Strain engineering of Tungsten Disulfide with Polymeric Micrometric Artificial Muscles

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UNIVERSITÀ DI PISA Dipartimento di fisica

corso di laurea magistrale in fisica

Andrea Morandi

STRAIN ENGINEERING OF WS2 WITH POLYMERIC

MICROMETRIC ARTIFICIAL MUSCLES

Tesi magistrale

Relatore:

Prof. S. Roddaro Prof. A. Tredicucci

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Chapter 1 Introduction

The isolation of graphene by mechanical exfoliation in 2004 by Novoselov et al. (awarded with the Nobel prize in 2010) [1] opened the door to a whole new eld of research: two-dimensional materials. Nowadays, be-yond graphene, there exist a host of two-dimensional materials with a wide range of dierent properties [2]; each of them is characterized by strong in-plane bonds and weak inter-layer van der Waals interactions which give them a layered structure. Over the years, many dierent tech-niques have been developed to produce two-dimensional materials; they can be grouped in two classes: in "top down" approaches, atomically thin layers are extracted by exfoliation, in "bottom up" method, they are grown starting from their chemical components. These materials are expected to complement graphene in applications for which it does not possess the optimal properties; for example, graphene's lack of band-gap hampers its application for digital electronics and photo-detection. Among all these, the group called "transition metal dichalcogenides" (TMDs) combines the incredible thinness of graphene with exceptional semiconducting proper-ties [3]. Unlike graphene a TMD monolayer is three atoms thick; each sheet consists of a layer of transition metal atom (Mo,W..) between two planes of chalcogen atoms (S,Se..). Furthermore they present other in-triguing physical properties:

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Figure 1.1: Transition-metal and chalcogen position in the periodic table. Adapted from [5].

• Direct band gap, so that can be used in electronics as transistors and in optics as emitters and detectors;

• No inversion center, which allows to access a new degree of free-dom of charge carriers: the K-valley index (often reered as pseu-dospin)[4];

• Large exciton binding energy, that allows room-temperature photo-luminescence;

• Strong spin-orbit coupling that leads to a spin-orbit splitting of hun-dreds meV in the valence band and a few meV in the conduction band.

The most studied TMD, that is molibdenum disulde (MoS2), has

al-ready been successfully used to to implement logic circuits, ampliers and photodetectors [6]. Apart from their outstanding electrical and optical performances, atomically thin semiconductors have also shown very inter-esting mechanical properties. For example while silicon typically breaks

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7

Figure 1.2: View of a bulk TMD and zoomed region of single layer. Adapted from [10].

at strain levels of ∼1.5%, MoS2(and also WS2) does not break until ∼11%

of strain levels [7][8]. This outstanding stretchability of 2D crystals can be used to realize devices with electronic properties that are tuned through the introduction of mechanical deformations, thus opening new promising opportunities for the eld of strain engineering. While MoS2 has received

the most attention due to its relative ease of mechanical exfoliation, there are many other TMDs that have intriguing properties to study: this the-sis work has been focused on tungsten disulde (WS2). This material

has demonstrated superior optical properties compared to MoS2 [9] and

thanks to its large spin-orbit coupling, can have interesting applications in optoelectronics and spintronics.

In this thesis work I demonstrate a new platform for strain engineer-ing that consists of an heterostructure of bilayer graphene and tungsten disulde and exploits Polymeric Micrometric Articial Muscles (MAMs) as actuators. Bilayer graphene is grown via thermal decomposition of Silicon Carbide and the tungsten disulde is grown directly on top of graphene via Chemical Vapour Deposition, avoiding transfer procedures

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that complicate the fabrication process. Heterostructures of 2D materials have recently shown a regime of lateral motion with ultra-low friction, called Superlubricity, that allows an easy lateral sliding of the akes [11]. The key idea behind this thesis work is to exploit this phenomena with WS2 on top of graphene to create a non suspended set-up in which

tung-sten disulde can be strained at will, responding to lateral forces like a free-standing membrane. The polymeric actuators are implemented via a PMMA layer that mechanically contracts when stimulated by electron radiation, via a phenomena call cross-linking, mimicking the response of a muscle to an electrical stimulus. The strain induced in the WS2 is

then characterized by photoluminescence (PL) and Raman spectroscopy, and the eect on the material are studied with Atomic Force (AFM) and Scanning Electron Microscopy (SEM).

The thesis is organized as follows: In Chapter 2 I will briey introduce transition metal dichalcogenides, focusing on tungsten disulde, describ-ing the crystal structure and the dierent growth processes. Among all techniques, chemical vapour deposition and thermal decomposition of sili-con carbide (SiC) are discussed in detail, since they have been used in this master thesis work. Chapter 3 is dedicated to TMDs' properties, rstly optical and mechanical, then discussing how the strain eects the photo-luminescence emission and nally introducing Superlubricity. In Chapter 4 the state-of-the-art of strain engineering in two dimensional materials is reviewed, pointing up the benets and drawbacks of each technique. Chapter 5 is dedicated to the description of the experimental techniques used. In Chapter 6 there is the description of the original work done in this master thesis: I will discuss the methods used and the results, discussing both the advantages and the drawbacks of this technique in comparison with the others used in literature. Finally in Chapter 7 there are the con-clusion summarized, followed by a discussion of the perspectives opened.

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Chapter 2 TMDs structure and growth

2.1 Crystal structure

Figure 2.1: WS2 structure and top view of the monolayer.

Transition metal dichalcogenides are a class of materials with the for-mula MX2 where M is a transition metal element from group IV (Ti, Zr,

Hf...) group V (Nb or Ta) or group VI (Mo, W and so on), and X is a chalcogen (S, Se or Te). These materials have structures of the form XMX, with thecChalcogen atoms in two hexagonal planes separated by a plane of metal atoms. They exist in bulk form as stacks of strongly bonded layers with weak interlayer attraction. Monolayer TMDs exhibit

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Figure 2.2: Schematics of the structural polytypes: 2H (hexagonal sym-metry, two layers per repeat unit, trigonal prismatic coordination), 3R (rhombohedral symmetry, three layers per repeat unit, trigonal prismatic coordination) and 1T (tetragonal symmetry, one layer per repeat unit, octahedral coordination). Adapted from [12].

only two polymorphic forms, distinguished by metal atom coordination: trigonal prismatic (1H) and octahedral (1T) phases. Depending on the combination of the metal and chalcogen elements, one of the two coordi-nation modes is thermodynamically preferred. Tungsten disulde usually is found in 1H-form (that displays semiconducting properties) while 1T-form is less stable and exhibits metallic behavior [13]. In bulk TMDs, instead, a variety of polytypes are found to occur, which dier in terms of stacking orders. The overall symmetry can be hexagonal, rhombohe-dral or tetragonal as depicted in Figure 2.2 [7]. Similarly to graphene, the in-plane Brillouin zone (BZ) is an hexagon characterized by the

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high-2.2. TOP-DOWN SYNTHESIS 11 symmetry points (see Figure 2.3):

