Dipartimento di Fisica
Corso di Laurea Magistrale in Fisica
TESI
DI
LAUREA
MAGISTRALE
Scalable synthesis of WS
2
on CVD graphene:
heterostructure properties and optoelectronic
potential applications
CANDIDATA
RELATORI
Giulia Piccinini
Dott.ssa Camilla Coletti
Prof. Alessandro Tredicucci
CORRELATORE
Dott. Filippo Fabbri
C
ONTENTS
INTRODUCTION 1
1 GRAPHENE AND OTHER 2D MATERIALS 7
1.1 GRAPHENEPROPERTIES 7
1.2 PROPERTIESOFTRANSITIONMETALDICHALCOGENIDES 13 1.3 VDWHETEROSTRUCTURESANDTHEIRAPPLICATIONS 16
2 EXPERIMENTAL TECHNIQUES 19
2.1 CVDSYNTHESISOFGRAPHENEANDTRANSFERONASELECTEDSUBSTRATE 19 2.2 CVDSYNTHESISOFTUNGSTENDISULPHIDE 23 2.3 RAMANSPECTROSCOPYANDPHOTOLUMINESCENCE 25
2.3.1 RAMANONGRAPHENE 26
2.3.2 RAMANONWS2 29
2.4 ATOMICFORCEMICROSCOPY 30 2.5 SCANNINGELECTRONMICROSCOPY 31
3 SCALABLE SYNTHESIS OF GRAPHENE/WS2 VDWH 34
3.1 STATEOFTHEART 34
3.2 DIRECTCVDOFTUNGSTENDISULPHIDEONCVDGRAPHENE 36
3.2.1 PMMARESIDUESASWS2GROWTHPROMOTERS 40
4 EFFECTS OF WS2 GROWTH ON GRAPHENE 44
4.1 CHANGESINGRAPHENEPROPERTIESAFTERWS2SYNTHESIS 44
4.2 EFFECTOFTHERMALANNEALINGONGRAPHENEPROPERTIES 47
5 FIELD-EFFECT TRANSISTOR PERFORMANCES 52
5.1 FETFABRICATION 53
5.2 TRANSPORTMEASUREMENTS 55 5.3 PHOTOTRANSISTORRESPONSIVITY 60
CONCLUSIONS AND PERSPECTIVES 63
ELECTRON-BEAMLITHOGRAPHYFOLLOWEDBYREACTIVEIONETCHINGORMETALLIZATION 65
INTRODUCTION
Graphene is undoubtedly emerging as one of the most promising nanomaterials in a wide range of applications because of its outstanding electronic, optical, thermal, and mechanical properties. The isolation of graphene in 2004 by Geim and Novoselov first showed the possibility to obtain stable, single-atom layer two-dimensional (2D) materials from their van der Waals (vdW) solids. Despite its short history, graphene has already revealed a cornucopia of new physics and potential applications. One of the most notable property of this material is that charge carriers in graphene are able to travel sub-micrometer distances without scattering, a phenomenon known as ballistic transport. Tests have shown that carrier mobility for free-standing graphene can reach 200 000 cm2/Vs at room temperature (limited by the scattering of electrons with graphene’s
acoustic photons), the highest value ever observed in pure semiconductors [1], [2]. However, the quality of the graphene and the substrate that is used are the limiting factors for high carrier mobility. Indeed, the first experimental studies of graphene exfoliated onto silicon dioxide (SiO2) substrates, the substrate of choice in
microfabrication of graphene for its availability and versatility, showed device mobilities of less than about 20 000 cm2/Vs. The electronic response is dominated by scattering from
charged surface states and impurities [3], substrate surface roughness [2], [4] and SiO2
surface optical phonons [4]–[6]. There is a growing need, therefore, to identify materials, to be used as graphene encapsulants, that allow to enhance graphene carrier mobility by minimizing extrinsic sources of scattering, coming from both interface with substrate and air. In this context, hexagonal boron nitride (h-BN), an insulating isomorph of graphite with atomically flat and nearly free charge trapping layers, has proved to be an excellent candidate. An increase in mobility up to ∼100 000 cm2/Vs was obtained when
encapsulating graphene between exfoliated h-BN flakes (i.e., with bottom and top encapsulation) [7].
All the studies reported above have been performed with graphene (and h-BN) flakes obtained by mechanical exfoliation which is not a viable route to produce scalable 2D
Introduction
2
crystals due to the limited dimensions of the exfoliated flakes (typically a few tens of micrometers). In order to grow suitable graphene (and encapsulants) for large-scale device production, a precise control over flake size, layer number and morphology is necessary. To move towards realistic applications, scalable graphene production techniques have been developed, with chemical vapor deposition (CVD) on transition metals being the most effective. However, suitable large area encapsulants capable of exploiting graphene’s remarkable electronic properties have still to be found. In particular, at present, a CVD process for obtaining high quality h-BN is not widely available.
Since the isolation of graphene, a wide range of other 2D materials have been isolated, studied and synthetized. In particular, transition metal dichalcogenides (TMDs) have attracted notable attention because of their diverse properties and natural abundance. One of the most studied TMDs, tungsten disulfide (WS2), was found to be a good
encapsulant for improving the electrical properties of CVD graphene [8]: high graphene mobilities of ∼ 60 000 cm2/Vs and high charge homogeneity can be obtained when WS
2
exfoliated flakes are used as substrate (with h-BN used as top-encapsulant). These results indicate that WS2 is an appealing alternative to h-BN, but its performance as scalable
encapsulant (i.e., by using large area WS2 and not exfoliated flakes) remains to be studied.
CVD processes for the synthesis of WS2 have been developed in the last few years, but the
(opto)electronic properties of CVD WS2 vertically stacked on top of graphene have not yet
been systematically assessed. Hence, it would be appropriate to understand if it is possible to grow WS2 directly on CVD graphene, and to investigate the heterostack’s
optoelectronic properties and potential.
Indeed, besides being interesting as atomically flat encapsulant for improving electronic properties of graphene, TMDs assembled in van der Waals heterostructures (vdWH) with graphene are also appealing for a wide number of applications. Although research on vdWH is just at its beginning, indeed, such heterojunctions have been already applied to form devices, such as photodiodes [9], [10], phototransistors [11], [12], tunneling devices [13] and memory devices [14]. Notably, recent developments in vdWH research suggest that such layered heterostructures might facilitate progresses in modern integrated circuits industry.
Introduction
3
Concerning the WS2/graphene heterostructure, by combining WS2 optical absorbance
with the outstanding transport properties of graphene, efficient photodetectors have been demonstrated [15]. The system has a serious appeal for many other applications, ranging from flexible and transparent electronics [13] to optospintronics [16]. Developing a CVD process for the scalable synthesis of WS2/graphene heterostacks would
mean the possibility to choose any desired substrate, paving the way for new large-scale applications.
The goals of this thesis are: (i) to demonstrate the synthesis of scalable graphene/WS2
heterostructures on top of a technologically relevant substrate (i.e., SiO2), (ii) to
investigate the properties of the heterostack and of the two constituting 2D materials and (iii) to fabricate devices for assessing the electronic and optoelectronic performance of such 2D heterostack. By synthesizing both materials via CVD, with WS2 grown directly on
graphene, I demonstrated that a fully and homogeneous coverage of WS2 on top of
graphene can be obtained. Moreover, the avoidance of WS2 transfer results in a
monolithic structure with an atomically clean interface and in a faster stacking process. I used spectroscopic and microscopic techniques to outline the structural features of the grown materials. Besides the importance of completely covering graphene with a flat and inert material for improving its quality, WS2 is interesting as active material. Therefore, a
knowledge of its thickness is fundamental to figure out its optical properties, which in the case of TMDs are deeply connected with the number of layers. Thanks to Raman spectroscopy, I investigated the effects of WS2 growth on the underlying graphene layer,
which resulted affected by both compressive strain and doping. I demonstrated that the main reason of such doping induced in graphene is the presence of temperature-activated (during CVD growth) interface states at the SiO2 interface. Graphene field-effect
transistors (GFETs) devices were fabricated and gate voltage-dependent electronic transport measurements were performed. P-type doping of WS2-coated
graphene was confirmed. Moreover, transfer characteristics showed asymmetry between hole and electron conduction, accompanied by a forward/reverse sweep hysteresis. These effects are due to defects in WS2 which act as a sink to graphene electrons. Also,
the heterostructure turned out to be photoresponsive and the estimated photoresponsivity is higher than the standing record of CVD WS2 on graphene (i.e.,
Introduction
4 The thesis outline is presented below.
Chapter 1 discusses the main structural and electronic properties of graphene and
transition metal dichalcogenides, WS2 in particular. The importance of the dispersion
relation and helicity conservation for high carrier mobility in graphene is stressed.
