Chemical Vapor Deposition of Two-Dimensional Crystals

19

Chemical Vapor Deposition of Two-Dimensional Crystals Zachary R. Robinson1, Scott W. Schmucker2, Kathleen M. McCreary2, Enrique D. Cobas3 1 AS EE POSTDOCTORAL FELLOW R ESIDING AT U.S. NAVAL RE SEARCH LABORATORY, SW WASHINGT ON, D C, USA; 2 NR C P OSTDOCTORAL FELLOW R ESIDING AT U.S. NAVAL RESEARCH LABO RA TORY , S W WASHINGTON, DC, US A; 3 MA TERIALS S CIENCE AND TECHNO LOGY DIVISION, U.S. NAVAL RESEARCH LABORATORY, SW WASHINGTON, DC, USA

CHAPTER OUTLINE 19.1 Introduction ............................................................................................................................ 785 19.2 Chemical Vapor Deposition of Graphene ........................................................................... 787 19.2.1 Graphene Growth on Ni............................................................................................ 790 19.2.2 Graphene Growth on Cu........................................................................................... 792 19.2.3 Graphene Growth on Cu Single-Crystal Substrates ................................................ 801 19.2.4 Graphene Growth on Iridium ................................................................................... 807 19.2.5 Graphene Transfer ................................................................................................................. 808 19.3 Chemical Vapor Deposition of Hexagonal Boron Nitride ................................................. 809 19.3.1 Precursors .................................................................................................................... 811 19.3.2 Monolayer Hexagonal Boron Nitride....................................................................... 814 19.3.3 Substrate Interactions ................................................................................................ 815 19.3.4 Other Deposition Methods ....................................................................................... 815 19.4. Chemical Vapor Deposition of Molybdenum Disulfide..................................................... 817 References........................................................................................................................................... 824

19.1 Introduction The experimental isolation of graphene in 2004 surprised the scientific community and unequivocally proved the existence and stability of atomically thin crystals [1], which had been theoretically predicted to be thermodynamically unstable [2,3]. It has been called the “mother of all graphitic forms,” because it is the structural building block upon which many carbon materials are based [4]. Handbook of Crystal Growth. http://dx.doi.org/10.1016/B978-0-444-63304-0.00019-6 Copyright © 2015 Elsevier B.V. All rights reserved.

785

786

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.1 Schematic representation of several two-dimensional crystals. (a) Molybdenum disulfide: sulfur atoms shown in yellow, molybdenum in teal. (b) Hexagonal boron nitride: boron shown in blue, nitrogen in peach. (c) Graphene: with carbon atoms in gray. Image generated by Dr Alejandro Suarez using Materials Studio software, used with permission.

In other ways, graphene is the mother of all two-dimensional (2D) materials. For decades, graphene was considered as a theoretical model used to better understand the interactions of stacked layers in bulk graphite crystals [5]. Its isolation has motivated investigations into a whole new class of materials: 2D crystals. This rapidly growing class of materials includes graphene, hexagonal boron nitride (hBN), and transition metal dichalcogenides (TMDs). The material properties exhibited by 2D crystals can vary widely, from wide band gap insulators (such as hBN) to semiconductors (such as TMDs) to zero gap semimetals (graphene). Despite this variety, all 2D crystals have strong intralayer bonds while concurrently demonstrating weak interlayer bonding. It is this weak van der Waals bonding that allows a single layer to be peeled away, or mechanically exfoliated, from the bulk material to produce small, high-quality 2D crystals. The structure of several common 2D crystals is shown schematically in Figure 19.1. Mechanical exfoliation is a quick and effective means to produce individual crystals for fundamental investigations. Unfortunately, this technique limits production to micron-sized fragments, which must be painstakingly sought out. In contrast, chemical vapor deposition (CVD) is a technique capable of providing large-area, scalable, costeffective and high-quality crystalline material while also being compatible with standard semiconductor processing techniques [6]. The synthesis of 2D crystals using CVD has seen rapid growth and success in recent years and continues to improve and evolve. In this chapter, we will review CVD-based growth of 2D crystals with emphasis on the various growth techniques, growth mechanisms, substrates, and resulting film qualities.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 787

The atomically thin structure of 2D crystals dictates that it is comprised entirely and inherently of surface. As a result, many characteristics that become boundary conditions or imperfections in three-dimensional crystals compose the fundamental nature of 2D crystals. The techniques used to characterize these materials therefore utilize surface science instruments such as X-ray photoelectron spectroscopy, low-energy electron diffraction (LEED), and scanning tunneling microscopy (STM) that are very effective at studying surfaces and interfaces of thin films. These characterization techniques have been developed over decades, and they have been used to study and understand the impact of crystal structure on various growth and adsorption processes in situ. In fact, the development of modern surface science techniques paralleled the early development of graphene growth. As a result of this, there are several studies that predate the Nobel prize-winning isolation of graphene. Observations of graphene growth on transition metals extend back to 1969, although graphitic contamination is likely as old as metallurgy. In early studies, accidental carbon deposition on and precipitation from metal substrates resulted in graphene formation. In many cases it was even noted that a 2D layer of “monolayer graphite” had formed on the metal substrate [8–11]. The films were studied initially because they were an irritation, as when their presence reduced the catalytic activity of metal surfaces [12]. Following the development of LEED, researchers observed segmented ring structures in the LEED patterns of platinum and ruthenium surfaces after heating [13,14]. Although originally attributed to a disordered surface structure, by 1968 these patterns were revealed to be carbon-based and reproducible upon hydrocarbon dosing of the metal surface at elevated temperature [15]. The final conceptual leap came from May in 1969, who identified the signature of rotationally disordered, polycrystalline monolayer graphite in these LEED rings [16]. CVD of graphene on Pt surfaces was better understood after the development of the STM, which allowed for imaging of monolayer graphene islands stitching into the polycrystalline film (Figure 19.2) [7]. In addition to this chapter, there are a number of thorough review articles on the topic of 2D material growth. One such article focuses on the surface science of graphene growth, and thoroughly explains the details of the growth on many different metal substrates [17]. Another thorough review was published in Surface Science, and can be found in reference [18]. Much of the knowledge and many of the techniques realized from graphene synthesis provide a strong foundation for the more nascent areas of hBN [19,20] and TMD growth [21–24]. As the focus of this chapter is the growth of 2D crystals, discussion of the electronic, thermal, optical, and mechanical properties that motivate much of the growth will be omitted. The material properties of graphene, hBN, and TMDs have been the topic of extensive review articles to which the interested reader is directed [4,23,25–29].

19.2 Chemical Vapor Deposition of Graphene Since the discovery that 2D materials could be isolated for electronic applications, the material that has received the most attention from nearly every perspective is graphene. Graphene’s atomic structure is composed of a hexagonal lattice with a two-atom basis,

788

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.2 Early atomic-resolution image of monolayer graphene grown on Pt(111) circa 1992. The 2.46 Å periodic structure is the graphene lattice, while the larger 22 Å periodic structure is a Moiré pattern arising from the overlay of the graphene lattice with that of the Pt(111) substrate. Reprinted with permission from Ref. [7].

which forms a honeycomb structure, as in Figure 19.3. The structure of graphene is that of a single atomic sheet of graphite, where graphite is formed by stacking many graphene layers. As the fundamental building block of graphite, which has been studied for many years [5], the electronic structure of graphene has also been known. Once the mechanical exfoliation technique was developed and electrical measurements of graphene could be made, there was a simultaneous effort by many groups to grow large area films of single-layer graphene. Sublimation-based growth on SiC remains one of the most successful of such techniques and has a chapter in this book dedicated to it. Growth on metal substrates has also had many successes with advantages (such as substrate cost) and disadvantages (such as the need to transfer the graphene to an insulating substrate for electronic applications) compared to SiC-based growth. Although many experiments utilize ultra-high vacuum (UHV) chambers with in situ characterization techniques to better understand the growth mechanism, much of the effort to produce graphene has been performed in a tube furnace style growth chamber. Typically, this type of furnace consists of a quartz tube placed inside a clamshell style heater with a vacuum pump on one end of the tube and a precursor gas inlet on the other. Hydrocarbon precursor gases, most commonly methane (CH4), are typically used as the carbon source and are diluted in a mixture of hydrogen and argon gases. The

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 789

FIGURE 19.3 Graphene is a honeycomb structure, which results from a hexagonal lattice with a two-atom basis. The C atoms in graphene are represented in black and the lattice points in gray.

hydrogen acts as a reducing agent for the metallic substrates, which often have a thin oxidized layer prior to growth due to atmospheric exposure. This type of growth has led to large area graphene films produced on various transition metal substrates, of which Ni, Cu, and Ir have been the subject of extensive, recent study and will be discussed in detail. Other studies have explored the formation of graphene on Pt [7,30–33], Pd [32], Re [9], Ru [34–36], and Co [32]. One of the primary factors that influences the particular type of growth that occurs is the solubility of C in the substrate (A plot of the solubility for each metal described in this chapter is in Figure 19.4). On metal substrates in which C solubility is high, notably Ni, carbon can absorb into the bulk and precipitate out as a graphitic film upon cooling [37]. In cases of low C solubility, notably Cu [37] and Ir [38], growth proceeds primarily by a surface decomposition process in which a gaseous carbon-containing precursor decomposes catalytically at the metal surface and growth propagates by C attachment to the free edge of the graphene film. For some carbide-forming transition metals, growth either coexists with [10,39] or competes with [40] carbide formation. In practice, the balance between these processes depends on both substrate material and growth conditions. In addition to solubility, each substrate also has a strong effect on the crystal quality of the graphene. Based on the strength of this interaction, different growth processes and growth surfaces can either result in a weak epitaxial relationship, where the substrate orientation and symmetry has a small effect on the resulting graphene film’s rotational orientation, or a strong epitaxial relationship, where the graphene film’s rotational orientation, and in some cases even the lattice constant, is dictated by the underlying

HANDBOOK OF CRYSTAL GROWTH

Carbon solubility (at%)

790

1

Ir (Arnould 1972) Ru (Arnould 1972) Ru (Yang 1990) Ni (Nateson 1973) Ni (Yang 1990) Pt (Siller 1968) Pt (Yang 1990) Cu (Lopez 2004) Molten Cu (Bever 1946)

0.1

0.01

1E-3

600

700

800

900 1000 1100 1200 1300

Temperature (°C) FIGURE 19.4 Plot of the solid solubility of C in various transition metals as a function of temperature with molten Cu also included. Solubility can be a strong indicator of the possible growth mechanisms for graphene on each substrate material. Data taken from Refs [38,41–44].

metal substrate. The specific effect of several different transition metal substrates will be discussed in detail in the following sections, where in each case the crystal-growth goal is to produce large area, high-quality graphene in a cost-effective and reproducible way.

19.2.1

Graphene Growth on Ni

The growth of graphene crystals on Ni substrates is one of the oldest examples of graphene CVD and as a result is one of the most extensively studied. However, this system is also among the most complex, and only recently has a fundamental understanding of the growth and stability of graphitic films on Ni been developed. Among the transition metals, Ni has a high solid solubility of C, absorbing approximately one atomic percent C at high temperatures (see Figure 19.4). Additionally, the Ni(111) surface is the most closely lattice matched to graphene of all transition metal growth surfaces. Unlike other high-solubility metals (e.g., Pt), Ni interacts strongly with the graphene overlayer, leading to substantial modulation of graphene’s electronic structure [45] and rotationally aligned epitaxial films [11]. While carbide formation is common on several metals (e.g. W), Ni has the singular characteristic that the carbide phase is stable only up to the relatively low temperature of 460
C, while graphene on Ni is stable to 650
C [46]. Above 650
C, the graphene will dissolve into the Ni due to the high C solubility. This provides a temperature range over which graphene growth can proceed in the absence of a carbide and also for dissolution and precipitation of graphene without formation of the carbide phase during cooling. The exact growth mechanism for graphene on Ni is therefore intimately dependent on the growth temperature. Three distinct temperature ranges have been identified and explored [40]. Typical growth at temperatures above 650
C in a tube furnace occurs by a dissolution/ precipitation mechanism. At temperatures above 650
C, the C-Ni(111) system forms a