Γ = (0; 0), K = 4π 3a(1, 0), M = 4π 3a(0, √ 3/2)

where a is the in-plane M-M or X-X distance and varies between 3.1 A◦ and 3.7 , depending on the size of the metal and chalcogen ions. For WS2

a ≈ 3.16

◦

A[14]. TMDs have shown interesting layer-dependent properties; above all in many semiconducting TMDs there is a transition from an indirect band-gap in the bulk to a direct band-gap in the monolayer. In fact, decreasing the number of layers (i.e. going from bulk to monolayer), these materials reveal a progressive connement-induced shift in the in-direct gap. On the contrary, the change in the in-direct gap is signicantly small; as a consequence the WS2 crystals exhibit a crossover from indirect

to direct gap in the monolayer limit [15][16]. Density functional theory calculations can give an explanation of this dierent behavior in Γ and K point of the band structure: conduction-band states at K-point are mainly due to localized d-orbitals on the W atoms, that are relatively unaected by interlayer coupling since are located in the middle of S-W-S sandwich; states near Γ point are, instead, given by a combination of pz-orbitals on

the S atoms and d orbitals on W atoms, so they're aected more by in-terlayer forces [15][5]. The resulting band structure for WS2 is depicted in

Figure 2.4. The band-gap value for this monolayer, settled to 2.1 eV, does not consider excitonic eect. In fact, the value of this bandgap is higher for the monolayer (∼ 2.4 eV); the optical emission frequency is lowered due to the high energy binding of the excitons (∼ 300 eV). In chapter 3, a detailed description of this phenomenology is reported, together with a discussion on the methods to calculate the correct band gap [17].

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Figure 2.3: Brillouin zone of WS2 showing six valleys and opposite

spinorbit splitting of the valence band at the K and K' (-K) points. The red and blue surfaces represent spinorbit split valence band maxima, each of which is associated with a particular electron spin. The green surfaces represent the conduction band minima or the valleys. Adapted from [5].

2.2 Top-down synthesis

2.2.1 Mechanical exfoliation

Reliable production of atomically thin 2D-TMDs is essential for translat-ing their new electronic and optical properties into application. As already anticipated, I will show the most relevant fabrication methods reported in the literature. The rst one that was demonstrated is mechanical exfolia-tion [1]: it takes advantage of the weak interlayer interacexfolia-tion and exploits peeling of the surface with a sticky tape. It basically consists of repeat-edly cleaving a bulk layered material to generate step by step sheets of a small thickness. This method is commonly known as Scotch tape method and, despite its simplicity, it has proven to be a powerful technique to obtain high-quality two-dimensional sheets. Novoselov's pioneering work demonstrated in 2005 [2] that it can produce high-purity akes of 2D ma-terials (graphene, boron nitride, dichalcogenides, etc..) from few-layers down to monolayer thick, which is typically the desirable one for funda-mental characterization. This method has beneted either from several studies that have demonstrated that the number of layers of a TMD can

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2.2. TOP-DOWN SYNTHESIS 13

Figure 2.4: Band structures calculated from rst-principles density func-tional theory (DFT) for bulk and monolayer WS2. The arrows indicate

the fundamental bandgap (direct or indirect). The top of the valence band (blue) and bottom of the conduction band (green) are highlighted. Adapted from [12]

be identied simply by optical contrast [18]. Unfortunately this method lacks of control on the number of the deposited layers: only small fraction of the transferred material is eectively monolayer. A further drawback of the method is that it is not scalable, so it is only appealing for the fabrication of individual prototypes but not for large scale production.

2.2.2 Liquid-phase exfoliation

An alternative top-down method is exfoliating bulk layered materials im-mersed in a liquid medium via dierent approaches [10]. I will dedicate briey to this method since it produces large quantities of atomically thin crystals but with small lateral size of the akes (0.2-1 µm), thus it's not

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Figure 2.5: Simple representation of mechanical exfoliation or Scotch tape method.

appealing for strain engineering studies. Liquid-phase exfoliation exploits various physico-chemical procedures like direct sonication in a solvent or intercalation/expansion of ionic species like lithium (recently employing an electrochemical cell) [19] [20]; besides, some of these techniques, pro-duce 1T-MX2 structurally dierent from usual 2H-TMDs and require an

annealing at 300◦_{C to restore the metal atom coordination [21][13]. This}

approach, however, remains intriguing for its low cost and could be prefer-able in applications where large quantities of materials are required, such as electrochemical energy storage, catalysis or sensing.

2.3 Bottom-up growth

Bottom up approaches are desirable for the possibility of synthesizing large areas and uniform layers which are crucial for applications and for the wafer-scale fabrication of electronic and opto-electronic devices. In par-ticular I will focus on two methods that are chemical vapour deposition (CVD) and epitaxial growth of graphene on SiC via thermal decomposi-tion, since they both have been used to realize the samples used in this master thesis. The device architecture studied, in fact, is constituted by CVD grown tungsten disulde (WS2) on bilayer graphene obtained by

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2.3. BOTTOM-UP GROWTH 15

Figure 2.6: Silicon Carbide tetrahedron and overall crystal structure.

2.3.1 Epitaxial growth on SiC via thermal

decompo-sition

Silicon Carbide is a crystal formed by silicon and carbon, constituted by a Si atom placed at the center of a tetrahedral structure with C atoms at the vertexes as reported in Figure 2.6. These structures can be stacked in many dierent ways, determining the vast variety of polytypes (more than 250) in which SiC can be arranged. I will not go through the analysis of dierent exsisting polytypes, however it is important to note that for symmetry reasons the most common one used to obtain graphene sheets are 4H and 6H-SiC. They display similar (but not identical) features and top and bottom crystal faces terminate only with either Si or C atoms. The mechanism to growth graphene on top of SiC involves the desorption of Si atoms from the topmost surface; this can be done by heating the sample in vacuum at temperatures above 1050◦_{. Thermal decomposition of graphene}

has been rst reported for the Si-face of SiC [22] and afterwards adapted also to the carbon face [23], each face has its benets and drawbacks, in this work we choose to use graphene grown on the Si-face for it guarantees a better thickness control. Thermal decomposition on the Si-face generates a layer of carbon atoms (after the Si atoms are desorbed) that, despite having the same topological conguration of graphene, it does not display

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Figure 2.7: 4H and 6H polytypes of SiC, with Si and C-faces.

Figure 2.8: Double layer graphene on SiC. Some of the covalent bonds between the buer layer and the SiC bulk underneath are visible. Adapted from [25].

all the typical properties of graphene. This layer, called buer layer, in fact, presents some atoms covalently bonded to the Si atoms of the SiC underneath (∼ 30% of the total number) [24]. To obtain free-standing graphene is indeed necessary to prolong the growth time, permitting other Si atoms to sublimate; this process generates another layer of carbon atoms that rearrange themselves as graphene and become the new buer layer while the previous one, now decoupled from the substrate, is nally a free-standing layer and displays all the amazing properites of graphene.

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2.3. BOTTOM-UP GROWTH 17 Detail

Double layer graphene used for this master project was grown in NEST facilities in the Graphene Growth Lab, following a standard procedure to obtain high quality materials. This process is developed to prepare the surface of commercially-available SiC wafers and it is done in three dierent steps:

• Piranha solution treatment (H2SO4 : H2O2 - 3:1 for 10 minutes) to

remove organic contaminants;

• Hydrouoridric acid step (HF:H2O solution 1:10 for 1 minute) to

remove oxide on the SiC surface;

• Etching at 1250◦ _{C for 5 minutes in a cold wall reactor (HT-BM,}

Aixtron) to atten the surface;

The same reactor is then used to grow graphene by heating the substrate at 1400◦ _{C in an Ar environment.}

2.3.2 Chemical vapor deposition

Chemical vapor deposition of TMDs is preferable compared to exfoliation technique since it is scalable and the most used to obtain large area and high quality crystals [26]. CVD is implemented in a variety of formats de-pending on dierent pressure regime (like atmospheric pressure APCVD) [27][28], or dierent approach to start the chemical reaction (for example metal organic MOCVD) [29][30]. The common element of all this tech-niques is the fact that the substrate is exposed to one or more volatile precursors that react on (or near) the substrate's surface to produce the desired material; I will focus on CVD growth of WS2 since it has been

used in this master thesis project. The chemical reaction is the following redox and substitution process:

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Figure 2.9: Scheme of chemical vapor deposition set-up with horizontal hot-wall furnace.

