ARTICLES
PUBLISHED ONLINE: 5 MAY 2013 | DOI: 10.1038/NMAT3633
Grains and grain boundaries in highly crystalline
monolayer molybdenum disulphide
Arend M. van der Zande1,2 *† , Pinshane Y. Huang3† , Daniel A. Chenet2† , Timothy C. Berkelbach4 ,
YuMeng You5 , Gwan-Hyoung Lee2,6 , Tony F. Heinz1,5 , David R. Reichman1,4 , David A. Muller3,7
and James C. Hone1,2
Recent progress in large-area synthesis of monolayer molybdenum disulphide, a new two-dimensional direct-bandgap
semiconductor, is paving the way for applications in atomically thin electronics. Little is known, however, about the
microstructure of this material. Here we have refined chemical vapour deposition synthesis to grow highly crystalline islands
of monolayer molybdenum disulphide up to 120 µm in size with optical and electrical properties comparable or superior to
exfoliated samples. Using transmission electron microscopy, we correlate lattice orientation, edge morphology and crystallinity
with island shape to demonstrate that triangular islands are single crystals. The crystals merge to form faceted tilt and mirror
twin boundaries that are stitched together by lines of 8- and 4-membered rings. Density functional theory reveals localized
mid-gap states arising from these 8–4 defects. We find that mirror twin boundaries cause strong photoluminescence quenching
whereas tilt boundaries cause strong enhancement. Meanwhile, mirror twin boundaries slightly increase the measured in-plane
electrical conductivity, whereas tilt boundaries slightly decrease the conductivity.
C
haracterizing the structure and properties of grains and
grain boundaries is critical for understanding and controlling material properties in the expanding array of twodimensional (2D) materials1 . Studies have mapped the grain
structure of large-area graphene grown by chemical vapour
deposition (CVD), characterized defects such as grain boundaries
at the atomic scale and demonstrated that these grain boundaries
can strongly affect graphene’s electrical, optical and mechanical
properties2–11 . Much less is known about the grain structure and
properties of defects in other 2D materials, such as molybdenum
disulphide (MoS2 ). MoS2 has been widely explored in the form
of nanotubes12 , nanoparticles13 and thin films14,15 owing to its
excellent tribological and catalytic properties. Recent work has
shown that exfoliated monolayer MoS2 is a 2D direct-bandgap
semiconductor16,17 . Moreover, large-area monolayer MoS2 can now
be synthesized using CVD (refs 18–21), making it a promising candidate for building atomically thin layered electrical22–24 , optical16,25
and photovoltaic26 devices.
Here, we grow high-quality crystals of monolayer MoS2
exhibiting grain sizes up to 120 µm and possessing optical and
electronic properties comparable or superior to those of exfoliated
samples. Our study applies a diverse set of characterization
methods, combining transmission electron microscopy (TEM)
of atomic and crystal structure with optical spectroscopy and
electrical transport to determine the influence of grain boundaries
on the properties of the material. Using diffraction-filtered and
atomic-resolution electron microscopy, we characterize singlecrystal triangular islands and polycrystals containing tilt and mirror
twin grain boundaries. We find mirror twin boundaries are stitched
together predominantly through lines of 8- and 4-membered
rings, which produce mid-gap states in DFT calculations. Finally,
we show that individual grain boundaries strongly affect the
photoluminescence observed in MoS2 monolayers and slightly alter
their in-plane electrical conductivity.
Growth and characterization of large-grain MoS2 crystals
Figure 1a,b shows optical images of monolayer MoS2 grown on
a Si/SiO2 substrate. The samples are grown by CVD with solid
MoO3 and S precursors18 . In contrast to previous work18 , we did
not use seeds to nucleate growth; instead, our best growths were
obtained with carefully cleaned substrates and by minimizing the
exposure of the precursors to air during storage (see Methods
and Supplementary Fig. S1 for details). Figure 1a shows an optical
image of a typical sample. Isolated islands (violet triangles) have
edge lengths ranging from 30 to 80 µm. On the right (nearest to
the solid-state precursors during growth), the islands merge into
a continuous film. Although non-uniformity in nucleation density
and crystal size is a limitation of our growth techniques, each
growth contains a ∼1 × 15 mm region where hundreds of similar,
isolated islands grow. Across different growths, the average size
of islands varies between 20 and 100 µm, with individual triangles
up to 120 µm (Supplementary Fig. S2). Most islands are uniform
monolayers, except for small bilayer or multilayer patches (visible
in the centre of Fig. 1b).
We first use photoluminescence to measure the quality and
thickness of these triangles. Figure 1c shows photoluminescence
spectra from monolayer and bilayer MoS2 . The photoluminescence
peaks at 1.84 and 1.95 eV respectively match the A and B directgap optical transitions16 . The integrated intensity of the bilayer
peak is much weaker (∼7%) than the monolayer peak, reflecting
1 Energy
Frontier Research Center, Columbia University, New York, New York 10027, USA, 2 Department of Mechanical Engineering, Columbia University,
New York, New York 10027, USA, 3 School of Applied and Engineering Physics, Cornell University, Ithaca, New York 14853, USA, 4 Department of
Chemistry, Columbia University, New York, New York 10027, USA, 5 Departments of Physics and Electrical Engineering, Columbia University, New York,
New York 10027, USA, 6 Samsung-SKKU Graphene Center (SSGC), Suwon, 440-746, Korea, 7 Kavli Institute at Cornell for Nanoscale Science, Ithaca,
New York 14853, USA. † These authors contributed equally to this work. *e-mail: av2466@columbia.edu.