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 791

clean Ni(111) surface with C absorbed into the bulk. Under these conditions, the catalytic decomposition of precursor molecules contributes C to the surface, which is then absorbed until saturation is achieved. With a carbon-saturated Ni substrate, typical of atmospheric-pressure CVD reactions, the resulting film is dominated by multilayer graphene precipitated from the bulk by phase separation during cooling. The growth of high-quality graphitic films on Ni substrates has been observed for decades. In early work, Banerjee et al. noted the influence of the substrate on the formation of pyrolytic carbon films formed from a C3O2 precursor [47]. It was recognized that the basal plane of graphite ran parallel to the surface, and the graphite film formed on Ni had crystalline ordering in three dimensions. The authors also reported the substantial interdiffusion of C and Ni at temperatures above 500
C. Karu and Beer subsequently grew graphite films as thin as 50 nm on annealed Ni foils using a CH4 precursor gas at temperatures between 800 and 1050
C [48]. In 2006, renewed efforts in graphene growth led to new interest in this field. Graphene CVD on Ni was explored by Somani et al., who deposited solid C10H16O onto Ni followed by pyrolyzation above 700
C [49], generating highly defective few-layer graphene (approximately 35 layers). This growth process was scaled towards monolayer films by Obraztsov et al., who formed few-layer graphene films on Ni [50]. Finally, several groups demonstrated the growth and transfer of large area, few-layer graphene films that contained substantial monolayer regions [51,52]. In each case, growth occurred within the temperature range where C absorption and precipitation dominates, but various techniques were employed to minimize the formation of multilayer regions: Ni thin films, dilution of C precursor gasses, and rapid sample cooling. Graphene transfer was performed with various polymer stamp materials and chemical etching of the Ni substrate (See Section 19.2.5 for a brief overview of this process). Due to the absorption of C into the Ni substrate at elevated temperatures, graphene does not form when the precursor is introduced at high temperatures, but instead precipitates as the metal cools to room temperature. This is in direct contrast with the dominant growth mechanism on Cu, and lower-temperature phase segregation-driven growth on Ni. This distinction was made clear in a study by Li et al. [37], in which isotopic C was employed (alternating 12C and 13C CH4 during growth) to contrast the growth mechanisms on Cu and Ni surfaces. On Cu, where growth proceeds by a surface decomposition process, isotopically pure rings of graphene formed during growth (discussed in detail in the following section). In contrast, on Ni films at a growth temperature of 900
C, isotopically pure C mixes in the bulk at high temperature and precipitates as graphene with a random distribution of 12C and 13C isotopes. The increased mass of 13C enriched graphene changes the various C–C vibrational modes of graphene in a predictable way, allowing each isotopically labeled region to be distinguished by Raman spectroscopy (Figure 19.5). In addition to dissolution/precipitation, at temperatures between 460
C and 650
C, a monolayer graphene film will phase segregate at the Ni(111) surface [53]. This segregation process was described by Shelton et al., as early as 1974 [53], who observed and described

792

HANDBOOK OF CRYSTAL GROWTH

(a)

(b)

SiO2

1563

Rel. 1/cm

1544

Intensity (a.u.)

(c)

G (1553)

2D (2631)

D (1325) 1250

1750

2250

2750

-1

Raman shift (cm ) FIGURE 19.5 Graphene growth on Ni showing (a) an optical image of the film, (b) a 2-dimensional (2D) Raman map showing a mixed 13C and 12C signal, and (c) a Raman spectra identifying the G and 2D peaks, without any indication of a D-peak. Reprinted with permission from Ref. [37].

the monolayer formation as an equilibrium condition occurring within this temperature range. This phase segregation process occurs in conjunction with any precipitation of C during cooling. In cases where the Ni substrate is carbon-saturated, monolayer graphene will form at the Ni surface at 460–650
C, but further phase separation of C from Ni drives the formation of multilayer graphene with the phase segregated monolayer. Under lowpressure growth conditions, where the Ni substrate is not fully saturated with C, the Ni substrate will contain excess C, and the equilibrium graphene monolayer can be preserved [46]. At high temperatures and pressures, the Ni substrate is likely to be fully saturated with C, leading to multilayer graphene with poorly controlled layer count. For temperatures below 460
C, on the other hand, it has been found that a stable Ni2C surface phase exists. Since the solubility of C at this temperature is low, the C atoms do not dissolve into the bulk and instead form a surface carbide that exists on the surface along with graphene. Extended annealing results in a surface completely covered with graphene as the C atoms from the precursor gas slowly exchange with the Ni atoms in the carbide [54].

19.2.2

Graphene Growth on Cu

The primary substrate that is currently used for growth of large area graphene films is Cu. One reason that graphene growth on Cu substrates has become such a popular growth technique is the relative ease with which single-layer films can be grown.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 793

An important reason for this can be seen in the solubility plot for C and Cu (Figure 19.4). It is clear that there is an extremely low solubility of C in Cu for temperatures below the melting point of Cu (1085
C). As will be shown later, it turns out that for these temperatures, the C solubility is negligible, resulting in a growth mechanism that is quite different than the dissolution/precipitation mechanism used for most Ni-based growth. The growth mechanism, as will be discussed in the following section, is what enables single-layer growth of graphene on arbitrarily large Cu foil substrates. Combined with the fact that Cu foil can readily be purchased in sufficient quantity to keep a research team active for many months for around $100 and that graphene can be grown in a standard quartz tube furnace, this technique has quickly become widespread. In order to grow arbitrarily large single-layer graphene films that can be transferred to an arbitrary substrate, a CVD process was developed by Li et al., in which predominately single layers of graphene could be grown [55]. In their study, a Cu foil substrate was heated to 1000
C in a quartz tube furnace in a gas mixture that contained both CH4 and hydrogen. After a growth time of 10 min, the authors found that the Cu foil was completely covered with graphene. By using a polymer-assisted transfer process, in which a polymer support was coated on top of the graphene and then the Cu foil was dissolved in a Cu etchant, the graphene films could be transferred to an arbitrary substrate (See Section 19.2.5). Following this process, the polymer material was then rinsed off with acetone, leaving exposed single-layer graphene, such as in Figure 19.6(c) and Figure 19.6(d), where the graphene was transferred to a 300 nm SiO2 on Si substrate and glass substrate, respectively. On this substrate, measurements of graphene thickness were carried out by correlating Raman spectroscopy (Figure 19.7(c)) with a scanning electron microscope (SEM) image (Figure 19.7(a)) and optical microscopy image (Figure 19.7(b)). The Raman spectra confirm that almost the entire surface consists of single-layer graphene, which the authors claim resulted from all of the growth conditions that were longer than 10 min (growths less than 10 min yielded incomplete graphene overlayer coverage). This suggests that the growth mechanism on Cu substrates is different than for growth on Ni and other transition metals with high levels of C solubility. As was later discovered and will be discussed in detail below, the growth of graphene on Cu substrates self-limits at a single layer. As part of the same study [55], a polymer-assisted transfer process was used to transfer the graphene to an arbitrary substrate for electrical measurements. This was necessary since electrical measurements directly on the Cu foil are not possible due to the conducting substrate. As in Figure 19.8(a), a Si substrate with a 300 nm SiO2 was used, and a top gate was patterned and deposited using an Al2O3 dielectric along with electrical (source and drain) contacts to the graphene. Electrical measurements yielded a device mobility of approximately 4000 cm2/Vs, which exceeds the typical values for large area graphene films grown on SiC (see Figure 19.8). The synthesis of uniform and continuous single layer graphene on copper substrates was considered a major breakthrough for the graphene community. Researchers could now control the shape and size of the grown graphene film simply by modifying the copper

794

HANDBOOK OF CRYSTAL GROWTH

(a)

(b)

(c)

(d)

FIGURE 19.6 Graphene on Cu foil substrate. (a) and (b) are SEM images, showing dark graphene grains on the Cu foil substrate. (c) and (d) are from graphene films transferred from the Cu foil to SiO2 and glass, respectively. Reprinted from Ref. [55]

(b)

(c) D Intensity (a.u.)

(a)

G

2D

3L 2L 1L

1300 1500 1700 1900 2100 2300 2500 2700

Raman shift (cm–1) FIGURE 19.7 (a) and (b) are real-space images from SEM and optical microscopy from the same region of graphene transferred from a Cu foil (a) to SiO2 (b). Highlighted regions were measured with Raman spectroscopy, with the respective spectra in (c). Reprinted from Ref. [55]

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 795

(a)

VBG= +10V VBG= +5V VBG=

0V

VBG= –5V VBG= –10V

FIGURE 19.8 (a) an optical microscope image of the device, which contains a global back gate from the Si substrate, Al2O3 dielectric, and metallized top gate and source and drain contacts. The resulting transfer curves are in (b). Reprinted from Ref. [55]

Resistance (kΩ)

(b)

0.9 0.8 0.7 0.6 –2

–1

0 VTG(V)

1

–1

0 1 VTG–VDIRAC.TG(V)

substrate. Furthermore, the polymer-assisted transfer process allowed graphene to be deposited onto any desired substrate. These advantages motivated many research groups that had previously performed growth on Ni substrates to rapidly and easily transition to the use of Cu substrates, which spurred the widespread use of this technique. Further evidence that the growth self-limits at a single layer was presented by Li et al. [37] in a study that alternated the precursor gas between CH4 composed solely of 12C and CH4 composed of 13C. The resulting graphene film was therefore isotopically labeled, which results in a distinct Raman signal based on the atomic weight of the C atoms in the graphene. This enabled formation of Raman maps, such as those in Figure 19.9(e) and Figure 19.9(f), where regions that are from graphene grown with 12C are clearly distinguishable from 13C. The reason that this image was so powerful was that prior to development of graphene growth on Cu, growth was primarily done on Ni substrates. As was discussed in the previous section, growth of graphene on Ni substrates occurred by a dissolution/ precipitation mechanism, in which the resulting graphene film’s thickness was difficult to control. For that type of growth, the signal from the 12C and 13C were found to be mixed in the graphene film because the 12C and 13C mix together in the Ni before precipitating as graphene during the cooldown (Figure 19.5). The data in Figure 19.9, however, provide strong evidence that the graphene growth occurs exclusively on the surface of the Cu because of the distinct 13C and 12C regions in the film, which correspond to the time during which the chamber was filled with each particular precursor. It

HANDBOOK OF CRYSTAL GROWTH

(b) Intensity (a.u.)

(a)

2D13 2D12

D13G13 D12G12 12C&13C 13C

(c) 200 100 0

12C

1200 1600 2000 2400 2800

–100

Raman shift (cm–1)

(d)

(e)

G13 G12 G13+12

300

CCD cts.

796

0

5

10

15

20

25

Distance (μm)

(f)

FIGURE 19.9 (a) an optical microscope image of the region mapped with Raman spectroscopy. Regions highlighted with colored circles have spectra in (b). (c) is a line scan from (d)–(f), in which maps from 1500 to 1620 per cm2, 1500–1560 per cm2, and 1560–1620 per cm2 are shown. Clearly distinct regions in (e) and (f) indicate growth that occurred by a surface-mediated growth mechanism. Reprinted with permission from Ref. [37].

also helps to explain why the growth self-terminates at a single layer. Since the Cu foil’s catalytic activity is the primary mechanism for the decomposition of the hydrocarbon precursor gas (usually CH4), once the graphene film grows large enough to cover the entire Cu surface, there are no longer any exposed Cu atoms to catalyze the decomposition reaction. Since graphene growth temperatures are insufficient to induce thermal decomposition of CH4 (i.e., without an exposed, active catalyst), once the graphene film covers the entire surface, the growth stops. It has also been shown that when a graphene island encounters a step in the underlying Cu surface, the graphene grows continuously over the surface structure [56]. While this technique is extremely effective for growing large areas of single layer graphene, which may be suitable for certain applications, it is also desirable to have graphene films that are composed of large graphene crystals. This is important for many electronic applications because grain boundaries between adjacent and rotationally misaligned graphene grains are a source of electron and hole scattering [57,58]. There are several possible techniques that can be employed to increase the grain size of graphene. One option is to perform the growth on a substrate that yields a single rotational orientation epitaxial relationship with the graphene. In that case, when graphene grains that nucleate in different regions of the substrate eventually coalesce, they will be rotationally aligned due to the templating effect of the substrate. This particular strategy

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 797

(a)

FIGURE 19.10 A Cu enclosure used to suppress nucleation for graphene films grown by chemical vapor deposition is shown in (a), with the enclosure crimped along three sides and then placed in a tube furnace for growth as in (b). Reprinted with permission from Ref. [59].

(b)

will be discussed in more detail later in this section. Another strategy for increasing the grain size of the graphene involves reducing the nucleation density, so each graphene grain can grow to a large size before coalescing. A technique for performing this type of growth was developed by Li et al. [59] in which a Cu foil was folded into a small “packet” as shown in Figure 19.10. The edges of the packet were folded over, but not sealed, so that inside the packet there was both an increased partial pressure of Cu vapor due to its sublimation during graphene growth, and also a reduced partial pressure of CH4. The net result of this procedure was a much reduced nucleation rate of graphene within the packet, and graphene islands 100’s of microns across were reported as in the SEM image of Figure 19.11. One study in which the atomic structure of graphene grown by CVD on Cu foils was studied utilized transmission electron microscopy (TEM) to reveal some interesting properties of the graphene films. The work, which was published by Huang et al. [60], required taking graphene grown by a technique similar to the one above [55] and transferring the graphene to a Cu or holey silicon nitride TEM grid. Once on the TEM grid, real-space images of the graphene lattice could be measured. In their paper, a region where two adjacent graphene grains met was imaged, and it was found to be composed of pentagon/heptagon graphene defects (see Figure 19.12). Interestingly, the graphene domains on either side of the grain boundary are rotated with respect to each other, suggesting that coalescence of individual graphene islands resulted in a polycrystalline graphene film.

798

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.11 Large graphene “island” resulting from the reduced nucleation growth technique in Figure 19.10. Reprinted with permission from Ref. [59].