Chapter 2 reports the CVD method for graphene growth and all the details about how
I synthesized graphene in our laboratories. The semi-dry transfer process I implemented to get graphene flakes on the desired substrate is also explained. This chapter gives also a presentation of the experimental techniques used in this work to characterize 2D materials: Raman spectroscopy and photoluminescence, atomic force microscopy and scanning electron microscopy. The section about Raman spectroscopy is followed by a brief presentation of the main Raman features of both graphene and WS2 and the phonon
modes from which they are generated.
Chapter 3 gives an overview about the state of the art of WS2 growth via CVD and the
recent progresses on stacking graphene with WS2 in a scalable manner. The CVD process
I implemented for WS2 synthesis is here reported in detail and the features of the resulting
material discussed. This work demonstrates that WS2 monolayer with uniform coverage
can be obtained on top of graphene, although three-dimensional WS2 crystals are
identified on top of the WS2 monolayer. Also, selective growth of WS2 on top of graphene,
with respect to the SiO2 surface, is observed. The reasons behind these findings are
reported and discussed.
Chapter 4 reports the Raman characterization of the underlying graphene layer.
Raman spectroscopy is exploited to understand how the synthesis of WS2 performed
directly on CVD graphene affects the graphene electronic properties. Correlations plots of the two characteristic Raman peaks of graphene are used to investigate the presence of strain and doping. A study of annealed graphene samples reveals that the doping induced in graphene after WS2 synthesis is due to the thermal process undergone by the
sample. Moreover, graphene native tensile strain is removed and a much stronger compressive strain induced after the top-encapsulation with WS2. This may be due to the
need of the graphene lattice to reach a more stable configuration by forming a moiré pattern with WS2.
Chapter 5 is centred on device fabrication and direct evaluation of the vertical
Introduction
5
measurements are reported and discussed. Furthermore, thanks to measurements performed both in the dark and with the sample illuminated, the photoresponsivity of such devices was demonstrated. An estimation of the heterostructure photoresponsivity and a possible explanation of the high value obtained is given.
Transport measurement were performed by Leonardo Martini (CNI@NEST Pisa).
Finally, in Conclusions and perspectives the major scientific findings of this work are summarized and a series of perspectives for additional research in the field are proposed.
6
1
GRAPHENE
AND
OTHER
2D
MATERIALS
1.1
GRAPHENE PROPERTIES
More than 70 years ago, Landau [18] and Peierls [19] argued that strictly two-dimensional (2D) crystals were thermodynamically unstable and could not exist. Their theory pointed out that a divergent contribution of thermal fluctuations in low-dimensional crystal lattices leads to displacements of atoms comparable to interatomic distances at any finite temperature. 2D materials were presumed not to exist until 2004, when Andre Geim and Konstantin Novoselov [20] managed to isolate and investigate graphene, leading to the beginning of the “graphene boom” and opening the door to new 2D materials.
Graphene is an allotrope of carbon consisting of a single layer of carbon atoms arranged in a honeycomb lattice. Such a lattice is the result of the 𝑠𝑝2 hybridization of the 2𝑠 and 2𝑝 orbitals of C atoms, while the 1𝑠 electrons remain more or less inert. The unit cell contains two atoms belonging to two sublattices, 𝐴 and 𝐵. Each atom of sublattice A is surrounded by three atoms of sublattice B, and viceversa.
Fig. 1.1 (a) Honeycomb lattice of graphene with a1 and a2 base vectors and δ1, δ2, δ3 vectors connecting each atom
to the nearest neighbour. Sublattices A and B are shown as blue and red. (b) Representation of graphene lattice in reciprocal space and some special points in the Brillouin zone [21].
The lattice vectors of the Bravais lattice are
a1 = 𝑎 2 (3, √3), a2 = 𝑎 2 (3, −√3), (1) (a) (b)
1 Graphene and other 2D materials
8
where 𝑎 ≈ 1.42 Å is the nearest-neighbour distance.
The vectors connecting each atom to the nearest neighbours are
δ1=
𝑎
2 (1, √3), δ 2= 𝑎
2 (1, − √3), δ 3= 𝑎 (−1, 0). (2)
The reciprocal lattice is represented by the vectors
b1=
2𝜋
3𝑎 (1, √3), b2 =
2𝜋
3𝑎 (1, −√3). (3)
In panel (b) of Fig. 1.1 is shown the Brillouin zone (BZ) with the high-symmetry points
𝜞 = (0,0), 𝑲 = (2𝜋 3𝑎, 2𝜋 3√3𝑎), 𝑲’ = ( 2𝜋 3𝑎, − 2𝜋 3√3𝑎), 𝑴 = ( 2𝜋 3𝑎, 0). (4)
Fig. 1.2 Band structure of graphene [22].
In Fig. 1.2 is shown the graphene band structure. The 1s core atomic orbitals of carbon atoms are strongly localized near the nuclei and give rise to dispersionless crystal core levels that do not appear in the figure. The 𝑠𝑝2 hybridized states (𝜎-states) give rise to occupied and empty bands (𝜎-bands) with a huge gap, whereas 𝑝𝑧orbitals form a single band (𝜋-band) with a conical self-crossing point at the 𝑲 point of the BZ (and by symmetry also at 𝑲’). This band structure was obtained for the first time by Wallace in 1947, by a simple tight-band model [23]. Only the 𝜋-states and nearest-neighbour hopping energy
1 Graphene and other 2D materials
9
(parameter 𝑡 ) are considered, so that there are only hopping processes from one sublattice to the other (nearest-neighbours belong to different sublattices). The basis of the electronic states contains two 𝜋-states per unit cell belonging to atoms of sublattices 𝐴 and 𝐵. The tight-binding Hamiltonian is therefore
𝐻̂(𝒌) = ( 0 𝑡𝐹(𝒌)
𝑡𝐹∗(𝒌) 0 ) (5)
where 𝐹(𝒌) denotes the sum
𝐹(𝒌) = ∑ 𝑒𝑖𝒌⋅𝜹 𝛅 = 2 exp (𝑖𝑘𝑥𝑎 2 ) cos ( 𝑘𝑦𝑎√3 2 ) + exp(−𝑖𝑘𝑥𝑎) (6)
The resulting energy spectrum is
𝐸(𝒌) = ±𝑡|𝐹(𝒌)| = ±𝑡√3 + 𝑓(𝒌) (7) where 𝑓(𝒌) = 2 cos(√3𝑘𝑦𝑎) + 4 cos ( √3 2 𝑘𝑦𝑎) cos ( 3 2𝑘𝑥𝑎) (8)
Since 𝐹(𝑲) = 𝐹(𝑲′) = 0 , the two bands cross at these points. By expanding the Hamiltonian near these points, we find
𝐻̂𝐾(𝒒) ≈ 3𝑎𝑡 2 ( 0 𝑞𝑥− 𝑖𝑞𝑦 𝑞𝑥+ 𝑖𝑞𝑦 0 ) (9)
where 𝒒 = 𝒌 − 𝑲. The prefactor in the right-hand side of Eq. (9) can be interpreted as ℏ𝜈, where
1 Graphene and other 2D materials 10 𝜈𝐹= 3𝑎𝑡 2ℏ (10)
is the electron velocity at the Fermi level. Experimentally, 𝜈𝐹≈ 𝑐
300≈ 10
6 m/s where 𝑐 is
the velocity of light.
