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Beam Interactions with Materials and AtomsNuclear Instruments and Methods in Physics Research, Section B.

Title of Paper: Alpha Particle Irradiation of Bulk and Exfoliated MoS2 and WS2

Membranes

Corresponding Author: Liam H. Isherwood (liam.isherwood@postgrad.manchester.ac.uk) Dr. Aliaksandr Baidak (aliaksandr.baidak@manchester.ac.uk)

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Title

Alpha Particle Irradiation of Bulk and Exfoliated MoS2 and WS2 Membranes

Names

Liam H. Isherwooda,b, Robyn E. Worsleya, Cinzia Casiraghia and Aliaksandr Baidaka,b.

Affiliations and Addresses

a School of Chemistry, University of Manchester, Oxford Road, M13 9PL, Manchester,

United Kingdom

b Dalton Cumbrian Facility, University of Manchester, Westlakes Science and Technology

Park, Moor Row, CA24 3HA, Cumbria, United Kingdom

Abstract

The properties of two-dimensional transition metal dichalcogenide (TMDC) nanosheets

have been intensively studied in recent years as these compounds have emerged as promising

materials for future electronic, photonic and sensor applications. Some of these applications

may require the nanosheets to be exposed to radiation fields; therefore, an understanding of

their interaction mechanisms with ionising radiation is required.

In our experiment, we administer 1.66 MeV helium nuclei to bulk and liquid phase

exfoliated MoS2 and WS2 membranes to two total absorbed doses. Raman spectroscopy shows

small changes, within the spectrometer resolution, in all samples. Although small, some

reproducible changes are observed such as a blueshift of the E2 g1 and A1 g modes in the bulk

MoS2 membrane irradiated to a high total absorbed dose; these shifts are accompanied by a

small broadening of both peaks. In bulk WS2 membranes, He2+ irradiation induces a blueshift

and monotonic mean peak width decrease of the 2LA(M) phonon mode with increasing

fluence. The structural changes associated to these peak shifts are currently unknown. Raman

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spectroscopy, in agreement with energy dispersive x-ray spectroscopy, shows that amorphous

carbon is present in the exfoliated MoS2 and WS2 membranes after irradiation. It is suggested

that this amorphous carbon could be formed by radiolytic amorphisation of residual solvent,

retained within the ripplocations of the exfoliated nanosheets.

Due to the morphology of the liquid phase exfoliated nanosheets, they exhibit greater

radiation stability than bulk TMDCs when exposed to heavy charged particle radiation. These

results differ from those previously reported by for monolayer and bulk MoS2 produced by

mechanical exfoliation. The relative radiation stability of the liquid phase exfoliated

nanosheets is expected to arise from attenuation and dissipation of the ion beam energy by the

residual solvent and the amorphous carbon subsequently produced.

1. Introduction

Transition metal dichalcogenides (TMDCs) are layered compounds that exhibit strong

covalent intralayer bonds and weak interlayer Van der Waals interactions. TMDCs are

structured such that each layer consists of three atomic planes: a triangular lattice of transition

metal atoms sandwiched between two triangular lattices of chalcogen atoms (S, Se or Te).

Similarly to other Van der Waals solids, such as graphite, TMDCs can be exfoliated down to

few and single-layers by micromechanical (MME) [1], liquid phase (LPE) [2] or other types

of chemical exfoliation methods. [3] Alternatively, atomically thin TMDCs can be obtained

via chemical vapour deposition (CVD). [4]

TMDCs are structurally similar to each other but have an array of electronic properties

ranging from semiconducting to metallic, depending on their chemical composition,

electronic density, geometry and thickness. [5] For example, single layer molybdenum

disulfide (MoS2) and tungsten disulfide (WS2) are semiconductors with direct bandgaps of

approximately 1.9 eV [6] and 2.05 eV [7] respectively, which make them ideal for opto-

electronics. [8] Because of their complementary properties compared to graphene, TMDCs

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and graphene can be assembled into heterostructures to create a wide range of devices. [9] In

addition, single and few-layer TMDCs have been shown to be effective materials for gas

sensor [10], piezoelectric [11], supercapacitor [12] and bio-sensor applications. [13]

For two-dimensional materials to be utilised in future applications, the exfoliation or

deposition method must be scalable, versatile, economically viable and yield high quality

material, i.e. it possesses a low concentration of defects. [14] Whilst MME produces material

with the highest carrier mobility, it is neither scalable nor economically viable. It is likely that

two-dimensional TMDCs produced by CVD and LPE will be utilised in future applications.

CVD can produce large area and high quality monolayer films. [4] However, the choice of

substrates is limited due to the high temperatures used. Moreover, the TMDC film can

become doped during the transfer procedure to arbitrary substrates, which could compromise

the quality of the material. [15]

In contrast, LPE is a cheap, scalable and versatile technique that enables typically few-

layer nanosheets to be obtained in large quantities via sonication of bulk material in a solvent

of the appropriate surface energy (~70 mJ m-2 for MoS2 and WS2). [2] The dispersed

nanosheets can then be deposited by drop-casting or vacuum filtration. Additionally, the

dispersions can be formulated into inks and printed directly onto flexible and temperature

sensitive substrates to produce electronic devices. [16,17] However, the nanosheets produced

by LPE are often more defective than material obtained via CVD and MME. This is partly

due to the small lateral dimensions of the nanosheets produced by LPE (few hundred nm)

compared with MME (typically micrometres) and CVD (several micrometre sized single

crystals up to wafer scale thin films). [1, 2, 4] The edges of all two-dimensional crystals are

defective; therefore, a TMDC flake with small lateral dimensions has a high ratio of edge

defects relative to the pristine coordination geometry within the basal plane.

A comprehensive understanding of the radiation damage mechanisms in TMDCs is a

prerequisite of their application in devices to be used in radiation environments such as space,

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or as radiation detectors. [18] Currently, only a few works have looked at this problem, and

selected publications are summarised in Table I. This table shows that previous studies have

been conducted on TMDCs produced by CVD and MME methods, whilst no study has been

conducted on solution processed TMDCs produced, for example, by LPE. [27]

The exfoliation or deposition protocol used to produce two-dimensional TMDCs can

influence their interactions with ionising radiation. For example, Ochedowski et al showed

that monolayer MoS2 exhibited different morphologies after swift heavy ion irradiation at

grazing incidence depending on the method of production. [28] MoS2 obtained via CVD

formed ~400 nm rifts whereas mechanically exfoliated monolayers folded to create “half-

nanotube” structures either side of the heavy ion trajectory. The interaction of the swift heavy

ions with water, trapped between the MoS2 and substrate during MME, is thought to

contribute to the different radiation-induced morphologies.

