancies, 77–55 (Stone–Wales) and 777– 555 (3-heptagon-3-pentagon) defects; the trapping energies of the former would be too large (according to calculations) for W to undergo de-trapping again. However, jumps of W-atoms on graphene were observed. The combination with calculations confirming the energetics that could lead to the observed jump frequency and length suggests that strain fields created in C–C bonds in hexagons near 777–555 defects act as the responsible trapping centres. Following the above experimental studies (e.g. in Refs. 73 and 66), the attractive interaction between transition metal atoms and vacancy-type defects in graphene has been proven theoretically via first principle studies by Krasheinnikov et al.74: nanometre-range attractive interactions between p-vacancy type defects originate from an interplay of local strain in the atomic network, created by the defect, and electronic adsorption effects. Furthermore, analysis of the band structure of graphene with defects showed that some defects open up a semiconductor gap. Giovannetti et al.75 investigate, via DFT modelling, metal contacting of graphene by depositing graphene on various metal substrates. When graphene is chemisorbed


2015

VOL

60

NO

3

139

Bangert and Zan


6 a Middle panel: noise-reduced HAADF lattice image of two- and three-layer graphene with Fe impurities (the white line shows the approximate position of the sheet edge). The left- and right-hand panels are enlarged views of the red frames, with overlaid model structures (solid lines indicate the surface layer) to clarify the position of the Fe atoms; these sit on the surface on B sites. Shown in the insets are HAADF simulations of three layers (left) and two layers (right). The spectrum inset shows the Fe L2,3 absorption edge obtained on the single, arrowed atom. b Noise-filtered HAADF image of a Cr atom on monolayer graphene; the HAADF image simulation is shown in the inset. c Same area as in b after repeat scanning, revealing a divacancy, where the Cr atom had been. Model structures are overlaid to show the sites of the defects. d Raw HAADF image of monolayer graphene patch (dark grey) bordered by hydrocarbons (lighter grey). Cr atoms sit on the hydrocarbon contamination; a chain of Cr atoms (arrowed; top) can be seen moving from an area of Cr clusters (white patch in the top right corner: the image is overexposed here because of the high Cr concentration) towards the edge of the top hole (black area) and to decorate the edge of the bottom hole (arrowed, bottom). Images are taken in a Nion VG STEM and a Nion UltraSTEM at 80 and 60 keV, respectively, with y80–190 mrad HAADF detector angles (taken from Ref. 71)

(on Co, Ni, and Pd) the graphene bands are strongly perturbed and acquire a mixed graphene–metal character. However, when the bonding is weak (on Al, Ag, Cu, Au, and Pt), the unique electronic structure of graphene is preserved, leading to shifts of the Fermi level with respect to the Dirac point of y0?5 eV, resulting in doping (n-type on Al, Ag, and Cu, and ptype on Au and Pt). This is not influenced by the crossover from p-type to n-type and does also not conflict with the metal work function, which is y5?4 eV, and ˚ , has a value at equilibrium separations of y3?3 A much larger than the graphene work function of 4?5 eV. Ugeda et al.76 observe, in a study, combining STM with DFT calculations that even in weakly coupled graphene–metal systems (i.e. graphene deposited onto Pt) the electronic properties of existing Cvacancies in the graphene layer strongly differ from the ones found for C-vacancies in a well-decoupled graphene layer, due to the increased interaction of the latter with the metallic surface, destroying magnetic

140


2015

VOL

60

NO

3

moments that these would introduce in decoupled, free-standing graphene. In a further STM study employing differential conductance, dI/dV, measurements with the STM tip held over individual Co adatoms, Brar et al.77 measure the energy-dependent local density of states (LDOS) of these adatoms. They find that their electronic structure can be tuned by application of the device gate voltage, and that the Co atoms can be reversibly ionised. They observe the formation of large screening clouds around ionised Co adatoms as well as around intrinsic graphene defects.

Interaction of graphene with metal adatoms For metals deposited onto perfect graphene, McCreary et al.78 report, based on in situ transport measurements, that n-type doping behaviour is inferred by transition metal clusters of Ti, Fe, and Pt against expectations that would arise from the higher work function than that of

Bangert and Zan

graphene. This provides experimental evidence for a strong interfacial dipole favouring n-type doping as predicted theoretically.75 Ren et al.79 register a shift in the Dirac point into the conduction band in graphene (via FET measurements in UHV) with Ag and Cu, resulting in effective n-doping, while Au causes a shift into the valence band; here the authors rationalise this as a result of the difference in work function values between each metal and graphene. Hence, doping behaviour appears to vary with the metal species as do suggestions to justify this behaviour. More research needs to be done to arrive at a more unified explanation. The effect of metallic doping for the case of an entire metal monolayer was investigated, using DFT calculations, by Uchoa et al.80 The authors emphasise that this is conceptually different from the adsorption of isolated adatoms in graphene, where the metallic orbitals strongly localise around the impurities. Their calculations for the case of K suggest that the Ks electrons delocalise, i.e. the metallic band of K hybridises with the carbon near-free electron band, and that charge transfer of 0?51e per K atom can take place, which diminishes, however, with increase in the effective charging energy (e.g. depending on the substrate). For a monolayer of 4d transition metals such as Pd, hybridisation of the graphene pz orbital with the localised d orbitals at finite effective charging energy can produce strong itinerant magnetism. Liu et al.81 compare the behaviour of metals on graphene obtained as epitaxial layer by annealing of SiC to results of a comprehensive theoretical study (although the calculations were done with suspended graphene). Using firstprinciples calculations, the authors investigate metals ranging from group I to IV metals, 3d and group 10 transition metals, to noble metals and rare earth metals. They found that the ratios of the adsorption and the bulk cohesive energy, as well as the diffusion barriers control the growth morphologies of the latter classes of metals (from the transition metals onwards) via the bond order. This trend can be attributed to various degrees of covalent bonding mainly from the d electrons of the metal adatoms. However, the correlation between the adsorption strength and the value of the bond order is not obvious for group I– IV metals on graphene. In the case of group I alkali metal adatoms on graphene, the interaction between the alkali adatom and graphene is mainly ionic. For group II metals on graphene, there exist various degrees of ionic bonding and physisorption. Similarly, the group III and IV metals on graphene exhibit a mixture of ionic and covalent bonds. Moreover, long range interactions such as dipole–dipole, elastic and indirect electron and magnetic interactions induced by the metal adsorption can also play an important role in the growth morphology. The calculations predict furthermore that 3d and group 10 transition metals, noble metals, and some rare earth metal nano-structures on graphene are thermally very stable and exhibit high stability against coarsening. The authors further claim that their calculation predictions are consistent with available experimental results, although most of those where obtained from STM on graphene on substrates (SiC). Hardcastle et al.69 have used the plane-wave density functional theory code CASTEP (with Tkatchenko– Scheffler van der Waals corrections and generalised gradient approximation exchange correlation functionals by Perdew, Burke and Ernzerhof) in combination with atomic resolution STEM HAADF imaging70,71 to