The Ĥ𝐾(𝑞) eigenvalues are a function only of the magnitude of 𝒒
𝐸(𝒒) = ±ℏ𝑣𝐹|𝒒| (11)
It is very suggestive to write the Hamiltonian in the form
𝐻̂𝐾(𝒒) ≡ ℏ𝑣𝐹𝝈̂ ⋅ 𝒒 (12)
where the components of the operator 𝝈̂ are the usual Pauli matrices.
The Hamiltonian of Eq. (12) is a 2D analogue of the Dirac Hamiltonian for ultra-relativistic (or massless) fermions, with the velocity of light 𝑐 replaced by the Fermi velocity 𝑣𝐹. Moreover, the internal degree of freedom, which is just spin for Dirac fermions, is the sublattice index 𝐴 or 𝐵 in the case of graphene and it is usually referred to as pseudospin. The linear energy dispersion in Eq. (11) leads to the fact that the total density of states is directly proportional to energy and the carrier density is proportional to energy squared. Indeed, 𝜌(𝐸) = 1 𝐿2∑ 𝛿(𝐸 − 𝐸(𝒒)) = ∫ 𝑔𝑠𝑔𝜈 2𝜋𝑞𝑑𝑞 (2𝜋)2 𝛿(𝐸 − 𝐸(𝑞)) = 2|𝐸| 𝜋ℏ2𝑣 𝐹2 𝒒 (13)
which is plotted in panel (c) of Fig. 1.3, where 𝑔𝑠= 2 and 𝑔𝜈= 2 account for spin and valley degeneracies, respectively. The carrier density is given by
𝑛(𝐸) = 1 𝐿2 ∑ 𝑔𝑠𝑔𝜈= 𝑔𝑠𝑔𝜈 𝑞𝐹2 4𝜋= 𝐸2 𝜋ℏ2𝑣 𝐹2 |𝒒|≤𝑞𝐹 (14)
1 Graphene and other 2D materials
11
Fig. 1.3 (a) Three-dimensional band structure of π-bands in tight-binding approximation. Blow up of the energy dispersion relation for graphene around the Dirac point in the Brillouin zone. Adapted from [24]. (b) Schematic of low-energy electronic band structure of graphene. The red (green) balls with rightward (leftward) pointers represent pseudospin up (down) [25]. (c) Density of states per unit cell as a function of energy (in units of 𝑡) and its zoom-in close to the neutrality point. The density of states close to the neutrality point can be approximated by 𝜌(𝐸) ∝ |𝐸|. Adapted from [21].
To find the eigenstates of Dirac Hamiltonian in Eq. (9), it is useful to write it in the term of the polar angle 𝜙𝒒 of the reduced momentum 𝒒
𝐻̂(𝒒) = ℏ𝑣𝐹𝒒 ( 0 𝑒−𝑖𝜙𝒒 𝑒𝑖𝜙𝒒 0 )
(15)
where 𝜙𝒒 = tan−1( 𝑞𝑥
𝑞𝑦). The eigenfunctions are,
(b) (a)
1 Graphene and other 2D materials 12 𝜓𝑒,ℎ𝐾 (𝒒) = 1 √2( 𝑒− 𝑖𝜙𝒒 2 ±𝑒 𝑖𝜙𝒒 2 ) (16)
corresponding to electron (𝑒) and hole (ℎ) states, that are just linear combinations of the states from the sublattices 𝐴 and 𝐵.
A complete low-energy Hamiltonian consists of 4x4 matrices taking into account, besides two sublattices, also two conical points (valley). On the basis that
Ψ = ( 𝜓𝐾𝐴 𝜓𝐾𝐵 𝜓𝐾′𝐴 𝜓𝐾′𝐵) (17)
where e.g. 𝜓𝐾𝐴 (𝜓𝐾′𝐵) labels the component of the electron wave function corresponding to valley 𝑲 (𝑲’) and sublattice 𝐴 (𝐵), the Hamiltonian is a 2x2 block super matrix
𝐻̂ = (𝐻̂𝐾 0 0 𝐻̂𝐾′)
(18)
Next, we are going to find eigenvalues of the helicity operator (a very important feature of Dirac particle) which here is defined as:
ℎ̂ = 𝝈̂ ⋅ 𝒑 |𝒑|
(19)
where 𝒑 = ℏ𝒒 is the electron momentum operator. To do that, it is convenient to exchange the spinor components at the 𝑲 point (for hole states) which means the following change of basis
1 Graphene and other 2D materials 13 Ψ = ( 𝜓𝐾𝐴 𝜓𝐾𝐵 𝜓𝐾′𝐴 𝜓𝐾′𝐵) → Ψ = ( 𝜓𝐾𝐴 𝜓𝐾𝐵 𝜓𝐾′𝐵 −𝜓𝐾′𝐴) (20)
so that the Hamiltonian in Eq. (18) becomes
𝐻𝐾(𝒒) = ℏ𝑣𝐹𝜏𝑧⊗ 𝝈𝒒 (21)
where 𝜏 are Pauli matrices representing the valley degree of freedoms called valley pseudospin. Using Eqs. (19) and (21)
𝐻𝐾(𝒒) = ℏ𝑣𝐹𝒒ℎ̂ (22)
we find that the helicity operator commutes with the Hamiltonian, so the projection of the pseudospin on the direction of momentum 𝒒 is a well-defined conserved quantity. Chirality can be either positive or negative, corresponding to pseudospin and momentum being parallel or antiparallel to each other. It is precisely because pseudospin is locked with momentum that backscattering is prohibited in graphene, making a so high carrier mobility possible. Indeed, for an electron to backscatter (i.e. changing 𝒑 to −𝒑) it needs to reverse its pseudospin, therefore, in absence of a perturbation which flips the pseudospin, backscattering is not possible.
1.2
PROPERTIES OF TRANSITION METAL DICHALCOGENIDES
The lattice structure of single and multilayer MX2 where M is a transition metal atom
(Mo, W, etc.) and X a chalcogen atom (S, Se or Te) is schematically shown in Fig. 1.4. In contrast to planar graphene or monolayer hexagonal boron nitride (h-BN), a monolayer TMD is not single-atom thick since it presents an inner layer of M atoms ordered on a triangular lattice, which is sandwiched between two layers of X atoms located on the triangular net of alternating hollow sites. Therefore, the MX2 crystal
nearest-1 Graphene and other 2D materials
14
neighbor in-plane M-M and X-X atoms and a separation 𝑢 ≃ 𝑎/2 between M and X planes. Qualitatively, it has the same lattice vectors of graphene, but bigger lattice constant 𝑎 (e.g. 3.16 Å for WS2).
Fig. 1.4 (a) Sketch of the atomic structure of MX2. The bulk compound has a 2H–MX2 structure with two
MX2 layers per unit cell, each of them being built up from a trigonal prism coordination unit. (b) top view of monolayer MX2 lattice. Green (blue) circles indicate M (X) atoms. The nearest neighbors (δi) and the next
nearest neighbours (ai) vector are shown in the figure [26].
The in-plane Brillouin zone is shown in panel (a) of Fig. 1.5. It is a hexagon that contains the high symmetry points
𝜞 = (0,0), 𝑴 = (4𝜋 3𝑎, 2𝜋 3√3𝑎), 𝑲 = ( 4𝜋 3𝑎, 0). (23)
The six 𝑸 valleys that correspond to the minimum of the conduction band for multi-layer samples are also shown.