Akin to water trapped during MME, the solvent used during LPE can be retained

between the layers of the nanosheets during sonication. [29] After deposition, low boiling

point solvents can be removed by annealing. However, the most widely used solvent for

obtaining high concentrations of dispersed nanosheets is N-methyl-2-pyrrolidone (NMP); this

solvent has a boiling point of 202 oC and even when post-processing methods are used,

residues are often present. [30] Therefore, it is expected that this residual dispersant may

influence the interactions between heavy charged particles and TMDC nanosheets produced

by LPE. The solvent could attenuate the energy of the incident ions and decrease the rate of

defect production; alternatively, it could facilitate defect formation by producing reactive

radiolytic species. To circumvent the problems of removing high boiling point solvents,

additives such as amphiphilic surfactants and polymers can be added to relatively low boiling

point dispersants such as water. [31, 32] However, it is anticipated that residual additives

could also influence the interactions between heavy charged particles and TMDC nanosheets.

5

As already discussed, TMDC nanosheets produced by LPE possess a higher ratio of

edge defect sites to basal plane centres compared with material obtained via MME or CVD.

Trace amounts of water, present during sonication, are known to stabilise Mo-terminated

edges of MoS2 nanosheets produced by LPE. [33] Radiolysis of these quasi-bound molecules

could influence the radiation damage mechanisms of the nanosheets. Alternatively, if the

nanosheets are annealed and the coordination sites of metal edge centres are vacant then they

could localise radiolytic oxide formation leaving the basal plane undisrupted. Such reactions

could occur preferentially with edge defects as opposed to basal plane centres due to their

higher reactivity. [34] Therefore, depending on how the coordination sites of metal and

chalcogen terminated edges interact with ionising radiation; nanosheets produced by LPE

could have quite different interactions from MME and CVD material.

In this work we present a He2+ irradiation study of membranes made from liquid phase

exfoliated (exfol) MoS2 and WS2, obtained using LPE [2] and vacuum filtration. [35] The

energy transfer mechanism between the incident He2+ ions and the membranes is

predominantly inelastic and the irradiation is carried out at normal incidence. The extent of

radiation damage in the samples of reduced dimensionality is then compared with that of

membranes made from bulk, i.e. unexfoliated MoS2 and WS2.

The nanosheets produced by LPE were characterised by UV-Vis spectroscopy, atomic

force microscopy (AFM) and transmission electron microscopy (TEM). The bulk and

exfoliated membrane thicknesses were measured using stylus profilometry. The anticipated

helium ion ranges and the electronic/nuclear contributions to the energy transfer mechanisms

were determined using the Stopping Range of Ions in Matter (SRIM) programme. [36] The

structural and compositional changes induced by the energetic helium nuclei are probed via

Raman and energy dispersive x-ray (EDX) spectroscopies.

2. Experimental

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Sample preparation: 90 mg of MoS2 (6 μm particle size, from Sigma-Aldrich) is added

to a 250 mL reagent bottle containing 30 mL of NMP (99% ACS reagent grade, from Sigma-

Aldrich). The suspension is sonicated for 7 days (600W, Hilsonic bath) and then centrifuged

at 3500 rpm (903 RCF) for 1 hour (Sigma 1-14K); the supernatant is then collected. In the

case of WS2 (150 mg, 2 μm particle size, from Sigma-Aldrich), the material is added to NMP

(50 mL), following the same procedure as used for MoS2.

UV-Vis spectroscopy: A Varian Cary-5000 UV-Vis-NIR spectrophotometer was used

to measure the absorption spectra of the MoS2 and WS2 dispersions and the concentration of

material was calculated using the Beer-Lambert law. [37] For MoS2 dispersions, an absorption

coefficient of 34 mL mg-1 cm-1 at 672 nm was used. [2] Similarly, WS2 dispersion

concentrations were calculated using an absorption coefficient of 27.56 mL mg-1 cm-1 at 629

nm. [2] Six MoS2 dispersions were combined and a concentration of 61.2 μg mL-1 was

calculated from the mean of the individual dispersion concentrations. The process was

repeated for three WS2 dispersions yielding a concentration of 94.3 μg mL-1.

AFM: Dispersions were diluted (20 μg mL-1) and drop-cast directly onto oxidised

silicon wafers (1 cm2, SiO2 thickness ca. 290 nm, from IDB Technologies). Drop-casted

samples were annealed (1 hr, 300 oC) under constant nitrogen flow. A Bruker Multimode 8

AFM operating in non-contact mode and equipped with a silicon nitride SCANASYST-AIR

tip was used to measure flake dimensions.

TEM: Diluted dispersions (10 μg mL-1) were drop-cast directly onto holey carbon film

400 mesh Cu TEM grids (from Agar Scientific). The grids were annealed (1 hr, 300 oC) under

constant nitrogen flow. A Phillips CM200 TEM was operated at 200 kV to obtain diffraction

patterns and dark-field images.

Membrane fabrication: Dispersions were centrifuged at 13,800 rpm (20440 RCF) for

2 hours (Sigma 3-18KS) to sediment the dispersed TMDC nanosheets. The exfoliated material

was retrieved and re-dispersed in an aliquot of the supernatant before being deposited onto

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polyvinylidene fluoride (PVDF) filters (25mm diameter, 0.22 μm pore size, from Durapore)

by vacuum filtration. The same deposition process was used to fabricate bulk TMDC

membranes using 6.5 mg of MoS2 or 8.5 mg of WS2 in 7 mL of NMP without any sonication.

Acetone (2 x 5 mL) was flushed through each membrane to displace residual NMP.

Profilometry: Qualitative membrane edge thickness measurements were conducted for

all exfoliated and bulk TMDC membranes using a Bruker DektakXT stylus profilometer.

Irradiation parameters: A 5 MV tandem pelletron ion accelerator was operated at

beam currents of ~162 nA for low dose samples, denoted as (L), and ~355 nA for high dose

samples, denoted as (H). It should be noted that non-irradiated control samples are denoted as

(N). 3.8 MeV He2+ ions were administered to samples for 30 minutes per membrane. Titanium

foil (8 μm thick) was used to attenuate the incident ion beam to a reduced energy of 1.66 MeV

and prevent irradiation of the PVDF filter.

Raman spectroscopy: A Renishaw InVia spectrometer (resolution ~2 cm-1) equipped

with a 100x objective lens and 1800 lines mm-1 grating was used. 514.5 nm and 488 nm

excitation wavelengths were used to characterise MoS2 and WS2, respectively. A laser power

of 0.15 mW was used to avoid any thermal damage during measurements.

Scanning electron microscopy (SEM) and EDX: The membranes were bisected and

the sample stub was fixed into a 70o tilted sample stage which was then tilted a further 20o to

allow cross-section imaging and EDX analysis. A field-emission environmental SEM (FEI

Quanta 250) was operated under high vacuum conditions using an accelerating voltage of 15

kV. AZtec EDX software and a windowless X-Max silicon drift detector (Oxford

Instruments) were used to analyse and collect EDX data, respectively.

