Low-Temperature Solution Synthesis of Transition Metal Dichalcogenide Alloys with Tunable Optical Properties

Article abstract of DOI:10.1021/jacs.7b04443

Nanostructures of layered transition metal dichalcogenide (TMD) alloys with tunable compositions are promising candidates for a broad scope of applications in electronics, optoelectronics, topological devices, and catalysis. Most TMD alloy nanostructures are synthesized as films on substrates using gas-phase methods at high temperatures. However, lower temperature solution routes present an attractive alternative with the potential for larger-scale, higher-yield syntheses of freestanding, higher surface area materials. Here, we report the direct solution synthesis of colloidal few-layer TMD alloys, Mo_xW_1-xSe₂ and WS_2ySe_2(1-y), exhibiting fully tunable metal and chalcogen compositions that span the MoSe₂-WSe₂ and WS₂-WSe₂ solid solutions, respectively. Chemical guidelines for achieving the targeted compounds are presented, along with comprehensive structural characterizations (X-ray diffraction, electron microscopy, Raman, and UV-visible spectroscopies). High-resolution microscopic imaging confirms the formation of TMD alloys and identifies a random distribution of the alloyed elements. Analysis of the tilt-angle dependency of the intensities associated with atomic-resolution annular dark field imaging line scans reveals the types of point vacancies present in the samples, thus providing atomic-level insights into the structures of colloidal TMD alloy nanostructures that were previously only accessible for substrate-confined films. The A excitonic transition of the TMD alloy nanostructures can be readily adjusted between 1.51 and 1.93 eV through metal and chalcogen alloying, correlating the compositional modulation to the realization of tunable optical properties.

Full text of DOI:10.1021/jacs.7b04443

Article
pubs.acs.org/JACS
Low-Temperature Solution Synthesis of Transition Metal
Dichalcogenide Alloys with Tunable Optical Properties
Yifan Sun,†,⊥ Kazunori Fujisawa,‡,⊥ Zhong Lin,‡,⊥ Yu Lei,§,⊥ Jared S. Mondschein,†
Mauricio Terrones,*,†,‡,§,⊥ and Raymond E. Schaak*,†,∥,⊥
‡
§
∥
†Department of Chemistry, Department of Physics, Department of Materials Science and Engineering, Materials Research
⊥
Institute, and Center for 2-Dimensional and Layered Materials, The Pennsylvania State University, University Park, Pennsylvania
16802, United States
S
* Supporting Information
ABSTRACT: Nanostructures of layered transition metal dichalcogenide
(TMD) alloys with tunable compositions are promising candidates for a
broad scope of applications in electronics, optoelectronics, topological devices,
and catalysis. Most TMD alloy nanostructures are synthesized as films on
substrates using gas-phase methods at high temperatures. However, lower
temperature solution routes present an attractive alternative with the potential
for larger-scale, higher-yield syntheses of freestanding, higher surface area
materials. Here, we report the direct solution synthesis of colloidal few-layer
TMD alloys, MoxW1−xSe2 and WS2ySe2(1−y), exhibiting fully tunable metal and chalcogen compositions that span the MoSe2−
WSe2 and WS2−WSe2 solid solutions, respectively. Chemical guidelines for achieving the targeted compounds are presented,
along with comprehensive structural characterizations (X-ray diffraction, electron microscopy, Raman, and UV−visible
spectroscopies). High-resolution microscopic imaging confirms the formation of TMD alloys and identifies a random distribution
of the alloyed elements. Analysis of the tilt-angle dependency of the intensities associated with atomic-resolution annular dark
field imaging line scans reveals the types of point vacancies present in the samples, thus providing atomic-level insights into the
structures of colloidal TMD alloy nanostructures that were previously only accessible for substrate-confined films. The A
excitonic transition of the TMD alloy nanostructures can be readily adjusted between 1.51 and 1.93 eV through metal and
chalcogen alloying, correlating the compositional modulation to the realization of tunable optical properties.
tune the physical properties by continuously modulating the
elemental composition. Substitution of a parent TMD
compound with different metal or chalcogen atoms alters the
band structure, leading to tunable excitonic transitions between
the valence band and the conduction band. Physical vapor
deposition (PVD),10 chemical vapor deposition (CVD),11
mechanical exfoliation of bulk TMD alloys prepared via
chemical vapor transport (CVT),12 and chalcogen exchange13
have been used to access TMD alloy nanosheets that are
investigated for electronics,14 optoelectronics,15 catalysis,16 and
topological devices.17 For example, TMD alloy monolayers
exhibit tunable photoluminescence emissions,18 as well as
carrier type modulation,14 and high on/off ratio in field effect
transistors (FETs).19 With favorable hydrogen adsorption free
energy (ΔGH) on the edges and chemical activation on the
basal planes induced by atom substitution, TMD chalcogen
alloys are attractive hydrogen evolution reaction (HER)
catalysts.16 Recent studies also indicate that MoxW1−xTe2 is a
Type II Weyl semimetal with exotic topological properties.17
As an alternative to the synthetic methods highlighted above,
solution-based routes to TMD alloy nanosheets are attractive
for potentially achieving greater scalability and yield, substrate-
INTRODUCTION
■
Layered transition metal dichalcogenides (TMDs) are of broad
interest as two-dimensional (2-D) beyond-graphene materials
due to their unique and diverse properties.1,2 Coupled by weak
van der Waals interactions, each TMD layer is composed of a
transition metal atom (M) covalently bonded to six chalcogen
atoms (X). Hexagonal (2H) phases composed of metal−
chalcogen layers stacked in a trigonal prismatic arrangement are
most commonly observed, although the identities of the metal
and chalcogen atoms and their coordination geometries can
lead to rich polymorphism, with multiple classes of TMD
structure types each having distinct properties.3 The anisotropic
nature of 2-D TMD crystals facilitates the formation of both
monolayer and few-layer nanosheets with unique properties
that depend on their thicknesses.4 The group VI TMD
compounds MoS2, WS2, MoSe2, and WSe2 have been studied
most extensively because of their narrow semiconducting band
gaps that overlap well with the solar spectrum, as well as the
emergence of an intrinsic indirect-to-direct band gap transition
at monolayer thicknesses.5
Approaches such as defect engineering,6 strain engineering,7
surface functionalization,8 and heterostructure fabrication9 have
been used to modify the optical properties of TMDs to realize
new properties and further expand the scope of applications. A
complementary approach is to form TMD alloys that can fine-
Received: May 1, 2017
© XXXX American Chemical Society
A
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
mmol) and mixed with 5 mL of oleylamine (15.2 mmol) in an argon-
flushed septum-capped vial. Upon injection, 0.24 mL of CS2 was
introduced to the vial, forming a homogeneous solution with a
temperature increase. The solution in the vial was subsequently
injected dropwise into the three-neck flask at a rate of 2 mL/h using a
syringe pump. Over a 30 min interval, the injection was stopped and
the heating mantle was removed. After cooling down to room
temperature, the products were washed three times with a 1:1 toluene/
ethanol mixture, collected with centrifugation, and kept as a powder
under argon. For Se-rich samples (y < 0.5), the synthetic approach was
identical to that for MoxW1−xSe2, except diphenyl disulfide and
diphenyl diselenide were dissolved together in oleylamine in the
beginning and reaction time was modified. Table S2 provides detailed
information about the experimental parameters. Scheme S1 in the
Supporting Information summarizes the synthetic routes for both the
metal and chalcogen alloys.