The band structures of single layer WS2 obtained from Density Functional Theory (DFT)
and from tight-binding calculations are shown in panel (b) of Fig. 1.5. Contrary to what happens in other 2D crystals, such as graphene or phosphorene, the valence and conduction bands of TMDs present a very rich orbital contribution. They are made by the hybridization of the 𝑑 orbitals of the transition metal atom M and the 𝑝 orbitals of the chalcogen atom X.
1 Graphene and other 2D materials
15
Fig. 1.5 (a) Representation of two-dimensional Brillouin zone (BZ) of MX2. The high symmetry points Γ, M
and K. The Q points (which are not high symmetry points) indicate the position of the edges of the conduction band in multi-layer samples. (b) Electronic band structure of single-layer MX2 from Density
Functional Theory (DFT) calculations (black lines) and from tight-binding (TB) (red lines) [26]. (c) PL spectra collected from WS2 monolayer (black) and bilayer crystals (red). The shift in the PL emission energy is due
to the shrinking of the band gap in the bilayer crystal. Adapted from [27].
TMDs have an indirect band gap in their bulk form (~ 1.3 eV for WS2), however, when
they are thinned down to a single layer, the band gap becomes direct (~ 1.95 eV for WS2),
with the gap lying at the two inequivalent 𝑲 and 𝑲’ points of the hexagonal Brillouin zone. In multi-layer sample, instead, the maximum of the valence band is placed at the 𝜞 point of the BZ, while the edge of the conduction band is at the 𝑸 point of the BZ.
The presence of a direct gap in monolayer TMDs is experimentally confirmed by their strong photoluminescence (PL) which is due to the recombination of an electron in the conduction band with a hole in the valence band (panel (c) of Fig. 1.5) [27].
(a) (b)
1 Graphene and other 2D materials
16
1.3
VDW HETEROSTRUCTURES AND THEIR APPLICATIONSFollowing the development of many novel 2D materials, investigations of vdWH have attracted significant attention due to their outstanding potential in electronic and optoelectronic applications, including photodetection and the possibility of integrating them with the existing semiconductor technology for the next-generation electronic and sensing devices. Unlike conventional semiconductor heterostructures, layered materials consist of in-plane covalently bonded atomic layers that interact weakly (around 10 meV/Å2) with each other in the out-of-plane direction thanks to the dangling-bond-free
face of 2D materials. The weak vdW forces that hold the layers together allow to readily form any desired heterostack, without the conventional “lattice mismatch” issue. Indeed, in vdWHs each single layer does not change at the atomic scale, so its properties are preserved. What is more, by combining the different properties of the building blocks, synthesized vdWH may show integrated properties with respect to the segregated materials. In particular, 2D materials offer potential for large-scale and low-cost integration into the current dominant optical fibre network and silicon complementary metal-oxide–semiconductor (CMOS) technology.
For example, such heterojunctions have been already adopted to build optoelectronic devices. Graphene photodetectors, indeed, typically have low responsivity, which is a consequence of the material’s low absorption coefficient. Other 2D materials, like TMDs, have shown great potential in nanophotonics due to their efficient light absorption, fast responsivity, wide spectral range and thickness-dependent bandgap [28]. By combining high-mobility graphene with highly photoresponsive TMD in a hybrid vdWH photodetector, high photoresponsivity and fast response times were obtained [29].
Another appealing property of monolayer TMD is the sizeable spin-orbit splitting of the bands together with the lack of space inversion symmetry. As a result, a single valley (and a single spin polarization) can be selectively excited with circularly polarized light. Combining graphene and single-layer TMDs in a vertical vdWH brings together massless Dirac particles with long spin lifetimes and strongly spin polarized electrons with great potential for spintronics. Luo et al. demonstrated, for the first time, the opto-valleytronic spin injection across a TMD/graphene interface, paving the way for multifunctional 2D spintronic devices for memory and logic applications [30].
1 Graphene and other 2D materials
17
Graphene potential has been also exploited in silicon photonics technology for enhancing device integration density and operation bandwidth, reducing cost and power consumption. Wavelength-independent light absorption and ultrahigh carrier mobility, indeed, contribute to overcoming the limitations of silicon photonic devices. Moreover, by integrating graphene with an optical waveguide, which greatly increases the interaction length through the coupling between the evanescent waves and graphene, the problem of the limited absorption of the monolayer graphene can be solved [31]. This implies that high-efficiency, high-speed, broadband photodetectors and modulators can be obtained using the graphene-on-waveguide configuration. Even more appealing performances would be achieved if also monolayer TMDs (e.g. with a TMD/graphene heterostack) were introduced in silicon photonic devices. Refractive index of a TMD
monolayer, indeed, can be highly tuned using complementary
metal−oxide−semiconductor (CMOS) compatible electrical gating. This enables, when combining the TMD with photonic structures, the efficiencies of optical absorption (reflection) to be tuned from 40% (60%) to 80% (20%) [32]. This can overcome the
limitation of today’s integrated photonics by providing electro-optic properties to
traditionally passive optical materials, paving the way toward the development of field-effect photonics in which the optical functionality can be controlled with CMOS circuits.
2D layered materials can be also exploited in advanced 2D tunnelling transistors. Although graphene shows potential in realizing 2D high speed bipolar field effect transistor (FET) devices due to its high carrier mobility and ambipolar effect, the zero band gap limits its application for FET due to low on/off current ratios. When thin isolation barriers (h-BN or TMD) are vertically stacked between two graphene electrodes, quantum tunnelling comes into play. An applied gate voltage effectively modulates the work function of graphene and, therefore, the Schottky-barrier height across the graphene/semiconductor interface, resulting in a large on/off ratio that cannot be achieved in conventional planar graphene transistors [33].
As previously mentioned, TMDs, but first h-BN, have been investigated also as passive materials for encapsulating graphene in order to improve its carrier mobility [7]. H-BN has been shown to act as an ideal substrate by greatly improving graphene transport properties, because of its atomically flat surface, low interlayer electronic coupling, and almost perfect reticular matching. Interesting results were also obtained with exfoliated
1 Graphene and other 2D materials
18
flakes of MoS2 or WS2 used as substrates for graphene (with h-BN used as
2
EXPERIMENTAL
TECHNIQUES
2.1
CVD SYNTHESIS OF GRAPHENE AND TRANSFER ON A SELECTED
SUBSTRATE
Technological applications that take advantage of graphene's extraordinary electronic transport properties require structurally coherent graphene on a large scale (e.g., wafer-scale), or large arrays of graphene flakes positioned with a unique azimuthal orientation on a substrate. Currently, the highest-quality graphene synthesis method that fulfils the requirements of film quality, cost and scalability is graphene CVD on metal (e.g., copper) substrates. It consists on the catalytic decomposition of carbon-containing gases on metal surfaces. The chemistry of graphene growth is straightforward. The carbon source, usually some carbon-containing hydrocarbon (e.g., methane), is flown within a high-temperature CVD growth reactor towards a heated surface. The heat will cause the carbon precursor to decompose into carbon clusters on the metal catalyst substrates, followed by diffusion-aggregation of these clusters to form graphene. The metallic surface behaves as a catalyser by lowering the activation energy for graphene formation. The edge attachment energy for carbon atoms can be lowered even more thanks to the presence of oxygen on the copper (Cu) surface [34].
The complexity of the process lies in how to control the layer uniformity and how to attain single-crystal domain over macroscopically large areas. During the common CVD process for graphene growth, substrate temperature, catalysts, carbon source and gas ratios are all of considerable importance and differing growth conditions can lead to significant differences on the grown samples.
The choice of the metal catalysts has a relevant influence on the as-grown graphene sample: depending on the solubility of carbon in different metals and the strength of metal-carbon bonding, different metal catalysts produce graphene films of different thicknesses and crystalline quality.