3. Methods

AFM: The height profile and lateral dimensions of 100 flakes were measured. The

nanosheet height was divided by the thickness of a single TMDC monolayer (0.7 nm) to give

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the number of layers comprising each flake. Histograms were fitted with a Gaussian

distribution and contour plots display flake lateral dimensions as x and y coordinates whilst

the number of TMDC layers comprising the flake is plotted as the z-coordinate. Each black

spot represents a flake measurement (length, width and height). We remark that this analysis

is only qualitative as residual solvent can give single-layer TMDC flakes a thickness larger

than the nominal value of 0.7 nm.

SRIM & Dosimetry: SRIM was used to calculate the reduced energy of the alpha

particles after attenuation. The energy lost via nuclear ( dEdx nuc) and electronic ( dE

dx elec ) transfer

processes was calculated for 3.8 MeV He projectiles travelling in titanium. Simulations

afforded a dEdx nuc

value of -2.13 x 10-1 keV μm-1 and a dEdx elec

value of -2.67 x 102 keV μm-1.

After attenuation by 8 μm thick titanium, the resulting 1.66 MeV He2+ ions have an expected

mean ion range of 2.2 μm into the MoS2 or WS2 layer and no ion has a range greater than 3

μm. However, bulk TMDC densities were used in the SRIM calculation and the porosity of

the membranes mean this depth will be ~2 μm greater for exfoliated membranes than

calculated (40 – 45 % porosity). [35] The accumulated charge gathered by a picoampermeter

was used to calculate the fluence and dose rate. Following this, the ion beam area (0.38 cm2)

and expected range of the alpha particles (3 μm) was used to calculate the mass of irradiated

TMDC and the corresponding total absorbed dose in Grays (Gy) assuming bulk TMDC

densities (5.06 g cm-3 for MoS2 and 7.5 g cm-3 for WS2).

Raman spectroscopy: The Raman spectra of MoS2 and WS2 show two characteristic

peaks corresponding to the doubly degenerate in-plane E2 g1 mode and the out-of-plane A1 g

vibrational mode. In bulk MoS2, these modes have frequencies of 382 cm-1 and 407 cm-1,

respectively. [38] Whereas for bulk WS2, the E2 g1 mode has a frequency of 356 cm-1 and the

A1 g peak is higher in energy at 420 cm-1. [39] In addition to these two characteristic peaks, the

spectrum of WS2 exhibits a third peak redshifted from the E2 g1 signal by ~5 cm-1. This peak is

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generated by a second order Raman overtone scattering event involving a longitudinal

acoustic mode from the M point of the Brillouin zone, hence it is denoted as 2LA(M). [39]

The E2 g1 and A1 g peaks are fitted with a Lorentzian line shape; multi-peak deconvolution is

also used fit the 2LA(M) mode. 10 measurements were taken for (N) membranes whilst 15

and 20 measurements were taken within the irradiated regions of the (L) and (H) membranes,

respectively. The mean average is calculated for the full width at half maximum (FWHM) and

position of each peak. Error bars represent +/- one population standard deviation.

EDX: Transition metal, oxygen, sulfur, carbon and fluorine signals were detected in

EDX spectra. When calculating stoichiometric ratios, the spectra were deconvoluted such that

only the relevant transition metal, oxygen and sulfur contributed to the elemental

composition. The mean values of ratios calculated from multiple measurements were used to

create scatter plots in which error bars correspond to +/- one population standard deviation.

4. Results & Discussion

Figure 1 shows the nanosheet thickness distribution of the TMDC dispersions. The

AFM data shows that the average number of TMDC monolayers comprising each nanosheet

is between 8 and 11 for both the MoS2 (Fig. 1a) and WS2 dispersions (Fig. 1b). This data

clearly shows that both TMDCs have been well exfoliated to few-layer nanosheets. However,

the fitted Gaussian distribution suggests mono- and bilayer flakes comprise a minor fraction

of the dispersions. Figure 2 shows the relationship between the thickness of a nanosheet and

its lateral dimensions. It is evident that thicker flakes possess larger lateral dimensions in both

the MoS2 (Fig. 2a) and WS2 dispersions (Fig. 2b). However, the average flake thicknesses,

between 8 and 11 monolayers, have a wide distribution of lateral dimensions.

Table II shows the profilometry results obtained for the membranes made of exfoliated

TMDCs: the thickness in all cases is around a few micrometres. Note that the exfoliated MoS2

membrane irradiated in the low dose regime (sample: exfol MoS2 (L)) has a thickness smaller

10

than the ~5 μm required from SRIM simulations and membrane porosity. However,

profilometry measurements are taken at the edge of the membrane, where the material is

thinner than in the centre. This is demonstrated in Figure 3, which shows the cross-section

SEM image of the exfol MoS2 (L) sample: it can be seen that the membrane is considerably

thicker than the value of 1.65 μm obtained by profilometry. The bulk MoS2 and WS2

membranes have average thicknesses of 17 μm, and 15 μm, respectively. Therefore,

regardless of the membrane porosities, the ion beam is not expected to interact directly with

the supporting PVDF filter in any of the samples. This enables the irradiation to be conducted

using relatively high ion energies, i.e. principally inelastic energy transfer, and at normal

incidence; whilst eliminating the need to consider radiolytic product formation or secondary

electron emission from the substrate.

Dosimetry results are shown in Table III. The He2+ ion beam deposits on average 2.3

times more energy within (H) membranes compared to (L) samples. As discussed, the ion

range will be greater than that given by SRIM simulations due to the porosity of the

membranes. However, this reduced density is expected to compensate for the larger irradiated

volume when calculating total absorbed doses (Gy). Thermal effects were caused by ion beam

charging in (H) samples and manifested themselves as black circles on the supporting PVDF

filter. These thermal effects must be considered when discussing spectroscopic changes to (H)

samples. No thermal effects were observed in the (L) samples.

Let us now analyse the Raman spectra of the bulk WS2 membranes. Figure 4 shows the

mean values of the peak position and FWHM of the in-plane E2 g1 vibrational mode for (N),

(L) and (H) samples. In all membranes the changes in position and FWHM are comparable

with the resolution of the spectrometer. The largest change is observed in the bulk WS2 (H)

sample, where the E2 g1 mode exhibits a FWHM increase of 1.6 cm-1 compared with the (N)

membrane. The bulk WS2 data shown in Figure 4 were fitted using a single Lorentzian peak.

However, the data were also fitted with multi-peak Lorentzian lines such that the E2 g1 peak is

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deconvoluted to accommodate the 2LA(M) mode. Figure 5a shows that the E2 g1 peak position

and FWHM of the (N) and (L) samples remain unchanged after multi-peak Lorentzian fitting;

whilst the FWHM increase of the bulk WS2 (H) sample is no longer observed. Figure 5b

shows the average frequency and FWHM of the 2LA(M) mode. We can observe that the

FWHM decreases monotonically and the peak position blueshifts with increasing dose.

Therefore, the E2 g1 peak broadening in Figure 4 can now correctly be attributed to a radiation

or thermally induced blueshift of the 2LA(M) mode which increases the FWHM of the E2 g1

when using a single Lorentzian peak fit. Figure 5c shows representative Raman spectra of

bulk WS2 to illustrate the blueshift of the 2LA(M) mode.

