UC Riverside UC Riverside Electronic Theses and Dissertations
Title Crystal Growth and Phase Engineering of Two-Dimensional Metal Chalcogenides
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Author Mutlu, Zafer
Publication Date 2016
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UNIVERSITY OF CALIFORNIA RIVERSIDE
Crystal Growth and Phase Engineering of Two-Dimensional Metal Chalcogenides
A Dissertation submitted in partial satisfaction of the requirements for the degree of
Doctor of Philosophy
in
Materials Science and Engineering
by
Zafer Mutlu
December 2016
Dissertation Committee: Dr. Cengiz S. Ozkan, Chairperson Dr. Sandeep Kumar Dr. Elaine D. Haberer Dr. Mihri Ozkan
Copyright by Zafer Mutlu 2016
The Dissertation of Zafer Mutlu is approved:
______
Committee Chairperson
University of California, Riverside
Acknowledgments
I would like to express my sincerest gratitude to all individuals who have assisted me morally, intellectually, or financially throughout my Ph.D.
I would first like to show my gratitude for my advisor and co-advisor Dr. Cengiz
S. Ozkan and Dr. Mihrimah Ozkan, respectively, for their guidance and support.
I would also like to thank my committee members Dr. Elaine D, Haberer and Dr.
Sandeep Kumar for taking the time to serve during my Ph.D. exams and providing me valuable feedback insights.
I am grateful for the opportunity I had to collaborate and discuss research with Dr.
Roger K. Lake, Dr. Darshana Wickramaratne, Dr. Andre K. Mkhoyan, Ryan J. Wu, Dr.
Sefaattin Tongay, Anupum Pant and Dr. Ilkeun Lee.
I thank all of my colleagues who have helped me during my Ph.D. in any way possible.
I would like to thank my parents for their undying support emotionally and financially.
Last but not least, I would like to thank my wife Gul for her passionate love, support, encouragement, and patience.
The text and figures in this dissertation, in part or in full, are a reprint of the material as it appears in the following publications:
Zafer Mutlu, Ryan J. Wu, Darshana Wickramaratne, Sina Shahrezaei, Chueh
Liu, Selcuk Temiz, Andrew Patalano, Mihrimah Ozkan, Roger K. Lake, K. A.
iv Mkhoyan, Cengiz S. Ozkan, “Phase Engineering of 2D Tin Sulfides”, Small,
12(22), pp.2998-3004, 2016. Reprinted with permission from the WILEY-VCH.
Zafer Mutlu, Mihrimah Ozkan, Cengiz S. Ozkan, “Large-Area Synthesis,
Characterization, and Anisotropic Etching of Two-Dimensional Tungsten
Disulfide Films”, Materials Chemistry and Physics, 176, pp.52-57, 2016.
Reprinted with permission from the ELSEVIER.
Zafer Mutlu, Sina Shahrezaei, Selcuk Temiz, Mihrimah Ozkan, Cengiz S Ozkan,
“Facile Synthesis and Characterization of Two-Dimensional Layered Tin
Disulfide Nanowalls”, Journal of Electronic Materials, 45(4), pp.2115, 2016.
Reprinted with permission from the SPRINGER.
Zafer Mutlu, Darshana Wickramaratne, Serol Turkyilmaz, Hamed H. Bay,
Zachary J. Favors, Mihri Ozkan, Roger K. Lake, Cengiz S. Ozkan, “Two-
Dimensional Layered Semiconductor Tungsten Disulfide and Molybdenum-
Tungsten Disulfide: Synthesis, Materials Properties and Electronic Structure”,
Journal of Nanoscience and Nanotechnology, 16(8), pp.8419-8423, 2016.
Reprinted with permission from the ASP.
Zafer Mutlu, Darshana Wickramaratne, Hamed H. Bay, Zachary J. Favors,
Mihrimah Ozkan, Roger Lake, Cengiz S Ozkan, “Synthesis, Characterization, and
Electronic Structure of Few‐Layer MoSe2 Granular Films”, Physica Status Solidi
(a), 2011(12), pp.2671-2676, 2014. Reprinted with permission from the WILEY-
VCH.
v Anupum Pant†, Zafer Mutlu†, Darshana Wickramaratne, Hui Cai, Roger K. Lake,
Cengiz S. Ozkan, Sefaattin Tongay. “Fundamentals of Lateral and Vertical
Heterojunctions of Atomically Thin Materials”, Nanoscale, 8(7), pp.3870-3887,
2016 († These authors have contributed equally to this work). Reprinted with
permission from the RSC.
The co-author Cengiz S. Ozkan Lake, listed in the above publications directed and supervised the research which forms the basis for this dissertation. The remaining co- authors listed provided technical expertise and support as collaborators.
vi ABSTRACT OF THE DISSERTATION
Crystal Growth and Phase Engineering of Two-Dimensional Metal Chalcogenides
by
Zafer Mutlu
Doctor of Philosophy, Graduate Program in Materials Science and Engineering University of California, Riverside, December 2016 Dr. Cengiz S. Ozkan, Chairperson
The ability to grow crystals of desired dimensions is an important aspect of the design of new functional materials. The crystals grown in two dimensions enable successful and efficient modification of a wide range of physical, electronic and structural properties tailored for smaller and more efficient electronics, optoelectronics, valleytronics, and spintronic applications. The control of phases in crystals is also an additional element in the design of functional materials. Two-dimensional (2D) atomic crystals can crystallize in a variety of crystal phases, which possess distinct properties.
Therefore, identifying approaches for controlling dimension and crystal phase is essential for tuning the properties of 2D materials to specific applications.
Herein, various unique methods of producing 2D atomic layered metal chalcogenide crystals (LMCs) including growth of group VI transition metal dichalcogenides (TMDs) by chalcogenization of pre-deposited metal-containing precursors, and phase-engineered growth of group IV LMCs by vapor-phase reaction of metal oxide and chalcogen powders are demonstrated. Scanning transmission electron
vii microscopy (STEM), scanning electron microscopy (SEM), atomic force microscopy
(AFM), optical microscopy, Raman spectroscopy, photoluminescence (PL) spectroscopy, ultraviolet photoelectron spectroscopy (UPS), X-ray photoemission spectroscopy (XPS), and X-ray diffraction (XRD) are all notable experimental techniques used to characterize these various materials. The experimental results are corroborated by density functional theory (DFT) calculations.
viii Contents Chapter 1 Introduction...... 1 1.1 Brief properties of 2D LMCs ...... 3 1.2 Synthesis...... 12 1.3 Characterization Methods ...... 20 1.4 Applications ...... 22 1.5 Grand Challenges and Opportunities ...... 23 1.6 Thesis Outline ...... 24 References ...... 26
Chapter 2 Seed-Assisted CVD Synthesis of 2D WS2 and MoS2-WS2 ...... 41 Abstract ...... 41 2.1 Introduction ...... 41 2.2 Experimental Section ...... 43 2.2.1 Materials Synthesis ...... 43 2.2.2 Materials Characterization ...... 44 2.2.3 DFT Calculations ...... 44 2.3 Results and Discussion ...... 45 2.4 Conclusion ...... 54 References ...... 54
Chapter 3 Large-Area Synthesis and Anisotropic Etching of 2D WS2 ...... 58 Abstract ...... 58 3.1 Introduction ...... 58 3.2 Experimental Section ...... 61 3.2.1 Materials Synthesis ...... 61 3.2.2 Materials Characterization ...... 62 3.3 Results and Discussion ...... 62 3.4 Conclusion ...... 70 References ...... 71 Chapter 4 Phase Engineering of 2D Tin Sulfides ...... 74 Abstract ...... 74 4.1 Introduction ...... 74 4.2 Experimental Section ...... 76
ix 4.2.1 Materials Synthesis ...... 76 4.2.2 Materials Characterization ...... 77 4.2.3 DFT Calculations ...... 78 4.3 Results and Discussion ...... 79 4.4 Conclusion ...... 89 References ...... 90
Chapter 5 Facile Synthesis of 2D SnS2 Nanowalls ...... 94 Abstract ...... 94 5.1. Introduction ...... 94 5.2 Experimental Section ...... 96 5.2.1 Materials Synthesis ...... 96 5.2.2 Materials Characterization ...... 97 5.3 Results and Discussion ...... 98 5.4 Conclusion ...... 103 References ...... 103
Chapter 6 Synthesis of MoSe2 via Rapid Thermal Processing and Laser Annealing ...... 106 Abstract ...... 106 6.1 Introduction ...... 106 6.2 Experimental Section ...... 109 6.2.1 Materials Synthesis ...... 109 6.2.2 Materials Characterization ...... 110 6.2.3 DFT calculations ...... 111 6.3 Results and Discussion ...... 111 6.4 Conclusion ...... 121 References ...... 122 Conclusion……………………………………………………………………………..125
x List of Figures
Figure 2.1 The morphology and Raman spectra of the WO3 crystals………………… 46
Figure 2.2 The morphology of the WS2 crystals……………………………………….48
Figure 2.3 AFM thickness analysis of the WS2 crystal………………………………...49
Figure 2.4 Raman-PL spectra and DFT-calculated electronic band structure of the WS2 crystals of different thickness……………………………………………………………51
Figure 2.5 Raman imaging of the WS2 crystals of different morphologies…………….52
Figure 2.6 Raman-PL spectroscopy characterization and DFT-calculated band-gap of the crystals…………………………………………………………………………………...53
Figure 3.1 AFM analysis of the as-deposited and annealed W films.
Figure 3.2 XPS analysis of the W and WS2 films.
Figure 3.3 The morphology of the WS2 films.
Figure 3.4 Raman-PL spectroscopy characterization of the WS2 films.
Figure 3.5 AFM analysis of the triangular pits formed on the WS2 films by air- annealing.
Figure 4.1 The growth set-up, morphology and atomic structure of the SnS2 and SnS crystals.
Figure 4.2 The photo, SEM and optical images of the crystals obtained by various recipes.
Figure 4.3 Raman spectroscopy characterization of the SnS2 and SnS crystals.
Figure 4.4 The ionization potentials for the SnS2 and SnS crystals.
xi Figure 5.1 The growth set-up, thermal process and growth process for the SnS2 NWs.
Figure 5.2 The morphology of the SnS2 NWs.
Figure 5.3 XRD and Raman spectroscopy analysis of the SnS2 NWs.
Figure 5.4 XPS and UPS analysis of the SnS2 NWs.
Figure 6.1 The ball-and-stick model and synthesis process for MoSe2.
Figure 6.2 The AFM height images of the MoSe2 films
Figure 6.3 The histograms of the particle size distribution of the MoSe2 films.
Figure 6.4 AFM phase images of the MoSe2 films
Figure 6.5 The optical images of the Mo/Se/Mo and MoSe2 films.
Figure 6.6 XPS analysis of the MoSe2 films.
Figure 6.7 Raman-PL spectroscopy characterization of the MoSe2 film.
Figure 6.8 DFT-calculated electronic band structures of MoSe2.
Figure 6.9 The RLA setup and laser-power dependent Raman spectra.
Figure 6.10 The formatıon of MoSe2 from SELs with certain patterns on SiO2/Si.
xii List of Tables
Table 1.1. Electronic and material properties of different LMCs. The bandgaps for the monolayer structures are from ab-initio theoretical (T) results, the bulk bandgaps and lattice constants are from experimental data (E). Cited references are included in square parentheses.
Table 4.1. Comparison of experimentally measured and ab-initio DFT calculated Raman spectra peak positions in SnS2 and SnS in cm-1.
xiii Chapter 1 Introduction
Layered metal chalcogenides (LMCs) belong to a broader class of van-der-Waal
(vdW) materials that have been the subject of intense research. Examples of vdW materials include graphene [1], hexagonal boron nitride (h-BN) [2], transition metal dichalcogenides (TMDs) [3], metal halides [4], phosphorene [5], germanene [6], silicene
[7] and transition metal carbides and carbonitrides (MXenes) [8]. These materials and their heterostructures have been used to investigate a wide range of devices including field-effect transistors (FETs) [9], PN junctions [10], photodetectors [11], supercapacitors
[12], spin FETs [13], photovoltaics [14], thermoelectrics [15]. This family of materials also provides the opportunity for the following physical phenomena to be studied at low dimensions: Dirac fermions [1], valley and spin physics [16], superconductivity [17], commensurate and incommensurate charge density waves (CDW) [18], excitons and trions [19], itinerant magnetism [20], Lifshitz transitions [21] and Mott transitions [22].
The performance of these devices and the physical phenomena studied are quantitatively different from the properties demonstrated by these materials in their bulk form.
The electronic, thermal, optoelectronic and mechanical properties of the bulk
LMCs can change in beneficial ways when they are reduced to a few monolayers or single monolayer in thickness. Examples of such changes include an increase in the photoluminescence (PL) intensity by 3 orders of magnitude when bulk MoS2 is confined to a single monolayer and variations in the thermal conductivity by up to 2 orders of magnitude between the basal and cross-plane axis of conduction. The sensitivity of the properties of LMCs to changes in dimensionality, polytypism, strain and elemental
1 composition and the ability to vertically stack these materials to create heterostructures have resulted in a resurgence in research interest in this family of materials.
Studies on LMC materials can be found dating back to the 1950s. The first report on earth abundant MoS2 dates back to 1953 with intercalation studies on bulk MoS2. [23]
The metallic and semi-metallic LMCs were the focus of several studies in the 1970s with the identification of superconducting and CDW phenomena in these materials. This was complemented by the emergence of electronic structure calculations of a number of LMC compounds which were used to explain the observed phenomena. The first studies of transistor type structures using the LMC were first conducted on the metallic LMCs to investigate CDW phenomena and other correlated phenomena. [24] In 2004, the first top gated thin film WSe2 transistor was fabricated [25], with the proposal of applying LMCs for flexible electronics. This was followed by transport studies on MoS2 and TaS2 nanopatches by Ayari et al. [26] in 2007. The first experimental investigations of transport in monolayer MoS2 were published in 2010 by Radisavljevic et al. [9], and the number of publications on this family of materials has been rapidly increasing ever since.
Changes in the properties of LMCs are driven by qualitative changes in the electronic and phononic dispersion of the materials. Alignment of monolayers with respect to each other and the surrounding dielectric environment can also lead to changes in the electronic, optical and thermal properties. Current research on the LMCs family of compounds is focused on understanding the fundamental properties of these materials and their heterostructures.
2 In this chapter, an overview of some of the experimental and theoretical investigations of these properties, synthesis of LMCs and their heterostructures and possible emerging applications that can be derived from vertically and horizontally stacked vdW compounds is provided.
1.1 Brief properties of 2D LMCs
In their bulk form, these vdW materials are covalently bonded along the basal plane of the monolayer, and the individual monolayers are held together via weak vdW forces perpendicular to the basal plane.
Each monolayer can be stacked in a number of different configurations in the
LMC compounds. For each of these stacking configurations, the monolayers cleave along the plane of the chalcogen atom. Each monolayer is hexagonally packed. A single metal chalcogenide monolayer is approximately 3Å to 5Å thick and the vdW gap between the adjacent planes of chalcogen atoms varies between 3Å to 4Å. The electronegativity of the chalcogen atom in a LMC compound and its coordination (octahedral vs. trigonal prismatic) determines the metal-chalcogen bond lengths. Prior studies have shown the octahedral coordination is preferred when intra-layer bonding is ionic, which leads to maximum separation distance between the negatively charged chalcogens. When intra- layer bonding is more covalent, the trigonal prismatic coordination is preferred which leads to maximum overlap between the metal and chalcogen atom. [27] Each metal chalcogenide has a preferred ground state stacking order. The layered nature of these compounds means alternate stacking orders can be achieved by overcoming an energy
3 barrier that is unique to each material, either through charge transfer intercalation or controlled synthesis methods.
The electrical and optical properties of LMCs cover a wide spectrum based on the coordination of the metal atom and the filling of the d-orbital. Trigonal prismatic coordination of the transition metal atoms in LMC compounds result in degenerate a1
2 2 (dz ), e (dx -y2, xy) and e' (dxz, dyz) orbitals with a gap between the a1, e and e' group of
2 2 2 orbitals. Octahedral coordination of the metal atom results in degenerate eg (dz , x -y ) and t2g (dyz, xz, xy) orbitals. When the d-orbital of the LMC compound is fully occupied, the material is a semiconductor. Partial occupation of the d-orbital leads to either semi- metallic or metallic behavior. The electronic and structural properties of the LMCs have been summarized in Table 1.1.