During this thesis work, the CVD growth of graphene on Cu foils was performed according to the process developed and optimized in the group and reported by Miseikis et al. in Ref. [35]. Copper is indeed the most suitable metal catalyst to grow single-layer
2 Experimental techniques
20
graphene instead of multilayer graphene due to its ultra-low carbon solubility, low cost and high availability.
Fig. 2.1 (a) temperature profile of a typical four-stage CVD growth process. (I) Temperature ramp-up, (II) annealing, (III) growth, (IV) cool-down [35]. (b) Quartz/graphite sample enclosure. (c) Dark field optical image of a typical graphene flake on copper foil. A grain boundary of polycrystalline copper is also shown.
The 25 μm thick Cu foil used in this work was supplied by Alfa Aesar (purity 99.8%). Morphology of copper surface such as roughness, grain boundary, defects, and impurity particles play a crucial role for forming nucleation seeds of monolayer and multilayer graphene. In order to smooth the Cu surface before growth, electropolishing, thermal annealing and in-situ Cu oxidation were performed. The foil was electropolished in an electrochemical cell made using a commercially available Coplin staining jar as the vessel to keep the foil flat and parallel to the counter electrode, a thicker copper plate as the cathode and an electrolyte solution (25 ml of water, 12.5 ml of phosphoric acid, 12.5 ml of ethanol, 2.5 ml of isopropanol and 0.4 g of urea). The foil was subsequently annealed in an inert atmosphere (i.e., argon) within a 4-inch cold-wall CVD system (Aixtron BM) and
(b) (c)
2 Experimental techniques
21
then synthesized in the same reactor at a pressure of 25 mbar. Methane (CH4) gas was
used as carbon precursors and argon (Ar) as carrier gas. A typical temperature profile of a CVD growth process is shown in panel (a) of Fig. 1.6. The annealing as well as the growth was performed at a temperature as close as possible to the Cu melting point (i.e., growth T ~ 1080°C): under these conditions, Cu surface atoms are highly mobile and this allows for improving the quality of the grown graphene. The annealing time was kept at 10 min in all cases. The gas flow during both the temperature ramp-up and the annealing stages was 1000 sccm. The gas flow rates during growth were typically set to 1 sccm of methane, 100 sccm of hydrogen and 900 sccm of argon. The growth time varied depending on the nucleation density and crystal size, from 20 min to 30 min. After the growth, the chamber was cooled in argon/hydrogen atmosphere to a temperature of 120 °C before introducing the samples to air. To reduce the effective gas flow, the sample was contained in a custom-made enclosure, comprising a quartz disk suspended 6 mm above the sample using graphite spacers (panel (b) of Fig. 2.1) [35].
The electrical properties of CVD graphene cannot be tested in-situ on the conductive metal substrates, thus processes to transfer graphene on an appropriate insulating substrate have been developed. The ability to select the host substrate independently of the sacrificial growth substrate is a major advantage for graphene grown on metals. At the same time, the transfer process often affects negatively graphene's integrity, properties, and performance. Wrinkle formation, impurities, graphene tearing and other structural defects, can occur during transfer.
The first transfer technique reported in literature for graphene grown on Cu foil consisted in a transfer process implemented by reinforcing the graphene layer with a polymer film, e.g., poly(methyl methacrylate) (PMMA), and by subsequently etching off the Cu growth substrate. However, the wet etching step leads to graphene damaging, metal residues production and serious environmental pollution. It also increases production cost and requires long treatment cycles of several hours. A technique that copes with this problem is the so-called bubble transfer, which uses the generation of electrochemically generated hydrogen (H2) bubbles at the graphene-metal interface to
release the polymer-supported graphene film without the need of lengthy under-etching [36], [37]. The interaction between graphene and Cu is relatively weak with a low binding energy of 33 meV per carbon atom [38], [39], comparable to the interplanar coupling
2 Experimental techniques
22
strength of graphite (25 meV per carbon atom) [40]. After release of the polymer-reinforced graphene film from the growth substrate the graphene is contacted with the dielectric device substrate, either using a liquid medium (wet transfer) or a flexible stamp (dry transfer) [41].
Fig. 2.2 (a) Electrolysis cell with the NaOH aqueous solution. Pt foil used as cathode (on the left) and the sample used as anode (on the right). H2 bubbles surroundings the Cu foil are also visible. (b) Transfer
set-up: substrate is placed on a micromechanical metallic stage heated at 120°C and the PMMA membrane carrying graphene flakes is suspended over the substrate thanks to a fixed support. Thanks to an optical microscope it is possible to see when substrate and membrane are aligned. On the computer screen one of the graphene flakes is showed.
In this thesis work, to transfer the CVD graphene on SiO2 a semi-dry transfer process
was implemented [42]. A thin PMMA carrier membrane (AR-P 679.02 Allresist GmbH) was spin-coated on the Cu foil with graphene. Then the PMMA/graphene/Cu was dipped into an NaOH aqueous solution and used as the cathode of an electrolysis cell with a constant current supply (panel (a) of Fig. 2.2). It is worth noting that the graphene would be easily oxidized if the PMMA/graphene/Cu was used as the anode. At the negatively charged cathode, a water reduction reaction took place to produce H2. The reaction can be
represented as follows:
2H2O(l) + 2e- → H2(g) + 2OH-(aq)
These H2 bubbles provide a gentle but persistent force to detach the graphene film
from the Cu foil at its edges, and the process is aided by the permeation of the electrolyte
2 Experimental techniques
23
solution into the interlayers as the edges delaminate. The PMMA/graphene layer was detached from the Cu substrate after a few minutes as a result of the formation of a large number of H2 bubbles at the interface between the graphene and Cu substrate. After
rinsing with deionized water, the floating PMMA/graphene layer was aligned to the target substrates using a micromechanical stage (panel (b) of Fig. 2.2). The target substrates (300 nm of SiO2 layer on p-doped Si) were mildly heated to improve the adhesion of graphene,
and the PMMA was finally removed in acetone (ACE) and isopropanol (IPA). In order to clean better the sample from polymer residues, it was dipped in PMMA remover (AR-P 600-71 Allresist GmbH) for 3 minutes and then rinsed in deionized water.
2.2
CVD SYNTHESIS OF TUNGSTEN DISULPHIDE
Considerable efforts have been made in the past years in developing controllable synthesis of 2D semiconducting TMDs. Several routes, including sulfurization (or selenization) of metal (or metal oxides) thin films [43], physical vapor phase transport [44] and CVD [45], have been developed. Among all these methods, CVD shows great promise towards high-quality TMD production [46]. With this technique, monolayer TMDs that have single domains with scalable size, controllable thickness and excellent electronic properties can be obtained. In a typical CVD process for TMD growth, the substrate is exposed to volatile precursors that react in vapor phase, by which the desired layered materials can be formed. Besides standard CVD, also a metal organic chemical vapor deposition (MOCVD) method has been developed [47]. The major difference is from the precursors, where gaseous reactants with a high equilibrium vapor pressure around room temperature are selected.
In this work WS2 was grown by implementing CVD synthesis from solid precursors. As
said above, this approach relies on the reaction of sulfur (S) vapour with a W-precursor (in this case tungsten trioxide WO3) at elevated temperatures within a horizontal quartz
tube reactor. Not having individual control over each precursor's temperature limits the parameter space of the chemical reaction. To overcome this limitation, S was placed in the upstream zone of the furnace where a resistive heating belt wrapped around the tube separately control its temperature. Both WO3 and the growth substrates were positioned
2 Experimental techniques
24
powder in the low-temperature zone was heated and carried downstream to the reaction zone using an argon flux. This leads to vapor-phase reaction that involves a redox and a substitution process. Starting from W and S, the reaction is the following:
WO3 + S → WO(3-x) + SOx
WO(3-x) + SOx → WS2
The SOx is volatile and pumped away while the WS2 molecules form the layer on the
substrate.