Moving to the Raman spectra of bulk MoS2, Figure 6 shows the average FWHM and

position of the E2 g1 and A1g modes for (N), (L) and (H) samples. The (L) sample shows no

peak shift or broadening. Whereas, the A1 g peak broadens and exhibits an average blueshift of

1.2 cm-1 after high dose irradiation (Fig. 6a). Although this shift is within the resolution of the

spectrometer, the error bars show that it occurs reproducibly. The same trend is observed for

the E2 g1 mode (Fig. 6b): low dose irradiation has a negligible effect on the peak parameters,

whereas the (H) sample exhibits a reproducible 1.2 cm-1 blueshift and a 0.85 cm-1 FWHM

increase. Figure 6c shows representative Raman spectra illustrating these changes.

The phonon confinement model suggests that if radiation introduces defects into the

bulk MoS2 crystal lattice then the reduction in the phonon lifetime should increase the FWHM

of the E2 g1 and A1 g modes. The relaxation of the fundamental Raman selection rule then

facilitates a redshift of the E2 g1 and a blueshift of the A1 g mode. [19] No divergence of peak

positions is observed in our data. Tensile strain causes a redshift and FWHM increase of the

E2 g1 in monolayer MoS2. [40] Therefore compressive strain could account for the blueshift of

the E2 g1 mode in our data; however it should cause negligible changes to the A1 g peak. Finally,

sulfur vacancy or oxide induced p-type doping effects are known to induce a blueshift of the

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A1g mode in MoS2 but decrease its FWHM due to electron-phonon coupling. [23] Hence,

defects, strain and doping fail to account for the small but reproducible blueshift of the peaks

in bulk MoS2 (H).

It is possible that combinations of the above processes are convoluted with thermal

effects due to ion beam charging of (H) samples. This makes the analysis of the (H) samples

particularly challenging. Alternatively, because the energy deposition of the He2+ ions is

relatively constant with regards to time, a large proportion of this energy will be localised as

their velocity decreases considerably, i.e. at the Bragg peak, which is located at least 3 μm

from the surface. Therefore, conventional Raman spectroscopy, which probes only a few

hundred nanometres under the surface, may not be sensitive enough to detect radiolytic

changes occurring in the deeper regions of the TMDC membranes; giving rise only to small

changes, within the resolution of the spectrometer.

To complement the discussion of radiation effects on the Raman spectra of bulk TMDC

membranes, we now consider the exfoliated WS2 samples. The peak position and FWHM of

the E2 g1 mode for (N) and irradiated samples, obtained by a single peak Lorentizian fit, is

shown in Figure 7a. Radiation causes no discernible changes for either the (L) or (H) doses.

Moreover, the invariance of the peak parameters for the (H) dose suggests that thermal effects

did not alter the vibrational properties of the exfoliated WS2. The A1 g mode shows the same

radiation and thermally resistant trend. For bulk WS2 it was the 2LA(M) mode which shifted

to higher energy and exhibited a monotonic FWHM decrease. Figure 7b shows the FWHM

and peak position of the 2LA(M) mode of exfoliated WS2 obtained after deconvolution. The

mean value of the 2LA(M) frequency blueshifts upon irradiation, however, it is not

proportional to the fluence as the (L) sample exhibits a larger average blueshift and narrower

distribution compared with the (H) sample. Conversely to the bulk material, the FWHM of the

2LA(M) mode in exfoliated (L) and (H) samples increases. However, it should be noted that

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the 2LA(M) peak parameters of the (N) sample have a large variance which is caused partly

by the error associated with the deconvolution procedure. [39]

Recently, Ma et al showed that 60 keV Ar+ ion irradiation decreases the E2 g1 ¿ A1 g

intensity ratio of monolayer WS2 produced by CVD; the magnitude of the reduction was

correlated with increasing fluence. [41] In Raman spectroscopy, the sample geometry strongly

influences the intensity of the signal produced. Samples produced by CVD and MME can be

reproducibly analysed in backscattering geometry meaning that the angle dependence of the

laser on the Raman intensity does not influence the measurement. However, in this work the

random orientation of flakes within each membrane means that peak intensities cannot be

used as a reliable metric for radiation damage in our samples.

Finally, the mean values and distributions of the A1 g peak frequency and FWHM for the

exfol MoS2 (N), (L) and (H) systems is shown in Figure 8. Although the average energy of

the A1 gmode increases with fluence, the magnitude of the change between (N) and (H)

membranes is less than 0.5 cm-1 and does not occur reproducibly. The E2 g1 peak parameters

show the same trend observed for A1g. Therefore, upon exfoliation to few-layer nanosheets,

there appears to be no radiolytically or thermally induced changes to the vibrational spectrum

of exfoliated MoS2 or WS2.

The magnitude of the peak shifts and FWHM changes in the Raman spectra of

irradiated bulk MoS2 and WS2 membranes are greater than those of exfoliated TMDCs;

however, the differences between their vibrational properties are small. The threshold at

which quantum confinement effects begin to alter the band structure of MoS2 occurs at ~2.8

nm thickness, corresponding to 4 monolayers. [42] Figure 1 shows that the majority of the

TMDC nanosheets in the exfoliated membranes have thicknesses greater than 4 monolayers.

Therefore, from an electronic and optical point of view, our results show that the bulk

TMDCs and exfoliated nanosheets have similar interactions with the energetic alpha particles

under our experimental conditions.

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Although the properties of the nanosheets are similar to the bulk material, the structures

of the membranes made from these two materials are very different. In the case of exfoliated

membranes, the porosity is 40-50% greater compared with that of bulk TMDCs and the

membranes made from bulk material are even more porous due to inefficient stacking of the

particles. [35] Furthermore, membranes that contain exfoliated nanosheets are likely to retain

some residual solvent. [29] In our case this solvent is NMP and its high boiling point means it

is very difficult to remove completely, even after postprocessing and annealing. [30]

Nevertheless, NMP was used in this study because it produces concentrated dispersions

and enables membranes of sufficient thickness to be fabricated. As a consequence, the

radiation chemistry of the substrate effects does not influence the ion-TMDC interactions. The

PVDF filters could not be annealed above 140 oC, therefore the membranes were flushed with

acetone after fabrication to displace residual NMP from internal pores of the membranes.

Another observation derived by Raman spectroscopy is shown in Figure 9 which plots

the extended Raman spectrum of the bulk and exfol MoS2 (H) samples. Two broad peaks are

observed at 1350 cm-1 and 1600 cm-1, respectively, which are associated to the D and G peaks,

typically observed in amorphous carbon. [43] These peaks are present in exfol MoS2 and WS2

membranes after irradiation yet are absent in the bulk TMDC (L) and (H) spectra.