Research into monolayer and few-layer MoS2 and other LMCs was popularized by the first demonstrations of the indirect to direct gap transition in MoS2 through the use of PL experiments by Mak et al. [28] and Splendiani et al. [29]. The band gap of bulk
MoS2 increases from 1.29 eV to 1.8 eV when thinned down to a single monolayer. This phenomenon is explained by the orbital composition of the different high symmetry points in the group VI TMDs. The large contribution of the pz orbitals of the S atoms at the Гv valley results in the largest interlayer coupling greatest sensitivity to the presence
2 of adjacent layers. The Kv valley, with no dz or pz components has very weak interlayer coupling and, therefore, is insensitive to the presence of adjacent layers. When two monolayers are brought together, the Гv valleys of the two layers couple and split. In bulk
MoS2, the interlayer distance is approximately 3.12 Å and the energy splitting at Гv is
4 approximately 600 meV. When two group VI dichalcogenide monolayers are brought into close proximity the interlayer coupling causes the Гv valley to rise above the Kv valley which drives the direct to indirect gap transition when the film thickness increases above a monolayer. Recent angle-resolved photoemission spectroscopy (ARPES) measurements on MoS2, MoSe2 and WSe2 [30] [31] [32] have confirmed this shift of the valence band at Г as the origin of the indirect to direct crossover as a function of film thickness in the group VI TMDs. The sensitivity of these valleys to variations in interlayer coupling and film thickness can lead to changes in degeneracy and barrier heights which can affect the performance of devices formed with heterostructures of the group VI TMDs. Recent ab-initio studies have demonstrated the near degeneracy of the conduction band Kc and Λc valleys as the film thickness increases above four monolayers. [33] The Kc valley is three fold degenerate and the Λc valley is six fold degenerate.
5 Table 1.1: Electronic and material properties of different LMCs. The bandgaps for the monolayer structures are from ab-initio theoretical (T) results, the bulk bandgaps and lattice constants are from experimental data (E). Cited references are included in square parentheses.
Stacking Electronic Lattice Bandgap Order Characteristics Constant (Å) (eV) Group V TMDs: Metal, VS2, VSe2, 1T, 2H superconducting, 3.18 - 3.47 N/A NbS2, NbSe2, CDW [77] TaS2, TaSe2 Group VI TMDs 3.15 - 3.28 Bulk: 1.09 - 1.35 (E) MoS2, MoSe2, 2H, 1T, 3R Semiconducting [78][79][80] [78][79][80] MoTe2, WS2, 1L: 1.77 - 2.88 (T) WSe2 [81] Group III Bulk: 1.20 - 2.20 (E) metal [82] [83][84][85] chalcogenides α, β Semiconducting 3.58 - 4.00 1L: 1.88 - 3.28 (T) GaS, GaSe, [82] [83][84][85] [21] [39][40] GaTe, InS, InSe
There is an increasing effort to identify new LMCs beyond the group IV to VI
TMDs. Examples of new low-dimensional layered chalcogenides that have been synthesized are the group III (gallium, indium) chalcogenides, tin sulfides and selenides.
[34] [35] [36] [37] [38] Confining these materials to few-monolayers and a monolayer leads to the emergence of new phenomena and possible application spaces in which these materials can be adopted.
When the group III sulfides, selenides and tellurides are confined from their bulk structure to a single monolayer the parabolic dispersion of the valence band edge at Г is deformed into a 'Mexican-hat' dispersion. [21] [39] [40] The presence of a 'Mexican-hat'
6 dispersion results in a Fermi-ring of states which can lead to novel phenomena including a 1/√E singularity in the two-dimensional (2D) density of states. [41] [42] This singularity in the density of states can lead to novel correlated phenomena such as itinerant magnetism, superconductivity and also have practical implications such as enhanced thermo-electric response [40] [42] [43] and photo response [44].
In addition to obtaining an understanding of the intrinsic properties of the LMC compounds in their monolayer limit, it is also necessary to understand how these properties change when heterostructures are formed between LMCs and other vdW or non-vdW materials. This includes but is not limited to changes in material properties as the film thickness of the material increases and approaches the bulk limit, interfaces formed between different vdW materials and interfaces between LMCs and metal contacts. Each of these interfaces can either have perfect alignment or be rotationally misaligned. Prior studies have shown rotational faults in bilayer graphene or heterostructures between graphene, and other vdW materials can either decouple the two layers [45] [46], increase contact resistance [47] [48], affect thermal conductivity [49] and modulate the intrinsic Fermi velocity of graphene [46][50]. Thus, a similar understanding of interfaces formed between LMC compounds is necessary.
The self-passivated surfaces of the LMCs, in principle, enable atomically sharp interfaces to be engineered. vdW epitaxy was proposed as an approach to achieve atomically sharp quantum well structures between vdW materials that are highly lattice mismatched. [51] Prior photoemission spectroscopy experiments demonstrated the possibility of achieving different band alignments in a number of bulk LMC
7 heterostructures grown using vdW epitaxy. [52] Type-II (staggered) band alignments were shown to form in bulk MoS2/SnSe2, WSe2/SnSe2 SnS2/MoS2. Type-III (broken gap) alignments occur in SnSe2/MoTe2, SnS2/MoTe2 and SnSe2/GaSe. For the broken gap alignments, the valence band (conduction band) offsets for these heterostructures were measured to be 0.94 eV (0.92 eV), 1.72 eV (0.55 eV) and 0.83 eV (1.82 eV) respectively.
A number of theoretical calculations have examined the natural band alignments and offsets between different monolayers of the LMCs. [53] [54] Type-III band alignments were predicted to occur between monolayers of Mo and W selenides and tellurides
(MoTe2, MoSe2, WSe2, WTe2) and zirconium and hafnium sulfides and selenides (ZrS2,
ZrSe2, HfS2, HfSe2). The valence band offset between MoS2 and WS2 is 386 meV, and the conduction band offset is 310 meV. The type-II alignment formed between monolayer MoS2/WS2 was investigated using PL spectroscopy. Future calculations and experimental investigations of band alignments formed between few-layer LMCs will have to account for quantum dipoles, which can lead to alterations in the conduction and valence band offsets. [52] [55] Broken gap alignments are beneficial for next-generation broken gap heterojunction tunnel field effect transistor (TFET) devices that utilize monolayers or few-layers of the LMCs or post-CMOS excitonic devices.
Interfaces between monolayers and few-layers of LMC compounds and the metal contacts used to fabricate vdW FET devices is a major bottleneck that impedes the transfer characteristics in these devices. A number of comprehensive experimental and theoretical studies have recently demonstrated the natural band lineup between the group
VI TMDs and a variety of metal work functions. [56] [57] However, experimental studies
8 have shown that this simplified picture is insufficient to predict the ideal metal contact that would lead to an Ohmic contact with ultra-thin LMC transistors. The ultra-thin geometry of these devices leads to Fermi level pinning which alters the predicted alignment of the metal work function with the conduction and valence band of the LMCs.
Das et al. [58] investigated the n-type transfer characteristics between a 10 nm film of
MoS2 interfaced with metals that have a variety of work functions (WF); Nickel (WF=5.0 eV), Platinum (WF=5.9 eV), Scandium (WF=3.5 eV) and Titanium (WF=4.3 eV). For each of these metals, the Schottky barrier (SB) height was measured to be approximately
150 meV, 230 meV, 30 meV and 50 meV respectively. The n-type transfer characteristics observed with a high work function metal such as nickel and platinum was identified as evidence of Fermi level pinning. [58] Based on these extracted SB heights, scandium was identified as an optimum contact for n-type transport.
The sensitivity of the band edges in the LMC compounds to interlayer interactions also introduces the possibility of tuning the electronic properties of the group VI TMDs using strain, intercalation and engineering layered heterostructures. Using a two-point bending apparatus and PL experiments it was shown uniaxial strain up to 1.51% in bilayer WSe2 induces an indirect to direct gap transition. [59] This was explained with the use of density functional theory (DFT) calculations to be due to the near degeneracy
(50 meV) of the Kc and ∑c valleys of the conduction band. This indirect to direct band gap transition is less likely to occur at similar strain levels in the other group VI TMDs given the larger offset between Kc and the nearest conduction band edge.
9 Applying pressure to the LMCs also results in quantitative changes in the crystal structure and electronic structure. Recent diamond anvil cell experiments on bulk MoS2 demonstrated a structural distortion followed by an insulating to metal transition at a pressure of 19 GPa. [60] This equates to the c/a lattice-parameter ratio of bulk MoS2 decreasing from its equilibrium value of 3.89 to 3.675.
The ability to manipulate the electrical and optical properties of these LMCs through mechanical strain and pressure is possible due to the robust mechanical properties of the LMCs. Measurements of the stiffness and breaking strength of monolayer MoS2 using atomic force microscopy (AFM) demonstrate an in-plane stiffness of 270 +/- 100 GPa and a of 15 +/- 3 Nm-1 which occurs at an effective strain between 6 and 11%. [61] Although these values are lower than the fundamental Young's modulus of
1000 GPa and breaking strength of 130 GPa for monolayer graphene, they are superior in value to commonly used substrates for flexible electronics such as polymide and polydimethylsiloxane (PDMS). This would suggest LMCs are suitable candidates for flexible electronics. [62]
Future studies on such heterostructures would have to take into account effects such as lattice mismatch which can lead to misorientation between the two atomic planes and heterostructures formed with few-layer films of the LMCs. The electronic and vibrational properties of the LMCs are sensitive to proximity of adjacent monolayers. In addition to strain, pressure and interlayer spacing, misorientation of one monolayer with respect to another can affect the fundamental properties of these materials. Given the surge of interest in heterostructures formed with dissimilar or similar layered materials,
10 such devices are often formed by layer by layer growth or mechanical stacking. These approaches are known to lead to naturally misoriented interfaces.
Recent experiments have studied changes in the electronic and vibrational properties of 'folded' and misoriented bilayer MoS2. [63][64][65] The electronic structure of the misoriented bilayer remains indirect with no evidence of the individual layers being completely decoupled. Instead, the indirect bandgap between the twisted bilayers was shown by PL measurements to increase as a function of twist angle by up to 140 meV. Steric effects between the adjacent layers of S atoms are predicted by DFT calculations to lead to repulsion between the individual monolayers and increase the interlayer separation. This in turn leads to a downward shift in the valence band at Г and an increase in the indirect bandgap. Since the orbital composition and location of the band extrema are the same for the bilayers of the other group VI TMDs, similar behavior of the electronic structure with rotation can be expected.
In addition, the electronic, optical and mechanical properties of the LMCs, the intrinsic vibrational and thermal conductivity properties of these materials are being studied. [66][67][68] The in-plane thermal conductivity of suspended and supported monolayer and few-layer MoS2 was measured experimentally using Raman thermometry technique. The thermal conductivity of suspended and supported MoS2 is 34.5 W/mK and increases to 52 W/mK for the few-layer sample. This is in good agreement with theoretical predictions of the thermal conductivity in free-standing monolayers of MoS2.
Anisotropy in the intra-layer and inter-layer bonding of the LMCs also leads to anisotropic thermal conductivity properties. The cross-plane thermal conductivity of the
11 layered TMDs was experimentally measured to be approximately 1 W/mK, which was attributed to the weak vdW bonds perpendicular to the strong covalent intra-layer bonding.
A number of LMCs also exhibit competing instabilities which can lead to new collective ground states. Examples of such instabilities include ferromagnetic, superconductivity and charge density wave phases. Although no experimental demonstrations of ferromagnetism have yet been observed in the LMCs, a number of theoretical predictions have been put forth for emergent ferromagnetic order either through dilute magnetic doping of LMCs or in electrostatically doped III-VI compounds that fulfill the Stoner criteria for magnetism. [69][70][71] The group V dichalcogenides
(TaS2, TaSe2, NbS2, NbSe2, VSe2) and TiSe2 have complex phase diagrams that undergo either charge density wave and/or superconducting phase transitions. [18] Recent studies have begun to explore the effect of dimensionality on these phase diagrams.
[72][73][74][75][76] A number of studies have shown the charge density wave transition temperatures can be tuned as a function film thickness, itinerant ferromagnetism can be induced by electron doping monolayer and few-layer LMCs and competing phase transitions can be suppressed.
1.2 Synthesis
To access the properties of the LMCs discussed in Section 1.1 and enable their integration in devices requires methods for controllable synthesis of atomic thick LMCs of both large-area and high quality. In this section, the recent progress made in the synthesis of the LMCs, e.g., MoS2, MoSe2, WS2, WSe2, GaS, GaSe, InSe, and SnS2 and
12 their heterostructures with vertical and lateral interfaces, e.g., WS2/MoS2, and
WSe2/MoSe2 is overviewed.
Various methods have been used to synthesize LMCs. The first isolation of monolayer MoS2, was successively reported by Mak et al. [28] using scotch tape, otherwise referred to as mechanical exfoliation (ME). This led to the isolation of other monolayer to few-layer LMCs e.g., MoSe2 [86], WS2 [87], WSe2 [88], GaS [89] [90],
GaSe [34], and InSe [91] [92], which were isolated based from bulk counterparts.
Although mechanical exfoliation is impractical [93] [94] for large-scale applications of
LMCs due to low throughput and lack of control over layer thickness, it is still the method of choice to explore and prototype novel properties and new device concepts.
Liquid-phase exfoliation which involves ion intercalation or solvent-based exfoliation, is an alternative to mechanical exfoliation. It has been also used to produce a variety of
LMCs. [95] [96] [97] The liquid phase exfoliation is promising for integration of LMCs into large-scale solution-based or printable electronics. [95]
In addition to mechanical or liquid exfoliation, a few unconventional techniques such as laser thinning [98], thermal annealing [99] [100] [101], gas-assisted etching [102] and plasma etching [103] have been reported to produce 2D LMCs. Castellanos-Gomez et al. [98] demonstrated synthesis of monolayer MoS2 from bulk MoS2 by laser sublimation of the upper layers via heat induced laser absorption. Monolayer or few-layer
MoS2 films were obtained by heating multilayer MoS2 in air [99], oxygen [100] or in an argon environment [101], which led to the top-most layers being etched. Chemical dry- etching of MoS2 multilayer flakes down to monolayer was also reported by using XeF2 as
13 a gaseous reactant. [102] Using a plasma etching technique, Liu et al. [103] demonstrated etching of thick MoS2 sheets down to monolayer with lower surface roughness under Ar+ plasma irradiation. The aforementioned methods are likely to provide novel insights into engineering the geometry of LMCs for specific applications. However, these techniques are still inhibited defect formation, incomplete sublimation or pinhole creation. [104]
[105]
Apart from the above-mentioned synthetic techniques, very recently, chemical vapor deposition (CVD) techniques have shown great promise to produce high-quality
2D LMCs with controllable flake size and thickness. Several research groups have reported CVD growth of atomically thin 2D LMCs, with the focus on MoS2, MoSe2, WS2 and WSe2. For example, typical CVD growth of MoS2 involves sulfurization/decomposition of their corresponding pre-deposited thin metal films [106]
[107] [108] [109] [110] and vapor phase reaction/deposition of gaseous metal and chalcogen feedstocks [111] [112] [113] [114] [115] [116] [117]. WS2 layers have also been grown by sulfurization of pre-deposited W [110] [118] or WO3 [119] [120] films, and reaction of elemental S and W-based powders [121] [122] [123] [124]. CVD has also been explored using Se powder (or pellet) and MoO3 powder as the reactants, yielding monolayer to few-layer triangular MoSe2 islands. [125] [126] [127] [128] The recent demonstration [129] of the p-type transistor behaviors in exfoliated monolayer WSe2 flakes has triggered extensive interest on the production of large-scale intrinsically p-type monolayer LMCs for realizing complementary digital logic applications. Monolayer
14 WSe2 flakes have been mainly produced by gas-phase reaction of WO3 and Se powders
[130] [131] or vapour phase transport of WSe2 powder [132] [133] [134].
GaSe is an emerging 2D material with novel electronic and optoelectronic properties discussed in Section 1.1. Recent reports have demonstrated the possibility of the synthesis of large-scale monolayer and few-layer GaSe flakes by using CVD-based approaches. Lei et al. [35] showed the growth large-scale few-layer GaSe crystals with the thickness down to 2 nm on SiO2/Si by a vapor phase transport method using GaSe powders as precursor and GaSe flakes as seeds. By using a vdW method, Zhou et al.
[135] synthesized monolayer and few-layer GaSe p-type nanoplates, with the lateral size of up to tens of micrometers on a flexible mica substrate. Most recently, Li et al. [136] reported on the synthesis of monolayer and thicker 2D GaSe crystals, with a lateral size up to ~ 60 μm, on SiO2/Si substrates using GaSe crystals and Ga2Se3 powder as source materials.
Lower temperature processes such as molecular beam epitaxy (MBE) and atomic layer deposition (ALD) are possible alternatives to CVD. MBE at very low vacuum and controlled deposition rates can produce epitaxial LMCs [31] [137]. Recently, Zhang et al.