Fig. 2.1 Sketch of the furnace for the CVD growth of WS2. S is placed inside the heating belt, whereas the
WO3 powder is placed next to the substrate downstream in reference to the direction of argon flux.
In general, the growth mechanism of TMDs via CVD depends on some important parameters. The partial pressures of the evaporated precursors have to be kept stable and high enough to enable mixing of atomic gases and transport the atomic species to the substrate. Growth temperature is also fundamental: too low temperatures do not allow adatoms to have enough kinetic energy to diffuse and find the lowest potential energy site and an amorphous film will form [48]. The substrate surface morphology also influences the atomic layer of 2D materials: it was reported that the surface energy of the substrate affects the nucleation and growth of 2D TMDs [49], indeed the same process performed on different substrates leads to very different results. Lastly, an increase of the growth time results in a higher coverage [50].
The process used in this thesis work to grow WS2 on CVD graphene will be detailed in
2 Experimental techniques
25
2.3
RAMAN SPECTROSCOPY AND PHOTOLUMINESCENCE
Raman spectroscopy is a spectroscopic technique fundamental for characterization of 2D materials because it gives important information on crystal quality, number of layers, doping and strain. Furthermore, Raman spectroscopy is quick, convenient, non-destructive and non-invasive. When monochromatic light, usually from a laser source, is inelastically scattered in a material, the energy of the scattered light either decreases by exciting an elementary excitation of the solid material, (i.e. a phonon) or increases by absorbing it. Raman spectra give the intensity of the scattered light as a function of the energy shift from the incident light (Raman shift). This shift provides information about vibrational and other low frequency transitions in both crystals and molecules.
Fig. 2.3 Scheme of the Raman active transitions. Monochromatic light with energy ℎ𝜈0 is inelastically
Stokes (Anti-Stokes) scattered releasing (absorbing) an energy ℎΔ𝜈 to (from) the system. Adapted from [51].
Actually, about 99.999% of all incident photons in spontaneous Raman undergo elastic Rayleigh scattering. However, if the system has Raman-active modes, Raman scattering occurs and it can be classified as two types, Stokes Raman scattering and anti-Stokes Raman scattering. In the first one, an electron is excited from the ground level and falls into an excited vibrational level, releasing energy to the system. Conversely, the anti-Stokes scattering consists in the excitation of an electron that is already in the excited
2 Experimental techniques
26
vibrational state. When the electron returns to its ground state, excessive energy is released and the resulting scattered light has higher energy than incident light.
When a sample is illuminated by a laser, beside Raman scattering, photoluminescence (PL) can occur. When light with energy greater than the energy gap is shone on the sample, the charged carriers will occupy an excited state adsorbing a photon and subsequently relax on the ground state emitting another photon. The PL intensity and the energy of the emitted photons give important information about the material.
In this work all Raman and PL measurements were carried on with a Renishaw inVia system equipped with a 532 nm green laser.
2.3.1 RAMAN ON GRAPHENE
The phonon dispersion of graphene (panel (a) of Fig. 2.4) plays a key role in interpreting its Raman spectra. Graphene has two atoms per unit cell that, moving in the three-dimensional space, generate six phonon branches, three acoustic (A) and three optical (O), with in-phase and out-of-phase displacements of the two unit-cell atoms, respectively. The optical phonons in the zone-centre (𝛤) and zone edge (𝐾 and 𝐾’) region are of particular interest, since they are accessible by Raman spectroscopy.
Fig. 2.4 (a) Phonon dispersion of graphene in the BZ. Adapted from [52]. (b) Eigenvectors for the phonons at the high-symmetry Γ and 𝐾(𝐾’) points of the BZ. Each of these twelve modes is labelled and their atom displacements are indicated. The other symbols are symmetry assignments according to group theory [53].
2 Experimental techniques
27
Panel (b) of Fig. 2.4 shows the twelve phonon eigenvectors at the 𝛤 e 𝐾 points, that are described by the acoustic and optical unit-cell displacements multiplied by a wavevector 𝒒, which defines the phase modulation along the unit cells in the crystalline lattice. The atomic displacements in the graphene plane (i for in-plane) can be longitudinal (L) or transversal (T) with respect to the phonon wavevector direction. The atomic displacements perpendicular to the graphene plane (o for out-of-plane) are transversal (T) phonons, and they generally exhibit lower frequencies because the out-of-plane restoring forces are much weaker than the in-plane ones.
Fig. 2.5 Sketch of the main Raman processes in graphene. (a) G band, (b) D band double resonant process involving a scattering from a defect and (c) 2D band generated through a second-order process that is either double resonant or (d) triple resonant. Resonance points are shown as dashed circles.
The in-plane longitudinal optical (iLO) and the in-plane transversal optical (iTO) modes correspond to the vibrations of the sublattice 𝐴 against the sublattice 𝐵, and these modes are degenerate at the 𝛤 , representing the only first-order Raman active mode in graphene, named G band. The relatively high frequency of this optical phonon (~ 0.2 eV) allows the use of Raman spectroscopy to probe small environmental perturbations, including variation in strain [54] and doping [55]. On the other hand, the 2D and D bands arise from a second-order process, involving two iTO phonons near the 𝐾 point for the 2D band or one iTO phonon and one defect in the case of the D band. The D band comes from a totally symmetric vibration mode (breathing mode) occurring near the corners (𝐾 or 𝐾’
(c)
(a) (b)
2 Experimental techniques
28
points) of the first Brillouin zone. This mode belongs to the iTO phonon branch [56] and its Raman activity is mediated by a double-resonance mechanism in which the photoexcited electron/hole is inelastically scattered by a phonon and elastically backscattered by a defect that provides momentum conservation in the process, as shown in panel (b) of Fig. 2.5. Because the phonon scatters an electron/hole from one Dirac cone to another, this process is called intervalley. The overtone of the phonon modes that gives rise to the D band originates the 2D band. Since the electron/hole is scattered twice by the same phonon (with opposite wavevectors), the 2D do not require the defects to be observed.
Considering laser excitation at 532 nm, the one that has been used in this work, the position of the 2D peak is centred at ~2700 cm-1 and the G peak at ~1584 cm-1. In the
case of a disordered sample, we can also see the disorder-induced D band, at about half of the frequency of the G band (around 1350 cm−1). A typical graphene Raman spectrum
of graphene is shown in Fig. 2.6.
Thanks to Raman spectroscopy a number of information about graphene properties can be extracted: position and full width at half maximum (FWHM) of the G and 2D peak depend on the quality of the material and its number of layers, but are also affected by the presence of strain [54], [57], [58] or doping [59].
2 Experimental techniques
29 2.3.2 RAMAN ON WS2
Monolayer WS2 has three atoms in its unit cell and their vibrations result in nine normal
modes, including three acoustic and six optical branches: one each for LA, iTA, oTA and two each for LO, iTO, oTO. Panel (a) of Fig. 2.7 illustrates the two types of optical phonons at the Brillouin zone centre (the Γ point): one type involves only relative motion of the two chalcogen atoms, and the other type is participated by both the transition metal and the chalcogen atoms.
Raman spectroscopy has been proved to be a good technique also to investigate TMDs, in particular it is a useful tool to evaluate the number of layers [60]–[62].
The Raman active modes of WS2 are the A1g mode, corresponding to the out-of-plane
vibration that involves only the chalcogen, and the E2g1 mode, corresponding to the
in-plane displacement in opposite directions of both transition metals and chalcogen atoms (panel (a) of Fig. 2.7). Besides these two first-order modes at the Brillouin zone centre, there is a zone-edge mode, activated by disorder, which is a second-order Raman mode due to LA phonons at the M point, identified as 2LA(M). For the laser excitation used in this work (i.e. 532 nm), the E2g1 peak is centred at ~ 355 cm-1 and is reported to be
significantly strong in intensity compared to the A1g (at ~ 417 cm-1) [62], furthermore it
partially overlaps with the 2LA(M) peak that is centred at ~ 350 cm-1, making quite
difficult to identify the right position of the two with a proper fit.