Figure 10 shows a dark-field TEM image of an exfoliated MoS2 flake; the bright

straight line across the flake is a ripplocation. [44] Such features are formed by the turbulent

forces administered to the bulk material during the sonication step of LPE. Without annealing,

NMP could remain trapped within these cavities between the MoS2 layers. Therefore, it is

suspected that the amorphous carbon detected in the exfoliated membranes after irradiation is

formed by radiation induced-amorphisation of residual NMP. Acetone was used to displace

residual NMP from the pores of both bulk and exfoliated membranes after fabrication.

However, Figure 10 shows that it is the morphology of the nanosheets produced by LPE that

affects their radiation interaction mechanisms and not the internal structure of the membranes.

15

This is supported by the absence of the D and G peaks of amorphous carbon in the Raman

spectra of bulk TMDC (L) and (H) samples.

Guo et al showed that the Raman spectrum of monolayer MoS2, produced by MME,

shows substantial changes compared with bulk MoS2 after irradiation in the electronic energy

transfer regime (1.23 GeV 209Bi ions). [23] Our work shows that the vibrational properties of

membranes made of liquid phase exfoliated nanosheets exhibit small changes compared with

bulk material. We propose that the radiation induced-amorphisation of NMP could account

for this inverse trend. Firstly, the organic solvent attenuates the energy of the incident heavy

charged particles, thus effectively reducing the dose administered to the exfoliated material.

Secondly, amorphous carbon was shown to reduce the rate of sulfur vacancy formation in

few-layer MoS2 under electron irradiation. [21] Furthermore, increasing sulfur vacancy

concentration has been reported to alter the Raman spectra of monolayer MoS2. [22]

Therefore, the amorphous carbon formed during the He2+ irradiation of exfoliated membranes

could reduce the rate of sulfur vacancy formation and account for the invariance of their

Raman spectra relative to the bulk membranes.

In order to assess the influence of alpha particle radiation on the chemical composition

of the MoS2 membranes we now consider the EDX data obtained from cross-section SEM

images of the irradiated region. Figure 11a shows the relationship between the S:Mo and

O:Mo stoichiometric ratios as a function of radiation exposure for the bulk MoS2 membranes.

The average S:Mo ratio decreases slightly after exposure to radiation. However, the bulk

MoS2 in both the (L) and (H) dose membranes is still of the correct stoichiometric ratio and

the small 2.5% reduction in sulfur is not statistically significant.

The O:Mo ratio of the bulk MoS2 (N) sample is approximately twice as large compared

with those of the irradiated samples. The cause of this large O:Mo ratio is shown in Figure

11b, where the elemental composition of bulk MoS2 membranes are expressed in atomic

percentages. The carbon content of the bulk MoS2 (N) sample is ~15 % greater than for

16

irradiated samples. Oxygen containing carbonaceous contamination is suspected to be the

cause of the large O:Mo ratio observed for the non-irradiated sample. The unknown O:C

stoichiometry of the organic contamination prevents correlation of O:Mo ratios with oxide

formation or the reduction of pre-existing oxygen moieties. The source of contamination in

the bulk MoS2 (N), (L) and (H) samples is threefold: it can be introduced during membrane

fabrication, adsorbed from the atmosphere post-fabrication and accommodated via diffusion

from the PVDF filter as supported by the small fluorine contribution.

Moving to the exfoliated MoS2 EDX data, Figure 12a shows the stoichiometric S:Mo

and O:Mo ratios for both irradiated and non-irradiated membranes. Despite the small 3.6%

reduction in the S:Mo ratio for the (L) sample, the large variance of the EDX data suggests

that alpha particle irradiation does not lead to considerable sulfur impoverishment of the

exfoliated MoS2 samples. The exfol MoS2 (L) sample exhibits the largest O:Mo ratio. Figure

12b shows the elemental composition of the exfoliated MoS2 membranes. The exfol MoS2 (L)

sample contains the largest proportion of carbonaceous material. This organic material is

suspected to be the source of the oxygen and account for the high O:Mo ratio of the sample.

Despite the carbon contamination being problematic with regards to the quantification

of radiolytic oxide formation, the 40% increase of the carbon signal in the exfoliated MoS2

samples, relative to the bulk membranes, supports the argument that a considerable amount of

NMP remains adsorbed within the exfoliated membranes, possibly within the ripplocations of

the few layer MoS2 flakes (Figure 10).

The EDX data of bulk and exfoliated WS2 membranes echoes the trends observed in the

MoS2 samples. Small decreases in S:W ratios are observed, however the distribution of the

data does not suggest any considerable thermal or radiation-induced chalcogen vacancy

formation. Organic contaminants compromise the O:W ratio analysis. Moreover, the

increased carbon content of exfoliated WS2 membranes, compared with bulk samples,

suggests that NMP is retained within the exfoliated membranes.

17

5. Conclusions

In this study, the effects of 1.66 MeV He2+ radiation on the vibrational properties and

elemental composition of bulk and exfoliated MoS2 and WS2 membranes has been explored.

We believe this work represents the first irradiation experiment investigating solution

processed TMDCs. The interaction of energetic helium nuclei with MoS2 and WS2 induce

small changes in the Raman spectra of bulk material whilst the vibrational properties of the

exfoliated material were found to be invariant. Our results are very different from those

previously reported for monolayer and bulk MoS2, obtained via MME, that showed exfoliated

flakes to be less radiation hard than the bulk material. The radiolytic production of amorphous

carbon in the liquid phase exfoliated TMDC membranes, supported by EDX and Raman

spectroscopy, is thought to contribute to the invariance of the 2LA(M), E2 g1 and A1 g

vibrational modes in this study. The effect of amorphous carbon is twofold: (1) the dose

administered to the exfoliated nanosheets is reduced. (2) the amorphous carbon reduces the

rate of radiation-induced sulfur vacancy formation.

However, further investigation must be undertaken to elucidate the extent to which this

amorphous carbon influences the principally electronic energy transfer mechanisms occurring

between the exfoliated TMDCs and the incident He2+ ions. Such an experiment would involve

transfer of the membranes onto temperature stable substrates before annealing to remove the

solvent. However, trace amounts of NMP often remain after post-processing. Therefore

consideration must be given to alternative solvents. Also, the physical mechanism responsible

for the vibrational changes in irradiated bulk WS2 and MoS2 remains to be understood;

particularly for the spectroscopic changes observed in the bulk MoS2 (H) system which

cannot be explained by previously observed defect, strain or doping effects.

6. Acknowledgments

18

LHI thanks Dr. A. Smith for operation of the ion accelerator, Dr. S. Shubeita for

modification of the sample holder to accommodate the titanium attenuator and Dr. G. Glodan

for advice regarding cross-section SEM and EDX measurements. REW acknowledges

the Hewlett-Packard Company for financial support in the framework of the Graphene

NowNano Doctoral Training Centre. The work reported here was partly funded by the Dalton

Cumbrian Facility programme, a joint initiative of the Nuclear Decommission Authority and

theUniversity of Manchester, United Kingdom. LHI thanks EPSRC for a PhD scholarship

provided through the Doctoral Training Partnership scheme (grant EP/M507969/1). AB is

supported by a research fellowship provided through the Dalton Nuclear Institute, the

University of Manchester (Dalton Fellowship).