[31] has successively grown epitaxial MoSe2 films with various thickness by MBE and studied the transition from indirect to direct band gap in monolayer films by using
ARPES. In addition to MBE, ALD has attracted increasing interest in the growth of large-scale ultrathin LMC materials. Song et al. [138] demonstrated the wafer-scale synthesis of WS2 nanosheets on SiO2 substrates via sulfurization of WO3 thin films by using plasma-enhanced ALD using WH2(iPrCp)2 as W precursor. The growth of
15 monolayer to multilayer MoS2 film was demonstrated by ALD based on self-limiting reactions of molybdenum hexacarbonyl (MoCl5) and H2S at 300 °C on a sapphire wafer.
[139] The post annealing of the as-deposited MoS2 films further improved its crystallinity. Furthermore, Jin et al. [140] reported on the growth of MoS2 by ALD using
MoCl5 and dimethyldisulfide (C2H6S2) as Mo and S precursors at 100 °C. The crystallinity of the as-grown amorphous MoS2 films was observed to improve on annealing at 900 °C. Although current production status is still in a nascent stage due to poor crystallinity, small grain size and low throughput [138] [139] [140], ALD can be a possible approach to achieve large-scale production of 2D LMCs with precise with thickness control.
In parallel with the recent advances in the synthesis of LMCs, the stacking of various 2D layered materials in vertical and lateral directions is being explored for device applications. Layer-by-layer stacking of 2D LMCs could be achieved by using mechanical transfer techniques. [142] [141] [143] For instance, vertically stacked
WS2/MoS2 heterostructures were deposited on SiO2/Si substrates from CVD-grown WS2 and MoS2 monolayers using conventional PDMS stamping method. [141] However, the use of transfer techniques to vertically stack layers with contamination-free interfaces is challenging. Furthermore, building lateral heterostructures with atomically sharp interfaces remains a major challenge.
CVD-based approaches [10] [144] [145] haven been proven to produce vertical and in-plane heterostructures with atomically clean and sharp interfaces. Gong et al. [10] demonstrated the growth of vertical and lateral WS2/MoS2 heterostructures by a CVD set-
16 up using W, MoO3 and S powders with the addition of Te. By lowering the temperature, they showed that WS2 layers preferentially grow laterally from MoS2 monolayers whereas vertical heterostructures dominate at high temperatures. In a subsequent study,
Huang et al. [146] reported the synthesis of the in-plane MoSe2-WSe2 heterojunctions without the addition of Te. More choices of 2D LMCs building blocks and their combinations are expected in the near future.
CVD synthesis of LMCs leads to polycrystalline films with a variety of grain boundary topologies. Grain boundaries are a source of scattering for electrons and phonons. Hence, understanding grain and grain boundary characteristics in CVD grown
LMCs is essential to improve the characteristics of transistors that adopt these materials.
Recently, high-resolution transmission electron microscopy (HR-TEM) methods have been used to examine the morphology of the grains and their boundaries in MoS2 islands.
[116] The coalescence of the grains leads to the formation of grain boundaries consisting of 5- and 7-membered rings. Coalescence of grain boundaries can also occur with overgrowth. Prior studies [116] have shown there is a competition between the formation of conventional grain boundaries and the overlapped grain boundaries. However, a clear on the degree of lattice orientation mismatch in the two grains was not observed. Zande et al. [111] demonstrated that CVD grown single crystal and polycrystalline triangular
MoS2 grains contain both tilt and mirror twin boundaries containing 8- and 4- membered rings. MoS2 grain boundaries strongly affect the PL and slightly alter the in-plane electrical conductivity. The mirror twin boundaries lead to the strong PL quenching and
17 improved in plane conductivity whereas the tilt boundaries lead to the strong PL enhancement and the lower conductivity.
To fully utilize 2D LMCs’ potential to 2D materials-based electronic devices, p- and n-type doping is required. In this regard, considerable efforts have been recently devoted to developing new doping strategies. Laskar et al. [147] demonstrated p-type
20 -3 conductivity in few-layer MoS2 films with a hole density of 3.1x10 cm using Nb as a substitutional impurity on the Mo site. The Nb-doped MoS2 films were grown by sulfurization of the thin films deposited using electron beam evaporation in the sequence of Mo/Nb/Mo. In another study [148], p-type conduction in the exfoliated MoS2 flakes with a hole density of ∼3×1019 cm-3 was demonstrated by substitutional Nb doping. The
Nb-doped MoS2 crystals were grown by a chemical vapor transport (CVT) method using iodine as the transport agent. Comprehensive theoretical studies [149] have investigated the thermodynamics of doping single monolayers of the LMCs. Substitutional doping of
Mo atoms by Nb introduces minimal change to the metal-chalcogen bond length since the atomic number of Nb (41) is one less than the atomic number of Mo (42). Using DFT calculations it was shown 4% Nb doping introduces a hole defect state approximately 150 meV below the valence band edge of monolayer MoS2. [149] Future studies of substitutional doping in LMC compounds will require the exploration of alternative p- and n-type dopants combined with theoretical studies of thermodynamic ground state, phase diagrams and kinetics of dopant atoms in a given LMC.
Fang et al. [129] have reported the chemical p-doping of monolayer WSe2. NO2 molecules were absorbed on both the WSe2 channel and contacts of the FETs, leading to
18 lower the metal contact resistance. In a similar study [150], n-doping of few-layer MoS2 and WSe2 layers with high motilities was shown by molecular adsorption of potassium
(K) vapor. Electron doping by K vapor leads to lower contact resistance and high motilities. Yang et al. [151] reported a chlorine (Cl) molecular doping technique for
MoS2 and WS2, which leads to high electron-doping density, thus a significant reduction of Schottky barrier width. Two recent studies demonstrated the chemical doping of few- layer MoS2 [152] and monolayer WS2 [153] by using tetrafluoro- tetracyanoquinodimethane (F4TCNQ) as a p-type dopant. It was shown that the more
F4TCNQ adsorption, the more electrons withdraw from MoS2 and WS2, resulting in significant PL enhancement in the layers.
Molecular chemistry was first applied to LMC compounds to study intercalation chemistry. It is now well known that the ion based intercalation of the group VI TMDs induces structural or phase transformations, which leads to a concomitant change in their electronic structure [154] [155] [156] Intercalation of Li ions causes MoS2’s crystalline structure to change from the stable semiconducting 2H-MoS2 to the meta-stable metallic
1T-MoS2 phase. The electron transfer from Li to the MoS2 structure was identified as the driving mechanism for this structural change. The 2H and 1T phases can be identified by
Raman spectroscopy. While both 1T-MoS2 and 2H-MoS2 show Raman active A1g and E2g
-1 -1 modes at 380 cm and 410 cm , 1T-MoS2 results in additional Raman active modes at
160 cm-1, 230 cm-1 and 330 cm-1. [157] However, since the 1T phase is meta-stable, it gradually converts to the trigonal prismatic coordination over time. [155] [156] This phase transformation from trigonal prismatic to octahedral coordination in MoS2 is
19 currently being explored as a platform to engineer low-resistance contacts for field-effect transistor devices utilizing the group VI TMDs. [157] [158]
Alternatively, intercalation of zero-valence ions into LMCs compounds preserves the morphology and structure of the host material. [159] Such an approach can be used to engineer the interlayer coupling between the individual layers of LMC compounds.
1.3 Characterization Methods
A variety of different characterization methods, such as Raman spectroscopy
[160], PL spectroscopy [120], AFM [107], STM [137], TEM [106], ARPES [31], second harmonic generation [SHG] [161], X-ray photoelectron spectroscopy (XPS) [162] and X- ray diffraction (XRD) [109] have been used to study the structure and properties of
LMCs.
In this section, an overview of a limited number of spectroscopy and measurement techniques that have been crucial to characterize the structural, electrical, optical and vibrational properties of few-layer LMCs is provided.
Raman measurements are a useful technique to understand the stacking order, interlayer coupling and thickness of Raman measurements on registered bulk and few- layer metal chalcogenides demonstrate two primary Raman modes, the A1g out of plane and E2g in plane Raman modes. These two distinct peaks that correspond to the in-plane and out of plane vibrational modes are the most intense peaks in the Raman spectra of these materials, determined by the selection rules of the Raman active vibrational modes.
When LMCs, such as MoS2 and MoSe2 are reduced to a single monolayer there are quantitative shifts in the Raman spectra. [86] The in-plane E2g mode frequency increases
20 and the A1g out of plane mode frequency decreases as the film thickness is reduced to a single monolayer. These modes have also been shown to be sensitive to changes in stacking order and misorientation between monolayers. [163] [64] Raman modes due to two-phonon modes have also been observed in single- and few-layer LMC. [164] These shifts in the out of plane Raman mode frequency as a function of film thickness indicates the sensitivity of these materials to sample thickness, interlayer coupling, rotation and stacking.
TEM, in particular, has proven to be an excellent tool for the characterization of
2D materials at the atomic level. Recent advances in the TEM techniques, such as selected-area electron diffraction (SAED) [115], bright-field conventional TEM (BF-
CTEM) [111] and high-angle annular dark field scanning TEM (ADF-STEM) [106], have shown that the layer number and stacking order can be quantitatively determined, and the individual atoms in a layer can be visualized. In a recent report on the growth and structural characterization of MoS2, George et al. [106] showed that monolayers as well as bi- and triple-layers of MoS2 can be identified by using high-angle annular dark-field scanning TEM (HAADF-STEM).
In addition to TEM, ARPES has enabled direct observation of the Fermi surface and underlying electronic structure of 2D LMCs. Prior studies have shown [31] [30] the thickness-dependent electronic band structures of LMCs can be experimentally determined using ARPES. Zang et al. [137] directly measured the electronic band structure of MoSe2 using ARPES.
21 1.4 Applications
The current application space of LMCs and their heterostructures ranges from next generation electronic [165] and optoelectronic devices [11] beyond the technology road map, flexible electronics [62], piezoelectronics [166] and energy storage and conversion [167].
Research into electronic devices that utilize LMC materials as the active material are driven by the search for new materials that would scale beyond the international technology roadmap for semiconductors (ITRS). The semiconducting LMC compounds have an inherent advantage over traditional CMOS materials at dimensions beyond the scaling roadmap (< 5 nm). The larger effective masses of the LMCs (~0.4 to 0.5 m0, m0 is the free electron mass) provide a higher density of states and hence a higher drive current while simultaneously having a low direct tunneling current which lowers the off- state leakage current at channel lengths below 5nm. [168] The robust mechanical properties of LMC compounds combined with mobilities that vary between 1 cm2/Vsec and 100 cm2/Vsec make them suitable candidates for future flexible electronics.
2D LMCs with high surface-to-volume ratio, flexibility, high sensing ability, robust photoluminescence and good electrical conductivity have emerged as new promising materials for chemical sensors and biosensors. [169] The interaction with an analyte changes the charge transport properties of atomically thick LMC materials as an active channel in a FET configuration, so the electrical properties are correspondingly changed upon exposure. In addition to sensing based on changes in electrical properties,
LMCs also provide an alternative avenue for sensing based on light emission. [170] The
22 charge transfer upon physisorption of the analyte leads to changes in the PL emission efficiencies. Application of gate bias and exposure to light have been found to enhance device sensing performance. [171] LMC materials can be also potentially adopted for biosensing applications. More recently, MoS2 has been successfully used as a sensing material for the detection of DNA and small molecules on the basis of its high fluorescence quenching ability and different affinity toward ssDNA versus dsDNA. [172]
Along with their interesting catalytic properties, LMCs hold great promise as catalysts for hydrogen evolution reaction (HER) in which hydrogen is produced through the process of electrocatalytic or photocatalytic water reduction. LMCs provide large edge to volume ratio and greater density of edge sites for high catalytic performance. The edge sites of 2D layers with unsaturated coordination atoms and dangling bonds could serve as active sites for hydrogen bonding. Several studies [173] [174] have demonstrated high catalytic activity in MoS2 which emerges as a low-cost alternative material to the traditional HER catalysts such as Pt, and Pd. Metallic T phase of MoS2 was found to be a more active HER catalyst than semiconducting 2H-MoS2. [175] The possibility of tuning the HER activity of LMCs through phase selection, engineering, nanostructuring and alloying is expected to further increase the substantial advantages of 2D layers for their catalytic applications.
1.5 Grand Challenges and Opportunities
The current focus of research on 2D LMCs is on establishing the fundamental properties of these materials and their heterostructures. These efforts are driven by the hope that amongst this large family of compounds is a solution for next generation low
23 power electronics, high capacity energy storage and conversion and broadband optical devices. The LMCs also provide a platform to study fundamental condensed matter phenomena such as superconductivity and charge density waves at low dimensions. This combination of applied and fundamental research with this family of materials is still in its infancy. The first paper on a monolayer MoS2 transistor appeared in 2011 and the field has been growing ever since. The combined effort on the growth of new layered materials, advances in the fabrication of pristine heterostructures and sophisticated theoretical predictions based on ab-initio theory suggest that further developments in the field are forthcoming that would lead to the adoption of LMCs in practical applications and also provide a platform to explore exotic physical phenomena.
1.6 Thesis Outline
This dissertation focuses on the various unique methods of producing 2D atomic
LMC crystals including growth of group VI transition metal dichalcogenides by chalcogenization of pre-deposited metal-containing precursors, and phase-engineered growth of group IV LMCs by vapor-phase reaction of metal oxide and chalcogen powders. STEM, SEM, AFM, optical microscopy, Raman spectroscopy, PL spectroscopy, ultraviolet photoelectron spectroscopy (UPS), XPS, energy dispersive X- ray spectroscopy (EDX), and XRD are all notable experimental techniques used to characterize these various materials, and the experimental results are corroborated by density functional theory (DFT) calculations.
In chapter 2, we developed a seed-assisted CVD approach for the growth of 2D
WS2 crystals and its alloys with MoS2 on SiO2/Si substrates using hydrothermally
24 synthesized WO3 seeds. This method produces highly crystalline semiconductor WS2 crystals with a bandgap of 1.94 eV in monolayer form. PL spectroscopy measurements in conjunction with ab-initio DFT calculations show that bulk WS2 undergoes a transition from an indirect bandgap to a direct one in monolayer limit. We also showed that the modification of the band gap is possible by changing the composition of Mo and W in
MoS2-WS2 heterogeneous layers, pointing to creating a range of bandgap values by alloying 2D semiconductors.
In chapter 3, we demonstrated the large-area growth of atomically thin WS2 films via sulfurization of the pre-deposited W-based films on SiO2/Si substrates. The fundamental morphologic, electronic, optical and chemical properties of the pre- deposited W-based films and WS2 films were investigated using various microscopy and spectroscopy techniques. We showed that the WS2 films show laser power dependent PL characteristics, where the PL peak shifts toward lower photon energy, and its intensity increases with laser power. Moreover, we introduced a simple method to etch WS2 films with well-oriented triangular pits with a preferential edge termination.
In chapter 4, the phase-selective growth of 2D tin sulfides (SnS2 and SnS) is demonstrated. Highly crystalline hexagonal SnS2 and orthorhombic SnS crystals were synthesized on SiO2 substrates by a simple vapor-phase method using SnO2 and S as the source materials. The structural, vibrational, and electronic properties of each phase are studied by both experimental characterizations and ab-initio DFT calculations.
In chapter 5, we reported on the growth of 2D SnS2 nanowalls (NWs) by a facile vapor-phase synthesis method on insulator substrates such as SiO2/Si and MgO using
25 SnO2 and S powders as precursors. The fundamental properties of the SnS2 NWs were studied by using various techniques. The results suggested that the grown films are highly crystalline SnS2 with an open network structure constituted of interconnected NWs.
In chapter 6, we demonstrated the synthesis of few-layer MoSe2 films in arbitrary shapes and patterns on SiO2 wafers by rapid thermal processing (RTP) and Raman laser annealing of the stacked elemental layers, which are deposited using electron beam evaporation in the sequence of Mo/Se/Mo. The morphology and microstructure of the films were characterized by AFM and SEM. Compositional and electronic structure analysis of the films were done by XPS, PL spectroscopy and Raman spectroscopy.
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40 Chapter 2 Seed-Assisted CVD Synthesis of 2D WS2 and MoS2-WS2
Abstract
Atomically thin WS2 crystals were synthesized on SiO2/Si substrates via chemical vapor deposition (CVD) with the assistance of hydrothermally synthesized WO3 seeds.
This seed-assisted CVD approach allows production of highly crystalline semiconductor
WS2 with a bandgap of 1.94 eV in monolayer form. Photoluminescence (PL) spectroscopy measurements in conjunction with ab-initio density functional theory (DFT) calculations showed that bulk WS2 undergoes a transition from an indirect bandgap to a direct one in monolayer limit. We also synthesized heterogeneous systems of MoS2 and
WS2 and characterized them by Raman and PL spectroscopy measurements. The DFT calculations verify the modification of the band gap by changing the composition of Mo and W in MoS2-WS2 heterogeneous layers, pointing to creating a range of bandgap values by alloying two-dimensional (2D) semiconductors.