Characterization. Powder X-ray diffraction (XRD) patterns were
collected using a Bruker-AXS D8 Advance diffractometer equipped
with Cu Kα radiation and a LynxEye 1-D detector. Simulated XRD
patterns of WS2, MoSe2, and WSe2 were generated using the
CrystalMaker/CrystalDiffract software package. Transmission electron
microscopy (TEM), high-angle annular dark-field scanning trans-
mission electron microscopy (HAADF-STEM) images, and energy
dispersive X-ray spectroscopy (EDS) data with element maps were
acquired using a FEI Talos F200X operating at 200 kV. High-
resolution ADF images were obtained using a FEI Titan3 G2 60/300
TEM with a spherical aberration corrector on both the probe and the
image-forming lens at an accelerating voltage of 80 kV. Lower voltage
was used at high resolution to decrease irradiation damage. To
enhance visibility and reduce noise, all acquired high-resolution images
were processed by Gaussian Blur filter using ImageJ software. STEM
ADF image simulation was done by QSTEM software developed by C.
Koch.28 All parameters for the ADF image simulation were
appropriately set according to experimental imaging condition,
including acceleration voltage, spherical aberration (C3 and C5),
convergence angle, and inner/outer angle for the ADF detector.
Thermal diffuse scattering (TDS) was taken into account and tilting
angle was applied by QSTEM software by rotating the sample
structure. Bruker ESPRIT 2 software was applied for EDS data
interpretation. Micro-Raman measurements were performed in a
Renishaw inVia confocal microscope-based Raman spectrometer with
514.5 nm laser. The laser spot size using the 100× objective lens was
approximately 1 μm. Alloy samples were dispersed in ethanol assisted
by sonication and then ultraviolet−visible (UV−vis) transmittance
measurements were performed using a LAMBDA 950 UV/vis/NIR
spectrometer (350−1200 nm, 0.5 nm data interval).
free samples with higher surface areas for catalysis, and mild
reaction conditions that can lead to the formation of different
phases.20 However, existing solution routes to TMD alloy
nanostructures require high-temperature annealing to homoge-
neously incorporate the alloyed elements,21,22 which is
detrimental to nanostructure retention, surface area, and
solution dispersibility, or are applicable to only a limited
range of alloy compositions.23,24 Additionally, little is known
about the atomic-level distribution of the alloyed elements in
solution-synthesized TMD nanostructures.21−25
Here, we report a low-temperature solution-phase approach
to directly synthesize few-layer colloidal TMD alloy nanostruc-
tures. Both metal [MoxW1−xSe2] and chalcogen [WS2ySe2(1−y)
]
alloys across the entire solid solution ranges were obtained via
isoelectronic substitution mediated by strategic selection of
reagents and reaction pathways. The colloidal MoxW1−xSe2 and
WS2ySe2(1−y) alloys exhibit continuously tunable optical proper-
ties, including excitonic transitions that range from 1.51 to 1.93
eV. In addition to providing chemical guidelines for achieving
complete access to the MoSe2−WSe2 and WS2−WSe2 solid
solutions using colloidal synthesis methods, we present a
detailed microscopic study that identifies the locations and
distributions of the alloying elements and vacancies. Such
atomic-level structural information has not been available
previously for colloidal TMD nanostructures, since the few-
layer nanosheets that comprise them are not flat on a substrate
as they are in more traditional CVD-grown films. The approach
used to obtain atomic-level microscopic insights is anticipated
to be broadly applicable across a wide range of colloidal
nanostructures that have buckled, crumpled, or curved
components.
EXPERIMENTAL SECTION
■
Materials. Oleylamine (technical grade, 70%), tungsten(VI)
chloride (WCl6, ≥99.9%, trace metal basis), diphenyl diselenide
(98%), diphenyl disulfide (99%), hexamethyldisilazane (HMDS,
reagent grade, ≥99%), and carbon disulfide (CS2, ≥99.9%, anhydrous)
were purchased from Sigma-Aldrich. Molybdenum(V) chloride
(MoCl5, 99.6%, metal basis) and oleic acid (technical grade, 90%)
were purchased from Alfa Aesar. Solvents, including toluene and
ethanol, were of analytical grade. All chemicals were used as received
without further purification.
Synthesis of MoxW1−xSe2 Alloy Nanostructures. All reactions
were carried out under an argon atmosphere using standard Schlenk
techniques and workup procedures were performed in air. In the
synthesis of MoxW1−xSe2, 39 mg of diphenyl diselenide (0.125 mmol)
dissolved in 15 mL of oleylamine (45.6 mmol) were added to a 100
mL three-neck flask and degassed for 30 min under vacuum at 120 °C.
Then, 0.5 mL of HMDS was injected into the flask after cooling to 100
°C under argon and the mixture was heated to 300 °C. Meanwhile,
0.125 mmol of MoCl5, WCl6, or a mixture of the two (Table S1) was
dissolved in 1 mL of oleic acid (3.1 mmol) with sonication. The metal
reagent was then injected dropwise into the hot mixture at a rate of 2
mL/h using a syringe pump. After 30 min, the injection was completed
and the heating mantle was removed to allow the reaction mixture to
cool to room temperature. The products were washed three times with
a 1:1 toluene/ethanol mixture and kept as a powder under argon.
Synthesis of WS2ySe2(1−y) Alloy Nanostructures. To prepare
sulfur-rich chalcogen alloys with y > 0.5, CS2 was used as the sulfur
source, as reported previously for the synthesis of colloidal WS2,26 and
diphenyl diselenide was used as the selenium source.27 Diphenyl
diselenide dissolved in 15 mL of oleylamine was added to a 100 mL
three-neck flask and degassed for 30 min under vacuum at 120 °C.
HMDS (0.5 mL) was then injected into the flask after cooling to 100
°C under argon and the mixture was heated to 300 °C. Meanwhile,
WCl6 (50 mg, 0.125 mmol) was dissolved in 0.3 mL of oleic acid (0.95
RESULTS AND DISCUSSION
■
Colloidal MoxW1−xSe2 Alloy Nanostructures. The TMD
alloys with variable metal element composition spanning the
MoSe2−WSe2 solid solution (MoxW1−xSe2) were synthesized
via the dropwise injection of metal reagents (MoCl5, WCl6)
dissolved in oleic acid into a solution of oleylamine, HMDS,
and diphenyl diselenide at 300 °C, as described in detail in the
Experimental Section. MoSe2 and WSe2 both adopt the
hexagonal 2H structure type and have closely matched lattice
constants and similar electronic structures,29 which helps to
facilitate the formation of MoxW1−xSe2 alloys rather than a
mixture of MoSe2 and WSe2. Colloidal MoxW1−xSe2 nanostruc-
tures with x = 0, 0.14, 0.35, 0.52, 0.77, and 1 were prepared by
modifying the amounts of MoCl5 and WCl6 dissolved in oleic
acid.