554
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b
a
100 µm
123 µm
c
e
d
Monolayer
2.5
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Bilayer
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I / ISi
2.0
1.5
1.0
0.5
0.0
Mo
1.4
1.6 1.8 2.0
Energy (eV)
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S
Figure 1 | Large-grain MoS2 growth. a, Optical reflection image of a CVD growth of a typical large-grain MoS2 on a SiO2 (285 nm)/Si substrate. The image
contrast has been increased for visibility; magenta is the bare substrate, and violet represents monolayer MoS2 . b, Optical image of a monolayer MoS2
triangle. The triangle is 123 µm from tip to tip. c, Photoluminescence spectra from monolayer (red) and bilayer (blue) MoS2 . Peak height is normalized to
the silicon Raman peak. The narrow spikes at high energy are the Raman transitions (see Supplementary Fig. S3a). d, High-resolution ADF-STEM image of
freely suspended monolayer MoS2 on a TEM grid. The bright spots are molybdenum atoms; the grey spots are two stacked sulphur atoms. The lattice is
composed of hexagonal rings alternating molybdenum and sulphur sites; top view and side views of the structure are overlaid. e, DF-TEM image of a large
triangle with the diffraction pattern inset. Together, the diffraction pattern and the DF-TEM image show that the triangle is a continuous single crystal. The
∼2–4 µm brighter and darker areas are rotationally aligned bilayers of MoS2 . Similar variations in contrast have been observed in bilayer graphene, where
they reflect differences in stacking order29 .
the expected evolution from a direct-gap to an indirect-gap
semiconductor. Surprisingly, the measured range of peak widths of
our CVD monolayer MoS2 (50–60 meV) is similar to that observed
for freely suspended samples of exfoliated MoS2 (also 50–60 meV)
and is much narrower than that of MoS2 exfoliated onto SiO2
(100–150 meV; ref. 16). These results suggest that the CVD-grown
samples are either of higher quality or are in a cleaner electrostatic
environment than exfoliated samples.
To characterize the crystal structure of these islands, we used a
combination of TEM, electron diffraction techniques and atomic
resolution STEM. Figure 1d shows an image of the atomic structure
of our CVD MoS2 monolayers taken by aberration-corrected
annular dark-field scanning transmission electron microscopy
(ADF-STEM). In ADF-STEM, a 60 keV, ångstrom-scale electron
beam scans over the sample, and an image is formed by collecting
the medium- to high-angle scattered electrons. The contrast of
ADF-STEM images scales roughly as the square of the atomic
number Z (ref. 27). Thus, the brightest spots in Fig. 1d are
the molybdenum atoms and the dimmer spots are the two
stacked sulphur atoms. The hexagonal lattice is clearly visible, as
indicated by the top- and side-view schematics in Fig. 1d. No point
defects, atomic substitutions or voids were initially observed within
single crystals. However, defects readily form under extended
imaging (see Methods)28 .
We used selected-area electron diffraction and dark-field
TEM (DF-TEM) to characterize the large-scale crystal structure
of the islands. Figure 1e, inset shows a selected-area electron
diffraction pattern of a triangular island roughly 45 µm across.
The six-fold symmetry in the position of the diffraction spots
demonstrates that the triangle contains no rotational boundaries,
to within our measurement accuracy of 0.5◦ . Figure 1e shows the
corresponding DF-TEM image: by using an aperture to select
a narrow range of diffracted beams, DF-TEM filters images
by crystallographic orientation2,11,29,30 , and shows single crystals
such as Fig. 1e in a single image. Similar analyses of dozens of
triangles show that triangular-shaped islands are predominantly
single crystals. Under continued growth, these islands merge
together to form aggregates or continuous sheets (Fig. 1a and
Supplementary Fig. S2).
The large-grain, highly crystalline MoS2 monolayers, with
optical properties equivalent or superior to exfoliated samples,
enable new experiments studying the growth, structure and
properties of MoS2 crystals and grain boundaries.
Grain structure and orientation of CVD MoS2
We use electron diffraction techniques to uniquely identify the
crystal structure and edge orientation of MoS2 triangles. Figure 2a,b
shows the bright-field image and diffraction pattern of a singlecrystal triangle. As seen in the schematic in Fig. 2a, the lattice
of monolayer MoS2 is divided into molybdenum and sulphur
sublattices, which reduces the hexagonal lattice from six-fold to
three-fold symmetry. As a result, the six [1̄100] diffraction spots
are broken into two families: ka = {(1̄100), (101̄0), (01̄10)} and
kb = −ka (Fig. 2b; ref. 31). Through both experiment and Bloch
wave simulations32 (see Supplementary Methods), we find the ka
spots are higher in intensity (∼10% for an 80 kV electron beam)
and point towards the molybdenum sublattice, as indicated by
the arrows in Fig. 2a,b.
We likewise observe this intensity asymmetry in DF-TEM.
Figure 2d shows the bright-field image and diffraction pattern of
two triangles. An aperture positioned at the magenta circle on the
inset in Fig. 2d forms the dark-field image in Fig. 2e. As the triangles
are rotated 180◦ from one another, the aperture simultaneously
collects the bright ka spot for the left triangle and the darker
kb spot for the right triangle. This produces a corresponding
contrast difference between the islands. We can thus use either
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NATURE MATERIALS DOI: 10.1038/NMAT3633
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a
2 µm
c
b
Experiment
ka
0.0
d
1 µm
[¬
2110]
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[1¬100]
kb
[¬
1100]
ka
¬
[2110]
Relative intensity
1.0
Simulation
kb
Position
e
kb
1 µm
ka
Figure 2 | Diffraction imaging of crystal orientation and edge terminations. a, Bright-field image of a single-crystal triangle with a Mo-zigzag edge
orientation. b, Diffraction pattern from a. The asymmetry of the Mo and S sublattices separates the [1̄100] diffraction spots into two families:
ka = {(1̄100),(101̄0),(01̄10)} and kb = −ka . c, A line profile through experimentally measured diffraction spots (black) and Bloch-wave simulations (red).
The higher intensity ka spots point towards the Mo sublattice, as indicated by the arrows in a,b. Both curves were normalized to the height of the [2̄110]
peaks, and the red curve was offset horizontally for visibility. d, Bright-field TEM image of two triangles with S-zigzag edge orientations. The curved
appearance of the crystal edges contrasts with the sharper crystal edges of the Mo-zigzag edges in a. The inset diffraction pattern shows the location of
the aperture used to form the image in e. e, Dark-field image of the region in d. As the triangles are rotated 180◦ from one another, the aperture
simultaneously collects the brighter ka spot for the left triangle and the darker kb spot for the right triangle. This produces a corresponding contrast
difference between the two islands that allows us to uniquely infer crystallographic orientation from dark-field images. To improve visibility, we have
clipped the bottom of the intensity range in this image.
the diffraction pattern or DF-TEM to uniquely identify the lattice
orientation of MoS2 crystals.