W O3+ S → W O2+ SO + W S2

Detail

To synthesize WS2 a classical vapor-phase approach was used with S

(Sigma Aldrich, 99.998 %) and WO3(Sigma Aldrich, 99.995%) powders as

the precursors. All the process were performed in a 2.5-inches horizontal hot-wall furnace (Lenton PTF). There are 2 zones at dierent tempera-ture: the central one, in which the WO3 powder is placed, can be heated

up to 950◦ _{C by the furnace; and an external colder area with a}

resis-tive heating belt to separately control the temperature of the sulfur kept there. Chemical vapor deposition is a growth process that can be aected by several factors that have non-trivial consequences on the result. The principal are:

• Substrate material and its supercial morphology, the roughness de-termines dramatically the mean free path on the surface;

• Temperature of the chamber: this aects the mobility of molecules on top of the substrate, if it is too low the reaction doesn't start, if

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2.3. BOTTOM-UP GROWTH 19 it is too high the surface diusion leads to formation of 3D structure since the molecules can explore all the surface and tend to accumu-late on the same nucleation site;

• Chamber pressure, determined by the quantity of precursors and ux and type of carrier gas;

• Duration time of the process (usually 1 hour);

• Distance and relative position between W O3 powder and substrate

(usually 2 cm, with the substrate downstream, but other works re-port the powder below with the sample lifted up);

The procedure is the following: rst the chamber is pumped down to 5x10−2 _{mbar, this value is important to set a common starting point to}

guarantee the reproducibility of dierent runs. Once the base pressure is reached the furnace starts to heat up the chamber at a rate of 10C/s and, during the ramp, Argon is uxed at 500 sccm (Standard Cubic Centimeters per Minute), leading to a pressure of 4.6 mbar, to prevent the evaporation of the sulfur. When the target temperature is reached, the ow rate is reduced and the belt starts heating the sulfur which evaporates initiating the WO3 sulfurization process. The Argon ux has a double function:

it sets the pressure of the chamber and carries the sulfur vapor to the target sample. In this master thesis work the WS2 was grown directly on

bi-layer graphene via CVD following a procedure developed by A.Rossi et al. [31]. The recipe is adapted to maximize the area of graphene covered by monolayer WS2, which is the target of this study, setting the

growth temperature of the hot zone at 880◦ _{and the Argon ux rate at 80}

sccm. A coverage of the graphene greater than 50% is obtained and the WS2 is mostly monolayer with some bi-layer regions while single-crystals

domain still remain of the order of hundreds of nanometers in size. This implies a complicated surface morphology in which the complete coverage

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of graphene areas is due to coalescence of dierent crystals leading to a monoatomic buy polycrystalline lm.

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Chapter 3 TMDs properties

3.1 Optical properties

One of the most important properties of a monolayer TMDs is the pres-ence of a direct gap, which leads to strong photoluminescpres-ence emission and is appealing for several optical and optoelectrical application. As already discussed these materials undergo a transition from from indirect-to-direct bandgap when going from bulk to monolayer form. This transition is evi-dent in the Photoluminescence (PL) emission; in fact the position and the intensity of the PL peak is modied by the number of the layers as de-picted in Figure 3.1. Monolayer WS2 displays a remarkable emission with

a quantum yield (6%) that is much greater than the corresponding one of monolayer MoS2 (0.1%) [9]. The optical properties of two-dimensional

crystals are dominated by excitonic eects. Excitons in semiconductors are bound states formed by an electron and a hole due to Coulomb attrac-tion; typically they result from photo-excitation. The energetic levels of an exciton mimic an hydrogen atom, and follow a Rydberg-like series where the electron and hole masses are used in place of the ones of the proton and of the electron. In fact the attractive potential U(r) is the same as an Hydrogen atom, corrected with the dielectric constant on the medium ε=ε0 εr (thus the correction is of a factor εr: the relative permittivity)

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Figure 3.1: Photoluminescence spectra for mono-bi-tri-layer and bulk WS2. The positions for the excitons A and B as well as the indirect band gap (I) are labeled, on the right part is reported the magnication factor. Adapted from [32]

U (r) = − e

2

4πε0εrr

Given this, we can rescale the solution of the Hydrogen atom to obtain energy levels of a exciton, knowing that the binding energy of the ground state of H amounts to 1 Rydberg = 13.6 eV = Ry(H). Thus we obtain:

EB(n) = (mr/m0) ε2 r 1 n2Ry(H) = EB n2

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3.1. OPTICAL PROPERTIES 23 EB = EB(1) = (mr/m0) ε2 r Ry(H) where mr= (m_m_ee_+mmh)

h is the reduced mass of the exciton.

Figure 3.2: a)Schematic representation of electron and hole for three-dimensional bulk material and for two-three-dimensional monolayer, with elec-tric eld line depicted. b)Impact on optical absorption spectrum of the change in the dimensionality; is evident the increase of both the exciton binding energy and the optical band-gap. Adapted from [17]

In traditional semiconductors the large dielectric screening (ϵ) and the small quasiparticle eective mass (compared to m0, the free electron mass)

results in exciton binding energy of the order of 1-10 meV [33], therefore at room temperature their eect on optical properties (like photolumines-cence, reectivty, etc) is not appreciable. In fact this bound energy is small compared to thermal uctuation (∼ 25 meV) and their eect can

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Figure 3.3: Photoluminescence spectrum of WS2 with reported the

con-tribution of A neutral exciton (red) and A- charged trion exciton (green). Adapted from [17]

only be seen at low temperature. Experimental and theoretical studies have demonstrated that the excitonic properties of a 2D-monolayer dier greatly from those of three-dimensional bulk semiconductors of the same material [34]. In particular electron and hole forming an exciton in a monolayer TMDs are strongly conned in the plane and experience a re-duced dielectric screening. This can be understood, in the rst place, with a simple semiclassical reasoning: most of the electric eld lines joining the electron and the hole extend outside the material and don't suer elec-trical screening (a representation of this phenomena in depicted in Figure 3.2). The eectiveness of the dielectric screening depends, then, on the separation between the electron and the hole in the real space; this causes a signicant change of the disposition of the energy levels of the excitonic states. This phenomena leads to a deviation from the classical hydrogenic exciton behaviour. As a matter of fact the excitonic series in monolayer WS2 is dierent respect to the simple Rydberg model, in particular for

the n=1,2 excitons, as can be seen in Figure 3.3. Chernikov et al. [17] have used this correction to the Rydberg series to extract an estimation of these eects in their work, concluding that:

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3.1. OPTICAL PROPERTIES 25 • The quasiparticle band-gap increases, mainly due to the dierent change of the density of state of the system, up to (2.41 ± 0.04) eV; • The exciton binding energy increase, due to the lower screening, up

to (0.32 ± 0.04) eV.