2.1 Introduction
2D layered transition metal dichalcogenides (TMDs), especially WS2, have become the focus of intense scientific interest due to their exotic physical properties and exciting prospects for a variety of applications. [1, 3] TMDs exhibit a large variety of electronic behaviors such as semiconductivity, superconductivity or charge density waves. [4-6]
Layered semiconductor WS2, with an indirect bandgap of 1.48 eV in bulk form and a direct band gap of 1.94 eV in monolayer form, exhibits tunable electronic properties allowing for the realization of advanced optoelectronic devices, transistors and
41 valley polarization devices, etc. [7-10] The WS2 bulk unit cell belongs to the hexagonal space group P63/mmc with lattice parameters of a=3.1532 Å and c=12.323 Å, and each
WS2 monolayer contains an individual layer of W atoms with a 6-fold coordination symmetry, which are hexagonally packed between two trigonal atomic layers of S atoms.
[11]
Scientists and researchers around the world have been racing to discover and engineer methods for the reliable synthesis of 2D materials. Although the isolation of atomically thin 2D flakes from their bulk crystals through exfoliation is still the method choice for fundamental studies, it is impractical for large scale applications due to the size limitations of the produced flakes. [12]
CVD has been considered to be one of the most promising approaches for large scale production of 2D materials. [13, 14] Typical CVD synthesis of atomically thin WS2 films includes sulfurization of W or WO3 films, which are pre-deposited on the substrates by using sputtering, e-beam deposition, thermal evaporation or atomic layer deposition
(ALD) methods, or vapor phase reaction of elemental S and WO3 powders at high temperatures. [15-17]
Recently, Orofeo et al. have reported on the synthesis of few-layers WS2 films by sulfurization of magnetron sputtered W films. [15] The synthesis of monolayer to few- layer WS2 nanosheets has been demonstrated through sulfurization of an ALD WO3 film.
[18] Zang et al. have presented the synthesis of atomically thin WS2 triangular flakes on sapphire via low-pressure CVD reaction of WO3 and S. [19] However, to date, no studies have examined the role of the seeding in the synthesis of WS2.
42 Here, we have developed a seed-assisted CVD procedure for the synthesis of 2D atomically thin WS2 and MoS2-WS2 heterogeneous layers using WO3 nanoplates (NPs) as nucleation promoters. The morphological, optical and electronic properties of the synthesized layers have been characterized by scanning electron microscopy (SEM), atomic force microscopy (AFM), optical microscopy, Raman spectroscopy and PL spectroscopy. Ab-initio DFT calculations have been performed to evaluate the electronic band structures.
2.2 Experimental Section
2.2.1 Materials Synthesis
WO3 NPs were synthesized by a simple hydrothermal process and used as the precursor for WS2 growth. 0.16 gr of sodium tungstate dehydrate (Na2WO4∙2H2O) was first dissolved in 2 mL of deionized (DI) water and 10 mL of nitric acid (HNO3) aqueous solution under magnetic stirring. The light yellow aqueous solution obtained after continuous 2 h of vigorous stirring was transferred to a 45 mL teflon-lined autoclave. The autoclave was sealed and kept at 180 ºC for 2 h. The final product was washed with DI water several times and dried in a vacuum at 100 ºC for 5 h. Si substrates with a thermally deposited 300 nm thick SiO2 layer were cleaned in successive ultrasonic baths of acetone and isopropanol (IPA) for 30 min each.
WO3 NPs were dispersed on SiO2/Si substrate, and the substrate was placed face- down above a polished quartz boat containing 0.5 mg of WO3 powder. The WO3-coated substrate with the boat was then located at the center of heating zone. A ceramic boat containing 100 mg of S was located upstream where S starts evaporating when the
43 temperature of the center of the tube is 525 ºC. The furnace was pumped down to 0.1
Torr in order to remove the air and then stabilized at 20 Torr with an Ar flow of 200 sccm. Next, the center of heating zone was heated to 525 ºC, then to 850 ºC with ramping rates of 10 ºC /min and 5 ºC /min, respectively. The furnace was kept at 850 ºC for 15 min, and then naturally cooled down to room temperature (RT).
For the synthesis of MoS2-WS2, the same experimental set-up and the recipe were used. The WO3 NPs were also employed as nucleation promoters. The only difference was the use of 0.5 mg of MoO3 powder instead of WO3.
2.2.2 Materials Characterization
Raman and PL spectra were collected using a 532 nm laser (<2 mW excitation power, 100x objective lens) at RT. The thickness measurements were performed by AFM in tapping mode. Microstructural analysis was done by SEM.
2.2.3 DFT Calculations
Our calculations were based on first-principles DFT using the projector augmented wave method as implemented in Vienna ab-initio simulation package
(VASP). The screened Heyd-Scuseria-Ernzerhof (HSE) hybrid functional was employed for this study. A Monkhorst-Pack scheme was adopted to integrate over the Brillouin zone with a k-mesh 9 x 9 x 1 (8 x 8 x 4) for monolayer and bulk structure. A plane-wave basis kinetic energy cutoff of 280 eV was used. The lattice constant for monolayer WS2 structure was obtained from a volume optimized bulk 2H-WS2 structure. The atomic coordinates for monolayer structure were optimized at this fixed lattice constant. Spin- orbit coupling (SOC) was included self-consistently within the band structure
44 calculations for monolayer WS2. For the HSE calculations, 25% short-range exact
Hartree-Fock exchange was used with the Perdew-Burke-Ernzerhof (PBE) correlation.
The HSE screening parameter, µ, was empirically set to 0.5 (1/Å) for bulk 2H-WS2 and
0.26 (1/Å) for monolayer WS2. For the Mo1-xWxS2 system, the calculations were conducted using the PBE GGA approximation of DFT as implemented in the software package VASP. A (5x5x1) Mo1-xWxS2 monolayer supercell was setup using the volume optimized bulk lattice constant of WS2 (3.179 Å). The following values of x were used:
0.0, 0.12, 0.20, 0.32, 0.40, 0.56, 0.60, 0.80 and 1.0. For each of these supercell structures the positions of the W atoms were randomly placed. Following a structure optimization of each supercell, the electronic band structure was calculated for each concentration of
W atoms. For each structure, the bandgap remained direct.
2.3 Results and Discussion
The morphology of WO3 NPs, which were hydrothermally synthesized, was investigated by SEM (Figure 2.1(a)). The thickness and lateral size of the NPs was estimated to be tens of nanometers and a few microns, respectively. WO3 NPs were further characterized by Raman spectroscopy analysis. Raman spectra of WO3 (Figure
2.1(b)) display three major Raman modes at 809 cm-1 and 708 cm-1 (O-W-O stretching modes) and 271 cm-1 (O-W-O bending mode), consistent with previous studies. [20]
45
Figure 2.1 The morphology and Raman spectra of the WO3 crystals. SEM image (a) and Raman spectra (b) of WO3 NPs.
The low (Figure 2.2(a)) and mid-high (Figure 2.2(b)) magnification optical images of a typical sample show that our growth method produces primarily triangular and multi-point star shaped WS2 domains which are several microns in size and have a clear contrast from SiO2 substrate. We can hypothesize that the diffusion of S into WO3
NPs results in WS2 nuclei in conjunction with chemisorbed sulfur species. [21] The thinner WS2 layers then expand from the edges of the nuclei upon further sulfurization with the supply of WO3 vapor.
It is known that any compositional inhomogeneity on SiO2/Si substrate surface can provide lower necessary activation energy for the formation of WS2 nuclei. [19]
Therefore, we employed WO3 NPs as nucleation promoters for the growth of WS2 crystals by bottom-up CVD process. We found that the distribution of WS2 domains is not homogenous on the substrate (Figure 2.2(a)) due to the heterogeneous distribution of
WO3 NPs over the substrate. This is a current limitation of our experimental approach that can be resolved technically in the future. We also observed that the dense WO3 seeds
46 promote a large number of WS2 domains to form aggregates and continuous sheets
(Figure 2.2(b)).
In the course of these experiments, we noticed that WO3 crystals with different morphologies can form (Figure 2.2(c,d)). Although these observations were not common in our experiments, it clearly shows that the morphology of WO3 dictates the size and shape of the grown WS2 domains, consistent with previous studies. [13] Moreover, since the activation energy for the nucleation is substantially low due to the roughness and edges provided by WO3 crystals, more layers are grown from the same origin, which creates thicker regions on the layers.
47
Figure 2.2 The morphology of the WS2 crystals. Low (a) and mid-high (b) magnification optical microscopy images of a typical WS2 film grown on SiO2/Si. Note that the alignment marks were made after the WS2 growth for the device fabrication.
SEM images (c, d) of WS2 domains with different morphologies.
AFM is a powerful method to determine the thickness of 2D materials with a precision of 5 %. [22] The AFM topography image (Figure 2.3) of a bilayer WS2 domain reveals the presence of the sharp edges forming a 60° angle. Their edges are along a specific crystallographic direction, and only one edge termination is energetically favored. [23] Although our experiments cannot determine whether the preferred edge is the W-edge or S-edge, it is expected that the preferred edge is the W-edge in WS2 domains owing to the sharp edges. [24] The thickness of the first layer was determined to
48 be 1.3 nm, which is higher than that of the second layer with a thickness of 0.7 nm
(Figure 2.3). [16, 19] This discrepancy arises from differences in the interactions of the tip with the sample and substrate. [22] Therefore, the thickness of a monolayer WS2 can be better determined by measuring the height of a second WS2 layer on a first layer since the tip-layer interactions are constant.
Figure 2.3 AFM thickness analysis of the WS2 crystal. AFM image of a bilayer WS2 crystal. The inset shows the height profile along the dashed line in the image.
Raman spectroscopy was used to confirm that the grown domains are indeed
WS2. The Raman spectrum (Figure 2.4(a)) of a monolayer WS2 using a laser with a
1 wavelength of 532 nm demonstrates the characteristic in-plane vibrational (E 2g) mode
-1 -1 and the out-of-plane vibrational (A1g) mode at 350.25 cm and 417.50 cm , respectively.
49 1 [25] The frequency difference between the E 2g and A1g modes was determined to be to
-1 be 67.25 cm , consistent with what was observed on the CVD-grown monolayer WS2.
[17]
The PL study showed that WS2 exhibits a transition from indirect band gap semiconductor in the form of bulk to direct band gap one in monolayers, similar to MoS2.
[26] A strong and sharp PL peak at 1.94 eV was found in monolayer WS2 (Figure
2.4(b)), which indicates an indirect to direct band gap transition in monolayer WS2. [16,
27] The PL intensity was found to be extremely weak on bulk WS2, consistent with an indirect band gap semiconductor in bulk form. To gain further insight, we computed the electronic band structures of monolayer and bulk WS2 using ab-initio DFT. According to our calculations (Figure 2.4(c,d)), monolayer WS2 has a direct band gap of 1.90 eV and bulk WS2 has an indirect band gap of 1.48 eV, which is consistent with prior reported studies on the electronic structure of WS2 and our PL data on monolayer WS2. [28, 29]
50
Figure 2.4 Raman-PL spectra and DFT-calculated electronic band structure of the
WS2 crystals of different thickness. Raman spectra (a) of monolayer WS2. PL spectra comparison (b) of monolayer and bulk WS2. DFT-calculated electronic band structures of monolayer (c) and bulk (d) WS2.
We also performed Raman mapping of the WS2 domains with different
1 morphologies by plotting the intensity ratio of E 2g and A1g peaks. It is known that intensity ratio mapping provides a more accurate characterization and better resolution for 2D layers with different thicknesses. [30] The Raman mapping image (Figure 2.5(a)) of the WS2 domain with perfect triangular shape revealed the presence of the second layer on the nucleation site, which is not visible in the optical microscopy image (Figure
2.5(b)). The second layer nucleates in the vicinity of the edges of the WO3 nucleus where the first layer grows, and it follows the growth direction of the underlying layer. [13] It was also found that a large WS2 domain can be formed by joining several layers from the same nucleus. Although the grain boundaries are not clear in the optical microcopy image
51 (Figure 2.5(d)) of the multi-point star WS2 domain, the Raman mapping image (Figure
2.5(c)) clearly shows the boundaries between the triangular shaped flakes, proving the capability of Raman spectroscopy in characterization of atomically thin 2D TMDs.
Figure 2.5 Raman imaging of the WS2 crystals of different morphologies. Raman mapping and optical images of the triangular (a, b) and multi-point star (c, d) shaped
WS2 domains, respectively.
We also synthesized MoS2-WS2 heterogeneous layers using WO3 NPs as nucleation promoters in order to study the electronic and optical properties of the
1 heterogeneous systems. The intensity ratio (E 2g/A1g) mapping image (Figure 2.6(a)) and
Raman spectra (Figure 2.6(b)) taken from different regions of a typical MoS2-WS2
52 domain confirm the co-existence of MoS2, WS2 and Mo1-xWxS2 layers. [31, 32] The PL measurements (Figure 2.6(c)) revealed that Mo1-xWxS2 has a relatively strong PL peak at
1.87 eV (663 nm), lying between those of MoS2 (1.85 eV) and WS2 (1.94 eV).
Figure 2.6 Raman-PL spectroscopy characterization and DFT-calculated band gap of the MoS2-WS2 heterogeneous crystals. Raman mapping image (a) and optical microcopy image (inset) of a MoS2-WS2 domain. The corresponding Raman (b) and PL
(c) spectra comparison of the P1 (MoS2), P2 (WS2), and P3 (Mo1-xWxS2) points in (a).
The variation in band gap as a function of the composition of W atoms in the Mo1-xWxS2 supercell by the DFT calculations (x=0 denotes a MoS2 supercell while x=1 denotes a
WS2 supercell) (d).
To verify the modification of the PL peak in the Mo1-xWxS2 system and understand the dependence of the band gap on Mo and W composition, we conducted the
DFT calculations on a Mo1-xWxS2 monolayer supercell and showed the variation in band
53 gap as a function of the composition of W atoms in monolayer Mo1-xWxS2 supercell
(Figure 2.6(d)). The calculation results are in close agreement with our experiment results and the previous reports. [31, 33]
2.4 Conclusion
In conclusion, we have demonstrated the synthesis of 2D semiconductor WS2 layers on SiO2/Si via CVD using hydrothermally synthesized WO3 NPs as seeds. The
DFT computational results show that an indirect-to-direct band gap transition occurs in
WS2 monolayers, in good agreement with PL spectroscopy measurements. AFM results confirm that this growth method produces atomically thin WS2 layers with a monolayer thickness of 0.7 nm. We have also synthesized MoS2-WS2 heterogeneous layers by using this method, showing our capability to fabricate a variety of different 2D crystalline materials and their heterogeneous systems. We hope that the results reported herein create a basis for the systematic exploration of these materials in research and potential applications.
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57 Chapter 3 Large-Area Synthesis and Anisotropic Etching of 2D WS2
Abstract
Emergent properties of WS2 at the quantum confinement limit hold promise for electronic and optoelectronic applications. Here we reported on the large-area synthesis of two-dimensional (2D) WS2 with strong photoluminescence (PL) properties via sulfurization of the pre-deposited W films. Detailed characterization of the pre-deposited
W films and WS2 films were performed using various microscopy and spectroscopy methods. By directly heating WS2 in air, we showed that the films tend to be etched into a series of triangular shaped pits with the same orientations, revealing the anisotropic etching behavior of WS2 edges. Moreover, the dimensions of the triangular pits increase with the number of layers, suggesting a thickness dependent behavior of etching in WS2.
This method offers a promising new avenue for engineering the edge structures of WS2.
3.1 Introduction
Graphene-like 2D semiconducting transition metal dichalcogenides (TMDs), especially WS2, with sizeable band gaps have showed vast potential for electronics and optoelectronics due to exotic physical and electronic properties, recently. [1, 2] WS2 is a layered semiconductor material with a direct band gap of ∼1.95 eV in monolayer form.
[3] For electronics, the presence of a band gap allows field-effect transistors (FETs) to have high on/off ratios, and its ultrathin nature allows the channel length to be reduced relative to those fabricated with conventional semiconductors. [4, 5] For optoelectronics, the direct band gap produces high photoconductivity, and strong PL. [3, 6] To explore
58 new fundamental properties and to further develop their electronic and optoelectronic applications, synthesis of large-area atomically thin WS2 layers with uniform properties by a facile and scalable method is an essential requirement.
Recent top-down approaches including mechanical exfoliation and liquid exfoliation to obtain high crystalline WS2 flakes have attracted considerable attention. [7,
8] However, lateral dimensions of the flakes synthesized by the exfoliation methods are limited to few microns, which limits their applications in large-scale electronics and optoelectronics.
