Synthetic approaches to two-dimensional transition metal dichalcogenide nanosheets

Accepted Manuscript Synthetic Approaches to Two-Dimensional Transition Metal Dichalcogenide Nanosheets Jack R. Brent, Nicky Savjani, Paul O'Brien PII: DOI: Reference:

S0079-6425(17)30067-1 http://dx.doi.org/10.1016/j.pmatsci.2017.06.002 JPMS 459

To appear in:

Progress in Materials Science

Received Date: Accepted Date:

25 May 2017 8 June 2017

Please cite this article as: Brent, J.R., Savjani, N., O'Brien, P., Synthetic Approaches to Two-Dimensional Transition Metal Dichalcogenide Nanosheets, Progress in Materials Science (2017), doi: http://dx.doi.org/10.1016/j.pmatsci. 2017.06.002

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Synthetic Approaches to TwoDimensional Transition Metal Dichalcogenide Nanosheets Jack R. Brent1†, Nicky Savjani2† and Paul O’Brien1,2* 1) Materials Science Centre, University of Manchester, Oxford Road, Manchester, M13 9PL, United Kingdom. 2) Department of Chemistry, University of Manchester, Oxford Road, Manchester, M13 9PL, United Kingdom. E-mail: Paul.O’[email protected], Tel: +44 161 306 2279 † Both authors contributed equally to this review

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Abstract Two-dimensional transition metal dichalcogenides (TMDCs) have become the focus of intense research due to their unique physical and chemical properties. These features arise from the anisotropic effects of the materials as it is thinned towards a monolayer. Several review articles have exhaustively examined the structures and applications of such materials, whereas relatively little critical analysis has been applied to the ever-expanding array of synthetic methods. This review thoroughly documents all synthetic methods to 2D-TMDCs (by both ‘top-down’ and ‘bottom-up’ approaches) published up to January 2015, by considering the quality and characteristics of the nanosheets produced. This work has been achieved by directly comparing the crystallinity, dimensions, yields and electronic properties of selected TMDC nanosheets with respect to each method. Finally this review concludes with an examination of the next generation 2D-TMDCs that explores the production of ternary TMDC nanosheets, heterojunctions between two TMDCs, and the production of patterned nanosheet surfaces.

Keywords:2D-materials, transition metal dichalcogenides, nanosheet exfoliation, nanosheet growth, FETs, doping, heterostructures.

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1 Introduction The discovery of graphene in 2004 has spawned entirely new fields of research.[1] The unique properties arising in highly confined two-dimensional carbon systems have the potential to revolutionise a number of industries and create new ones. Specifically, graphene displays unparalleled strength[2], conductivity[1] and flexibility,[3] compared to alternatives like indium-tin oxide (ITO), and is more transparent than metal films, meaning the 2Dmaterial has the potential to enable a wide range of flexible, ultrathin electronic and photovoltaic arrays.[4-7] Graphene has already been shown to be a highly suitable material for the development of next-generation energy storage, bioimaging, sensing, electronic and catalytic devices (among others).[8-11] Graphene can also be processed from solution bringing inherent benefits of using more efficient printed and roll-to-roll manufacturing approaches.[12] The exceptional conduction arises as a result of graphene’s extended network of sp2 hybridized carbons, a feature which also gives rise to zero-bandgap, metallic properties. Metallic characteristics prevent the use of graphene in photovoltaic and logic devices without considerable modification of the pristine sheet to generate the desired semiconductor properties.[13-15] Considerable effort has been focussed on the synthesis of 2D-nanosheets of a wide variety of more suitable materials, leading to the successful syntheses of a vast array of organic and inorganic compounds, fabricated using a wide range of synthetic methods.[16-20] Perhaps the most promising and best-explored class of ultrathin nanosheets are those derived from the transition metal dichalcogenide (TMDC) families of materials (MX2 where M is a metal and X is sulfur, selenium or tellurium). These span the full range of electronic behaviour from metallic to wide band-gap insulating properties.[21, 22] The structures and electronics of these compounds have been exhaustively described elsewhere[21] but briefly, the crystals consist of atomic layers of hexagonally-packed layer of metal atoms sandwiched 3

4 between two layers of chalcogen atoms.[23] This sandwich structure results in valencesatisfied atoms, absent of reactive sites, rendering the c axis relatively inert, when compared to the edge sites. The precise coordination geometry about the metal centre (Oh or TP) is primarily determined by the metal d-electron count, but also on the relative sizes of the constituent elements.[24] The d-electron count also determines the electronic and magnetic properties of each material,[24, 25] with a significant contribution from the chalcogen.[23] As with graphene,[26] these

two-dimensional materials exhibit quantum

confinement effects[27] when thinned to a few-layer sheet. It has been shown that these effects increase anisotropy and can produce a host of enhancements in the electronic,[28-30] optical,[31] structural,[32-34] thermal[35] and

mechanical[36] properties that are

significantly different from those of the bulk material. Such effects are well illustrated by a number examples: The semimetallic-to-semiconducting transition observed in ultrathin TiS2,[37] the metallic-to-insulating transition of TaS2 [22] and effects on the band-gaps indirect-to-direct transition,[38] widening of the bandgap by up to ~50 %[29, 39-41]) For Mo and W TMDCs thinning from bulk to a single layer and dramatic improvements in catalytic activity.[42, 43] It is hoped that successful exploitation of the properties arising in ultrathin TMDCs will allow their use in a variety of electronic,[44-46], solar,[47, 48] energy storage[25, 49, 50] and biomedical[51] applications.[52, 53] MoS2 is by far the most commonly synthesised MX2 nanomaterial and the material most commonly incorporated into devices. MoS2 also displays a number of well-documented electronic properties whose theoretical limits have been thoroughly explored.[54, 55] In nanomaterials syntheses, the crystallinity is a crucial factor determining its quality, however it is difficult to measure the crystal structures of the thin TMDC nanosheets (< 10 nm) using standard in-plane X-ray diffraction;[56] high-resolution electron microscopy is generally required to directly determine these features. We believe that, although a number of factors 4

5 contribute to measured electronic properties,[57] the carrier mobility (µ e) and on/off ratios (ION/IOFF) of untreated MoS2 field-effect transistors (FETs) (often created in many works on MoS2 nanosheet production) indirectly correlates to the crystallinity of the nanosheets used, therefore may be employed as a useful yardstick for the comparison of the quality of nanosheets produced. A summary of the performances of MoS2-based FETs, as well as typical nanosheet dimensions isolated through the methods documented in this review are shown in Table 1 andTable 2. It is clear that the incorporation of nanosheets into working devices has been well-explored elsewhere[51, 58] but a full examination into the synthetic routes to these nanosheets is still lacking. For this reason we here examine the synthetic methods that have been published (up to January 2015) for the production of TMDC nanosheets, focussing on the factors which can be used to tune the thickness, lateral dimensions and crystallinity of the product. In addition, we analyse each procedure in terms of its ease of scale up and suitability for specific applications (catalytic, photovoltaic etc.). We hope to collate the details of a synthetic toolkit, using which it will be possible to selectively synthesise nanosheets with desired properties in ways which allow their incorporation into truly novel applications. The crystallisation of TMDCs into nanoflowers and other hierarchical arrangements, and the formation of other low-dimensional metal chalcogenide materials are beyond the scope of this review.

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2 ‘Top-down’ Methods for TMDC Nanosheet Production 2.1 Introduction As discussed in Section 1, lamellar TMDC crystals have strong intralayer covalent bonds and weak interlayer van der Waal forces.[59] By imparting a shear force to the bulk crystal, one can overcome the weak van der Waals forces within the crystal lattice and produce thinner materials. Repeating this thinning process can exfoliate the lamellar flakes further and eventually allow the isolation of single layer nanosheets.[60] The ways to affect this process can be separated into several different types (Figure 1): Micromechanical exfoliations rely on the introduction of a peeling force to separate the layers, a shear-induced separation or by compressive delamination. Exfoliation in a liquid medium generally operates by applying a vibrational energy to the crystal or exposing it to forces generated within a cavitation field. Chemical intercalation and exfoliation works on the basis of the insertion of a small atom or molecule within the interlayer space of the bulk TMDC, before forced expansion of the intercalate separates the crystals. Finally, thinning by etching techniques require not the separation of sheets, but rather the destructive removal of unwanted layers from a bulk crystal to produce a one- or few-layered nanosheets on a substrate. Below we examine the routes to TMDC nanosheet fabrication by ‘top-down’ exfoliation methods with specific attention paid to the quality of nanosheets produced.

2.2 Micromechanical Exfoliation The first use of what would become the ‘Scotch Tape method’ to thin layered compounds was reported by Frindt in 1966 and resulted in the isolation of flakes of MoS2 approximately 10 nm thick.[61] More recently, Geim, Novoselov and co-workers,[1, 62] famously went further, greatly improving micromechanical cleavage of lamellar crystals. This allowed the isolation and thorough scientific assessment of single- and few-layer TMDC 6

7 nanosheets. The method is based on the adhesion of a lamellar crystal between two pieces of an adhesive tape, before peeling the tape apart to cleave the crystal. The adhesion between the tape and the basal plane of the crystal is significantly stronger than the weaker van der Waals forces between the layers, allowing isolation of increasingly thin flakes. Repeating this process on the freshly peeled flakes has been found to produce single- and few-layer nanosheets that can be transferred onto a substrate for characterisation and device fabrication. Since the Scotch Tape method was shown to produce single-layer graphene,[1] this method has been used in many occasions in the production of nanosheets of a host of lamellar crystals, including many TMDCs, where its parent lamellar crystal is available (Figure 2).[12, 60, 62, 63] Castellanos-Gomez found that viscoelastic stamps can be used to provide the same effect as Scotch tape in exfoliating lamellar crystals. Nanosheets of NbSe2,[64] TaSe2[65] and MoS2[64] have been obtained by repeated peeling of the TMDC crystals using a poly(dimethylsiloxane) stamp. Flakes can be re-cleaved onto clean sections of the stamp to produce ultrathin and monolayer sheets which can be transferred to a substrate. The viscoelastic stamp method has additional benefits over the Scotch tape method as the stamp does not leave traces of adhesive bound to the surface of the sheet, which can be problematic when truly pristine nanosheets are required.[65] Shortly after publication of the Scotch Tape method, the Geim and Novoselov also found that simply rubbing a lamellar crystal (such as NbSe2 and MoS2, amongst other examples) onto a desired substrate (it was suggested that virtually any solid surface is suitable) can provide the shear forces required to delaminate the nanosheets from its parent crystal.[62] The discovery of the simple 'Chalkboard' method led to attempts to increase the volume of nanosheets which could be produced. One improvement was the development of a ‘sandpaper-assisted’ rubbing technique, which uses the friction of fine-grain sandpaper on the 7

8 TMDC powder to generate the shear forces required for exfoliation. Initially, the use of sandpaper allowed the production of nanosheets with thicknesses >100 layers.[66] The procedure was refined by the introduction of a smooth rubbing step whereby thicker material was deposited into a film, followed by further sandpaper-assisted thinning (Figure 3a) to produce a film of single- and few-layer thick nanosheets. The freshly produced nanosheets can then transferred onto a solid or flexible substrate for further investigation.[67] Gacem and colleagues tackled the micromechanical exfoliation of TiS2, NbSe2 and MoS2 by the development of a widely-applicable, electrostatic-assisted exfoliation method.[68] Subjecting a glass substrate and TMDC crystal to an electric field at raised temperature and pressure promotes the decomposition of Na2O to relatively mobile Na+ and immobile O2- ions within the substrate. Migration of cationic sodium ions produces a negative space charge and induced electrostatic field within the lamellar crystal. Sufficient attraction between the crystal surface and the substrate allows cleavage of TMDC layers which are perfectly adhered to the glass. In most cases the isolation of a single layer via this method is not possible, but the adhered flakes can then be further thinned using the available micromechanical methods discussed before. This method was found to produce significantly larger nanosheets; the micromechanical forces are being applied more uniformly across the crystal surface when overcoming the interlayer van der Waals forces. Tearing of the nanosheets is therefore inhibited and thin flakes of greater lateral dimensions may be obtained. This is shown by the significantly larger single-layer graphene nanosheets obtained (cross-sectional dimension increases of 100’s of µm were documented), as well as the more moderate increases in dimensions in single-layer MoS2 and NbSe2 (to up to 30 µm observed). All of the simple, small scale micromechanical methods mentioned in Section 2.2 are facile, cheap and can rapidly isolate a small number of nanosheets. As relatively lowenergy processes are used, micromechanical exfoliation minimises unwanted scission or 8

9 breakage of sheets, allowing isolation of (1 to 20) µm-scale mono- and few-layer sheets of pristine, crystallographically pure nanosheets; a higher purity material than those obtained from any other fabrication method. The pristine quality of the TMDC nanosheets produced is evidenced by the highest recorded values of µe and ION/IOFF in devices produced.[69, 70] For these reasons the micromechanical methods are most regularly employed by research groups studying the fundamental properties of such ultrathin materials[71-101] or for proof-ofconcept device fabrication.[57, 69, 70, 102-155] However, a number of problems are regularly documented with the use of micromechanical exfoliation, the most important issue being that TMDC monolayers produced from the processes are a tiny minority of the flakes isolated, often found as monolayer regions within flakes made up of bilayer, trilayer and thicker areas, leading to flakes containing numerous sections with different thicknesses. In addition, these single layer areas are also commonly found in small, irregular shapes, complicating their identification and incorporation into devices. As a consequence, significant time is expended assaying substrates for viable nanosheets for further use. These issues make the standard micromechanical processes impractical as a scalable method for nanosheet production. These methods discussed in the following sections of this review are essentially an attempt to deal with the lack of thickness control and scalability inherent in the above methods. In an effort to improve the scalability of mechanical exfoliation, a wet grinding method was developed to promote the potentially scalable micromechanical exfoliation of TMDC crystals via the formation of colloidal aqueous dispersions of MoS2,[156] WS2[156] and NbSe2[157] nanosheets. The wet grinding method causes the delamination of TMDC crystals by generating both shear and compression forces to produce the nanosheets within the aqueous medium (Figure 3b).[158] Chu et al. state that this method is capable of generating MoS2 and WS2 nanosheets in high yields;[156] though precise figures were not 9

10 reported. It is clear that this process is not capable of producing monolayers; the average number of layers in MoS2 nanosheets were found to be 3-5 layers, with lateral dimensions of 500 nm to 1 µm (Figure 3c and d);[156] exfoliated WS2[156] and NbSe2[157] nanosheets were both isolated with smaller cross-sectional dimensions than for the molybdenumanalogue. Considerations of the mechanism of delamination observed during the grinding method led to the use of inorganic salts (such as NaCl) in a solvent-free grinding medium.[159] The rough surface of the inorganic additive promotes the repeating compression- and shear-induced smearing of the large TMDC chunks into smaller nanoflakes during the milling process. The advantage of the ease of additive removal of NaCl by a simple water washing step to leave uncontaminated TMDC nanosheets makes this an attractive avenue for further exploration. The milling times had shown control of the production of few- or single-layer MoS2, with 3 hours milling was found to produce singlelayer MoS2 in yields of up to 95%, but with the consequence of a significantly reduced lateral dimensions 30-200 nm (Figure 4). Such grinding techniques are clearly scalable processes, capable of producing large amounts of high-surface area material, however such a drastic reduction in lateral dimensions leaves sheets unsuitable for most electronic applications.

2.3 Liquid-Mediated Exfoliation Looking beyond mechanical methods, many recent strategies for exfoliation have been based on the ultrasonication of a liquid sample containing the bulk TMDC crystals, which results in energetic agitation.[12] The waves propagate through the solvent causing alternating high and low pressure cycles that aid exfoliation. These cycles generate two different forms of energy inputs; namely vibration and cavitation forces.[160-162] The process for liquid exfoliation relies on the continuous exposure of the crystals to these energetic forces to continually thin them, until dispersed TMDC nanosheets can be isolated from the unexfoliated material by centrifugation.[163, 164] 10

11 It is thought that the liquid-mediated exfoliation process is a practical approach to large-scale nanosheet production. The ease with which these colloidal nanosheet suspensions can be used to produce air-stable inks[165], thin films[166, 167], and both hybrid and composite materials[168] that can be used to further understand and take advantage of the nanosheets produced has resulted in significant interest in the method. Before the process can be scaled up, a number of conditions must be fulfilled to ensure scalable exfoliation is achieved. In particular, maximising the efficiency in the modes of exfoliation via the vibration and cavitation processes is paramount. Until recently, liquid-mediated exfoliation was synonymous with ultrasonic exfoliation. The use of an ultrasonic bath or probe remains the most common approach to the formation of colloidal nanosheet dispersions, though recently developed jet cavitation [169] and high-shear mixing [170] methodologies (discussed below) are also highly promising. In traditional methods, ultrasonic energy, when transferred into the liquid by either an ultrasonic probe or horn, generates both vibrational and cavitation modes within the solvent during the exfoliation procedure. The vibrational energy arises from the generation and propagation of a transverse wave through the solvent which causes vibrational modes within the lamellar crystals. With enough ultrasonic power, these vibrational modes are able to overcome the van der Waals forces within TMDC flakes, as the combination of the compressive wave on the interface of the layered material and a tensile stress wave reflected back on the material body is able to peel nanosheets from the bulk.[12] Increasing the ultrasonic power also allows for the generation of cavitation forces, which carry more energy than vibrational forces. The cavitation forces are formed by the in situ generation and implosion of a cavitation bubble within the dispersing solvent, by the intrinsic low-high pressure cycles observed within the dispersing medium. The implosion of the cavitation bubble generates fast and powerful shockwaves that generate high local pressures (KPa to 11

12 MPa scale) and temperatures of 100’s to 1000’s °C.[161, 162]The cavitation forces generated from the bubble-collapse mechanism promote shear-exfoliation of the crystals.[160] As well as promoting exfoliation, both energy inputs often causes breakages of the nanosheets when exfoliating[171]; a slight ‘chipping’ effect is commonly observed with vibrational-induced exfoliation, specifically breakages of the outer surfaces of the nanosheets exposed to vibrational modes during exfoliation.[172] The collapse of cavitation bubbles near the surface of the dispersed flakes can lead to full scission of nanosheets perpendicular to the basal plane. The generation of such high-energy cavitation forces therefore has a detrimental effect on the lateral dimensions of the nanosheets.[169, 172, 173]. However, cavitation has been shown to exfoliate nanosheets far more effectively than vibration, leading to higher concentrations of dispersed material;[169, 174, 175] at the expense of lateral dimensions. Utilisation of vibrational forces helps maintain sheet size but fails to exfoliate as efficiently.[173] The balance between the two modes will therefore depend on the desired characteristics of the final material.

2.3.1 Factors to Consider for Liquid-Mediated Exfoliation 2.3.1.1 Solvent Selection The selection of a suitable solvent is essential to allow for effective exfoliation to occur. The solvent plays three roles in the exfoliation process: (1) as a medium that can effectively transfer the acoustic power from the sonotrode,[173] (2) as a solvent with the correct parameters that minimises the enthalpy of mixing between the liquid and the interlayer faces of the lamellar crystal, to allow for solvent intercalation[164, 176] and (3) to stabilise the exfoliated sheets by providing a steric barrier to inhibit reaggregation.[172, 177] A number of research groups have contributed to the assessment of a wide range of solvents

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13 and to examine their solubility parameters, both theoretically and in practice.[164, 173, 176, 178, 179] Although other parameters, such as the surface tensions and relative permittivity can be considered to approximately judge the effectiveness a number of solvents as dispersing media,[176] the use of solvents that can closely match the Hansen solubility parameters calculated for the group VI family of TMDCs (general values of δD ∼ 18 MPa1/2, δP ∼ 8.5 MPa1/2 and δH ∼ 7 MPa1/2) feature heavily in the synthesis of TMDC dispersions.[176, 180] Typically, good solvents for the effective exfoliation of TMDC nanosheets were found to be nitrogen-functionalised liquids that are denser than water. Solvents based on the family of Nalkylpyrrolidones, such as N-methylpyrrolidone (NMP)[181] and N-cyclohexylpyrrolidone (CHP),[176] are particularly useful, as they lead to nanosheet dispersions at high concentrations and are known to be stable for months[164, 182] – a important factor when introducing this process to a commercial field. However, the solvents of choice are high boiling and often hazardous, making these solvents undesirable in the consideration of liquid exfoliation on industrial-scales.[164] In an attempt to avoid the use of high boiling solvents, Zhang and colleagues investigated the use of ethanol and water to form stable MoS2 and WS2 dispersions.[178] From both theoretical and experimental data, they confirmed that combining two, benign, low boiling solvents unsuited to nanosheet stabilisation (the Hansen solubility parameters of: water - δD 15.5 MPa1/2, δP 16.0 MPa1/2 and δH 42.3 MPa1/2; ethanol - δD 15.8 MPa1/2, δP 8.8 MPa1/2 and δH 19.4 MPa1/2[177]) produces a mixture better able to match to those values found for the better suited solvents (such as NMP and CHP).[164, 176] It was found that 35 % and 45 % (w/w) ethanol solutions gave the highest concentration of dispersed WS2 and MoS2, respectively (Figure 5). WS2 was exfoliated to a concentration of 0.032 mg mL-1 and MoS2 dispersions reached a maximum concentration at 0.018 mg mL-1. These concentrations 13

14 are significantly lower than those presented by Coleman et al.,[181] however it is the first reported instance of successful exfoliation of either material in an inert, aqueous medium.