This results are in good agreement with the measured photolumines-cence of the A exciton measured in literature [9]. Accurate studies reveal that the A exciton peak is not the only one appearing in the PL signal, but there exist also the contribution of a trion exciton, which is a localized excitation consisting of three charged quasiparticles [35]. If the Coulomb-bound electron-hole pair is strong enough, excitons can capture additional charges and form charged excitons (known as trions). A negative trion consists of two electron and a hole while a positive trion is composed by two holes and a single electron. Then the two peaks revealed in photo-luminesce spectrum are due to a neutral A exciton at 2.036 eV and a A-trion charged exciton at 1.995 eV [32]. This A-trion emission is usually (in the absence of high doping levels) one order of magnitude smaller than the exciton one, and thus the overall peak seems symmetric, even if it has a small contribution of the trion emission at a slightly lower frequency (due to an extra binding energy of ∼ 40 meV). Several works in the liter-ature report the electrical tuning of these excitons, thanks to the dierent nature of the charged A- and neutral A peak [36].

3.1.1 Raman signature of few-layer WS

2

Raman spectroscopy is fundamental for the analysis and identicaition of two-dimensional materials. It has proven to be a powerful and non-destructive technique to recognize dierent materials and to determine the thickness (i.e. the number of the layers) of the two dimensional crystals. It basically consist in mapping the phonon energy of the sample by measur-ing the wavelength dierence between a laser excitation and the radiation

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Figure 3.4: Room-temperature Raman spectra from a monolayer WS2, us-ing the 514.5 and 488 nm laser excitation, respectively, includus-ing lorentzian peak ts for 514.5 nm data. The spectrum collected for 514.5 excitation revals many second order peaks. The inset shows the atomic displace-ments for the in-plane phonon mode E2G(Γ)and the out-of-plane phonon

mode A1G(Γ) for two adjacent layers. Adapted from [identication]

emitted after inelastic scattering. In fact, when the sample is exposed to the laser beam, the most part of the photons is elastically scattered by the sample (Rayleigh scattering), while a small fraction is inelastically scattered. The photons resulting from this process have a frequency that is dierent from those of the incident photons. More likely to occur is the Stokes Raman shift in which the emitted photons have lower frequency than the incident beam (this process generates a phonon in the medium). If, instead, the emitted light has reduced its frequency, than the

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electro-3.2. MECHANICAL PROPERTIES 27 magnetic eld has absorbed a phonon (Anti-Stokes Raman shift) from the material. In both cases evaluating the frequncy dierence with the inci-dent radiation is, then, a reliable method to extract information about the vibrational modes of the sample. Experimental description of this method will be given in chapter 5. For laser excitation of 514.5 nm the Raman spectrum of WS2 becomes very rich revealing many second-order peaks

that are stronger than those observed in the bulk material. In particu-lar the second order peak 2LA(M) at 352 cm−1 _{became twice stronger}

in intensity then the rst order A2G(Γ) peak. This makes the 2LA(M)

peak of Raman spectrum of WS2 a clear ngerprint of the monolayer as

reported in Figure 3.4. Since the work of Terrones et al. [identication] this technique has become standard to identify single layer WS2.

3.2 Mechanical properties

Figure 3.5: Simple stress vs strain plot, the Young's modulus corresponds to the slope of the curve. For moderate strain levels, while the elastic approximation holds, the value of the elastic modulus is constant. On the right side a schematic representation of strain in a cylindrical rod.

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Among all the fascinating properties that have raised up interest in two-dimensional materials, mechanical properties deserve a specic dis-cussion. The most important parameter that characterizes the mechani-cal properties is the Young (or Elastic) modulus E, which describes the tensile elasticity and is dened as the ratio of tensile stress σ to tensile strain ε.

E = σ(ε) ε

The Young's modulus has the dimension of a pressure, being the ratio of a force over an area and the adimensioned strain (dened as the elonga-tion of the material divided per its length). A further crucial parameter usually encountered in mechanical deformation is the Poisson ratio (ν) and quanties the deformation in the transverse direction respect to the applied strain.

ν = −εtrans εaxial

The physical meaning can be easily seen in a cylindrical rod subject to ten-sion: the Poisson ratio quanties the changing in the diameter of the rod. Most materials have Poisson's ratio values ranging between 0.0 and 0.5 [7]; the minus sign in the formula comes from the fact that common materials tends to contract in the directions transverse to the direction of stretch-ing. In the study of 2D-material properties it is important to introduce the the 2D-Young modulus E2D, which is more intrinsic and corresponds

to the 3D value multiplied by the nominal thickness of the atomic layer. In particular 2D-material experiments and simulations usually measure the 2D Young's modulus and the conversion to the conventional 3D is useful for comparison with other bulk materials. In 2D-materials the thickness of the bulk is determined by the number of the layers multiplied by the interlayer distance. Therefore, one needs to divide the 2D value of the monolayer by the interlayer distance in order to convert it into the 3D Young's modulus. Graphene is the strongest material ever studied, hav-ing the Young modulus ∼ 1TPa; Transition Metal Dichalcogenides have

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3.3. STRAIN ENGINEERING 29

E2D E3D

Graphene (349 ± 12) N/m (1025 ± 35)GPa WS2 (177 ± 12) N/m (272 ± 18) GPa

Table 3.1: Table of 2D and 3D-Young's modulus for graphene and Tung-sten Disulde. The interlayer distance is 0.35 nm for graphene and e 0.65 nm for WS2 [7].

shown remarcable values of approximately 1/3 of graphene [8] [37]. In the work done by Liu et al. [7] there is both an experimental and theoreti-cal evaluation of 2D and 3D Young's modulus, the values are reported in Table 3.1.

3.3 Strain engineering

Figure 3.6: WS2 monolayers under isotropic tensile strain and

compres-sion. The equilibrium structures are marked with red circles. a[A◦] indi-cates the lattice parameter of the crystal. Adapted from [38]

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Figure 3.7: a)Photoluminescence spectra of WS2 for dierent strain levels.

b) Integrated intensities for A and A- exciton peak. c) PL peak position for increasing strain values. Is represented the Indirect peak for strain levels ∼ 2.5 %. Adapted from [39]

One of the most attractive facts about TMDs is the intriguing cor-relation that exists between their optical and mechanical properties. In particular the band structure can be tuned with strain, opening the per-spectives of strain engineering studies. Several works have studied the correlation between strain and optical band-gap [40] (see details about the experimental technique in Chapter 4). Uniform (uniaxial or biaxial) tensile strain is expected to induce a linear decrease of the band gap, and reaching a direct-to-indirect gap transition, with the valence band edge shifting from the K to the Γ point (see Figure 3.6). Further tensile strain

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3.4. SUPERLUBRICITY 31 eventually closes the gap, leading to a semiconducting-to-metal transition for extreme elongations (of the order of 11 %) [38]. This transition is ex-pected to be possible in the lab, since for such strain the metal-chalcogen bonds are stretched, but not yet broken. For experimental reasons is more complicated to induce compressive strain, in fact, due to small bending rigidity, suspended TMDs monolayer will bend and create wrinkles. In the work done by Wang et al. [39] the electronic bandgap, and thus the consequent PL emission, is studied via DFT calculation and PL measure-ment. The photoluminescence peak linearly redshift with a rate of 11.3 meV/% for uniaxial tensile strain, that means that both A and A- exci-tons emission frequency is reduced at same rate. This is well expected since the strain reduces the electron bandgap while leaving the exciton binding energy unchanged (except a small eect on the reduced mass). In the cited work it is shown that for strain higher than ∼ 2.5% there is a broadening of the PL peak due to the competition between the direct and the indirect emission, in agreement with the DFT calculations that demonstrate a direct-to-indirect transition for ∼ 2.6% strain levels [39].