study metal-atom behaviour on suspended graphene. They find ‘binding energies’ (defined as the difference in enthalpy of the composite system supercell and that of the sum of the isolated metal and geometry-optimised graphene supercells) of a fraction of an electron volts for Cr (y0?5 eV) and Au (y0?3 eV) and of y1 eV for Al while Liu et al.81 find adsorption energies, defined as the difference between the energy of the relaxed adatom– graphene system and that of the isolated perfect graphene sheet and an isolated adatom of ,0?1 eV for, e.g. Cr and Au, and of y1 eV for Al. Although presenting somewhat different values (and the authors of Ref. 81 investigating a larger range of metals with, interestingly, a quite different behaviour for 3d metals, as mentioned above, because of bond order and d-shell filling state), the authors of both Refs. 81 and 69 predict the same trend in mobilities, because of low diffusion barriers for noble metal atoms, on the order of thermal energies (around 0?01 eV), higher for, e.g. Al (0?1 eV), and much higher for rare earths (few electron volts in Ref. 81). A note of caution has to be given here with regards to the calculation and modelling results: reality is likely to defy these results and thus the prospects of controlled functionalisation via metals. The contamination on graphene (1) prevents the formation of continuous thin metal layers (as assumed for the calculations in Ref. 80) and (2) will interfere with the dynamics of the metal atoms as predicted in various calculations. This has been pointed out and proven in observations by the authors in Refs. 70 and 71. It should furthermore be mentioned from an experimental point of view that attempts have been made to stabilise highly mobile metal atoms (Au, Pt) on graphene having undergone hydrogenation70 or nitrogen doping,82 to provide anchoring points. This resulted in a finer cluster dispersion in the case of Au, all clusters, however, are still sitting in the contamination, and in a dispersion of single atoms and atomic clusters in the case of Pt, sustained primarily at graphene edges with few atoms and clusters stable on terrace sites.

Metal-atom mediated etching of graphene As already indicated in the previous sections, etching of graphene can be catalysed by metal atoms and considerable effort has gone into investigations of these etching effects in graphene. This is of interest not only for a general understanding of the metal–graphene interaction with respect to consequences for electrical contacting, but also for nano-sculpting and, thereby, functionalisation. Wang et al.83 etch graphene loaded with Pt particles at 1000uC in a hydrogen atmosphere. This results in the formation of nano-structured defect sites, including trenches, nano-ribbons, islands, and holes, which, in turn, lead to an increase in the number of unsaturated carbon atoms and, consequently, enhance the interaction of CO2 molecules presumably with the graphene edges, thus resulting in a high capacity for storing CO2. Ramasse et al.44 show via atomic resolution high angle ADF imaging in combination with EELS, carried out in an aberration-corrected scanning transmission electron microscope on suspended, single-layer graphene, onto which the metals Cr, Ti, Pd, Ni, Al, and Au atoms had been deposited that nanoscale holes were etched into


2015

VOL

60

NO

3

141

Bangert and Zan


7 Scanning tunnelling microscopy (STM) imaging of nitrogen dopants: a STM image of the most common doping form observed on N-doped graphene on copper foil, corresponding to a single graphitic N dopant. (Inset) Line profile across the dopant shows atomic corrugation and apparent height of the dopant (Vbias50?8 V, Iset50?8 nA); b Simulated STM image of graphitic N dopant (Vbias50?5 V), based on DFT calculations. Also superposed is a ball-and-stick model of the graphene lattice with a single N impurity; c STM image of N-doped graphene on copper foil showing 14 graphitic dopants and strong intervalley scattering tails. (Inset) FFT of topography shows atomic peaks (outer hexagon) and intervalley scattering peaks (inner hexagon, indicated by red arrow) (Vbias50?8 V, Iset50?8 nA); d Spatial distribution of N–N distances from eight samples on copper foils with different N concentrations. The distributions are all fit well by a quadratic power law (expected error bands in grey) over all length scales indicating that N dopants incorporate randomly into the graphene lattice; e dI/dV curves taken on an N atom (bottom) and on the bright topographic features near the nitrogen atom on N-doped graphene on copper, offset vertically for clarity. The top curve is the dI/dV spectrum taken y2 nm away from the dopant. (Inset) Positions where the spectra were taken (Vbias50?8 V, Iset51?0 nA) (taken from Ref. 18)

graphene. These initiated at sites where single atoms of all the metal species except for gold come into close contact with the graphene (initially metals accumulate in the abundant contamination on the graphene surface). The ebeam scanning process is instrumental in promoting metal atoms from clusters, formed during the original metal deposition process, onto the clean graphene surface, where they catalyse the hole-forming process. These observations concur with literature calculations by Boukhvalov and Katsnelson,84 predicting a much lowered vacancy formation in graphene when metal adatoms are present. The authors in Ref. 44 also address the requirement and importance of oxygen atoms in this process, although not predicted by such calculations in Ref. 84. Si-impurity atoms, especially when present in SiOx molecules, dramatically decrease the vacancy formation energy. This observation is supported by calculations.44 It is also suggested that this metal-mediated etching of graphene in a STEM in UHV at 60 kV could be exploited in controlled nano-manipulation and selfassembly processes for future graphene-based devices.

Non-metal p- and n-dopants in graphene The authors will now focus on the effects of single, inlattice group III and V substituents as acceptors and donors. Doping graphene and CNTs with nitrogen has been shown to result in n-type behaviour.85,86 Not all nitrogen defects, however, change the electronic structure in the same way. For example, while substitutional nitrogen defects in single-walled CNTs result in n-type doping, pyridinic nitrogen defects have been predicted to

142


2015

VOL

60

NO

3

produce p-type doping.87 Synthesis conditions and resulting defects can affect the way in which nitrogen is incorporated into CVD graphene, hence exact knowledge of nitrogen atom sites is vital in understanding and optimising the effect of the doping. Zhao et al.18 used STM, Raman spectroscopy, X-ray spectroscopy and first principles calculations to characterise individual nitrogen dopants in CVD-grown monolayer graphene on a copper substrate. They found that individual nitrogen atoms were incorporated as graphitic dopants, and that a fraction of the extra electron on each nitrogen atom was delocalised into the graphene lattice. The electronic structure of nitrogen-doped graphene was strongly modified, but only within a few lattice spacings of the site of the nitrogen dopant. Zhao et al. furthermore observe spatial random occupation with N-atoms in the vicinity of each other occupying the same sub-lattice with N-concentration between 0?23 and 0?35%. This fraction together with the observed Dirac point shift infers that each graphitic N dopant contributes (on average) y0?42¡0?07 mobile carriers to the graphene lattice. Localised charge occurs at ˚ away the N-centres with electronic perturbations up to 7 A from the dopant (Fig. 7). Charge distribution changes on the single atom level arising from substitutional nitrogen atoms in graphene have also been observed via HRTEM88 through contrast changes, which, in turn, are primarily due to a change in the electronic configuration on the neighbouring carbons rather than on the nitrogen atom itself. Lv et al.89 investigate CVD grown graphene, N-doped by addition of NH3. Raman indicates a shift of the Fermi level upwards by y350 meV; STM dI/dV measurements