2.3.1.2 Sonication Power and Type There are two types of acoustic probes available: an ultrasonic bath or a sonoprobe tip. Studies into their uses for graphene exfoliation in NMP[172] found that sonoprobe tips can generate a higher energy output and the power can be applied directly to the dispering medium, these sonotrodes were found to generate significantly higher concentrations of graphene dispersions at relatively shorter times (~2 mg/mL after 6 hours) compared to ultrasonic baths (0.6 mg/mL after 100 hours); however these dispersions from the sonoprobe were found to contain graphene sheets that were smaller than those observed from bath sonication. The increased power input presumably promotes nanosheet scission. Similar findings were later observed in the exfoliation of MoS2.[164] It seems likely that these effects arise as a result of the relative importance of cavitation and vibration forces from the dispersing medium. The selection of the acousic power input be based on the nature of the TMDC nanosheets dispersions required in its applications. A small number of reports have tackled the refinement of the ultrasonic power employed, using sonoprobes, to optimise the exfoliation. Zhong and colleagues investigated the concentration and morphology of exfoliated MoS2 nanosheets in NMP at different ultrasonic powers of an ultrasonic horn probe. They found a non-linear relationships between both the concentration and the average lateral dimensions of the nanosheets within the MoS2 dispersions with the range of ultrasonic powers applied (Figure 6a and b).[174] The optimal power within their system was obtained within the region of 250 to 320 W; at lower power, cavitation is inhibited, producing less exfoliation energy, whereas at above 320 W the high concentration of cavitation bubbles within the liquid medium suppresses the cavitation forces

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15 generated, effectively shielding the nanosheets to these forces. The largest flake dimensions were obtained at low power but with the trade-off result of producing relatively thick sheets (Figure 6c-f). Similar experiments were carried out by Yang et al.. They found that increasing power increased the concentration of exfoliated material, while decreasing both the thickness and lateral size of dispersed sheets.[175]

2.3.1.3 Sonication Time Many investigations into the effect of reaction time have been carried out on numerous occasions by the Coleman group.[164, 176, 181] Using a number of solvents, the effective exfoliation of a wide range of TMDCs has been documented; by using an ultrasonic horn probe, MoS2 dispersion concentrations as high as 40 mg/mL in NMP can be produced by a prolonged sonication time up to 200 hours,[181] but care has to be taken as any increase in sonication times also leads to the increase in nanosheet scission, reducing both the thickness (closer to monolayer) and lateral dimensions (towards the sub-100 nm scale) of the sheets exfoliated. Optimised concentration/nanosheet size correlations are typically seen after 48-60 hours sonication (average dimensions of 500-800 nm). Similar observations were made using ultrasonic baths,[164] highlighting the judicious control of exfoliation times necessary to control the thickness/size distribution of the exfoliated TMDCs produced, whilst producing appreciable dispersion concentrations.

2.3.1.4 The Use of Additives Further attempts to carry out exfoliation reactions in benign, low-boiling solvent have led to the widespread use of solvent/surfactant mixtures. Such methods have the added benefits of reducing costs and environmental impact.[183] The role of the additive/surfactant is twofold: (1) it modulates the surface tension of the solvent to improve the dispersibility of exfoliated flakes[184] and (2) it adsorbs onto the surface of the nanosheets creating an

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16 additional repulsive force (by electrostatic or steric means) that inhibits reaggregation.[185, 186] Many groups have focused on the use of additives and surfactants to stabilise nanosheets in water and low-boiling solvents (an examples is shown in Figure 7).[187, 188] The type of surfactant does not appear to be overly important as long as it is capable of matching the surface tensions of both the liquid phase and the nanosheet faces within the solid-liquid interface to promote colloidal dispersion.[183, 189] The development of surfactant-assisted exfoliation has led to the investigation of a wide variety of surfactants as additives for the stabilisation of nanosheets. A large variety of small coordinating molecules,[187, 188] surfactants,[190] polymers,[36, 183, 191-195] and copolymers[196] have been used to produce dispersions of many TMDCs in appreciable concentrations. With the correct selection of surfactant, dispersions of many TMDCs can be produced in solvents better suited to scale-up processes.[183] The concentration of dispersed material is highly dependent on the concentration of the surfactant; as the nanosheets are exfoliated, adsorption of the surfactant causes the surface tension of the liquid phase to increase, limiting the extent of exfoliation. The addition of a constant supply of surfactant allows better maintenance of the surface tension of the liquid phase and promotes continued exfoliation and an improved yield.[196, 197] Although the inclusion of surfactant/additives does provide benefits, it can be difficult to remove the stabiliser completely in order to obtain a film of pristine nanosheets, and the surfactant often remains as part of the component within electronic assemblies.[183] This could potentially perturb the nanosheet properties, as well as diminishing their utility in potential applications. Alternatively, rather than addressing the difficulty in surfactant removal, there is considerable scope for the deliberate inclusion of the surfactant molecule in order to complement or even enhance the anisotropic properties of the TMDC nanosheets. For example, the retention of the luminescent properties of restacked WS2 was achieved by 16

17 coating the nanosheets with the polymeric surfactant PEO-PPO-PEO (polyethylene oxidepolypropylene oxide-polyethylene oxide), as a consequence of the increased interlayer distance the nanosheets remain electronically decoupled from one another.[197] This technique of flocculation in the presence of heteromolecules has received attention following its application in stabilising the nanosheets from reaggregation, as well as generating nanosheet-based materials that can be used in a wide array of applications.[168, 198]

2.3.1.5 A Size Selection Process Generally, the nanosheets produced from liquid-mediated exfoliation under optimised conditions (sonoprobe, 48-60 hours sonication in NMP[181]) produces an MoS2 dispersion of around 10 mg/mL; with an average thickness of 3-4 layers. However, these dispersions consists of nanosheets with a very broad distribution of lateral dimensions, with nanosheets ranging from 50 nm to 2 µm within a single sample.[164] Coleman et al. introduced a lateral size-selection regime by a controlled centrifugation process, coupled with sediment recycling, to isolate nanosheets within better-defined size ranges (Figure 8).[181] This approach potentially can be useful, for isolation of nanosheets if specific sizes are required. [181, 199, 200]

2.3.1.6 The Source of Precursor In almost all cases, the source of the precursor is the pure powders of TMDCs which are commercially available. However it has been shown that MoS2 nanosheets can potentially be exfoliated from unrefined, natural molybdenite, obtained from pan-global molybdenum mining localities.[182] The sonication of the parent mineral under the ultrasonic bath conditions in NMP, analogous to the Coleman method[164] was found to produce few layer MoS2 nanosheets of high crystallinity and consistent lateral sizes (200 nm – 1 µm) to those produced from pure MoS2 powder. Although the nanosheets themselves were of high quality,

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18 STEM-EDX mapping revealed that the surfaces of the nanosheets have nanometre scale calcium carbonate aggregates, which are present in the raw material (Figure 9). This is likely to seriously affect the electronic properties of the exfoliated material but may allow considerably cheaper fabrication of high surface-area materials for other applications.

2.3.2 Electrical Properties of Nanosheets by Liquid-Mediated Exfoliation Only a few investigations into the electronic characteristics of solvent-exfoliated TMDC have been reported.[201-203] Specifically, electrical measurements of MoS2 nanosheets indicate a reduction in both ION/IOFF (102 to 104) and µe (0.1 – 0.4 cm2 V-1 s1

),[202, 203] compared with the optimised theoretical values and devices fabricated from

micromechanically exfoliated samples. Among a number of other factors considered, it is thought that the reduction in flake size as a result of scission, and reduction of electrical contact connectivity by adsorbed solvent or surfactant residues may be the primary causes of the reduction in the desired properties. However it is this same adhesion of stabilisers which facilitates exfoliation and allows the formation of colloidally stable dispersions.

2.3.3 Liquid-Mediated Exfoliation: Further Developments The continued development of liquid mediated exfoliation has led to attempts to increase the scalability of nanosheet production. High-shear mixing and jet cavitation have both been shown to be viable alternatives to traditional exfoliation methods. Further work from Coleman et al. has shown the potential of high-shear mixing for the exfoliation of MoS2, MoSe2 and WS2, amongst other 2D-materials.[170, 204] Rather than relying on vibration and cavitation forces to promote exfoliation, the group has harnessed shear forces directly by using both a standard, laboratory high-shear mixer [204] and a common household blender (Figure 10a), to efficiently produce 2D-nanosheet dispersions [170]. This has a considerable impact, as it is a considerably cheaper, more easily scalable process which

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19 has the potential to produce concentrated TMDC nanosheet dispersions in the >100’s L scale (Figure 10b). In view of this, Coleman et al. have thoroughly explored the optimal conditions for the exfoliation of TMDC nanosheets, with a number of parameters considered. The highshear exfoliation method results in the production of (amongst other TMDCs) MoS2 nanosheet dispersions of reasonable concentrations (≤ 0.5 mg/mL), albeit with smaller nanosheet dimensions comparable to those seen by ultrasonic-assisted exfoliation (2-12 layers thick, 40-220 nm across). However with an optimised exfoliation rate of ca. 1.3 mg/min, the high-shear method can produce concentrated nanosheet dispersions in appreciable volumes in as little as 90 min, rather than the 10’s of hours observed before. The exfoliation process, driven purely by cavitation processes has also been explored by Shen et al..[169]The use of a jet cavitation device promotes a high rate of cavitation bubble generation. The apparatus used has been shown to provide sufficient force to exfoliate MoS2 to nanosheets (to monolayer thickness) in reaction times as short as 30 minutes, a dramatic improvement over ultrasonic exfoliation.[181] Due to the combination of a short reaction time and the principle of cavitation shielding, scission is not so significant a factor, as evidenced by the micron-scale nanosheets produced. The researchers suggest that the method is suitable for scale-up to industrial production; however no data on the dispersed concentration or yields of nanosheets have been reported, nor has any statistical analysis of the lateral size distribution been carried out.

2.4 Chemical Intercalation and Exfoliation Intercalation-assisted exfoliation methods work via the insertion of species into the interlayer cavity, thus expanding the interlayer distance and weakening dispersion forces between layers. In some cases, weakening of the van der Waals forces coupled with ultrasonication is sufficient to form TMDC nanosheet dispersions.[205] More often,

19

20 intercalating species which can react to form gases are chosen. The expansion of gas within the interlayer void provides a peeling force to cleave layers, often providing a far higher yield of dispersed nanosheets than is possible with other exfoliation methods. Lithium is by far the most commonly used intercalatant, though acid intercalation and direct insertion of gaseous species has both been explored.

2.4.1 Intercalation with Lithium Over several decades there had been sustained development in the intercalation chemistry of TMDCs in both the fundamental understanding of the materials produced and their potential applications.[49, 206-211] The intercalation of small alkali metal salts, in particular, has become a promising route for the exfoliation of TMDC nanosheets. Several studies have shown that the small lithium metal ion, is by far the most effective intercalatant into bulk TMDCs.[212-214] It is this simple chemistry that resulted in the isolation of singlelayer MoS2 18 years before the discovery of graphene.[215] In virtually all variations on the lithium intercalation and exfoliation method, it is the reaction of intercalated lithium ions with water (aided by ultrasonication) to form LiOH and H2 which creates the driving force for exfoliation; the evolution and expansion of hydrogen gas between the lamellar crystals produces an internally-generated peeling force which allows separation of TMDC nanosheets. Some TMDCs undergo a number of structural and electronic changes as a result of the intercalation process and are often referred to as chemically exfoliated (ce-)TMDCs, in order to distinguish them from nanosheets produced by other methods. For this reason it is important to discuss these changes and their effects, as well as the methods by which intercalation may be achieved.

20

21

2.4.2 Effects of Lithium Intercalation Although the intercalation of small molecules within the interlayers of a TMDC crystal is known to increase the interlayer space (and in turn weakening of interlayer forces within the crystal),[216] lithium intercalation is not a purely mechanical process but also a chemical one (Figure 11);[217, 218] an energy input into the crystal lattice takes place by a single electron transfer process from the metal ion to the lowest lying unoccupied energy levels of the in the metallic d-bands of the MX2 crystal.[219] The effect of this charge transfer on the electronic band structure of the intercalated TMDC crystals is dependent on the resulting d-electron count of the metal centre. Group VI MX2 materials undergo particularly significant changes. When the crystals of MoS2, MoSe2, WS2 and WSe2 are intercalated with n-BuLi,[50, 63, 220-230] both the electron transfer and subsequent coordination of the small lithium cation results in the reduction of the TMDC crystal to a [Li+xMX2δ-]n species.[230-232] The energetic penalty encountered due to the d3 electron count at the metal centre is minimised by forcing a change in the metal coordination geometry to a charge-stabilising distorted octahedral (1T’-)geometry.[220] The subsequent reaction of [Li+xMX2δ-]n with water removes both the metal cations and the formal charge within the nanosheets and promotes exfoliation to give the neutral MX2 nanosheets in the higher energy, 1T’-geometry.

The electronic changes experienced by the group VI TMDCs give rise to significant alterations in the properties of the material.[230] Intercalation (and the conversion to the octahedral geometry) removes the intrinsic band gap allowing for a metallic character in the 1T’-phase. It has also been noted that 1T’-MoS2 is more catalytically active in hydrogen evolution reactions due to its increased conductivity.[231, 233-240] The extent of conversion

21

22 to the 1T’-structure is determined by the extent of lithium intercalation; The 1H-1T’ interfaces that are formed within the partially-intercalated metastable group VI TMDC nanosheets have been noted as unique electronic heterojunctions across a chemically homogeneous layer, indicating a potential route to molecular electronic devices (Figure 12).[222, 241, 242] Unfortunately longer-term advancements in ce-MoS2 have been hindered by the instability of the 1T’-phase.[220] The conversion of sheets back to their native 2H-structure takes place over several days, but may be accelerated by annealing in an inert atmosphere at around 400°C to restore the semiconducting properties of the nanosheets.[243, 244] Therefore ce-MoS2 sheets can potentially be used in semiconductor devices if thermally treated. Due to its conducting nature , no information on the µe or ION/IOFF values of ce-MoS2 have been documented from this class of material.[224, 245] As a consequence of the crystal phase change, ce-MoS2 in the 1T’-phase can be readily dispersed in aqueous solutions[215, 246, 247] which are typically stable for several months;[248] this allows far easier handling compared to 2H-MoS2 which is only readily dispersed in heavy, organic solvents or aqueous/surfactant media. Generally, such obvious changes in the properties of nanosheet are not seen in group IV, V and VII TMDCs; the donated electron density from the alkali metal to the crystal results in a minimal effect on the electronic and crystal structure of the TMDC; the lowest energy system in the [Li+xMX2δ-]n after electron donation is often the retention of the geometry (and therefore the crystal field) of the TMDC before intercalation. For example, the intercalation of ReS2 produces only a slight phase-change, consisting of an small increase in disorder and Re cluster spacing.[249] The only exception to this is the full 1T’- to 2Htransition that is observed in the lithium intercalation of TaS2 crystals.[250] 22

23

2.4.3 Methods of Lithium Intercalation The original process developed by Morrison and colleagues[215, 246] is still widely used as a method to produce single layer MoS2 nanosheets for research purposes (Figure 13).[251, 252] The [Li+xMoS2-] intermediate is formed by soaking bulk MoS2 powder in a solution of n-BuLi under an inert atmosphere at elevated temperatures (80 °C+) for up to four days. A solvent-free method for the intercalation of lithium has been achieved via the heating of a TMDC and LiBH4 under an argon envelope.[231] The intercalated material is then immersed in water and ultrasonicated to produce MoS2 nanosheet dispersions. Although solvent-free intercalation and dispersion has only been carried out with MoS2,[231] WS2[253] and ReS2,[249] the Morrison method has been shown to be applicable to a variety of layered materials,[205] including hydritic-TiS2,[254] TaS2,[255] NbSe2,[256] MoSe2[234] and WS2.[252] These two methods above allow for almost complete exfoliation of many bulk MS2 crystals to nanosheets; far greater yields can be achieved via lithium intercalation and exfoliation than through any other top-down method,[243] with examples often quoting near quantitative conversions of the bulk material to nanosheets. The sheets obtained, normally isolated in the 100’s nm to sub-micron scale in terms of lateral dimensions,[257] tend to be considerably thinner than those obtained by other top-down methods; it is possible to form almost exclusively monolayer dispersions of MoS2.[215] Exfoliated nanosheets have been shown to be stable for months in aqueous and low-boiling solvent dispersions (Figure 14).[258] Dispersions in aqueous media also allows easy functionalization of ce-TMDC surfaces with organic ligands[194, 259] and formation of novel composite materials[260] in much the same way as liquid exfoliated dispersions (see section 2.3.1.4) for numerous applications.[244, 261-263] There are, however, a number of issues with the use of the intercalation and exfoliation method. Many issues with the methodology stem from the 23

24 source of lithium ions: n-BuLi has a number of hazards associated with its use and the intercalation mechanism takes place at elevated temperatures, creating problems for largescale production. The long lithation times necessary has also been cited as an issue in its use.[264] Although MS2 dispersions can be readily produced, lithium intercalation has been found to be less effective for the selenide and telluride families of materials.[181, 265] Difficulties also arise in determining and controlling the extent of intercalation, as the incomplete intercalation has been found to give low yields of single layer materials, whereas over insertion results in the decomposition of the [LixMX2] intermediate to respective metal nanoparticles and elemental chalcogens.[229, 266] Milder conditions for promoting lithium intercalation have therefore been sought. One of the most promising is an electrochemical-assisted process, which allows fine control over the extent of intercalation within the bulk crystals.[245, 267] The method uses lithium foil as the anode and a TMDC crystal as the cathode in an electrochemical cell (Figure 15a). Monitoring the galvanostatic discharge allows the level of the migration of lithium ions from the anode to the lamellar material to be measured directly.[268] The intercalated material can then be ultrasonicated in water or ethanol to exfoliate to single-layer products by the same mechanism discussed in Section 2.4.1. This technique has been widely used to obtain nanosheets of molybdenum, tungsten, titanium, tantalum and zirconium sulfides[245, 269] and for niobium and tungsten selenides.[267, 270] The exfoliations can be conducted at room temperature in a significantly shorter reaction time than the basic intercalation process (completed in hours, rather than days). The use of metallic lithium in the electrochemical process significantly reduces the hazards associated with the intercalation process and permits far greater control over the extent of intercalation, therefore maximising the nanosheet size and yield. Zhang and colleagues have reported NbSe2 crystals up to 6.4 µm[267] and the formation of MoS2 dispersions in which 92 % of the dispersed sheets are 24

25 single-layer.[245] The quantity of nanosheets produced using this method depends on the effective area of the electrode, with test reactions carried out on a 1.1 cm2 copper electrode producing 2 mg of single- or few-layer MoS2 nanosheets with sizes ranging from 500 nm to over 6 µm, depending on the electrical input used. However, upward scaling of this process is limited by the volumetric resistance exhibited in the battery-type cells used.[245] Again, due to the electron-transfer process between the lithium atom and the TMDC nanosheets, the exfoliated products undergo a reversible phase transition to the metallic 1T’-phase, resulting in semiconducting-to-metallic electronic transformation. A slightly modified procedure from the basic lithation process was described by Yang and Zhong.[271] Rather than a lithium ion, their method promotes the intercalation of neutral lithium atoms (sourced from LiOH) in ethylene glycol under hydrothermal-type conditions. It was found that in this reaction that ethylene glycol plays a dual role: as a solvent and as a reductant of the lithium ions in solution prior to intercalation. Therefore the zerovalent lithium-intercalated MoS2 species (Lix0 MoS2) produced is the active intermediate. The reaction of the intermediate with water produces MoS2 nanosheets in high yields, with diameters averaging 200 nm and a uniform, average thickness of 4.7 nm, corresponding to few-layer MoS2. One important feature of this reaction was the retention of the 2Hcrystallinity within the nanosheets, as the reduction of the lithium ion inhibits any form of charge transfer to the nanosheets.

2.4.4 Other Intercalates and Chemical Modification Processes It is clear that the risks associated with the large scale use of organolithium salts present difficulties for the scale-up of intercalation methods. Until 2009 there existed no reports of chemical intercalation and exfoliation of layered materials without lithium, since

25

26 then there has been a sustained drive to intercalate small molecules that can expand when prompted (by decomposition or otherwise) within the interlayers of TMDCs. Zhao and colleagues carried out the exfoliation of MoS2[272] and WS2[273] nanosheets by acid intercalation. Concentrated sulfuric acid and oleum solutions were used to intercalate the interlayers of pre-milled MoS2 and WSe2 powders (by solvent-free NaCl grinding), respectively, for 8-24 h at 90 °C. H2SO4-intercalated TMDCs were successfully produced. After washing, the intercalated nanosheets were redispersed in water (with an additive to aid dispersibility[272]) and ultrasonicated for a short period to promote the decomposition of the intercalated SO4- and OH- groups to generate expanding SO2 and O2 gases. This process was found to produce high yields of MoS2 and WS2 nanosheet dispersions (62 and 90% respectively). A high proportion of the MoS2 sheets obtained were monolayers, whereas WS2 sheets were commonly found as bilayers. However, the average lateral dimensions of these nanosheets were on the order of 60-80 nm, with little material found to be larger than 100 nm in diameter. Acid intercalation is a purely mechanical process and does not induce a change within the TMDC lattice; the exfoliated MoS2 and WS2 sheets retain their 2H-structure. An extension of the above method has shown the possibility of producing large-area films from intercalation methods. Patil et al. have demonstrated that a pellet, produced from the compression of a H2SO4-intercalated MoS2 powder,[274] when in controlled, gradual contact with the surface of water promoted the exfoliation of MoS2 (Figure 16d). The nanosheets remain on the surface of the liquid medium as the surface tension of the water prevents the exfoliated sheets from sinking. The resulting film of free-floating hexagonalMoS2 on the surface of the aqueous layer is comprised of 200 nm to 2 µm wide nanosheets, the majority of which were monolayer. These films can then be transferred to substrates by a simple dip coating process. 26

27 As with lithium intercalation, electrochemical routes for acid intercalation have also been explored. Building on research carried out over thirty years ago,[275, 276] a simple, yet facile acid-assisted electrochemical route was developed for the exfoliation of MoS2 in 2014.[277, 278] The process occurs by applying a controlled, low voltage bias through an electrochemical cell that uses a bulk MoS2 crystal as an anode in an in a sulphate-based (0.5 M H2SO4 or Na2SO4) electrolyte (Figure 16a and c). The voltage bias promotes two reaction stages: the initial low bias promotes the oxidation of water to produce hydroxyl- and oxygen radicals[279] which are thought to react with edge sites and grain boundaries of the MoS2 flakes, creating defect sites at these locations. These defects open space for more efficient intercalation of anionic groups. After a short exfoliation cycle (30 minutes) the average thickness of exfoliated MoS2 nanosheets (with the expected 2H-structure) produced was found to be 1.8 nm (corresponding to 3-4 layers). Lateral dimensions were exceptionally high, ranging from 5 to 50 µm (Figure 16b), the procedure does result in some oxidation of the MoS2 surface (though this has been limited by the optimisation of the process) and the high quality of the exfoliated 2H-MoS2 nanosheets is reflected in the excellent electronic properties of the devices fabricated from these materials. The FET produced shows a µe value of 1.2 cm2 V-1 s-1 and ION/IOFF above 10 6, values which are comparable to those of micromechanically-exfoliated sheets, but the concentrations of the aqueous MoS2 dispersions obtained were in the region of 0.007-0.014 mg/mL (yields of between 5-9 %), corresponding to a low throughput process. Supercritical CO2 intercalation[280] has been shown to be a potentially effective method for the exfoliation of MoS2 and WS2. The procedure involves the introduction of CO2 at 10 MPa pressure into a pressurised reaction chamber with an ultrasonic probe horn containing the lamellar materials to intercalate the supercritical fluid within the interlayer spaces (Figure 17a). Ultrasonication aids CO2 insertion by promoting the atomisation of the 27