3.4 Superlubricity

Superlubricity is a mechanical regime of in which friction becomes to small to be appreciated; its denition is still quite vague because there is not an accepted limit value for which it applies. In this work I will use it to describe the sliding of crystalline akes in a state of ultralow friction (for example triggered by an AFM tip with a reasonably low load). The rst observation of superlubricity in microscale graphite structure was re-ported in 2012 [41], this paper rst brought superlubricity from a merely academic topic, achievable only in idealized condition, to an phenomenol-ogy suitable for application. The fact that bulk graphite (but also TMDs) excels as a lubricant material is a consequence of this low friction between

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dierent layers. In fact, also the mechanical exfoliation process exploits this low inter-layer Van der Waals attraction. TMDs (principaly MoS2 but

also WS2) are nowadays widely studied for their peculiar nanotribological

properties and there are several studies that demonstrate self-retraction and superlubricity eects [42][43]. Van der Waals heterostructure of TMDs with graphene seems to be an optimal platform to study (and maybe ex-ploit) this fascinating nanotribological properties.

In the work done by A.Rossi et al. in his PhD thesis [25], (which has been fundamental for the growth procedure adopted in this master thesis, and for the properties of the grown heterostructure), there is both an accurate theoretical analysis and some experimental evidence of superlubricity in VdW-heterostructure of graphene and WS2. The whole platform is made

as described in Chapter 3: graphene from thermal decomposition of SiC and CVD-grown WS2 on top, tuning the recepy to obtian small single

akes; in his work the sliding of small akes of WS2 on top of graphene is evident. The trigger of this movement can be either the STM tip (see Fig-ure 3.8) or the AFM tip in contact mode (see FigFig-ure 3.9), while when the AFM is set in tapping mode there is no evident sliding of the triangular akes. The sliding under the AFM tip is particularly interesting since it shows superlubricity in ambient pressure condition (while STM operates in vacuum atmosphere) and it can open the road to advanced devices (like nanomotors or nanoelectromechanical systems (NEMS).

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3.4. SUPERLUBRICITY 33

Figure 3.8: Observation of displacement of the akes of WS2 on graphene on SiC due to superlubricity.(a-c) Three STM images of the same area in which is evident the displacement of a single triangular WS2 ake. Adapted from [25]

Figure 3.9: Observation of displacement of the akes of WS2 (highlighted with white circles) on graphene on SiC during an AFM scan in contact mode.a)b) Two AFM images (phase image) in which is evident the sliding of a single WS2 ake. Adapted from [25]

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Chapter 4 State of the art

4.1 Strain engineering in 2D-crystals

A wide range of possible techniques to induce strain in two-dimensional materials exists. These go from very simple ideas to more elaborated se-tups. I will describe rst the most used one in the literature: elongating an elastic substrate, bending of exible substrates, piezoelectric stretch-ing and substrate thermal expansion. These methods are mostly borrowed from well-known material science mechanical tests. In fact, they involve mechanical deformation of the substrate and take advantage of its adhe-sion to the two-dimenadhe-sional crystal to transfer the strain pattern. The interaction with the substrate and the clamping with the sample are not simple to control, and these studies could only demonstrate low strain levels. In the following I will talk about techniques used to induce strain in graphene suspended membranes: with uniform dierential pressure and micro-actuators (rst mechanical then polymeric). One of the great ad-vantages is that it is possible to induce custom pattern of strain; this is particularly intriguing to generate anisotropic strain prole that could generate giant pseudomagnetic eld.

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Figure 4.1: Picture of optical measure on elongated substrate which con-sists of a slab of PDMS. In the middle a WS2 triangular ake enlightened by a laser source is depicted. Adapted from [44].

4.1.1 Elongating the substrate

Uniaxal tensile strain can be applied by depositing the 2D material on an elastic substrate and elongating it with a straining stage. This method for example has been used to apply strain to monolayer WS2 by Wang

et al. [39]. In this specic case the as-grown CVD WS2 thin layers on

Si/SiO2 substrate were transferred on to a exible substrate such as PET

(polyethylene terephthalate). The adhesion between the substrate and the material is non-trivial and is aected by many factors like the chosen plastic support, the quality of the WS2 ake and the adopted transfer procedure. While some groups were able to reach strain levels up to 4% [39], others claim that the van der Walls interactions are strong enough to clamp the ake and prevent the slippage only for strain levels less then 2% [44].

4.1.2 Bending of exible substrates

Similarly to the previous technique, uniaxal tensile strain can be easily applied to 2D materials by depositing the akes onto a target substrate

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4.1. STRAIN ENGINEERING IN 2D-CRYSTALS 37

Figure 4.2: a)Flat substrate with 2d material on top. b)Two-point bending apparatus. c)Four-point bending apparatus. Adapted from [45].

Figure 4.3: Picture of the apparatus of the bending test. R is supposed to be much greater than the substrate thickness h and the 2D material length L, allowing to consider the strain on the 2D sheet uniaxial.

and doing a sort of bending test. Obviously the substrate is chosen with a large length-to-width ratio to allow uniform bending and reversibility; moreover the size of the 2D material layers is orders of magnitude smaller than the substrate length (at least 103_{times smaller) in order to guarantee}

uniform strain on the sample. Is it possible to realize either two-point or four-point bending experiments (see Figure 4.2).

Given these conditions, which permits us to neglect the strain in other directions, the longitudinal strain can be expressed as:

ϵ = h 2R

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Figure 4.4: Device representation of tri-layer MoS2 under piezoelectric

compressive biaxial strain. Adapted from [46]

where h is the thickness of the slab and R the radius of curvature (see Figure 4.3). The maximum reported uniaxial strain achieved by this tech-nique ranges from 0.5% to 2.5%, depending on the exact experimental conditions and the clamping technique used.

4.1.3 Piezoelectric stretching

Piezoelectric substrates are very appealing to stretch two-dimensional ma-terials since one can easily regulate the deformation of the substrate by varying the applied voltage to the two faces of the material. A simpli-ed model of this device is depicted in Figure 4.4. The main drawback of this technique is that only moderate strain levels are achievable and a maximum 0.2% strain has been reported in literature) [46]. Hui et al. have studied the eect of compressive strain of MoS2 onto a piezoelectric

substrate, benetting from the stronger eect on the band structure of TMDs of biaxial strain than uniaxial strain.