Figure S1 shows powder XRD data for MoxW1−xSe2
nanostructures with x = 0, 0.14, 0.35, 0.52, 0.77, and 1, as
well as reference patterns for bulk MoSe2 and WSe2 for
comparison.30,31 It is well-known that irregular stacking of
nanosheets, small domain sizes, and lattice strain induced by
B
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
defects and curvature lead to significantly broadened diffraction
peaks in solution-synthesized TMD nanostructures. However,
peaks corresponding to the (002), (100), (103), (105), (110),
(008), (200), and (203) planes can be identified in the XRD
patterns for all of the MoxW1−xSe2 nanostructures. The lattice
constants for the end members are closely matched: for MoSe2,
a = 3.288 Å and c = 12.931 Å;30 while for WSe2, a = 3.282 Å
and c = 12.96 Å.31 Therefore, the XRD patterns are nearly
identical and peak positions do not shift in a significant or
measurable way as x changes, which makes it difficult to
distinguish between mixtures of separated MoSe2 and WSe2 vs
the targeted MoxW1−xSe2 alloys.
To better characterize the TMD alloy nanostructures, a suite
of microscopic and spectroscopic techniques were used. As is
typical for solution-synthesized TMD nanostructures, the
MoxW1−xSe2 nanosheets agglomerate to produce nanoflower-
like morphologies,2 as shown by TEM in Figure S2. As a
representative example, the STEM and TEM images of
Mo0.35W0.65Se2 in Figures 1a and 1b reveal uniform nanoflowers
Figure S3, indicate that the Se signal is nearly constant while
the Mo and W signals increase and decrease in intensity,
respectively, as x increases. Figure 1g shows a plot of the
relative amounts of Se, W, and Mo in each MoxW1−xSe2 sample,
as derived from the EDS data in Figure S3. All samples contain
∼65% Se and 35% metal (W + Mo). The Mo/W ratios vary
linearly across the entire solid solution, indicating the Mo/W
ratio can be readily tuned by varying the metal reagent
stoichiometries. Additionally, the nearly constant chalcogen/
metal ratio of approximately 1.8 implies that all of the
MoxW1−xSe2 nanostructures are chalcogen-deficient. Sub-
strate-confined TMD nanosheets synthesized using traditional
gas-phase deposition routes commonly contain chalcogen
defects due to intrinsic grain boundaries, atomic vacancies, or
substitutional doping.6 Ex situ treatments, such as mechanical
processing and heterolayer stacking,6 can also result in the
formation of vacancies. For the colloidal MoxW1−xSe2
nanostructures, defects are most likely to arise from domain
boundaries in the polycrystalline nanosheets,32 or chalcogen
vacancies induced by Se dissolution in oleylamine at elevated
temperatures.33
Figure 2, which shows ADF-STEM images of the
Mo0.35W0.65Se2 nanoflowers, provides high-resolution, atomic-
level structural and compositional information about the TMD
alloys. The ADF-STEM image in Figure 2a corresponds to the
edge of a Mo0.35W0.65Se2 nanoflower; the top and left sides are
closer to the center of the nanoflower, while the majority of the
panel shows a terminal nanosheet that is exposed at the edge of
the nanoflower. The ADF-STEM image shows that the TMD
layer thickness decreases moving from the interior to the edge
of the nanoflower, ultimately reaching monolayer thickness at
the edges with domain sizes of 5−10 nm. While monolayer,
bilayer, and trilayer domains can be differentiated by contrast,
ADF-STEM has the ability to also distinguish atoms based on
Z-contrast (Z = atomic number).34 For MoxW1−xSe2, in
addition to the Mo vs W site occupancies, three chalcogen
configurations exist: diselenium (Se2), monoselenium vacancy
(VSe), and diselenium vacancy (VSe2). An ADF line scan was
acquired from the high-resolution STEM image, as indicated by
the green arrow in Figure 2b. The corresponding line scan
profile, shown in Figure 2c, reveals the Mo vs W site
occupancies as well as the chalcogen configurations. Simu-
lations of the STEM-ADF intensity profile based on MoSe2,
WSe2, and multiple vacancy configurations was carried out to
unambiguously assign the identities of the atoms (Figure 2d
and e). Based on the assumption that the dodecagonal-shaped
vacancies belong to chalcogen sites, as supported from EDS
quantification, simulation data, and previous reports,35 the left
hump of the two-peak pair was assigned as the chalcogen
sublattice, and the right one was assigned as the metal
sublattice. Taken together, this analysis allows identification of
the Mo and W atom locations, as well as the observation of
both Se2 and VSe vacancies.
Figure 1. (a) HAADF-STEM image, (b) TEM image, and (c−f)
STEM-EDS element maps (100 nm scale bars) of Mo0.35W0.65Se2
nanostructures. (g) Plots of elemental composition (green) and metal/
chalcogen ratio (blue) as a function of x in the MoxW1−xSe2 alloy
nanostructures.
A growing number of reports describe the solution synthesis
of TMD nanostructures, and while the colloidal TMD
nanostructures are typically crumpled, buckled, and/or
aggregated, they remain of interest as potential materials for
catalysis, particularly as discrete nanoflowers that expose a high
density of catalytically active edge sites.36 The high surface areas
contributed by the 2-D layers are also of broad interest for
applications in energy storage and conversion, as well as gas
sensing.37 It is thus important to be able to fully characterize
the atomic distributions and vacancies in TMD alloy
with diameters of approximately 200 nm. The nanoflowers are
comprised of curled 2-D nanosheets that protrude from the
thicker core to form a three-dimensionally structured
architecture. Figure 1c−f also shows STEM-EDS mapping
data for the Mo0.35W0.65Se2 nanostructures, indicating that Mo,
W, and Se are uniformly distributed throughout the nano-
flowers. EDS spectra for each MoxW1−xSe2 sample, shown in
C
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
Figure 2. (a) High-resolution ADF-STEM image highlighting a region of a Mo0.35W0.65Se2 nanostructure, showing few-layer, bilayer, and monolayer
domains with overall decreasing layer thickness toward the edges. (b) Atomically resolved ADF-STEM image of Mo0.35W0.65Se2 nanostructures,
where dodecagonal-shaped chalcogen vacancies are circled. (c) Experimental ADF intensity curve (gray) corresponding to the line scan indicated by
the green arrow in (b). The atoms are identified according to the simulated profiles (colored). Simulated structures of the diselenium configuration
(Se2, black) and the monoselenium vacancy (VSe, red) are shown for (d) MoSe2 and (e) WSe2, along with the corresponding ADF intensity curve
with 0° tilting angles. (f) Top and (g) side views of the hexagonal TMD structure, where metal (pink) and chalcogen (yellow) atoms are covalently
bonded with trigonal prismatic coordination.
nanostructures, as these features are known to significantly
impact properties. Taking MoS2 as an example, both p- and n-
type transfer behaviors have been identified from sulfur and
molybdenum vacancies for MoS2, even for different regions of
the same sample.38 On the other hand, chemical treatment for
monolayer MoS2 induces an enhanced photoluminescence
emission with near-unity quantum yield, due to the elimination
of defect-mediated recombination.39 Substrate-confined TMD
nanosheets are flat and therefore able to be adequately
characterized using microscopic techniques that rely on analysis
of intensity profiles. However, TMD nanoflowers and related
nanostructures that are not flat contain 2-D nanosheets that
span a wide range of orientations, and as a result adopt various
tilt angles relative to the imaging plane, which can significantly
impact the intensity profiles.