28 supercritical fluid. Once the reaction pressure is relieved, the mass expansion of the small molecules within the interlayer surface of the selected TMDC promotes exfoliation. Characterisation shows that a large proportion of the nanosheets are of monolayer thicknesses, with lateral sizes ranging between 0.5 µm and 2 µm. The relatively large size of the products indicates that sonication-induced scission of flakes is minimised by the short reaction time applied from the relatively low power sonotrode. No structural transformations are induced during the procedure, therefore 2H-polymorph is retained. The process was only conducted in with 100 mg loadings of the TMDCs, with discussions on exfoliation yields and rates unavailable (Figure 17b). Ammonia has been used to assist the exfoliation of VS2.[281, 282] Rather than allowing ammonia to diffuse through layers in order to intercalate, lamellar VS2·NH3 is formed directly via the decomposition of Na3VO4 in the presence of thioacetamide which provides both a sulfur and ammonia source. Sonication of this precursor in ice water has been shown to yield a dispersion of VS2 nanoflakes by the expansion of the ammonia intercalate, causing delamination. The thickness of the obtained product is estimated at 4-5 layers, with typical dimensions in the hundreds of nanometres. A

two-step

intercalation

process

to

produce

nanosheets

was

recently

introduced,[213] where hydrazine hydrate is initially used to expand the TMDC interlayer spaces (Figure 18a). It is thought that upon intercalation the N2 H4 undergoes oxidation to N2H5+ which breaks down to give N2, NH3 and H2 gases upon heating, promoting the expansion of the interlayer spacing. The resulting volume of the TMDC crystals (exfoliation of MoS2, WS2, TiS2, TaS2 and NbS2 has been carried out) was over 100 times the volume of the starting material ((Figure 18b and c), which allows the much larger alkali metal naphthalide solution to intercalate successfully to further weaken the interlayer bonding. Because the interlayer space is expanded to such an extent, it is possible to include lithium, 28

29 sodium or potassium cations into the structure (Figure 18d). As a consequence of the increased reactivity with water, sodium-intercalated TMDCs produces exceptionally large, high-quality flakes, of which 80% were monolayer, in near-quantitative yields (up to 90%). Average lateral dimensions of 10 µm found for MoS2 - around 10-15 times larger than those seen with basic lithium intercalation. The other TMDCs tested produced nanosheets 1-10 µm in length. XPS studies showed that sodium and lithium intercalation induce different structural and electronic changes in the host material as a result of their different complexation characteristics. Indeed the nanosheets produced by Na intercalation were found to retain the 2H-phase of the bulk MoS2 crystal. The reason for the retention of the 2Hgeometry was initially unclear. Only recently has the electrochemical intercalation of MoS2 with sodium ions been studied;[283] (S)TEM and in situ XRD investigations indicate that there is a intercalatant concentration-dependence on the phase transition of [NaxMoS2]n: no structural transition is observed from 2H-MoS2 to 2H-[Na0.5MoS2]n, before a phase change to 1T’-[Na1MoS2]n. This is presumably attributed to the more diffuse charge applied from the sodium cation to the crystal. As sodium is less effective at promoting full electron transfer and the TMDC does not undergo the expected crystal phase change. A novel intercalation-exfoliation route to hybrid MoS2/MoO3 nanosheets has been investigated by Joh and colleagues, using uses hydrogen peroxide as the intercalatant species (Figure 18e).[284] Initially, the H2O2 partially oxidises the edge sites and dangling bonds of the (bulk) MoS2 sheet. This expands the interlayer distance at the crystal edges to allow the insertion of H2O2 into the interlayer spaces where further oxidation occurs, leading to the formation of MoO3 nanoparticles decorating MoS2 sheets which can be separated by simple sonication. A similar study has found that the use of a smaller amount of peroxide allows greater control over the extent of oxidation.[264] These milder conditions still cause selective oxidation of dangling bonds and defect sites but the formation of discreet MoO3 nanoparticles 29

30 is inhibited. The partial oxidation of sheets is sufficient to allow separation of nanosheets under stirring or ultrasonic treatment. Under these mild conditions, exfoliated nanosheets were found to be 2-5 micron in length and typically less than 3 nm thick. Optimised conditions lead to a yield of around 60 %.[264]

2.5 Hybrid Exfoliation Techniques There are issues with the micromechanical, liquid-mediated and intercalationexfoliation methods detailed here. Attempts have been made to combine the techniques in a potentially synergistic approach to exfoliation. It is thought that the introduction of a micromechanical component can promote the exfoliation of large-area single- and few-layer nanosheets, whereas exfoliation in a liquid medium can promote a larger throughput and yield, and an intercalation process can reliably produce monolayer nanosheets. A number of micromechanical milling/liquid-mediated exfoliations have been investigated as potential avenues to a hybrid system. A two-step procedure of micromechanical ball-milling of MoS2 in an aqueous surfactant (sodium dodecylsulfate) solution, followed by ultrasonication was found to readily produce aqueous MoS2 dispersions.[171] During the thinning process, the low-energy ball milling does not cause significant damage to the in plane structure of 2D nanosheets and generates fewer defects and impurities as compared to sonication alone.[158] The nanosheets from the hybrid route were found to have lateral dimensions up to 700 nm (a small improvement over those seen in liquid exfoliation alone) and thicknesses between 1.2 nm (bilayer) and 8 nm within a significantly shorter time (3-14 hours). The dispersions produced were found to have concentrations below 1 mg mL-1; although this method does show improvements over standard liquid exfoliation it is unlikely to become widely used.

30

31 A more successful procedure involves the grinding of an MoS2 paste in NMP in a ceramic mortar for selected times (30 minutes to 2 hours), followed by removal of solvent, redispersion in the 45% (w/w) ethanol dispersion system developed by Zhou et al.[178] and subjected to a short sonication step for a further 2 hours to produce a mixture of single- and few-layer MoS2 nanosheets (Figure 19a).[285, 286] This method drastically improves the concentration obtained in the mixed solvent from the 0.018 mg mL-1 reported by Zhou to over 26 mg mL-1, comparable to the concentrations obtained via the Colman method of liquid exfoliation after 100 hours of sonication,[181] but with the significant disadvantage of producing nanosheets with small lateral dimensions of 20-70 nm (Figure 19b and c). Despite significant investigations into the effect of grinding times,[286] and grinding methods,[287] the first serious evaluation of grinding solvents was published only in 2014; Nguyen and colleagues examined the use of a variety of solvents during the initial delamination phase of MoS2 nanosheet production.[288] After the grinding step, the samples were dried and subjected to ultrasonic treatment in ethanol to produce MoS2 nanosheets with lateral dimensions in the 10’s to the 100’s nm. Although NMP was found to be the most effective for producing larger sheets with a high aspect ratio (lateral dimensions/thickness), it is noteworthy that the use of acetonitrile, hexane and cyclohexane as the dispersing medium in the grinding stage allowed the formation of ethanol dispersions with higher concentrations; usually these are considered poor solvents for MoS2 exfoliation (Figure 19d). In 2014, Wei et al. introduced a novel method of micromechanical exfoliation by liquid N2-quenching, followed by ultrasonic-assisted exfoliation.[289] The principle of liquid N2-quenching relies on flash cooling the MoS2 flakes (heated to 80 °C in a KOH solution beforehand). The surface of the hot flakes of MoS2 cools faster than their interior due to poor interlayer thermal conduction. This causes shear stress as a result of rapid contraction of surface layers while also minimising scission and vibrational chipping during the 31

32 ultrasonication stage. This hybrid method results in the production of pristine MoS2 nanosheets with thicknesses ranging from 1.5 (bilayer) to 3.5 nm, and lateral dimensions of 0.5-3.5 µm (Figure 19e) – significantly larger than those isolated from other liquid-assisted methods described above, competing with the micromechanical means in terms of dimensions, but the report is lacking on production yield. Applying the liquid-mediated exfoliation process to a pre-intercalated TMDC [216]) could dramatically increase the production rates of TMDC materials by liquid-exfoliation. The process relies on the addition of a metal salt to a solution prior to exfoliation. Choi and Choi[212] have neatly shown this principle by the addition of an alkali metal hydroxide (LiOH or NaOH) to an NMP solution containing MoS2 powder. After two hours of sonication and removal of the metal hydroxide and unexfoliated material, the yields of the stable dispersions produced were 65-81% from LiOH and NaOH, respectively (Figure 20). These yields are significantly higher than the best values obtained from both the additive-free NMP exfoliation (~2.4 %)[231] and the aqueous dispersion containing a surfactant (~10 %)[190] in similar timeframes. The nanosheets produced from LiOH-assisted exfoliation were found to be well exfoliated, with 85% of measured sheets containing 1-3 layers. Lateral dimensions ranged from 50 nm to 1 µm. Although the process hasn’t been directly observed, it is thought that the alkali metal cations interact in a similar fashion to the basic intercalation mechanism described before; the negatively charged hydroxide ions appears to display some Lewis basic character that ‘open up’ the crystal edge for intercalation by coordinating to the edge-situated Mo ions. Although lithium has been shown to be more effective than sodium as an intercalating cation, we agree with the authors that the use of NaOH will be beneficial, as it allows for the intercalation-exfoliation reaction to be carried out without the need of moisture exclusion.

32

33 A slightly different process was described by Zheng and colleagues,[290] where a selected TMDC powder (MoS2, MoSe2, WS2 or WSe2) was added to an aqueous NaCl or CuCl2 solution, with a surfactant to aid nanosheet dispersion. The water was then removed by boiling and the TMDC/salt/surfactant composite then subjected to ultrasonic treatment. The group reports that the salts insert into the interlayer during the boiling stage, before exfoliation by bath-sonication in water, although there is no direct evidence confirming that intercalation takes place. It is therefore not yet clear whether this is a true hybrid method, or simply the first use of inorganic additives to the sonication-assisted process. Regardless, this method has been shown to provide few-layer TMDC nanosheets with lateral dimensions in the 100’s nm to micron scale, with a 30 fold increase in the concentrations of TMDC dispersions obtained compared to those commonly observed without the uses of salts.

2.6 Destructive Thinning and Etching Techniques All of the exfoliation techniques described in Sections 2.2 to 2.5 are based on the delamination of a bulk TMDC crystal to produce a large number of nanosheets that are considerably thinner. Thinning and etching techniques are considered to be an ‘on demand’ method to further reduce (or in cases modify) the z-dimensions of a pre-exfoliated nanosheet to match the application of the materials required. The destructive removal of unwanted layers from a few-layer crystal leaves a nanosheet with a specific thickness.[291] These methods also introduce nanosheet patterning to generate multilayer heterostructures not available by any other means.[291] A few studies have been carried out on the thinning of few-layer MoS2 (itself produced by a micromechanical method) by the controlled thermal annealing in the presence[292, 293] or absence of oxygen.[294] The procedure relies simply on thermal energy adsorption within the layers of the nanosheets: Poor thermal transfer between layers,

33

34 due to the weak van der Waals interactions, promotes the sublimation of the upper layers, while layers beneath remain intact. These techniques can, in the absence of oxygen, reliably leave monolayer MoS2 adhered to the substrate (given suitable temperature controls; Figure 21a-d). The SiO2/Si substrate acts as an effective heat sink, dissipating excess energy and protecting the bottom layer, with which it is in intimate contact. The monolayer nanosheets produced from these annealing processes were found to have smooth surfaces, but also contained triangular pits (sizes of the holes in the region of 100’s nm), arising from anisotropic etching (Figure 21e-h). The thinned nanosheets retain the 2H-crystal lattice of MoS2. Extended annealing cycles was found to propagate the anisotropic etching process to produce single layer MoS2 meshes (Figure 21i-l).[292] The presence of oxygen in the process during the annealing process has a profound effect on the potential application of the nanosheets obtained. In a similar procedure to that described above, the sublimed MoS2 was found to be oxidised to MoO3 in an oxygencontaining atmosphere. The MoO3 product then deposits back onto the thinned nanosheets produced, converting the nanosheets to a p-type semiconducting material, albeit with poor µe and ION/IOFF measurements (~0.02 cm2 V-1 s-1 and ~103, respectively).[292] This technique was also used in an attempt to promote layer-by-layer thinning[294], but found that an inhomogeneous surface appeared after a short thinning cycle – the example provided thinned an eight-layer nanosheet to a nanosheet with parts thinned to six- and seven-layers, as the anisotropic etching was also found to be more prevalent in thicker nanosheet thinning (Figure 22). The use of the scanning laser on a Raman microscope has allowed the thinning of micromechanically exfoliated MoS2 flakes to monolayer thickness.[295] The process is similar to thermal etching, whereby heat (induced by a scanning Raman laser causes the sublimation of the upper layers, whereas the lower layer (in contact with the substrate – a 34

35 heat sink) remains intact (Figure 23a). This method has been found to rapidly produce monolayer MoS2 at a rate of 8 µm2/min and can be used to produce large area monolayer nanosheets from both CVD-grown and micromechanically exfoliated MoS2 nanosheets (Figure 23b and c). The real time collection of in situ Raman spectra allows for precise control to identify the thinning process, and to confirm the presence of monolayer MoS2. Although laser ablation does not induce oxidation of the lower layers, the resulting samples were found to be considerably rougher than the pristine, exfoliated precursor. This may be due to the presence of adlayer impurities or incompletely removed upper layers. It is thought that the presence of these impurities result in µe values of a thinned MoS2 monolayer being an order of magnitude lower (between 0.04−0.49 cm2 V−1 s−1) than the benchmarks set by the micromechanically exfoliated MoS2-based FETs. Argon plasma has also been used to similar effect in the thinning of micromechanically exfoliated MoS2 samples (Figure 23d-g).[296] Whereas laser thinning relies on indiscriminate ablation to a single layer, this thinning technique relies on controlling the irradiation time to thin nanosheets layer-by-layer. The process works by the irradiation of the top layer with Ar+ plasma (which is monitored by photoluminescence), which causes disorder in the top layer. Continual irradiation dissipates the upper layer until the disordered component is sublimed, signified by an abrupt PL intensity increase in real-time measurements, corresponding to the thickness reduction within the material. The use of a chemical etching agent to form single-layer MoS2 has also been examined.[291] The use of XeF2 promotes the oxidation of MoS2 on the surface of the sheet. Time-controlled reactions was found to promote layer-by-layer thinning on few layer nanosheets to desired thicknesses, although a significant roughening of the surface, compared to the initial micromechanically exfoliated sample, was observed (Figure 23h and i). Although this method indiscriminately thins the whole sample, it has been shown that 35

36 poly(methyl methacrylate) or exfoliated graphene may be used as an etching mask, allowing the formation of thinned regions of MoS2 nanosheets within a sample.

36

37

3 ‘Bottom-up’ Methods for TMDC Nanosheet Growth 3.1 Introduction Bottom-up syntheses require the formation of TMDCs by the reaction of metal and chalcogen containing precursors. These reactions, which can be carried out in the solid, liquid and gaseous phase, have become increasingly popular in the production of TMDC nanomaterials. In almost all cases the aim is to form reactive Mn+ and nX2- ions which react directly to form the desired material. Typically the chemistry is robust and well-known, and there exist an almost infinite number of possible precursors. The novelty (and difficulty) lies in developing sufficient understanding of the energetics and control over the reaction kinetics.[297] Fundamentally, nanosheets are high-surface area, high energy, and unstable structures, therefore this understanding allows researchers to impose reaction conditions which promote the selective formation of nanosheets over more stable low-dimensional structures or simple bulk crystals. The inherent anisotropy of layered MX2 crystals simplifies the process of forcing two-dimensional nanosheet growth (compared to the difficulties in growing non-layered TMC nanosheets), but with the aggressive conditions required (high energy-processes, need of a reducing atmosphere, long reaction times), the growth of other low-dimensional materials (including many-layered nanosheets) must be deliberately hindered (Figure 24). Although gas-phase reactions have been well explored and continue to receive considerable attention there is a distinct lack of literature on liquid- and solid-phase syntheses. This may be due to difficulties in isolating nanosheets from the reaction media while preventing flocculation or aggregation (even if exclusively monolayer product is synthesised, it will be most likely isolated as a restacked network of sheets), though the recent developments discussed below have gone some way to addressing this problem. Gas37

38 phase syntheses clearly do not suffer from these issues, though promoting lateral (epitaxial) growth was initially problematic.[42] Deposition methods now regularly produce sheets with comparable electronic properties to those obtained from micromechanical exfoliation. Below, we discuss the fabrication of nanosheets via gas-, liquid-and solid-phase processes.

3.2 Gas-Phase Syntheses Gas-phase synthesis of 2D TMDCs is a promising research area due to their inherent scalability and compatibility with existing technologies[298, 299]. Reliable growths of single layer nanosheets of lateral sizes ranging from a few atomic units[300] to several centimetres[301, 302] have recently been reported. The maximum size of a TMDC nanosheet formed via top-down methods is limited by the grain size of the starting material, gas-phase syntheses can, in principle, generate nanosheets that are only limited by the substrates used and ingenuity employed.[303] In addition, the current-generation of gas phase-grown TMDC nanosheets exhibit crystal-lattice structures of similar purity to those obtained from micromechanical cleavage (and therefore comparable electrical characteristics) but in nanosheets that are up to 106 times larger[304, 305] The fabrication of such high-quality, large-area sheets has become more common and there has been increased interest in using grown, rather than exfoliated sheets to examine the fundamental properties of TMDC nanosheets.[305] There is a large variety of parameters that need to be considered in the synthesis of TMDC in gas-phase reactions, as the formation, morphology, thickness and sizes of these nanomaterials are hugely dependant on the specific conditions employed in the reaction. The generation of the TMDC nanosheets within gas-phase reactions falls into three relatively broad categories: The two major methodologies are the physical vapour deposition (PVD) and chemical vapour deposition (CVD) reactions (Figure 25a and b). PVD and CVD

38

39 modes generate the TMDC monomer units in the vapour phase required for production, but in differing ways. Regardless of the TMDC vapour source, the nucleation and growth characteristics, as well as parameters governing nanosheet size, thickness, morphology, orientation, stoichiometry etc. appear to follow similar trends. It therefore makes little sense to examine the two methods individually. The third method is less common but nevertheless has been well investigated, and relies on the direct sulfurisation of metallic and metalcontaining thin films.

3.2.1 LP-PVD and CVD Methods 3.2.1.1 TMDC Vapour Generation PVD methods rely on the sublimation of selected bulk TMDC materials (most commonly in powder form) by heating under reduced pressures. The requirement of high temperatures (800-950 °C) and low pressures (< 100 Torr) are common, in order to volatilise appreciable quantities of the TMDC monomer in the reaction chamber.[306, 307] Such methods are simple and highly reliable routes to the generation of Group VI TMDC, including MoS2[308-310] and WSe2[311] vapours, whereas other groups have rarely been studied. TMDC vapour production by CVD is generally based on the volatilisation and thermolysis of precursors containing the required elements. The reactive metal and chalcogen ions generated in the gas phase can then react to form the TMDC monomer in the gas phase. A well-known example in the formation of TMDCs is the reaction of a sulfur vapour (produced directly from elemental sulfur) with metal oxides. In particular, Group VI oxides (such as MoO3 and WO3) are routinely selected as cheap, easily handled precursors with relatively low melting and evaporation temperatures.[312] The MO3 sources are easily volatised to their suboxide MO3-x,[313, 314] making the nanosheet growth by CVD means 39

40 accessible in both low pressure (LP-) and atmospheric pressure (AP-) systems; as discussed below, control over pressure can maximise the potential of CVD systems to generate the desired large-area monolayer films. Further reaction with sulfur-rich vapours converts the gaseous metal suboxide units of the desired TMDC in the gas phase. Although this procedure has mainly been investigated for the production of MoS2,[305, 315, 316] the growth of MoSe2,[304, 317, 318] WS2[319] and WSe2[320] nanosheets have also been formed by similar means. Other CVD based monomer generation procedures have used different precursors: The synthesis of NbS2,[321] MoS2[322] and WS2[323] nanosheets was achieved by the vapour-phase reaction of their respective metal chlorides with sulfur vapours at 800850 °C. A single-source precursor route has also been developed for the formation of MoS2; the decomposition of [H4N]2[MoS4], within a DMF aerosol, has been found to reliably form MoS2 nanosheets on graphene substrates.[307]

3.2.1.2 Deposition of TMDC vapours on substrates Once the TMDC monomers have been generated in the gas-phase, these vapours can be deposited on substrates within the reaction cool zone. Transport of monomers is facilitated by an inert carrier gas.[305, 306, 312, 313, 324-326] The nature of deposition directly determines the morphological characteristics of the resulting TMDC nanosheets. The basic premise for nanosheet growth from gas-phase reactions combines the Vapour-Solid and the Frank-van der Merwe growth mechanisms,[297] where the TMDC vapours initially deposit and form nanocrystals by random nucleation on the substrate surface After nucleation, the gaseous source will tend to diffuse to the edges of the nuclei and the crystal will prefer to grow along the lateral dimension of the nanocrystal, leading to the formation of discreet nanosheet islands on the substrate surface. Continued growth eventually results in the coalescence of nanosheet islands to produce continuous films with clear grain boundaries (Figure 25c and d). The high growth anisotropy arises due to the relative chemical anisotropy 40

41 of the atoms in the TMDC nanocrystals at the edges, compared to its surface;[321] The edge atoms have dangling bonds whereas the surfaces are passivated with a layer of relatively inert chalcogen atoms. This growth profile often leads to the formation of highly symmetrical triangular domains with energetically favoured zigzag edge-terminal structure.[305, 306] The growth profiles of group V TMDC nanosheets have been examined in detail. Their growth anisotropy increased by the position of electrostatic potential within the sheet structure; recent calculations have shown that the atomic charges are located at the vertices of a growing TMDC nanosheet significantly exceeds that of the charges located edge sites and the basal plane by approximately 33% and 400% respectively.[327] Therefore group V TMDCs grow such that the distance between vertices is maximised, resulting in highly symmetrical triangular sheets. As the properties of these novel materials are largely dependent on quantum confinement effects, it is essential to develop methodologies by which one can selectively produce sheets with specific thickness, size and crystal quality to match their potential applications.[328, 329] It is therefore necessary to develop a full understanding of the seeding and growth profiles of TMDC nanosheets. Perhaps the most important factor in the usefulness of the gas-phase growth strategy is the orientation of sheets on a substrate. TMDCs can grow in either a perpendicular (edge-maximising) or parallel direction (epitaxial), with respect to the substrate. TMDC films grown in these differing orientations exhibit vastly different properties.[321] In the simplest of cases, deposited materials will grow perpendicular to the surface of the substrate in order to minimise the energetic penalty experienced in the interactions of the two surfaces,[42] unless specific experimental parameters are chosen to reduce this effect and obtain laterally-oriented nanosheets.[42, 312] In this situation, a number of conditions are required to promote epitaxial growth of TMDC