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4.1. STRAIN ENGINEERING IN 2D-CRYSTALS 39

Figure 4.5: Exploitation of thermal coecient mismatch, with heating of the substrate and local heating via laser. Adapted from [48]

4.1.4 Exploiting substrate thermal expansion

Thermal expansion mismatch between lm of 2D crystal and substrate is considered to be a drawback since it produces undesired strain during the dierent thermal processes needed for the fabrication of micro-devices. However, is it possible to exploit this mismatch to induce biaxial strain in a controllable way. 2D semiconductor usually have small (or negative) thermal expansion coecient [47], so one can choose a substrate with large positive thermal expansion coecient and then it's sucient to heat up the substrate to cause strain in the 2D semiconducting crystal. This technique has been used both to induce homogeneous strain on the sample by warming up the whole substrate and locally increasing the temperature with a laser beam attaining non-homogeneous strain proles. Even though the choice of specic substrate (like poly-dimethil siloxane) with large thermal coecient, instead of conventional dielectric ones, has magnied the eect of this method, the strain level achieved are still modest (up to 0.23%) [48].

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4.2 Strain engineering on graphene suspended

membranes

To study the mechanical properties of graphene (and now also other 2D-semiconductors) several eorts have been done in realizing suspended membranes of these materials. These setups permit an accurate analy-sis of the material, decoupling it from the interaction with the substrate, and allow the implementation of inhomogeneous strain prole. This could be very interesting for two main reasons: rstly anisotropic strain pro-le can aect greatly the electronic states, changing the gaps and in-ducing other intriguing phenomena; for example it is predicted that tri-axial strain prole is equivalent to the application of a strong pseoudo-magnetic eld and zero-eld quantum Hall eects [49] . Secondly, local variation of strain could be exploited to attract excitons to specic part of the sample, a phenomena, known in literature as funneling eect, which con be appealing for energy harvesting application [50]. Suspended se-tups, though, are very fragile and require elaborated procedures to grow and transfer the samples. There are several techniques used to transfer two-dimensional crystals from the growth substrate to the target device [51] [52]. Usually CVD-grown tungsten disulde is transferred via a Poly methyl methcrylate-assisted wet technique that consist:

• Spinning PMMA over the sample; • Soaking in hot base;

• Delamination of the polymer lm in water; • Fishing the material onto target substrate;

This is the backbone of the procedure; the recipe, then, can be changed adapted varying temperature, concentration and chemical components of

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4.2. STRAIN ENGINEERING ON GRAPHENE SUSPENDED MEMBRANES41

Figure 4.6: Tranfer procedure for WS2 grown on SiO2. Adapted from [28]

the base solution, type of polymer used as vector, timing of baking, delam-ination method, timing and so on. To transfer graphene membrane from Copper (which is the one used in the next paragraphs) the procedure is quite dierent, it is called "bubbling transfer technique": it involves the PMMA as a vector as before, but the detachment from Copper exploits water splitting phenomena (a base solution and voltage generator). Al-though in the recent years transfer procedure has proven to be reliable, they can cause unintentional doping and complicate the fabrication pro-cess. [53]

4.2.1 Straining using dierential pressure

Non-isotropic strain proles can be obtained in free-standing graphene membranes clamped on the edge of a non radially symmetric hole, and subjected to a uniform mechanical load given by a pressure dierence between the two faces of the ake (see Figure 4.7). In the work done by Settembrini et al. [54] it is shown that free-standing membranes are obtained depositing graphene layers onto a micropatterned SiN support, adapting the transfer procedure described in the previous chapter. In the

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Figure 4.7: Device prole, in the middle is depicted the bended graphene membrane, the whole device can be put under a optical microscope to perform raman measurement, to evalute the strain level. Adapted from [54]

cited work anisotropic components in the strain prole induced by the pressure are demonstrated in elliptical holes. This implementation can be pushed further on since alternatives clamping geometries could be used to design more complicated non-uniform and non-isotropic strain patterns basically changing the boundary condition of the suspended membrane. In this specic case a very appeling opportunity is to implement triaxial strain proles that, as previously said, give rise to pseudo-magnetic eld. While this device architecture is promising for all the explained reasons, it has two main drawback: primarily the set-up requires much eort for fabrication and it's very fragile and, secondly, is quite dicult to evaluate the strain given by the graphene adhesion on the vertical sidewalls of the SiN hole.

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Figure 4.8: Device picture. Adapted from [55]

4.2.2 Micromechanical actuators

Another approach to induce strain in graphene membranes is to use micro mechanical actuators. This kind of devices belongs to the promising fam-ily of Micro ElectroMechanical System (known as MEMS) which exploits microfabrication technology to create small-scale mechanical actuators, pumps or motors. In the work done by CITE the device is a MEMS-based micromaterial test platform that consist of two suspended shuttle-beams bridged by the graphene, one of the shuttle is thermally actuated and the other is xed to calibrated springs that allow measuring the force pulling it.

The working principle of the device is the following: Joule heating is induced in the thermal beams by applying a voltage dierence, then ther-mal expansion of the beams cause a controllable motion of the shuttle. The device is depicted in gure 4.8. This technique, which has reached record strain level (close to theoretical limit of 11%) in a controllable and reversible way, is suitable to perform test on graphene's properties. It also has the benet of being very compact (circa 1mm2 _{of surface), such as}

the whole device can be easily put under a optical microscope in a Ra-man set-up to take measurement of strain levels. Moreover a recent work, by Goldsche et al. [56], has adapted MEMS to achieve triaxial strain in a set-up that consists of three shuttles, and has thus extended the

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ex-Figure 4.9: Device picture zoomed in the central zone. There are repre-sented the two epoxy clamps to stick graphene sheet to the shuttle. On top of it there is the scheme with laser source used for Raman spectroscopy measurement. Adapted from [55]

perimental possibility of this technology. The greatest advantage of using MEMS devices is that they guarantee a precise direct control and measure-ment of the pulling force applied to the sample. Despite these strengths, MEMS-based set-ups like this require elaborated fabrication procedures from precise transfer of graphene membrane between the shuttles, to the use of epoxy patches to prevent the graphene from slipping away when experiencing high tensile forces. The last part, despite its experimental complexity, it is of crucial importance since without proper clamping be-tween substrate and graphene, extreme strain level couldn't be reachable. In summary, micro-actuators are suitable to perform precise mechanical test on two-dimensional materials but remain challenging to produce and measure procedures involves non-trivial clamping mechanism.

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4.2.3 Polymeric articial micro muscles

A fully new approach adopted by Colangelo et al. in 2017 [57] to con-trol the strain in suspended membranes of graphene involves polymeric micrometric articial muscles (MAMs). The cited work, developed in the NEST laboratory, was the rst to exploit PMMA as an actuator and con-stitutes the starting point for this matesr thesis project. This new device architecture is thought to improve the setups with the standing graphene membranes strained with the pressure load (previously described) allow-ing completely custom strain prole. In this case the active MAMs active material is PMMA, this choice as been made to take advantage of PMMA properties since it can be used both as positive and negative resist. In fact PMMA is rst spinned over the graphene suspended membrane (the same conguration as the one strained with pressure loads), then removed everywhere but selected areas to create the "muscles ber" and nally crosslinked with high electron dose to induce the strain. A detailed de-scription of the MAMs realization and utilization will be given in chapter 6, since the same polymeric micro muscles actuators have been used for this master project. Here I will stress MAMs benifts: their geometry can be fully customizable leading to any desired strain geometry, furthermore this technique, is obtained by e-beam lithography and, thus, permits in-situ imaging via scanning electron microscopy.

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Figure 4.10: Representation of deposition of graphene and PMMA on the substrate with the hole, removal of PMMA on the suspended graphene to obtain free-standing membrane, creation of the 'muscles' with e-beam irradiation. Adapted from [57].