References

[1] K.S. Novoselov, D. Jiang, F. Schedin, T.J. Booth, V.V. Khotkevich, S.V. Morozov and

A.K. Geim, Proc. Natl. Acad. Sci. U. S. A., 102 (2005) 30.

19

[2] J.N. Coleman, M. Lotya, A. O’Neill, S.D. Bergin, P.J. King, U. Khan, K. Young, A.

Gaucher, S. De, R.J. Smith, et al., Science, 331 (2011) 6017

[3] H. Li, J. Wang, Q. Luo, H. Peng, C. Luo, R. Qi, R. Huang, J. Travas-Sejdic and C. Duan,

J. Alloys Compd., 699 (2017)

[4] F. Reale, K. Sharda and C. Mattevi, Appl. Mater. Today, 3 (2016)

[5] M. Chhowalla, H.S. Shin, G. Eda, L. Li, K.P. Loh and H. Zhang, Nat. Chem., 5 (2013) 4

[6] K.F. Mak, C. Lee, J. Hone, J. Shan and T.F. Heinz, Phys. Rev. Lett., 105 (2010) 13

[7] H.R. Gutiérrez, N. Perea-López, A.L. Elías, A. Berkdemir, B. Wang, R. Lv, F. López-

Urías, V.H. Crespi, H. Terrones and M. Terrones, Nano Lett., 13 (2013) 8

[8] K.F. Mak and J. Shan, Nat. Photonics, 10 (2016) 4

[9] A.K. Geim and I.V. Grigorieva, Nature, 499 (2013) 7459

[10] H. Li, Z. Yin, Q. He, H. Li, X. Huang, G. Lu, D.W.H. Fam, A.L.Y. Tok, Q. Zhang and

H. Zhang, Small, 8 (2011) 1

[11] W. Wu, L. Wang, Y. Li, F. Zhang, L. Lin, S. Niu, D. Chenet, X. Zhang, Y. Hao, T.F.

Heinz, et al., Nature, 514 (2014) 7523

[12] N. Choudhary, M. Patel, Y. Ho, N.B. Dahotre, W. Lee, J.Y. Hwang and W. Choi, J.

Mater. Chem. A, 3 (2015) 47

[13] D. Sakar, W. Liu, X. Xie, A.C. Anselmo, S. Mitragotri and K. Banerjee, ACS Nano, 8

(2014) 4

[14] K.S. Novoselov, V.I. Fal’ko, L. Colombo, P.R. Gellert, M.G. Schwab and K. Kim,

Nature, 490 (2012) 7419

[15] Y. Chen, X. Gong and J. Gai, Adv. Sci., 3 (2016) 8

[16] D. McManus, S. Vranic, F. Withers, V. Sanchez-Romaguera, M. Macucci, H. Yang, R.

Sorrentino, K. Parvez, S. Son, G. Lannaccone, et al., Nat. Nano., 12 (2017) 4

[17] A.G. Kelly, T. Hallam, C. Backes, A. Harvey, A.S. Esmaeily, I. Godwin, J. Coelho, V.

Nicolosi, J. Lauth, A. Kulkarni et al, Science 356 (2017) 6333

20

[18] J. Lee, M.J. Krupcale and P.X.L. Feng, Appl. Phys. Lett., 108 (2016) 2

[19] S. Mignuzzi, A.J. Pollard, N. Bonini, B. Brennan, I.S. Gilmore, M.A. Pimenta, D.

Richards and D. Roy, Phys. Rev. B., 91 (2015) 19

[20] H. Komsa, J. Kotakoski, S. Kurasch, O. Lehtinen, U. Kaiser and A.V. Krasheninnikov,

Phys. Rev. Lett., 109 (2012) 3

[21] E. Rotunno, F. Fabbri, E. Cinquanta, D. Kaplan, M. Longo, L. Lazzarini, A. Molle, V.

Swaminathan and G. Salviati, 2D Mater., 3 (2016) 2

[22] W.M. Parkin, A. Balan, L. Liang, P.M. Das, M. Lamparski, C.H. Naylor, J.A.

Rodríguez-Manzo, A.T.C. Johnson, V. Meunier and M. Drndić, ACS Nano, 10 (2016) 4

[23] H. Guo, Y. Sun, P. Zhai, H. Yao, J. Zeng, S. Zhang, J. Duan, M. Hou, M. Khan and J.

Liu, Appl. Phys. A, 122 (2016) 375

[24] L. Madauß, O. Ochedowski, H. Lebius, B. Ban-d’Etat, C.H. Naylor, A.T.C. Johnson, J.

Kotakoski and M. Schleberger, 2D Mater., 4 (2016) 1

[25] S. Bertolazzi, S. Bonacchi, G. Nan, A. Pershin, D. Beljonne and P. Samorì, Adv. Mater.,

29 (2017) 18

[26] S. Tongay, J. Suh, C. Ataca, W. Fan, A. Luce, J.S. Kang, J. Liu, C. Ko, R.

Raghunathanan, J. Zhou, et al., Sci. Rep., 3 (2013)

[27] Y. Hernandez, V. Nicolosi, M. Lotya, F.M. Blighe, Z. Sun, S. De, I.T. McGovern, B.

Holland, M. Byrne, Y.K. Gun’Ko, et al., Nat. Nano., 3 (2008) 9

[28] O. Ochedowski, H. Bukowska, V.M.F. Soler, L. Brökers, B. Ban-d’Etat, H. Lebius and

M. Schleberger, Nucl. Instrum. Methods Phys. Res., Sect. B, 340 (2014)

[29] U. Halim, C.R. Zheng, Y. Chen, Z. Lin, S. Jiang, R. Cheng, Y. Huang and X. Duan, Nat.

Commun., 4 (2013) 2213

[30] E.P. Nguyen, B.J. Carey, T. Daeneke, J.Z. Ou, K. Latham, S. Zhuiykov and K. Kalantar-

Zadeh, Chem. Mater. 27 (2015) 1

21

[31] M. Lotya, Y. Hernandez, P.J. King, R.J. Smith, V. Nicolosi, L.S. Karlsson, F.M. Blighe,

S. De, Z. Wang, I.T. McGovern et al, J. Am. Chem. Soc., 131 (2009) 10

[32] A.S. Wajid, S. Das, F. Irin, H.S.T. Ahmed, J.L. Shelburne, D. Parviz, R.J. Fullerton, A.F.