FIGURE 19.12 TEM image of graphene transferred from Cu foil to a TEM grid. The grain boundary in (a) is depicted in (b) as a series of pentagon/heptagon defects. Adapted from Ref. [60]

(a)

(b) 27°

In addition to performing real-space imaging, the authors also observe a diffraction pattern from the graphene, which shows a circular ring with varying intensity (Figure 19.13(b)). This immediately indicates that the graphene film is indeed composed of a variety of different rotational orientations. By using an aperture to select only a single rotational orientation at a time, a dark-field image was created and color coded according to rotational orientation (see Figure 19.13(g)). The average grain size was calculated to be 250  11 nm, whereas slight changes to the growth process resulted in films with grains on the order of several microns. In order to study the growth process in situ, Wofford et al. grew graphene on a Cu foil substrate in a chamber containing a low energy electron microscope (LEEM) [61]. Growth in the LEEM chamber enables simultaneous measurement of the graphene film in both real space and k-space in situ. Because the LEEM operates in an UHV, they were unable to do growth by catalytic decomposition of a hydrocarbon, such as CH4, due to the high pressures that are needed to suppress Cu sublimation during growth [62].

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 799

(a)

(f)

(b)

(c)

(d)

(e)

(g)

FIGURE 19.13 (a)–(e) depict the process used to create a dark-field image of the graphene using selected-area diffraction. (f) and (g) are dark-field images, showing a graphene film composed of many different rotational orientations of graphene. Reprinted from Ref. [60]

Instead, they directly evaporated C using an electron beam, and grew the graphene film by providing a flux of atomic C to the Cu substrate. Although in this case the Cu was not catalyzing the decomposition of a hydrocarbon precursor, the interaction between the C and Cu substrate are expected to be similar. Based on their LEED measurements, the authors found that ordered graphene films were only obtained when the Cu foil was heated to temperatures of approximately 800
C or higher, suggesting a lower limit for expected high crystal quality graphene films grown on Cu. The authors also found that after heating the Cu foil to temperatures typically used for CVD growth of graphene, the Cu foil crystallizes with a (100) surface. Individual graphene “islands” were imaged with dark-field imaging and found to each be composed of multiple rotational domains of graphene, as in Figure 19.14. This distinction, where we refer to individual graphene regions as “islands” rather than “grains” or “domains,” is important because of the clear rotational misorientation that exists between different grains within the same island. The work presented in this study clearly demonstrates that a single nucleation site for graphene on Cu foil can result in a single graphene island that upon a closer diffractionbased study, is actually composed of several different rotational orientations. Figure 19.14 shows a single graphene island that is imaged by inserting an aperture and imaging only a single rotational orientation at a time. This shows that a four-lobed graphene island grows from what appears to be a single nucleation spot and actually consists of four distinct rotational orientations. This has several implications, most notably that there appears to be a very weak interaction between the Cu foil surface and the graphene, resulting in a weak or negligible epitaxial relationship between the substrate and overlayer. Also, it suggests that a more detailed understanding of the interaction between the Cu surface and graphene is necessary in order to optimize and better understand the growth process.

800

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.14 A single graphene island is shown in (a) with each rotational orientation that it is composed of shown in (b)–(e). Reprinted with permission from Ref. [61].

Since graphene grown on Cu foil is expected to include grain boundaries [60,61], their effect on electrical transport has been studied both theoretically and experimentally. By using scanning potentiometry, a study was published by Ji et al. [63], in which epitaxial graphene on SiC was used. It was found that for the case of SiC, substrate steps and regions where the graphene film thickness was nonuniform caused significant scattering in the film. Similar measurements have also been performed on graphene films grown by CVD on a Cu substrate, where it was found that grain boundaries, regardless of the angle between the adjacent grains, have an effective conductivity that is one-third the value (or less) of bulk graphene [58]. Despite overwhelming evidence that the graphene films grown by CVD on metal foils are polycrystalline, this growth technique has been widely adopted and is used regularly by many groups due to the relative ease with which single-layered films can be grown. Most of the Cu foils are manufactured by a cold-rolling process, which, due to the plastic deformation that the Cu undergoes during manufacturing, results in a surface that is terminated with a (100) plane [64]. Electron backscatter diffraction (EBSD) images of a Cu foil surface can be seen in Figure 19.15(a) and Figure 19.15(b) for “as-received” Cu foil and Cu foil post graphene growth, respectively. Figure 19.15(d) and Figure 19.15(e) are inverse pole plots, indicating both surfaces are terminated with a plane very close to the (100) [65]. However, following graphene growth several small (111) grains can also be seen in Figure 19.15(b). One other EBSD study of graphene growth on Cu foil shows the presence of substantial variation in the surface texture. This difference is likely because the studies employed Cu foil produced by different manufacturers with potentially different manufacturing techniques [66]. Since Cu is a face-centered cubic material, its two lowest energy surface terminations are the (100) and (111) planes, which have square and hexagonal symmetry, respectively. The Cu(111) surface is close-packed and is the lowest energy surface. Both Cu(100) and

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 801

(a)

(d)

(b)

(e)

(c)

(f)

FIGURE 19.15 EBSD data showing a Cu foil surface before graphene growth (a) and (d) and following conventional CVD growth (b) and (e). The growth was performed by the same group as in the original work [55]. Image (a) is 1.4  1.1 mm and (b) is 2.5  1.9 mm. Reprinted with permission from Ref. [65].

˚ , which is 3.5% larger than the graphene Cu(111) have a surface lattice constant of 2.55 A lattice constant. Since both Cu(100) and Cu(111) are well lattice-matched to graphene, it may be possible to develop growth conditions that encourage a sufficiently strong interaction between the substrate and overlayer to epitaxially grow graphene on the Cu surface. In this context, the term epitaxial is defined using the common definition where the overlayer’s rotational ordering is affected by the growth substrate. In the case of a strong interaction between the graphene and Cu(100), the graphene growth process is expected to result in a two-domain graphene film due to the two equivalent ways that a hexagonal material (such as graphene) can be grown on a surface with square symmetry (such as Cu(100)) [67]. Epitaxial growth on Cu(111) is expected to result in single-domain graphene if the interaction is strong enough, because both the substrate and overlayer have the same symmetry. In order to gain further insight into the specific effect that the Cu substrate may have on the graphene growth process, detailed UHV-based studies have been performed on Cu single-crystal substrates.

19.2.3

Graphene Growth on Cu Single-Crystal Substrates

In order to study the effect of Cu surface symmetry on the graphene growth process, there have been only a few studies performed by various research groups on graphene growth on Cu single-crystal substrates. The primary difficulty that all of the groups encountered was developing a growth technique that would allow for in situ measurements to be performed, which typically require UHV (1010 Torr), while being able to achieve growth in the same chamber, which requires pressure in the mTorr regime and temperatures in excess of 800
C. Since typical UHV heating is done by either radiative heating (passing a large current through either a foil or filament that is very close to the sample) or electron-beam heating (passing a current through a filament that is biased approximately 1 kV with respect to the sample so that electrons from the filament

802

HANDBOOK OF CRYSTAL GROWTH

bombard the sample), the main challenge for these studies was developing growth techniques that enabled heating while maintaining the growth chambers at high pressure. The two substrates that will be discussed are Cu(111), which is the lowest energy surface for Cu, and Cu(100), which is the second lowest energy surface. One of the first reported studies where graphene was likely grown on a Cu(100) surface was actually published well before Geim and Novoselov’s Nobel prize-winning paper was published in 2004 [1]. In the paper, Alstrup et al. [8] studied decomposition of CH4 on a clean Cu(100) and oxygen predosed Cu(100) single crystal. Importantly, the authors note that while exposing the hot Cu(100) surface to CH4, all of the filaments in their chamber, including their ion gauge, were turned off in order to prevent dissociation of CH4 and direct deposition of atomic carbon by the hot filament. Following exposure of the Cu(100) at 1000 K to up to 10 Torr of CH4, the authors found that approximately one monolayer of graphitic carbon was on the surface. They noticed that the graphitic carbon (graphene) deposition started around 800 K, and they were apparently the first group to determine that a graphene film would self-terminate at a single layer on Cu(100). In the same study, the authors found that the energy barrier for the initial decomposition reaction was reduced by 40% by adsorbing one-half of a monolayer of oxygen (compared to the atomic density of the Cu(100) surface) on the surface. This suggests that the inclusion of oxygen in the graphene growth process could be used to encourage graphene growth on Cu substrates that otherwise require temperatures very close to the Cu melting temperature. In order to study graphene growth on clean and oxygen predosed Cu(100) in more detail, Robinson et al. [67] conducted a study in which graphene was grown on Cu(100) in a UHV system that was equipped with in situ LEED, which enabled measurement of the epitaxial relationship between the graphene and Cu(100). LEED images for graphene growth on the clean Cu(100) surface are shown in Figure 19.16(b) and Figure 19.16(c) for growth temperatures of 800
C and 900
C, respectively. At 800
C, a rotationallydisordered graphene film is found with four weak preferred orientations with respect to the substrate. However, at 900
C, 12 very sharp graphene spots are visible in the LEED image, indicating a well-ordered two-domain graphene growth. This suggested that the interaction between the graphene and Cu(100) is sufficient for the Cu substrate to

(a)

(b)

(c)

FIGURE 19.16 LEED images from (a) clean Cu(100), (b) graphene on Cu(100) grown at 800
C, and (c) graphene on Cu(100) grown at 900
C. Images taken at 70 eV. Reprinted with permission from Ref. [67].

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 803

(a)

(b)

(c)

pffiffiffiffiffiffiffiffiffiffiffiffiffi pffiffiffi FIGURE 19.17 LEED images from (a) 2  2 2R45
reconstruction on the Cu(100), (b) growth on the oxygen
reconstructed surface at 800 C, and (c) growth on the oxygen-reconstructed surface at 900
C. Images taken at 70 eV. Reprinted with permission from Ref. [67].

influence the ordering of the graphene surface over very large length scales, since conventional LEED has a spot size of approximately 1 mm. Additionally, the authors studied growth ofpffiffiffigraphene on a Cu(100) surface that pffiffiffiffiffiffiffiffiffiffiffiffiffi was predosed with oxygen, resulting in a 2  2 2R45
reconstruction on the surface. This reconstruction is well-known to be the result of approximately half of a monolayer of oxygen on the Cu(100) and was studied due to the likelihood of small amounts of residual oxygen that may be present in tube furnaces during growth. It was found that the extremely small amount of oxygen on the surface of the Cu(100) dramatically influenced the graphene ordering, as shown in Figure 19.17. The small amount of oxygen also influenced the nucleation density of the graphene, with the presence of oxygen reducing the nucleation density by about a factor of 10. Around the same time of publication [67], Hao et al. [68] published a report that growth of large-grain graphene was enabled by performing growth on Cu foil in a dilute oxygen background. In this work, centimeter-sized graphene islands were grown in a conventional tube furnace, where dilute O was used to passivate the surface of the Cu foil. This resulted in a much reduced nucleation density, allowing each graphene island to grow to macroscopic scales before film coalescence. This technique is one possible route towards production of large area, single crystals of graphene. All of these data present an interesting story, and show that there is room for significant improvement for graphene growth on Cu foils, which recrystallize with a (100) surface upon annealing. However, since even in the best case this is expected to result in a two-domain graphene film due to the symmetry of placing a hexagonal lattice on a square lattice, graphene growth on Cu(111) would be highly desirable. For the case of Cu(111) single-crystal substrates, Gao et al. were one of the first groups to report on successful graphene growth [69]. In order to achieve growth, a first attempt was made by directing a gas nozzle in their UHV system at the face of a Cu(111) crystal and opening a valve to let ethylene into their chamber. The chamber pressure was

804

HANDBOOK OF CRYSTAL GROWTH

maintained at 105 Torr, with the expected pressure near the Cu(111) surface much higher, while the crystal was heated to 1000
C. Using this technique, almost no growth was observed on the surface. This was probably caused by Cu sublimation, which competes with graphene growth at high temperature, as will be discussed later [62]. In order to achieve graphene growth, the Cu(111) crystal was thermally cycled between room temperature and 1000
C with a constant background pressure of ethylene. The authors speculate that the ethylene adsorbs on the surface at low temperature and forms graphene during the high temperature anneal. Once a film was grown, the authors performed STM on the surface and observed small graphene islands on the surface. Two predominant rotational orientations of the graphene film were observed with rotational angles of 0
and 7
with respect to the underlying Cu(111). It is unclear what the result of the thermal cycling growth process was on the rotational orientation of the resulting film. In order to study the nucleation processes more closely, Nie et al. [70] exposed Cu(111) and Cu(100) substrates to a flux of atomic carbon and monitored the graphene growth with in situ LEEM/m-LEED. They found that for growth on Cu(111), single crystals of graphene were only grown when the C-flux was high enough to induce nucleation on terrace sites. When graphene nucleated at step edges or other defects in the substrate, the graphene grain was typically polycrystalline, as observed in their similar work on Cu foils [61]. Graphene growth on Cu(111) was also found to be strongly affected by the growth temperature, where both the morphology and crystal structure are improved considerably when growth is performed at temperatures of 950
C rather than 690
C. This study also reported that as a graphene grain grows across a step bunch in the underlying Cu surface, the rotational orientation of the grain can change. For instance, in Figure 19.19, a graphene grain undergoes a 21
rotation as it crosses a step bunch in the Cu surface. In order to further understand the effect that the Cu(111) surface had on the graphene growth process, Robinson et al. modified a UHV chamber to incorporate a button-heater style heated stage [62]. This enabled growth at relatively high pressure (>mTorr), which was followed by a quick return to UHV after growth. The UHV chamber was equipped with in situ LEED. It was found that performing growth in 5 mTorr of ethylene at 800
C resulted in very faint arcs, corresponding to a small percentage of graphene coverage. Growth at 900
C, on the other hand, resulted in no graphene on the Cu(111) (Figure 19.20(a) and (b)). Since at 900
C the Cu is expected to be even more catalytically active than at 800
C, it was determined that Cu sublimation during the growth was competing with the graphene nucleation and growth process. The sublimation rate at 800
C is very small, while at 900
C it is approximately one-third of a monolayer per second. The high sublimation rate of the underlying Cu prevents graphene nucleation and possibly explains why other groups performing growth at high temperature and low pressure were unable to observe graphene growth [69]. In order to suppress the sublimation of Cu during growth, a second set of experiments were performed in which both 5 mTorr of ethylene and 45 mTorr of Ar were backfilled into the chamber during growth. The role of the Ar was to

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 805

(a)

(b)

FIGURE 19.18 (a) and (b) are SEM images of graphene on clean Cu(100) and oxygen-reconstructed Cu(100), respectively. Reprinted with permission from Ref. [67].