Fig. 2.7 (a) Optical phonons at the Brillouin zone centre of monolayer TMDs (metal atom in orange and chalcogen atoms in green). Atomic displacement 2) corresponds to the E2g1 Raman active mode, while 3)
to the A1g mode. (b) WS2 Raman spectrum taken with 514 nm excitation light. All the main peaks and their
overtone are highlighted. Adapted from [60].
2 Experimental techniques
30
Several ways to extrapolate the thickness of the WS2 crystal looking at Raman
behaviour have been studied in recent years. Berkdemir et al. reported that the 2LA mode increase in intensity with decreasing the number of layers and reach a maximum for the monolayer, while the intensity of the A1g mode decrease [60]. As indicated in [60], we
studied the I2LA/IA1g ratio to determine the thickness of our WS2 flakes. In the same work,
they also reported that the A1g mode blueshifts when increasing the number of layers,
consistent with the increasing restoring force caused by van der Waals interactions established among layers. This is in agreement with previous results reported for MoS2
[63]. The E2g1 and 2LA phonon modes, instead, exhibit very subtle redshifts when
increasing the number of layers. Similar anomalous behaviour of the E2g1 mode has been
reported in few-layered MoS2 films [64] and might be caused by stronger dielectric
screening of the long-range Coulomb interactions between the effective charges in thicker samples.
2.4
ATOMIC FORCE MICROSCOPY
Atomic force microscopy (AFM) is a high-resolution type of scanning probe microscopy (SPM), with vertical resolution in the order of fractions of a nanometer.
The AFM consists of a cantilever with a sharp tip (probe) at its end that is used to scan the specimen surface. The cantilever is typically silicon or silicon nitride with a tip radius of curvature on the order of nanometers. When the tip is brought into proximity of a sample surface, vdW forces between the tip and the sample lead to a deflection of the cantilever according to Hooke's law. The vertical and lateral deflections of the cantilever are measured through a laser that is reflected off the back end of the cantilever and directed towards a position sensitive detector (panel (a) of Fig. 2.8).
AFM can work in two main modes, contact and tapping. In contact mode, the sample is moved under the tip and the contours of the surface are measured either using the deflection of the cantilever directly or, more commonly, using the feedback signal required to keep the cantilever at a constant position. In tapping mode, the cantilever is driven to oscillate up and down at or near its resonance frequency. This oscillation is commonly achieved with a small piezo element in the cantilever holder. When the morphology of the substrate changes, the vibration frequency changes accordingly, since
2 Experimental techniques
31
it depends on the distance from the substrate. The phase of the oscillation changes when changing the probed material.
In this work all the AFM measurements were performed with with a Bruker Dimension Icon microscope used in ScanAsyst tapping mode.
Fig. 2.8 (a) Schematic illustration of the principles of AFM. The scanner is composed of three piezo components, which control the horizontal (x and y) and vertical (z) movement of the sample. (b) Qualitative profile of the Lennard-Johns potential that describe the interaction between tip and sample. In contact mode, the force on the tip is repulsive and, by maintaining a constant cantilever deflection, the force between the probe and the sample remains constant and an image of the surface is obtained. In tapping mode, the cantilever is oscillated at its resonant frequency. By maintaining a constant oscillation amplitude, a constant tip-sample interaction is maintained and an image of the surface is obtained.
2.5
SCANNING ELECTRON MICROSCOPY
A scanning electron microscope (SEM) is a type of electron microscope that produces images of a sample by scanning the surface with a focused beam of electrons. The electrons interact with atoms in the sample, producing secondary electrons, backscattered electrons as well as X-rays that contain information about the sample's surface topography and composition. The electron beam is scanned in a raster scan pattern, and the beam's position is combined with the detected signal to produce an image. SEM can achieve resolution better than 1 nanometer.
The electrons are emitted from a field emission gun by placing a cathode filament at a huge electrical potential gradient, so large that the work function of the material is overcome and electrons are drawn off of the filament.The electrons are then accelerated and the beam is focused on the specimen by electrostatic lens. The impinging electron
(b) (a)
2 Experimental techniques
32
beam penetrates the substrate in a pear–like shaped volume (panel (a) of Fig. 2.9), whose size depends on the incoming electron energy, atomic number and density of the specimen. The image is reconstructed collecting the outcoming electrons that can be distinguished according to the different kind of interaction that they experience.
Fig. 2.9 (a) Pear-like interaction volume and interaction effects between the impinging electron beam and the sample. Sketches of the origin of secondary electrons and backscattering electrons are also displayed. (b) Schematic drawing of the SEM illustrating placement of electron generation, collimation process, sample interaction and electron detection.
The two principal detection modes are backscattering and secondary electrons. Backscattered electrons experience elastic scattering with the nucleus. The number of backscattered electrons emitted from a specimen increases with an increasing atomic number, therefore different materials can be easily distinguished. These electrons have an energy range from 50 eV (electron volts) up to nearly the incident beam energy. Elastic scattering can change the path of the primary electron up to 180 degrees: collecting them at different angles allows also reconstructing a topological image of the system. Secondary electron are electrons emitted by core levels of the investigated material,
(a)
2 Experimental techniques
33
originated by inelastic scattering with the incident beam. The energy range for secondary electrons states that the electron must have an energy greater than the 0 eV but less than 50 eV. This energy range limits the escape depth of the electrons. As their energy is < 50 eV, they come from the first few nm of the material and deliver information about the topography of the surface.
In this work SEM measurements were performed at 5 keV using a Zeiss Merlin microscope, equipped with a field emission gun.
3
SCALABLE
SYNTHESIS
OF
GRAPHENE/WS
2VD
WH
3.1
STATE OF THE ART
As discussed in the introduction, chemical vapor deposition (CVD) has emerged as the most effective and reproducible way to obtain high quality, transferrable, size-and thickness-controllable graphene, which is industrially compatible. Miseikis et al. demonstrated a fast and high-controllable method to grow high-mobility large-crystal graphene in deterministic positions utilizing CVD on patterned copper substrates [42], providing a significant advance for adoption of graphene in wafer-scale fabrication.
Most recently, CVD has been also successfully used for the synthesis of WS2 film at
large scale reaching flake sizes of hundreds of micrometers. Gutierrez et al. and Cong et al. reported the synthesis of single- and few-layered 2D triangular microplatelets of WS2
via the sulfurization of ultrathin WO3 films deposited on Si/SiO2 substrates [61], [65].
Similar results were achieved by several groups, which deposited WS2 by direct reaction
of gaseous precursors on the substrate surface, rather than by substrate metallization and subsequent sulfurization: this offers important advantages in terms of process scalability [45], [66], [67].
Remarkably, vapor phase deposition can also be employed to synthesize multilayer vdWH, which pave the way for exploring new devices and physics based on 2D materials. Indeed, thanks to the possibility of combining a wide range of 2D building blocks in a scalable way, most of the critical components in modern electronics can be redesigned and fabricated.
With regard to the heterostack studied in this thesis, the first WS2/graphene
heterostructures were made in literature by growing the two materials separately via CVD and by then transferring CVD WS2 on CVD graphene [68]. Using this approach, Tan et al.
demonstrated lateral Gr−WS2−Gr photodetectors exhibiting high photoresponse (∼ 3.5
A/W) under ambient atmosphere and high input visible light at 532 nm [68].
However, staking of 2D materials by transfer methods often results in impurities trapped at the heterojunction interface [69], which can act as scattering centres, strongly affecting transport and optical properties of atomic-layered materials.