While the STEM-ADF analysis in Figure 2 provides good
agreement between the elemental assignments and composi-
tions for the TMD metal alloy nanostructures, the relative ADF
intensity ratios between the chalcogen sites and the metal sites
do not match quantitatively with the simulated values. For
example, simulations indicate that the Se2 configuration should
have a higher ADF intensity than W. In contrast, W appears to
have a higher intensity than Se2 in the ADF line-scan profile in
Figure 2. To understand this deviation between the simulated
and observed intensities, we investigated the effect of tilt on
ADF intensity. When a hexagonally structured 2H-type TMD
nanosheet such as MoxW1−xSe2 is flat, the two chalcogen atoms
of the Se2 configuration are stacked in a column.40 The ADF
intensity ratio is therefore sensitive to the relative angle
between the electron probe and the sample plane, because off-
axis samples will be tilted relative to the electron probe, which
will decrease the ADF intensity and broaden the ADF intensity
profile at chalcogen positions. Figure 3a shows high-resolution
ADF-STEM images for a region of a colloidal MoxW1−xSe2
nanosheet spanning tilt angles that range from 0° to 20°.
Included in the imaged region are adjacent Se2/W, VSe/W, and
VSe/Se2 configurations. As shown in Figure 3b, along with the
corresponding plot of ADF intensity region vs tilt angle, the
Se2/W and VSe/Se2 configurations exhibit a strong tilt angle
dependency, while the value of VSe/W remains nearly constant.
At high tilt angles approaching 20°, the intensity from the
monoselenium vacancy VSe becomes close to that from the
diselenium Se2 configuration. At a tilt angle of 8°, the simulated
ADF intensity ratios of Se2/W, VSe/W, and VSe/Se2 are 0.79,
0.31, and 0.40, respectively, which agree well with exper-
imentally obtained values of 0.78, 0.35, and 0.44. This suggests
that the imaged nanosheets are indeed tilted, and confirms that
tilt angle analysis is important for quantitatively matching
D
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
Figure 4. Raman spectra of the as-prepared MoxW1−xSe2 nanostruc-
tures, acquired using a 514.5 nm excitation laser. The spectra are
normalized to the intensity of the A1g-like peak. Raman modes are
labeled based on bulk notation.
few-layer WSe2, both modes are degenerate at around 250 cm−1
to contribute to a combined A1g-like peak.42 The A1g-like mode
of MoxW1−xSe2 shows a prominent blue shift from 237.8 to
250.0 cm−1 that correlates linearly with W content (Figure S5),
as interpreted with the modified random element isodisplace-
ment (MREI) model based on effective mass and out-of-plane
force constants.19 The structural evolution is also accompanied
by some weak-vibration second- and higher-order modes, as
evident in Figure 4. The insights gained from analysis of the
Raman spectra further confirm the formation of a composition-
tunable series that spans the MoSe2−WSe2 compositions.
Colloidal WS2ySe2(1−y) Alloy Nanostructures. Compared
with MoxW1−xSe2 alloys that continuously vary the metal while
keeping the chalcogen constant, the related chalcogen counter-
parts, WS2ySe2(1−y), are more synthetically challenging.
Although theoretical calculations indicate that both
MoxW1−xSe2 and WS2ySe2(1−y) solid solutions are energetically
favorable relative to phase-separated mixtures,43 chemical issues
such as precursor incompatibility,44 chalcogen replacement,13
and product decomposition10 make it difficult to tune
chalcogen content across the entire solid-solution range using
colloidal synthesis methods. For gas-phase synthesis methods,
experimental parameters including temperature gradient,18
partial vapor pressure,10 reaction time,13 and precursor position
in the tube15 can be tailored to achieve thermodynamic or
kinetic control and accomplish the atomic-level mixing of S and
Se within solid-solution TMDs. Similar full-range TMD alloys
have not been accessible previously as solution-synthesized
colloidal nanostructures.
The successful synthesis of colloidal WS2ySe2(1−y) nanostruc-
tures requires careful consideration of reagent reactivity. While
the MoCl5 and WCl6 reagents exhibit similar reactivities and
therefore can be mixed together in the desired stoichiometries
to obtain targeted MoxW1−xSe2 alloys, the available sulfur
reagents evaporate and/or decompose rapidly, which substan-
tially diminishes the utilization of sulfur during the reaction.
Therefore, in contrast to the stoichiometric synthesis of
MoxW1−xSe2, an excess of S is required to produce y-tunable
WS2ySe2(1−y) ternary compounds. For sulfur-rich samples
[WS2ySe2(1−y) having y = 1, 0.81, 0.70 and 0.60], diphenyl
diselenide could be added directly to an existing protocol for
Figure 3. Effect of tilting on the STEM-ADF intensity in WSe2
nanostructures. (a) Simulated ADF images of the structure of WSe2 at
tilting angles ranging from 0° to 20°. (b) Calculated STEM-ADF
intensity ratios of VSe/Se2, Se2/W, and VSe/W at different tilting
angles. Se2 exhibits higher ADF intensity when the tilting angle is
smaller than 6°, while the relationship is inverted when tilted more
than 7°.
simulated and observed ADF intensity data that is used to
determine the locations of the metal and chalcogen atoms in
colloidal TMD nanostructures. Additional strain and curvature
within the nanosheets may also impact the ADF intensity
simulation, although to a lesser extent than the tilt angle.
Overall, the combined analyses confirm that the colloidal TMD
MoxW1−xSe2 nanostructures are true solid-solution alloys with
atomic-level mixing of Mo and W rather than segregated phases
of MoSe2 and WSe2.
Raman scattering spectra probes the dynamic interaction
between the metal and chalcogen atoms, which provides more
information about atom substitutions in low-dimensional
materials.41 Figure 4 shows the Raman spectra of the
MoxW1−xSe2 nanostructures measured with 514.5 nm laser.
For MoSe2, the out-of-plane mode A1g and in-plane mode E21g
are centered at 237.8 and 289.1 cm−1, respectively, while for
E
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
synthesizing colloidal WS2 nanostructures using CS2 as the
sulfur reagent.26 Increasing the amount of diphenyl diselenide
beyond WS1.20Se0.80 does not result in a substantial increase of
selenium incorporation, while reducing the amount of diphenyl
disulfide results in amorphous products. Therefore, simple
modification of the chalcogen precursor amounts cannot
achieve full tunability across the WS2−WSe2 solid solution,
limiting the formation of WS2ySe2(1−y) from the WS2 end
member (i.e., sulfur-rich compositions) to a maximum of y =
0.60. Similar situations also occur for gas-phase CVD or PVD
syntheses of TMD chalcogen alloys.16,45
To access Se-rich WS2ySe2(1−y) alloys, we began by adding the
sulfur reagent to the WSe2 end member. Diphenyl disulfide was
used as the sulfur reagent instead of CS2 due to its volatility and
fast decomposition at relatively low temperatures. As shown in
Table S2, the sulfur utilization is much lower than that of
selenium, so excess of the sulfide reagent was used. At the same
time, the typical 30 min reaction time resulted in the formation
of Se-containing nanoplate impurities (Figure S6). We
hypothesize that excess sulfur in solution might consume the
tungsten reagent at a faster rate than selenium does, leaving
unreacted selenium available for undesired side reactions.
Therefore, to achieve the desired reactivity and access the
targeted Se-rich WS2ySe2(1−y) alloys, shorter reaction times were
used along with tuning the S/Se reagent composition.