(a)

(d)

(b)

(c)

(e)

FIGURE 19.19 (a)–(c) LEEM progression of graphene growing on Cu(111), with colored circles in (c) shown in the LEED image of (d), indicating a 21
rotation after graphene crossed the step bunch indicated with the red dotted line. Schematic representation in (e). Reprinted from Ref. [70]

806

HANDBOOK OF CRYSTAL GROWTH

(a)

(c)

(b)

(d) Intensity (a.u.)

2.0

1.5

Cu(111)

1.0

0.5

Graphene (×2)

0.0 0

100

200

300

400

Angle (degrees)

FIGURE 19.20 (a) and (b) are for growth on Cu(111) in 5 mTorr of ethylene at 800
C and 900
C, respectively. (c) is growth at 900
C in 5 mTorr of ethylene with 45 mTorr of Ar to supress Cu sublimation during growth. (d) is a circumferential intensity profile for both the Cu(111) and graphene. Reprinted from Ref. [62].

suppress the Cu sublimation at the growth temperature. In this case, for growth at 900
C, single-domain, rotationally aligned graphene formed on the surface of the Cu(111), as can be seen in Figure 19.20(c). A LEEM/m-LEED study was performed ex situ, which enabled real-space imaging of the graphene/Cu(111) and measurement of the epitaxial relationship between each individual graphene grain and the underlying substrate. These measurements further confirmed that graphene was grown with a single rotational orientation on the Cu(111) surface. To conclude, graphene growth on Cu substrates has been shown to be an extremely useful growth technique for making large area films in an extremely cost-effective way. For many applications, this growth technique can provide the quality films that are required. However, growth on foils typically results in polycrystalline graphene growth where the quality of the film is impacted by the underlying Cu surface. Techniques have been developed, though, for controlling the nucleation density and therefore increasing the size of individual graphene single-crystal islands [59,68]. Growth on Cu single-crystal substrates has been extremely effective at determining the growth mechanism and

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 807

limiting factors but falls short of being a manufacturing solution due to the preventatively high cost of single-crystal substrates. Therefore, further development is needed in order to produce single-crystal quality graphene films.

19.2.4

Graphene Growth on Iridium

Iridium is another transition metal studied extensively as a graphene growth surface. Graphene growth on Ir is often characterized by low carbon solubility, leading to predominantly monolayer growth. However, the solubility of C in Ir is higher than in Cu. Although graphene interacts weakly with the Ir(111) surface [71], single-domain epitaxial graphene growth on Ir, where the graphene overlayer is in registry with the underlying substrate, has been demonstrated [72]. Also, graphene islands on Ir(111) will grow both up and down across atomic steps on the surface [73]. This is in contrast to Ru, where growth proceeds only down across the atomic steps of the surface [34]. Two distinct growth mechanisms have been explored for graphene growth on Ir: CVD and temperature programmed growth (TPG). Under CVD, the Ir substrate is heated and exposed to a hydrocarbon precursor gas, which cracks by the catalytic action of the substrate. In contrast, under TPG, the substrate is first exposed to a precursor, which adsorbs at low temperature and is then flashed to elevated temperature, at which point the precursor decomposes, forming a partial graphene layer at the surface. The dosing and flashing procedure is then repeated until the desired coverage is achieved, each cycle saturating 22% of the exposed Ir surface [74]. Although many processing cycles are required, the TPG growth process leads to increased epitaxial registry between the graphene film and Ir substrate. These growth processes are similar to those employed on Cu substrates. The most common growth recipe for graphene on Cu is a CVD process performed at elevated temperature, but some early demonstration employed a recipe closer to TPG [69], which led to two-domain epitaxial graphene growth on the Cu(111) surface. Like Cu, under optimized conditions, graphene CVD on Ir will proceed by surface decomposition due in part to the low carbon solubility of Ir (See Figure 19.4). This was demonstrated in several early papers. In an early LEEM study, Loginova et al. explored graphene growth on both Ru(0001) and Ir(111) [73]. The cracking of various precursor molecules was restricted to the metal surfaces in both cases with the graphene forming rotationally aligned epitaxial films, as noted in early work [75]. Additional STM studies have elucidated the nucleation and growth of graphene islands on the Ir(111) surface for both CVD and TPG processes [76–78]. As shown in Figure 19.21, graphene is grown by CVD at 847
C, forming graphene islands at atomic steps that subsequently grow and coalesce until full surface coverage is achieved. The growth rate decreases as surface coverage approaches 100% due to the need for exposed Ir to facilitate catalytic decomposition of the precursor molecule. The derivative of each STM images is shown to clarify the details of the highly stepped surface.

808

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.21 Derivative scanning tunneling microscopy topographic images of graphene islands forming on Ir(111) by CVD growth at 847
C. Each large-scale image is 1  1 mm in dimension. Growth pressure was 5  10-10 mbar and growth times were 20 s (a), 40 s (b), 160 s (c), and 320 s (d), respectively. In (a) and (b), graphene sheets on upper and lower terraces are indicated in the inset by blue (lower) and red (upper) coloration. In (c) and (d), the graphene sheets have coalesced across terraces and are simply indicated with blue and red striping. Figure from Ref. [76].

(a)

(b)

(c)

(d)

The use of the TPG process provides optimum control of alignment between graphene and Ir substrate and of coverage. Growth proceeds in a manner similar to the CVD procedure, where graphene islands form and ultimately coalesce. However, such island formation is highly dependent on the TPG conditions. In particular, as shown in Figure 19.22, the size and distribution of islands is a function of TPG annealing temperature with higher temperature annealing producing larger, sparser islands with regular edge structure dependent on the graphene and Ir atomic structures [76].

19.2.5 Graphene Transfer To access the electronic properties of CVD-grown 2D materials, it is often necessary to transfer from metallic growth surfaces onto insulating platforms. Several procedures have been developed to facilitate this transfer, including wet-chemical etching procedures [55, 79], often employing a polymer stamp, as outlined in [Figure 19.23]. The polymer stamp is commonly composed of poly(methyl methacrylate) (PMMA) or polydimethylsiloxane (PDMS). The metal/2D material/polymer stack is then transferred to a wet chemical bath which selectively dissolves the metal, leaving pristine 2D material on a polymer support floating atop the solution. The 2D material can then be extracted from solution onto an arbitrary substrate and the polymer dissolved in a solvent bath. The process is simple,

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 809

(b)

(c)

(d)

(e)

(f) n (Islands / site)

(a)

0.01

1E-3 1E-4 1E-5 1000

1200 1400

T (K)

FIGURE 19.22 Derivative scanning tunneling microscopy topographic images of graphene islands on Ir(111) using temperature programmed growth. Each large image is 250  250 nm and each corresponds to a 20 s annealing cycle to a controlled temperature: (a) 597
C, (b) 697
C, (c) 847
C, (d) 1047
C, and (e) 1197
C. In (f), the relationship between annealing temperature and island density is shown, where island densities fall at elevated TPG annealing temperature. Figure from Ref. [76].

inexpensive, and generalizable to many materials. Among the disadvantages, wet chemistry can lead to contamination of 2D materials and also requires the use of a sacrificial metal growth substrate. In the case of polycrystalline Cu foils, the costs are low, but this limits the use of high-quality, single-crystal metal surfaces, on which CVD growth is better controlled. Also, the wet transfer of 2D materials leads to the introduction of folds and wrinkles into the film. These wrinkles can be undesirable, as when they create rotational disorder in transferred films [177], or desirable, as when they introduce stretchability into otherwise inelastic films [178]. Alternate techniques include electrostatic transfer of graphene [179] and dry transfer to polymers with the assistance of an azide linker molecule to promote adhesion [180].

19.3 Chemical Vapor Deposition of Hexagonal Boron Nitride Hexagonal boron nitride is an inorganic, electrically insulating analogue to graphite. It exhibits a similar layered crystal structure where strongly bonded atoms form sheets with a hexagonal lattice while adjacent sheets are weakly bonded via van der Waals forces and has a tendency to form in lamellar microstructures (See Figures 19.24 and 19.25). The lattice is hexagonal with AB stacked layers, so its unit cell is comprised of two ˚ , c ¼ 6.61 A ˚ ) giving an inter-layer spacing of 3.2 A ˚ [81]. AB stacked layers (a ¼ 2.504 A

810

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.23 Procedure for patterned growth and transfer of graphene on Ni thin films. Prepatterned Ni films on SiO2 are heated in a precursor gas, with C absorbing into the Ni film. Upon cooling, carbon precipitates out as few-layer graphene. Graphene is then transferred by wet-chemical etching of the support layer with an optional polymer stamp. Reprinted with permission from [52]

The different electron affinities of boron and nitrogen make hBN a covalently bonded insulator with a large band gap [82,83]. Its hardness, high thermal conductivity, thermal stability, low shear strength, and chemical inertness have led to interest in hBN anticorrosive coatings [19] and solid lubricants [84]. More recently, its insulating nature and surface smoothness have led to interest in using hBN films for electronic applications, including ultraviolet emitters [85,86], tunnel barriers [87], and as an inert, flat, and insulating substrate for graphene [19,20,83–85,88,89] with which it has a small (1.8%) lattice mismatch [90]. Several different methods [19] have been explored for deposition of hBN films with recent efforts most focused on single or few-layer films. CVD is the most common method for synthesis of thin hBN films for several reasons. Under some conditions, it offers the ability to limit hBN film growth to one or a few monolayers. The hardware is relatively accessible compared with other methods, requiring only inexpensive gas flow controls, a vacuum pump, and a method for substrate heating, such as a furnace. In addition, the technique and instrumentation is shared with the most frequently used method for synthesis of monolayer graphene [55], such that there is a significant investment in CVD knowledge and hardware at research institutions worldwide. Indeed, efforts to grow vertical graphene/hBN heterostructures [89], patterned in-plane

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 811

FIGURE 19.24 Drawing of the crystal structure of hBN. Reprinted with permission from [81].

graphene/hBN heterostructures [92] (see Figure 19.26), and (BN)xC(1-x) films [93] have each met with promising preliminary success. Magnetron sputtering [94,95], molecular beam epitaxy [96,97] (MBE), and ion beam assisted electron beam deposition [98,99] (IBAD) have also been generally used to deposit thicker hBN films and are briefly discussed later in this section. The typical CVD apparatus for deposition of hBN thin films consists of a precursor container connected to a deposition chamber, as shown in Figure 19.27. The precursor container may be a borazine bubbler (often refrigerated to slow decomposition), which provides borazine vapor carried by an inert gas (Ar or N2) or a heated ammonia borane container, which provides hydrogen gas and a mixture of precursor gases to the deposition chamber [20].

19.3.1

Precursors

Gaseous precusors for CVD of hBN are combinations of ammonia (NH3) with diborane [101,102] (B2H6) or boron trichloride [103] (BCl3). However, due to the toxicity of these gases, most academic researchers employ polymeric, liquid, or solid starting precursors with a 1:1 B:N ratio, principally borazine [104–110] (B3N3H6) or ammonia borane [89,100,111,112] (BH3NH3), to deliver boron and nitrogen to the target substrate. Borazine (B3N3H6) (#4 in Figure 19.28) is the BN analog to benzene and is a flammable and air-sensitive liquid with a high vapor pressure at room temperature. When heated (as low as 70
C), it undergoes dehydrogenation to form polyborazylene while

812

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.25 Scanning electron micrograph of hBN powder displaying its characteristic 2D plate-like microstructure. Adapted with permission from [91].