41

42 nanosheets on substrate. Below we examine each variable in turn in an attempt to understand the processes. 3.2.1.2.1 Promoting Epitaxial Growth on Substrates It has been noted that the selection of the correct substrates is critical, as poorly suited substrates can produce TMDC films with vertically-aligned nanosheets,[321] or poorly formed or small epitaxially grown nanosheets[320, 330] It seems that 3 major factors must be considered in order to minimise the energetic penalty of the lateral growth of a nanosheet.[307, 313, 331] (1) Obtaining a substrate with an atomically flat and inert surface, free of dangling bonds. (2) Minimisation of the atomic-lattice mismatch between the substrate and TMDC surface. (3) The promotion of van der Waal interactions between the substrate and crystal along the basal plane. Zhang, Liu

and

colleagues have explored the use of fluorophlogopite

(KMg3AlSi3O10F2) mica substrates to promote growth of monolayer MoS2 by the sulfurization of MoO3 powder by LP-CVD,[313] as the mica is considered to be a van der Waals epitaxy substrate for nanosheet growth. It has a pure-2D surface and a lattice mismatch with MoS2 of just 2.7 % (Figure 26a). This allows the growth of extended MoS2 nanosheets. In addition, several groups have reported successful lateral TMDC nanosheet growth on sapphire, despite relatively a high lattice mismatch.[311, 332] A recent LP-CVD-promoted MoS2 deposition study has found that although the lattice mismatch between the two components was significant, the in-plane periodicity of the Al2O3 lattice is almost exactly 1.5 times greater than that of MoS2, allowing the formation of a (3 x 3) MoS2 on (2 x 2) sapphire superlattice that aids epitaxial growth (Figure 26b).[331] This observation has yet to be unambiguously confirmed, but limited DFT calculations and the propensity of the MoS2 domains to align non-randomly on the substrate (as reported by Ji et al.) support the theory.[331] 42

43 Promotion of van der Waals dispersion forces is often achieved by the use of 2D materials as templating substrates to generate bilayer vertical heterojunctions. [333] Large area MoS2 flakes were grown on CVD-grown graphene on a copper foil[307] and highly orientated pyrolytic graphite (HOPG),[332] at 400 °C by LP-CVD to form MoS2/graphene hybrid structures. In addition, single-layer WS2 nanosheets have been grown by the AP-CVD process on micromechanically exfoliated h-BN flakes with its WSe2 analogue grown by LPCVD on epitaxially grown graphene nanosheets on SiC.[323, 334] In the examples descried above, the deposited MoS2, WS2 and WSe2 flakes grew parallel to the substrate despite a relatively high lattice mismatch (up to 23% in WSe2/graphene heterostructures).[307, 323] It has been shown that in this instance the h-BN sheet edge sites provide better nucleation for the growth of WS2, presumably due to π-π stacking effects similar to those seen in graphene substrates.[323] Similarly Ago et al. have found that AP-CVD-grown NbS2 nanosheets tended to grow perpendicular to the SiO2/Si substrate, whereas CVD-grown graphene on SiO2/Si promotes lateral sheet growth, as the atomically smooth surface and both the van der Waals and π-interactions facilitate NbS2 vapour adsorption that in turn promotes the growth along the graphene surface.[321, 335] It seems that a 2D substrate, with unsaturated faces that are free of dangling bonds, allows favourable columbic and π-stacking interactions for the deposition of further, layered products. In the absence of templating substrates discussed above, seeding promoters are often used to impart greater uniformity and crystallinity on deposited films and to promote lateral growth on non-templating substrates (such as SiO2/Si). The addition of large, aromatic molecules (such as perylene-3,4,9,10-tetracarboxylic acid tetrapotassium salt (PTAS) and 3,4,9,10-perylene-tetracarboxylicacid-dianhydride (PTCDA)) and reduced graphene oxide to the reactions, has been found to be generally as useful as seeding promoters for the growth of MoS2 monolayers in both LP- and AP-CVD formats (Figure 27).[312, 336] Further work 43

44 surveyed a host of aromatic seeding promoters, including metal-bearing macrocyclic ligands, can facilitate the growth of large area monolayer MoS2.[324] The use of such molecules, provides further evidence for the potential of π-stacking effects to direct lateral growth on the often non-templating substrates. The inclusion of seeding promoters is thought to have two effects: (1) Increasing the adhesive force of the substrate to the MoS2 monomer, promoting the anisotropic growth of an MoS2 monolayer over thicker nanosheet islands mediated by adatom cohesive forces and (2) offering a heterogeneous nucleation site to promote formation of an MoS2 nucleus with a lower energetic boundary than that for homogeneous nucleation.[324] The concentration of the promoter on the substrate is crucial to the formation of large-area ultrathin sheets over smaller, thicker materials.[324] Conversely van der Zande and colleagues have indicated that in the fabrication of MoS2 films with poor control over the concentrations of seeding promoters can reduce the quality of TMDCs nanosheets produced, though the reasons for this were not explored.[305] Despite the findings above, a number of reports have described the growth of single layer TMDCs on substrates that are not generally considered to be suitable templates for promoting epitaxial growth in the absence of templating/seeding agents.[309, 317, 325, 332] These articles describe the growth of MoS2 and MoSe2 nanosheets (by LP-PVD, AP-CVD and LP-CVD means) on glass and SiO2/Si substrates, which do not contain atomically smooth surfaces or promote dispersion interactions. It is unclear why the nanosheets grow parallel to these types of substrate; though some studies have shown that, in the absence of templating substrates, small bilayer or multilayer patches are often present within the grown nanosheets, which could act as a nucleus to promote the lateral growth.[305] Some reports also indicate that pre-treatment of the substrates with a Piranha etching solution or by exposure to O2 plasma (also see Section 3.2.1.2.3) [305] before use may also promote epitaxial growth on SiO2 and other substrates.[337] The mechanism by which the growth of 44

45 TMDCs on treated substrates occurs is not well understood, though it is likely that the surface is activated by the presence of additional hydroxyl groups after such a treatment.[338] 3.2.1.2.2 The Mechanism for Nanosheet Growth To fully understand the growth process, Cao et al. have gone some way in determining the mechanism of deposition in TMDC systems by studying the deposition of MoS2 nanosheets by the LP-CVD reaction of MoCl5 and sulfur gases (Figure 28).[332] They found that in these gas phase reactions, the rate limiting step is the precipitation of MoS2 monomers (MoS2(g)
MoS2(s)) onto the substrate. This itself is primarily governed by the relative partial pressure of the MoS2(g) gas and the vapour pressure of the growing MoS2(s) nanosheet. Thermodynamics dictate that in order to achieve net growth of nanosheets, the partial pressure of the gas must exceed that of the vapour pressure. Precise control of the partial pressure of the gaseous species, with respect to the vapour pressure of the precipitate, is imperative. Too large a difference in pressures leads to monomer saturation in the cool zone, followed by rapid and uncontrolled deposition of thicker nanosheets, and the loss of epitaxial-growth control as the energetic penalty of the interaction with the nanosheet surface is overcome.[42] Therefore control over the rate of precursor sublimation, the overall pressure of the system, carrier gas flows, and the substrates used can produce variations in both the partial pressure of MoS2(g) and the vapour pressure of the growing nanosheet to dictate the nature of the nanosheets produced. Below are the conditions that have been applied to observe the characteristics of the nanosheet growth, when parameters described earlier have been met. 3.2.1.2.3 Controlling Nanosheet Thickness In the synthesis of monolayer films, the presence of bilayer, trilayer and thicker nanosheets are effectively impurities and can adversely affect the properties of the material as 45

46 a whole.[308] In addition, control of the desired crystal lattice of MoX2 (X = S, Se) during the growth process of few-layer nanosheets is lacking, as both the rotation of the adjacent crystal lattices[306] and the mixture of 2H- and 3R-stacking[317] have been observed. Exerting close control over thickness was initially difficult in gas-phase syntheses; however it is now possible to reliably and selectively produce crystalline films of monolayer and fewlayered MoS2 under gas-phase growth conditions.[332, 339] Thickness control is generally exerted in one of two ways: (1) by selecting conditions which allow such reliable growth that it is possible to stop the reaction when a desired thickness is obtained. (2) The development of self-limiting synthetic conditions for the selective formation of monolayer sheets. Maintaining a controlled lower pressure in the reaction chamber is favourable for obtaining monolayer, or few layer films; ambient-pressure reactions have been shown to produce TMDC nanosheets however, control over partial pressure is inherently more difficult.[321, 339] This was neatly shown by Cao et al. by demonstrating that the overall pressure of the system can alter the partial pressure of the gaseous MoS2 in the precipitation step (Figure 29). Reaction pressure modulation leads to a wide range of MoS2 nanostructures, ranging from micron-scale nanosheet islands that consist of many layers, to uniform monoand few-layer films.[332] They also provided similar results in thickness control by varying the amount of the volatilised molybdenum precursor available in the reaction (by altering the amount of MoCl5 placed within the reaction chamber) prior to sulfurisation. Duan et al.[340] have shown that increasing the Argon gas flow within a LP-PVD reaction for the growth of MoSe2 from 100 to 200 sccm generates a higher vapour pressure of WSe2 during the precipitation stage, which in turn overcomes the crystal growth anisotropy and produces nanosheets that are generally thicker than those obtained at a lower carrier gas flow rate. It is also found that the interaction of the TMDC with its supporting substrate during growth was found to play a critical role in the self-limiting growth of the nanosheets, granting 46

47 access to nanosheets with controlled mono- or bilayer thickness. This is dictated by the interaction of the MoS2 layer with the selected substrate affecting the vapour pressure (the capability of TMDC units escaping from the solid to the gas phase) of the nanocrystal. In a number of cases, this interaction between these two surfaces has been found to decrease the vapour pressure of the surface of the nanosheet, allowing for growth in the z-direction. This interaction with substrates is expected to relax with the layer number, resulting in the increasing of the vapour pressure in the TMDC nanosheet with increased thickness. This has been shown in practice in the growth of MoS2 (by LP-CVD means) onto a HOPG and sapphire substrates under identical experimental conditions; the MoS2 nanosheets were found to be monolayer on sapphire, whereas the one grown on HOPG was bilayer.[339] Other work has shown that the use of fluorophlogopite mica at a growth temperature of 530 °C promotes the exclusive growth of monolayer MoS2 films only.[313] Excess vapours contribute to the production of a small amount of 0D and 1D MoS2-nanomaterials on the surface of the monolayer, but presence of the low dimensional impurities rather than additional film layers shows that nanosheet growth on the substrate is reliably limited to a single layer. Increasing the growth temperature to 700 °C significantly increases the partial pressure of MoS2 in the system to produce ~ 7.5 nm thick MoS2 films.[313] This would imply that the vapour pressure on the surface of the newly grown monolayer films on sapphire and fluorophlogopite mica, and bilayer films on HOPG, has increased sufficiently to inhibit further z-growth, with the differing thicknesses dictated by the properties of the substrates used.[332] The resulting nanosheet growth on these substrates is thought to follow the Stranski-Krastanov growth mode, with a critical layer number of 1 (for monolayers) or 2 (for bilayers).[341] Interestingly, the variation of the growth modes with different substrates does suggest a non-innocent interaction between the grown nanosheets and the substrate is prominent. This prospect was partially confirmed by DFT studies of the interaction of a MoS2

47

48 nanosheet with a sapphire substrate.[331] The calculations indicate that there is significant charge transfer between sapphire and MoS2 sheets and very little contribution from dispersion forces. In contrast MoS2/metal and MoS2/graphene interactions are typically driven by pd[342] and S3p-C2p[343] orbital hybridization respectively. It has therefore been found that many measured properties and potential applications of the supported nanosheets will exhibit a significant contribution from the substrate used. For example, it has already been shown that of MoS2/graphene composites show significantly different catalytic activity from either of the parent materials.[343] Recently, high-quality, monolayer MoS2 nanosheets, from the LP-CVD reaction of MoO3 and S vapours, can be reliably grown onto a polycrystalline gold film (Figure 30a).[314] In this case, it appears that the substrate has a more active role in promoting epitaxial growth. The surface gold atoms on the film absorb a significant proportion of the sulfur precursor in the gas phase, limiting the availability of the chalcogen source to only allow conversion of the metal suboxide in the plane of the of the chalcogen source; i.e. the plane of the substrate. This restriction of the chalcogen source inhibits the formation of MoS2 nanocrystals thicker than a monolayer, therefore assuring the growth of nanosheet islands (up to 50 µm wide) and films that are exclusively monolayer (Figure 30b-e). Layer-controlled LP-CVD growth of MoS2 by pre-treatment of the SiO2/Si substrate with an oxygen plasma prior to TMDC nanosheet growth has recently been reported.[344] Selected lengths of oxygen plasma treatments (for 90, 120 and 300 sec) of the substrates prior the LP-CVD process generates nanosheet films (in the cm2 scale) of mono-, bi- and trilayer MoS2, respectively (Figure 31). This suggests that a simple modulation in the partial pressure of the SiO2-supported MoS2 nanosheets is possible. This self-limiting phenomenon is thought to be dictated by the increased interaction of the MoS2 nanosheet with the increased Si-O or SiOH bonding nodes on the substrates surface, similar to the proposed effect of piranha 48

49 solution on SiO2 substrates.[338] This interaction subsequently lowers the localised vapour pressure of the nanosheet to facilitate growth in the z-direction. However there no examination of large-area film uniformity has been presented, therefore it is unclear how plasma treatment of SiO2 affects the crystallinity of grown films. 3.2.1.2.4 Controlling Nanosheet Dimensions After establishing the need to control the vapour pressure to facilitate the growth of mono- or few layer TMDCs on substrates, it is necessary to maximise the lateral dimensions of the nanosheets produced, whilst keeping within the vapour-phase parameters to retain the precise layer number. The simplest of means to control nanosheet size is by limiting the reaction time, therefore limiting the TMDC vapour source. This is clearly observed by Zhang and Liu in the LP-CVD growth of MoS2 on fluorophlogopite mica.[313] Increasing the growth time from 35 min to 2.5 h results in the formation of nanosheet islands with edge lengths increasing from 0.32 µm to 2.66 µm, while longer reaction time results in the full coverage of the substrate with a uniform MoS2 monolayer (Figure 32a-d). The high nucleation density observed in the growth of MoS2 on the mica limits the island crystallite size; therefore the procedure produces polycrystalline monolayer films rather than macroscopic single crystals. It is evident that within a continuous nanosheet film, the crystallite size is inversely proportional to the number of nucleation sites. As nanosheets nucleate and then grow until island domains merge (forming grain boundaries), fewer nucleation sites leads to a relative increase in the length of the growth stage, resulting in fewer, larger crystallites. Brief, but important discussions on the conditions for optimising nucleation density in the LP-PVD and LP-CVD reaction of MoS2,[314, 325, 345] WS2[319] and WSe2[320] on non-templating substrates have appeared. Controlling nucleation density in the growth 49

50 process is determined by a large number of conditions, including: temperature, the use of reducing gas, flow rate, partial pressure of the precursor at the surface at the substrate and total pressure of the system. Anecdotally nucleation is most favourable on imperfections (scratches, substrate edges, dust particles and rough areas) on the substrate surface. This observation has led to the use of patterned SiO2 substrates to promote nucleation at regular intervals, allowing continuous film growth between nucleation sites (Figure 32e).[306] In the first instance, fine control of the reaction temperature is needed to promote the growth of larger nanosheets, as small temperature variations can dramatically alter the nucleation density and, therefore the lateral sizes of the nanosheets grown, as well as facilitate growth of thicker nanosheets. In a number of cases, it is clear to see that the lowering of the growth temperature, typically results in a higher nucleation density.[334] Examples include the growth of 2-5 µm or 10-50 µm WSe2 at 750 or 850 °C respectively[320] and the small temperature changes required for the growth of many 1-3 µm MoS2 nanosheets, or fewer, larger (~15 µm) domains by adjusting the growth temperature from 935 to 960 °C (Figure 33).[345] There is, however, an example where the reaction temperature indirectly dictates the growth of MoS2 nanosheets on Au, as a consequence of the adsorption amount of sulfurised precursors upon the substrate surface (see Section 3.2.1.2.3).[314] The lateral dimensions of MoS2 nanosheets obtained from this process showed a progression from ~200 nm to ~55 µm, to ~30 µm by performing the reaction at 530°C, 750 °C and 850 °C, respectively. The reduction in size was thought to arise from the dramatically decreased local concentration of the sulfurized precursor at the higher temperatures. Duan and colleagues have provided a large amount of quantitative data on the effects of temperature on nucleation density and nanosheet dimensions.[340] They note that in their LP-PVD reaction that monolayer domains of WSe2 formed at 750 °C nucleated at a 50

51 density of 4200 sites/mm2 and grew only to a size of around 300 nm (edge length). Increasing the substrate temperature by just 15 °C increased edge length to around 10 µm and reduced the nucleation density by almost 75 %. Increasing the temperature further resulted in reduced nucleation density and increased domain size, as well as the formation of bi- and multilayer sheets (Figure 34a-d).

The use of hydrogen in LP-CVD reactions has also been explored. Its functions are threefold: (1) removal of carbonaceous deposits from breakdown of precursors,[346] (2) to avoid oxidation of MoS2 by residual oxygen and unavoidable O2 leakage[347] and (3) as a more effective reducing agent (with respect to sulfur), promoting more effective reduction of the metal oxides (therefore promoting a richer MO3-x-rich environment for sulfurisation and subsequent precipitation).[320] Ajayan[304] and Xiong[317] have reported that the change in H2 flow rates can also alter the nucleation density of MoS2 vapours; Fu and Zhang discovered that the slow increase of H2 flows during the LP-CVD reaction of MoO3 and S2 vapours garners MoS2 nanosheets with larger lateral dimensions grown on Au (of up to 80 µm). They also show that excess H2 within the reaction chamber results in the etching of the alreadygrown nanosheet islands reducing the overall size of the nanosheet produced (Figure 34ej).[348] It was also shown that that the use of hydrogen in the synthesis of MoSe2 nanosheets grown by LP-CVD is essential; the lower chemical reactivity of elemental selenium generally means that the low partial pressure of MoSe2(g) cannot overcome the vapour pressure of MoSe2(g) on SiO2 substrates in the absence of a reducing gas.[304, 317] Other reports of triangular TMDCs nanosheet islands and films by controlling both the overall pressure of the reaction system and the partial pressure of precursors within the reaction chamber during the sulfurization of metallic species by CVD.[306, 325] It was found 51

52 that varying the overall pressures in the reactor system can be achieved by modulating the rate of sulfur volatilisation.[306] As a result, a pressure dependency of MoS2 growth was observed, with respect to both the density of nucleation and single-crystalline growth rates were affected by the pressure of the system. As the pressure within the system increases from 3.8 to 900 Torr, less MoO3 is volatilised, slowing and ultimately stopping the sulfurization and deposition processes (Figure 35a-f). At reaction pressures of 7.5-30 Torr the growth of small triangular crystals is observed. These extend to large area films in the pressures range of 30-75 Torr. At these pressures both supplies of MoO3 and sulfur are sufficient for large area growth, resulting in the coalescence of triangular domains. At slightly higher pressures, 75-300 Torr sulfur is plentiful but lower evaporation of MoO3 limits the supply of Mo leading to isolated but very large, 50-80 µm, triangular islands. At higher pressures, 300-600 Torr, these isolated triangles shrink in size and their morphology begins to change. At the highest pressures investigated, 600-900 Torr, the lack of the precursor sources in the atmosphere results in the growth of star shaped islands; a possible indication of substoichiometric nanosheet production (see section 3.2.1.2.5). Similarly, the temperature at which the MoO3 source is heated at (and therefore the rate of volatilisation of the metal oxide) under LP-CVD-type conditions is found to lead to MoS2 nanosheet films with size-selective grains.[325] The reduced molybdenum feedstock promoted by the lower MoO3 evaporation temperatures reduces the probability of generating nucleation sites on the SiO2 substrates, therefore allows the growth of the MoS2 nanosheets grains from an average of 20 nm (at 650 °C MoO3 evaporation temperature) to up to 600 nm (at 540 °C evaporation temperature) (Figure 35g-n). 3.2.1.2.5 The Stoichiometry of the Nanosheets The production of TMDC materials via a bottom-up process relies on the formation of a species with a 1:2 metal-to-chalcogen ratio, and a number of articles in have shown that 52

53 obtaining the correct stoichiometry can be a difficult task. Hersam and Lauhon was the first to consider the importance of the precursor feeds within a MoO3/S-based LP-CVD system and generated nanosheet islands with a controlled stoichiometry.[349] By controlling the volatilisation rates of the sulfur source by temperature and exposure restrictions within the reaction, they were able to produce three groups of nanosheets that exhibited significantly different physical properties (Figure 36). XPS revealed that the standard procedure, carried out with the sulfur source at 170 °C produced stoichiometric, intrinsic MoS2 (I-MoS2). Reducing the temperature led to the formation of a sulfur deficient structure, (D-MoS 2) with MoO3-x impurities. Further depletion of sulfur content was found to produce deformed MoS2-x nanosheets with significant MoOx content (MoS xOy). It was found that reducing stoichiometry produces a profound effect on the photoluminescence characteristics of the flakes produced. These observations highlight the need to fully understand the composition of the nanosheets as the reducing sulfur content in the nanosheets results in a negative effect on both the uniformity and intensity of its intrinsic properties, such as their photoluminescence profiles.[349, 350] 3.2.1.2.6 Understanding the Structural Defects in TMDC Nanosheets Grown in the Gas Phase. The variables in the gas-phase growth processes discussed above have been analysed thoroughly, it is necessary to critically evaluate the quality of nanosheets produced by the gas-phase reaction and growth mechanisms. The large area nanosheet films, produced using the gas phase reactions, appear to be uniform and highly crystalline when imaged by macroscopic methods. However when studied microscopically,[305-307, 313, 351] it is clear that that the films are comprised of many monolayer grains, TMDCs islands, with coalescence at the edges of the grains.[306] In 53