Figure 5a shows the XRD patterns for chalcogen alloy series.
Comparing the crystal structures of end members of
WS2ySe2(1−y), the hexagonal lattice is expanded by ∼4% laterally
and ∼5% vertically for WSe2 vs WS2.31 The change is much
larger than that for metal substitution (MoxW1−xSe2), as Se is
larger than S while Mo and W are more closely matched in size.
Accordingly, XRD peaks corresponding to (100), (103), (105),
and (110) planes for the WS2ySe2(1−y) samples exhibit a
progressive shift to higher 2θ values as y increases. Figure S7a
and b, which shows HAADF-STEM and TEM images of the
WS2ySe2(1−y) nanostructures with y = 0.44, reveals nanosheet
aggregates having morphologies similar to the MoxW1−xSe2
samples but with less-defined and lower-density nanoflower
architectures due to the shorter reaction times needed to
synthesize the chalcogen alloys relative to the metal alloys. As
shown in Figure S8, the average diameter of the WS2ySe2(1−y)
nanostructures decreases as the amount of incorporated sulfur
increases. The STEM-EDS element maps in Figure S7c−f
confirm the uniform distribution of W, S, and Se within the
WS1.56Se0.44 nanostructures. The EDS spectra for all
WS2ySe2(1−y) nanostructures, shown in Figure S9 and
summarized in the plot in Figure 5b, confirm the tunable S/
Se composition ratio across the entire WS2−WSe2 solid
solution, and indicate the chalcogen deficiency with a constant
metal/chalcogen ratio (∼1.8) as for the MoxW1−xSe2 metal
alloys.
Similar to the data in Figure 2 for the Mo0.35W0.65Se2 metal
alloy sample, Figure 6 shows ADF-STEM images highlighting
the edges of a WS0.88Se1.12 chalcogen alloy nanostructure.
Monolayers, bilayers, and few-layer regions are visible, with an
average domain size of ∼5−10 nm (Figure 6a). At higher
magnification (Figure 6b), a line scan across the monolayer
region reveals information about the atom distributions in the
WS0.88Se1.12 sample. Using the strategy described in detail for
analysis of the Mo0.35W0.65Se2 sample in Figure 2, we were able
to identify distinct chalcogen configurations among the
possibilities of diselenium (Se2), disulfur (S2), monosulfur
plus monoselenium (S + Se), monoselenium vacancy (VSe),
Figure 5. (a) Powder XRD patterns of WS2ySe2(1−y) nanostructures
with y = 0, 0.06, 0.13, 0.44, 0.60, 0.70, 0.81, and 1. Simulated
diffraction patterns based on the crystal structures of bulk WS2 and
WSe2 are provided at top and bottom, respectively, for comparison.
(b) Plots of elemental composition (orange) and metal/chalcogen
31
ratio (blue) as a function of y in the WS2ySe2(1−y) alloy nanostructures.
monosulfur vacancy (VS), and double-chalcogen vacancy (V)
via ADF intensity simulations (Figures 6c and 7a). These
results confirm atomic-level mixing of S and Se and therefore
the formation of a S:Se solid solution, and provide
unprecedented insight into atomic distributions in colloidally
synthesized TMD nanostructures. The different chalcogen
configurations have different ADF intensities, as expected, as
well as distinct characteristic peak-to-peak distances. As shown
in Figure 6c, convolution of the smaller S atom results in a
broad shoulder next to the W peak due to its shorter W−S
bond length relative to W−Se, while the larger Se atom next to
a W atom can be resolved as two distinct W and Se peaks. As
for the metal alloy samples, the ADF intensity for the chalcogen
alloys is also dependent on the tilt angle during imaging (Figure
7b). Accordingly, at a tilt angle of 10°, the experimentally
measured ADF intensity ratios [Se2/W = 0.71, S2/W = 0.18, (S
+ Se)/W = 0.59, and VSe/W = 0.50)] become comparable to
the calculated ratios [Se2/W = 0.72, S2/W = 0.24, (S + Se)/W
= 0.42, and VSe/W = 0.34].
Figure 8 shows composition-dependent Raman spectra for
the WS2ySe2(1−y) samples. WS2 nanoflowers show the out-of-
plane A1g mode at 419.5 cm−1, the peak contributed by the in-
plane E12g mode at 350.0 cm−1, and the second-order 2LA(M)
F
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
Figure 7. (a) Simulated ADF images of monosulfur plus
monoselenium (S + Se) and disulfur (S2) configurations for TMD
chalcogen alloys with 0° tilting angles. Calculated ADF intensity curves
for diselenium (Se2, black), monosulfur plus monoselenium (S + Se,
blue), disulfur (S2, dark red), monoselenium vacancy (VSe, red), and
monosulfur (VS, orange) are shown on the panel to the right. The
calculated ADF intensity is noticeably decreased by the presence of
monoselenium vacancy (VSe), which is close to that of disulfur (S2).
(b) Calculated STEM-ADF intensity ratios of (S + Se)/W and S2/W
at different tilting angles.
Figure 6. (a) High-resolution ADF-STEM image highlighting a region
of a WS0.88Se1.12 nanostructure, showing few-layer, bilayer, and
monolayer domains at the edge. (b) Atomically resolved ADF-
STEM image of WS2ySe2(1−y) nanostructures. (c) Experimental ADF
intensity curve (gray) corresponding to indicated by the green arrow
in (b). The atoms are identified according to the simulated profiles
(colored).
mode at 356.5 cm−1 deconvoluted by multipeak Lorentzian
fitting, as well as the first-order LA(M) mode at 174.5 cm−1
(Figure S11). Under excitation of a 514.5 nm laser, the Raman
intensity ratio between the 2LA(M) mode and the A1g mode
increases with thinner samples as the double resonance is only
active for monolayer WS2.46 The fingerprint ratio in our case is
around 0.7, which is between 2.2 (monolayer WS2) and 0.47
(bulk WS2), corroborating the few-layer structure observed by
electron microscopy. The Raman peaks shift progressively upon
incorporating Se. Using WS0.88Se1.12 as a representative
example, the A1g mode red shifts to 405.2 cm−1 and the peak
from the E12g and 2LA(M) modes remains around 353 cm−1
(Figure S12). The WSe2 feature of degenerate A1g and E21g
modes red shifts from 261.3 cm−1 for the y = 0.81 sample
(WS1.62Se0.38) to 249.7 cm−1 for the y = 0 sample (WSe2). The
shift could result from tensile strain and structural distortions
caused by incorporating the larger-diameter Se vs S atoms.16
Atomic disorder also contributes to the broadening of Raman
peaks for alloy samples, in contrast to the sharper and more
distinct peaks for the WS2 and WSe2 end members.
Figure 8. Raman spectra of the as-prepared WS2ySe2(1−y) nanostruc-
tures, acquired using a 514.5 nm excitation laser. The spectra are
normalized to the peak with the highest intensity for each sample.
Raman modes are labeled based on bulk notation.