FIGURE 19.26 Monolayer lateral graphene/hexagonal boron nitride (hBN) heterostructures created by patterned conversion of graphene to hBN. Schematic (left), optical micrograph showing differing optical contrasts on SiO2 substrate (center) and Raman map graphene’s 2700 cm1 Raman peak (right). Reprinted with permission from [125].

releasing hydrogen gas [113]. At temperatures of 800
C and above, the dehydrogenation process will continue and polyborazylene will transform to hBN [20,108,113–115]. Metal surfaces can speed up the formation of hBN, as discussed later in this section. Polyborazylene can also be synthesized ex situ, spun-cast onto the target substrate [115,116], and thermally decomposed to form hBN films but is most commonly formed in situ by decomposition of borazine. Ammonia borane (BH3NH3) (AB, also known as borazane) can be viewed as a convenient in situ source of B- and N- containing vapors, including borazine. The cost and high reactivity of borazine and polyborazylene to ambient air make the use of ammonia borane an attractive option for CVD of hBN films. AB is a stable solid under ambient conditions but begins to decompose at temperatures as low as 70
C [117] to yield several gases with 1:1 B:N stoichiometry, including borazine. However, the decomposition process and resulting mixture of gases is more complex than that using borazine. Several studies have examined the thermal decomposition of AB, focusing on applications in hydrogen storage, in particular the works of Wolf, Baitalow, and Baumann

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 813

Ammonia borane (BH3NH3) 130 °C

H B H N

Polyiminoborane Borazine H -[BH=NH]H N

H B B H

H N n

Ammonia borane (BH3NH3)

N

B H

Hydrogen (H2)

H N

B H

Hydrogen (H2) B H N H

1100 °C

Boron nitride

Borazine

B H N H

Pt foil

Borazine, hydrogen

Pt foil

H2 Furnace (1100 °C)

Hot plate (130 °C) FIGURE 19.27 Rendering of a typical ambient pressure or low-pressure hBN CVD reactor using ammonia borane precursor. Reprinted from [100].

FIGURE 19.28 Thermal decomposition pathway of ammonia borane (1) to Aminoborane (2) and borazine vapor (4). Reprinted from [91].

[117–119] and more recently Frueh et al. [91]. Decomposition of AB proceeds via rather complex sequential dehydrogenation to Aminoborane (NH2BH2), Polyaminoborane (NH2BH2)x, Polyiminoborane (NHBH)x, borazine (H3NBH3), Diborane (B2H6), and boron nitride (BN), as well as hydrogen gas (See Figure 19.28) [91]. For CVD of hBN films, the gaseous products are of interest: Aminoborane, Diborane, and borazine. Although the reactions leading to each decomposition product may be distinct, their activation energies are so similar that competing reactions occur simultaneously, producing a mixture of decomposition products at all but the slowest heating rates (0.05
C min1)

814

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.29 Mole yield of hydrogen, Aminoborane, borazine, and Diborane from the thermal decomposition of ammonia borane (Borazane) up to 227
C at different heating rates. Reprinted with permission from Ref. [118].

[117] (See Figure 19.28). This complicates the analysis of AB as a precursor for hBN films. Furthermore, Aminoborane is extremely reactive [120] and may polymerize spontaneously or in reaction with its own ammonia borane precursor [121], leading to downstream deposition of larger BN particles and nanostructures [122,123]. Despite these complications, several research groups [111,122] have demonstrated deposition of ultrathin hBN films using AB precursor, and recent efforts to selectively filter out the more reactive gas species have also demonstrated some success in reducing particulates on hBN thin films [112].

19.3.2

Monolayer Hexagonal Boron Nitride

For various electronic applications, a single atomic layer of hBN (thickness about 0.4 nm) with an atomically smooth surface is highly desired. Self-limited monolayer hBN growth, where the first inert hBN layer inhibits adsorption of more precursor molecules, has been observed by several groups [106,109,110,124]. Borazine vapor at very low pressures (w106 Torr) has been observed to result in self-limited monolayer hBN growth on Pt(111) [114], Ru(001) [105,109,114], Rh(111) [105,106], and Ni(111) [105]. Other recent efforts have focused on obtaining single-layer hBN films in more easily accessible pressure regimes (w100 mTorr to ambient). Kim et al. [122] showed monolayer hBN deposition on Cu foil under low pressure (350 mTorr) using an ammonia borane precursor (See Kim et al.) [100]. These two results and the earlier UHV processes point strongly at pressure and precursor delivery rate as critical parameters for the control of thickness of hBN films deposited on metal foils. Most recently, Gong et al. [125] have demonstrated continuous and complete conversion of graphene on Cu foils to both hBN and intermediate (BN)xC(1-x) films using boric acid and ammonia precursors at low pressure. For electronic applications, crystallinity, grain size, film thickness, porosity, and surface roughness are the most important physical parameters. Surface imaging techniques (scanned probe microscopy and electron microscopy) are normally used, in addition to optical microscopy. Raman spectroscopy, as discussed in relation to graphene earlier in the chapter, is also a crucial indicator of hBN film crystallinity [126]. Films (especially those deposited on metal substrates) are often transferred to standard ultra-flat surfaces like oxidized Si wafers for characterization. The film transfer

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 815

Parameter (A) Mismatch (%)

hBN (0001) 2.50

C (0001) 2.46

Co (0001) 2.51

Ni (111) 2.49

Cu (111) 2.50

Ru (0001) 2.71

Rh (111) 2.69

Pd (111) 2.75

Pt (111) 2.77

--

1.6

-0.4

0.4

0.0

-8.4

-7.6

-10.0

-10.8

FIGURE 19.30 In-plane lattice parameters for various hexagonal close-packed metal surfaces and corresponding lattice mismatch of each with bulk hexagonal boron nitride.

techniques are common among graphene, hBN, and other 2D crystals and were presented in Section 19.2.5 of this chapter.

19.3.3

Substrate Interactions

The interaction between neighboring hBN layers is generally weaker than that of hBN to different adjacent material or substrate. Thus, the substrate surface properties and structure can have an important influence on the growth of hBN films. Although hBN is inert, its chemical precursors are very reactive, especially at the high temperatures used for hBN formation, and can form interface compounds with reactive substrates. Oxygenrich surfaces like SiO2 and Al2O3 can oxidize the boron in precursors, leading to thermally stable interface compounds [20]. Several metal surfaces efficiently catalyze hBN formation, in particular Ni(111), which has been shown to produce complete hBN surface coverage with one-tenth the precursor dose required to achieve the same coverage in Cu(111) [127] despite the similar lattice parameters. The Nickel surface has a stronger interaction with hBN than other metal surfaces do, exhibiting larger work function lowering [128]. Both Ni(111) and Cu(111), as well as Co(0001) and graphite offer good in-plane lattice match for the epitaxial growth of hBN (See Figure 19.31). Monolayer hBN can also grow epitaxially on Rh(111) and Ru(0001), but the additional lattice mismatch leads to the formation of a regular corrugation of hBN with a periodicity of 3.2 nm [105,106,124], as shown in Figure 19.31. Lastly, hBN exhibits a peculiarly small and negative thermal expansion coefficient [129]. This leads to expansion during cooling of hBN films deposited at high temperature, just as the underlying metallic substrate contracts. The result is the formation of wrinkles in hBN film, as seen in Figure 19.32.

19.3.4

Other Deposition Methods

Sputtering, MBE and IBAD provide some alternative methods for deposition of hBN. JinXiang et al. [94] first deposited hBN films by magnetron sputtering of a bulk hBN target, and in a more recent and detailed study performed by Sutter et al. [95], boron was sputtered onto a Ru(0001) epitaxial film in a nitrogen/argon atmosphere. The few-layer hBN film can be clearly seen in cross-sectional TEM (Figure 19.33).

HANDBOOK OF CRYSTAL GROWTH

Height (nm)

816

(a)

(b)

(c)

(d)

0.10

0.10

0.05

0.05

0

5 10 Distance (nm)

15

0

5 10 Distance (nm)

15

FIGURE 19.31 Scanning tunneling microscopy image and height profile of monolayer hexagonal boron nitride corrugation formed by decomposition of borazine on Rh(111) (a,c) and Ru(0001) (b,d). Reprinted with permission from Ref. [105].

FIGURE 19.32 Atomic Force Microscopy image of an incomplete monolayer hexagonal boron nitride islands grown by CVD on a Cu substrate and transferred to SiO2 showing consistent thickness of 0.42 nm (left) and optimal surface roughness of 0.181 nm (right) excluding w1 nm wrinkles arising from thermal expansion mismatch. Reprinted with permission from Ref. [122].

Fewer works have focused on MBE growth of hBN films so far. Gupta et al. [96] deposited hBN films on sapphire (0001), using ammonia (NH3) as a nitrogen source and an effusion cell as the boron source. The resulting films were confirmed to be hBN via IR spectroscopy but showed a highly polycrystalline texture in electron diffraction. The technique was more recently improved upon by Tsai et al. [97] by using nitrogen plasma and selecting a lattice˚ and RHEED patmatched substrate, Ni(111). The surface RMS values of the film was 9.5 A terns confirmed a correlated Ni(111)-hBN system (Figure 19.34). Efforts to use IBAD have focused on the synthesis of cubic boron nitride (cBN) rather than the hexagonal phase but often result in a combination of cBN and hBN phases,

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 817

FIGURE 19.33 Cross-sectional transmission electron microscope image of a few-layer hexagonal boron nitride thin film sputtered onto epitaxial Ru(0001)/Al2O3(0001). Reprinted with permission from [95].

FIGURE 19.34 RHEED patterns of a 100 nm hexagonal boron nitride film on a Ni(111) surface at [110] (a) and [-1-12] (b) azimuth of Ni(111). Adapted with permission from [97].

depending on the deposition conditions. In this method, boron is evaporated via electron beam onto a substrate (e.g., hydrogen-passivated Si) held at high temperature (w400
C) that is simultaneously bombarded with argon and nitrogen ions from an ion source [98,128–130] (See Figure 19.35). Under some conditions (namely high boron flux ˚ /s), the resulting films exhibit only spectroscopic signatures corresponding rates of 0.5 A to hBN [98]. Independent fluxes of boron and nitrogen at the sample surface can create nonstoichiometric films [130], a hindrance when attempting to deposit intrinsic hBN.

19.4. Chemical Vapor Deposition of Molybdenum Disulfide It is often overlooked that atomically thin MoS2 was “discovered” nearly concurrently with graphene. Within months of Geim’s seminal work, which isolated 2D graphene [1], researchers also successfully isolated and characterized single layers of MoS2, as well as NbSe2 and Bi2Sr2CaCu2Ox [131]. Immediately following the successful isolation of 2D crystals, many experimental and theoretical efforts focused on graphene due to its exotic properties, as mentioned in previous sections. In contrast to the linear, gapless energy dispersion of graphene, single-layer MoS2 is a direct-gap semiconductor with reasonable electronic mobility [132,133] and high optical responsivity [134], making it a promising material for electronic components [135], optoelectronic applications [136], and chemical sensing [137], which has stimulated a resurging interest in the field of thin MoS2 crystals. While MoS2 in bulk form has been investigated for many years as a dry

818

HANDBOOK OF CRYSTAL GROWTH

5 8

4 7

9

3 6 1

2

FIGURE 19.35 Schematic of an ion beam assisted electron beam deposition system for BN film deposition noting the cryopump (1), Ar/N2 ion beam source (2), beam neutralizer (3), electron gun with boron crucible (6), thickness monitor (7), and substrate (8) along with optional beam profile monitor (4), mass spectrometer (5), and Auger electron analyzer (9). Reprinted with permission from [130].

lubricant [138] and as an industrial hydrodesulfurization catalyst [139], atomically thin layers offer a new suite of physics and possible applications to explore. Interestingly, an abrupt transition from indirect-gap to direct-gap semiconductor occurs at the single monolayer thickness [134]. Additionally, due to inequivalent K and K0 valleys at the Brillouin zone edge, it is possible to achieve full valley polarization by exciting carriers with circularly polarized light [140]. This suggests interesting phenomena and applications in spin- and valley-tronics for single MoS2 layers. MoS2 is only one member of the large group of materials referred to as the TMDs. All TMDs can be described by the chemical formula MX2, where M is a transition metal atom (from group 4–10 of the periodic table) and X is a chalcogen species (e.g. S, Se, Te). There are over 40 known TMDs that exist in bulk form [141,142], with many exhibiting a layer-type structure. Bulk TMDs have been investigated for many years [143–145] and demonstrate a wide array of electronic properties, including insulators such as HfS2, semiconductors such as MoS2, to true metals such as NbS2 and TiSe2. Low-temperature studies have also uncovered exotic properties such as superconductivity, charge density waves, and Mott transitions in certain TMDs [146]. Thinning of materials to very few or single atomic layers induces confinement effects, which lead to new characteristics. Although in its infancy, many investigations and synthesis methods of atomically thin TMDs have been concentrated on MoS2. We therefore focus the majority of our attention on molybdenum disulfide and will briefly mention synthesis techniques for additional

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 819

FIGURE 19.36 The crystal structure of MoS2. The top view shows a hexagonal structure with Mo and S occupying alternate corners. The trigonal prismatic unit cell of a single layer of MoS2 consists of three atomic sheets in an S-Mo-S stack. Reproduced from [147].