54 both nanosheet films and islands it is the grain boundaries that are one of the major sources of defects controlling properties in these materials. The structure of grain boundaries needs scrutiny, in order to understand how these features form and how their effects can be minimised (or utilised) to produce tailored macroscopic 2D-materials.[305, 306] A number of MoS2[305, 306, 352] and MoSe2[317] nanosheets, grown by CVD means, have been studied by atomic-resolution electron microscopy (HAADF-STEM, ADFSTEM) in order to provide images of grain boundaries in both nanosheet islands and films. In these separate studies, it was confirmed that three classes of boundaries are often present: Grain boundaries are normally formed as two growing monolayer domains coalesce and form chemical bonds that terminate growth at these coalescence points. The atomic configuration at the grain boundary is often strained, resulting in atomic displacements to allow for such chemical interactions between TMDC grains. This results in the formation of dislocation cores at these boundaries. The nature of the boundary, consisting of unusual inorganic ring networks, is dependent on the twist in geometry between the merging nanosheet grains (Figure 37a and c). Tilt boundaries can have merged grains with a tilt anywhere from 0 < x < 60°; the grain boundaries exhibit different dislocation core compositions and structures depending on the precise tilt-angle (Figure 38a-e). Alternatively grain boundaries can arise when growth of two grains proceeds from a single, imperfect nucleation site; promoting the growth of mirror-twin polycrystals with grains composed of 180° rotated triangles (termed twin-boundaries; (Figure 37b and d). Twin-boundary crystals are also interconnected by ringnetworks, but due to the fixed grain tilt of 180°, the shape and nature of these ring-networks are found to be of a more consistent form (Figure 38f and g). Finally, on the interaction of two grains, one grain could continue to grow on top of the other without forming chemical bonds between the two grains, resulting in the formation of a bilayer region (bilayer overlap; Figure 37e, Figure 38h). This overlap is a little observed process as the formation of a grain 54

55 boundary is favourable. Thorough discussions on the mechanism of the formation of the grain boundaries are available [305, 306, 352]. The presence of grain boundaries have been found to cause significant distortions on the properties of the material as a whole. van der Zande et al. discussed how the presence of the monolayer grain boundaries can affect the photoluminescence and electrical properties of the MoS2 nanosheets (Figure 39a-d).[305] The twin-boundary interface strongly modifies the photoluminescence with a 50 % quenching of intensity, as well as a slight blue-shift and broadening of the peak energy (by 8 and 5 meV, respectively), whereas a tilt boundary (of 40°) shows a 100 % enhancement in emission strength, as well as a blue-shift of the signal and broadening of the peak energy (26 and 5 meV, respectively). Both grain boundaries also cause a shift in the characteristic MoS2 Raman bands (Figure 39e and f). These changes of properties at the grain boundaries are thought to occur as a consequence of localised doping: The 4|8-dislocation ring motifs in the 180° twin-boundaries are molybdenum-rich whereas the 5|7-motifs commonly seen in the 40° tilt-boundaries have excess sulfur. The relative quantities of these structures therefore provide a tendency toward n- and p-doping respectively at the grain boundaries. The strain of the boundaries may also contribute to bandgap modification, or cause the boundary region to lift off the electrically disordered SiO2 substrate to enhance photoluminescence emissions. Nonetheless, the inclusion of these grain boundaries causes an impurity of the optical and photoelectronic properties of the nanosheets in the macroscale. As well as affecting the intrinsic properties of the material, the grain boundaries are thought to exhibit a higher reactivity to oxidation. Zhang,[319] Lee[353] and Zhang[325] have independently developed simple (but destructive) methods to image the grain boundaries in monolayer group VI chalcogenide nanosheet islands and films. These methods rely on the mild oxidation of the films under moisture-rich, ambient conditions, promoted by 55

56 mild thermal or UV radiation treatments, to allow imaging of grain boundaries (Figure 40a). The images provided by the Zhang et al.[319] confirm that the irregular shapes observed often observed in the growth of nanosheets islands are indeed different WS2 nanosheets grown from a single, imperfect nucleus (Figure 40b-m). This points out that in the synthesis of TMDC nanosheets, the controlled growth of the nuclei, an area which has received little scrutiny, is imperative to produce pristine nanosheets that are free of defects. The work by the Ly[353] and Zhang[325] research groups also confirm that the WSe2 and MoS2 films produced consist of nano-to-micron scale grains that are chemically bound together by these grain boundaries (Figure 40n-s). The second harmonic generation (SHG) of TMDC nanosheets, strongly affected by the crystallinity and symmetry of the material, can be used to identify different crystal orientations within a monolayer film that appears uniform under optical microscopy. This method is seen as a non-destructive method to give both quantitative and qualitative information of the nanosheet films produced. In the example described by Zhang et al., the monolayer MoS2 film grown by the LP-CVD reaction of MoO3 and S vapours, were found to contain a number of 20-50 µm irregularly-shaped grains, shown by the reduced SHG intensity at these grain boundaries.[354] Point defects within 2D nanosheets are noticeable in all gas phase-based productions, as imperfections in the growth processes are inherent to the methodology.[352, 355] The Zhou[352] and Jin/Ji/Yuan[356] groups have provided an in-depth analysis of the point defects commonly observed in CVD-grown MoS2 nanosheet islands by ADF-STEM (Figure 41a-f). The defects are a consequence of vacancies within the lattice (due to a loss of an atom or molecular unit), or the antisite defects (where an Mo atom substitutes a disulfur column and vice versa).[352] DFT calculations have shown detrimental effects to the electronic properties of the regions containing the point defect, and are thought to affect its 56

57 other intrinsic properties reducing the CVD-production route potential as to supply the material. The presence of both point defects and grain boundaries are not always detrimental to the properties of the TMDC nanosheets produced. The generation of such defects in a controlled manner could prove to be another avenue for tuning nanomaterials,[357, 358] as these features could potentially generate an increased density of active sites useful for electrical performance,[359] chemical sensing,[360] and catalysis.[361] Alternatively, the filling of the vacancies with foreign atoms may provide new opportunities to tune the properties of the MoS2 nanosheet and generate a new family for functionalised 2D semiconductors.[117, 352, 359, 362, 363] Nan and co-workers have already shown that defect engineering can result in a significant increase in the photoluminescence of MoS2 samples.[357] During our survey of the literature we have observed a number of nanosheets, mostly produced by CVD, which have dendritic adlayers on their surfaces.[306, 313, 323] The effects of these little-discussed by-products on MoS2 nanosheets have only recently been considered (Figure 41g-l).[364] It appears that the photoluminescent and conductive properties of the MoS2 are severely affected by the locations of the dendritic by-products and ultimately there is a negative effect on the properties of the whole nanosheet. The need for both the inhibition of adlayer formation and removal of the adlayers is made clear in this report, but no details have been presented how this problem may be tackled.

3.2.1.3 Electronic Properties of CVD/PVD-Grown MoS2 Nanosheets 3.2.1.3.1 MoS2 Nanosheet Islands The monolayer MoS2 nanosheet islands grown by CVD and PVD methods appear to be of a high crystallinity, exemplified by the high µe and ION/IOFF numbers observed, that are 57

58 comparable to the values obtained from the MoS2 FETs constructed with micromechanically exfoliated MoS2 nanosheets.[8, 305, 306, 308-310, 313, 324, 349, 352, 365-367] However, care needs to be taken with this conclusion, as some of the MoS2 FETs tested do not consider the true composition/structural integrity of the nanomaterial employed. Such findings of the non-uniformity of nanosheets (such as point defect, formation of layer boundaries, nonstoichiometric composition) have only been explored by few research groups. Najmaei et al.[368] have shown that the construction of a MoS2 FET which target the electrical properties observed in the presence of grain boundaries present can have significant effects on the overall performance (Figure 42a and b). The average two-point µe values of MoS2 nanosheet islands (grown by CVD) that measured through the grain boundaries of tilt-boundary and twin-boundary nanosheets were found to be 4.2, and 8.9 cm2 V-1 s-1 respectively, compared to 5.6 cm2 V-1 s-1 for a pristine nanosheet. A significant perturbation of µe within the tilt-boundary containing FET was observed when measured parallel with the boundary, but not with the twin-boundary crystal (averages of 2.0 and 5.5 cm2 V-1 s-1, respectively) suggesting that the tilt-grain boundaries contain intrinsic and extrinsic scattering centres that retard the potential µe values of the material. This is enforced by the effect of the photoluminescence at these types of grain boundaries discussed in Section 3.2.1.2.6. We have already discussed the effect of stoichiometry on the PL of nanosheets, but also note that both µe and ION/IOFF are also modulated by the composition of the nanosheet.[349] The nanosheets tested from stoichiometric MoS2 gave the poorest values (average of 0.5 cm2 V-1 s-1 and 104), while these values were significantly improved by lowering of the sulfur content. The µe number increased to 2.7 and 15.3 cm2 V-1 s-1 and ION/IOFF improved to ~10 5 and ~106 for sulfur deficient and MoOx-doped nanosheets, respectively. The measurements described for the partially oxidised sheets correspond to the 58

59 highest reported values of µe and ION/IOFF for as-synthesised unencapsulated nanosheets without an annealing phase. It is thought that both the structural defects, due to the presence of S vacancies, results in the nanosheets produced exhibiting a more metallic nature, with a net-positive contribution to the conductivity of the MoS2 nanosheets; in effect it is an example of n-doping. It is clear then, that tuning stoichiometry can be used to manipulate the electronic properties of nanosheets. Tongay et al. have also found that doping against the native character of a material can also force a switch between n- and p-type behaviour in MoS2.[369] Current LED, solar cell and photodiode technology relies on stable n-type and ptype conduction.[369] In order to fully utilise 2D-materials in these fields it is necessary to devise ways to achieve both n- and p-conduction. LP-CVD-grown MoS2 is a native n-type semiconductor due to sulfur vacancies.[81] Producing p-type behaviour in MoS2 or n-type behaviour in WS2 has traditionally been achieved through electronic engineering[142] surface functionalization[79] or charge transfer from physisorbed molecules,[370] rather than by imparting intrinsic properties to a material. Suh and colleagues were able to develop a route to p-type MoS2 via cation doping.[369] Using a LP-PVD technique similar to those described in above, multilayer Nb-doped MoS2 single crystals were obtained which showed p-type semiconductor characteristics. In a similar vein, chloride doping of nanosheets has been used to reduce resistance of MoS2 and WS2 nanosheets.[371] 3.2.1.3.2 MoS2 Thin Films It has been shown before that polycrystalline films of MoS2 nanosheets, grown on the centimetre scale, can contain a significant number of grain boundaries. To further understand the effects of grain boundaries on the films electronic properties, Najmaei et al.[368] created a series of FETs on an MoS2 film (with grains observed in the micron-scale) with an increasing distance between the two contacts MoS2, therefore an increase of 59

60 polycrystallinity, and in turn the number of grain boundaries the wider the separation is (Figure 42c-e). As a result, a systematic degradation in both µe and ION/IOFF were observed as the channel was widened from 4 to 76 µm. In an extension of the selective growth of MoS2 films with varying grain boundaries, Zhang et al.[325] explored the electrical properties of the numerous films made that contain grains varying from 20 to 600 nm (which in turn results in a decrease in the number of grain boundaries that accompany the semiconducting component; Figure 35). As observed before, the increase of the grain boundaries within the MoS2 film results in a degradation of its electrical efficacy (0.5 cm2 V-1 s-1 for 20 nm grain films to 7 cm2 V-1 s-1 for 600 nm grain films) Indicating that films with larger grains are a more desirable material to obtain. It has been noted that, after the production of a CVD- or PVD-grown polycrystalline TMDC film, an annealing step can greatly improve the electrical properties.[243] Heating of the film up to 1000 °C (in a Ar/S or Ar/H2S atmosphere to inhibit unwanted oxidation) has been found to greatly increase the crystallinity of the nanosheets, as the structural rearrangements at such a temperature reduces the number of defects and grain boundaries within the film. Studies of such have shown the µe values of poorly-formed MoS2 films (with initial µe values of 0.003 to 0.3 cm2 V-1 s-1)[325] can be improved to up to 17 cm2 V-1 s-1 by its annealing under inert conditions.[372]

3.2.2 Sulfurisation of Thin Metal-Containing Films 3.2.2.1 Preparation of Ultrathin TMDC Films by Metal or Metal Oxide Sulfurisation An alternative to the gas phase reaction and transport mechanism is the direct sulfurization of substrates coated with a thin layer of a precursor containing the targeted metal, negating the need for the transport of the TMDC within the reaction. Most commonly, 60

61 the metal oxides, such as MoO3 and WO3,[301, 302, 319, 372] are deposited onto substrates at varying thicknesses (from 0.3 to 10’s of nm) for further treatment, but other sources, such as metal films[373-376] and Mo/Au alloys[377] can also be used (Figure 43). These thin films are then directly sulfurized by the introduction of a sulfur vapour within the inert gas medium at elevated temperatures (either in the form of vaporised sulfur or H2S gas) to produce the thin desired nanosheet films. The mechanism for converting metal oxides to the TMDC discussed in Section 3.2.1.1 also applies in the sulfidation of such oxide films, but without the transport step. The mechanism for TMDC growth from metal films is by volatilisation of the partially sulfurised material on top of the surface of the metal, before complete sulfidation on the vapour phase above the substrate and redeposition.[375, 376] Sulfurization reactions of this type have been shown to produce thin films of MoS2[316, 372374, 377] and WS2[301, 302] on the cm2 scale, limited only by the size of the substrates and reaction furnaces used. Due to the non-transport mechanism, the morphology of the precursor tends to have a profound effect on that of the resulting TMDC film. It has also been shown that the use of a highly reactive sulfur source can allow the formation of MoS2 sheets at dramatically lower temperatures than equivalent reactions; H2S has been used in the synthesis of MoS2 from Mo metal films at temperatures as low as 250 °C.[376] The thickness of the precursor film directly determines both the average thickness and the relative orientation to the substrate of the MX2 nanosheets obtained.[42, 301, 378] This was neatly shown in a recently published syntheses of MoS2 and WS2 by the sulfurization of metal films of varying thicknesses on SiO2/Si.[328, 379, 380] In separate studies, the Cha and Choi groups provided evidence that discontinuous Mo (and W) films that were ≤ 0.5 nm thick produced horizontally-grown MoS2 (and WS2) nanosheets that were mono- to few-layers thick, whereas thicker, continuous seed layers (≥ 4 nm) promoted the growth vertically aligned stacks of nanosheets (Figure 44a-c). This horizontal-to-vertical 61

62 growth transition is considered to be a consequence of the competition of horizontal growth to expose low-surface energy basal planes versus the steric restriction of the crystal expansion with heavily-coated substrates. Accurate control of the thicknesses of nanosheet films produced has been found to be difficult, but a level of control can be exercised. Examples show that the controlled deposition of 1 nm WO3 onto SiO2/Si wafers (either by ALD[381] or sputtering[301]), followed by sulfurisation produces wafer-scale monolayer WS2 films. The treatment of thicker WO3 films with sulfur vapours (ranging from 1.6 to 3 nm[301, 381]) can give sufficient control to produce 2 to 4 layer thick WS2 films (Figure 44d-f).[301, 381] As these reactions employ no directing gas flow, the deposition of monomer units tends to be random, difficulties therefore arise in the promotion of homogeneity across the whole film, in terms of film thickness and crystallite size. Upon closer study of the MoS2 and WS2 films produced by these methods, it is clear to see that the many of the films produced consists of very small (10-50 nm) grains, resulting in nanosheet films with a very high density of grain boundaries and defects (Figure 44g). As a consequence, the electronic properties of these sheets are found to be particularly poor, exhibiting significantly reduced µe to those observed for CVD-grown and mechanically-exfoliated nanosheet counterparts.[373] The growth temperature selected in the sulfidation of Mo films is found to critically impact on both the quality and the dimensions of the film grown at lower pressures. Work by Rajan et al.[56] has shown that increasing the reaction temperature from 500 to 1100 °C results in an evolution of the surface morphology of the formed MoS2 nanosheets (Figure 45a-d). At lower temperatures, the monolayer nanosheet films produced highly polycrystalline with a significant number of grain boundaries. At the optimal temperature of

62

63 900 °C, a high-crystallinity, thin monolayer film is produced, while higher temperatures create micron-wide MoS2 flakes that are 10’s nm thick. A small effect in on nanosheet grain density was noted by changing the Mo precursor to be sulfurised. It has been noted that the lateral sizes of monolayer nanosheets grown from Mo films are significantly lower than if MoO3 is used (Figure 45e-f).[378] In addition, the grain density of nanosheets grown from MoO3 films can be further improved by increasing the film thickness of the precursor, but at the expense of an increase in thickness of the resultant film.

3.2.2.2 Electronic Properties Electrical measurements carried out on the as-obtained nanosheets produced by the sulfurisation of Mo/MoO3 films showed µe values of between 0.004 and 0.04 cm2 V-1 s1.[312, 367, 372, 373, 375, 382] In methods similar to those described in section 3.2.1.3, the measured µe of the TMDC films can be improved to as high as 6.5 cm2 V-1 s-1 by the introduction of an annealing step to ~1000 °C either by experiment procedural design[375] or by post-reaction treatment.[243, 373, 383]

3.2.2.3 Developing Methodologies for Film Sulfurisation Reactions The preparation of a uniform, atomic-thick layer of the precursor film (e.g. molybdenum metal) over a large surface area is currently one of the most difficult hurdles to overcome for large-scale production.[377] The process of Choi et al. allows the formation of a thin molybdenum layer by the low temperature reaction of Mo(CO)6 and a pre-deposited Au film to form a Mo-Au surface alloy with a controllable Mo-content. The Mo/Au alloy can then be sulfurised to produce MoS2 nanosheets with controllable thicknesses, before isolation of the nanosheet by Au etching (Figure 46a).[377] As the shape of the MoS2 film is determined by that of the Au substrate, patterning of the MoS2 film is possible by the 63

64 sulfidation a pre-lithographed Mo/Au pattern(Figure 46b). The method remains imperfect, with a significant number of MoS2 clusters disrupting the surface of the film(Figure 46c and d). The literature indicates that these clusters are likely formed during the sulfidation process and grow through the MoS2 film as it forms,[384] reducing the film quality. The poor structural uniformity of the films produced is reflected in the extremely poor carrier mobility µe (0.004 cm2 V-1 s-1). In the majority of cases described in Section 3.2.2.1, the metal precursor is prepared by directly sputtering a metallic species onto the desired substrate; the metallic layer is treated with vaporised chalcogen species in a separate reaction. A recently reported alternative is to sputter metal atoms through a reactive sulfur atmosphere, therefore achieving the formation of TMDC films in a one-step process. This approach has been successfully utilised in the production of large-area, near-pristine MoS2 monolayers.[303] The obtained films display high µe (12.2 cm2 V-1 s-1), but a significantly reduced ION/IOFF of 103. As the reaction takes place at 700 °C the films are effectively annealed during the procedure. The films also show intrinsic p-type character as a result of slightly increased sulfur content. A recent development in the gas phase synthesis of MoS2 nanosheets is the use of atomic layer deposition (ALD) as reported by Loh and colleagues.[383] Using a pulsed growth cycle on a sapphire substrate exposed to MoCl5 vapours, before purging the system with N2 and then introducing H2S gas at 300 °C (Figure 47a). This method allows precise control over the precursor addition into the system, as 5 ALD cycles can generate 1 micron wide triangular MoS2 monolayers, whereas 10 and 20 cycles produce films with consistent thicknesses (1.7 and 3.2 nm, corresponding to 2 and 4 layers, respectively; Figure 47b-d) and a larger number of ALD cycles can reliably produce nanosheet films with controlled thicknesses of up to 10 layers. However, the nanosheets produced from the initial 300 °C reaction are of poor crystallinity, as observed by the poor uniformity of the sample by AFM 64

65 and the loss of the characteristic photoluminescence signals of single-layer MoS2. These poor quality crystals can undergo reconstruction by high temperature annealing at 800 °C to improve the crystallinity sufficiently to restore the expected MoS2 monolayer properties. Currently, there is no information on the electronic properties of the nanosheet islands or films produced by this method. Instead of using amorphous metal/metal oxide films, Jiao and co-workers have promoted the synthesis of MoS2 nanosheets by the sulfurization of MoO2 microcrystal templates.[316] These rhomboidal microplates of MoO2 (produced by the LP-CVD-based reduction of MoO3 and deposition on SiO2/Si substrate) were then annealed in a sulfur vapour at 850-950 °C to reliably produce mono- to few-layer MoS2 on the surface of the microplate with lateral dimensions >10 µm, which can then be transferred onto a clean substrate (Figure 48a). The nanosheets produced have a significantly higher crystallinity, SAED and HR-TEM show that the high-crystallinity of the template is retained during sulfurization. The films exhibit ION/IOFF and µe values approximately one order of magnitude higher (~10 6 and 0.1-0.7 cm2 V-1 s-1, respectively) than those observed from the sulfurization of MoO3 and Mo metallic films (before annealing). The mechanism for the sulfurization of MoO2 microplates is thought to follow that observed in the sulfurization of MoO2 fullerenes and nanotubes;[385, 386] during sulfurization, the first layer of the oxide surface is sulfurised quickly, followed by a slower diffusion-controlled sulfurization of the lower layers. As this process is limited by the rate of diffusion of the sulfur vapours through the already existing MoS2 outer walls the process becomes slower with increasing film thickness. It is this diffusion control that allows for MoS2 with single- or few-layer nanosheets to be reliably obtained by control of the annealing time (from 0.5 h to 6h; Figure 48b and c).

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66

3.3 Nanosheet Growth in the Liquid- and Solid-Phase The liquid-phase syntheses of TMDC nanosheets involve the decomposition of a precursor mixture by either thermal decomposition within a solution or injection into a hot liquid (hot injection synthesis). The TMDC nanosheets produced in this fashion can then be isolated as colloidal dispersions in common, low boiling solvents. Solid-phase reactions (based on the 'melt' reaction) generally involve the deposition of a single-source precursor onto the desired substrate, which is then heated slowly to melt the precursors, followed by subsequent formation of the TMDCs within the molten mixture. Both liquid- and solid-phase (or solventless) reactions have received considerably less attention than gas-phase syntheses; however these procedures constitute an emerging field in TMDC nanosheet production.

3.3.1 Liquid-Phase Reactions Although solution-phase synthetic methods are considered highly promising in the production of many nanomaterials as they allow for the use of simple synthetic procedures that can assert control on specific growth parameters,[19] their synthetic implementation for single-layer 2D TMCD production has been challenging. Numerous reports have observed the production of TMDC nanosheets in a liquid medium, however such reactions often result in the formation of non-monolayer products and within array of hierarchical structures,[208] especially in the absence of directing/templating agents.[387, 388] Some reports show the possibility of obtaining freestanding TMDC 2D-nanomaterials,[389] however, these methods were also limited in the fact that only multi-layer (10’s nm thick)[390] or monolayer nanosheets with lateral dimensions in the 10’s nm[391] were produced. The principles controlling TMDC nanosheet growth in the liquid phase are similar to those at work in gas-phase growth. Random nucleation forms seed nuclei which promote the growth of larger nanosheets and conditions must be chosen in order to ensure that growth is

66

67 favourable in the xy-planes and unfavourable in the z-plane. In liquid-phase reactions this is achieved through the selection of coordinating ligands[387, 392, 393] performing a similar role to templating substrates in gas-phase syntheses) and through judicious control over reaction kinetics. Whereas the role of templating substrates in gas-phase growth is partly to reduce the energetic penalty of parallel growth,[320] the role of capping agents/coordinating ligands in the liquid phase is to bind preferentially to the c-axis and thus raise the energetic penalty of growth in the z-plane. As the growth of nanosheets proceeds from the edge of seed nuclei, it is imperative that the seed crystal itself being of a monolayer in order to obtain a monolayer product. Heavy solvents such as oleylamine are regularly used both to reach the temperatures required for reaction and as efficient coordinating ligands to enforce nanosheet growth.[391] In particular, oleylamine has been shown to be essential for the formation of nanosheets over other nanostructured morphologies.[394] However the solvent becomes chemically bound to the nanosheet surface, forming ligand-capped nanosheets.[391] Removal of these capping agents from the nanosheet is problematic, and often results in the restacking of the freshly produced nanosheets.[392] Nonetheless liquid-phase growth of TMDC nanosheets is an intriguing prospect in future works. Two dual-source methods have been developed for the synthesis of group IV and V sulfides, most recently highlighted by the work by Cheon et al..[393] Both of these processes are based on the sulfidation of the corresponding metal chloride in oleylamine, but are performed using contrasting methods. The contrasting results provide some insight into the reaction kinetics of growth and the morphologies which arise.