In contrast, chemical vapor deposition (CVD) techniques have great potential in producing large-area WS2 over macroscopic sizes, which are ideal for integration with current CMOS platform. Typically, CVD growth of WS2 includes vapor phase reaction or deposition of gaseous metal and chalcogenide feed stocks, and sulfurization or decomposition of the pre-deposited films. [9-12]
Among the CVD techniques, sulfurization of the pre-deposited thin films is emerging as a quick and easy way to obtain large-area atomically thin WS2 films on insulating substrates. Recently, several studies have reported direct sulfurization of WO3 films deposited either by thermal evaporation or by atomic layer deposition (ALD). [11,
12] However, to the best of our knowledge, large-area growth of atomically thin WS2 films on SiO2 substrates by sulfurization of W films deposited by e-beam deposition has not been reported yet.
In parallel with the recent advances in the synthesis of TMDs, numerous methods have been developed to enable their identification and characterization. PL spectroscopy
59 is a powerful method of probing the electronic structure of TMDs. [13] Recent studies report that MoS2 shows laser-dependent PL spectra, where the PL intensity and position are affected by varying the duration and intensity of laser irradiation. [14] To the best of our knowledge, there is no systematic study concerning the effect of laser power on PL spectra of WS2, which is essential to probing the PL characteristics of WS2.
Recently, there has been an increasing interest in the nanoscale control of edge structures of TMDs. Nanoscale control of edge structures offers new pathways toward fine tuning the electronic, optical, chemical, magnetic, and catalytic properties of TMDs.
[15-18] One way to engineer edge structures of TMDs is heating them in an O2 environment. [19, 20] For example, high density triangular pits with the Mo or S terminated zigzag edges on the surface of MoS2 sheets have been obtained upon annealing them in air, which might arise from the anisotropic etching of the active MoS2 edge sites. [19] Such edge terminated MoS2 structures find applications in diverse catalytic reactions. [21, 22] As mentioned above, most studies so far have focused on the oxidative etching of MoS2. However, to the best of our knowledge, the etching behavior of WS2 films has yet to be experimentally studied.
In this study, we demonstrate the large-area synthesis of atomically thin WS2 films on SiO2 substrates. Briefly, the W films are deposited on SiO2 substrate by using e- beam deposition method. In the next step, the as-deposited W films are annealed at 500
°C and then sulfurized at 850 °C to obtain WS2 films. The fundamental morphologic, electronic, optical and chemical properties of the pre-deposited W films and WS2 films are investigated using optical microscopy, atomic force microscopy (AFM), scanning
60 electron microscopy (SEM), X-ray photoemission spectroscopy (XPS), Raman spectroscopy and PL spectroscopy. We also systematically investigate the dependence of
PL spectra of WS2 films on laser power. Furthermore, we introduce a simple method to etch WS2 films with well-oriented triangular-shaped pits by heating them in air.
3.2 Experimental Section
3.2.1 Materials Synthesis
WS2 films were synthesized via three steps: i) deposition of the W films, ii) annealing of the W films, and iii) sulfurization of the annealed W films. First of all, the
W films with a thickness of ∼15 nm were deposited on SiO2 substrates (300 nm thick
SiO2 layer deposited thermally on Si wafer) by e-beam deposition method. The deposition of the films was carried out in a load-lock chamber at a base pressure of ∼10-6
Torr and deposition pressure of ∼10-5 Torr, with an average growth rate of 1 Å/s. Next, the as-deposited W films were annealed at 500 °C in open air for 60 min in order to oxidize the films in a single-zone vacuum tube furnace with 2-inch diameter quartz tube.
The annealed W films were then put at the center of the heating zone, and 250 mg of S powder was put in a ceramic boat at the upstream where S starts evaporating when the temperature of the center of the tube is 525 °C. High-purity Ar gas with a constant flow of 100 sccm was used as the carrier gas. After that, the furnace was pumped down to remove the air, and the pressure was stabilized at 250 Torr. Further, the center of heating zone was heated to 850 °C with a heating rate of 5 °C/min, and it was kept at that temperature for 10 min. Finally, it was naturally cooled down to room temperature. The
61 etching of WS2 films was performed in the same furnace by simply heating them in air at
500 °C for 45 min.
3.2.2 Materials Characterization
The morphology of WS2 films was investigated using SEM. Raman and PL spectra were collected using a 532 nm laser (<2 mW excitation power, 100x objective lens) at RT. Surface chemical analysis was carried out by using a XPS system equipped with an Al Kα monochromated X-ray source and a 165-mm mean radius electron energy hemispherical analyzer. The vacuum pressure was kept below 3 × 10-9 torr, and the neutralizer was applied during the data acquisition. The thickness measurements were performed by AFM in tapping mode.
3.3 Results and Discussion
The morphology and thickness of the W films before and after annealing are determined by AFM analysis. The AFM height image (Figure 3.1(a)) of the as-deposited films reveals island type morphology with a height of about 15 nm. The simultaneously recorded AFM phase contrast image (Figure 3.1(b)) reveals an additional contrast on the surface of islands, which may be attributed to the ‘native oxide layer’.
62
Figure 3.1 AFM analysis of the as-deposited and annealed W films. AFM height and corresponding phase images for the as-deposited (a, b) and annealed (c, d) W films on
SiO2 substrates, respectively. The AFM height and phase contrast images are taken in an area of 2x2 μm2. The insets in (a) and (c) show the height profiles for the island morphologies.
The as-deposited W films are annealed in open air for 60 min at 500 °C.
Annealing is found to induce changes in the morphology and thickness of the films. The
AFM height image (Figure 3.1(c)) of the annealed W films reveals roughly triangular island type morphologies with a height of about 6 nm. While the density of the islands increases, the thickness of the islands decreases upon annealing. This could be explained as follows: It is known that when Tammann temperature (∼599 °C for WO3) is achieved,
63 the evaporation of metal oxide thin films can be initiated even though the processing temperature is lower than their melting point (∼1473 °C for WO3). [23] It is possible that surface oxidation of the as-deposited W films may occur during deposition. Thus, the as- deposited W films with the native oxide layer might become thinner possibly due to the partial evaporation of the metal oxide layer upon annealing. The AFM phase contrast image (Figure 3.1(d)) reveals layered-type island morphologies with one single phase contrast, which can indicate a completed oxidation.
The as-deposited and annealed W films are investigated by XPS analysis. The high-resolution XPS spectra (Figure 3.2(a)) of the as-deposited films reveal four major peaks at 37.99 eV, 35.87 eV, 33.38 eV and 31.2 eV. The typical doublet W 4f peaks of W are clearly visible in the spectra, which are at 33.38 eV and 31.2 eV. [9] The two upper binding energy peaks at 37.99 eV, 35.87 eV can be attributed to WO3, confirming the existence of WO3 in the as-deposited films. [24] It is also found that there is a shoulder at upper energy side of each W peak, and the shoulders are related to WO2 and WOx. [24]
We can conclude that W, WO3 and sub-stoichiometry tungsten-oxides exist in the as- deposited film. The high-resolution XPS spectra (Figure 3.2(b)) of the annealed films reveal only two W 4f peaks at 37.97 eV and 35.87 eV, which are assigned to WO3. After annealing, the metallic W peaks and the shoulder peaks related to WO2 or Wx completely disappear. This evidence strongly supports the stoichiometry phase evolution of the as- deposited films.
64
Figure 3.2 XPS analysis of the W and WS2 films. High resolution XPS spectra of W 4f core levels of the as-deposited (a) and annealed (b) W films. High resolution XPS spectra of W 4f (c) and S 2p (d) core levels of WS2 films.
The XPS data of WS2 films are also captured to analyze surface composition of the WS2 films. The two peaks of each core level appear due to spin orbital splitting of W
4f (Figure 3.2(c)) and S 2p (Figure 3.2(d)) core levels. The peaks at 38.89 eV, 35.40 eV, and 33.2 eV are ascribed to W 5p3/2, W 4f5/2, and W 4f7/2 orbitals, respectively. The binding energies for the S 2p1/2 and S 2p3/2 are located at 164.0 eV and 162.82 eV, respectively. The binding energies of the W and S elements are consistent with the W
(+4) and S (-2) oxidation states in WS2, further confirming the formation of a pure WS2 phase. [25]
To study the effect of annealing on the growth of WS2 films, we have synthesized
WS2 films by sulfurizing the as-deposited and annealed W films. It is observed that
65 sulfurization of the as-deposited W films leads to thick WS2 films with poor coverage
(Figure 3.3(a,b)). However, thinner WS2 films with high coverage are obtained by sulfurization of the annealed W films (Figure 3.3(c,d)). For clarification, monolayer, few-layer and thick regions of WS2 films and the substrate are indicated by arrows on the high magnification SEM images (Figure 3.3(b,d)), as confirmed by Raman and PL spectroscopy. We can conclude that the WS2 growth is substantially improved by annealing the as-deposited W films before sulfurization.
Figure 3.3 The morphology of the WS2 films. Low and high magnification SEM images of WS2 films obtained by sulfurizing the as-deposited (a, b) and annealed (c, d)
W films, respectively.
The growth mechanism of WS2 films can be explained as follows. The reaction of
S vapor with WO3 islands result in WS2 nuclei. This reaction includes a transition to
WO(3-x) species along with the subsequent formation of oxisulfides. As the reaction proceeds, a complete conversion to WS2 occurs, and thinner WS2 layers expand from the edges of WS2 nuclei, in a similar manner as in MoS2. [26]
66 We have taken Raman and PL spectra from monolayer and few-layer regions of a
WS2 film as indicated by arrows on the optical microcopy image (Figure 3.4(a)) of the film. The Raman spectra (Figure 3.4(b)) of monolayer region of the WS2 film
1 demonstrate the characteristic in-plane vibrational (E 2g) mode and the out-of-plane
-1 -1 vibrational (A1g) mode at ∼350 cm and ∼416 cm , respectively. [27, 28] The frequency difference between the two modes is found to be ∼66 cm-1. These results are consistent with what has been reported for WS2 in the previous studies. [29] The PL spectra (Figure
3.4(c)) of monolayer and few-layer regions of the WS2 film are compared. For monolayer
WS2, the PL peak has a maximum at 1.96 eV, which falls in the range of the reported PL peak positions for monolayer WS2. [7, 29] The PL peak of monolayer WS2 is also higher and sharper than that of the few-layer WS2, which indicates the presence of a direct band gap in monolayer WS2. [30] For few-layer WS2, the PL peak shifts toward lower photon energy (∼1.93 eV), and its intensity decreases dramatically, revealing the thickness dependent PL characteristics of WS2.
67
Figure 3.4 Raman-PL spectroscopy characterization of the WS2 films. Optical microscopy image (a) of a WS2 film. Monolayer and few-layer regions of the WS2 film and the substrate are indicated by arrows. Raman spectra (b) of monolayer region of the
WS2 film. PL spectra (c) of monolayer and few-layer regions of the WS2 film. Laser power dependent PL spectra (d) of monolayer region of the WS2 film.
We have also investigated the effect of laser power on PL spectra of monolayer region of the WS2 film with increasing laser power ranging from 0.1 mW to 0.5 mW with a step of 0.05 mW, and from 0.5 mW to 5 mW with a step of 0.5 mW (Figure 3.4(d)). It is observed that the PL intensity increases with increasing laser power, consisted with the results reported for MoS2. [14] Moreover, the PL peak shifts toward lower photon energy as laser power increases. The shift in the PL peak can be attributed to the defect related optical transition. [31] Thus, we conclude that laser power must be carefully controlled in
68 the applications of PL spectroscopy in order to draw reliable conclusions about the PL characteristics of TMDs.
Furthermore, by directly heating a WS2 film in air at 500 °C, we show that the
WS2 film is tend to be etched with a series of the triangular-shaped pits with the same orientations (Figure 3.5). The triangular shape of the pits indicates that the WS2 film is etched preferentially along a specific crystallographic direction that is parallel to the edges of the pits with a preferential edge termination (W-edge or S-edge). [19] The etching is likely to be initiated at the intrinsic structural defects in the WS2 film. [19] We also show that the size of the pits is greatly affected by the WS2 film thickness, where the sizes of the pits increases with the thickness. The sizes of the triangular pits on a multilayers region (with a thickness of about 10.2 nm) of the WS2 film (Figure 3.5(a)) is found to be larger than these of the pits on a few-layer region (with a thickness of about
3.2 nm) of the film (Figure 3.5(b)). This suggests the thickness dependent behavior of the etching on the WS2 film. [32] These results are similar to previous reports on the etching of MoS2. [19, 32, 33]
69
Figure 3.5 AFM analysis of the triangular pits formed on the WS2 films by air- annealing. AFM height images showing the formation of the triangular pits on multilayers (a) and few-layer (b) regions of a WS2 film. The AFM images are taken in an area of 2x2 μm2. The insets in (a) and (b) show the height profiles for the triangular pits.
3.4 Conclusion
In summary, we report a facile method for the large-scale synthesis of atomically thin WS2 films with strong PL properties via sulfurization of the pre-deposited W films.
Thermal annealing of the W films before sulfurization is found to substantially improve the quality of WS2 films in terms of thickness and uniformity. We also find that WS2 films show laser power dependent PL characteristics, where the PL peak shifts toward lower photon energy, and its intensity increases with laser power. Furthermore, we introduce a simple method to etch WS2 films with well-oriented triangular pits with a preferential edge termination. The edge sizes of the pits increase with the thickness of
WS2 films. This etching method enables successful formation and modification of the edge structures of WS2 films with potentially tunable electronic, optical and magnetic properties, which promise various applications ranging from catalysis to energy harvesting.
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73 Chapter 4 Phase Engineering of 2D Tin Sulfides
Abstract
Tin sulfides can exist in a variety of phases and polytypes due to the different oxidation states of Sn. A subset of these phases and polytypes take the form of layered
Two-dimensional (2D) structures that give rise to a wide host of electronic and optical properties. Hence, achieving control over the phase, polytype, and thickness of tin sulfides is necessary to utilize this wide range of properties exhibited by the compound.
This study reports on phase-selective growth of both hexagonal tin (IV) sulfide SnS2 and orthorhombic tin (II) sulfide SnS crystals with diameters of over tens of microns on SiO2 substrates through atmospheric pressure vapor-phase method in a conventional horizontal quartz tube furnace with SnO2 and S powders as the source materials. Detailed characterization of each phase of tin sulfide crystals is performed using various microscopy and spectroscopy methods, and the results are corroborated by ab-initio density functional theory (DFT) calculations.
4.1 Introduction
2D tin sulfides belong to a broad class of van-der-Waal (vdW) materials that have been the subject of intense research. Their constituent elements Sn and S are inexpensive and abundant in nature. The atomically thin geometry combined with the ability to tune the electronic structure of 2D tin sulfides as a function of thickness makes them attractive candidates to be used in nanoelectronics and optoelectronic devices. [1] These sulfides can crystallize in a variety of crystal phases, such as SnS2 and SnS, which possess distinct physical and electronic properties. [2] Therefore, identifying approaches for controlling
74 crystal phase is essential for tailoring the properties of these materials to specific applications.
Tin sulfides possess desirable properties for numerous electronic and optical applications. SnS2 is an intrinsic n-type layered semiconductor with a bandgap of 2.18-
2.44 eV. [3-6] Potential applications of SnS2 include novel electronics devices, flexible electronics, optoelectronics, and energy storage. [7-10] The high electron affinity of SnS2 makes it a suitable candidate to achieve broken-gap heterostructures which enable the design and operation of 2D tunnel field-effect transistor (TFETs). [7, 11] Moreover, SnS2 crystals show high charge carrier mobility (~ 230 cm2/Vs) combined with high on/off current ratios (>106) in field-effect transistors (FETs), fast photocurrent response time (~
5 μs) in photodetectors, and high theoretical specific capacity (~ 920 mAh/g) in lithium- ion batteries (LIBs) as the anode material. [1, 9, 12] On the other hand, SnS is an intrinsically p-type layered semiconductor with an indirect and a direct bandgap at ~ 1-
1.1 eV and ~ 1.3-1.5 eV, respectively. [3, 13-16] The desirable bandgap, high absorption coefficient (>104 cm-1) and sensitivity of its properties to changes in dimensionality, and elemental composition make SnS a promising material for optoelectronic applications.
[15, 17-21] Furthermore, combined SnS and SnS2 heterostructures have been reported to show improved photocatalytic performance on photodegrading organic dyes, compared with that of as-synthesized pure phases of SnS and SnS2. [4]
Selective pure phase synthesis of single crystalline tin sulfides is a challenging task due to the high sensitivity of the different phases to growth conditions. Tin sulfides have been previously synthesized mostly in the form of bulk single crystals or thin films.