TMD materials at the conclusion of the section. As evident in Figure 19.36 [147], a single layer of MoS2 is actually composed of three atomic layers: a top and bottom sulfur layer sandwich the center Mo sheet. From the bird’s eye view, this results in a hexagonal structure, similar to graphene. The growth of MoS2 crystals provides challenges that are unique to the synthesis processes of TMDs. As opposed to graphene and hBN synthesis, there are no gaseous precursors that can be utilized, necessitating the use of solid source or liquid precursors. Additionally, the majority of MoS2 synthesis methods to date are catalyst-free. In the absence of catalyst-mediated synthesis, minor perturbations in parameters such as precursor-to-substrate distance, growth temperature, growth time, and surface treatment may strongly impact growth dynamics and crystalline quality. Synthesis of a continuous and uniform single layer of MoS2 remains a challenge. Additionally, due to sulfur’s reactivity with most metallic species at elevated temperatures, synthesis of MoS2 is traditionally performed on insulating substrates, making characterization with STM a challenge. While photoluminescence spectroscopy (PL) and Raman spectroscopy are not often considered surface science techniques, they have become seminal techniques in the characterization of thin films of MoS2. The transformation from indirect to direct-gap semiconductor at monolayer thickness is accompanied by a significant enhancement in the PL intensity as well as a strong emission peak near 1.9 eV [134]. As evident in Figure 19.37(a) and Figure 19.37(b), the PL spectra of single-atomic MoS2 layers are easily differentiated from multilayers. The peak separation between the characteristic Raman peaks, E2g1 and A1g, is dependent on the layer number and is found to

820

HANDBOOK OF CRYSTAL GROWTH

FIGURE 19.37 Photoluminescence (PL) (a,b) and Raman (c,d) characterization of thin MoS2 films. (a) A strong decrease is evident in PL intensity as layer number increases. (b) Multilayer films develop a low energy PL emission peak, associated with the indirect band gap. (c,d) The separation between E2g1 Raman peaks decreases with decreasing thickness. Source: Figure (a) and (b) from ref [134] and (c) and (d) from ref [148]

decrease with decreasing thickness [148], as seen in Figure 19.37(c) and Figure 19.37(d), providing a fingerprint of layer number. We discuss several of the methods that have been utilized to synthesize thin films of MoS2. The majority of growth processes can be performed in a quartz tube furnace and have implemented one of the following precursors: (a) thin films of Mo metal, (b) liquid solutions containing ammonia tetrathiomolybdate, or (c) MoO3 powder. A schematic of the growth processes associated with the various precursors is shown schematically below in Figure 19.38. Single to few-layer MoS2 films can be obtained directly by exposing thin films of Mo metal predeposited onto a desired substrate to sulfur vapor while at elevated temperatures

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 821

FIGURE 19.38 Schematic description of common MoS2 synthesis techniques. (a) Thin films of predeposited Mo are exposed to sulfur at elevated Temperatures. (b) Substrates are dipped in (NH4)2MoS4 then subjected to a two-step thermal anneal. (c) Sulfur reduces MoO3 precursor to form MoS2. Reproduced with permission from [149].

[150–153]. The MoS2 formation is a product of elemental reactions between Mo and S, as sulfur is known to form sulfides with most metals. Stoichiometric and continuous films of nonuniform thickness are attainable on a variety of substrates by implementing this technique. The substrate dimensions determine the size of the synthesized film, providing a straightforward means to scale up the MoS2. The thickness of predeposited Mo metal provides rough control over the MoS2 layer number, although thinner Mo films can result in discontinuous islands of MoS2 [150]. Process temperatures, as well as substrate, can dramatically affect crystal quality, with higher temperatures and sapphire substrates demonstrating highest quality growth [151,152]. The sulfur source may also strongly impact the uniformity of the resulting MoS2. Through the use of H2S gas, as opposed to the more commonly used vaporized elemental sulfur, uniform films from 2 to 12 layers were synthesized [152]. The H2S gas is believed to decompose more readily than elemental sulfur, contributing to the improved uniformity. Typical grain sizes are found to be on the order of 10–30 nm, considerably smaller than that of mechanically exfoliated MoS2 crystals.

822

HANDBOOK OF CRYSTAL GROWTH

The thermal decomposition of ammonium tetrathiomolybdate (NH4)2MoS4 has also been utilized to achieve MoS2 films and takes place in two distinct steps [154]. Above temperatures of 155
C, (NH4)2MoS4 decomposes into molybdenum trisulfide following [Eqn 19.1]. (NH4)2MoS4 / 2NH3 þ H2S þ MoS3


(1)

At further elevated temperatures, above 335 C, MoS3 decomposes into elemental sulfur and MoS2. The dissolution of (NH4)2MoS4 in N, N-dimethylformamide provides a liquid precursor that can be coated on a substrate [155] or bubbled through in order to transport vapor to the substrate [156]. Large area and uniform trilayer MoS2 was synthesized in a two-step thermolysis of (NH4)2MoS4 solution [155]. In this technique, an insulating substrate (SiO2 or sapphire) was immersed in the solution then slowly withdrawn to achieve uniform coverage prior to annealing. Raman spectroscopy confirmed the presence of MoS2 following the initial anneal at 500
C, albeit of poor quality. A second anneal at 1000
C was necessary to improve crystallinity. Resulting films were polycrystalline, exhibiting grain size on the order of 160 nm. In spite of the small grain size, films exhibited high on-off ratio and electronic mobility comparable to exfoliated material. As an alternative to coating substrates in solution, a second technique demonstrated the liquid precursor could be transported to the substrate by bubbling argon through the liquid [156]. CVD-grown graphene was selected as the substrate due to its lack of dangling bonds and weak out of plane van der Waals forces. The interface properties strongly impact the resulting growth, as no MoS2 was observed on bare Cu foil or SiO2. At the relatively low temperature of 400
C, hexagonal nanoflakes, having lateral dimensions as large as several microns with thicknesses on the order of several nm, were synthesized. The morphology of MoS2 varied from isolated nanoflakes, to loosely packed nanoflakes connected by a thin MoS2 film, to closely packed MoS2 flakes, as evident in Figure 19.39. A large fraction of MoS2 grains were epitaxially grown and all tended to orient as closely as possible to the underlying graphene grains, stressing the importance of the graphene substrate in this particular growth technique. A technique that has been gaining popularity utilizes solid sources of MoO3 and elemental sulfur in a standard quartz furnace [24,149,155,157–162]. Typically, the desired substrate (SiO2, sapphire, or mica) is positioned in close proximity to MoO3 powder at the center of a quartz tube furnace with sulfur located upstream in a low-temperature zone. Although exact recipes vary from lab to lab, processes are often performed in inert atmosphere (Ar or N) at ambient conditions, with temperatures in the range of w600
C–1000
C. In this temperature range, sulfur vapor partially reduces MoO3 to form volatile MoO3-x. MoO3-x is subsequently adsorbed onto the substrate where it diffuses across the surface and further reacts with sulfur to produce MoS2 as described by the equation: S2 þ MoO3-x/MoS2 þ SO2. Single-layer equilateral triangles, several to tens of microns in length, are reliably produced with this technique (Figure 19.40(a)). The triangular structure is expected for sulfur-terminated crystals and is an indication of high quality, single grains. Under the right conditions, the triangles can coalesce to form a

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 823

(a)

(e)

(b) Wrinkle

Wrinkle

0 nm

(c)

0

20 μm

0

0 nm

flakes with MoS2 thin film in between

Stage 2: Loosely packed MoS2

5 μm

300 nm

300 nm

Stage 3:

5 μm

Closely packed MoS2 flakes

0 nm

0 nm

MoS2 flakes

flakes formation

(d)

MoS2 film

0

Stage 1: Isolated MoS2

300 nm

300 nm

0

5 μm

MoS2 flakes

Graphene

Cu foil

FIGURE 19.39 MoS2 nanoflakes synthesized by Shi et al. The AFM images show the progression from isolated nanoflakes at low precursor exposure (a,b) to closely packed flakes with a background MoS2 thin film for larger precursor exposure. (e) A schematic diagram of the various stages of growth. Reprinted with permission from [156].

FIGURE 19.40 The growth of MoS2 using MoO3 precursor. (a–d) SEM images showing the growth of single layer MoS2 from small triangles to continuous films [149]. Reprinted with permission.

continuous, polycrystalline film (Figure 19.40(d)). The progression of MoS2 growth on the substrate is nucleation-limited, meaning growth typically occurs at rare and specific sites on the substrate and then progresses outward from these locations. In an effort to enhance the number of nucleation sites, researchers have modified the growth substrates with “seeding particles” [163,164] such as reduced graphene oxide or perylene3,4,9,10-tetracarboxylic acid tetrapotassium. The higher density of nucleation sites provided by these seeding particles promotes formation and lateral growth of single layer MoS2. Artificial edges on the substrate created via conventional lithography techniques have also been explored as a means to control nucleation sites [149]. Not all synthesis techniques fall into the categories described above. An alternative synthesis technique utilizes MoCl5 as the source of Mo metal [165]. Single-layer, large area uniform and continuous films are synthesized, having grains on the order of hundreds of nm.

824

HANDBOOK OF CRYSTAL GROWTH

Thicker films can be controllably synthesized by modifying growth pressure and precursor amount. A physical vapor transport technique, using only MoS2 powder, has produced highquality triangular crystals on the order of several mm [166]. Gold films have been used as a catalyst for MoS2 synthesis by which Mo(CO)6 decomposed into an Au-Mo alloy, which then transformed into few-layer MoS2 upon H2S exposure [167]. While the synthesis of MoS2 has received the most attention thus far, several other TMDs have recently been synthesized. Analogous to the techniques discussed above, WS2 thin films have been synthesized by sulfurization of predeposited WO3 [168,169] and W films [170], as well as through vapor phase reaction of WO3 and sulfur [171,172]. Interestingly, Zhang et al. observed improved edge structure when hydrogen gas was introduced during the synthesis process, which may be a result of the highly reductive nature of hydrogen. Due to the reduced reactivity of Se compared to S, fewer procedures have been reported for MoSe2 and WSe2 synthesis. Very recently, MoSe2 monolayers and few-layer films were achieved using MoO3 and Se precursors [173–175]. Similarly, vapor phase reaction of WO3 and Se produced crystals of monolayer WSe2 [169,176]. In both cases, the presence of hydrogen was necessary to activate the selenization of the metallic species. The synthesis of 2D TMDs has clearly experienced rapid progress and growth in a very limited time. With possible applications in nanoelectronics, optoelectronics, and spintronics on the horizon, it is expected that this trajectory will continue into the future. In this nascent field, there are many exciting opportunities to develop high-quality materials, as well as expand the fundamental understanding of 2D crystals.

References [1] Novoselov KS, Geim AK, Morozov SV, Jiang D, Zhang Y, Dubonos SV, et al. Electric field effect in atomically thin carbon films. Science 2004;306:666–9. http://dx.doi.org/10.1126/science.1102896. [2] Landau LD. Zur Theorie der phasenumwandlungen II. Phys Z Sowjetunion 1937;11:26–35. [3] Peierls R. Quelques proprie´te´s typiques des corps solides. Ann Inst Henri Poincare´ 1935;5:177–222. [4] Geim AK, Novoselov KS. The rise of graphene. Nat Mater 2007;6:183. [5] Wallace PR. The band theory of graphite. Phys Rev 1947;71:622. http://dx.doi.org/10.1103/ PhysRev.71.622. [6] Cao H, Yu Q, Jauregui LA, Tian J, Wu W, Liu Z, et al. Electronic transport in chemical vapor deposited graphene synthesized on Cu: quantum Hall effect and weak localization. Appl Phys Lett 2010;96:122106–122106. [7] Land TA, Michely T, Behm RJ, Hemminger JC, Comsa G. STM investigation of single layer graphite structures produced on Pt(111) by hydrocarbon decomposition. Surf Sci 1992;264:261–70. http:// dx.doi.org/10.1016/0039-6028(92)90183-7. [8] Alstrup I, Chorkendorff I, Ullmann S. The interaction of CH4 at high temperatures with clean and oxygen precovered Cu(100). Surf Sci 1992;264:95. http://dx.doi.org/10.1016/0039-6028(92)90168-6. [9] Gall, Mikhallov, Rut’kov, Tontegode AY. Nature of the adsorption binding between a graphite monolayer and rhenium surface. Sov Phys J Solid State 1985;27:1410–4. [10] Nagashima A, Nuka K, Satoh K, Itoh H, Ichinokawa T, Oshima C, et al. Electronic structure of monolayer graphite on some transition metal carbide surfaces. Surf Sci 1993;287:609–13.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 825