The first method produces MX2 nanosheets by the hot-injection of CS2 into an oleylamine solution containing MClx at elevated temperatures (up to 300 °C; Figure 49).[395] 67

68 The second method involves the heating of an oleylamine solution containing the MClx and elemental sulfur precursors (from room temperature to up to 250 °C) for prolonged reaction times.[396] Both robust procedures have been used to produce nanosheets of 1T-TiS2, 1TZrS2, 1T-HfS2, 1T-VS2, 2H-NbS2 and 2H-TaS2; the TMDC nanosheets are produced in their most stable crystal structures.[393-396] The TMDC nanosheets isolated were found to be small (< 100 nm diameter), however, the report indicates that precise size control (from 20 to 100 nm) can be garnered by modulating the precursor concentrations. Although the products from the two processes described above show a number of similarities, there is one key difference: The well-shaped nanosheets produced by the CS2 hot injection method were found to be non-monolayer with typical thicknesses in the few-to-10’s of nm range (Figure 50), whereas the thermal reaction with elemental sulfur has been found to produce nanosheets of poor quality, in terms of shape, size and crystallinity, but produced a monolayer product. Investigations found the reasoning behind the differences in the products isolated from these two methods:[393] In both processes, the chalcogen precursors employed generate the highly reactive H2S species, which in turn reacts with the metal chlorides at the elevated temperatures imposed. The hot injection of the highly reactive CS2 precursor generates a burst influx of H2S in the liquid medium that promotes a burst nucleation step and produces small molecular clusters that are multilayer, from which the larger nanosheets grow.[393] On the other hand, the heating of the sulfur solution produces the H2S intermediate[394] in a slower, gentler manner resulting in the bulk nucleation of monolayer seed crystals. However, the generation of sulfur radicals promotes degradation of the structural integrity of 2D layered nanocrystals via the thermal decomposition route. In contrast with problems associated with the use of elemental sulfur in the synthesis of metal sulfides, the thermal reaction of MClx with elemental selenium can prepare higher quality 1T-TiSe2, 1T-VSe2, 2HNbSe2 and 2H-TaSe2 (as well as ZrSe3 and HfSe3) nanosheets that are 3-20 layers thick. The 68

69 reason for this resides in the fact that, under the conditions imposed, elemental selenium is more reactive, and generates relatively milder radical species compared to its sulfur counterpart. In both cases, the fact that the nucleation process occurs relatively rapidly means that a large number of seed crystals are produced at the start of the reaction. Rapid nucleation results in a relative shortening of the growth stage as the as precursors stock is rapidly consumed. Nanosheet sizes are generally limited to <100 nm before exhaustion of the precursors. In the view of this, Cheon et al. developed a methodology that promotes the growth of Group IV MS2 (M = Ti, Zr, Hf) nanosheets in the liquid phase that were isolated as exclusively monolayer, TMDC single crystals with lateral dimensions in the 100’s of nanometres.[327, 397] This was achieved by the restriction of the chalcogen source during the nucleation and growth processes. The use of 1-dodecanethiol (1-DDT) in place of CS2 in the reaction with the metal chloride is essential to control both size and thickness in nanosheet growth (Figure 51a). 1-DDT has been found to generate H2S 86 times slower than from CS2 under the conditions imposed. The slower reactivity of 1-DDT garners a level of control in the nucleation step, ensuring the production of a significantly smaller number of monolayer seed crystals, allowing for the extension of the growth stage and thus generating significantly larger nanosheets of 1T-TiS2, 1T-ZrS2 and 1T-HfS2 (Figure 51b-e). Similar controls in the nucleation and growth process have been shown by the direct treatment of TiCl4/oleylamine solutions with H2S gas; a slow, controlled flow of H2S gas into the solution at 230 °C garners single layer nanosheets with dimensions between 300 nm and 1 µm, whereas oversaturation of the gas promotes burst nucleation, and subsequent formation of 3layer nanodisks that are smaller than 50 nm in diameter (Figure 51h-j). In contrast to the examples earlier, it is clear that there is currently a lack of control in the formation of Group VI sulfides by liquid growth, as all examples known on the growth 69

70 of MoS2 and WS2 produce exceptionally small nanosheets. Separation of nucleation and growth stages of group IV TMDCs clearly presents some challenges not present in other systems. To date no group has reported the lateral dimensions of MoS2/WS2 nanosheets in excess of 100 nm. The first working synthesis of this type was the formation of 1H-MS2@oleylamine (M = Mo, W) reported by Altavilla et al..[391] Using the thermal decomposition process with a single-source precursor ([NH4]2[MS4]), single-layer MX2 nanosheets were formed in oleylamine (Figure 52a), however, these sheets were invariably small, (< 15 nm; Figure 52c). Similar results with the hot injection process have been described by the Liu/Li et al..[392] 1H-WS2 can be made by the injection of a sulfur/oleylamine solution into a 300 °C solution of WCl6 in oleylamine/octadecene. The nanosheets produced by the hot injection processes were found to be exclusively monolayer, and were extremely small (< 25 nm). In this synthesis, the nanosheets isolated are colloidally unstable; the nanomaterials, capped with oleylamine, are found in random arrangements within 100’s nm to micron-scale aggregates (Figure 52b). It is clear from these reactions that, although the products obtained appear to be exclusively monolayer, the nucleation step in the reaction is dominant in the process, inhibiting continued nanosheet growth due to precursor starvation. Any future endeavours in this area will require sufficient control in the nucleation step, to propagate the larger nanosheet growth desired in many fields of interest. Until recently, all such methods of TMDC production relied on the use of oleylamine as the liquid medium, as the high-boiling solvent promoted the production of TMDC in nanosheet-shapes,[394]. The impact of coordinating ligands similar to those used to control morphologies in other 2D-material growth,[398] as well as quantum-dot growth synthesis[399] was relatively understudied. Ozin et al. challenged this by introducing strongly coordinating ligands to the precursor mixtures in an effort to produce WS2 70

71 nanosheets.[400] They observed that the thermal breakdown of WCl6 and CS2 in oleylamine produced a mixture of mono- and few-layer 2H-WS2 nanosheets that are colloidally-unstable, similar to the results obtained by Liu and Li (Figure 52d, f).[392] However, the introduction of oleic acid or 1-DDT to the reaction resulted in the production of (mean diameter 100 nm) exclusively monolayer 1T-WS2 nanosheets that are colloidally stable (Figure 52e). The 1T/2H-polymorphism observed in the production of WS2 is driven by the reactivity of the tungsten precursor; a highly reactive tungsten intermediate yields the 2H-structure, whereas introducing a strongly coordinating ligand reduces its reactivity sufficiently to produce the 1T-structure. Supressing the effect of the coordinating ligand was successfully achieved by introducing hexamethyldisilazane (HMDS) to the oleic acid-promoted growth, as the two species react, stripping the coordinating ligand from the growing crystal, resulting in the production of 2H-WS2. The lack of reactivity of HDMS with 1-DDT is replicated in the addition of the silazane to the 1-DDT-promoted reaction, as only the 1T-WS2 is produced. They also reasoned with the colloidal stability of 1T-WS2 and instability of 2H-WS2; the distorted 1T-WS2 is stabilized by the electrostatic interactions formed between a stable [WS2]- species and the ligated oleylammonium cation on both surfaces of the nanosheet. This charged structure at the same time prevents the aggregation and thickness growth of the nanosheets. On the other hand, the 2H-WS2 nanosheets are considered to be neutral and coordinated weakly with the oleylamine ligands, therefore nucleation and growth processes can occur, resulting in the morphologies observed.

3.3.2 Solventless Reactions As with the liquid phase of nanoparticle growth, the ability to control the growth of TMDCs to produce freestanding nanosheets within solid phase reactions has been challenging. Many examples are known for the growth of MoS2 nanomaterials in the solid state that exhibit a large array of micro- and macrostructures.[401] To date, there are only a 71

72 few available procedures that have been able to produce TMDC nanosheets (namely MoS2) with the desired two-dimensionality. As discussed above, there is little scope for the inclusion of capping agents into the mixture and the near-instantaneous reaction results in a high nucleation-density. Generally this results in the production of small, non-monolayer TMDC sheets.[402] To date, only MoS2 films have been synthesised, using the solventless thermolysis of substrates coated with [NH4]2[MoS4] (which was prepared by dip-coating[402] or spincoating[403] of aqueous solutions containing the precursor) in a H2S-rich environment (Figure 53a and b). Both the dip-coating and spin coating processes have been found to produce large-area (> 1 cm2) films, consisting of crystallites with a varied thickness (1 to 3 layers have been reported; Figure 53c and d). However, the nanosheet films are constructed from a large number of small (< 50 nm) MoS2 grains within the films, and subsequently a large number of grain boundaries, imaged by high-resolution electron microscopy. There are no electrical measurements for the films prior the annealing steps, but after annealing in a sulfourous environment at 1000 °C (as previously described in Section 3.2.1.3.2) the MoS2 films showed greater crystallinity and higher µe (0.1-4.7 cm2 V-1 s-1) and ION/IOFF (ranging from 103 to 105), comparable to that produced by the micromechanically exfoliated MoS2based FETs.

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4 Advances in TMDC Nanosheet Engineering The work we have covered so far has focussed on efforts to develop routes to highly pure 2D nanosheets. The long term goal is not necessarily the use of these pristine materials in device manufacture, rather the aim is to develop a toolkit of routes to reliably produce nanosheets which can be combined and patterned to create truly novel 2D- or 3Dnanomaterials with tunable properties in complex and pre-defined architectures.[404] The formation of these next-generation materials has been made possible by understanding the processes required to form binary materials and applying these lessons to allow the formation of alloyed nanosheets, construction of heterojunctions between nanosheets, as well as postproduction etching of nanosheets to produce complex patterned structures.

4.1 Development of Ternary TMDCs for Developing Tuneable Materials. In order to maximise the potential of TMDCs (as well as other 2D-materials) in many areas, precise control of a number of properties of the nanosheets is of crucial importance. Specifically, the ability to manipulate the band gap of each material to match that required for specific applications allows greater versatility in the creation of electronic and photonic devices. One of the current exploration routes of TMDC nanosheets is based on the tuning of bandgaps to achieve selected spectral responses in optoelectronic devices.[405] A number of routes to enable the tuning of properties (including of the band-gap) have been explored, including chemical treatment and doping,[47, 363, 406-412] dimension tuning,[413-415] strain- and pressure-mediated,[416-419] defect engineering[359, 420] and dielectric screening,[421] among other more complex proposals.[422] All of these methods have been successfully shown allow the adjustment of the band gap of certain materials, however the extent to which the gap can be changed is severely limited. Alloying is the exception to this rule as it allows precise band gap tuning over a wider spectral range, limited

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74 to the band gaps of the constituent materials.[423, 424] As the electronic properties of the members of the TMDC family spans the full range of behaviour from insulating to metallic,[23] precise band-gap engineering is likely to be possible via alloying TMDC materials. Moreover, most single-layer TMDCs generally adopt the same (trigonal-prismatic or hexagonal) structures and have similar lattice dimensions, allowing relatively strain-free inclusion of heteroatoms within a framework.[406] Therefore the inclusion of heteroatoms (either metals or chalcogens) leads to the formation of stable, ternary materials with composition-dependant band gaps.[423] Further to this, the doping and alloying of wellunderstood materials may yet yield an improved performance of the properties regularly investigated in binary TMDC nanosheets. In practice, band gap tuning of TMDCs via alloying was initially demonstrated in bulk TMDC crystals[425-427] and has more recently been shown to be applicable to ultrathin materials.[428] It is critically important to develop a strategy to reliably produce monolayer TMDC alloys to overcome the limitations experienced in other bandgap modulating processes.[429-433] It is likely that alloyed nanosheets will obey the same general rules as those made from binary compounds, therefore the quality of individual sheets and the presence of defect structures will have a significant impact on the observed properties of any prepared material. Therefore as well as designing syntheses for ultrathin alloys, it is essential that reaction conditions allow the number of defects to be minimised in order to provide reliable and useful materials. Similarly the photoluminescence, Raman scattering and electronic properties tend to lie between those of the constituent materials, however to our knowledge whether such characteristics are linearly dependent upon composition is yet to be systematically studied.

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4.1.1 From Top-Down Methods In one of the first groups to examine the properties of ultrathin ternary nanosheets, Huang and Xie[342, 428] have taken advantage of the previous work done in the growth of a series of (Mo xW1-x)S2 and (Mo xW1-x)Se2 single crystals (where x ranges from 0 to 1) by CVD.[426] By applying the Scotch Tape method of micromechanical exfoliation to the bulk alloys, (MoxW1-x)S2[428] and (MoxW1-x)Se2[434] monolayer nanosheets were produced (Figure 54a and b). The compositions of the nanosheets match those of its parent crystals, identified by HAADF STEM imaging (atomic-resolution mapping of the heterometals; Figure 54c),[428] Raman modes (monitor the A1g and E2g bands of the MoX2 and WX2 units within the ternary nanosheets; Figure 54d)[434] and the photoluminescence signatures (show a blue-shift in the peak of the material as an indirect consequence of the bandgap modification by increasing the tungsten content; Figure 54e).[428, 434] In principle such methods may be applied to any bulk, lamellar alloy, however the composition of each nanosheet will be determined solely by the composition of the corresponding bulk powder. It is therefore vital that there exists no variation in composition within a sample or single crystal in order to obtain a homogenous mixture of exfoliated sheets, as little control may be exercised during the thinning stage.[434] The method also suffers from all of the advantages and issues highlighted in Section 2.2. Using similar methods to the Coleman process,[176] Mo(SxSe1-x)2 nanosheet dispersions (where x = 0, 0.5, 0.6, 0.75, 1) have been produced by liquid exfoliation of preprepared alloy crystals.[435] The few-layer nanosheets from each alloy (with lateral dimensions in the 100’s nm), were found to retain the homogenous compositions seen within their parent crystals (Figure 55a-f). Studying the catalytic activity of the alloy has revealed a

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76 systematic structure-activity relationship. It has been found that Mo(S0.50Se0.50)2 nanosheets exhibit a synergetic increase in HER activity over the binary MoS2 and MoSe2 catalyst films. Tremel et al.[436] described the lithium intercalation-exfoliation synthesis of (NbxW1-x)S2 nanosheets from (Nb xW1-x)S2 coin-roll nanowires (Figure 55d). The coin-roll nanowires were pre-fabricated by the sulfurization of WOx@Nb 2O5 nanorods with H2S gas at 850 °C (in the high-energy, metastable 1T-lattices; Figure 55e). Characterisation by STEM and EDX found the nanosheets were small (mean size was 300 nm), adopted the poorlystable 1T-structure and were relatively impure, with the composition of individual nanosheets ranging from WS2 to (Nb0.5W0.5)S2 (Figure 55f). Such non-homogenous mixtures are likely to be difficult to incorporate into efficient devices due to band gap variations throughout the sample.

4.1.2 From Bottom-Up Methods The most promising routes to ternary nanosheets appear to be based on the gasphase bottom-up syntheses, as these methods are thought to provide a tuning procedure to produce nanosheets tailored to their needs. The Zhang and Xie research groups[345] employed the LP-PVD growth method of TMDCs within an argon/hydrogen carrier gas mixture. The co-deposition of MoS2 and MoSe2 vapours onto a SiO2/Si substrate was achieved by independently heating MoS2 and MoSe2 sources (Figure 56a and b). Independent volatilisation of each component allows partial control over the extent of volatilisation of each precursor and therefore the composition of Mo(SxSe1-x)2 films grown. Both cm-scale single-layered films and triangular nanosheets with average dimensions of 2.5 µm have been produced (dictated by the reaction temperature and flow rate of H2; (Figure 56c and d), with precise control of the rate of

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77 alloying giving rise to uniformly grown nanosheets with the composition ranging from MoS2 to Mo(S0.60Se0.40)2 (Figure 56e and f), and giving specific and reproducible band gaps. Liu and co-workers had followed the sulfidation of a thin metal film route, using a co-sputtered Mo/W film on sapphire.[378] The film precursors used had elemental W/(Mo + W) ratios of 0, 0.19, 0.55, 0.85 and 1, and were sulfidised by sulfur vapours within a CVDtype setup. Sulfidation at increasing temperatures up to 950 °C produced a thin film of (Mo xW1-x)S2, with the elemental composition closely matching that of the initial co-sputtered Mo/W films used. A survey of the triangular nanosheets within the films produced showed that individual sheets have an average thickness of two-to-four layer nanosheets, but with an average grain size of only 50 nm. Two research groups have targeted the synthesis of Mo(SxSe1-x)2 by the LP-CVD reaction of MoO3 with a mixture of sulfur and selenium vapours onto SiO2/Si substrates. Duan and Pan[437] had achieved the synthesis by a slight modification of the CVD synthesis of MoS2 by described in Section 3.2.1.1, by introducing both elemental selenium and sulfur vapours within the argon carrier gas. By altering both the furnace temperature and the proportions of sulfur and selenium powders introduced, the researchers were able to obtain bi-to-trilayer nanosheets of Mo(SxSe1-x)2 (where x was between 0.1 and 0.94) with lateral dimensions of 30-80 µm. Each individual nanosheet produced was deemed pure; there was no observable compositional difference within nanosheet islands, but the composition of the film did vary in relation to the position of the substrates to the centre of the reaction zone of, as the substrates nearest to the centre tended to be almost exclusively MoS2, with an increasing selenium-content the further away from the hot zone. Bartels and co-workers[346] had looked to introduce the sulfur and selenium sources in a different manner. They discovered that the injection of a liquid chalcogen source 77

78 (thiophenol dissolved in THF as the sulfur source, with diphenyl diselenide added as the selenation agent) via a long syringe needle near the region of the tube furnace containing the MoO3 powder and a SiO2/Si substrate at 650-700 °C in a Ar/H2 environment (Figure 57a). The process is known to produce both thin films and islands of the resulting single layer Mo(SxSe1-x)2 alloys. As with Li et al,[437] both temperature and liquid S/Se ratios allows precise alloying control, as shown by the range of compositions which were isolated and tested (Figure 57b). Both the thin films (imaged at an approx. 70 by 70 µm area), and 10 µm islands produced show exceptional homogeneity, with only a very slight variation (1 %) in the S/Se ratio detected within samples (Figure 57c-e). Another avenue for the generation of band-gap tuneable materials is by post growth manipulation of the nanosheets. Li et al. have explored this with the chalcogen exchange reactions of MS2 and MSe2 monolayers (M = Mo, W), grown by standard LP-CVD techniques, by their exposure to selenium and sulfur vapours, respectively, at elevated temperatures.[347, 438] This is achieved by the transport of the vaporised selenium/sulfur (in a mixture of argon and hydrogen carrier gases) towards the pre-grown MS2/MSe2 nanosheets on sapphire within a furnace at temperatures of 600-900 °C (Figure 58a-f). The temperature to which the TMDC nanosheets are heated is critical for controlling the rate of chalcogen exchange, as the thermal energy required to overcome the reaction barrier enables the breakage of M-X bonds, subsequently promoting the chalcogen exchange-process. With the selection of the reaction temperature, the nanosheets can have the tailored composition anywhere between MS2 and MSe2 nanosheet alloys, although the composition are often not strictly uniform, as shown by HRTEM[438] (for MoSe2 sulfurisation; Figure 58g-k) and photoluminescence[347] measurements (in MoS2 selenation; Figure 58n); The former showed that the sulfurisation of MoSe2 occurs through discreet channels on the nanosheets, this in turn giving block MoS2/MoSe2 alloyed cross-sections. The latter describes the 78

79 emission wavelengths observed at the corners and edges of the resulting selenium-treated MoS2 nanosheets, which are longer than those obtained at the centre-site (a difference of 10 nm is observed from the alloyed nanosheet treated at 900 °C), that indicates that selenisation is preferable an edge sites and defects, with the central sites remaining comparatively inert to the selenisation process. This results in a band gap difference of 17 meV at these two sites. Finally, a two step procedure has also been developed that yields a tunable Mo(SxSe1-x)2 alloy films from CVD-grown MoS2.[439] Exposure of the MoS2 film to a sputtered Ar+ beam (mild enough to not cause damage down into the Mo core layer) gently removes the top layer of sulfur atoms to yield S-vacancies within the film. This is immediately followed by the treatment with diphenyl diselenide vapours and subsequent annealing. Repeated sputtering/selenation cycles of the targetted film have been coupled with photoluminescence spectroscopy to determine the ratio of selenation after each sequence. After the 6th sequence, the photoluminesnce signal has redshifted by 27 nm (or 0.07 eV), corresponding to a the formation of a Mo(S0.79Se0.21)2 alloy film, a composition confirmed by further XPS analysis. The photoluminescence intensity after treatment shows little deterioration , indicating that the process does not generate defects within the selenated films, but we believe the observed broadening of the photoluminescence spectrum of the Mo(S0.79Se0.21)2 alloy indicates an inhomoenous selenation; the surface exposed to the mild sputtering processes is susceptible to selenation, whereas the surface bound to the SiO2/Si substrate will be expected to be selenium-free.