Tunable Optical Properties. UV−vis absorption spectra
for all MoxW1−xSe2 and WS2ySe2(1−y) samples dispersed in
ethanol were acquired to characterize the tunable optical
properties achievable through systematically modulating the
metal and chalcogen compositions (Figure 9a). Figure 9b
shows UV−vis absorption spectra for the MoxW1−xSe2 metal
alloy samples. Absorption peaks are labeled according to the
convention proposed by Wilson and Yoffe.47 For MoSe2, A and
B excitonic features arising from the valence band splitting at
G
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
Figure 9. Tunable optical properties for colloidally synthesized TMD alloys. (a) Photograph of MoxW1−xSe2 and WS2ySe2(1−y) samples suspended in
ethanol. (b) UV−vis absorbance spectra, (c) the A excitonic peaks obtained through Lorentzian peak fitting, and (d) corresponding excitonic
transition energy for MoxW1−xSe2 nanostructures. (e) UV−vis absorbance spectra, (f) the A excitonic peaks obtained through Lorentzian peak
fitting, and (g) corresponding excitonic transition energy for WS2ySe2(1−y) nanostructures. UV−vis spectra are normalized according to the highest-
energy excitonic peak.
the K and K′ points of the Brillouin zone appear at 823 nm
(1.51 eV) and 741 nm (1.67 eV), respectively. The A excitonic
transition is located at 768 nm (1.61 eV) for WSe2, while the
broad feature around 605 nm (2.05 eV) is considered as the
overlap of the B and A′ excitonic transitions, which could be
caused by additional optical transitions and interlayer coupling
for MoS2, WS2, MoSe2, and WSe2 multilayers.48 Higher-lying
interband transitions, including the C peak for MoSe2 and the
B′ peak for WSe2, require further optical measurements and
band structure calculations, and are therefore out of the scope
of our work.49 Lorentzian peak fitting was utilized to determine
the position of the A exciton transition. The curved background
was linearly modified to partially reduce scattering effects for
the samples dispersed in solution (Figure S13).50,51 As
indicated in Figure 9c and d, the A excitonic transition ranges
from 1.61 to 1.51 eV with a downward curve as the percentage
of Mo in MoxW1−xSe2 increases, which could be due to uneven
changes of the conduction band minimum (CBM) and the
valence band maximum (VBM) through alloying.52 The CBM
states at the K point are mainly comprised of out-of-plane dz2
orbitals and localized around Mo with lower energy. In
contrast, the VBM states are mainly composed of in-plane dxy
2
2
and dx −y orbitals, which introduce a uniform distribution of
VBM states with wave function delocalization. As a
consequence, the CBM drops rapidly as Mo is incorporated
into WSe2, while VBM decreases linearly, leading to the
observed bowing toward the Mo-rich side of MoxW1−xSe2.
As displayed in Figure 9a, a gradient color change from
brown to yellow is observed for the WS2ySe2(1−y) series. The
UV−vis spectra in Figure 9e reveal that the A excitonic
transition exhibits a noticeable red shift with increasing Se
content. The A exciton peak, which changes more significantly
with composition for the chalcogen alloy system relative to the
metal alloy, varies from 642 to 768 nm, corresponding to 1.93−
1.61 eV, as Se incorporation increases (Figure 9f and g). As Se
gradually replaces S, the bonding between the metal and the
chalcogen becomes more covalent in nature, leading to
broadening of the VBM dz2 orbital and thus a smaller energy
gap at the K point.16 In addition, the gap variation of the A
excitonic transition for the chalcogen alloys is more uniform
H
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
Notes
and closer to linearity compared with that for MoxW1−xSe2. The
smaller bowing effect for WS2ySe2(1−y) relative to MoxW1−xSe2 is
attributed to larger lattice mismatch and chemical potential
difference between diselenides and disulfides.43
The authors declare no competing financial interest.
ACKNOWLEDGMENTS
■
Y.S. and R.E.S. were supported by the U.S. National Science
Foundation Grant No. DMR-1607135. J.S.M. was supported by
funds from Penn State University. K.F., Z.L., Y.L., and M.T.
acknowledge support from the U.S. Army Research Office
MURI Grant W911NF-11-1-0362, the Air Force Office of
Scientific Research MURI FA9550-12-1-0471, as well as the
National Science Foundation for 2DARE-EFRI-1433311 and
2DARE-EFRI-1542707. M.T. and R.S. also acknowledge the
Center for 2-Dimensional and Layered Materials at the
Pennsylvania State University. Electron microscopy was
performed at the Electron Microscopy Facility at the Materials
Characterization Lab of the Penn State Materials Research
Institute. We thank Jennifer Grey and Ke Wang for the
assistance with TEM characterization.
CONCLUSIONS
■
Given the broad scope of potential applications for colloidal
TMD alloys with tunable compositions, it is important to
identify synthetic routes capable of achieving complete metal
and chalcogen solid solutions, understand how the elements are
distributed throughout the nanostructures, identify the types of
vacancies that are present, and correlate compositional
modulation with tunable properties. Here, TMD alloy
nanostructures spanning the entire MoSe2−WSe2 and WS2−
WSe2 solid solutions have been synthesized directly in solution,
without the need for additional high-temperature annealing.
Chemical guidelines, complementary to existing gas-phase
synthesize methods, have been identified for achieving full
compositional modulation in both metal [MoxW1−xSe2] and
chalcogen [WS2ySe2(1−y)] TMD alloys. Alloying was confirmed
by multiple characterization techniques, and composition-
dependent excitonic transitions were observed. Solution-
synthesized colloidal TMD alloys, which are of interest for
their high surface areas, substrate-free nature, and higher mass
yields, are more challenging to characterize at the atomic level
relative to substrate-confined films made by gas-phase
deposition methods because of the buckled, crumpled, and
semiaggregated nature of the colloidal nanostructures. Using
ADF-STEM coupled with analysis of the tilt-angle dependency
of the ADF intensity, important information about the atomic
distributions and types of vacancies was obtained, bridging the
gap between the knowledge that is already available for
substrate-bound TMD alloy films and the largely unknown
atomic-level details of colloidal analogues. Such fundamental
insights are significant for advancing the emerging applications
of TMD nanostructures in the fields of catalysis, gas sensing,
and optics. Additionally, the solution synthesis capabilities
demonstrated here may in the future permit the formation of
ordered TMD alloys, which have been computationally
predicted to have a small energetic preference to random
substitutional alloys and therefore may be accessible using such
low-temperature methods.53,54 The versatility demonstrated for
tuning both the metal and chalcogen compositions suggests
that Se−Te ternary TMD alloys and quaternary TMD alloys,
such as MoxW1−xS2ySe2(1−y), may also be accessible.
REFERENCES
■
(1) Chhowalla, M.; Shin, H. S.; Eda, G.; Li, L.-J.; Loh, K. P.; Zhang,
H. Nat. Chem. 2013, 5, 263−275.
(2) Lin, Z.; McCreary, A.; Briggs, N.; Subramanian, S.; Zhang, K.;
Sun, Y.; Li, X.; Borys, N. J.; Yuan, H.; Fullerton-Shirey, S. K.;
Chernikov, A.; Zhao, H.; McDonnell, S.; Lindenberg, A. M.; Xiao, K.;
́
LeRoy, B. J.; Drndic, M.; Hwang, J. C. M.; Park, J.; Chhowalla, M.;
Schaak, R. E.; Javey, A.; Hersam, M. C.; Robinson, J.; Terrones, M. 2D
Mater. 2016, 3, 042001.