Bangert and Zan


8 Calculated formation energy, experimental and simulated scanning tunnelling microscopy (STM) images of different Ndoping configurations: a Formation energies of different N-doping configurations in N-graphene sheets (as illustrated in insets) computed using ab initio calculations. b–e Simulated STM images depicting two different atomic configurations for double substitution of nitrogen dopants b–c and for two pyridine-like N-dopants d–e. The biases are as follows: b 21?0; c 21?0; d 20?7; e 20?7 eV. The carbon and nitrogen atoms are illustrated using grey and cyan balls, respectively. The superscript B9 is used to differentiate between two N atoms as being first-nearest neighbours (N2AB) or third-nearest neighbours (N2AB9) (taken from Ref. 89)

at the doping site confirm the presence of a high-intensity peak located at ca. 0?2 eV, and DFT calculations performed close to the scanning points reveal a projected density of states at 0?4 eV above the charge neutrality point, which is caused by the nitrogen pz orbital. The presence of N was further confirmed by XPS, showing substitutional and pyridine-like N-incorporation. Scanning tunnelling microscopy revealed additional defects consisting of two N-atoms substitutionally integrated beside a C-atom, hence, sitting on the same sublattice, and including various numbers of vacancies (Fig. 8).

The authors point out that the doping imbalance between the two sublattices could have a strong impact on the electronic transport properties. Furthermore, a novel and outstanding Raman enhancement of RhB molecules was demonstrated when using N-doped graphene sheets as a substrate. These findings support prospects of using chemical doping with group III and V elements, such as N, B or P, to tailor the electronic and chemical properties of graphene that might lead to selective and efficient molecular sensors. Similarly, Zhao et al.90 observe by STM and X-ray both, B and N, introduced into graphene

9 a Scanning tunnelling microscopy image of 30630 nm area exhibiting multiple defect forms. Red and green triangles indicate the graphitic B dopants in different sublattices. Blue arrows indicate complicated defect forms; b Scanning tunnelling microscopy images of different defect forms associated with pentagon–heptagon pairs. [(19) and (29)] Proposed structures for features in inset (1) and inset (2), respectively. Vbias520?5 V, Iset50?5 nA; c dI/dV spectra taken across the features in inset (2) of c with the red spectrum taken at the centre of the feature. Vbias520?5 V and Iset50?3 nA; d dI/dV spectra taken across the feature in inset (6) of c with red spectrum taken at the centre of the feature. Vbias520?5 V and Iset50?3 nA (taken from Ref. 90)


2015

VOL

60

NO

3

143

Bangert and Zan


10 a Processed medium angle annular dark field image of a dopant atom in graphene; b Carbon K-edges obtained from the region (outlined by a box) surrounding the dopant atom shown in a and from a region of pristine graphene. (The peaks in the spectra at y297 and 302 eV are artefacts and are likely to be due to bad pixels on the CCD camera); c Experimental and modelled EELS spectra from pristine and defect regions. Images are taken in a Nion UltraSTEM at 60 keV and 40–195 mrad HAADF detector angle (taken from Ref. 91)

144


2015

VOL

60

NO

3

grown on Cu via CVD. The authors found that B, like N, incorporates into the carbon lattice primarily in the graphitic form. Density functional theory calculations indicate that boron dopants interact strongly with the underlying copper substrate while nitrogen dopants do not. The local bonding differences between graphitic boron and nitrogen dopants lead to large scale differences in the dopant distribution. The distribution of dopants is observed to be completely random in the case of boron, while nitrogen displays strong sublattice clustering, as seen in Ref. 89. Furthermore, nitrogen-doped graphene is relatively defectfree while boron-doped graphene films show a large number of Stone–Wales defects (Fig. 9). These defects do, however, not electronically dope the graphene film. Individual substitutional nitrogen dopant atoms in graphene were also visualised in HAADF images obtained by Nicholls et al.91 in a STEM in combination with EELS and modelling. Electron energy loss spectroscopy revealed that there is a change in the carbon K-edge in the region surrounding the nitrogen dopant atom compared to pristine graphene (Fig. 10). Modelling shows that the change in the carbon K-edge near to the nitrogen defect is due to the influence of the nitrogen atom on the electronic structure of the neighbouring carbon atoms, and the extra feature can be attributed to the C–N bonds. Whereas the above authors could not show the existence of nitrogen via the N-edge in EEL spectra directly, a recent study by Bangert et al.22 reveals single Nas well as B-atoms introduced via low energy ion implantation into graphene by atomic resolution HAADF imaging as well as directly by energy loss spectrum imaging using the N- and the B-K EEL absorption edge, respectively (see Figs. 11 and 12). N incorporates predominantly substitutionally at implantation energies of 25 eV (Fig. 11a and b) with a small fraction of N-atoms captured beside vacancies (Fig. 11d). Zhao et al.21 have observed, although via non-spatially resolved methods (XPS), that the pyridinic N-fraction increases, when increasing the implantation energy from 25 to 150 eV. This suggests that higher energy ion implantation results in an increasing number of defects in the graphene. Before the experiments by Bangert et al.,22 the possibility of atomic N- and B-doping of graphene via low energy ion-implantation had been theoretically explored by ˚ hlgren et al.,92 who combined classical molecular A dynamics simulations and density-functional-theory total-energy calculations. Their simulations show that the optimum irradiation energy is 50 eV with substitution probabilities of 55% for N and 40% for B. The authors further estimate probabilities for different defect configurations to appear under B and N ion irradiation. This is in good agreement with the experimental results in Ref. 22 where N-ion implantation doping at 20–30 eV ion energy achieved a retention of .15% of substitutional N within the pristine graphene lattice. Hence, successful n-type doping of graphene with nitrogen can be achieved, not only via chemical methods, but directly via non-chemical processes used in commercial semiconductor technologies with prospects for largescale controlled and flexible doping of graphene and, generally 2D materials, providing a method compatible with those in IC technologies. Importantly, the authors in Ref. 22 show that this is also possible with implantation of B-ions, providing potential substitutional p-type doping (Fig. 12), which so far has been questionable.

Bangert and Zan


11 Atomic resolution HAADF image of graphene implanted at 25 eV with N to a dose of 61014 cm22: a Raw data and; b data obtained by deconvolution of image a as described in the text. The patches in the top left and bottom right corners are ubiquitous contamination; N-atoms (here all substitutional) shown as brighter contrast spots (encircled) in a and as orange spots in the false colour image in b, where colours are assigned on a temperature scale with increasing intensity represented by the colour sequence ranging from blue (thin regions or low atomic number) over orange to yellow/white (thick regions or high atomic number); c intensity histogram of C- and N-atoms, the latter having intensity values .2s larger than C-atoms (set at the value of 1), the inset shows an extended view of the distribution; d pristine graphene patch (deconvolved image as in b) with an N-atom beside a vacancy and a substitutional N-atom. Model sketches are shown below. An intensity trace along the dotted line is overlaid on image d, showing the expected 1?46 enhancement at the N-atom. Images are taken in a Nion UltraSTEM at 60 keV with a 86–190 mrad HAADF detector angle (taken from Ref. 22)

This is in contrast to the above papers,89,90,18 where dopants were introduced chemically during the CVD process, resulting in associated defect sites. Ion implantation, in contrast to Ref. 90 showed, for example, no defects in the case of B-doping.