75 The bulk single crystals have been grown by the Bridgman method and chemical vapor transport (CVT); the thin films have been synthesized using various deposition methods.
[2, 5, 22-27] However, relatively few studies have reported the synthesis of 2D crystals of SnS2 and SnS using vapor-phase methods. [3, 9, 28, 29]
In this study, we demonstrate phase-controlled synthesis of 2D SnS2 and SnS crystals using an atmospheric pressure vapor-phase method. The fundamental structural and electronic properties of each crystal phase are comprehensively investigated using optical microscopy, atomic force microscopy (AFM), scanning electron microscopy
(SEM), scanning transmission electron microscopy (STEM), Raman spectroscopy, and ultraviolet photoelectron spectroscopy (UPS) and compared with results of ab-initio DFT calculations.
4.2 Experimental Section
4.2.1 Materials Synthesis
SnO2 and S powders were used as source materials in a conventional single-zone horizontal quartz tube (1 in.) furnace. In a typical experiment, SnO2 (~ 15 mg) powders in a quartz crucible were placed in the center of the tube furnace. Also, S (~ 350 mg) powders in a quartz crucible were placed 13 cm upstream from the center of the tube furnace. The degenerately boron-doped (0.001-0.005 Ω cm) Si substrates (~ 1×1 cm2) with 285 nm SiO2 capping layer were cleaned by sonication in acetone and isopropyl alcohol (IPA) baths, sequentially for ~ 1 min each, followed by drying with N2 gas. The
SiO2 substrates were placed downstream to the center at a distance of about 9 (where the temperature was ~ 550 °C) and 11.5 cm (where the temperature was ~ 425 °C),
76 respectively, for the growth of SnS and SnS2 crystals. The temperature of the substrates was determined by using a type K thermocouple. The tube furnace was pumped down to remove the air and then filled with a high-purity Ar gas (~ 100 sccm) to atmospheric pressure. The center of heating zone was heated to 705 °C within 35 min. When the center of heating zone reached 705 °C, the temperature of the S powders was measured to be ~ 350 °C by using the thermocouple. The tube furnace was kept at 705 °C for 40 min, and cooled down to RT.
4.2.2 Materials Characterization
To transfer the as-grown crystals from a SiO2 substrate to a Cu TEM grid, an all- dry transfer method that employs a polydimethylsiloxane (PDMS) film was developed.
Briefly, the PDMS film was placed on the as-grown crystals on the SiO2 substrate and then peeled off. To ensure the transfer of the crystals from the SiO2 substrate to the
PDMS film, the film was slightly rubbed on by a q-tip. Next, the TEM grid was placed on the area of the film and then peeled off. FEI Titan G2 60-300 aberration-corrected STEM equipped with a CEOS DCOR probe corrector was used in this study. Annular dark-field scanning TEM (ADF-STEM) images (2048×2048 pixel2) were acquired with the STEM operated at 200 keV using a dwell time of 6 μs per image pixel at a camera length of 130 mm. The beam convergence angle, αobj, was measured to be 26 mrad. The ADF detector inner and outer angles of collection were measured to be 54 and 317 mrad, respectively.
The measured probe size was ~ 0.8 Å.
UPS characterization was carried out by using a XPS system equipped with a He I of UV source. The aperture for UPS has a diameter of 110 μm. Vacuum pressure was
77 kept below 3×10-9 Torr, and neutralizer was applied during the data acquisition. For UPS spectra comparison, commercial available bulk SnS2 crystals were used.
Microstructural analyses were done by SEM. Raman spectra were collected using a 532 nm laser (<2 mW excitation power, 100x objective lens). For Raman spectra comparison purposes, commercial available bulk SnS2 crystals were used. Thickness measurements were performed using AFM in tapping mode.
4.2.3 DFT Calculations
The calculations were based on the first-principles DFT and density functional perturbation theory using the projector augment wave method as implemented in the software package VASP. [42] For the electronic structure calculations, a Monkhorst-Pack scheme was adopted to integrate over the Brillouin zone with a k-mesh of 12 × 12 × 8 (8
× 8 × 1) for the bulk (few-layer) structures. A plane-wave basis kinetic energy cutoff of
500 eV was used. vdW interactions in bulk and few-layer SnS2 were accounted for using a semi-empirical correction to the Kohn Sham energies when optimizing each structure.
[43] The optimized lattice constants for bulk SnS2 and SnS were a=3.70 Å, c=11.85 Å and a=4.35 Å, b=3.99 Å and c=11.21 Å, respectively. This study used the Heyd-Scuseria-
Ernzerhof (HSE) hybrid functional to obtain accurate values for the bandgaps and ionization potentials for each phase. [44] The HSE calculations incorporated 25% short- range Hartree-Fock exchange and the screening parameter was set to 0.2 Å-1.
Calculations of the vibrational frequencies at Γ point relied on density functional perturbation theory.
78 4.3 Results and Discussion
2D tin sulfide crystals with lateral sizes over tens of microns were synthesized by atmospheric pressure vapor-phase growth method in a horizontal quartz tube furnace
(Figure 4.1(a)). In a typical synthesis, SnO2 and S powder precursors were located at the center and upstream of the tube furnace, respectively. SiO2 substrates were placed close to each other at downstream of the tube furnace. The 2D crystals were obtained on the
SiO2 substrates by heating the tube furnace to 705 °C and maintaining this temperature for 40 min before cooling down to room temperature (RT) (see ‘Experimental Section’ for the details).
79
Figure 4.1 The growth set-up, morphology and atomic structure of the SnS2 and SnS crystals. The growth setup (a) for the growth of SnS2 and SnS crystals. Optical microscopy images of SnS2 (b) and SnS (c) crystals grown on SiO2 substrates. Low magnification ADF-STEM images of SnS2 (d) and SnS (e) crystals transferred to a TEM grid. Filtered atomic-resolution ADF-STEM image of SnS2 (f) and SnS (g). Ball-and- stick models are overlapped to highlight atomic positions.
Optical images of the SiO2 substrates show that SnS2 crystals (Figure 4.1(b)) grow at a lower substrate temperature of ~ 425 °C while SnS crystals (Figure 4.1(c)) form at a higher temperature of ~ 550 °C. SnS2 crystals show a trapezoidal shape, and
SnS crystals have a rectangular shape.
80 Since SnS and SnS2 possess different crystal structures, atomic resolution STEM imaging can be used to directly verify the phases obtained from the growth. SnS2 and SnS crystals obtained by the previously described recipe were first transferred to separate
TEM grids and imaged in a STEM (see ‘Experimental Section’ for the details). Figure
4.1(d,e) show low magnification ADF-STEM images of SnS2 and SnS crystals, respectively. Figure 4.1(f,g) show filtered high-resolution ADF-STEM images of both phases obtained along the [001] zone axis where the crystal structure of the two phases can be distinguished. The image filtering procedure followed that used by Wu et al. [30]
SnS2, as seen in the ADF-STEM image, appears to crystallize in the 1T polymorph. This arrangement is similar to that observed for 1T-MoS2. [31, 32] On the other hand, the
ADF-STEM image of SnS shows its orthorhombic arrangement, which can be clearly differentiated from the image of SnS2.
The growth of the 2D crystals can be explained by the reaction of SnO2 and S.
When the temperature of the tube furnace is above the Tammann temperature of SnO2
(TTam= ~ 400 °C), SnO2 evaporation will be initiated, even though the applied temperature is lower than its melting point in powder form ( Tmelt= ~ 1350 °). [33, 34] At the early stage of the growth (~ 400 °C), SnO2 vapor from solid SnO2 powders makes contact with the substrate and forms small SnO2 clusters which act as nucleation sites for tin sulfide crystals. [35] SnO2 vapor from solid SnO2 powders reacts with S vapor to initiate the formation of SnS2 according to the equation (1) [3]
SnO2 (s) + 1.5S2 (g) → SnS2 (s) + SO2 (g) (4.1)
81 As the reaction temperature increases, the growth of SnS2 will eventually stop due to its phase instability on the high temperature. [2, 36, 37] At the same time, S source will be eventually depleted, which results in a reducing atmosphere in the tube. As a result, SnS growth will be favored according to the equation (2) [3]
SnO2 (s) + S2 (g) → SnS (s) + SO2 (g) (4.2)
To help further understand the underlying growth mechanisms of tin sulfide crystals, we performed a series of experiments by using three different growth recipes.
Recipe I served as the control experiment where the furnace with the quartz tube was heated from RT to 705 °C and then kept at this temperature for 40 min, followed by cooling down to RT. In recipe II, the furnace with the quartz tube was heated from RT to
625 °C, and immediately cooled down to RT without further heating. For recipe III, the quartz tube was initially held outside the furnace. While the furnace was being heated to
705 °C, the quartz tube was immediately slid inside the furnace at 625 °C. Then, the furnace with the quartz tube was kept at 705 °C for 40 min, followed by cooling down to
RT. Only one long rectangular SiO2 substrate was used for each experiment, and the growth parameters such as heating rate, gas flow rate, pressure, precursors, and substrates were kept constant for all the experiments.
Recipe I yielded both SnS2 and SnS crystals, but the two were spatially separated on the substrate (Figure 4.2(a)). SEM imaging (Figure 4.2(b,c)) and Raman spectroscopy (Figure 4.3(c,f,i)) confirmed that SnS was grown on the hotter side of the substrate compared to SnS2. On the other hand, recipe II resulted in primarily SnS2 crystals at the lower temperature side of the substrate while recipe III gave mainly SnS
82 crystals, which appears only in the hotter region. These results indicate that the growth of
SnS2 crystals is initiated mostly during the early stage of heating, while the formation of
SnS crystals starts in the later stage of the heating. Hence, the formation of SnS occurs at higher temperatures compared to that of SnS2, which is consistent with results of previous reports on the growth of SnS2 and SnS crystals. [2, 36, 37]
We also studied the trace reaction products of SnO2 (Figure 4.2(d)) and S found in the SnO2 crucible. It was found that recipe II produced yellow transparent SnS2 powders (Figure 4.2(e)), while recipes I and II yielded dark grey SnS powders (Figure
4.2(f)) as confirmed by Raman spectroscopy (Figure 4.3(c,f,i)). For recipe II, since the growth was performed at low temperature for a short period of time, SnO2 powders was sulfurized to form solid SnS2 as a result of the reaction according to the equation (1). For recipes I and III, as the reaction proceeds, solid SnS2 dissociation occurs as a result of the incongruent sublimation, which is promoted by both high temperature and S depletion.
[38] The dissociation of solid SnS2 results in the phase transition from SnS2 to SnS by the equation (3) [38]
SnS2 (s) → SnS (s) + 0.5 S2 (g) (4.3)
In this series of experiments, we also observed a growth of multi-stacked SnS2 crystals (Figure 4.2(g-i)). The interlayer S vdW interactions between consecutive SnS2 layers can result in misorientation of the layers during growth. 2D multistacked SnS2 crystals with varying configurations may possess some novel electronic and optical properties. Controlling stacking is beyond our scope in this study, but we point out that
83 this simple method can be further optimized in order to achieve the control over the stacking.
We performed Raman spectroscopy on the SnS2 and SnS crystals obtained by the discussed growth recipes. The Raman spectra of an 8L SnS2 crystal (~ 5.08 nm as measured in AFM with monolayer thickness of 0.6 nm [1]) show a strong Raman peak at
-1 314.9 cm , corresponding to the A1g phonon mode of SnS2 (Figure 4.3(a-c)). [39] In
-1 addition to the A1g mode, an intra-layer Eg mode with peak at ~ 206.2 cm is observed in multilayer (ML) SnS2 crystal with a thickness of ~ 17.94 nm (30L) (Figure 4.3(d-f)).
[39] The absence of the Eg mode in the 8L SnS2 is presumably due to the undetectably weak rejection of the Rayleigh scattered radiation. [8] The Raman spectra from a ML
SnS crystal (Figure 4.3(g)) was also measured. The thickness of the SnS crystal was estimated to approximately 30L (determined by cross-sectional STEM image in Figure
4.3(h)). The Raman spectra of the SnS crystal show a very different Raman pattern
(Figure 4.3(i)) with the peaks centered within the range of 50-250 cm-1. The peaks at
-1 -1 -1 95.6 cm , 192.0 cm , and 218.3 cm can be assigned to the Ag mode, and the peak at
-1 163.9 cm corresponds to the B3g phonon mode of SnS. [3, 40]
84
Figure 4.2 The photo, SEM and optical images of the crystals obtained by various recipes. Photo image (a) of a long SiO2 substrate after the growth. A stark boundary on the substrate surface of 2D SnS crystals in the high temperature region of the tube furnace to 2D SnS2 crystals in the cooler region can be seen. SEM images of an SnS2 crystal (b) and an SnS crystal (b) obtained from the surface of the SiO2 substrate in (a).
Photo images of SnO2 powders (d), SnS2 powders (e) obtained by recipe II, and SnS powders (f) obtained by recipes I and III. Optical images (g-i) of multi-stacked SnS2 crystals.
To support results of the experimental Raman spectroscopy, we calculated the zone-center vibrational modes of bulk SnS2 and SnS using ab-initio DFT (see
‘Experimental Section’ for the details). The vibrational modes at Γ point gave us insight
85 into the allowed Raman active modes for each phase. For bulk SnS2, we calculated the vibrational modes that correspond to the out-of-plane A1g and in-plane Eg Raman active
-1 -1 modes. The calculated values of 313 cm and 204 cm obtained for the A1g and Eg modes, respectively, are in good agreement with the Raman spectrum we obtained for the
2D SnS2 crystal (Table 1).
Table 4.1. Comparison of experimentally measured and ab-initio DFT calculated Raman
-1 spectra peak positions in SnS2 and SnS in cm .
Phases SnS2 SnS
Raman modes A1g Eg Ag Ag Ag B3g
Experimental 314.9 206.2 95.6 192 218.3 163.9
Calculated 313 204 95 196 216 163
For bulk SnS, we also calculated vibrational modes that correspond to the Ag and
B3g Raman active modes of the bulk orthorhombic structure of SnS. Based on the
-1 -1 displacements of these modes, we determined Ag modes at 95 cm , 196 cm , and 216
-1 -1 cm and the B3g mode at 163 cm . These values are also in good agreement with our assignment of the Raman modes using the measured Raman spectra (Table 1) and prior
Raman scattering studies on SnS. [40] The agreement between our theoretical calculations of the Raman modes for each phase of tin sulfides and our experimental
Raman spectra is further demonstration of the controlled growth we have over each phase. Moreover, we do not observe Raman peaks at 474 cm-1, 632 cm-1, or 774 cm-1
86 associated with SnO2. [41] This implies negligible chemical reaction between the crystals and substrate. [35]
Figure 4.3 Raman spectroscopy characterization of the SnS2 and SnS crystals.
Optical image, AFM image and Raman spectra of an 8L SnS2 crystal (a-c) and a 30L
SnS2 crystal (d-f). Optical image, cross-sectional ADF-STEM image and Raman spectra of a ~ 30L SnS crystal (g-i).
The ionization potential for each tin sulfide phase can be determined by UPS.
Using the He source at photon energies of 21.22 eV, we determined the ionization potential, φ, by subtracting the width of the UPS spectra from 21.22 eV. We determined the ionization potential of our as grown SnS2 and SnS crystals and compared them to the measurements on commercial SnS2 crystals. The ionization potential of our as-grown
SnS2 was found to be 7.51 eV (Figure 4.4(a)), close to that of commercial SnS2 crystal
87 (7.7 eV) (Figure 4.4(b)). The ionization potential of SnS was determined to be 5.78 eV
(Figure 4.4(c)).
Figure 4.4 The ionization potentials for the SnS2 and SnS crystals. UPS spectra acquired for our as-grown SnS2 crystal (a), commercial SnS2 crystal (b), and our as- grown SnS crystal (c).
Using DFT we evaluated the ionization potentials for SnS and SnS2 using an explicit slab model. 8L slabs of (001) SnS2 and (100) SnS were predicted to have bandgaps of 2.31 eV and 1.07 eV, respectively, which are close to the bandgap values for
88 the respective bulk structures. Calculations of each 8L structure were then used to obtain the difference between the average of the electrostatic potential in the bulk and the eigenenergies with respect to the vacuum potential. This provided an ionization potential of 7.51 eV and 4.9 eV for SnS2 and SnS, respectively, which is in good agreement with the ionization potential of 7.51 eV and 5.78 eV we obtain for SnS2 and SnS, respectively.