(3) Voiry, D.; Mohite, A.; Chhowalla, M. Chem. Soc. Rev. 2015, 44,
2702−2712.
(4) Lv, R.; Robinson, J. A.; Schaak, R. E.; Sun, D.; Sun, Y.; Mallouk,
T. E.; Terrones, M. Acc. Chem. Res. 2015, 48, 56−64.
(5) Wang, Q. H.; Kalantar-Zadeh, K.; Kis, A.; Coleman, J. N.; Strano,
M. S. Nat. Nanotechnol. 2012, 7, 699−712.
(6) Lin, Z.; Carvalho, B. R.; Kahn, E.; Lv, R.; Rao, R.; Terrones, H.;
Pimenta, M. A.; Terrones, M. 2D Mater. 2016, 3, 022002.
(7) Conley, H. J.; Wang, B.; Ziegler, J. I.; Haglund, R. F.; Pantelides,
S. T.; Bolotin, K. I. Nano Lett. 2013, 13, 3626−3630.
(8) Voiry, D.; Goswami, A.; Kappera, R.; Silva, C. C. C.; Kaplan, D.;
Fujita, T.; Chen, M.; Asefa, T.; Chhowalla, M. Nat. Chem. 2015, 7,
45−49.
(9) Gong, Y.; Lin, J.; Wang, X.; Shi, G.; Lei, S.; Lin, Z.; Zou, X.; Ye,
G.; Vajtai, R.; Yakobson, B. I.; Terrones, H.; Terrones, M.; Tay, B. K.;
Lou, J.; Pantelides, S. T.; Liu, Z.; Zhou, W.; Ajayan, P. M. Nat. Mater.
2014, 13, 1135−1142.
(10) Feng, Q.; Mao, N.; Wu, J.; Xu, H.; Wang, C.; Zhang, J.; Xie, L.
ACS Nano 2015, 9, 7450−7455.
(11) (a) Lin, Z.; Thee, M. T.; Elías, A. L.; Feng, S.; Zhou, C.;
́
Fujisawa, K.; Perea-Lopez, N.; Carozo, V.; Terrones, H.; Terrones, M.
ASSOCIATED CONTENT
APL Mater. 2014, 2, 092514. (b) Li, H.; Zhang, Q.; Duan, X.; Wu, X.;
Fan, X.; Zhu, X.; Zhuang, X.; Hu, W.; Zhou, H.; Pan, A.; Duan, X. J.
Am. Chem. Soc. 2015, 137, 5284−5287. (c) Li, H.; Wu, X.; Liu, H.;
Zheng, B.; Zhang, Q.; Zhu, X.; Wei, Z.; Zhuang, X.; Zhou, H.; Tang,
W.; Duan, X.; Pan, A. ACS Nano 2017, 11, 961−967.
■
S
* Supporting Information
The Supporting Information is available free of charge on the
ACS Publications website at DOI: 10.1021/jacs.7b04443.
(12) Chen, Y.; Xi, J.; Dumcenco, D. O.; Liu, Z.; Suenaga, K.; Wang,
D.; Shuai, Z.; Huang, Y.-S.; Xie, L. ACS Nano 2013, 7, 4610−4616.
(13) Su, S.-H.; Hsu, Y.-Te; Chang, Y.-H.; Chiu, M.-H.; Hsu, C.-L.;
Hsu, W.-T.; Chang, W.-H.; He, J.-H.; Li, L.-J. Small 2014, 10, 2589−
2594.
Additional synthetic details and characterization data and
analyses (XRD, EDS, TEM, HAADF-STEM, Raman, and
UV−vis) (PDF)
(14) Li, X.; Lin, M.-W.; Basile, L.; Hus, S. M.; Puretzky, A. A.; Lee, J.;
Kuo, Y.-C.; Chang, L.-Y.; Wang, K.; Idrobo, J. C.; Li, A.-P.; Chen, C.-
H.; Rouleau, C. M.; Geohegan, D. B.; Xiao, K. Adv. Mater. 2016, 28,
8240−8247.
(15) Gong, Y.; Liu, Z.; Lupini, A. R.; Shi, G.; Lin, J.; Najmaei, S.; Lin,
Z.; Elías, A. L.; Berkdemir, A.; You, G.; Terrones, H.; Terrones, M.;
Vajtai, R.; Pantelides, S. T.; Pennycook, S. J.; Lou, J.; Zhou, W.;
Ajayan, P. M. Nano Lett. 2014, 14, 442−449.
AUTHOR INFORMATION
■
Corresponding Authors
*mut11@psu.edu
*res20@psu.edu
ORCID
Raymond E. Schaak: 0000-0002-7468-8181
I
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX

Journal of the American Chemical Society
Article
(16) Fu, Q.; Yang, L.; Wang, W.; Han, A.; Huang, J.; Du, P.; Fan, Z.;
Zhang, J.; Xiang, B. Adv. Mater. 2015, 27, 4732−4738.
(17) Belopolski, I.; Sanchez, D. S.; Ishida, Y.; Pan, X.; Yu, P.; Xu, S.-
Y.; Chang, G.; Chang, T.-R.; Zheng, H.; Alidoust, N.; Bian, G.;
Neupane, M.; Huang, S.-M.; Lee, C.-C.; Song, Y.; Bu, H.; Wang, G.;
Li, S.; Eda, G.; Jeng, H.-T.; Kondo, T.; Lin, H.; Liu, Z.; Song, F.; Shin,
S.; Hasan, M. Z. Nat. Commun. 2016, 7, 13643.
(18) (a) Li, H.; Duan, X.; Wu, X.; Zhuang, X.; Zhou, H.; Zhang, Q.;
Zhu, X.; Hu, W.; Ren, P.; Guo, P.; Ma, L.; Fan, X.; Wang, X.; Xu, J.;
Pan, A.; Duan, X. J. Am. Chem. Soc. 2014, 136, 3756−3759. (b) Wu,
X.; Li, H.; Liu, H.; Zhuang, X.; Wang, X.; Fan, X.; Duan, X.; Zhu, X.;
Zhang, Q.; Meixner, A. J.; Duan, X.; Pan, A. Nanoscale 2017, 9, 4707−
4712.
L.; Mallouk, T. E.; Balicas, L.; Pimenta, M. A.; Terrones, M. Sci. Rep.
2014, 4, 4215.
(43) Kang, J.; Tongay, S.; Li, J.; Wu, J. J. Appl. Phys. 2013, 113,
143703.
(44) Xie, L. M. Nanoscale 2015, 7, 18392−18401.
(45) Feng, Q.; Zhu, Y.; Hong, J.; Zhang, M.; Duan, W.; Mao, N.; Wu,
J.; Xu, H.; Dong, F.; Lin, F.; Jin, C.; Wang, C.; Zhang, J.; Xie, L. Adv.
Mater. 2014, 26, 2648−2653.
(46) Berkdemir, A.; Gutier
́
rez, H. R.; Botello-Men
́
dez, A. R.; Perea-
ez-
Lopez, N.; Elías, A. L.; Chia, C.-I.; Wang, B.; Crespi, V. H.; Lop
́
́
Urías, F.; Charlier, J.-C.; Terrones, H.; Terrones, M. Sci. Rep. 2013, 3,
1755.
(47) Wilson, J. A.; Yoffe, A. D. Adv. Phys. 1969, 18, 193−335.