[11] Gamo Y, Nagashima A, Wakabayashi M, Terai M, Oshima C. Atomic structure of monolayer graphite formed on Ni(111). Surf Sci 1997;374:61–4. http://dx.doi.org/10.1016/S0039-6028(96) 00785-6. [12] Rut’kov EV, Tontegode AY. A study of the carbon adlayer on iridium. Surf Sci 1985;161:373–89. http://dx.doi.org/10.1016/0039-6028(85)90816-7. [13] Hagstrom S, Lyon HB, Somorjai GA. Surface structures on the clean platinum (100) surface. Phys Rev Lett 1965;15:491–3. http://dx.doi.org/10.1103/PhysRevLett.15.491. [14] Grant JT, Haas TW. A study of Ru(0001) and Rh(111) surfaces using LEED and Auger electron spectroscopy. Surf Sci 1970;21:76–85. http://dx.doi.org/10.1016/0039-6028(70)90064-6. [15] Morgan AE, Somorjai GA. Low energy electron diffraction studies of gas adsorption on the platinum (100) single crystal surface. Surf Sci 1968;12:405–25. http://dx.doi.org/10.1016/0039-6028(68) 90089-7. [16] May J. Platinum surface LEED rings. Surf Sci 1969;17:269–70. [17] Batzill M. The surface science of graphene: metal interfaces, CVD synthesis, nanoribbons, chemical modifications, and defects. Surf Sci Rep 2012;67:83–115. http://dx.doi.org/10.1016/j. surfrep.2011.12.001. [18] Wintterlin J, Bocquet M-L. Graphene on metal surfaces. Surf Sci 2009;603:1841–52. http://dx.doi. org/10.1016/j.susc.2008.08.037. [19] Arya SPS, D’Amico A. Preparation, properties and applications of boron nitride thin films. Thin Solid Films 1988;157:267–82. http://dx.doi.org/10.1016/0040-6090(88)90008-9. [20] Bresnehan MS, Hollander MJ, Wetherington M, Wang K, Miyagi T, Pastir G, et al. Prospects of direct growth boron nitride films as substrates for graphene electronics. J Mater Res 2014;29: 459–71. http://dx.doi.org/10.1557/jmr.2013.323. [21] Ganatra R, Zhang Q. Few layer MoS2–A promising layered semiconductor. ACS Nano 2014;8: 4074–99. [22] Ariwala D, Sangwan VK, Lauhon LJ, Marks TJ, Hersam MC. Emerging device applications for semiconducting two-dimensional transition metal dichalcogenides. ACS Nano 2014;8:1102–20. [23] Xu M, Liang T, Shi M, Chen H. Graphene-like two-dimensional materials. Chem Rev 2013;113: 3766–98. [24] Balendhran S, Walia S, Nili H, Ou JZ, Zhuiykov S, Kaner RB, et al. Two-dimensional molybdenum trioxide and dichalcogenides. Adv Funct Mater 2013;23:3952–70. [25] Castro Neto AH, Guinea F, Peres NMR, Novoselov KS, Geim AK. The electronic properties of graphene. Rev Mod Phys 2009;81:109. http://dx.doi.org/10.1103/RevModPhys.81.109. [26] Geim AK. Graphene: status and prospects. Science 2009;324:1530–4. http://dx.doi.org/10.1126/ science.1158877. [27] Sarma S Das, Adam S, Hwang EH, Rossi E. Electronic transport in two-dimensional graphene. Rev Mod Phys 2011;83:407–70. http://dx.doi.org/10.1103/RevModPhys.83.407. [28] Avouris P. Graphene: electronic and photonic properties and devices. Nano Lett 2010;10:4285–94. http://dx.doi.org/10.1021/nl102824h. [29] Balandin AA, Ghosh S, Bao W, Calizo I, Teweldebrhan D, Miao F, et al. Superior thermal conductivity of single-layer graphene. Nano Lett 2008;8:902. http://dx.doi.org/10.1021/nl0731872. [30] Weinberg WH, Deans HA, Merrill RP. The structure and chemistry of ethylene adsorbed on platinum (111): CFSO-BEBO calculations and experimental studies. Surf Sci 1974;41:312–36. http://dx. doi.org/10.1016/0039-6028(74)90311-2. [31] L B. A LEED study of the deposition of carbon on platinum crystal surfaces. Surf Sci 1975;53: 317–29. http://dx.doi.org/10.1016/0039-6028(75)90132-6.

826

HANDBOOK OF CRYSTAL GROWTH

[32] Hamilton JC, Blakely JM. Carbon segregation to single crystal surfaces of Pt, Pd and Co. Surf Sci 1980;91:199–217. http://dx.doi.org/10.1016/0039-6028(80)90080-1. [33] Abon M, Billy J, Bertolini JC, Tardy B. Carbon on Pt(111): characterization and influence on the chemisorptive properties. Surf Sci 1986;167:1–17. http://dx.doi.org/10.1016/0039-6028(86) 90782-X. [34] Sutter PW, Flege JI, Sutter EA. Epitaxial graphene on ruthenium. Nat Mater 2008;7:406–11. [35] Sutter E, Albrecht P, Sutter P. Graphene growth on polycrystalline Ru thin films. Appl Phys Lett 2009;95:133109. http://dx.doi.org/10.1063/1.3224913. [36] Sutter E, Acharya DP, Sadowski JT, Sutter P. Scanning tunneling microscopy on epitaxial bilayer graphene on ruthenium (0001). Appl Phys Lett 2009;94:133101. http://dx.doi.org/10.1063/1. 3106057. [37] Li X, Cai W, Colombo L, Ruoff RS. Evolution of graphene growth on Ni and Cu by carbon isotope labeling. Nano Lett 2009;9:4268–72. http://dx.doi.org/10.1021/nl902515k. [38] Arnoult WJ, McLellan RB. The solubility of carbon in rhodium ruthenium, iridium and rhenium. Scr Metall 1972;6:1013–8. [39] Aizawa T, Souda R, Otani S, Ishizawa Y, Oshima C. Anomalous bond of monolayer graphite on transition-metal carbide surfaces. Phys Rev Lett 1990;64:768–71. http://dx.doi.org/10.1103/ PhysRevLett.64.768. [40] Lahiri J, Miller TS, Ross AJ, Adamska L, Oleynik II, Batzill M. Graphene growth and stability at nickel surfaces. New J Phys 2011;13:025001. http://dx.doi.org/10.1088/1367-2630/13/2/025001. [41] Natesan K, Kassner TF. Thermodynamics of carbon in nickel, iron-nickel and iron-chromiumnickel alloys. Metall Trans 1973;4:2557–66. [42] Siller RH, Oates WA, McLellan RB. The solubility of carbon in palladium and platinum. J Common Met 1968;16:71–3. [43] Lo´pez GA, Mittemeijer EJ. The solubility of C in solid Cu. Scr Mater 2004;51:1–5. http://dx.doi.org/ 10.1016/j.scriptamat.2004.03.028. [44] Bever MB, Floe CF. Solubility of carbon in molten copper. AIME Trans 1946;166:128–41. [45] Nagashima A, Tejima N, Oshima C. Electronic states of the pristine and alkali-metal-intercalated monolayer graphite/Ni(111) systems. Phys Rev B 1994;50:17487–95. http://dx.doi.org/10.1103/ PhysRevB.50.17487. [46] Addou R, Dahal A, Sutter P, Batzill M. Monolayer graphene growth on Ni(111) by low temperature chemical vapor deposition. Appl Phys Lett 2012;100:021601. http://dx.doi.org/10.1063/1.3675481. [47] Banerjee BC, Hirt TJ, Walker PL. Pyrolytic carbon formation from carbon suboxide. Nature 1961; 192:450–1. http://dx.doi.org/10.1038/192450a0. [48] Karu AE, Beer M. Pyrolytic formation of highly crystalline graphite films. J Appl Phys 1966;37: 2179–81. http://dx.doi.org/10.1063/1.1708759. [49] Somani PR, Somani SP, Umeno M. Planer nano-graphenes from camphor by CVD. Chem Phys Lett 2006;430:56–9. http://dx.doi.org/10.1016/j.cplett.2006.06.081. [50] Obraztsov AN, Obraztsova EA, Tyurnina AV, Zolotukhin AA. Chemical vapor deposition of thin graphite films of nanometer thickness. Carbon 2007;45:2017–21. http://dx.doi.org/10.1016/j. carbon.2007.05.028. [51] Reina A, Jia X, Ho J, Nezich D, Son H, Bulovic V, et al. Large area, few-layer graphene films on arbitrary substrates by chemical vapor deposition. Nano Lett 2008;9:30–5. [52] Kim KS, Zhao Y, Jang H, Lee SY, Kim JM, Kim KS, et al. Large-scale pattern growth of graphene films for stretchable transparent electrodes. Nature 2009;457:706–10. http://dx.doi.org/10.1038/ nature07719.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 827

[53] Shelton JC, Patil HR, Blakely JM. Equilibrium segregation of carbon to a nickel (111) surface: a surface phase transition. Surf Sci 1974;43:493–520. http://dx.doi.org/10.1016/0039-6028(74)90272-6. [54] Lahiri J, Miller T, Adamska L, Oleynik II, Batzill M. Graphene growth on Ni(111) by transformation of a surface carbide. Nano Lett 2011;11:518–22. http://dx.doi.org/10.1021/nl103383b. [55] Li X, Cai W, An J, Kim S, Nah J, Yang D, et al. Large-area synthesis of high-quality and uniform graphene films on copper foils. Science 2009;324:1312–4. http://dx.doi.org/10.1126/science. 1171245. [56] Rasool HI, Song EB, Allen MJ, Wassei JK, Kaner RB, Wang KL, et al. Continuity of graphene on polycrystalline copper. Nano Lett 2010;11:251–6. http://dx.doi.org/10.1021/nl1036403. [57] Yazyev OV, Louie SG. Electronic transport in polycrystalline graphene. Nat Mater 2010;9:806–9. http://dx.doi.org/10.1038/nmat2830. [58] Clark KW, Zhang X-G, Vlassiouk IV, He G, Feenstra RM, Li A-P. Spatially resolved mapping of electrical conductivity across individual domain (Grain) boundaries in graphene. ACS Nano 2013; 7:7956–66. http://dx.doi.org/10.1021/nn403056k. [59] Li X, Magnuson CW, Venugopal A, Tromp RM, Hannon JB, Vogel EM, et al. Large-area graphene single crystals grown by low-pressure chemical vapor deposition of methane on copper. J Am Chem Soc 2011;133:2816–9. http://dx.doi.org/10.1021/ja109793s. [60] Huang PY, Ruiz-Vargas CS, van der Zande AM, Whitney WS, Levendorf MP, Kevek JW, et al. Grains and grain boundaries in single-layer graphene atomic patchwork quilts. Nature 2011;469:389–92. 10.1038nature09718. [61] Wofford JM, Nie S, McCarty KF, Bartelt NC, Dubon OD. Graphene islands on Cu foils: the interplay between shape, orientation, and defects. Nano Lett 2010;10:4890–6. http://dx.doi.org/10.1021/ nl102788f. [62] Robinson ZR, Tyagi P, Mowll TR, Ventrice CA, Hannon JB. Argon-assisted growth of epitaxial graphene on Cu(111). Phys Rev B 2012;86:235413. http://dx.doi.org/10.1103/PhysRevB.86.235413. [63] Ji S-H, Hannon JB, Tromp RM, Perebeinos V, Tersoff J, Ross FM. Atomic-scale transport in epitaxial graphene. Nat Mater 2012;11:114–9. http://dx.doi.org/10.1038/nmat3170. [64] Barrett CS, Massalski TB. Structure of metals. 3rd ed. McGraw–Hill; 1966. [65] Robinson ZR, Tyagi P, Murray TM, Ventrice CA, Chen S, Munson A, et al. Substrate grain size and orientation of Cu and Cu–Ni foils used for the growth of graphene films. J Vac Sci Technol Vac Surf Films 2012;30:011401. http://dx.doi.org/10.1116/1.3663877. [66] Wood JD, Schmucker SW, Lyons AS, Pop E, Lyding JW. Effects of polycrystalline Cu substrate on graphene growth by chemical vapor deposition. Nano Lett 2011;11:4547–54. http://dx.doi.org/10. 1021/nl201566c. [67] Robinson ZR, Ong EW, Mowll TR, Tyagi P, Gaskill DK, Geisler H, et al. Influence of chemisorbed oxygen on the growth of graphene on Cu(100) by chemical vapor deposition. J Phys Chem C 2013; 117:23919–27. http://dx.doi.org/10.1021/jp410142r. [68] Hao Y, Bharathi MS, Wang L, Liu Y, Chen H, Nie S, et al. The role of surface oxygen in the growth of large single-crystal graphene on copper. Science 2013;342:720–3. http://dx.doi.org/10.1126/ science.1243879. [69] Gao L, Guest JR, Guisinger NP. Epitaxial graphene on Cu(111). Nano Lett 2010;10:3512–6. http:// dx.doi.org/10.1021/nl1016706. [70] Nie S, Wofford JM, Bartelt NC, Dubon OD, McCarty KF. Origin of the mosaicity in graphene grown on Cu(111). Phys Rev B 2011;84:155425. http://dx.doi.org/10.1103/PhysRevB.84.155425. [71] Busse C, Lazi c P, Djemour R, Coraux J, Gerber T, Atodiresei N, et al. Graphene on Ir (111): physisorption with chemical modulation. Phys Rev Lett 2011;107:036101.