Further functionalisation of graphene – molecular self-assembly, intercalation and nano-sculpting Extending out from reviewing the effects of dopants/ adatoms on graphene, it is worth mentioning that research is increasingly focusing on forming assemblies of inorganic clusters and super-molecules for extended, hierarchical self-assembled nano-structures on epitaxial graphene superstructures. These arise due to a commensurate match between molecular dimensions and the moire´ periodicity of the respective graphene/layeredsubstrate superstructure. Using moire´s (often used such superstructures are graphene/Ir or graphene/Ru) as active templates enables growth of two-dimensional cluster superlattices through a universal cluster-binding mechanism resulting in high superlattice order and extent, thermal stability and tunability of the cluster size. The creation of various metal-cluster superlattices was demonstrated, e.g. by Pan et al.93 (Pt) and N’Diaye et al.94 (Ir, Pt, W, Re). Preferred metal cluster nucleation sites within the moire´ supercell are fcc and hcp sites. For, e.g. Ru- and Pt-clusters on a graphene/Ru substrate

these were reported to be fcc sites93,94 whereas for Ir, Pt, W and Re on a graphene/Ir substrate these are hcp sites36 (see Fig. 13). The sizes of the deposited metal clusters and the distribution patterns appear to further depend on the metal–graphene lattice match and the metal–carbon (M–C) bond strength, the bond strength of the metal-cluster and the metal-substrate atoms (i.e. on either side of the graphene layer), as well as metal cohesive energies within the cluster. There seem to be competitive effects: when the M–C bond dissociation energy is higher than the cohesive energy within the metal cluster (in the case of Ir), and also when the interaction of the cluster with the underlying substrate is weak (for Ru), 2D cluster are seeded, whereas when the cohesive energy dominates (at higher coverages) 3D clusters tend to form (Ref. 39); for some metals they tend to retain the assembly regularity, for others the assembly loses long-range order. The latter situation is also reported to occur at increased coverage, e.g. in the case of Ni on graphene/Ru. Here, it was also seen that RT Ni-deposition (as opposed to deposition at 150 K) leads to a dramatic size increase in the clusters, which are, however, well-matched in terms of shape and orientation, to the moire´ registry.95 It is agreed, however, that the mechanisms at play in the various cases are not fully understood yet. Pollard et al.96 report the trapping of perylene tetracarboxylic diimide molecules in energy minima associated with the moire´ pattern produced in the super structure (see Fig. 14). Such assemblies are also interesting in relation to graphene electronics, as it


2015

VOL

60

NO

3

145

Bangert and Zan


12 a Double Gaussian filtered atomic resolution HAADF image of a substitutional B-atom in graphene (for process details see ‘‘Methods’’ section in Ref. 22); b raw HAADF signal obtained from the framed area in a simultaneously with the EEL spectrum image in c. The B-atom can be seen to have a lower intensity than the C-atoms, according to the Z 2 intensity relationship in HAADF images; c EEL spectrum image intensity in an energy window (190–220 eV) around the B–K edge, representing an atomic boron map (‘temperature’ colour scale: blue/green, low intensity; yellow/white, high intensity); d sum of spectra extracted from pixels of the spectrum image represented in part c of the area around and including the B-atom. Images are taken in a Nion UltraSTEM at 60 keV with a 86–190 mrad HAADF detector angle (taken from Ref. 22)

has been demonstrated that molecules can act as molecular dopants97 and also that superstructures resulting from moire´ patterns of graphene grown on Ir(111) give rise to the formation of a bandgap.98

There exists a considerable number of papers dealing with consequences of intercalation of metals with layered graphene structures or graphene-heterostructure systems (i.e. metal insertion between two layers of graphene or between graphene and a non-metallic substrate) for electronic doping effects and bandgap tuning, the authors of, e.g. Ref. 99 interpret STM results obtained of the graphene/SiC system intercalated with a mono-layer of gold as indicative of hole-doping. Sheu and Yang100 describe bandgap opening and Dirac cone and Fermi level shifts depending on Au-coverage in a theoretical study of Au-intercalated bi-layer graphene. More in-depth reporting of these interesting modifications of graphene structures, however, exceeds the scope of this article. Finally, there is increasing literature on nano-sculpting of graphene for the purpose of electronic and chemical functionalisation (e.g. Ref. 101) including manipulation of graphene by the probe of the imaging/ analysing instrument (e.g. electron beam, STM tip) itself, followed by imaging,62,102 however, including results of such activities, again, is beyond the scope of this article, and further information can be found in Refs. 62, 83, and 101 and references therein.

Conclusions 13 a Atomic resolution scanning tunnelling microscopy (STM) topograph of graphene on Ir(111). The rhombic moire´ unit cell is indicated by lines. Tunnelling voltage applied to tip Ut5z0?2 V; tunnelling current It523 nA; b Scanning tunnelling microscopy topograph after deposition of 0?02 ML Ir on graphene at 350 K; Ut5z0?2 V; It58 nA; c Schematic illustration of the density function theory (DFT) optimised C(10610)/ Ir(969) unit cell. Shading of the C atoms corresponds to their heights as calculated by DFT. First, second, and third layer Ir atoms are coloured cyan, red, and green. Hcp-type region, full circle; fcc-type region, short-dashed circle; atop-type region, dashed circle segments (taken from Ref. 36)

146


2015

VOL

60

NO

3

This review is aimed at giving an overview on studies with high localisation of the effects of impurity atoms and molecules and few-atom clusters adsorbed on or incorporated into graphene. In the respective reviewed literature, sites of foreign atoms were imaged down to the atomicscale via (S)TEM and STM and bonding as well as electronic doping effects investigated by correlation with highly spatially resolved spectroscopies – EELS in the (S)TEM and current–voltage spectroscopies in the STM. Attachment of foreign atoms occurs favourably at vacancy defects in graphene; these maybe involuntarily or voluntarily introduced by the imaging or by external charged particle probes. Foreign atoms, e.g. metals, exhibit also significantly increased binding probability at graphene edges. Furthermore, topographic effects (ripples) as well as