Figure 4.11: SEM image of a circular membrane with the MAMS. There are indicated the polymer-free graphene and the two area of PMMA ex-posed to e-beam which contracts. Adapted from [57].

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Chapter 5 Experimental technique

5.1 Atomic Force Microscopy

Atomic Force Microscopy is a a high-resolution scanning probe technique that allows obtaining subnanometer-scale information about the surface morphology and the thickness of 2D-materials. The method is based on probing the near-eld forces acting on a sharp tip placed in contact or inproximity with the studied sample. This interaction can be described by the Lennard-Jones potential:

VLJ = ϵ [( r0 z )12 − 2( r0 z )6]

where ϵ is the depth of the potential well depicted in Figure 5.1 a) and r0 is

the equilibrium distance. This is a widely used semi-empirical description in which potential is constituted by two terms. The rst high exponent term describes the short-range interaction of the almost impenetrable core of the atoms of the sample. The second term with the inverse six power law corresponds to the uctuating dipole-dipole interaction, which is at-tractive. The measurement principle is rather simple: a cantilever with a sharp tip is scanned over the sample and the supercial morphology of it are measured recording the cantilever deection using a laser beam, the optical lever arm of the cantilever and a four-segment photodetector (see

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Figure 5.1: a) Lennard-Jones potential with highlighted the area of inter-est for dierent modes in which is possible to take measurement: in red contact mode, in yellow non-contact mode and tapping-mode with the black arrow. b) Simple representation of an AFM set-up. c) Dierences between contact and tapping mode.

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5.2. SCANNING ELECTRON MICROSCOPY AND E-BEAM LITHOGRAPHY49 Figure 5.1 b)). The three main modes in which an AFM can be used are

the following (depicted in Figure 5.1 c)):

• Contact Mode: the tip is dragged across the sample and the sur-face height is measured either using the deection of the cantilever directly or, more commonly, using the feedback signal required to keep the cantilever at a constant position;

• Non-contact Mode: the tip does not contact the surface instead is oscillated around its resonant frequency over the sample and expe-riences change in frequency due to WdV forces;

• Tapping Mode: tip is oscillated at xed frequency and the surface is mapped via the change in amplitude due to interaction with the substrate.

In this work the AFM set-up used is a Bruker Dimension Icon Atomic Force Microscope.

5.2 Scanning Electron Microscopy and E-Beam

Lithography

Scanning Electron Microscopy (SEM) is one of the most common tech-niques in nanotechnology. It consists in scanning the surface of the sample with a focused beam of electrons that are accelerated in high vacuum to a few keV. The electrons interact with atoms in the sample, producing various signals that contain information about surface topography and composition of the sample. The image is reconstructed collecting the out-coming electrons that can be distinguished according to the dierent kind of interaction that they experience. They can be divided in two categories: backscattered and secondary electrons. The rst ones experience elastic scattering from the nucleus and their collection at dierent angles allows

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a reconstruction of the image of the system. The second one, originated by inelasting scattering, have low energy <50 eV, and deliver information about the topography of the surface. The image generated by a SEM is black-and-white, and is constructed associating to every pixel the number of electron collected by the detector in a small amount of time while the electron beam is targeted to that position.

The same focused electron beam can be used to perform Lithography drawing custom shapes on a surface covered with an electron-sensitive lm called resist. This is a common application of EBL called direct writing, the system basically consists of a source of electrons, a focusing optics set, a blanker to turn the beam on and o, a deection system for scanning the beam over the sample, and a stage for a coarse positioning of the studied sample (see Figure 5.2). The resist materials, poly-methy-methacrylate (PMMA) being one, are at the core of fabrication process. A solution of PMMA can be spinned on the planar surface of the substrate, with a thickness controlled via the rate of the spinning. Then a short baking is needed to let the solvent evaporate thus leaving only the polymer on the substrate. The resist can be divided in two categories, depending on how they react to light exposition, electron beams or other external stimuls:

• Positive resist are broken into smaller organic units that are much more easily dissolved by solvents that to not attack the unexposed material.

• Negative resist are polymerized further and harden so that solvents subsequently remove those parts that have not been exposed. In this work PMMA as been exploited both as positive and negative resist as described in the work of Zanier [58], depending on the irradiation dose of exposure, and the lithography and the scanning electron microscopy were performed in the same experimental setup. The primary advantage of electron-beam lithography is that it can draw custom patterns with

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5.3. MICRO-RAMAN SPECTROSCOPY 51 a sub-10 nm resolution, avoiding the use of a mask. It is important to underline that imaging with the SEM the sample is always followed by unintentional deposition on the sample of ions (usually of carbon) cre-ated by the electron beam. Despite the vacuum chamber, in fact, there are molecules of CO2 and other substances near the sample that can be

ionized and than attracted by the charge deposited on the sample. Car-bons and other elements deposited by this phenomena hamper the PL signal of WS2. Therefore the samples in which we have studied the PL

emission, have not been imaged by scanning electron microscopy after the MAMs contraction [59]. In this thesis project the SEM measurement and EBL procedures were performed using a ZEISS Ultraplus FEG Scanning Electron Microscope equipped with an interferometric stage.

5.3 Micro-Raman spectroscopy

Raman spectroscopy is a technique used to observe vibrational mode in a system. It relies on inelastic scattering of monochromatic light, usually from a laser in the visible spectrum. When laser light interact with the material, three dierent process can occur (see Figure 5.3):

• Rayleigh scattering is the elastic scattering of photons on the sample, the frequency of the photon emitted is equal to the incident one. This part of the spectrum constitutes the most intense one, as reported in the spectrum in Figure 5.3;

• Stokes Raman scattering is inelastic scattering of photons in which the emitted photon frequency is smaller than the incident, since it has generated a phonon in the sample;

• Anti-Stokes Raman scattering the process analogous to the previ-ous described in which a phonon is absorbed from the material and results in a photon energy bigger than the incident photon;

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Figure 5.2: Scheme of a Scanning Electron Microscope with all the com-ponents to generate, focus and then collect the electron backscattered.

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5.3. MICRO-RAMAN SPECTROSCOPY 53

Figure 5.3: Picture of the dierence of Rayleigh, Stokes Raman and Anti-Stokes Raman scattering. The lower gure represents the spectrum of the scattered photons with realistic intensity. Is evident that the most intense part is due to elastic scattering, and needs to be properly ltered to reveal the other components of the spectrum.

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Figure 5.4: Scheme of Raman set-up used. There is the laser source (532 nm in our Renishaw set-up), the dierent optical lens, grating to separate dierent frequency components and the CCD detector that collects the signal.

Electromagnetic radiation from the illuminated spot is collected with a lens and sent through a grating and than a detector. The most commonly used are a Charge-Coupled Devices (CCD) that convert incoming photons into electron charges, quantifying the radiation collected. Elastic scattered radiation at the wavelength corresponding to the laser line (Rayleigh scat-tering) is ltered out, while the rest of the collected light is dispersed onto the CCD. Important to perform micro-Raman spectroscopy is the possi-bility of focusing the laser in a spot at the diraction limit (for visible light ∼ 500nm) to achieve at least micrometric resolution. The sample should furthermore on a stage that can be moved with submicrometric precision, allowing to create maps of data. Stokes and Anti-Stokes Raman scattering are typically very weak compared to Rayleigh scattering. In Raman spec-troscopy wavenumber is commonly used (also called Raman Shift) instead of wavelength, dened as the dierence of the incident wavenumber and the detected one.