Figure 2 also shows the dominant morphologies of MoS2
triangles: molybdenum zigzag (Mo-zz) in Fig. 2a and sulphur
zigzag (S-zz) in Fig. 2d (refs 33,34). These two edge terminations
are commonly observed in MoS2 nanocrystals33 , and they are the
two most energetically stable edge orientations34 . We consistently
observe that Mo-zz triangles have sharper, straighter edges than S-zz
triangles. This morphological difference allows us to easily identify
crystal orientation in optical images. High-resolution STEM images
show that despite their sharp appearance in DF-TEM, the edges
of both our S-zz and Mo-zz crystals are rough at the nanoscale
(Supplementary Fig. S4).
In addition to single-crystal triangles, the CVD process produces
polycrystalline aggregates that can be characterized by DF-TEM.
Figure 3a shows two S-zz triangles intersecting at an angle of
40◦ ± 0.5◦ . Figure 3b shows a corresponding composite DF-TEM
image of the crystal structure. The composite image is constructed
by overlaying colour-coded DF-TEM images2 acquired from the
crystallographic orientations shown in the diffraction pattern of
Fig. 3a. The intersection of the grains forms an abrupt, faceted tilt
boundary. Low-quality growths can produce the types of film in
Fig. 3c, which shows small, misshapen, crystals. The corresponding
diffraction pattern and dark-field overlay in Fig. 3d shows that
the disordered crystals are randomly oriented polycrystalline
aggregates. As indicated by the arrows in Fig. 3c, adjacent MoS2
grains sometimes overlap to form rotationally misaligned bilayer
regions. Similar combinations of abrupt and overlapped crystal
interfaces have been seen in CVD-grown graphene2,5,35 . Unlike
in graphene, the grain boundaries in MoS2 tend to be strongly
faceted on the micrometre scale. These facets may originate from
the relative surface energy of different edge configurations36 ,
which will be more directionally dependent in MoS2 than in
556
homoelemental structures. Similar phenomena probably underlie
the differences in edge roughness we see in Mo-zz versus S-zz
triangles and the faceting at the edges of large crystals such
as in Fig. 1b.
In addition to tilt boundaries, we observed a second common
line defect in MoS2 crystals: mirror twin boundaries. In monolayer
MoS2 , mirror twins are the intersections of two MoS2 crystals
with a relative in-plane rotation of 180◦ that effectively swaps
the positions of Mo and S lattice sites. Figure 3e shows the
bright-field image of two triangles meeting at 180◦ , and Fig. 3f
shows the corresponding DF-TEM image. The mirror twin appears
clearly as an intensity change in DF-TEM images, through the
same diffraction asymmetry seen in Fig. 2e. Unlike tilt boundaries,
mirror twin boundaries such as in Fig. 3e,f grow within nucleation
islands. Figure 3g,h and Supplementary Fig. S5 show another
commonly observed shape that contains mirror twin boundaries:
a symmetric, six-pointed star containing cyclic twins that resemble
other inorganic and geological crystal growths37 .
The atomic structure of MoS2 grain boundaries
We next examine the atomic structure of MoS2 grain boundaries
using ADF-STEM. Figure 4a,b show a mirror twin boundary in
a 6-pointed star. Whereas the grain boundary follows the zigzag
direction on the micrometre scale, it exhibits faceting of ±20◦
relative to this direction on the nanometre scale. This nanoscale
faceting is surprising because it avoids the highly symmetric
zigzag boundary structure that would correspond to the microscale
alignment seen in Fig. 3h. Figure 4b shows the atomic structure of
the grain boundary. The overlaid purple and green polygons show
that the boundary is formed primarily from 4- and 8-membered
rings, with a recurring periodic 8-4-4 ring motif.
STEM images in an ArXiv preprint posted during the review
process of this work show a MoS2 tilt boundary formed from
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a
b
10 µm
c
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5 µm
d
e
f
1 µm
g
5 µm
h
Figure 3 | Tilt and mirror twin grain structures. a, Bright-field TEM image of two triangles that have grown together. Inset diffraction pattern shows the two
crystal orientations are 40◦ ± 0.5◦ apart (measured between the red and cyan lines, which are at equivalent spots in the diffraction pattern).
b, Colour-coded overlay of DF-TEM images corresponding with the two red- and cyan-circled spots in a shows a tilt grain boundary as a faceted line
connecting the two triangles. c, Bright-field image of a region containing irregularly shaped MoS2 islands, with diffraction pattern inset. Red arrows indicate
regions where adjacent grains overlap, forming rotationally misaligned bilayers. d, Coloured DF-TEM overlay shows that the irregular shapes are
polycrystalline aggregates. The crystals are connected both by faceted, abrupt grain boundaries, and ∼1 µm overlapped bilayers. e, Bright-field image of a
mirror twin composed of 180◦ -rotated triangles, with diffraction pattern inset. f, Dark-field image corresponding with the orange circle in e shows the two
triangles have different diffraction intensity, similar to the 180◦ -rotated triangles in Fig. 2d,e. This intensity change indicates the presence of a mirror twin
boundary between the darker and lighter regions. The small triangle in the centre is multilayer MoS2 . g, Bright-field image of a 6-pointed star, with inset
diffraction pattern. h, DF-TEM image corresponding with the orange circle in g shows that the star contains several rotationally symmetric mirror twins,
forming a cyclic twin.
a line of 5–7 rings38 . This finding complements our atomicresolution images of mirror twin structures: both tilt and mirror
twin boundaries are commonly found in our films. These results
are consistent with a recent theory paper that predicts that in MoS2 ,
either odd- or even-membered rings can form grain boundaries
depending on the tilt angle and stoichiometry, among several
factors39 . For comparison, graphene’s tilt boundaries are most
commonly formed by 5- and 7-membered rings2,40 , whereas twin
boundaries have been observed with 8-5-5 motifs4 . Our data also
suggest that a priori simulation of boundary structures may be
challenging: for example, our 8-4-4 motif does not match the
theoretically predicted structure in ref. 39, mostly owing to its
surprising faceted structure.
In our observed 8-4-4 motif, neighbouring 4-rings meet at a
sulphur site, which appears to have four nearest neighbours rather
than the usual three. This change in coordination requires either the
S atoms to have dangling bonds, or the Mo atoms to change their
oxidation state. Using first-principles density functional theory41
(DFT), we find that after relaxation, the 8-4-4 boundary structure
is at a local energetic minimum (Fig. 4c).