4.2 Construction of Heterojunction contacts Between TMDC Nanosheets Following the advancement on the applications of semimetallic TMDC nanosheets, a considerable effort has been made to understand the van der Waal interactions of the nanosheets with both encapsulating agents and other TMDC nanosheets as potential electrical

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80 juntions within devices.[440-452] However the heterojunctions formed commonly suffer from high and poorly controlled resistances due to the contact arising from these relatively inert van der Waal interactions. These effects are commonly and significantly detrimental to the device characteristics. In view of this, there has been a surge of articles which attempt to improve the connectivity by promoting the formation of heterojunctions between TMDC nanosheets in both the lateral and vertical dimentsions, as theoretical studies suggest the formation of such van der Waals heterostructures will significantly impact the value of µe and other electronic and optical properties.[453, 454] The first series of materials of this type were formed via the stacking of monolayer TMDCs to form bilayer or multilayer vertical heterojunctions. A number of reports study the heterojunctions of artificially-stacked MoS2 bilayers,[417, 446] but a growing number of examples bearing MX2/M’X’2 heterojunctions (including MoS2/WS2,[443] MoS2/WSe2[455, 456] (Figure 60a) and MoSe2/WSe2[447, 457]) have been produced. The stacking of nanosheets is commonly carried out through dry-transfer of mechanically exfoliated[455, 457] or gas-phase grown sheets.[417, 446, 447] Composites and heterostructures can also be formed from liquid-exfoliated sheets either by solvent removal or deliberate flocculation of colloidal sheets.[450] Formation of heterostructures by these methods leads to arbitrary alignment of sheets, producing an interlayer twist (Figure 60b). It has been shown that, the optical and electronic properties in artificially-stacked TMDC bilayers are extremely sensitive to the imposed interlayer twist, garnering enhanced (or detrimental) properties that could be utilised in the future depending on orientation.[417] Although this imposted twist can encourage an element of tuning of certain properties, that include second harmonic generation[446] (Figure 60c-l) and bandgap-tuning (Figure 60m),[417, 458] the significant issue of the non-reliable production and twist-control of such artificially-stacked bilayer

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81 nanosheets means that such materials are only ever likely to be of use to the research community. In contrast to the restacking method, the growth of bilayer (and multilayer) TMDCs, grown by many gas- and liquid- phase reactions, always exhibit a far greater regimented structure, as the growth of the second layer is stringently templated by the base monolayer.[332] These multilayer crystals are often found with a tilt angle of 60°, akin to the 2H-orientation (a tilt angle of 0 degrees (3R-) is ocassionally observed). Although the ability to select a tilt is lost, this method can reliably produce TMDC bilayer nanosheets given the right conditions are met (Section 3.2.1.2.3). Until late 2014, this process was limited to the two layers being of the same composition. Zhou and Ajayan[459] were the first to devise a single-step CVD method for the production of a MoS2-WS2 vertically heterostructured bilayer nanosheet. The CVD reaction from MoO3, and a W+Te mixture (Te to aid volatility of W; Figure 60a) exposed to sulfur vapours at 850 °C reliably grew WS2 nanosheets grown on top of monolayer MoS (Figure 60b and c). Importantly, within the stacked structure the 2H-arrangement was preseved (shown by the atomic-resolution TEM images; Figure 60f and g), indicating that the growth of WS2 on top of MoS2 is via a templated, non-random, epitaxial mechanism and therefore more reproducible than the mechanical transfer of layers described prevoisly. There is evidence that the stacked junctions exhibit a stronger interlayer coupling than from those created from artificially-restacked materials, exhibited by a second low-energy photoluminescence peak (in addition to the photoluminecscence signature of its parent nanosheets). Such an emission had not been previously reported and is thought to be direct evidence of a theoretically predicted[454] low-energy, direct band-gap. In addition to the vertical heterosturctures, Zhou/Ajayan et al. showed that lowering the reaction temperature to 650 °C produced a nanosheet with a significantly different

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82 morphology;[459] the production of a MoS2-WS2 heterostructure junction within the plane of the monolayer lattice (Figure 60d and e). The principle of gas-phase growth from the Zhou/Ajayan-CVD method bear some relations to the Duan(CVD and PVD)[460] and Wu/Sanchez-(PVD)[461] growth systems: the inclusion of two differing metal chalcogenide precursors placed within the same reaction, with the sequential volitalisation and deposition of each metal chalcogenide onto the same substrate; the successive growth of the second material at the reactive edge of the existing domain on the first-grown nanosheet. The process has to occur in one step, as the reactive edge-sites are easily passivated upon termination of the initial growth cycle and exposure to the environment, leaving them unlikley to act as nucleation sites for the second growth stage. The very small lattice mismatch of the family of TMDCs (for example, MoS2 and MoSe2, or WS2 and WSe2 mismatches of ~4 %) means that the lateral heterostructures produced are close to pristine, with few defects observed around the heterojunction. Although the general process is concurrent, the methods on which the differing processes produces the metal chalcogenides vapours in order. Zhou/Ajayan relies on the relative volatilasion of the MoO3 and W + Te mixtures upon sulfurisation and deposition, finding that the generation of higher local vapour pressure of MoS2 than than of WS2 on the substrates surface – only the MoS2 growth will be promoted until theprecursor feed has starved and is replaced by WS2 vapour. The method developed by Wu/Sanchez relied on the relative volatilities of MoSe2 and WSe2 within its mixture during a LP-PVD-type experiment. The rapid evaporation of MoSe2 at 950 °C stifles the volitalisation of WSe2 until the MoSe2 precursor becomes selenide-deficient and loses volatility, allowing for a sequential TMDC vapour production. Duan et al. uses both a LP-CVD and LP-PVD system to generate MoS2MoSe2 (via MoO3 with S and Se powders) and WS2-WSe2 (from bulk WS2 and WSe2), respectively. The Duan systems relies on the insitu mechanical switching of each of the 82

83 different precursor feedstocks into the correct zone of the furnace to promote the generation of TMDC gases and subsequent deposition. All the examples described in this section produce highly uniform, monolayer, triangular domains consisting of a TMDC monolayer surrounded by a different TMDC nanosheet. The nanosheets producedare typically 10-20 µm in diameter, and remain monolayer in thickness. The atomic resolution images of the MoS2-WS2 lateral heterostructures, described by Gong et al. shows that the interface of the heterojuction consists of a series of sharp zigzag or armchair steps, with a seamless connection at the interface. The fact that the interface exhibits a crystal lattice analogous to a typical TMDC nanosheet (the expected hexagonal crystal orientation) strongly suggests that WS2 sheet grows epitaxially from the edge of the pre-deposited MoS2 nanosheet. Such lateral heterostructures of a WS2-WSe2 heterostructure (each of the respective nanosheets having a p- and n-type characteristics), forms a natural heterojucntion p-n diode.[460]

4.3 Patterned TMDC Nanosheets from the Destructive Thinning and Etching processes. In revisiting the destructive thinning and etching methodologies, the control made available by the laser-,[295] plasma-[296] and chemical-induced[291] techniques, opens the possibility of engineering the thickness of a multi-layered flake to generate desired patterns or even to section TMDC nanosheets. The considerations of complex-nanosheet structures generation make these little-explored production methods particularly intriguing. As the MoS2 band structure depends on the number of layers, the control over the thickness allows the research community to fabricate nanostructured devices with spatially dependent electronic properties, or tailored shapes to meet any potential demand.[291, 404] As a result, 83

84 a number of examples of basic geometric structures have been produced, that could provide the impetus in engineering such novel nanosheet-based hierarchal structures.

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5 Conclusions and Perspectives Research into two-dimensional TMDCs has progressed rapidly since interest was sparked by the discovery of graphene; in little over a decade synthetic techniques have progressed from the laborious peeling of individual flakes to gas-phase deposition on the scale of square centimetres. A wide variety of syntheses have been developed, producing MX2 nanosheets clearly suited to specific applications. Despite such progress the field remains in its infancy, with the influence of many synthetic parameters on the structure and properties remaining poorly understood. The related study of graphene and its derivatives, especially the inclusion of these materials into working devices, has received more attention and is more fully developed. Exfoliation of TMDC nanosheets from bulk materials has developed immensely since the first micromechanical exfoliation reports in 2004. New and developing methods that are based on liquid-mediated exfoliation and small molecule intercalation are potential as industrial processes, but pristine nanosheets are elusive. In this respect, gas-phase growth (such as CVD and PVD) of the ultrathin materials has the most potential to produce highquality nanosheets, using existing technologies in appreciable quantity. However, the lack of a method for production in industrial scale, coupled with the relative underdevelopment of liquid- and-solventless reactions indicates that there remains considerable scope for further investigation before useful techniques emerge. We feel that the research community is still exploring the many avenues of TMDC nanosheet production to strive for a process that can produce high-quality 2D materials in appreciable quantities. In addition, the recent explorations in the formation of heterojunctions and the inclusion of dopants to form ternary compounds remain exciting areas of research, opening the door to new materials with truly novel properties.

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6 Acknowledgements NS and POB thank the Dinswade trust for supporting the PDRA position. JRB acknowledges funding from the Engineering and Physical Sciences Research Council (Core Capability in Chemistry, EPSRC grant number EP/K039547/1).

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7 Figures and Tables

Figure 1. Representation of the forces used to exfoliate nanosheets in ‘top-down’ methods.

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Figure 2. Optical images of (a) TaSe2,[65] (c) MoS2 [64] and (d) WSe2[57] nanosheets (on SiO2/Si substrates) produced by micromechanical exfoliation. (b) The topographic profile of the TaSe2 nanosheets shown in (a) were identified by AFM, indicating the thickness of the flakes. Reprinted by permission of Springer, AIP Publishing LLC and the American Chemical Society.

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Figure 3. Diagrams showing the principles of (a) the sandpaper-assisted rubbing[67] and (b) the wet grinding[156] methods of micromechanical exfoliation. (c and d) The AFM image and height profile depicts the average dimensions of MoS2 nanosheets obtained from the wet-grinding process. Reprinted by permission of Nature Publishing Group and the Royal Society of Chemistry.

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Figure 4. (a) AFM image and (b) topographic line profile of 1H-MoS2, prepared by solvent-free grinding with NaCl, deposited on the surface of freshly cleaved mica from the dispersion in EtOH.[159] Reprinted by permission of the Royal Society of Chemistry.

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Figure 5. (a and b) Photographs of MoS2 and WS2 dispersions in various ethanol/water mixtures. (c and d) The absorbance of the MoS2 and WS2 suspensions in ethanol/water mixtures with different composition are shown as dots.[178] Reprinted by permission of John Wiley and Sons.

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Figure 6. (a) Photograph of the MoS2 dispersions in NMP, prepared by probe sonicaltion at different ultrasonic power intensities. (b) The plot of the intensity of A excitonic peak (after background subtraction) in the UV-vis spectrum as a function of ultrasonic power. (c-f) Typical bright field TEM images of MoS2 nanoflakes prepared with 100, 250, 320 and 400 W, respectively.[174] Reprinted by permission of the Royal Society of Chemistry.

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Figure 7. (a) Photographs of the supernatants of exfoliation of MoS2 flakes in (1) ethanol in the absence of PVP, (2) ethanol in the presence of PVP, and (3) water in the absence of PVP. (b) Photographs of MoS2PVP dispersed in different solvents.[183] Reprinted by permission of John Wiley and Sons.

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Figure 8. (a) Photograph of the dispersions after size selection (these dispersions have been diluted by a factor of 100 to emphasize the colour change). (b) Both the mean (black and red) and maximum (blue) flake sizes observed as a function of final centrifugation rate. The dashed lines represent the sizes of the flakes in the normal starting dispersion, which has been sonicated for 50 h and centrifuged once at 1500 rpm for 45 min. (c-e) Sample TEM images of flakes after size selection for different final centrifugation rates.[181] Reprinted by permission of the American Chemical Society.

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Figure 9. (a) Photograph of the raw molybdenite slag, and its exfoliated nanosheets in NMP. (b) HAADFSTEM image and (c-f) EDX mapping, showing the distribution of Mo, S, Ca and O, respectively.[182] Reprinted by permission of the Royal Society of Chemistry.

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Figure 10. (a) Photograph of the Vitamix blender used by Coleman et al.. (b) Litre-scale exfoliated dispersions of (left to right) MoS2, BN, and WS2.[170]. Reprinted by permission of the American Chemical Society.

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Figure 11. Mechanistic representation of the phase changes that are observed in the exfoliation of MoS2 by lithium intercalation.

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Figure 12. (a) STEM and (b and c) Fourier transform images of a region where a boundary between 2H(left) and 1T’-(right) phases is observed. (d) The phase boundary is parallel to {-1100} planes and marked by a dashed line.[222] Reprinted by permission of the American Chemical Society.

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Figure 13. A schematic representation of the lithium intercalation and exfoliation process of MoS2 nanosheets.[271] Reprinted by permission of Elsevier.

100

101

Figure 14. A photograph of ce-MoS2 dispersed in various solvents after settling for one month.[258] Reprinted by permission of the Royal Society of Chemistry.

101

102

Figure 15. Schematic illustration of the processes for TMDC nanosheet exfoliation by electrochemical lithium intercalation.[267] Reprinted by permission of John Wiley and Sons.

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Figure 16. (a) A photograph of the acid-assisted MoS2 electrochemical exfoliation experiment setup. [277] (b) A TEM image of an exfoliated MoS2 nanosheet obtained from the electrochemical exfoliation process; the inserted image is a higher resolution TEM image of a small section of the sheet. (c) A schematic illustration of the acid intercalation and exfoliation process.[278] (d) Schematic illustration of the careful exfoliation of an acid-intercalated MoS2 pellet.[274] Reprinted by permission of Elsevier, the American Chemical Society and the Royal Society of Chemistry.

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Figure 17. (a) A schematic drawing of the supercritical CO2 intercalation device. (b) A photograph of the low concentration ethanol solutions of the exfoliated MoS2 and WS2 .[280] Reprinted by permission of the American Chemical Society.

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Figure 18. Diagrams briefly explaining the production of TMDC nanosheets by (a) N2H4 treatment and preexfoliation (the results of the preexfoliation are shown in photographs (b) and (c)) followed by (d) the reaction of pre-exfoliated MoS2 with metal naphthelide salts to form an intercalation sample, then exfoliation to single-layer sheets in water.[213] (e) The schematic of exfoliation mechanism of the exfoliation of MoO3-decorated MoS2 nanosheets.[284] Reprinted by permission of Nature Publishing Group and John Wiley and Sons.

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Figure 19. (a) A photograph of MoS2 dispersions (after centrifugation) by 0.5 hour of grinding and different durations of sonication. Left: 3 minutes, middle: 10 minutes and right: 20 minutes of sonication. (b) AFM of MoS2 nanosheets exfoliated by combined grinding (3 h) and sonication (2 h); scale bar is 100 nm. (c) AFM height profile across the MoS2 nanosheets in panel (b).[286] (d) UV−vis spectra of MoS2 samples exfoliated with grinding in various solvents. Inset: photo of the suspensions ground with (i) NMP, (ii) acetonitrile (iii) hexane and (iv) cyclohexane.[288] (e) A TEM image of a MoS2 sample processed from a dispersion of liquid N2-quenched exfoliated MoS2; inset: SAED pattern of the monocrystalline MoS2 in the recorded area.[289] Reprinted by permission of John Wiley and Sons, the American Chemical Society and the Royal Society of Chemistry.

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Figure 20. Photographs of (a) as-prepared MoS2 dispersions after bath sonication for 2 h using different alkali metal hydroxides (LiOH, NaOH, and KOH) and (b) the same dispersions after standing for one week.[212] Reprinted by permission of American Chemical Society.

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Figure 21. (a-d) Optical microscopy images of a single MoS2 nanosheet after (a) 0 h, (b) 2 h, (c) 4 h and (d) 7 h annealing.[294] (e-h) AFM images of mono- and bilayer MoS2 oxidized at 320 °C for (e) 1, (f) 3, (g) 4, and (h) 6 h. The scale bars are 2 µm.[293] Optical photographs of a bilayer MoS2 nanosheet before (i) and after (j) the simultaneous etching and thinning into a monolayer MoS2 nanomesh by thermal annealing at 330 ° C for 15 h. (k and l) AFM height and phase images of the MoS2 nanosheet after thermal annealing.[292] Reprinted by permission of the Royal Society of Chemistry, American Chemical Society and John Wiley and Sons.

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Figure 22. AFM images of (a) a pristine eight-layer MoS2 and (b) thinned MoS2 after one cycle with its accomapying heigh profiles (c and d).[294] Reprinted by permission of the Royal Society of Chemistry.

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Figure 23. (a) An illustration for the laser-thinning of a multilayer MoS2 sheet to a monolayer. (b and c) Optical microscopy images of a multi-layered MoS2 nanosheets (b) before and (c) after scanning a 514 nm laser (10 mW incident power).[295] Optical (d and e) and AFM (f and g) images of a bilayer MoS2 before and after Ar+ plasma thinning.[296] (h and i) Optical images of a MoS2 nanosheet (h) before and (i) after etching to produce a monolayer section.[291] Reprinted by permission of American Chemical Society and Springer.

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Figure 24. Schematic respresentation of the growth of a nanosheet from molecular precursors.

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Figure 25. (a and b) The reaction conditions for the growth of TMDC nanosheets by (a) LP-PVD[309] and (b) LP-CVD[304] methods. Optical images of (c) a monolayer MoSe2 nanosheet island[304] and (d) the edge of a MoS2 nanosheet film,[352] obtained by gas-phase syntheses. Reprinted by permission of the American Chemical Society.

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Figure 26. (a and b) The probable overlap of MoS2 units on (a) fluorophlogopite mica[313] and (b) sapphire[331] substrates during growth. (c and d) SEM images showing the role of the templating substrate in the growth of NbS2 nanosheets, using (c) graphene and (d) Si substrates.[321] (e) WS2 nanosheets grown on a h-BN flake.[323] Reprinted by permission of the American Chemical Society and the Royal Society of Chemistry.

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Figure 27. (a and b) Typical growth results obtained (a) with and (b) without using PTAS as a seeding promoter. The insets are AFM images of the MoS2 nanosheet films (in a) and MoS2 nanoparticles (in b, supported by corresponding heigh profiles).[324] Reprinted by permission of the American Chemical Society.

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Figure 28. A representative mechanism in the processes observed in typical CVD reaction.[332] Reprinted by permission of Nature Publishing Group.

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Figure 29. (a-d) Optical images of the MoS2 materials grown with widely different total pressures (2 Torr, 50 Torr, 250 Torr and 750 Torr, respectively). (e) Comparison of the Raman A1g and E2g peak shift for the same samples.[332] Reprinted by permission of Nature Publishing Group.

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Figure 30. (a) Photograph of the batch production feature of MoS2 on Au foils. (b-e) SEM images of the MoS2 triangular flakes grown at various temperatures displaying different domain sizes; all are produced as only monolayers.[314] Reprinted by permission of the American Chemical Society.

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Figure 31. (a) Schematic diagram of the layer-controlled CVD growth of large-area MoS2 films after the surface plasma treatment. (b-d) Optical images of uniform (b) mono, (c) bi, and (d) trilayer MoS2 films. (e-g) Corresponding AFM images of the marked areas of the films.[344] Reprinted by permission of the Royal Society of Chemistry.

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Figure 32. (a-c) SEM images showing the LP-CVD growth of MoS2 on mica between 35-150 min growth. Inset in (b) is the AFM profile of the MoS2 flake (scale bar 200 nm). (d) Statistical relation of surface nucleation density with flake size obtained from samples shown in (a-c) (left to right, N = 60); numbers above the graphs are the averaged flake sizes.[313] (e and f) Growth using patterned SiO2 substrates (long rectangular bars of few nm-thick SiO2) where the nucleation is clearly affiliated to these bars prior to nanosheet film growth. The scale bars are set at 40 µm.[306] Reprinted by permission of the American Chemical Society and Nature Publishing Group.

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Figure 33. (a-c) Optical images of as-synthesized MoS2 monolayers at different MoS2 evaporation temperatures.[345] Reprinted by permission of John Wiley and Sons.

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Figure 34. (a-h) Optical microscope images of WSe2 samples grown at different temperatures under a designed flow rate for 20 min: (a) 750 °C, 100 sccm (inset: SEM image of same size sample, indicating high density; scale bar, 5um); (b) 765 °C, 100 sccm; (c) 780 °C, 100 sccm; (d) 795 °C, 100 sccm. All of the scale bars are 20 µm.[340] (e-j) SEM images of MoS2 grown on Au foils under various H2 flow rates (0-10 sccm, growth temperature 680 °C, 50 sccm Ar carrier gas).[348] Reprinted by permission of the American Chemical Society and John Wiley and Sons.

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Figure 35. (a-f) The effect of pressure on the growth processes. Optical images of MoS2 films grown at (a) 0.5-1 kPa, (b) 1-4 kPa, 4-10 kPa, 10-40 kPa, 40-80 kPa and (f) 80-120 kPa.[306] (i-n) AFM images for monolayer MoS2 nanosheet films grown at differing evaporation temperatures ((i) 650 °C, (j) 600 °C, (k) 600 °C, (l) 560 °C, (m) 540 °C and (n) 540 °C). The inserted shapes represent the different grain sizes observed in each film (20, 80, 100, 200, 400 and 600 nm, respectively).[325] Reprinted by permission of Nature Publishing Group and the American Chemical Society.

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Figure 36. (a-c) Integrated photoluminescence intensity maps for samples of (a) I-MoS2, (b) D-MoS2 and (c) MoS2-x class nanosheets.[349] Reprinted by permission of the American Chemical Society.

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Figure 37. (a and b) TEM images of (a) tilt-growthand (b) mirror-twin grain boundaries.[305] Insets show the diffraction pattern, indicating the two crystal orientations are 40° and 0° (± 0.5°) apart, respectively. (c) Colour-coded overlay of DF-TEM images of the tilt-growth crystal, corresponding with the two red- and cyan-circled spots in a shows a tilt grain boundary as a faceted line connecting the two triangles. (d) Dark-field image shows the mirror-twin crystal have different diffraction intensity; The small triangle in the centre is a multilayer MoS2 nanoparticle. (e) False-colour dark-field TEM image showing the presence of two adjacent grains. The white solid line marks the grain boundaries, the yellow dashed line marks the overlapped grain junction.[306] Reprinted by permission of Nature Publishing Group

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Figure 38. (a-c) ADF-STEM images of (a) an 18.5° MoS2 grain boundary consisting of 5|7 and 6|8 dislocation rings, which are magnified in (b and c). (d and e) Two types of 4|6 dislocation rings, often observed in 17.5° grain boundary, are also magnified.[352] (f and g) An ADF-STEM image of (f) a mirror-twin (60°) MoS2 grain boundary consisting of (g) 4|4 and 4|8 dislocation rings.[305] (h) ADFSTEM image of an overlapped junction between two MoS2 grains with the inserted FFT of the image showing two sets of MoS2 diffraction spots.[306] Reprinted by permission of the American Chemical Society and Nature Publishing Group.