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5.4. PHOTOLUMINESCENCE 55

Figure 5.5: Renishaw set-up used in this master thesis both for Raman and PL measurement, in the lab of material characterization. It is equipped with a 532 nm laser and four dierent lens 5X, 20X, 50X, 100X visible in the inset. RamanShif t = ∆ω =( 1 λL − 1 λs )

This typical unit of measurement of Raman shift is the cm−1_{. Every}

material can be detected via his typical Raman spectrum, thus, this tech-nique has proven to be a useful non-destructive tool in material science.

5.4 Photoluminescence

Photoluminescence is light emission from matter after the absorption of electromagnetic radiation. Usually in semiconductors light with energy

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Figure 5.6: Two band model for photoluminescence. There is the repre-sentation of the incident light at hν > EGAP, then the relaxation towards

the bottom of the conduction band for electrons (top of valence band for holes), and in red the emitted photon.

hν > EGAP is shone on the material, a carrier will occupy an excited

state via the absorption of a photon and then it will relax on the ground state emitting a photon of dierent energy. The frequency of the emitted photon gives information about the bandgap of the material and in two-dimesional TMDs on the energy binding of the exciton formed via the interaction of the electron and the hole. In fact the frequency of PL peak corresponds to the optical bandgap reduced by the binding of the exciton. The experimental set-up is similar to the one used to perform Raman spectroscopy (see Figure 5.4), but in this case the measurement is directly the frequency of the emitted photon.

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Chapter 6 Experimental results

The starting idea of this master project is to exploit superlubricity between tungsten disulde and graphene to realize a platform for strain engineering studies. WS2 is grown by CVD on top of graphene and benets from the

ultralow friction with his substrate responding to horizontal forces like a free-layer of 2D material. The goal is to combine this "superlubric table" [11] with polymeric micro-actuators [57] to create a playground for strain studies in a non suspended set-up. In this chapter there is the description of the original work developed during this master thesis project. Firstly there is the characterization of the device studied via SEM, AFM, Raman and PL spectroscopy. Then there is description of the design and fabrication of PMMA actuators followed by the demonstration and the quantication of their contraction. After that, the strain induced in WS2 is demonstration and the evaluated via the measurement of PL

shift. The last part of the chapter discuss the relaxation of the muscles with time.

6.1 Device architecture

The device studied in this master thesis (represented in Figure 6.1) has been fabricated in NEST laboratory in Pisa by the Graphene Growth

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Lab. The SiC substrate is a commercially available 6H-SiC(0001) wafer. Bilayer graphene on top of it is grown via thermal decomposition of SiC and WS2 is CVD grown. Both these steps have been obtained following

the procedure already illustrated in Chapter 2. On top of the structure a 100nm-thick PMMA layer is spin coated on sample.

Figure 6.1: Representation of the whole heterostructure. At the bottom of the gure, the SiC substrate is visible, on top of which bilayer graphene is grown by thermal decomposition. The rst layer is indicated as "Buer layer", since 30% of its C atoms are covalently bonded with the SiC. On top of it there is Tungsten Disulde grown by chemical vapour depostion. The whole heterostructure is coated with a 100nm-thick layer of PMMA.

6.2 Structural and morphological

characteri-zation

6.2.1 AFM characterization

AFM shows that the growth process does not yield a monocrystalline monolayer sheet of WS2. Rather, WS2 only partially cover the bilayer

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6.2. STRUCTURAL AND MORPHOLOGICAL CHARACTERIZATION59 terraces, where multilayer WS2 is also present. In Figure 6.2 are evident

both the multi-layer growth and the fractures in WS2.

Figure 6.2: AFM image. There are the SiC terraces and on top the tung-sten disulde supercial morphology, with multi-layer regions and cracks. The darker areas in gure are without WS2while brighter areas are

multi-layer WS2 preferentially grown at the edges of SiC terraces.

6.2.2 SEM characterization

The surface of the heterostructure has been characterized by SEM imag-ing, to investigate the supercial morphology produced by the growth steps. WS2 is typically found to cover approx approximately 50 % of the

sample, mostly being monolayer, with some bi-tri-layer enclosures. This can be seen in Figure 6.3 where, in the inset, is described the correlation between the grey-scale of the SEM image with the presence of WS2 and

graphene. Being able to recognize the presence of WS2 and his integrity

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Figure 6.3: SEM image. The greyscale has been correlated to the presence of the Tungsten Disulde and his thickness: white is graphene, dark gray is monolayer WS2 (1L−WS2), light grey is WS2 bi-trilayer (ML−WS2).

that could be strained with the polymeric actuators. In order to obtain these information the data of the SEM image have been crosschecked with Raman and PL measurement; this process is described in the next section.

6.2.3 Raman characterization

Several works in literature (one among all, the article of M. Terrones et al. [60]) has individuated the ngerprint of the monolayer WS2 when

the ratio of the 2LA(M) is about twice the A1G(Γ) peak (for selected

excitation wavelength). The typical Raman spectrum of monolayer WS2

in our device is reported in Figure 6.4. The spectrum revealed is very rich; it is composed by many vibrational modes of WS2 and their second

order modes. Thanks to four lorentzian tting (of the peaks indicated in Figure 6.4 with black arrows) we extrapolated the intensity of dierent

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6.2. STRUCTURAL AND MORPHOLOGICAL CHARACTERIZATION61

Figure 6.4: Typical Raman spectrum of WS2 on bilayer graphene on

Sil-icon Carbide. 100X lens, 0.5 µW with an integration time of 3 seconds. The rich variety of peaks is due to the resonance of the excitation wave-length, but the ngerprint of the monolayer WS2 is evident being the

2LA(M) peak much greater than A1G(Γ). Typical peaks transverse (TO)

and longitudinal optical (LO) phonons of the SiC are evident in the right part of the spectrum.

peaks. This is not trivial for peak 2LA(M) since it partially overlaps with E2G(M) at around 355 cm−1. Given the intensity value of each peak

was than possible to compute the intensity ratio I2LA/IE2G(M ) for each

spectrum acquired. Collecting a Raman spectrum for an array of points (with the set-up described in the previous chapter) we were able to map the sample, distinguishing zone with monolayer, from multilayer and areas without WS2 with a resolution of the order of the focal spot (in our case

∼ 500nm , with 532nm laser source). Crosschecking this information with the SEM maps, we were able to correlate the greyscale of the SEM with

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the material on the surface, realizing a rough (but fast) way to investigate the material surface.

6.2.4 PL characterization

Figure 6.5: Typical photoluminescence spectrum, taken with 100X lens, and laser excitation of 532 nm. Thanks to gaussian tting it is evident that the asymmetric peak is due to the contribution of both the neutral A exction and the A- trion.

Photoluminescence measurement are carried on in the same experimen-tal set-up; the typical spectrum for the monolayer is reported in Figure 6.5. This spectrum displays all the features reported in literature for tung-sten disulde, intense excitonic peak at around 625 nm composed by two contribution: major peak corresponds to neutral A exciton and a lower contribution by a charged exciton A- (trion exciton). Peak values are reported in gure and are in good agreement with values reported in lit-erature. The t is gaussian since the complicated supercial morphology

Strain engineering of Tungsten Disulfide with Polymeric Micrometric Artificial Muscles

Contents

Chapter 1

Introduction

Chapter 2

TMDs structure and growth