Figure 4d shows the DFT-calculated density of states (DOS) of
defect-free monolayer MoS2 (black curve). The light green band
indicates the bandgap of pristine MoS2 . When a grain boundary
is implanted into the MoS2 , the DOS develops mid-gap states
(pristine + GB, red dashed curve) that appear mainly in the
projected DOS of the atoms in the boundary (GB alone, filled
blue curve). Figure 4e shows a spatial plot of the local DOS of
pristine + GB in the plane of the Mo atoms, integrated in the
energy range of the pristine bandgap; the mid-gap states are spatially
localized at the boundary. Analogous calculations on alternative
mirror twin boundary structures (Supplementary Fig. S6) confirm
that this behaviour is generic. These localized mid-gap states, typical
of defects in semiconductors, are important because they can affect
the optical and transport properties of the material4,42 .
Optical properties of MoS2 grain boundaries
Using the understanding of the grain structure developed above, we
can systematically explore the effects of individual grain boundaries
on the optical and electronic properties of MoS2 .
We measure the optical properties of MoS2 using
photoluminescence mapping, where we step a focused excitation
laser over the sample and record photoluminescence spectra at
each point. Figure 5a,e shows optical images of a mirror twin island
(Fig. 5a) and a tilt boundary island (Fig. 5e). Figure 5b–d,f–h shows
corresponding maps of integrated photoluminescence intensity,
peak position and peak width. Away from the boundaries, both
islands show similar emission intensities, peak positions (1.84 eV)
and peak widths (57 meV in the mirror sample and 60 meV in
the tilt sample, reflecting typical sample-to-sample variation). At
the grain boundaries, the islands show strong photoluminescence
modifications: the mirror boundary is a clearly visible straight line
with 50% quenching of intensity (Fig. 5b), an 8 meV upshift in peak
energy (Fig. 5c) and a 5 meV peak-broadening (Fig. 5d). The tilt
boundary, which has a faceted structure similar to Fig. 3b, shows a
surprising 100% enhancement in emission strength, 26 meV upshift
and 5 meV broadening. The width of the signal change around
the tilt boundary is also larger than around mirror boundaries.
Supplementary Fig. S3b,c shows corresponding Raman maps
of the tilt boundary.
The data demonstrate that photoluminescence is strongly
affected by the presence of grain boundaries and that these
changes depend on the grain boundary angle. Several factors may
contribute. In the simplest theory, photoluminescence quenching
arises from defects in semiconductors, such as mid-gap states
at the boundaries, which can act as centres for non-radiative
recombination43 . Our calculations indicate that these effects,
even after accounting for exciton diffusion, are too small to
explain the photoluminescence results (Supplementary Methods).
Two further aspects of the grain boundaries probably play a
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NATURE MATERIALS DOI: 10.1038/NMAT3633
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a
20°
zig-zag
20°
5 nm
b
c
1 nm
d
DOS (eV¬1/atom)
2.5
e
Pristine
Pristine + GB
2.0
GB alone
1.5
1.0
0.5
0.0
¬2
¬1
0
Energy (eV)
1
2
Local midgap DOS
Figure 4 | Grain boundary atomic structure. a, High-resolution ADF-STEM image of a mirror twin boundary. The boundary is visible just below the
annotated line. The annotation indicates the nanoscale faceting of the boundary at ±20◦ off of the zigzag direction. b, Zoomed-in image of the grain
boundary shows a periodic line of 8-4-4 ring defects. c, An atomistic model of the experimental structure shown in b. Energy minimization with DFT
confirms that this boundary is locally stable. d, The total DOS of defect-free, that is pristine, MoS2 (black), the total DOS of MoS2 with the grain boundary
(red dashed) and the projected DOS of the atoms along the grain boundary (blue filled). The dashed grey line denotes the Fermi energy of pristine MoS2 ,
and the light green shaded area indicates the pristine bandgap. All states have been given a Gaussian broadening of 0.07 eV. e, A 2D spatial plot of the
local mid-gap DOS (integrated in the plane of the Mo over a 1.7 eV range about the Fermi energy of the pristine MoS2 ). The colour scale of the density is
0–0.05 states per bohr3 .
role in controlling photoluminescence: doping and strain. First,
local changes in doping may occur at grain boundaries. For
example, our observed 8-4-4 defects in the mirror boundary
are molybdenum rich, which would n-dope the boundary,
whereas 5–7 defects observed at a tilt boundary are sulphurrich38 , which would p-dope the boundary. Previous investigations
have shown that photoluminescence is strongly affected by
charge density, with increased/decreased electron density causing
photoluminescence quenching/enhancement44 . The influence of
charge is thus consistent with the observed trends in the strength
of the photoluminescence at these different types of boundary.
Other sources of doping can include spatial differences in growth
kinetics or the preferential attachment of adsorbates at boundaries.
Second, strain around the boundaries may modify the bandgap45 or
cause the boundary region to lift off the electrically disordered SiO2
surface, enhancing photoluminescence emission16 .
Electrical properties of MoS2 grain boundaries
An important question for device applications is whether grain
boundaries disrupt or modify electrical transport. We fabricated
field-effect transistors (FETs) from islands containing a single grain
boundary (Fig. 5i, inset). The FET channels followed three different
configurations: within a grain (pristine), across the grain boundary
(perpendicular) and along the grain boundary (parallel). Figure 5i,j
shows plots of conductance versus gate voltage (transfer curves) for
four representative devices fabricated on mirror twin islands: two
pristine devices (black and magenta), one parallel device (orange
curve) and one perpendicular device (cyan curve).
558
First, we compare the variability of pristine, single-crystal
devices. In devices from different growths, room-temperature,
field-effect mobilities ranged from 1 to 8 cm2 V−1 s−1 , and typical
on/off ratios ranged between 105 and 107 . Multiple devices
fabricated within each single crystal yielded mobilities that were
identical to within our measurement error. Specifically, the pristine
devices in Fig. 5i,j (black and magenta curves) show n-type
behaviour with mobilities of 3–4 cm2 V−1 s−1 . These ranges of values
are comparable to those reported for back-gated FETs fabricated
with mechanically exfoliated MoS2 on SiO2 (refs 22,24,46,47;
in the absence of high-K dielectrics) and equivalent to the best
reported values for monolayer CVD MoS2 (refs 18,19). Our present
mobilities are probably limited by substrate effects and the local
electrostatic environment48,49 .