4.4 Conclusion
Highly crystalline 2D SnS2 and SnS crystals with lateral sizes over tens of microns have been directly synthesized on SiO2 substrates using a simple atmospheric pressure vapor-phase method. High-resolution ADF-STEM imaging reveals the phases of
2D hexagonal SnS2 (1T) and orthorhombic SnS crystals. The phase purity of SnS2 and
SnS crystals is confirmed by the analysis of the Raman and UPS spectra, and the results are corroborated by ab-initio DFT calculations. A series of growth experiments are performed to determine the growth process of the 2D crystals. It is found that temperature determines the phase of each crystal. SnS2 crystals grow at lower temperatures compare to the case of SnS crystals. Moreover, we have showed that the hexagonal SnS2 phase can be transformed into orthorhombic SnS phase upon high temperature treatment during growth. This temperature controlled process is advantageous for the purpose of avoiding highly toxic byproducts such as H2S in H2 reduction. Overall, these results are useful for future studies of the structural, electronic, and optical properties of 2D hexagonal SnS2 and orthorhombic SnS and may offer important guidance for realizing SnS2- and SnS- based electronic and optoelectronic devices.
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93 Chapter 5 Facile Synthesis of 2D SnS2 Nanowalls
Abstract
Two-dimensional (2D) layered metal chalcogenides (LMCs), especially tin sulfides, have recently received great interest due to their enticing physical and chemical properties and hold promise for various applications. We reported on synthesis of phase- pure 2D SnS2 nanowalls (NWs) by a facile vapor-phase synthesis method on insulator substrates such as SiO2 and MgO using SnO2 and S powders as precursors. The synthesized SnS2 NWs were characterized to study their fundamental properties by using various techniques such as scanning electron microscopy (SEM), x-ray diffraction
(XRD), Raman spectroscopy, x-ray photoelectron spectroscopy (XPS), and ultraviolet photoelectron spectroscopy (UPS). The synthesized films have an open network structure constituted of very uniform interconnected NWs with high crystallinity.
5.1. Introduction
Recently, 2D layered LMCs have been the subject of intense research due to their exotic physical and chemical properties. [1, 2] Tin sulfides are particularly interesting members of the family of LMCs. Their constituent elements, Sn and S, are inexpensive, environmentally friendly, and abundant in nature. Sn sulfides can exist in diverse crystal phases, such as orthorhombic SnS and hexagonal SnS2, owing to the oxidation characteristics of the elements Sn and S. [3]
SnS2 is a layered compound semiconductor with CdI2-type crystal structure. It is composed of layers of Sn atoms sandwiched between two closed-packed layers of S atoms. SnS2 has been reported to be an intrinsic n-type semiconductor. [3] SnS2 can exist
94 in various polytypes which differ in the stacking sequence of the sandwich layers. [4] The three common polytypes of SnS2 are 2H, 1T, and 4H. The distinctions between polytypes of SnS2 are mostly revealed in physical and electronic properties of the crystals. [5]
Figure 5.1 The growth set-up, thermal process and growth process for the SnS2
NWs. Illustration of experimental setup (a), thermal process (b), and synthesis process
(c).
Along with their interesting properties, nanostructured SnS2 crystals with different morphologies hold great promise for various applications ranging from energy storage to photocatalytic degradation. The high electrical conductivity and theoretical capacity of
SnS2 NWs make them a promising material for anodes of lithium (Li)-ion batteries
(LIBs). [6] The open spaces between neighboring SnS2 NWs can accommodate volume changes associated with insertion and extraction of Li ions and allow easy diffusion of electrolyte solution in LIBs. SnS2 nanostructures can be annealed to form tin dioxide
95 (SnO2), which opens avenues for applications such as gas sensing of chemicals such as ethanol and n-butanol. [7] It has also been shown that SnS2 nanosheets can be used as counterelectrode material for dye-sensitized solar cells owing to their excellent catalytic activity for triiodide reduction. [8]
To date, synthesis of SnS2 nanocrystals with various morphologies, such as nanoparticles, nanoplates, nanosheets, nanotubes, and nanoflowers, has mainly been based on reactions in liquid media or chalcogenization of thin films, and scale-up of current processes has been limited due to expensive or corrosive precursors. [9-13] To the best of our knowledge, large-area growth of SnS2 NWs on insulator substrates from powder precursors has not been reported yet.
Herein, we report a scalable facile vapor-phase synthesis method for growth of well-aligned SnS2 NWs on insulator substrates including silicon dioxide (SiO2) and MgO using SnO2 and S as vapor sources. The vapor-phase synthesis of SnS2 is based on the reaction of SnO2 vapor with S vapor at elevated temperature. The advantage of this method lies in its use of precursors that do not require conversion of a reactant solid phase to product. The fundamental properties of the SnS2 NWs have been studied by using various techniques. The results suggest that the grown films are indeed highly crystalline SnS2 with an open network structure constituted of interconnected NWs.
5.2 Experimental Section
5.2.1 Materials Synthesis
A facile vapor-phase synthesis method was used to synthesize SnS2 NWs on insulator substrates such as SiO2 (300-nm-thick SiO2 layer deposited thermally on Si
96 wafer) and MgO. S powder (99.5%, Alfa Aesar) and SnO2 powder (99.9%, Alfa Aesar) were selected as precursors. The reaction took place in a single-zone vacuum tube furnace with a 1-inch quartz tube (MTI OTF-1200X-S50-2F) (Figure 5.1(a)). High- purity Ar gas at constant flow of 60 sccm was used as carrier gas. In a typical growth process, 25 mg SnO2 powder was put in a polished quartz boat at the center of the heating zone, and 250 mg S powder was put in another polished quartz boat at the upstream. The substrates (~ 1x1 cm2) were cleaned by sonication in acetone and isopropyl alcohol (IPA) baths, sequentially for 1 min each. Immediately following the cleaning procedure, the substrate was placed at the downstream of the furnace. The furnace was pumped down to remove air, and the pressure was stabilized at ~ 0.1 Torr. Next, the center of the heating zone was heated to 710 oC within 30 min. The furnace was kept at 71 oC for 10 min, then naturally cooled down to room temperature (Figure 5.1(b)).
5.2.2 Materials Characterization
The morphology of the NWs was investigated using scanning electron microscopy (SEM, FIB NNS450). X-ray diffraction (XRD) studies were carried out using a Philips X’Pert diffractometer with Cu Ka radiation in Bragg-Brentano geometry.
Raman spectra were collected using a Horiba system with a 532-nm laser (<2 mW excitation power, x100 objective lens). All Raman spectra were calibrated using the
520.5 cm-1 line of a Si wafer. For Raman spectra comparison purposes, commercially available bulk SnS2 crystals (2D Semiconductors Inc.) were used. X-ray photoelectron spectroscopy (XPS) and ultraviolet photoelectron spectroscopy (UPS) characterization were carried out using a Kratos AXIS ULTRADLD XPS system equipped with an Al Ka
97 monochromated x-ray source, He I ultraviolet (UV) source (21.22 eV), and a 165-mm mean radius electron energy hemispherical analyzer. The slot size for XPS was 300 µm x
700 µm, and the aperture for UPS had a diameter of 110 µm. The vacuum pressure was kept below 3 9 10-9 torr, and the neutralizer was applied during data acquisition. The C 1s peak at 284.8 eV was used for binding energy calibration.
5.3 Results and Discussion
The formation of SnS2 NWs includes the following three major steps (Fig. 1c):
The first step involves sublimation and reaction of S and SnO2 to produce gaseous SnS2
o species. As the temperature is about 425 C, which is the Tammann temperature of SnO2, evaporation of SnO2 will be initiated, even though the applied temperature is lower than its melting point. [14] S also starts evaporating when the temperature at the center of the
o tube is 425 C. As a result, SnO2 vapor will react with S vapor to form SnS2 species. In the experiments, we used excess amount of S (mass ratio of S to SnO2 of about 10) to ensure far more vapor of S than of SnO2. As a result, we can reasonably assume that the reaction of SnO2 is complete. Next, SnS2 species will be transferred to the downstream by
Ar gas. The final step includes diffusion of SnS2 species from the gas phase onto the receiving substrates and finally precipitation.
98
Figure 5.2 The morphology of the SnS2 NWs. Low- and high-magnification SEM images of SnS2 NWs/SiO2 (a, b) and SnS2 NWs/MgO (c, d).
Low-magnification (Figure 5.2(a)) and high-magnification (Figure 5.2(b)) SEM images of synthesized SnS2 NWs on SiO2 substrates (SnS2 NWs/SiO2) reveal that the
SnS2 films have an open network structure constituted of interconnected, well-ordered
NWs. The thickness of the NWs and the film is estimated to be several tens of nanometers and a few microns, respectively, from the high-magnification SEM image of the NWs (Figure 5.2(b)). To further examine the role of the interaction with the substrate, we studied growth on different receiving substrates such as SiO2 and MgO. We grew SnS2 films on these substrates under identical experimental conditions. The morphology of SnS2 NWs grown on MgO substrates (SnS2 NWs/MgO) (Figure 5.2(c))
99 was found to be similar to that of SnS2 NWs/SiO2 (Figure 5.2(a)). However, the size of the SnS2 NWs/MgO was noticeably smaller (Figure 5.2(d)).
Figure 5.3 XRD and Raman spectroscopy analysis of the SnS2 NWs. XRD pattern of
SnS2 NWs/SiO2 (a). Reference data are also shown. Raman spectra of SnS2 NWs/SiO2
(b), SnS2 NWs/MgO (c), and commercially available SnS2 crystals (d). The left insets in
(c, d) show the magnified Eg peaks.
The crystallographic structure of SnS2 NWs/SiO2 was determined by XRD analysis. In the XRD pattern (Fig. 3a), the (001) (100), (101), (102), (003), (110), (111),
(103), (201), and (202) diffraction peaks of SnS2 NWs are observed at 2θ angles of
15.029o, 28.200o, 32.125o, 41.887o, 46.122o, 49.961o, 52.452o, 54.962o, 60.621o, and
o 67.154 , respectively. All the diffraction peaks can be indexed to hexagonal SnS2 [Joint
Committee on Powder Diffraction Standards (JCPDS) card file no. 23-677], indicating
o synthesis of pure SnS2. [15] The strong and sharp peak at 2θ = 15.029 indicates preferred growth orientation along the (001) plane of hexagonal SnS2. Such observation of preferred growth orientation of the SnS2 crystals along the (001) plane can be
100 explained by considering the surface energies of SnS2 crystal planes. [16] Surface energy theory suggests that the equilibrium form should be such that the crystal has minimum surface energy per unit volume. As reported in literature, the (001) plane of SnS2 may possess the lowest surface energy and thus adsorb more SnS2 species than the other planes of SnS2 during growth. [17]
Raman spectroscopy has proven to be a versatile tool for studying various 2D materials. [1] We performed Raman spectroscopy characterization on SnS2 NWs/SiO2,
SnS2 NWs/MgO, and commercially available SnS2 crystals (for comparison purposes).
The two typical Raman-active modes of SnS2 NWs/SiO2 appear in the spectrum (Figure
-1 5.3(b)), namely a strong A1g Raman mode at 313.3 cm and a weak Eg Raman mode at
-1 204.4 cm , confirming that the synthesized NWs are indeed SnS2. [18] The intensity of the Eg peak is much lower than that of the A1g peak, presumably due to the reduction in the scattering centers for in-plane scattering. [3,19] Similar to SnS2 NWs/SiO2, the
-1 Raman spectra (Figure 5.3(c)) of SnS2 NWs/MgO show the A1g peak at 313.8 cm and
-1 the Eg peak at 204.8 cm , indicating a negligible effect of the substrate on the Raman features of the SnS2 NWs. We also performed Raman spectroscopy on commercially available SnS2 crystals for comparison purposes. The Raman spectra (Figure 5.3(d)) of
-1 -1 SnS2 crystals show the A1g peak at 313.8 cm and the Eg peak at 204.9 cm , being very similar to those of SnS2 NWs/SiO2 and SnS2 NWs/MgO. Thus, we can conclude that the synthesized SnS2 NWs are similar in quality to SnS2 crystals.
101
Figure 5.4 XPS and UPS analysis of the SnS2 NWs. Survey XPS (a) and high- resolution XPS spectra for Sn 3d (b) and S 2p (c) binding energies of SnS2 NWs. (d) He-I
UPS spectra of SnS2 NWs.
XPS was used to investigate the chemical surface composition of the NW samples. The survey XPS spectrum (Figure 5.4(a)) of the samples revealed two strong peaks indicating presence of Sn and S. Oxygen, whose O 1s peak is located at ~ 532 eV
(Figure 5.4(a)), was only present in negligible amounts. The high-resolution spectra of the samples were used to determine the oxidation states of the elements. Two peaks appear for each core level due to spin-orbit splitting of Sn 3d (Figure 5.4(b)) and S 2p
(Figure 5.4(c)) core levels. The binding energies for Sn 3d3/2 and Sn 3d5/2 are 495.7 eV and 487.3 eV, respectively, with energy separation of 8.4 eV, being near the value
IV reported for S in SnS2. [20] The S 2p binding energies of the SnS2 NWs are 163.6 eV
102 and 162.4 eV, for S 2p1/2 and S 2p3/2, respectively. The results demonstrate that no elemental sulfur (164.05 eV) is present in the product. [20] We also used UPS to determine the ionization potential of the SnS2 NWs. The ionization potential was measured to be 6.66 eV, which is close to the ionization potential value of SnS2 crystals reported by Schroeder et al. [21] This result further confirms that the crystals are indeed
SnS2.
5.4 Conclusion
We demonstrated facile, large-scale synthesis of highly oriented SnS2 NW films on insulator substrates by reaction between SnO2 and S vapors. The SnS2 NWs exhibited high crystallinity with preferred growth orientation along (001) plane and an open network structure constituted of interconnected NWs. Such highly crystalline SnS2 NWs may find applications in energy storage, solar cells, sensors, and catalytic applications where specific surface area plays an important role in overall productivity.
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105 Chapter 6 Synthesis of MoSe2 via Rapid Thermal Processing and Laser Annealing
Abstract
Few-layers MoSe2 granular films were prepared using rapid thermal processing
(RTP) of stacked elemental layers (SELs) deposited using electron beam evaporation in the sequence of Mo/Se/Mo. RTP was carried out at various temperatures ranging from
550 to 750 °C, and durations ranging from 90 to 150 s. Morphology and microstructure of granular thin films were characterized by atomic force microscopy (AFM) and scanning electron microscopy (SEM). Compositional and electronic structure analysis of the films were done by X-ray photoelectron (XPS), photoluminescence (PL) and Raman spectroscopy. The results confirmed the formation of four-layer MoSe2 granular thin films at 750 °C for 90 s. Density functional theory (DFT) calculations showed that the direct gap at K is near degenerate with the calculated indirect band gap in the 4L structure. In addition to RTP, Raman laser annealing (RLA) was performed on SELs with various laser powers. Raman spectroscopy results revealed the local formation of MoSe2 with increasing laser energy.
6.1 Introduction
Semiconducting phases of transition metal dichalcogenides (TMDs) have attracted great attention recently because of their fascinating physical properties. [1]
MoSe2 is a semiconductor phase of TMDs and crystallizes in a strongly covalent bound
Se-M-Se sandwich structure (Figure 6.1(a)) Consecutive Se-Mo-Se layers are weakly bound by van-der-Waal (vdW) interaction. MoSe2, with an indirect band gap of 1.1 eV in
106 the bulk form and a direct band gap of 1.55 eV in the monolayer (ML) form, exhibits tunable electronic properties, allowing for the realization of its potential in optoelectronics and spintronics. [2-4] The potential applications of MoSe2 have motivated the development of large scale synthesis of atomically thin MoSe2 films.
Figure 6.1 The ball-and-stick model and synthesis process for MoSe2. The ball-and- stick model (a) of MoSe2. The schematic (b) of SELs deposited on SiO2/Si substrate. The schematic view (c) of RTP system working principles. The illustration (d) of the thermal process performed for MoSe2 growth.
Recent top-down approaches such as mechanical and liquid exfoliation to obtain large area MoSe2 thin films have received considerable attention. [5, 6] Nevertheless, the lack of uniformity in thickness and size undermines the viability of such approaches.
Selenization of molybdenum trioxide (MoO3), which is a conventional bottom-up approach for synthesis of MoSe2, has been recently performed to grow MoSe2 films in a
107 chemical vapor deposition (CVD) system. Making use of the elemental Se and MoO3 powders' low melting and evaporation temperatures, Shaw et al. and Lu et al. have demonstrated that elemental Se and MoO3 are suitable precursors for CVD of MoSe2 thin films. [7, 8] Although CVD methods are promising techniques for future efficient production of MoSe2, current production status of MoSe2 by the aforementioned methods is still in a nascent stage. Alternative MoSe2 synthesis techniques have focused on the direct selenization of Mo-containing films. Pouzet and Bernede have synthesized MoSe2 films by solid phase selenization of Mo thin films in a CVD furnace. [9] However, each of these growth techniques is hampered by the formation of thick films, including perpendicularly grown three-dimensional (3D) structures, with small grain sizes and poor crystallinity. Alternatively, Lee et al. have reported on the synthesis of thick MoSe2 films during rapid thermal processing (RTP) of Mo-coated glass under Se pressure. [10] We extend this study by reporting on the formation of few-layers (FL) MoSe2 granular films following RTP of stacked Mo/Se/Mo layers.