Bangert and Zan


14 Scanning tunnelling microscopy images: Acquired following deposition of a DP-PTCDI and b DB-CTCDI on a graphene monolayer formed on Rh(111)/YSZ Si(111). c, d Diagram of junctions between DP-PTCDI dimers c and trimers d stabilised, respectively, by two and three C5O…NH hydrogen bonds between neighbouring molecules with dimer centre–centre spacing of d and trimer vertex to molecule centre spacing r. e Placement of DP-PTCDI trimers and dimers. f DB-CTCDI trimer junction analogous to d with vertex to molecule centre spacing r and placement of DBCTCDI trimer on the graphene superstructure. g, h STM images of DB-CTCDI showing chirality of junctions and intramolecular detail of molecules. The hexagons in h highlight the chirality of the molecular arrangement. Imaging parameters: a 21?0 V, 100 pA; b 1?0 V, 100 pA; g 1?5 V, 50 pA; h 21?5 V, 50 pA (taken from Ref. 96)

hydrogenation and fluorination facilitate attachment of atoms and molecules. A general observation is that foreign species are, however, mobile on graphene; it is not clear whether the imaging probe plays a role in this. A large number of DFT calculations have been carried out of the energetics involving different binding sites; the outcome of the various studies is somewhat diverse, a number of articles agree, however, in that binding energies at many defect types should be sufficient to keep the foreign species anchored at room temperature. Emphasis in this review is put on metal atom capture and anchoring in the graphene lattice itself and as adatoms on the surface; these issues are very important for contacting graphene and use in plasmonics. An emerging issue in a number of studies is metal mediated etching of graphene. Again the role of the imaging probe in the hole-formation process is not entirely clear, and neither is that of the metal in the catalytic reaction. Despite impressive time sequence imaging following the movement/defect reconstruction involving individual metal atoms there is strong evidence that oxygen atoms have to be involved, too. Claims that the metal atoms at graphene vacancy defects and edges, as shown in atomically resolved TEM images, are alone responsible for the carbon atom dissociation and re-arrangement in these

studies should be taken with caution. Oxygen has been detected, e.g. by the authors of this article, in the presence of metal atoms in studies including atomically resolved EELS; HREM phase contrast images alone do not easily reveal the presence of oxygen. The incorporation of group III and V elements for nand p-doping is another focus of this review. The most progress so far has been achieved not only with Ndoping, chemically into graphene grown on substrates but also recently via low-energy ion implantation into freely suspended graphene, resulting in substitutional (graphitic) N as well as N-incorporation at defect sites (pyridinic, pyrrolic), and leading to different electronic behaviours, but predominantly to n-type doping. First success with B as p-type dopant has also been reported, again in films on substrates as well as in freestanding films. Substrate influences seem to generally affect the local bonding and the distribution of dopants. An important emerging perspective in the reviewed literature is that there appear to be ceaseless possibilities of atomic-scale, controlled functionalisation of graphene when carried out with nano-probes and under nanomanipulation, e.g. in atomic-resolution scanning microscopes; whether such manipulations can be transferred


2015

VOL

60

NO

3

147

Bangert and Zan


to and integrated into large-scale technologies will be the next big question and challenge to tackle. Other 2D materials, especially 2D forms obtained from transition metal dichalcogenides103–105 have sparked wide interest in both, device physics and technological applications at the atomic monolayer limit. In contrast to graphene, these materials have a bandgap. And as with graphene doping and dosing are of primary importance for directed and controlled functionalisation. These issues warrant a review article in its own right; the number of publications on 2D dichalcogenides is rapidly expanding and will soon reach the same extent as with graphene. The authors conclude this review by referring to a few examples in the current research focus on MoS2, and WSe2. These materials are under investigation for use as single-layer transistors with excellent electrical performance,106–109 namely, high on/off current ratios and suggested high carrier mobilities. Transistors from 2-D materials are expected to overcome the severe short channel effects, which limit the performance and operation of future (sub-5 nm gate length) devices. They are also promising for LEDs. Furthermore, 2D monolayers can be stacked together with precise control to form novel Van der Waals heterostructures for new functionalities. An outstanding challenge for realising the application potential of 2D dichalcogenides – like in the cases of graphene – is to engineer Ohmic contacts. Transition metals deposited on 2Ds tend to cluster. The extent of the clustering does, however, depend on the metal, as previously pointed out for the case of graphene. Especially in the case of MoS2, the deposition of metal contacts may affect the structure and electronic properties substantially.110,111 The high reactivity of metal atoms involved in the deposition process can strongly perturb phonons and electrons in a monolayer MoS2 film. It is, therefore, important to investigate the metal deposition and interaction to better understand and control the electronic transport and heat dissipation; so, for example, it is essential to achieve uniform wetting of the 2-D surfaces by metals. High hole mobilites in WSe2 monolayer-FETs are achieved with degenerately doped contacts: the use of heavily p-doped contacts proved essential in lowering the resistances of metal contacts to WSe2 by orders of magnitude, and to enable the demonstration of p-FETs with effective peak mobilities of y250 cm2 V21 s21, and high ON/OFF ratios of .106. Currently, this is achieved by NO2 surface doping of the source/drain regions (Ref. 112, and references therein). The authors emphasise the necessity of surface doping for obtaining high performance monolayer-FETs, and in this regard the exploration of other dopant species for both n- and p-doping is needed in the future.

References 1. F. Schedin, A. K. Geim, S. V. Morozov, E. W. Hill, P. Blake, M. I. Katsnelson and K. S. Novoselov: Nat. Mater., 2007, 6, 652–655. 2. L. Ju, B. Geng, J. Horng, C. Girit, M. Martin, Z. Hao, H. A. Bechtel, X. Liang, A. Zettl, Y. R. Shen and F. Wang: Nat. Nanotechnol., 2011, 6, 630–634. 3. F. Scha¨ffel, J. H. Warner, A. Bachmatiuk, U. Queitsch, B. Rellinghaus, B. Bu¨chner, L. Schultz and M. H. Ru¨mmeli: Phys. Status Solidi, 2010, C 7, 2731–2734. 4. X. Wang and H. Dai: Nat. Chem., 2010, 2, 661–665. 5. L. P. Biro´ and P. Lambin: Carbon, 2010, 48, 2677–2689. [review]. 6. A. Venugopal, L. Colombo and E. M. Vogel: Appl. Phys. Lett., 2010, 96, 013512.