Jankowski, R.C. Hedden and M.J. Green, Carbon, 50 (2012) 2

[33] A. Gupta, V. Arunachalam and S. Vasudevan, J. Phys. Chem. Lett. 7 (2016) 23

[34] J. Xie, H. Zhang, S. Li, R. Wang, X. Sun, M. Zhou, J. Zhou, X.W. Lou and Y. Xie, 25

(2013) 40

[35] U. Khan, A. O’Neill, M. Lotya, S. De and J.N. Coleman, Small, 6 (2010) 7

[36] J.F. Ziegler, M.D. Ziegler and J.P. Biersack, Nucl. Instrum. Methods Phys. Res., Sect. B,

268 (2010)

[37] D.F. Swinehart, J. Chem. Educ., 39 (1962) 7

[38] C. Lee, H. Yan, L.E. Brus, T.F. Heinz, J. Hone and S. Ryu, ACS Nano, 4 (2010) 5

[39] A. Berkdemir, H.R. Gutiérrez, A.R. Botello-Méndez, N. Perea-López, A.L. Elías, C.

Chia, B. Wang, V.H. Crespi, F. López-Urías, J. Charlier, et al., Sci. Rep. 3 (2013)

[40] H.J. Conley, B. Wang, J.I. Ziegler, R.F. Haglund, S.T. Pantelides and K.I. Bolotin, Nano

Lett., 13 (2013) 8

[41] L. Ma, Y. Tan, M. Ghorbani-Asl, R. Boettger, S. Kretschmer, S. Zhou, Z. Huang, A.V.

Krasheninnikov and F. Chen, Nanoscale, 9 (2017)

[42] A. Splendiani, L. Sun, Y. Zhang, T. Li, J. Kim, C. Chim, G. Galli and F. Wang, Nano

Lett., 10 (2010) 4

[43] C. Casiraghi, A.C. Ferrari and J. Robertson, Phys. Rev. B., 72 (2005) 8

[44] A. Kushima, X. Qian, P. Zhao, S. Zhang and J. Li, Nano Lett., 15 (2015) 2

Ref.

Material & (production

method)

Radiation source & (angle of

incidence)

Main Characterisation

TechniquesMain Results

19Monolayer

MoS2

(MME)

25 keV Mn+

(45o) Raman and DFTE' redshifts, A1

' blueshifts and the FWHM of both modes increases with decreasing defect spacing. Defect activated peaks between 140 & 420 cm-1.

20 Monolayer MoS2

80 keV electrons

DFT and atomic resolution TEM

Displacement threshold energies calculated. Formation of single & double sulfur vacancies and

22

(MME) (normal) impurity atom inclusion observed.

21Few-layer

MoS2

(MME)

200 keV electrons (normal)

EDX and CLSulfur vacancies form non-radiative centres.

Presence of amorphous carbon contamination negates sulfur impoverishment.

22 Monolayer MoS2 (CVD)

200 keV electrons (normal)

EDX, Raman and in situ electrical measurements

Correlation between sulfur vacancy percentage and magnitude of E' redshift. Reduction in

channel current with increasing fluence.

231-4 layer and

bulk MoS2

(MME)

≤ 1.23 GeV 209Bi (normal)

AFM, Raman and XPS

Hillock formation. FWHM decrease and blueshift of A1 gfor mono and few-layer flakes (bulk

unaffected). MoIV is oxidised to MoVI (from XPS).

24Monolayer

MoS2 (MME and CVD)

93 and 91 MeV Xe & 84 MeV Ta (≤ 5o)

AFM and atomic resolution TEM

Folding and nano-incission formation depends on incidence/azimuthal angles and if substrate is

suspended/supported.

25Monolayer

MoS2

(MME)

500 eV Ar+

(58o)

Raman, PL and electrical

measurements

Ar+ induces sulfur vacancies. PL decrease in intensity, E' redshift and A1

' blueshift. Spectroscopic and electrical changes are reversed

after exposure to alkanethiol vapours.

26

Monolayer MoS2,

MoSe2 & WSe2

(MME)

3 MeV He2+

(normal)

PL, DFT, Raman and nano-Auger

spectroscopy

Radiation induces single and di-chalcogen vacancies. Free and bound exciton PL signals

increase when measured in a nitrogen atmosphere due to charge transfer interactions between the N2

molecules and the TMDC defect sites.

Table I

Selected publications regarding electron and ion beam irradiation of TMDCs. The key

findings of each work are detailed alongside irradiation parameters, characterisation methods

used and the material investigated. Abbreviations not already defined are as follows: density

functional theory (DFT), full width at half maximum (FWHM), transmission electron

microscopy (TEM), cathodoluminescence (CL), Atomic Force Microscopy (AFM), x-ray

photoelectron spectroscopy (XPS) and photoluminescence (PL).

Thickness (μm)

Exfoliated Membrane (N) (L) (H)

MoS2 3.32 1.65 3.94

23

WS2 4.32 5.01 3.21

Table II

The edge thicknesses of the exfoliated TMDC membranes obtained via profilometry. The

non-irradiated samples are denoted as (N); whilst (L) and (H) represent the membranes

irradiation to low and high total absorbed doses, respectively.

24

MembraneRadiantFluence

(ions cm-2)

Dose Rate(ions cm-2 s-1)

TotalAbsorbed

Dose(MGy)

BulkMoS2 (L) 2.37 x 1015 1.31 x 1012 416

BulkMoS2 (H) 5.20 x 1015 2.88 x 1012 912

ExfolMoS2 (L) 2.19 x 1015 1.21 x 1012 383

ExfolMoS2 (H) 5.11 x 1015 2.83 x 1012 896

BulkWS2 (L) 2.14 x 1015 1.18 x 1012 253

BulkWS2 (H) 4.97 x 1015 2.76 x 1012 588

ExfolWS2 (L) 2.08 x 1015 1.15 x 1012 246

ExfolWS2 (H) 4.85 x 1015 2.69 x 1012 574

Table III

Dosimetry calculation results for TMDC membranes irradiated to low (L) and high (H) total

absorbed doses.

Figure Captions

25

Figure 1: Frequency histograms fitted with a Gaussian distribution showing the AFM

thickness measurements of 50 MoS2 flakes (Fig. 1a) and 50 WS2 (Fig. 1b) flakes. The

thickness is expressed as the number of individual TMDC monolayers comprising each flake.

Figure 2: Contour plots showing the relationship between the number of MoS2 (Fig. 2a) or

WS2 (Fig. 2b) monolayers that comprise a flake, plotted on the z axis, and its lateral

dimensions (length and width), plotted on the x and y axes. Each black spot represents one

flake measurement (thickness, length and width).

Figure 3: Cross section SEM image of the exfol MoS2 (L) sample showing that the central

irradiated region is considerably thicker than the value of 1.65 μm given by profilometry. The

scale bar represents the maximum range of a He2+ ion according to SRIM simulations and

membrane porosity.

Figure 4: The mean FWHM of the E2 g1 mode plotted against its average peak position for

bulk WS2 membranes using a single Lorentzian fit.

Figure 5: The mean FWHM of the E2 g1 (Fig. 5a) and 2LA(M) (Fig. 5b) modes plotted against

their average peak positions for bulk WS2 membranes using a multi-peak Lorentzian fit.