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Figure 39. (a and b) Optical images of MoS2 nanosheet islands containing a twin- and 40° tilt-boundaries, respectively. (c and d) The space-resolved photoluminescence quantum yields measurements for the islands in (a and b). A 50 % quenching and 100 % enhancement was observed for the islands with twinand tilt-boundary, respectively. (e anf f) Maps of the peak position for the two Raman modes for the same monolayer tilt boundary represents the E2g mode. Both modes show an upshift of 1 cm‐1 at the grain boundary. Scale bar 5 µm.[305] Reprinted by permission of Nature Publishing Group.

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Figure 40. (a) Schematic of the ultraviolet-assisted oxidation of nanosheet films, developed by Ly et al.[353] (b to g) Typical AFM images for 200, 400, and 600 nm grain-sized MoS2 thin films before (b-d) and after (e-g) oxidation at 250 °C in air for 15 min.[325] SEM images of as-grown monolayer WS2 islands before (h-k) and after (l-o) a mild oxidation process, with (p-s) models of the interplay between grain boundary propagation and flake shape evolution for the four typical flakes.[319] Reprinted by permission of the American Chemical Society.

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Figure 41. (a-f) Atomic resolution ADF images of various intrinsic point defects present in monolayer MoS2 grown by LP-CVD.[352] (g-l) Numerous spectroscopic views of a MoS2 nanosheet island containing adlayer impurities. (g and k) Raman map of the separation between the A1g and E2g peaks of one MoS2 flake. (h and k) Map of the photoluminescence intensity on the same MoS2 domain. (i) AFM and (j) MIMIm images of the same domain. (l) Line cuts of the height (blue) and MIM-Im signal (red) in (i) and (j). All scale bars are 5 µm.[364] Reprinted by permission of the American Chemical Society.

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Figure 42. (a and b) Optical images of MoS2 FETs assembled (a) along the grain of a tilt-boundary island and (b) perpendicular through the grain of the twin-boundary crystal. (c) An optical image of an electronic device, created from an MoS2 nanosheet film, with varying channel widths designed for the channel length dependency measurements. (d and e) Estimated values in µe and ION/IOFF of devices in (c) as a function of channel length.[368] Reprinted by permission of the American Chemical Society.

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Figure 43. Schematic illustrations for the synthesis of TMDC nanosheet films by the sulfurisation of MO3 films (produced by (a) the evaporation of a MoO3 solution on a sapphire substrate[372] and (b) the ALDgrowth of a WO3 film[301]). Reprinted by permission of the Royal Society of Chemistry and the American Chemical Society.

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Figure 44. (a) A photograph of as-grown MoS2 on SiO2/Si substrates grown with Mo seed layer of (left-toright) 20, 15, 10, 7, 5, 2, and 0.3 nm. (b and c) TEM images of MoS2 grown with (b) 10 nm Mo and (c) 0.3 nm Mo seed layers.[380] (d-f) Optical and AFM images of the (d) mono-, (e) bi- and (f) tetra-layer ALDbased WS2 nanosheets produced. (g) An atomic resolution TEM image of the WS2 film in (d), exhibiting a high level of polycrystallinity and defects. The insert show the fast Fourier transformation (FFT) of the corresponding TEM image.[301] Reprinted by permission of the American Chemical Society.

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Figure 45. (a-d) Atomic force micrographs of MoS2, produced by Mo film sulfidation at (a) 500, (b) 700, (c) 900 and (d) 1100 °C.[56] (e and f) AFM images of the nanosheets (and accompanying grain height profiles) produced from the sulfidation of (e) MoO3 and (f) Mo metal films on SiO2/Si.[378] Reprinted by permission of AIP Publishing LLC and the Royal Society of Chemistry.

132

133

Figure 46. (a) A schematic for the synthesis of MoS2 films from a Mo-Au surface alloy: treatment of Mo(CO)6 onto a gold thin film forms the Mo-Au surface alloy, which further reacts with H2S to form MoS2 atomic layers in both a large scale and (b) a specific pattern. (c) AFM micrograph showing the surface morphology of a sulfurised Mo-Au surface alloy. The white arrows indicate small dimples formed upon the coalescence of Au vacancies. (d) The representation of the MoS2 film (containing a nanosheet film and clusters) formed on top of the reacted Au surface.[377] Reprinted by permission of John Wiley and Sons.

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Figure 47. (a) Schematic illustration of one growth cycle of an ALD-synthesised MoS2 film. (b-d) AFM images and height profiles of (b) 20, (c) 10 and (e) 5 ALD growth cycles on a sapphire substrate, with measured heights of 3.2, 1.7 and 0.6 nm, respectively.[383] Reprinted by permission of the Royal Society of Chemistry.

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135

Figure 48. (a) schematic for the controlled synthesis and cleavage of mono-to-few layer MoS2 from MoO2 microplates. (b and c) Optical and (insets) AFM images of transferred MoS2 flakes synthesized by annealing for 0.5 and 6 h, respectively.[316] Reprinted by permission of the American Chemical Society.

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136

Figure 49. Schematic representation of the effects of the sulfur source on the products formed in the growth of Group IV and V TMDCs in the liquid phase.

136

137

Figure 50. (a) HRTEM side-view image of multiply stacked ZrS2 trilayers, with a regularly spaced gap of ∼1.5 nm corresponding to the length of oleylamine surfactant layers. (b-d) TEM images of assembled ZrS2 trilayer pillars (grown by the ZrCl4/CS2 hot-injection reaction), with to orientation of the pillars assessed by rotating the TEM holder from (b) 0 to (c) 30 and (d) 60°.[395] (e) Side-on and (f) top-down views of the TEM-imaged monolayer TiS2 nanosheets in colloidal pillar networks (34 and 50 nm across, respectively), grown from the TiCl4/S thermal reaction.[396] (g) Top and (h) side views of ~ 30 nm thick TiSe2 nanocrystals, grown from the TiCl4/Se reaction.[393] Reprinted by permission of the American Chemical Society and John Wiley and Sons.

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Figure 51. (a) Equation showing the formation of single- and multilayer ZrS2 nanosheets with the use of different chalcogen precursors. (b-f) TEM images of the single-layer (b) ZrS2, (d) TiS2 and (f) HfS2 nanosheets produced by the reaction of MCl4 with 1-DDT in oleylamine; the inserts (c, e and g) are pseudocolor images of a small area to enhance thickness contrast. (h) The experimental setup for the growth of ZrS2 nanosheets with a regulated H2S gas supply, accompanied by TEM images of the ZrS2 nanodisks (i) and nanosheets (j) produced by using rapid and slow H2S supply, respectively.[397] Reprinted by permission of the American Chemical Society.

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Figure 52. (a) The mechanisms for the growth of WS2 nanosheet aggregates by the hot injection reaction of WCl6/S. (b) SEM and (c) TEM images of WS2 nanoflakes produced by the reaction in oleylamine.[392] (d) A schematic for the hot injection process to produce 2H- and 1T’-WS2. (e and f) Photographs of the as-prepared dispersion. HRTEM images, confirming the crystal structure of the nanosheets obtained.[400] Reprinted by permission of John Wiley and Sons and the American Chemical Society.

139

140

Figure 53. (a and b) Schematic illustration of the formation of MoS2 films by (a) dip-coating[402] and (b) spin-coating[403] substrates with [H4N]2[MoS4]. (c and d) TEM images show the polycrystallinity of the films obtained by these methods, before annealing in sulfur-rich environments. Reprinted by permission of the American Chemical Society and John Wiley and Sons.

140

141

Figure 54. (a and b) Optical micrographs of exfoliated (a) (Mo0.47W0.53)S2[428] and (b) (Mo0.66W0.34 )Se2 [434] flakes deposited on the 300 nm SiO2/Si substrate. (c) A representative STEM image of a (Mo0.47W0.53)S2 monolayer; the z-contrast allows for determination of the metal content in the TMDCs. (d) Raman spectra of (Mo1-x Wx )Se2 monolayers with different metal compositions.[434] (e) Normalised PL spectra of (Mo1-xWx )S2 monolayers with different metal compositions. [342] Reprinted by permission of the Royal Society of Chemistry and the American Chemical Society.

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Figure 55. (a-c) SEM images of films produced from few layer (a) Mo(S0.5Se0.5)2 , (b) MoS2 and (c) MoSe2 nanosheet dispersions and (d-f) corresponding EDS data.[435] (g) Synthesis of the (NbxW1-x)S2 nanosheets (by lithium intercalation and exfoliation) from pre-fabricated (NbxW1-x)S2 coin-roll nanowires. (h and i) TEM images of the (h) as-synthesized (NbxW1-x)S2 coin roll nanowires and (i) the resulting (NbxW1-x)S2 nanosheets.[436] Reprinted by permission of the Royal Society of Chemistry.

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Figure 56. (a) Illustration of three-zone furnace for the growth of the Mo(Sx Se1-x)2 monolayer by MoS2 and MoSe2 vapours. (b) A photograph of bare SiO2/Si substrate, a MoS2 and Mo(S0.80Se0.20 )2 monolayer film on SiO2/Si substrates. (c and d) Optical images of as-grown Mo(S0.80Se0.20)2 monolayers as (c) islands and (d) a film. (e) Low and (f) high magnification HAADF-STEM image in false colour. Mo sites: yellow; S2 sites: light blue; SSe sites: yellow, and Se2 sites: red. The red square in (e) outlines the part of the image shown in (f).[345] Reprinted by permission of John Wiley and Sons.

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Figure 57. (a) A representation of the process for the reaction of MoO3 with liquid-based chalcogen sources for the formation of Mo(SxSe1−x)2. (b) a photoluminescence spectrum of the complete range of Mo(SxSe1−x )2 nanosheets produce by the exposure of MoO3 to different ratios of S and Se vapours in the gas phase; the insert shows a typical photoluminescence map of a single ternary nanosheet (scale bar, 7 um) showing uniformity.[437] (c-e) Photoluminescence spectra obtained on a regular grid across a 20 µm × 20 µm areas of a (c) Mo(S0.21Se0.79)2, (d) Mo(S0.85Se0.15)2 and (e) Mo(S0.97Se0.03 )2 alloy monolayers produced by this method, showing no meaningful spatial variation in the emission energy is observed within each sheet.[346] Reprinted by permission of the American Chemical Society and John Wiley and Sons.

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Figure 58. (a and d) Schematic illustrations of the chalcogen exchange reactions; vaporized selenium(sulfur) is carried to the heated MoS2(MoSe2) flakes. Optical and (inset) AFM images of the nanosheets before (b and e) and after (c and f) the chalcogen exchange process at 800°C, on sapphire. (gk) TEM images of MoS2 flakes selenised at (g) 600 °C and (h) 700 °C; (i-k) MoSe2 sulfurized at 700°C.[438] (l) Photoluminescence spectra collected from numerous locations on a nanosheet crystal obtained from MoS2 selenisation at 900 °C, the inhomogeneity is observed by the red line.[347] Reprinted by permission of Fronteirs and John Wiley and Sons.

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Figure 59. (a) Optical microscope image of a WSe2/MoS2 hetero-bilayer on a Si/SiO2 substrate. (b) The selected-area electron diffraction pattern of the hetero-bilayer, with the pattern of MoS2 (green) and WSe2 (blue) indexed.[462] (c-g) Optical images for a series of MoS2 bilayers with a stacking angle distributed from θ = 0 to 60. (h-l) The corresponding false color-coded second harmonic intensity mappings of flakes shown in (c-g). The scale bar is 5 µm.[446] (m) Photoluminescence spectroscopy for 532 nm excitation for the bottom (as grown, grey line) and top (transferred, green line) monolayers and tBLs at various twist angles θ. The spectra are offset for clarity and color-coded by θ as indicated with black corresponding to 0° and yellow to 60°. The black dashed lines highlight relevant tuning features.[417] Reprinted by permission of Proceedings of the National Academy of Sciences of the United States of America and the American Chemical Society.

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Figure 60. (a) A schematic of the synthesis process for both heterostructures based in the Zhou/Ajayan synthesis protocol. (c) Schematic and (c) optical images of the vertically stacked WS2/MoS2 heterostructures synthesized at 850 °C, showing the bilayer feature. (d) Schematic and (e) optical images of theWS2/MoS2 in-plane heterojunctions grown at 650 °C. (f) The Z-contrast image of the step edge of the WS2 /MoS2 bilayer. The green dashed line indicates the step edge, and the two triangles indicate the 60° orientation of the MoS2 (top part of image) and WS2 (bottom part) layers, indicative of a 2Hstructure. (g) Schematic of vertical and horizontal MoS2 /WS2 heterojunctions.[459] Reprinted by permission of Nature Publishing Group.

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Figure 61. (a, and b) SEM image of MoSe2-WSe2 monolayer heterojunctions grown by the Wu/SanchezPVD process (scale bars 10 µm). (c) An ADF-STEM image of an area of the interface between MoSe2 (darker) and WSe2 (brighter); scale bar, 2 nm.[461] (d) Atomic-resolution Z-contrast STEM images of the in-plane interface between the (top) MoSe2 -WS2 (scale bar 2nm) and MoS2-WS2 monolayer heterojunctions (scale bar 1 nm). In both, a small roughness in the transition can be seen as a consequence of monomer mixing between the two growth phases. Both parts of the heterojunctions retain the same atomic planes within the structures.[459] (e-j) Raman maps at wavenumbers intrinsic to the bands observed in WS2 (419 and 665 cm-1) and WSe2 (256 and 775 cm-1), demonstrating that WS2 is localized at the centre region and WSe2 is located at the periphary of the triangular domain; scale bar, 5 µm.[460] Reprinted by permission of Nature Publishing Group.

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Figure 62. (a-c) Optical images of a multi-layered MoS2 flake at different steps during the laser-controlled thinning and patterning process: (a) initial flake, (b) after laser-thinning, (c) after patterning it into ribbons with different widths.[295] (d and e) AFM micrograph of the single-layer MoS2 ribbons patterned by laser. (f) An optical image of an MoS2 flake covered by few-layer exfoliated graphene (pre-patterned with hexagonal holes by e-beam lithography and oxygen-plasma etching). (g) Hexagonal pits grown in a MoS2 flake from graphene holes after XeF2 etching. (h) Image of one pattern on a MoS2 flake.[291] (i) An optical image of a pristine bilayer MoS2 and (k) periodical MoS2 heterostructure obtained by the plasma thinning of a PMMA-masked MoS2 bilayer (depicted in j). (l) A Raman map generated by the intensities of the E2g peak of the heterostructure. Inset is the cross section of the peak intensity across the sample.[296] Reprinted by permission of the American Chemical Society and Springer.

149

150 Table 1. Summary of selected measurements of MoS2-based FETs and typical nanosheet dimensions isolated through top-down methods.

Theoretical + Top down Theoretical calculations Monolayer MoS2 Bulk MoS2

Electronic properties MoS2 FETs Nanosheet thickness in FET

μe (cm V s-1)

-

Dimensions of nanosheets produced

Total reaction time

Maximum yielda

References

-

-

[463] [464]

Minutes

Very low

ION/IOFF

Range of thicknesses

% monolayer

Average lateral dimension

410 ~ 100

-

-

-

-

1 layer

0.1-22

10 -10

1 layerb

217

106

-

<1 to 15 μm -

-

-

170-80

10

7

-

-

-

-

-

27-25

107

-

-

-

-

-

6

1 to 10’s of layers 1 to few layers

-

<1 to 30 μm 30 to 200 nm 500nm to 1 μm

-

-

< 1 hour

Very low

[68]

1 hour

Quantitative

[159]

8 hours

-

[156]

0.43 mg/mL (14 %) 2 mg/mL (2.4 %) 7.6 mg/mL (8 %)

[164, 181, 202, 203]

2

-1

Micromechanical Basic methods Basic methods with encapsulation for FET performance optimisation

1 layer

c

2 layersc c

4

7

3 layers

15-10

10

Electrostaticassisted

1 layer

0.96

10

Milling with NaCl

-

-

-

Milling by grinding

-

-

-

3

1 to 100’s of layers -

3 to 5 layers

< 0.1%

< 0.1% n/a 0%

[62, 64, 85, 102, 107, 112, 113, 149, 465] [69] [70, 78]

Liquid-mediated exfoliation 2 hours Sonication in NMP

20 layers (approx.)

0.1-0.4

2

4

10 -10

1 to 10’s layers

<10 to 30 %d

50nm to 3 μmd

12 hours 60 hours

150

151 140 hours Sonication in 45% EtOH/H2O High-shear exfoliation

-

-

-

3 to 4 layers

-

-

-

-

1 to 12 layers

~ 8%

> 90 %

100 nm to 2 μm 30 to 180 nm

40 mg/mL (40 %) 0.018 mg/mL (< 1 %) 0.5 mg/mL (1 %)

[178]

2-4 days

up to 60 mg/mL (Quantitative)

[215, 231, 241, 243, 247, 258]

<6 hours

Quantitative

[245]

24 hours

~ 50 %

[271]

2 hours

0.023 mg/mL (1 %)

[171, 285, 286]

12 hours

4 mg/mL (40 %)

< 36 hours

-

2 hours

0.81 mg/mL (81 %; M = Li) Up to 32.5 mg/mL (depending on MoS2 loading ~ 65 %; M =

8 hours

(1 mg/min)

[170]

Lithium Intercalation Basic methods

-

-

-

Mostly 1 layer

Electrochemicallyassisted

-

-

-

1 layer

>95 %

With zerovalent Li

-

-

-

2 to 10 layers

0%

-

-

-

1 to 10 layers

<1 to 10 %

20 to 700 nm

-

-

-

0%

500 nm to 3.5 μm

Hybrid exfoliation methods Grinding/milling and liquid exfoliation Liquid N2 quenching and liquid exfoliation

MOH intercalation and liquid sonication

151

-

-

-

2 to 5 layers

1 to 9 layers

~85 % for 1-3 layers

200 to 800 nm

500 to 700 nm 400 to 600 nm

50 nm to 1 μm

[171, 285, 286] [289]

[212]

152 Na) Acid intercalation Basic intercalation and exfoliation Electrochemicallyassisted

Up to 100 %

50 to 100 nm

< 30 h

~65 %

[272]

-

5 to 50 μm

< 1 hour

0.014 mg/mL (9 %)

[277, 278]

1 to few layer

~ 20 %

500 nm to 2 μm

< 2 hours

Quantitative

[280]

-

1 to 2 layers

> 90 %

1 to 10 μm

-

[213]

-

-

1 to 4 layers

~ 60 %

[264]

1 layer

0.02

10

Selective (from 1 to few layer)

n/a

n/a

1 hour

n/a

[292]

Raman laser ablation

1 layer

0.04-0.49

-

1 layer

n/a

n/a

8μm/min

n/a

[295]

Ar + plasma ablation

-

-

-

n/a

n/a

Up to 2 min

n/a

[296]

Chemical etching

-

-

-

n/a

n/a

Up to 2 min

n/a

[291]

-

-

-

2 layer

1.2-2

10 -10

Supercritical CO2

-

-

-

H2NNH2/ sodium naphthalide

-

-

Hybrid MoS2/MoO3

-

Thermal etching

3

6

Mostly 1 layer 1 to few layers

Other Intercalation

2 to 5 μm

~ 60 hours 1 to 5 hours

Thinning methods 3

Selective (from 1 to few layer) Selective (from 1 to few layer)

(a) For liquid-mediated exfoliations, the value in the parentheses is the yield relative to the starting weight of the bulk MoS2 precursor. (b) MoS2 nanosheet was encapsulated with a 30nm HfO2 layer before FET fabrication. (c) MoS2 nanosheet was encapsulated with a 40-50nm Al2O3 layer before FET fabrication. (d) The sizes and thicknesses of MoS2 nanosheets produced are dependent on a large number of exfoliation conditions.

152

153 Table 2. Summary of selected measurements of MoS2 -based FETs and typical nanosheet dimensions isolated through bottom-up methods.

Electronic properties MoS2 FETs

Bottom up

Nanosheet thickness in FET

2

-

μe (cm V 1 -1 s )

ION/IOFF

Dimensions of nanosheets produced

Total Reaction time

Maximum yield

References

20 min – 4 hours

-

[8, 305, 306, 308-310, 313, 324, 349, 352, 365, 366, 466]

> 1 cm

2 – 12 hours

-

[307, 332, 339, 467-469]

-

-

-

-

[452]

Selective (from 1 to few layer)

-

> 1 cm

30 min – 3 hours

Quantitative

[312, 328, 372, 373, 375, 379, 380, 382]

2 to few

>0%

> 1 cm

< 3 hours

-

[377]

Range of thicknesses

% monolayer

Average lateral dimension

Selective (from 1 to few layer)

Up to 100 %

< 1 to 135 μm

Selective (from 1 to few layer)

Up to 100 %

-

Gas phase TMDC growth: Nanosheets

c

1 layer

0.2-15.3

Films

1 layer

0.003-0.3 (pre anneal) 0.1-17 (after anneal)

Gas-phase growth with encapsulation for FET performance optimisation

1 layer

b

4

6c

10 -10

4

5

10 -10 (pre anneal)

35.5

-

0.0040.04 (pre anneal) 0.1-12 (after anneal)d 0.004

10 -10 (pre anneal) 104-106 (after anneal) -

Other gas phase syntheses

Mo/MoO3 film sulfidation

1 – 5 layers

Mo-Au alloy

2 layers

153

4

5

154 sulfidation MoCl5 sulfidation and annealing

-

-

-

MoO2 microcrystal sulfidation

1 layer

0.3

10

-

-

-

Liquid-phase syntheses Thermal breakdown of [NH4]2[MoS4]

layers Selective (from 1 to 10 layers) 1 to few layers

6

~ 100 %

1 μm sheets or > 1 cm films

< 1 hour

-

[383, 470]

~ 100 %

5 to 20 μm

30 min – 6 hours

-

[316]

1 layere

~ 100 %

< 25 nm

2 hours

Quantitative

[391]

2 to 3 layers

Low

> 1 cm

2 hours

-

[402]

2 layer

Low

> 1 cm

2 hours

-

[403]

Solid-phase TMDC growth from [NH4]2[MoS4] Dip coating and anneal

2 to 3 layers

Spincoating and anneal

2 layer

0.1-4.7 (after anneal) 0.1 (after anneal)

3

6

10 -10 3

10

(a) All carrier mobility measurements were for monolayer MoS2-based FETs unless stated otherwise. (b) MoS2 nanosheet was encapsulated with a 30nm HfO2 layer before FET fabrication. (c) Study by Kim and Lauhon discovered that an influence on stoichiometry gave a variation on the carrier mobilities and on/off ratios observed: sulfur rich MoS2+x – 0.47 cm2 V-1 s-1, 104; stoichiometric MoS2 – 2.66 cm2 V-1 s-1, 105; MoS2-x – 15.3 cm2 V-1 s-1.[349] (d) the high value of 12 cm2 V-1 s-1 was obtained from a 3-4 layer nanosheet.[382] (e) nanosheets produced contains a coating of oleylamine.

154

155

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