148


2015

VOL

60

NO

3

7. F. Xia, V. Perebeinos, Y.-M. Lin, Y. Wu and P. Avouris: Nat. Nanotechnol., 2011, 6, 179–184. 8. J. Lee, S. Shim, B. Kim and H. S. Shin: Chem. Eur. J., 2011, 17, 2381–2387. 9. F. Schedin, E. Lidorikis, A. Lombardo, V. G. Kravets, A. K. Geim, A. N. Grigorenko, K. S. Novoselov and A. C. Ferrari: ACS Nano., 2010, 4, (10), 5617–5626. 10. J. Lee, K. S. Novoselov and H. S. Shin: ACS Nano., 2011, 5, (1), 608–612. 11. S. Entani, S. Sakai, Y. Matsumoto, H. Naramoto, T. Hao and Y. Maeda: J. Phys. Chem. C, 2010, 114, 20042–20048. 12. W. X. Wang, S. H. Liang, T. Yu, D. H. Li, Y. B. Li and X. F. Han: J. Appl. Phys, 2011, 109, 07C501. 13. K. S. Novoselov, A. K. Geim, S. V. Morozov, D. Jiang, Y. Zhang, S. V. Dubonos, I. V. Grigorieva and A. A. Firsov: Science, 2004, 306, 666–669. 14. T. O. Wehling, K. S. Novoselov, S. V. Morozov, E. E. Vdovin, M. I. Katsnelson, A. K. Geim and A. I. Lichtenstein: Nano Lett., 2008, 8, (1), 173–177. 15. A. Castellanos-Gomez, M. Wojtaszek, A. N. Tombros and B. J. van Wees: Small, 2012, 8, (10), 1607–1613. 16. L. Liu, S. Ryu, M. R. Tomasik, E. Stolyarova, N. Jung, M. S. Hybertsen, M. L. Steigerwald, L. E. Brus and G. W. Flynn: Nano Lett., 2008, 8, (7), 1965–1970. 17. Z. Jin, J. Yao, C. Kittrell and J. M. Tour: ACS Nano., 2011, 5, 4112–4117. 18. L. Zhao, R. He, K. T. Rim, T. Schiros, K. S. Kim, H. Zhou, C. Gutie´rrez, S. P. Chockalingam, C. J. Arguello, L. Pa´lova´, D. Nordlund, M. S. Hybertsen, D. R. Reichman, T. F. Heinz, P. Kim, A. Pinczuk, G. W. Flynn and A. N. Pasupathy: Science, 2011, 333, 999–1003. 19. D. Usachov, O. Vilkov, A. Gru¨neis, D. Haberer, A. Fedorov, V. K. Adamchuk, A. B. Preobrajenski, P. Dudin, A. Barinov, M. Oehzelt, C. Laubschat and D. V. Vyalikh: Nano Lett., 2011, 11, 5401–5407. 20. T. Wu, H. Shen, L. Sun, B. Cheng, B. Liu and J. Shen: New J. Chem., 2012, 36, 1385–1391. 21. W. Zhao, O. Ho¨fert, K. Gotterbarm, J. Zhu, C. Papp and H.-P. Steinru¨ck: J. Phys. Chem. C, 2012, 116, 5062–5066. 22. U. Bangert, W. Pierce, D. M. Kepaptsoglou, Q. Ramasse, R. Zan, M. H. Gass, J. A. Van den Berg, C. B. Boothroyd, J. Amani and H. Hofsa¨ss: Nano Lett., 2013, 13, 4902–4907. 23. C. Mangler and J. C. Meyer: C. R. Phys., 2014, 15, 241–257. 24. B. Kabius, P. Hartel, M. Haider, H. Mu¨ller, S. Uhlemann, U. Loebau, J. Zach and H. Rose: J. Electron Microsc., 2009, 58, 147–155. 25. J. Zach: Philos. Trans. R. Soc. London, Ser. A, 2009, 367, 3699. 26. O. L. Krivanek, N. Dellby, M. F. Murfitt, M. Chisholm, T. J. Pennycook, K. Suenaga and V. Nicolosi: Ultramicroscopy, 2010, 110, 935–945. 27. A. Hashimoto, K. Suenaga, A. Gloter, K. Urita and S. Iijima: Nature, 2004, 430, 870–873. ¨ . Girit, J. C. Meyer, R. Erni, M. D. Rossell, C. Kisielowski, 28. C ¸. O L. Yang, C.-H. Park, M. F. Crommie, M. L. Cohen, S. G. Louie and A. Zettl: Science, 2009, 323, 1705–1708. 29. Z. Liu, K. Suenaga, P. J. F. Harris and S. Iijima: Phys. Rev. Lett, 2009, 102, (1–4), 015501. 30. P. Koskinen, S. Malola and H. Ha¨kkinen: Phys. Rev. B, 2009, 80, (1–3), 073401. 31. J. Kotakoski, J. C. Meyer, S. Kurasch, D. Santos-Cottin, U. Kaiser and A. V. Krasheninnikov: Phys. Rev. B, 2011, 83, (106), 245420. 32. J. H. Warner, E. Roxana Margine, M. Mukai, A. W. Robertson, F. Giustino and A. I. Kirkland: Science, 2012, 337, 209–212. 33. A. Chuvilin, J. C. Meyer, G. Algara-Siller and U. Kaiser: New J. Phys., 2009, 11, (1–9), 083019. 34. W. Zhou, M. D. Kapetanakis, M. P. Prange, S. T. Pantelides, S. J. Pennycook and J.-C. Idrobo: Phys. Rev. Lett., 2012, 109, 206803. 35. Q. M. Ramasse, C. R. Seabourne, R. Zan, D. M. Kepaptsoglou, U. Bangert and A. J. Scott: Nano Lett., 2013, 13, 4989–4995. 36. A. T. N’Diaye, S. Bleikamp, P. J. Feibelman and T. Michely: Phys. Rev. Lett., 2006, 97, 215501. 37. M. Ishigami, J. H. Chen, W. G. Cullen, M. S. Fu˜hrer and E. D. Williams: Nano Lett., 2007, 7, 1643–1648. 38. G. Li, A. Luican and E. Y. Andrei: Phys. Rev. Lett., 2009, 102, 176804. 39. Z. Zhou, F. Gao and D. W. Goodman: Surf. Sci., 2010, 604, L31– L38. 40. R. Zan, C. Muryn, U. Bangert, P. Mattocks, P. Wincott, D. Vaughan, X. Li, L. Colombo, R. S. Ruoff, B. Hamilton and K. S. Novoselov: Nanoscale, 2012, 4, 3065–3068. 41. G. Algara-Siller, O. Lethinen, A. Turchanin and U. Kaiser: Appl. Phys. Lett., 2014, 104, 153115.