Raman spectroscopy is widely used as a very efficient and non-destructive method to study two-dimensional (2D) materials such as graphene and TMDs. [11-13]
Recently, laser annealing using Raman spectroscopy has been performed for thinning of
2D materials. Gomez et al. have reported on the laser-thinning of multilayer MoS2 down to a ML 2D crystal using Raman spectroscopy. [14] An optical microscope-focused laser beam system has been used to construct reduced multilayer graphene oxide sheets. [15]
Laser direct writing of graphene patterns on SiO2 has been shown by Fan et al. [16]
108 However, the formation of MoSe2 from stacked elemental layers (SELs) using Raman laser annealing (RLA) has not been reported.
In this study, we have successfully synthesized FL MoSe2 granular films using
RTP of SELs in the sequence of Mo/Se/Mo on SiO2/Si. The morphology of the films has been characterized by optical microscopy, AFM and SEM. The phase composition of the films has been studied by XPS. PL and Raman spectroscopy have been performed to investigate the optical and electronic properties of the films. Our PL data has been compared to the results of electronic band structures of ML and four layer (4L) MoSe2 using first-principles DFT calculations. Moreover, we have also investigated the formation of MoSe2 from SELs by RLA using Raman spectroscopy.
6.2 Experimental Section
6.2.1 Materials Synthesis
Films subjected to RTP were prepared using e-beam evaporation of 0.5 nm Mo,
34 nm Se, and 0.5 nm Mo films (Figure 6.1(b)) on p-type (100) oriented Si wafers
2 (∼1 cm ) with 300 nm thermal SiO2, which was previously cleaned by sequential sonication in toluene, acetone and isopropyl alcohol (IPA) for 30 min, respectively. The purpose of the Se layer between two Mo layers was to prevent the loss of Se during the thermal step. The deposition was carried out in a load-lock chamber at a base pressure of
∼10-7 Torr and deposition pressure of ∼10-6 Torr, with an average growth rate of 0.1 Å/s.
Before deposition, the vacuum chamber was completely covered with aluminum foil to prevent Se contamination on the walls of the chamber. After deposition, the foil was coated by a thicker nickel layer, and properly disposed of.
109 RTP of deposited films was then performed in a RTP system at various temperatures (550 to 750 °C) and durations (90 to 150 s). The temperature of the system was controlled by a pyrometer, exposed to the back of the sample holder (a 4″ Si wafer), and calibrated by means of a thermocouple in contact with it (Figure 6.1(c)). The RTP system was ramped to 550-750 °C with a heating rate of ∼100 °C/s, and the temperature was maintained for 90-150 s. Annealing was followed by a rapid cooling step down to
∼200 °C with a cooling rate of ∼100 °C/s, and then to room temperature (RT) by a slow cooling rate of 20 °C/s. The purpose of the rapid heating and cooling steps was to prevent loss of Se and induce the recrystallization process. The thermal cycle is illustrated in
Figure 6.1(d) for clarity. The MoSe2 flakes were mechanically exfoliated from the bulk
MoSe2 crystals by a conventional scotch tape method. [2]
6.2.2 Materials Characterization
AFM imaging and thickness measurements were done in tapping mode.
Microstructural analysis was performed by SEM. The compositional analysis of the films was done by a XPS system equipped with an Al Kα monochromated X-ray source and a
165-mm electron energy hemispherical analyzer. The vacuum pressure was kept below
3 × 10-9 Torr and a neutralizer was applied during the data acquisition. Raman and PL spectra of the MoSe2 film and MoSe2 flake were collected using a 532 nm laser (100x objective lens) at RT. The laser beam, with a laser power of ∼0.06 mW, was focused on a
0.8 µm diameter spot. The effect of laser power on RLA process of SELs was also investigated by varying the laser power from 0.0065 mW to 32.5 mW. All measurements were performed at RT and in air.
110 6.2.3 DFT calculations
Our calculations were based on first-principles DFT using the projector augmented wave method as implemented in the software package VASP. [17] Standard approximations of DFT at the LDA or GGA level of theory are known to underestimate the band gaps in the TMD family of materials. To account for this, the screened Heyd–
Scuseria–Ernzerhof (HSE) hybrid functional was employed for this study. A Monkhorst-
Pack scheme was adopted to integrate over the Brillouin zone with a k-mesh 9x9 x1 for the ML and FL structures. A plane-wave basis kinetic energy cutoff of 280 eV was used.
Spin orbit coupling was included self-consistently. The lattice constants for the FL
MoSe2 structures were obtained from a volume optimized bulk 2H-MoSe2 structure. The atomic coordinates for the thin film structures were optimized at this fixed lattice constant. Spin-orbit coupling (SOC) was included self-consistently within the band structure calculations. For the HSE calculations, 25% short-range exact Hartree–Fock exchange was used with the Perdew-Burke-Ernzerhof (PBE) correlation. The HSE screening parameter, µ, was empirically set to 0.3 (1/Å).
6.3 Results and Discussion
AFM was used to characterize the surface morphology and thickness of MoSe2 granular thin films. Figure 6.2(a-d) show the AFM images of MoSe2 films on SiO2/Si after annealing at 550 °C for 90 s (referred to as sample A), 750 °C for 90 s (referred to as sample B), 750 °C for 120 s (referred to as sample C) and 750 °C for 150 s (referred to as sample D). The surface of all the samples appears to be granular.
111
Figure 6.2 The AFM height images of the MoSe2 films at 550 °C for 90 s (a), 750 °C for 90 s (b), 750 °C for 120 s (c) and 750 °C for 150 s (d).
Figure 6.3 show the histogram of the heights of the MoSe2 islands measured from the AFM images (Figure 6.2(a-d)). The heights of MoSe2 islands from the substrate are
2.75-3.25 nm, 3-4 nm, 6.25-8 nm, and 15.5-20 nm for sample A, B, C, and D, respectively. The height of the MoSe2 islands increases as the annealing temperature is raised. For the annealing step at 750 °C, the density of the number of islands decreases and the height of the islands increases with increased annealing time, possibly due to particle agglomeration. [18]
112
Figure 6.3 The histograms of the particle size distribution of the MoSe2 films. AFM height histograms of MoSe2 films at 550 °C for 90 s (a), 750 °C for 90 s (b), 750 °C for
120 s (c), and 750 °C for 150 s (d).
Figure 6.2(b) shows that there are some flat MoSe2 regions with a length of
∼300 nm on the surface of sample B. The height of the flat regions is determined to be
∼3.5 nm for sample B. This corresponds to a thickness of four monolayers (4L) assuming each monolayer is approximately 9 Å. [2] The surface roughness values, Ra and Rms, are determined to be 1.2 and 1.48 nm, respectively, for sample B. The surface roughness is relatively higher than expected since defects and contamination are possibly introduced during the rapid heating and cooling process.
113
The AFM phase images (Figure 6.4) of the samples reveal that the surface of the
SiO2/Si substrate is not completely covered by MoSe2 possibly due to the loss of Se and
Mo and agglomeration during the rapid heating and annealing. [19] It has been reported that annealing temperatures higher than 700 °C lead to enhanced layer growth of MoSe2.
[18]
Figure 6.4 AFM phase images of the MoSe2 films at 550 °C for 90 s (a), 750 °C for 90 s (b), 750 °C for 120 s (c) and 750 °C for 150 s (d).
Figure 6.5 shows the optical microscopy images of sample B with a scratch before and after RTP. Since the color contrast of MoSe2 on SiO2/Si strongly depends on the oxide thickness of the substrate, lens magnification, and thickness of the film, surface
114 sensitive spectroscopy techniques are essential to characterize the quality of the synthesized films.
Figure 6.5 The optical images of the Mo/Se/Mo and MoSe2 films.
It is well-known that XPS is a very useful method for the determination of chemical compositions and chemical states of material surfaces. The chemical composition of sample B was further investigated by XPS analysis (Figure 6.6). Mo 3d5/2 and Mo 3d3/2 peaks are found at 229.2 and 232.3 eV, confirming that Mo is in its Mo (IV) state; and the 3d peak of Se element is split into the well-defined 3d5/2 and 3d3/2 peaks at
54.8 and 55.7 eV, respectively (Figure 6.6(a,b)). [20, 21] The stoichiometric ratio
(Se/Mo) estimated from the respective integrated peak area of the XPS spectra is approximately 2. This result further confirms the formation of MoSe2. More detailed analysis results reveal some contributions from MoO3, exhibiting two characteristic peaks at 236.2 and 233.4 eV, attributed to the 3d3/2 and 3d5/2 binding energies for the Mo (VI) state. [22]
High-resolution XPS with peak deconvolutions of the O 1s binding energy and Si
2p binding energy spectrum of sample B are given in Figure 6.6(c,d). The small peak at
115 531.1 eV is associated with MoO3. [22] MoO3 formed during the growth process may be due to the presence of trace amounts of O2 in the system.
Figure 6.6 XPS analysis of the MoSe2 films. High-resolution XPS with peak deconvolutions of the Mo 3d binding energy (a), Se 3d binding energy (b), O 1s binding energy (c), Si 2p binding energy (d) spectrum of the MoSe2 film synthesized at 750 °C for 90 s.
Raman spectroscopy, a powerful and non-destructive tool for characterizing
TMDs, was employed to study the lattice dynamics of sample B. [23-25] In Figure
6.7(a), we compare the Raman spectra of the as-deposited, post-RTP and exfoliated
MoSe2 flakes. The Raman spectra reveal that the MoSe2 film has two characteristic A1g
(out-of-plane) and inline image (in-plane) Raman modes, confirming the formation of
-1 MoSe2. [26] The A1g and the inline image modes are identified at ∼243.0 and 283.7 cm , respectively, which is consistent with prior studies. [13] Moreover, the wavelengths and intensity ratios of the A1g and the inline image peaks are similar in both the resultant
116 MoSe2 and exfoliated MoSe2 flake, showing that optical properties of the grown samples are comparable to those of the exfoliated single crystal MoSe2 flakes. The details on
MoSe2 flakes are given in Figure 6.7(c,d).
Figure 6.7 Raman-PL spectroscopy characterization of the MoSe2 film. Raman (a) and PL (b) spectra comparison of the as-deposited MoSe2, post-RTP MoSe2 film (sample
B) and exfoliated MoSe2 flake. The optical image (c) and corresponding AFM height image (d) of the exfoliated MoSe2 flake.
PL spectroscopy was used to measure the quality and the thickness of sample B
[4, 13]. Figure 6.7(b) illustrates the PL spectra comparison of the as-deposited, post-
RTP, and exfoliated MoSe2 flakes. It has already been reported that ML MoSe2 has a very sharp and intense PL peak at 1.55 eV, corresponding to a direct band gap. [2, 7] We
117 have also observed a peak at ∼1.55 eV. However, the intensity of the peak is low and broadened, which is an indication of an indirect band gap in the resultant MoSe2. [7]
The PL intensities and energies are close in both sample B (1.55 eV) and exfoliated MoSe2 flakes (1.57 eV), showing the PL properties of the grown samples are comparable to those of exfoliated single crystal MoSe2 flakes. Although the 4L has an indirect band gap the presence of a PL peak at 1.50-1.60 eV in the few-layers MoSe2 is quite interesting, and can be explained by the DFT calculations. [13]
To gain further insight, we compute the band structure of 1L and 4L MoSe2 using ab-initio DFT as shown Figure 6.8. Our prior studies on the electronic structure and lattice parameters of bulk MoSe2 are good in agreement with prior experimental and theoretical studies. [27] The lattice parameters from this study are used for our simulations on 1L and 4L MoSe2. Our calculations show monolayer MoSe2 is a direct gap semiconductor at the K symmetry point with a band gap of 1.624 eV. For the 4L structure, the band gap becomes indirect with the valence band maxima shifted to Г with a band gap of 1.600 eV. In the 4L structure, the direct gap at K is near degenerate with the calculated indirect gap, with a band gap of 1.605 eV.
Figure 6.8 DFT-calculated electronic band structures of MoSe2. DFT calculations of the 1L (red) and 4L (blue-broken) MoSe2.
118 The effect of laser-induced heating on various materials such as carbon nanotubes, graphene, and silicon has been studied in the past two decades. [28, 29]
Herein, we have studied the formation of MoSe2 from SELs by RLA using a Raman spectroscopy system (Figure 6.9(a)). The Raman spectra variation of SELs has been examined when the laser power is tuned to 0.0065 mW (P1), 0.065 mW (P2), 0.65 mW
(P3), 6.5 mW (P4), 16.25 mW (P5), and 32.5 mW (P6). Figure 6.9(b) shows the Raman spectra comparison of the films with various laser powers. For annealing with P1 and P2, no Raman peaks of MoSe2 are observed. As the laser power is increased, the Raman features of the MoSe2 and substrate become clear. For annealing with a laser power of
P6, the A1g and the inline image peaks are clearly observed, confirming the formation of
MoSe2. The A1g and the inline image peaks are very similar to those of exfoliated MoSe2 flakes. [2] However, we did not observe any PL possibly due to high thickness and poor crystallinity which is expect to be improved by tuning the thickness of the SELs and laser parameters. The intensity of the Si peak is also increased by the increased laser power, which is an indication of the heating and thinning effect of the laser beam. [14] Beyond
P6, the MoSe2 layer is completely removed from the substrate.
119
Figure 6.9 The RLA setup and laser-power dependent Raman spectra. The schematic diagram of RLA setup (a). Raman spectra of the films annealed at different laser powers (b).
Raman mapping has also been performed to form MoSe2 from SELs with certain patterns on SiO2/Si wafers as illustrated in Figure 6.10. The results reveal the potential of the laser at wavelength of 532 nm in forming of MoSe2 patterns from the stacked layers when the laser power is adjusted. Our goal is to develop a laser technology that allows for precise formation and structuring of atomically thin MoSe2 films on SiO2/Si wafers.
120
Figure 6.10 The formatıon of MoSe2 from SELs with certain patterns on SiO2/Si.
Optical microscopy image (a) of a Mo/Se/Mo film deposited on SiO2/Si without heat treatment. Optical microscopy image (b) of a film after scanning a laser with incident laser power of 32.5 mW in the area marked by a dashed square (10x10 points). MoSe2 and Mo/Se/Mo regions are marked with 2 in the blue-circle and 1 in the black-circle, respectively. Optical microscopy image (c) of a film after scanning a laser with incident laser power of 32.5 mW in the area marked by a dashed square (30x30 points). MoSe2 and Mo/Se/Mo regions are indicated with blue and black arrows, respectively. The scale bar in all images is 5 µm. Raman spectra (d) of a MoSe2 region on the film marked by a blue arrow in (c).
6.4 Conclusion
In summary, we have developed a procedure to fabricate few-layers MoSe2 granular films. From their Raman and photoluminescence spectroscopy analyses, we
121 conclude that the semiconducting properties of rapid thermal processed few-layers granular films of MoSe2 resemble that of single crystal MoSe2 flakes. This method could now be implemented for the synthesis of other dichalcogenides materials. Moreover, we have shown that RLA can be employed to produce MoSe2 in arbitrary shapes and patterns on SELs. Following the optimization of the thickness of SELs and laser parameters such as wavelength, power and exposure time, the quality of the laser processed MoSe2 can be improved.
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124 Conclusion
Herein, we have demonstrated various novel growth strategies for the production of several 2D materials, including growth of group VI transition metal dichalcogenides
(TMDs) by chalcogenization of pre-deposited metal-containing precursors, and phase- engineered growth of group IV LMCs by vapor-phase reaction of metal oxide and chalcogen powders are demonstrated. Detailed characterization of each phase of tin sulfide crystals was performed using various microscopy and spectroscopy methods, and the results were corroborated by ab-initio density functional theory calculations. The significant results of this dissertation as following:
1. The large-scale production of 2D crystals can be obtained via by chalcogenization
of pre-deposited metal-containing precursors.
2. Vapor-phase reaction of metal oxide and chalcogen powders with the help of the
patterned seeds can provide ultrathin thin films with controlled thickness and
morphology.
3. Crystal phase engineering of the crystals can be achieved by fine tuning of the
reaction temperature.
4. Air annealing can be utilized to open active edges with well-defined shapes on the
2D crystals.
5. Laser annealing can be used to make nanopatterns of 2D metal chalcogenides on
SiO2/Si chips.
125