Next we compare transport across grain boundaries. In the
mirror twin device shown in Fig. 5i,j, the perpendicular device (cyan
curve) shows nearly identical performance to the pristine devices
(magenta and black curves), indicating that the mirror twin boundary has little effect on channel conductivity. However, the electrical
characteristics of the parallel devices differed measurably from the
pristine devices. Although the transfer curves of the parallel device
in Fig. 5i,j (orange curves) are qualitatively similar to that of the
pristine devices, they have a 25% larger on-state conductivity and
60% larger off-state conductivity. This change is only slightly larger
than the device-to-device variation, but was consistent across four
parallel devices measured on mirror twins (particularly in the off
state). The data impose a rough upper limit for conductivity along
the one-dimensional grain boundary of 1.5 µS-µm in the on-state
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Optical image
a
b
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Photoluminescence
quantum yield
Photoluminescence
peak position
c
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peak width
d
Mirror boundary
10 µm
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f
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h
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100
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35
70
10¬5
¬70
Parallel to GB
Perpendicular to GB
Pristine
Pristine
¬35
0
Vg (V)
35
70
Figure 5 | Optical and electronic properties of mirror and tilt boundaries. a–d, Optical measurements of an island containing a mirror twin boundary.
e–h, Corresponding measurements for an island containing a tilt boundary. a,e are optical images; b–d and f–h are colour plots of photoluminescence. In
b,f, red is the relative quantum yield, with colour scale 0–1100 a.u. We see 50% quenching at the mirror twin boundary and a 100% enhancement at the tilt
boundary. In c,g, green is the peak position, with colour scale 1.82–1.87 eV. There is an upshift of 8 meV at the mirror twin boundary, and a much stronger
26 meV upshift in the tilt boundary. In d,h, cyan is the peak width with colour scale of 53–65 meV. The peak broadens from 55 to 62 meV at the boundary
in both samples. i,j, Linear and logarithmic electrical transport transfer curves of 4 FETs fabricated from the mirror twin MoS2 island shown in the inset of i.
The curves correspond with pristine regions (magenta and black), and regions containing a grain boundary running perpendicular (cyan) and parallel
(orange) to the flow of electrons. All data were measured at room temperature, using the Si growth substrate as a back-gate and a source–drain bias of
500 mV.
(Fig. 5j at Vg = +70 V) and 120 pS-µm in the off state (Fig. 5j at
Vg = −70 V). These values can be compared to overall sheet conductivity for pristine devices of 2.2 µS/ in the on state and 56 pS/
in the off state. This analysis indicates that the few-atom-wide twin
boundaries, although still semiconducting, have similar conductivity of up to a 1-µm-wide strip of pristine material. This result is consistent with the mid-gap states we predict in the material. The full
effect of mid-gap states, however, is probably limited by the disorder
in the grain boundary, which will prevent a continuum of states.
In contrast to our measurements of the mirror twin, FET
devices of tilt boundaries show a decrease in conductance in
both the parallel and perpendicular configurations. On one tilt
boundary, shown in Supplementary Fig. S8, the on-state boundary
conductance of both parallel and perpendicular devices decreases by
30% compared with pristine crystals. However, measuring six angle
devices yielded no consistent change in conductance, decreasing
by 5–80% compared with pristine. This variability is unsurprising:
in graphene, electrical measurements are highly dependent on
the structure of the boundary5 . We expect these variations to be
wider and more complex in semiconducting MoS2 , where the
electronic structure of the boundary should depend on the tilt
angle and atomic structure39 . Significantly, in all mirror twin and
tilt boundaries we measured, the change in conductance from any
single grain boundary is only slightly more than the device-todevice variation—an important conclusion for the applicability of
MoS2 -based electronics.
In this work, we produced monolayers of large-area, highly
crystalline MoS2 on insulating surfaces and imaged the crystal edges,
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NATURE MATERIALS DOI: 10.1038/NMAT3633
ARTICLES
grain structure and grain boundaries with electron microscopy.
We identified the location of individual grain boundaries on the
surface and systematically compared the optical and electronic
properties to electronic structure calculations. In particular, we
used the repeatability of the mirror twin structure to correlate
measurements of the structure, photoluminescence and electronic
transport of a single grain boundary structure. In tilt boundaries,
photoluminescence and electrical transport measurements vary
markedly for different boundaries. Controlling grain boundaries
will be essential for optimizing MoS2 electronic and photovoltaic
devices, where defects in large-area heterostructures can lead to
interlayer leakage or unwanted exciton recombination.
MoS2 is just the first member of a large family of layered
transition metal dichalcogenides to be synthesized as a monolayer50 .
The above techniques are crucial for exploring synthesis strategies
to optimize the grain properties, and provide a template for studies
of the microscopic and macroscopic impact of grain structure on
other monolayer membranes. Our results are thus a significant step
forward in realizing the ultimate promise of atomic membranes in
electronic, mechanical and energy-harvesting devices.
Methods
Growth procedure. We grew monolayer MoS2 by CVD using a method similar
to ref. 18, but without seeds to nucleate growth. Growth substrates were Si with
285 nm of SiO2 . Substrates were cleaned in acetone and isopropanol, followed
by 2 h in H2 SO4 /H2 O2 (3:1) and 5 min of O2 plasma. They were then loaded
into a 2-inch CVD furnace and placed face-down above a crucible containing
14 mg of MoO3 (≥99.5% Sigma Aldrich #100932642) with another crucible
containing 120 mg of sulphur (≥99.5% Sigma Aldrich #101144903) located
upstream. The CVD growth is performed at atmospheric pressure while flowing
ultrahigh-purity nitrogen. The growth recipe is: sit 4 h at 105 ◦ C with 500 s.c.c.m.,
ramp to 700 ◦ C at 15 ◦ C min−1 with 10 s.c.c.m., sit 5 min at 700 ◦ C, cool to
570 ◦ C without feedback (∼20 min) with 10 s.c.c.m., open furnace and flow 500
s.c.c.m. for rapid cooling.
High yield is achieved by using ultraclean substrates and with fresh precursors.