Bangert and Zan

42. J. C. Meyer, O. C. Girit, M. F. Crommie and A. Zettl: Nat. Lett., 2008, 454, 319–322. 43. J. H. Warner, M. H. Ru¨mmeli, A. Bachmatiuk, M. Wilson and B. Bu¨chner: ACS Nano., 2010, 4, 470–476. 44. Q. M. Ramasse, R. Zan, U. Bangert, D. W. Boukhvalov, Y.-W. Son and K. S. Novoselov: ACS Nano., 2012, 6, (5), 4063–4071. 45. R. Zan, U. Bangert, Q. Ramasse and K. S. Novoselov: ‘invited Perspective’, J. Phys. Chem. Lett, 2012, 3, (7), 953–958. 46. R. Erni, M. D. Rossell, M.-T. Nguyen, S. Blankenburg, D. Passerone and P. Hartel: Phys. Rev. B, 2010, 82, 165443. 47. A. Thust, W. M. J. Coene, M. Op de Beeck and D. Van Dyck: Ultramicroscopy, 1996, 64, 211–230. 48. X. Gao, Y. Wang, X. Liu, T.-L. Chan, S. Irle, Y. Zhao and S. B. Zhang: Phys. Chem. Chem. Phys., 2011, 13, 19449–19453. 49. D. W. Boukhvalov and M. I. Katsnelson: J. Phys. Chem. C, 2009, 113, 14176–14178. 50. W. Bao, F. Miao, Z. Chen, H. Zhang, W. Jang, C. Dames and C. N. Lau: Nat. Nanotechnol., 2009, 4, 562–566. 51. H. Wang, K. Li, Y. Cheng, Q. Wang, Y. Yao, U. Schwingenschloegl, X. Zhang and W. Yang: Nanoscale, 2012, 4, 2920–2925. 52. F. Schaeffel, M. Wilson and J. H. Warner: ACS Nano., 2011, 5, 9428–9441. 53. F. Banhart, J. Kotakoski and A. V. Krasheninnikov: ACS Nano, 2011, 5, 26–41. 54. J. C. Meyer, C. Kisielowski, R. Erni, M. D. Rossell, M. F. Crommie and A. Zettl: Nano Lett., 2008, 8, (11), 2582–2586. 55. M. H. Gass, U. Bangert, A. L. Bleloch, P. Wang, R. L. R. Nair and A. K. Geim: Nat. Nanotechnol., 2008, 3, 676–681. 56. Y. Kim, J. Ihm, E. Yoon and G.-D. Lee: Phys. Rev. B, 2011, 84, 075445. 57. J. Kotakoski, A. V. Krasheninnikov, U. Kaiser and J. C. Meyer: Phys. Rev. Lett., 2011, 106, 105505. 58. R. Zan, Q. M. Ramasse, U. Bangert and K. S. Novoselov: Nano Lett., 2012, 12, 3936–3940. 59. H. Wang, Q. Wang, Y. Cheng, K. Li, Y. Yingbang, Q. Zhang, C. Dong, P. Wang, U. Schwingenschlo¨gl, W. Yang and X. X. Zhang: Nano Lett., 2012, 12, 141–144. 60. A. W. Robertson, B. Montanari, K. He, J. Kim, C. S. Allen, Y. A. Wu, J. Olivier, J. Neethling, N. Harrison, A. I. Kirkland and J. H. Warner: Nano Lett., 2013, 13, 1468–1475. 61. J. Lee, W. Zhou, S. J. Pennycook, J.-C. Idrobo and S. T. Pantelides: Nat. Commun., 2013, 4, 1650. 62. J. H. Warner, Z. Liu, K. He, A. W. Robertson and K. Suenaga: Nano Lett., 2013, 13, 4820–4826. 63. J. A. Rodrı´guez-Manzo, O. Cretu and F. Banhart: ACS Nano, 2010, 4, (6), 3422–3428. 64. J. J. Palacios, A. J. Pe´rez-Jimeńez, E. Louis, E. SanFabiań and J. A. Verge´s: Phys. Rev. Lett., 2003, 90, 106801. 65. N. Nemec, D. Tomańek and G. Cuniberti: Phys. Rev. Lett., 2006, 96, 076802. 66. A. V. Kosinnikov, P. O. Lehtinen, A. S. Foster, P. Pyykko¨ and R. M. Nieminen: Phys. Rev. Lett., 2009, 102, 126807. 67. Y. Gan, L. Sun and F. Banhart: Small, 2008, 4, 587–591. 68. H. L. Zhuang, G. Zheng and A. K. Soh: Comput. Mater. Sci., 2008, 43, 823–828. 69. T. Hardcastle, C. R. Seabourne, R. Zan, R. M. D. Brydson, U. Bangert, Q. M. Ramasse, K. S. Novoselov and A. J. Scott: Phys. Rev. B, 2013, 87, 195430. 70. R. Zan, U. Bangert, Q. Ramasse and K. S. Novoselov: Small, 2011, 7, (20), 2868–2872. 71. R. Zan, U. Bangert, Q. Ramasse and K. S. Novoselov: Nano Lett., 2011, 11, (3), 1087–1095. 72. J. A. Rodriguez-Manzo and F. Banhart: Nano Lett., 2009, 9, 2285–2289. 73. O. Cretu, A. V. Krasheninnikov, J. A. Rodrı´guez-Manzo, L. Sun, R. M. Nieminen and F. Banhart: Phys. Rev. Lett., 2010, 105, 196102. 74. A. V. Krasheinnikov and R. M. Nieminen: Theor. Chem. Acc., 2011, 129, 625–630. 75. G. Giovannetti, P. A. Khomyakov, G. Brocks, V. M. Karpan, J. van den Brink and P. J. Kelly: Phys. Rev. Lett., 2008, 101, 026803. 76. M. M. Ugeda, D. Fernańdez-Torre, I. Brihuega, P. Pou, A. J. Martıńez-Galera, R. Pe´rez and J. M. Go´mez-Rodrı´guez: Phys. Rev. Lett., 2011, 107, 116803. 77. V. W. Brar, R. Decker, H.-M. Solowan, Y. Wang, L. Maserati, K. T. Chan, H. Lee, C ¸ . O. Girit, A. Zettl, S. G. Louie, M. L. Cohen and M. F. Crommie: Nat. Phys., 2011, 7, 43–47. 78. K. Pi, K. M. McCreary, W. Bao, W. Han, Y. F. Chiang, Y. Li, S.-W. Tsai, C. N. Lau and R. K. Kawakami: Phys. Rev. B, 2009, 80, 075406.