(48) Li, Y.; Chernikov, A.; Zhang, X.; Rigosi, A.; Hill, H. M.; van der
Zande, A. M.; Chenet, D. A.; Shih, E.-M.; Hone, J.; Heinz, T. F. Phys.
Rev. B: Condens. Matter Mater. Phys. 2014, 90, 205422.
(49) Zhao, W.; Ghorannevis, Z.; Chu, L.; Toh, M.; Kloc, C.; Tan, P.-
H.; Eda, G. ACS Nano 2013, 7, 791−797.
(19) Zhang, M.; Wu, J.; Zhu, Y.; Dumcenco, D. O.; Hong, J.; Mao,
N.; Deng, S.; Chen, Y.; Yang, Y.; Jin, C.; Chaki, S. H.; Huang, Y.-S.;
Zhang, J.; Xie, L. ACS Nano 2014, 8, 7130−7137.
(20) Han, J. H.; Lee, S.; Cheon, J. Chem. Soc. Rev. 2013, 42, 2581−
2591.
(50) Smith, R. J.; King, P. J.; Lotya, M.; Wirtz, C.; Khan, U.; De, S.;
O’Neill, A.; Duesberg, G. S.; Grunlan, J. C.; Moriarty, G.; Chen, J.;
Wang, J.; Minett, A. I.; Nicolosi, V.; Coleman, J. N. Adv. Mater. 2011,
23, 3944−3948.
(51) Fan, X.; Xu, P.; Li, Y. C.; Zhou, D.; Sun, Y.; Nguyen, M. A. T.;
Terrones, M.; Mallouk, T. E. J. Am. Chem. Soc. 2016, 138, 5143−5149.
(52) Tongay, S.; Narang, D. S.; Kang, J.; Fan, W.; Ko, C.; Luce, A. V.;
Wang, K. X.; Suh, J.; Patel, K. D.; Pathak, V. M.; Li, J.; Wu, J. Appl.
Phys. Lett. 2014, 104, 012101.
(21) Meiron, O. E.; Houben, L.; Bar-Sadan, M. RSC Adv. 2015, 5,
88108−88114.
(22) Zhang, J.; Kang, W.; Jiang, M.; You, Y.; Cao, Y.; Ng, T.-W.; Yu,
D. Y. W.; Lee, C.-S.; Xu, J. Nanoscale 2017, 9, 1484−1490.
(23) Xu, C.; Peng, S.; Tan, C.; Ang, H.; Tan, H.; Zhang, H.; Yan, Q.
J. Mater. Chem. A 2014, 2, 5597−5601.
(24) Yang, Y.; Wang, S.; Zhang, J.; Li, H.; Tang, Z.; Wang, X. Inorg.
Chem. Front. 2015, 2, 931−937.
(25) Gong, Q.; Cheng, L.; Liu, C.; Zhang, M.; Feng, Q.; Ye, H.;
Zeng, M.; Xie, L.; Liu, Z.; Li, Y. ACS Catal. 2015, 5, 2213−2219.
(26) Mahler, B.; Hoepfner, V.; Liao, K.; Ozin, G. A. J. Am. Chem. Soc.
2014, 136, 14121−14127.
(53) Gan, L.-Y.; Zhang, Q.; Zhao, Y.-J.; Cheng, Y.; Schwingenschlogl,
̈
U. Sci. Rep. 2014, 4, 6691.
(54) Tan, T. L.; Ng, M.-F.; Eda, G. J. Phys. Chem. C 2016, 120,
2501−2508.
́
(27) Jung, W.; Lee, S.; Yoo, D.; Jeong, S.; Miro, P.; Kuc, A.; Heine,
T.; Cheon, J. J. Am. Chem. Soc. 2015, 137, 7266−7269.
(28) Koch, C. T. Ph.D. Thesis, Arizona State University, Tempe, AZ,
2002.
(29) Huang, B.; Yoon, M.; Sumpter, B. G.; Wei, S.-H.; Liu, F. Phys.
Rev. Lett. 2015, 115, 126806.
(30) James, P. B.; Lavik, M. T. Acta Crystallogr. 1963, 16, 1183.
(31) Schutte, W. J.; De Boer, J. L.; Jellinek, F. J. Solid State Chem.
1987, 70, 207−209.
(32) Sun, Y.; Wang, Y.; Sun, D.; Carvalho, B. R.; Read, C. G.; Lee, C.-
H.; Lin, Z.; Fujisawa, K.; Robinson, J. A.; Crespi, V. H.; Terrones, M.;
Schaak, R. E. Angew. Chem., Int. Ed. 2016, 55, 2830−2834.
(33) Sun, D.; Feng, S.; Terrones, M.; Schaak, R. E. Chem. Mater.
2015, 27, 3167−3175.
(34) Krivanek, O. L.; Chisholm, M. F.; Nicolosi, V.; Pennycook, T. J.;
Corbin, G. J.; Dellby, N.; Murfitt, M. F.; Own, C. S.; Szilagyi, Z. S.;
Oxley, M. P.; Pantelides, S. T.; Pennycook, S. J. Nature 2010, 464,
571−574.
(35) (a) Azizi, A.; Zou, X.; Ercius, P.; Zhang, Z.; Elías, A. L.; Perea-
Lopez, N.; Stone, G.; Terrones, M.; Yakobson, B. I.; Alem, N. Nat.
́
Commun. 2014, 5, 4867. (b) Azizi, A.; Wang, Y.; Stone, G.; Elias, A. L.;
Lin, Z.; Terrones, M.; Crespi, V. H.; Alem, N. Nano Lett. 2017, 17,
2802−2808.
(36) Jaramillo, T. F.; Jørgensen, K. P.; Bonde, J.; Nielsen, J. H.;
Horch, S.; Chorkendorff, I. Science 2007, 317, 100−102.
(37) Huang, X.; Zeng, Z.; Zhang, H. Chem. Soc. Rev. 2013, 42, 1934−
1946.
(38) McDonnell, S.; Addou, R.; Buie, C.; Wallace, R. M.; Hinkle, C.
L. ACS Nano 2014, 8, 2880−2888.
(39) Amani, M.; Lien, D.-H.; Kiriya, D.; Xiao, J.; Azcatl, A.; Noh, J.;
Madhvapathy, S. R.; Addou, R.; KC, S.; Dubey, M.; Cho, K.; Wallace,
R. M.; Lee, S.-C.; He, J.-H.; Ager, J. W.; Zhang, X.; Yablonovitch, E.;
Javey, A. Science 2015, 350, 1065−1068.
(40) Azizi, A.; Wang, Y.; Lin, Z.; Wang, K.; Elias, A. L.; Terrones, M.;
Crespi, V. H.; Alem, N. Nano Lett. 2016, 16, 6982−6987.
(41) Zhang, X.; Qiao, X.-F.; Shi, W.; Wu, J.-B.; Jiang, D.-S.; Tan, P.-
H. Chem. Soc. Rev. 2015, 44, 2757−2785.
(42) Terrones, H.; Del Corro, E.; Feng, S.; Poumirol, J. M.; Rhodes,
D.; Smirnov, D.; Pradhan, N. R.; Lin, Z.; Nguyen, M. A. T.; Elías, A.
J
DOI: 10.1021/jacs.7b04443
J. Am. Chem. Soc. XXXX, XXX, XXX−XXX