The yield is significantly lowered by dirty substrates and old precursors. In these
cases, any monolayers will be smaller and misshapen, as outlined in Supplementary
Fig. S1. Closest to the source, the film becomes a mixture of multilayer stacks
and other crystal allotropes, whereas furthest from the source, nothing grows (see
Fig. 1a and Supplementary Fig. S2).
TEM sample preparation. For the TEM grids, we use 10-nm-thick amorphous
SiN windows (TEMwindows.com) for DF-TEM samples and holey Quantifoil
grids (Ted Pella) for ADF-STEM. Polymer transfer layers were prepared by
spinning poly(methyl methacrylate) (PMMA) A2 onto finished MoS2 on silicon
oxide/silicon chips at 4000 r.p.m. for 60 s to form 50-nm-thick layers. The chips
were floated on 1 M KOH. The KOH etched the silicon oxide epi-layer, causing
the chips to fall off, and leaving the polymer and MoS2 -coated polymer membrane
floating on the surface. The membrane was transferred to deionized water several
times, and then scooped onto a TEM grid and dried. The PMMA was removed by
baking the TEM grids in an ultrahigh-vacuum heater at 300 ◦ C for 10 min or an
atmospheric pressure Ar/H2 gas flow for 4 h. Before atomic-resolution imaging, we
baked the samples overnight in an ultrahigh vacuum at 130 ◦ C.
ADF-STEM. ADF-STEM imaging was conducted using a NION UltraSTEM100.
Although our 60 kV beam is below threshold for damage in MoS2 (ref. 28),
defects readily form under extended imaging, possibly catalysed by surface
contaminants from the transfer process. Imaging conditions were similar
to those in ref. 27. Using a 25–30 mrad convergence angle, our probe size
was close to 1.3 Å. STEM images were acquired with the medium-angle
annular dark-field detector with acquisition times of between 16 and
40 µs per pixel. The atomic resolution images in Fig. 4 were smoothed to
improve contrast.
DF-TEM. TEM imaging and diffraction were conducted using an FEI Technai
T12 operated at 80 kV, which did not cause any apparent damage to the MoS2
membranes at these lower dose densities. Acquisition times for DF-TEM
images (both displaced-aperture and centred DF-TEM) were 5–50 s per
frame. The intensity ranges of DF-TEM images were generally clipped to
improve contrast.
DFT modelling. DFT calculations were performed with the PW91 generalized
gradient approximation for the exchange-correlation functional and ultrasoft
pseudopotentials, as implemented in the Quantum Espresso electronic structure
560
package41 . Structural relaxations were carried out at the gamma point, and
single-point calculations were performed with appropriately converged
Monkhorst-type k-point grids. Supercells were built with about 10 Å separation
between replicas to achieve negligible interaction.
Photoluminescence. The photoluminescence measurements were performed in
a Renishaw InVia Raman Microscope using a 532 nm laser and a 1,800 l mm−1
grating. The curves in Fig. 1c were taken at 100 µW laser power for 10 s. Although
the indirect bandgap transition is not visible in the bilayer, this is not unusual on
a SiO2 substrate. Photoluminescence maps were taken at 100 µW laser power for
1 s with a 300 nm step size.
Electrical measurements. We fabricated back-gated FETs directly on the
285-nm-thick SiO2 dielectric on degenerately doped silicon growth substrates.
Grain boundaries were identified by photoluminescence as shown in Fig. 5. The
MoS2 was shaped by patterning a PMMA mask with electron-beam lithography and
etching with O2 plasma for 20 s at 50 W with 20 s.c.c.m. flow rate. The electrodes
were then patterned by electron beam lithography with a bilayer PMMA stack and
a subsequent evaporation of Al/Cr/Au (50 nm/5 nm/50 nm). We did not anneal the
samples after fabrication.
All devices were measured at room temperature under 10−5 torr inside
a Lakeshore probe station, using the silicon substrates as a global backgate.
We used an Agilent 4155C semiconductor parameter analyser to perform
2-point conductance measurements. FET transfer curves shown in Fig. 5 were
taken at 500 mV drain bias. The devices showed ohmic contacts from −500 to
500 mV source–drain bias.
Received 2 November 2012; accepted 15 March 2013;
published online 5 May 2013
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Acknowledgements
Overall project coordination, sample growth, and electrical and optical characterization
were supported as part of the Center for Re-Defining Photovoltaic Efficiency
Through Molecular-Scale Control, an Energy Frontier Research Center funded
by the US Department of Energy (DOE), Office of Science, Office of Basic
Energy Sciences under Award DE-SC0001085. A.M.v.d.Z was supported by
the EFRC as a research fellow. Electron microscopy was performed at and
supported by the Cornell Center for Materials Research, an NSF MRSEC (NSF
DMR-1120296). P.Y.H. was supported under the National Science Foundation
Graduate Research Fellowship Grant No. DGE-0707428. D.A.C. was supported
by a Columbia University Presidential fellowship and a GEM PhD Fellowship
sponsored by the Center for Functional Nanomaterials at Brookhaven National
Lab. T.C.B. was supported under the Department of Energy Office of Science
Graduate Fellowship Program (DOE SCGF), administered by ORISE-ORAU
under Contract No. DE-AC05-06OR23100. The authors thank S. Gondarenko,
R. Hovden, I. Meric, J. Ravichandran, L. Wang, J. Richmond-Decker and P. Kim
for helpful discussions.
Author contributions
A.M.v.d.Z. supervised and coordinated all aspects of the project. MoS2 growth was carried
out by D.A.C. and A.M.v.d.Z. Electrical characterization and analysis was carried out by
A.M.v.d.Z., D.A.C. and G-H.L. under the supervision of J.C.H. Electron microscopy and
data analysis were carried out by P.Y.H. with D.A.M.’s supervision. Optical spectroscopy
and data analysis were carried out by A.M.v.d.Z. and Y.Y. under T.F.H.’s supervision.
DFT simulations were carried out by T.C.B. under D.R.R.’s supervision. A.M.v.d.Z.,
P.Y.H., D.A.C., T.C.B., Y.Y., T.F.H., D.A.M. and J.C.H. wrote the paper.
Additional information
Supplementary information is available in the online version of the paper. Reprints and
permissions information is available online at www.nature.com/reprints. Correspondence
and requests for materials should be addressed to A.M.v.d.Z.