79. Y. Ren, S. Chen, W. Cai, Y. Zhu, C. Zhu and R. S. Ruoff: Appl. Phys. Lett., 2010, 97, 053107. 80. B. Uchoa, C.-Y. Lin and A. H. Castro Neto: Phys. Rev. B, 2008, 77, 035420. 81. X. Liu, C. Z. Wang, M. Hupalo, W. C. Lu, M. C. Tringides, Y. X. Yao and K. M. Ho: Phys. Chem. Chem. Phys, 2012, 14, 9157– 9166. 82. S. Stambula, N. Gauquelin, M. Bugnet, S. Gorantla, S. Turner, S. Sun, J. Liu, G. Zhang, X. Sun and G. A. Botton: J. Phys. Chem. C, 2014, 118, 3890–3900. 83. M.-X. Wang, Q. Liu, Z.-F. Li, H.-F. Sun, E. A. Stach and J. Xie: J. Phys. Chem. Lett., 2013, 4, 1484–1488. 84. D. W. Boukhvalov and M. I. Katsnelson: Appl. Phys. Lett., 2009, 95, 023109. 85. M. Terrones, A. Jorio, M. Endo, A. M. Rao, Y. A. Kim, T. Hayashi, H. Terrones, J.-C. Charlier, G. Dresselhaus and M. S. Dresselhaus: Mater. Today, 2004, 7, 30–45. 86. H. Wang, T. Maiyalagan and X. Wang: ACS Catal., 2012, 2, 781– 794. 87. E. Cruz-Silva, F. Lo´pez-Urıás, E. Munõz-Sandoval, B. G. Sumpter, H. Terrones, J.-C. Charlier, V. Meunier and M. Terrones: ACS Nano, 2009, 3, 1913–1921. 88. J. C. Meyer, S. Kurasch, H. J. Park, V. Skakalova, D. Ku¨nzel, A. Groß, A. Chuvilin, G. Algara-Siller, S. Roth, T. Iwasaki, U. Starke, J. H. Smet and U. Kaiser: Nat. Mater., 2011, 10, 209–215. 89. R. Lv, Q. Li, A. R. Botello-Meńdez, T. Hayashi, B. Wang, A. Berkdemir, Q. Hao, A. L. Elıás, R. Cruz-Silva, H. R. Gutie´rrez, Y. A. Kim, H. Muramatsu, J. Zh, M. Endo, H. Terrones, J.-C. Charlier, M. Pan and M. Terrones: Sci. Rep., 2012, 2, 586. 90. L. Zhao, M. Levendorf, S. Goncher, T. Schiros, L. Pa´lova´, A. Zabet-Khosousi, K. T. Rim, C. Gutie´rrez, D. Nordlund, C. Jaye, M. Hybertsen, D. Reichman, G. W. Flynn, J. Park and A. N. Pasupathy: Nano Lett., 2013, 13, 4659–4665. 91. R. J. Nicholls, A. T. Murdock, J. Tsang, J. Britton, T. J. Pennycook, A. L. Koo´s, P. D. Nellist, N. Grobert and J. R. Yates: ACS Nano, 2013, 7, 7145–7150. ˚ hlgren, J. Kotakoski and A. V. Krasheninnikov: Phys. 92. E. H. A Rev. B, 2011, 83, (1–4), 115424. 93. Y. Pan, M. Gao, L. Huang, F. Liu and H.-J. Gao: Appl Phys Lett., 2009, 95, 093106. 94. A. T. N’Diaye, T. Gerber, C. Busse, J. Myslivecˇek, J. Coraux and T. Michely: New J. Phys., 2009, 11, 103045. 95. M. Sicot, S. Bouvron, O. Zander, U. Ru¨diger, Y. S. Dedkov and M. Fonin: Appl. Phys. Lett., 2010, 96, 093115. 96. A. J. Pollard, E. W. Perkins, N. A. Smith, A. Saywell, G. Goretzki, A. G. Phillips, S. P. Argent, H. Sachdev, F. Mueller, S. Hoefner, S. Gsell, M. Fischer, M. Schreck, J. Osterwalder, T. Greber, S. Berner, N. R. Champness and P. H. Beton: Angew. Chem., 2010, 49, (Int. Ed), 1794–1799. 97. W. Chen, S. Chen, D. C. Qi, X. Y. Gao and A. T. S. Wee: J. Am. Chem. Soc., 2007, 129, 10418–10422. 98. I. Pletikosic´, M. Kralj, P. Pervan, R. Brako, J. Coraux, A. T. N. Diaye, C. Busse and T. Michely: Phys. Rev. Lett., 2009, 102, 056808. 99. B. Premlal, M. Cranney, F. Vonau, D. Aubel, D. Casterman, M. M. De Souza and L. Simon: Appl. Phys. Lett., 2009, 94, 263115. 100. S.-Y. Sheu and D.-Y. Yang: Carbon, 2014, 71, 76–86. 101. Z.-J. Jiang, Z. Jiang and W. Chen: J. Power Sources, 2014, 251, 55–65. 102. Q. Xu, M.-Y. Wu, G. F. Schneider, L. Houben, S. K. Malladi, C. Dekker, E. Yucelen, R. E. Dunin-Borkowski and H. W. Zandbergen: ACS Nano, 2013, 7, 1566–1572. 103. K. S. Novoselov, D. Jiang, F. Schedin, T. J. Booth, V. V. Khotkevich, S. V. Morozov and A. K. Geim: Proc. Natl. Acad. Sci. U.S.A, 2005, 102, 10451–10453. 104. K. F. Mak, C. Lee, J. Hone, J. Shan and T. F. Heinz: Phys. Rev. Lett., 2010, 105, 136805. 105. A. Splendiani, L. Sun, Y. Zhang, T. Li, J. Kim, C.-Y. Chim, G. Galli and F. Wang: Nano Lett., 2010, 10, 1271–1275. 106. B. Radisavljevic, A. Radenovic, J. Brivio, V. Giacometti and A. Kis: Nat. Nanotechnol., 2011, 6, 147–150. 107. S. Das, H. Y. Chen, A. V. Penumatcha and J. Appenzeller: Nano Lett., 2013, 13, 100–105. 108. M. S. Fu¨hrer and J. Hone: Nat. Nanotechnol., 2013, 8, 146–147. 109. H. Wang, L. L. Yu, Y. H. Lee, Y. M. Shi, A. Hsu, M. L. Chin, L. J. Li, M. Dubey, J. Kong and T. Palacios: Nano Lett., 2012, 12, 4674–4680. 110. I. Popov, G. Seifert and D. Tomanek: Phys. Rev. Lett., 2012, 108, 156802. 111. W. Chen, E. J. G. Santos, W. G. Zhu, E. Kaxiras and Z. Y. Zhang: Nano Lett., 2013, 13, 509–514. 112. H. Fang, S. Chuang, T. C. Chang, K. Takei, T. Takahashi and A. Javey: Nano Lett., 2012, 12, (7), 3788–3792.


2015

VOL

60

NO

3

149

Electronic functionalisation of graphene via external doping and dosing

Electronic functionalisation of graphene via external doping and dosing

Suggest Documents

Electronic Strengthening of Graphene by Charge Doping

Electronic Strengthening of Graphene by Charge Doping

Electronic doping of graphene by deposited transition metal atoms

X-ray induced electrostatic graphene doping via defect charging in

Doping Graphene via Organic Solid-Solid Wetting Deposition - arXiv

Synergetic Effects of Al3+ Doping and Graphene

Dynamical polarization of graphene at finite doping

A FACILE DOPING OF GRAPHENE OXIDE NANOSHEETS

Single step, complementary doping of graphene - Squarespace

Structural and Electronic Properties of Graphene and Graphene-like

Electronic transition from graphite to graphene via controlled ... - arXiv

Electronic properties of graphene and graphene nanoribbons with ...

Diastereoselective functionalisation of benzo

Electronic and Transport Properties of Graphene - Springer

Electronic properties of bilayer and multilayer graphene

Structural and Electronic Properties of Oxidized Graphene

Structural and electronic properties of graphene ... - Refubium

Electronic properties of graphene - II

Superconducting Graphene: the conspiracy of doping and strain

Formation, Energetics, and Electronic Properties of Graphene

Mechanical and electronic properties of graphene

electronic properties of functionalized graphene

Electronic properties of monolayer and bilayer graphene

Electronic and Transport Properties of Graphene - Springer