3 Scalable synthesis of graphene/WS2 vdWH
35
The synthesis of vdWH, performed growing one material directly on top of the other, can provide cleaner interfaces. Rossi et al. managed to synthesize via CVD continuous atomic-thick WS2 films directly on 2D epitaxial graphene (EG) obtained on silicon carbide
and subsequently demonstrated an entirely scalable WS2/EG photodetector with a
maximum photoresponsivity R ~ 220 A/W under continuous-wave illumination [17]. Although this result represents an important step forward the implementation of scalable photoresponsive devices based on 2D materials, SiC-based technologies are geared for a niche market (which can justify the cost).
One challenge in fabricating scalable optoelectronics devices toward the integration on a silicon photonics platform is growing photoactive 2D materials, such as TMDs, directly on CVD graphene. As mentioned above, this will allow to get the heterostack on top of any substrate of interest. First results in this direction were provided by Bianco et al., who showed how to deposit atomic layers of WS2 on CVD poly-crystalline graphene
(supported on SiO2/Si substrate) using elemental sulfur (S) and tungsten hexacarbonyl
(W(CO)6) as precursors [70]. They obtained flakes with dimensions of a few hundreds of
nanometers, grown mainly on bilayer graphene. Monolayer graphene was almost uncovered by WS2, with exception of wrinkles. Moreover, extending the growth time, they
observed a thickening of the already present flakes, instead of an increase of nucleation of new material on monolayer graphene.
Therefore, a process for the homogeneous growth of WS2 on CVD graphene has to be
developed, investigating the effect of the most important parameters on the controlled growth, including the choice of precursors, tube pressure, growing temperature, holding time, the amount of sulfur powder, and gas flow rate.
To the best of our knowledge, this thesis work is the first study that shows how to deposit atomic layers of WS2 on CVD single-crystal graphene. In fact, the absence of grain
boundaries, that are sites of preferable adsorption of external species, reduces appreciably the scattering of charge carriers. In this way, photoactive layered heterostructures would be built in a scalable manner, while retaining high electrical mobility.
3 Scalable synthesis of graphene/WS2 vdWH
36
3.2
DIRECT CVD OF TUNGSTEN DISULPHIDE ON CVD GRAPHENE
To synthesize WS2 directly on CVD graphene/SiO2 substrate, the CVD approach was
adopted, using as precursors tungsten WO3 (Sigma Aldrich, 99.995%) and S (Sigma Aldrich,
99.998%) powders in a 1:50 ratio (3 mg of WO3 and 150 mg of S). The process was
performed within a 2.5-inches horizontal hot-wall furnace (Lenton PTF) (Fig. 3.1).
Fig. 3.1 Lenton furnace used to synthetize WS2 with highlighted components. S powder is placed inside the
heating belt, whereas the WO3 powder is placed upstream next to the substrate.
The furnace comprises a central hot-zone, where a crucible with the WO3 powder was
placed 2 cm distant from the growth substrate, and an inlet zone heated by a resistive belt in which the S powder was positioned. After the chamber was pumped down to a pressure of ~ 5 x 10-2 mbar, the temperature ramp-up was started with a rate of 5°C/min,
paying attention to increase the pressure inside the chamber to a value high enough (4.6 mbar) to keep the sulfur solid. To do that, a flux of 500 sccm of Ar was flowed during the temperature ramp-up. During the process the belt temperature was set to 200°C to evaporate sulfur, while the temperature within the reaction zone was set to 900 °C. After reaching 900°C, the Ar flux was suddenly reduced to 8 sccm, which leads the furnace pressure to drop immediately to 0.6 mbar. In these conditions sulfur starts to evaporate and to sulfurize WO3. The growth time was set to 1 hr. Finally, the furnace was naturally
cooled down to room temperature and the sample was removed from the tube.
Panels (b)-(d) of Fig. 3.2 report representative SEM images obtained when growing with parameters of pressure, temperature and time mentioned above. First of all, it can
3 Scalable synthesis of graphene/WS2 vdWH
37
be seen that WS2 is grown preferentially on graphene, which appears almost fully covered
(panels (a)-(b) in Fig.3.2). A single crystalline triangular flake is clearly visible (with dark contrast) in the zoomed image in panel (d): each WS2 crystal has dimensions of ~ 300 nm.
The WS2 crystals grow with a preferential orientation or rotated by 60°, reasonably
indicating the most stable stacking with graphene. Additional WS2 nanostructures within
the triangular flakes are clearly observed (especially in panel (d)). The nanostructures density is higher at the centre of each flake, most likely at the nucleation point of the WS2
flake. They have an average lateral dimension of ~10 nm and the average height is around 2-3 nm, as measured via AFM. A high crowding of these nanocrystals is also present along graphene wrinkles (as seen in panel (c)), which indicates that the morphology of the underlying graphene layer strongly affects WS2 nucleation.
Fig. 3.2 (a) Raman map of the A1g peak intensity and (b) SEM image taken at the edge of a graphene flake.
While on SiO2 there is no evidence of WS2, graphene is almost entirely covered. (c) Graphene wrinkling
provides the nucleation centres for upper layer material growth. This SEM image is taken on a less covered graphene area to show better this phenomenon. (d) Close up of a WS2 triangular flake (in black).
Nanostructures (in light grey) are grown radially starting from the centre of the flake.
(a) (b)
3 Scalable synthesis of graphene/WS2 vdWH
38
In order to determine the number of layers of the material grown, Raman spectroscopy was performed. Laser spot has a size of about 1 µm, therefore it is important to bear in mind that all the collected data are the result of signals averaged.
Fig. 3.3 (a) Raman spectrum of the heterostructure in the range of WS2 and Si peaks by using 532 nm laser
excitation. (b)-(d) Histograms of the intensity ratio between the 2LA and the A1g peak, the A1g peak position
and the intensity ratio between A1g and Si TO peaks. Data were taken on the same 15x15 µm area of the
heterostructure.
Panel (a) of Fig. 3.3 reports the typical WS2 Raman spectrum measured on the
synthesized sample with the representative WS2 peak discussed in Section 2.3.2. As visible
in the histogram reported in panel (b) an intensity ratio between the 2LA and the A1g larger
than 2 was retrieved, which is indicative of monolayer WS2 [60]. The position of the A1g
peak, which is expected to sensibly blueshift when increasing the number of layers (as discussed in section 2.3.2), was also determined. Panel (c) shows that the A1g peak
position value is consistent with that reported by Berkdemir et al. for single-layer WS2 (~
(a) (b)
3 Scalable synthesis of graphene/WS2 vdWH
39
417.2 cm-2). In agreement with Gutiérrez et al. [61], an intensity ratio I
A1g /ISi TO < 1 was
observed, further corroborating the monolayer thickness of our flakes (panel (d)). The strongest evidence of the monolayer nature of the grown material came from photoluminescence spectra. The band gap structure of TMDs is highly dependent on their thickness. Since monolayer WS2 is a direct band gap semiconductor, only strong
direct-transition (DT) emission located at ∼ 625 nm (1.95 eV) can be observed in its PL spectrum. As the WS2 layer number increases, indirect-transition (IT) emissions should show up in
the higher wavelength side of the DT ones [71]. In Fig. 3.4 is displayed the PL spectrum of the grown material: a strong photoluminescence (PL) response is visible (PL intensity is more than 7 times higher than Si peak intensity, therefore more than 9 times higher than A1g peak intensity) and there is no sign of additional peaks relative to transitions at lower
energy.
Fig. 3.4 Photoluminescence of WS2 (laser excitation 532 nm). The presence of a single peak centred to the
wavelength corresponding to the WS2 bandgap confirms that monolayer material is grown.
For what concerns the additional nanostructures, there was no doubt that they are WS2 nanocrystal: Raman spectra indeed confirmed the absence of other materials (e.g.
unreacted tungsten oxides). Owing to their bulk nature, WS2 nanocrystals do not
contribute to PL radiative recombination. Moreover, also the Raman signal that they generate, is so low [60] that this is completely covered by monolayer signal. The lower