Competing financial interests
The authors declare no competing financial interests.
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Grains and grain boundaries in highly crystalline
Grains and grain boundaries molybdenum
in highly-­crystalline monolayer molybdenum monolayer
disulphide
disulfide Supplementary Figure S1: Commonly-­observed shapes in MoS2 CVD. Optical micrographs of various CVD MoS2 crystal shapes found in different growths. a) Mo-­‐z-­‐z triangles and 6-­‐point star grown on “clean” Si/SiO2 substrates. Note: the small gold marks are metal alignment marks that were deposited after growth. b) Mo-­‐z-­‐z mirror twin crystal used for electrical devices in Figures 5i-­‐j. c) S-­‐z-­‐z triangles and 5-­‐ and 6-­‐point stars. d) Hexagons. e) Gear-­‐like polycrystalline structures grown on “dirty” substrates. f) 3-­‐point stars grown on “dirty” substrates. a-­‐c) show the types of crystals achieved in large grain growth, while d-­‐f) show the types of crystals grown as a result of dirty substrates or old precursors. 1
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Supplementary Figure S2: Continuous sheet Figure S2a-­‐b show optical images from two different, yet typical, samples similar to Figure 1a. The key difference between the samples is the grain size. On the top left in each image is bare oxide with sparse crystals. In both samples, CVD MoS2 crystals can grow together to form continuous monolayer sheets (bottom right). The 2
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gradients in nucleation density and grain size reflect the substrate proximity to the solid MoO3 source. The small gold and orange dots, indicated by a red circle, are alignment marks placed after growth. Figure S2c shows two representative histograms of island size (as the square root of the area) extracted from portions of Figure 1a and Figure S2a. In the larger grain growths, such as in Figure S2b, the triangles grow together more to form aggregates and continuous sheets, making accurate measurement of grain size impossible without crystalographically sensitive techniques; For this reason only isolated islands in smaller-­‐grain growths are included in Figure S2c, and regions near continuous sheets are excluded from the measurement. The spread in the histogram demonstrate the wide range of grain sizes and spatial inhomogeneity in all of our samples. 3
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Supplementary Figure S3: Raman spectra and mapping a) Raman spectra showing the E2g1 and A1g vibrational modes1 for monolayer (red curve) and bilayer (blue curve) MoS2 corresponding with the photoluminescence spectra from Figure 1c. b-­‐c) Maps of the peak position for the two Raman modes for monolayer tilt boundary from Figure 5 where b) represents the E2g1 mode. and c) represents the A1g mode. Both modes show an upshift of 1 cm-­‐1 at the grain boundary. While this shift may indicate a change in strain or doping at the boundary, it is difficult to interpret compared to the more marked changes to the photoluminescence seen in Figure 5 of the main text. Scale bar 5 µm. 4
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Supplementary Figure S4: High resolution images of edge roughness a) Dark Field TEM image of a single Mo-­‐z-­‐z triangle on a holey amorphous carbon substrate. b-­‐d) STEM images at increasing magnifications from the region inside the red box in a). In c), higher magnification reveals ~ 10 nm edge roughness as highlighted by the orange curve. In d) the MoS2 lattice is visible, as well as atomic scale edge roughness. The non-­‐uniform background variation in all images is due to the amorphous carbon TEM grid support (the perforated sheet in a). 5
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Supplementary Figure S5: DF-­TEM imaging of cyclic twin a) Bright-­‐field image of a 6-­‐pointed star. b) Full diffraction pattern shows star has no rotational boundaries. c-­‐e) Dark-­‐field TEM images corresponding with the c) red, d) blue, e) orange spots in the diffraction image. The red and blue [-­‐1100] spots show opposite intensity because they swap whether the ka or kb spot is captured for each region. The orange spot shows an even intensity over the entire star; this occurs because unlike the [-­‐1100] spots, the intensity of the [-­‐2110] spots are 6-­‐fold symmetric (See Figure 1c). f) Bright-­‐field TEM image with overlaid edge orientations extracted from DF-­‐TEM. The outer edges are oriented along the Mo-­‐z-­‐z direction, which demonstrates that the grain boundaries are oriented along the S-­‐z-­‐z directions. 6
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Supplementary Figure S6: Simulations of alternate mirror-­twin boundaries a) Geometry optimized structure for the experimentally observed grain boundary, as determined by DFT (a). b) The local DOS (LDOS) integrated in the plane of the Mo over a 1.7 eV window inside the band gap of pristine MoS2 confirms the spatial localization of mid-­‐gap states. c) The energy-­‐resolved density of states for pristine MoS2 (black curve), the DOS of the structure shown in (a) (red curve), and the DOS projected just onto the atoms in the grain boundary (blue curve). In (c), we have subtracted off the contribution to the DOS arising from the unphysical edge atoms. 7
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Analogous calculations are shown for an armchair grain boundary d-­‐f) and for a zig-­‐
zag grain boundary g-­‐i); both which are entirely periodic in two dimensions. In panels (b), (e), and (h), the colorscale indicates the magnitude of the integrated LDOS, from 0 (dark) to 0.025 bohr-­‐3 (light). Supplementary Figure S7: Time resolved photoluminescence Time-­‐resolved photoluminescence measurements obtained by time-­‐correlated single photon counting with femtosecond excitation by 400-­‐nm laser pulses. Results for exfoliated and CVD MoS2 samples on oxide layers, after accounting for the instrument response function, yield nearly identical time constants of τexfoliated=44 ps, and τCVD=42 ps. The measured emission decay found to varied considerably from sample to sample for both the exfoliated and CVD grown 8
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materials, so the similarity of the curves in this figure should not be regarded as having fundamental significance. Supplementary Figure S8: Electrical measurements on a tilt boundary a) Linear and b) logarithmic electrical transport transfer curves of 3 FETs fabricated from the tilt boundary MoS2 island shown in the inset of (a), which has a tilt angle of 42° (Scale bar 10 μm). The FETs containing the perpendicular (black) and parallel (orange) boundary orientation to the flow of electrons are both 30% lower in conductance than the pristine region (cyan) in the “on” state (gate voltage = +70 V). 9