Representative Raman spectra of the bulk WS2 membranes showing the blueshift of the

2LA(M) mode with increasing dose (Fig. 5c).

Figure 6: The average peak positions of the A1 g (Fig. 6a) and E2 g1 (Fig. 6b) modes plotted

against their mean FWHM values for bulk MoS2 membranes. Representative Raman spectra

of the bulk MoS2 membranes showing the FWHM increase and blueshift of the A1g and E2 g1

modes in the bulk MoS2 (H) sample (Fig. 6c).

Figure 7: The average peak positions of the E2 g1 (Fig. 7a) and 2LA(M) (Fig. 7b) modes

plotted against their mean FWHM values for exfoliated WS2 membranes.

Figure 8: The mean FWHM of the A1 g mode for exfoliated MoS2 membranes plotted against

its average peak position.

26

Figure 9: Representative extended Raman spectra of the exfol MoS2 (H) and bulk MoS2 (H)

samples showing the D and G peaks of amorphous carbon in the exfoliated membrane.

Figure 10: Dark-field TEM micrograph of an exfoliated MoS2 flake showing the morphology

of the solution processed material. The bight straight line on the bottom left corner of the

flake is a ripplocation. Electron diffraction patterns show the characteristic hexagonal motif of

crystalline MoS2 circled in red (inset).

Figure 11: Scatter plot showing the relationship between the mean S:Mo and O:Mo

stoichiometric ratios for bulk MoS2 membranes (Fig. 11a). The average chemical composition

of the bulk MoS2 membranes expressed in terms of atomic percentage (Fig. 11b).

Figure 12: Scatter plot showing the relationship between the mean S:Mo and O:Mo

stoichiometric ratios for exfoliated MoS2 membranes (Fig. 12a). The average chemical

composition of the exfoliated MoS2 membranes expressed in terms of atomic percentage (Fig.

12b).

27

List of Figures

Figure 1a

Figure 1b

28

Figure 2a

29

Figure 2b

Figure 3

30

Figure 4

354 354.5 355 355.5 356 356.5 357 357.5 3582

2.5

3

3.5

4

4.5

5

5.5

6

6.5

7

(N) (L)

(H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

354 354.5 355 355.5 356 356.5 357 357.5 3582

2.5

3

3.5

4

4.5

5

5.5

6

6.5

7

(N) (L)

(H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

Figure 5a

354 354.5 355 355.5 356 356.5 357 357.5 3582

3

4

5

6

(N) (L)

(H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

31

354 354.5 355 355.5 356 356.5 357 357.5 3582

3

4

5

6

(N) (L)

(H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

Figure 5b

346 348 350 352 354 3567

7.5

8

8.5

9

9.5

10

10.5

11

11.5

12

(N) (L)

(H)

2LA(M) Position (cm-1)

2LA(

M) F

WHM

(cm

-1)

346 348 350 352 354 3567

7.5

8

8.5

9

9.5

10

10.5

11

11.5

12

(N) (L)

(H)

2LA(M) Position (cm-1)

2LA(

M) F

WHM

(cm

-1)

32

Figure 5c

Figure 6a

407.8 408 408.2 408.4 408.6 408.8 409 409.2 409.43.5

3.7

3.9

4.1

4.3

4.5

4.7

4.9

5.1

5.3

(N) (L)

(H)

A1g Position (cm-1)

A1g F

WHM

(cm

-1)

33

407.8 408 408.2 408.4 408.6 408.8 409 409.2 409.43.5

3.7

3.9

4.1

4.3

4.5

4.7

4.9

5.1

5.3

(N) (L)

(H)

A1g Position (cm-1)

A1g F

WHM

(cm

-1)

Figure 6b

382.4 382.6 382.8 383 383.2 383.4 383.6 383.8 3843

3.5

4

4.5

5

5.5

6

(N) (L)

(H)

E2g Position (cm-1)

E2g F

WHM

(cm-

1)

382.4 382.6 382.8 383 383.2 383.4 383.6 383.8 3843

3.5

4

4.5

5

5.5

6

(N) (L)

(H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

34

Figure 6c

Figure 7a

355.5 355.7 355.9 356.1 356.3 356.5 356.7 356.95

5.5

6

6.5

7

7.5

8

(N) (L) (H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

35

355.5 355.7 355.9 356.1 356.3 356.5 356.7 356.95

5.5

6

6.5

7

7.5

8

(N) (L) (H)

E2g Position (cm-1)

E2g F

WHM

(cm

-1)

Figure 7b

347 348 349 350 351 35210

11

12

13

14

(N) (L)

(H)

2LA(M) Position (cm-1)

2LA(

M) FW

HM (c

m-1)

347 348 349 350 351 35210

11

12

13

14

(N) (L)

(H)

2LA(M) Position (cm-1)

2LA(

M) FW

HM (c

m-1)

36

Figure 8

407.9 408 408.1 408.2 408.3 408.4 408.53.5

4

4.5

5

5.5

(N) (L)

(H)

A1g Position (cm-1)

A1g F

WHM

(cm

-1)

407.9 408 408.1 408.2 408.3 408.4 408.53.5

4

4.5

5

5.5

(N) (L)

(H)

A1g Position (cm-1)

A1g F

WHM

(cm

-1)

Figure 9

37

Figure 10

Figure 11a

0 0.1 0.2 0.3 0.42.05

2.1

2.15

2.2

(N) (L)

(H)

O:Mo Ratio

S:Mo R

atio

38

0 0.1 0.2 0.3 0.42.05

2.1

2.15

2.2

(N) (L)

(H)

O:Mo Ratio

S:Mo R

atio

Figure 11b

(N) (L) (H)0%

10%

20%

30%

40%

50%

60%

70%

80%

90%

100%

MoSFOC

Bulk MoS2 Membrane

Element

al Com

positio

n (Atom

ic %)

39

(N) (L) (H)0%

10%

20%

30%

40%

50%

60%

70%

80%

90%

100%

MoSFOC

Bulk MoS2 Membrane

Element

al Com

positio

n (Atom

ic %)

Figure 12a

0.2 0.3 0.4 0.5 0.6 0.7 0.81.8

1.9

2

2.1

2.2

2.3

(N) (L)

(H)

O:Mo Ratio

S:Mo R

atio

0.2 0.3 0.4 0.5 0.6 0.7 0.81.8

1.9

2

2.1

2.2

2.3

(N) (L)

(H)

O:Mo Ratio

S:Mo R

atio

40

Figure 12b

(N) (L) (H)0%

10%

20%

30%

40%

50%

60%

70%

80%

90%

100%

MoSFOC

Exfol MoS2 Membrane

Eleme

ntal Co

mpositi

on (At

omic %

)

(N) (L) (H)0%

10%

20%

30%

40%

50%

60%

70%

80%

90%

100%

MoSFOC

Exfol MoS2 Membrane

Eleme

ntal Co

mpositi

on (At

omic %

)

41