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Supplementary Methods: Electron Diffraction Simulation: We calculated the diffraction pattern for a monolayer of MoS2 using Bloch wave simulations to account for the complex scattering that allows the breaking of Friedel’s rule and produces this asymmetry (Figure 2). The asymmetry of the Mo and S sublattices separates the [-­‐1100] diffraction spots into two families : ka={(-­‐
1100), (10-­‐10), (0-­‐110)} and kb= -­ka. Our bloch-­‐wave simulations show that the ~10% higher intensity ka spots point toward the Mo sublattice, as indicated by the arrows in Figures 2(a-­‐b). We double-­‐checked this result with high-­‐resolution imaging to confirm that the asymmetric diffraction pattern in Fig 2(b) corresponds to the indicated orientation of the MoS2 lattice. Identifying triangle edge terminations: Via TEM analyses, we consistently observe that Mo-­‐z-­‐z triangles (Figure 2(b)) have sharper and straighter edges than S-­‐z-­‐z triangles (Figure 2(d)). This morphological difference allows us to rapidly identify the crystal edges and orientation of triangles on the growth substrate simply by optical microscopy. In doing so, we also observe that all crystals from the same growth run have the same morphology, i.e., triangles from a given run will either be dominated by Mo-­‐z-­‐z or by S-­‐z-­‐z edges, a preference we attribute to kinetic effects. Both triangle morphologies exhibit the same range of average sizes from 30-­‐70 µm. These classifications are important for understanding 10
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growth dynamics and suggest the possibility of refined control of edge morphology by tuning the CVD process. DFT calculations: Density functional theory (DFT) calculations were performed with the PW91 generalized gradient approximation for the exchange-­‐correlation functional and ultrasoft pseudopotentials, as implemented in the Quantum Espresso electronic structure package5. Supercells were generated with about 10 Å separation to ensure negligible interactions between replicas. Structural relaxations were carried out at the gamma point until all components of all forces were less than 0.001 a.u. Pristine MoS2 (3.12 and 2.32 Å for Mo-­‐Mo and Mo-­‐S bond lengths, respectively) energy calculations were done with a 16x16 Monkhorst-­‐type k-­‐point grid, confirming the material's direct band gap with a predicted energy of 1.9 eV. Preliminary calculations employing a finer k-­‐point grid for the structural relaxation show minor quantitative but not qualitative changes to the grain boundary geometry and electron structure6. Because the direction of the experimentally observed 8-­‐4-­‐4 grain boundary is incommensurate with the periodicity of the underlying crystal, a system periodic in two dimensions cannot be constructed. Thus we employed the system shown in Figure S6(a) which is periodic along the direction of the grain boundary, but finite in the orthogonal direction such that the edge, terminated by S dimers, is about 10 Å 11
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away from the grain boundary. Energy calculations for this 87-­‐atom supercell were performed with 5 k-­‐points in the periodic direction. Local density of states (LDOS) analysis, shown in Figure S6(b) confirms that the electronic effects of the artificial edges are physically confined along the perimeter and so should not affect the properties along the grain boundary. These conclusions are also corroborated by a negligible change in bond length, compared to the bulk, for atoms away from the grain boundary. Analogous calculations on 2D periodic systems with pure armchair and zig-­‐zag grain boundaries, shown in Figure S6 (d-­‐i) similarly yield mid-­‐gap states localized along the boundary, further indicating that the effect is generic and not an artifact of the finite strip size. Estimating non-­radiative recombination: Photoluminescence quenching commonly arises from defects in semiconductors, such as the predicted midgap states at the boundaries, which can act as centers for non-­‐radiative recombination7. While the amount of material structurally modified by the boundary is small compared with the 500-­‐nm laser spot size, the effect can be enhanced by the diffusion of photogenerated excitons to the boundary, which effectively increases the boundary width. Such a process would provide a natural explanation for the strongly reduced PL observed from some boundaries in our samples, since non-­‐radiative recombination may be more efficient in the presence of the localized structures and states of the boundary. Here we estimate the potential impact of this effect using measured values for the carrier mobility and exciton lifetime. 12
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DOI: 10.1038/NMAT3633
The electrical measurements in this paper yielded an electron mobility of µe = 3 cm2V-­‐1s-­‐1 for typical CVD-­‐grown samples. From the Einstein relation for a temperature of T = 300 K, we then obtain an electron diffusivity of = 8 x 10-­‐2 cm2s-­‐1. Since the electron and hole masses are similar, we assume that the exciton diffusivity is half of the electron diffusivity, i.e., Dexc = 4 x 10-­‐2 cm2s-­‐1. The scattering mechanisms for electrons and excitons need not be the same, so this should only be considered as an estimate. The characteristic length for diffusion of excitons towards the boundary from the 2-­‐
D bulk material over a time τ is given by . For a time interval of τ = 40 ps reported above for the exciton lifetime, we then infer a diffusion length of L = 24 nm. This length, while not negligible, is small compared to the 500 nm spot size of the excitation laser. The quenching of the PL by 50% observed at some grain boundaries must consequently rely primarily, as discussed in the main text, on mechanisms other than exciton diffusion. Supplementary References: 1 2 3 4 Lee, C. et al. Anomalous Lattice Vibrations of Single-­‐ and Few-­‐Layer MoS2. ACS Nano 4, 2695-­‐2700. Huang, P. Y. et al. Grains and grain boundaries in single-­‐layer graphene atomic patchwork quilts. Nature 469, 389-­‐392 (2011). Kim, K. et al. Grain Boundary Mapping in Polycrystalline Graphene. ACS Nano 5, 2142-­‐2146 (2011). Lahiri, J., Lin, Y., Bozkurt, P., Oleynik, I. I. & Batzill, M. An extended defect in graphene as a metallic wire. Nature Nanotechnology 5, 326-­‐329 (2010). 13
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5 6 7 14
DOI: 10.1038/NMAT3633
Paolo, G. et al. QUANTUM ESPRESSO: a modular and open-­‐source software project for quantum simulations of materials. Journal of Physics: Condensed Matter 21, 395502 (2009). J. Kunstmann, T. B., D. Reichman. Unpublished results. Neamen, D. A. Semiconductor physics and Devices. (McGraw-­‐Hill, 2002). NATURE MATERIALS | www.nature.com/naturematerials
© 2013 Macmillan Publishers Limited. All rights reserved.