828

HANDBOOK OF CRYSTAL GROWTH

[72] Coraux J, Busse C, Michely T. Structural coherency of graphene on Ir (111). Nano Lett 2008;8: 565–70. [73] Loginova E, Bartelt NC, Feibelman PJ, McCarty KF. Factors influencing graphene growth on metal surfaces. New J Phys 2009;11:063046. http://dx.doi.org/10.1088/1367-2630/11/6/063046. [74] N’Diaye AT, Engler M, Busse C, Wall D, Buckanie N, Meyer zu Heringdorf F-J, et al. Growth of graphene on Ir(111). New J Phys 2009;11:023006. http://dx.doi.org/10.1088/1367-2630/11/2/ 023006. ¨ rmer K, Bartelt NC, McCarty KF. Defects of graphene on Ir(111): rotational [75] Loginova E, Nie S, Thu domains and ridges. Phys Rev B 2009;80:085430. http://dx.doi.org/10.1103/PhysRevB.80.085430. [76] Coraux J, N’Diaye AT, Engler M, Busse C, Wall D, Buckanie N, et al. Growth of graphene on Ir(111). New J Phys 2009;11:023006. http://dx.doi.org/10.1088/1367-2630/11/2/023006. [77] N’Diaye AT, Bleikamp S, Feibelman PJ, Michely T. Two-dimensional Ir cluster lattice on a graphene moire´ on Ir(111). Phys Rev Lett 2006;97:215501. http://dx.doi.org/10.1103/PhysRevLett.97. 215501. [78] Pletikosic I, Kralj M, Pervan P, Brako R, Coraux J, N’Diaye AT, et al. Dirac cones and minigaps for graphene on Ir(111). Phys Rev Lett 2009;102:056808. http://dx.doi.org/10.1103/PhysRevLett.102. 056808. [79] Kim KS, Zhao Y, Jang H, Lee SY, Kim JM, Kim KS, et al. Large-scale pattern growth of graphene films for stretchable transparent electrodes. Nature 2009;457:706–10. [80] Ohta T, Robinson JT, Feibelman PJ, Bostwick A, Rotenberg E, Beechem TE. Evidence for interlayer coupling and moire´ periodic potentials in twisted bilayer graphene. Phys Rev Lett 2012;109:186807. http://dx.doi.org/10.1103/PhysRevLett.109.186807. [81] Pease RS. An X-ray study of boron nitride. Acta Crystallogr 1952;5:356–61. http://dx.doi.org/10. 1107/S0365110X52001064. [82] Doni E, Parravic GP. Energy bands and optical properties of hexagonal boron nitride and graphite. Nuovo Cimento Della Soc Ital Fis B-Gen Phys Relativ Astron Math Phys Methods 1969;64:117. http://dx.doi.org/10.1007/BF02710286. [83] Taylor R, Coulson CA. Studies in graphite and related compounds III: electronic band structure in boron nitride. Proc Phys Soc Sect A 1952;65:834. http://dx.doi.org/10.1088/0370-1298/65/10/307. [84] Deacon RF, Goodman JF. Lubrication by lamellar solids. Proc R Soc Lond Ser Math Phys Sci 1958; 243:464–82. http://dx.doi.org/10.1098/rspa.1958.0013. [85] Watanabe K, Taniguchi T, Kanda H. Direct-bandgap properties and evidence for ultraviolet lasing of hexagonal boron nitride single crystal. Nat Mater 2004;3:404–9. http://dx.doi.org/10.1038/nmat1134. [86] Kubota Y, Watanabe K, Tsuda O, Taniguchi T. Deep ultraviolet light-emitting hexagonal boron nitride synthesized at atmospheric pressure. Science 2007;317:932–4. http://dx.doi.org/10.1126/ science.1144216. [87] Britnell L, Gorbachev RV, Jalil R, Belle BD, Schedin F, Mishchenko A, et al. Field-effect tunneling transistor based on vertical graphene heterostructures. Science 2012;335:947–50. http://dx.doi.org/ 10.1126/science.1218461. [88] Dean CR, Young AF, Meric I, Lee C, Wang L, Sorgenfrei S, et al. Boron nitride substrates for highquality graphene electronics. Nat Nanotechnol 2010;5:722–6. http://dx.doi.org/10.1038/nnano. 2010.172. [89] Liu Z, Song L, Zhao S, Huang J, Ma L, Zhang J, et al. Direct growth of graphene/hexagonal boron nitride stacked layers. Nano Lett 2011;11:2032–7. http://dx.doi.org/10.1021/nl200464j. [90] Yankowitz M, Xue J, Cormode D, Sanchez-Yamagishi JD, Watanabe K, Taniguchi T, et al. Emergence of superlattice dirac points in graphene on hexagonal boron nitride. Nat Phys 2012;8: 382–6. http://dx.doi.org/10.1038/nphys2272.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 829

[91] Frueh S, Kellett R, Mallery C, Molter T, Willis WS, King’ondu C, et al. Pyrolytic decomposition of ammonia borane to Boron nitride. Inorg Chem 2011;50:783–92. http://dx.doi.org/10.1021/ ic101020k. [92] Levendorf MP, Kim C-J, Brown L, Huang PY, Havener RW, Muller DA, et al. Graphene and boron nitride lateral heterostructures for atomically thin circuitry. Nature 2012;488:627–32. http://dx.doi. org/10.1038/nature11408. [93] Ci L, Song L, Jin C, Jariwala D, Wu D, Li Y, et al. Atomic layers of hybridized boron nitride and graphene domains. Nat Mater 2010;9:430–5. http://dx.doi.org/10.1038/nmat2711. [94] Jin-Xiang D, Xiao-Kang Z, Qian Y, Xu-Yang W, Guang-Hua C, De-Yan H. Optical properties of hexagonal boron nitride thin films deposited by radio frequency bias magnetron sputtering. Chin Phys B 2009;18:4013. http://dx.doi.org/10.1088/1674-1056/18/9/066. [95] Sutter P, Lahiri J, Zahl P, Wang B, Sutter E. Scalable synthesis of uniform few-layer hexagonal boron nitride dielectric films. Nano Lett 2013;13:276–81. http://dx.doi.org/10.1021/nl304080y. [96] Gupta VK, Wamsley CC, Koch MW, Wicks GW. Molecular beam epitaxy growth of boroncontaining nitrides. J Vac Sci Technol B 1999;17:1246–8. http://dx.doi.org/10.1116/1.590731. [97] Tsai CL, Kobayashi Y, Akasaka T, Kasu M. Molecular beam epitaxial growth of hexagonal boron nitride on Ni(111) substrate. J Cryst Growth 2009;311:3054–7. http://dx.doi.org/10.1016/j.jcrysgro. 2009.01.077. [98] Plass MF, Fukarek W, Kolitsch A, Kreibig U. Growth and characterisation of boron nitride thin films. Surf Coat Technol 1996;84:383–7. http://dx.doi.org/10.1016/S0257-8972(95)02781-5. [99] Yokota K, Kimura H, Miyashita F. Crystalline lattice phase-conversion on thin boron nitride films deposited on silicon wafers by an ion beam assisted deposition method. Nucl Instrum Methods Phys Res Sect B Beam Interact Mater At 2007;257:459–62. http://dx.doi.org/10.1016/j.nimb.2007. 01.097. [100] Kim G, Jang A-R, Jeong HY, Lee Z, Kang DJ, Shin HS. Growth of high-crystalline, single-layer hexagonal boron nitride on recyclable platinum foil. Nano Lett 2013;13:1834–9. http://dx.doi.org/ 10.1021/nl400559s. [101] Ismach A, Chou H, Ferrer DA, Wu Y, McDonnell S, Floresca HC, et al. Toward the controlled synthesis of hexagonal Boron nitride films. ACS Nano 2012;6:6378–85. http://dx.doi.org/10.1021/ nn301940k. [102] Rand MJ, Roberts JF. Preparation and properties of thin film boron nitride. J Electrochem Soc 1968; 115:423–9. http://dx.doi.org/10.1149/1.2411238. [103] Takahashi T, Itoh H, Takeuchi A. Chemical vapor deposition of hexagonal boron nitride thick film on iron. J Cryst Growth 1979;47:245–50. http://dx.doi.org/10.1016/0022-0248(79)90248-3. [104] Auwa¨rter W, Kreutz TJ, Greber T, Osterwalder J. XPD and STM investigation of hexagonal boron nitride on Ni(111). Surf Sci 1999;429:229–36. http://dx.doi.org/10.1016/S0039-6028(99)00381-7. [105] Goriachko A, He Y, Knapp M, Over H, Corso M, Brugger T, et al. Self-assembly of a hexagonal Boron nitride nanomesh on Ru(0001). Langmuir 2007;23:2928–31. http://dx.doi.org/10.1021/ la062990t. [106] Corso M, Auwa¨rter W, Muntwiler M, Tamai A, Greber T, Osterwalder J. Boron nitride nanomesh. Science 2004;303:217–20. http://dx.doi.org/10.1126/science.1091979. [107] Kim KK, Hsu A, Jia X, Kim SM, Shi Y, Dresselhaus M, et al. Synthesis and characterization of hexagonal boron nitride film as a dielectric layer for graphene devices. ACS Nano 2012;6:8583–90. http://dx.doi.org/10.1021/nn301675f. [108] Shi Y, Hamsen C, Jia X, Kim KK, Reina A, Hofmann M, et al. Synthesis of few-layer hexagonal Boron nitride thin film by chemical vapor deposition. Nano Lett 2010;10:4134–9. http://dx.doi.org/10. 1021/nl1023707.

830

HANDBOOK OF CRYSTAL GROWTH

[109] Sutter P, Lahiri J, Albrecht P, Sutter E. Chemical vapor deposition and etching of high-quality monolayer hexagonal Boron nitride films. ACS Nano 2011;5:7303–9. http://dx.doi.org/10.1021/ nn202141k. [110] Nagashima A, Tejima N, Gamou Y, Kawai T, Oshima C. Electronic dispersion relations of monolayer hexagonal boron nitride formed on the Ni(111) surface. Phys Rev B 1995;51:4606–13. http:// dx.doi.org/10.1103/PhysRevB.51.4606. [111] Song L, Ci L, Lu H, Sorokin PB, Jin C, Ni J, et al. Large scale growth and characterization of atomic hexagonal Boron nitride layers. Nano Lett 2010;10:3209–15. http://dx.doi.org/10.1021/nl1022139. [112] Han J, Lee J-Y, Kwon H, Yeo J-S. Synthesis of wafer-scale hexagonal boron nitride monolayers free of aminoborane nanoparticles by chemical vapor deposition. Nanotechnology 2014;25:145604. http://dx.doi.org/10.1088/0957-4484/25/14/145604. [113] Fazen PJ, Beck JS, Lynch AT, Remsen EE, Sneddon LG. Thermally induced borazine dehydropolymerization reactions. synthesis and ceramic conversion reactions of a new high-yield polymeric precursor to boron nitride. Chem Mater 1990;2:96–7. http://dx.doi.org/10.1021/ cm00008a004. [114] Paffett MT, Simonson RJ, Papin P, Paine RT. borazine adsorption and decomposition at Pt(111) and Ru(001) surfaces. Surf Sci 1990;232:286–96. http://dx.doi.org/10.1016/0039-6028(90)90121-N. [115] Fazen PJ, Remsen EE, Beck JS, Carroll PJ, McGhie AR, Sneddon LG. Synthesis, properties, and ceramic conversion reactions of polyborazylene. a high-yield polymeric precursor to Boron nitride. Chem Mater 1995;7:1942–56. http://dx.doi.org/10.1021/cm00058a028. [116] Kho J-G, Moon K-T, Nouet G, Ruterana P, Kim D-P. Boron-rich boron nitride (BN) films prepared by a single spin-coating process of a polymeric precursor. Thin Solid Films 2001;389:78–83. http:// dx.doi.org/10.1016/S0040-6090(01)00760-X. [117] Wolf G, Baumann J, Baitalow F, Hoffmann FP. Calorimetric process monitoring of thermal decomposition of B–N–H compounds. Thermochim Acta 2000;343:19–25. http://dx.doi.org/10. 1016/S0040-6031(99)00365-2. [118] Baitalow F, Baumann J, Wolf G, Jaenicke-Ro¨ßler K, Leitner G. Thermal decomposition of B–N–H compounds investigated by using combined thermoanalytical methods. Thermochim Acta 2002; 391:159–68. http://dx.doi.org/10.1016/S0040-6031(02)00173-9. [119] Baumann J, Baitalow F, Wolf G. Thermal decomposition of polymeric aminoborane (H2BNH2)x under hydrogen release. Thermochim Acta 2005;430:9–14. http://dx.doi.org/10.1016/j.tca.2004.12.002. [120] Bo¨ddeker KW, Shore SG, Bunting RK. Boron-Nitrogen chemistry. I. syntheses and properties of new cycloborazanes, (BH2NH2)n1,2. J Am Chem Soc 1966;88:4396–401. http://dx.doi.org/10.1021/ ja00971a017. [121] Zimmerman PM, Paul A, Zhang Z, Musgrave CB. Oligomerization and autocatalysis of NH2BH2 with ammoniaborane. Inorg Chem 2009;48:1069–81. http://dx.doi.org/10.1021/ic801731s. [122] Kim KK, Hsu A, Jia X, Kim SM, Shi Y, Hofmann M, et al. Synthesis of monolayer hexagonal Boron nitride on Cu foil using chemical vapor deposition. Nano Lett 2012;12:161–6. http://dx.doi.org/10. 1021/nl203249a. [123] Wang Y, Yamamoto Y, Kiyono H, Shimada S. Effect of ambient Gas and temperature on crystallization of boron nitride spheres prepared by vapor phase pyrolysis of ammonia borane. J Am Ceram Soc 2009;92:787–92. http://dx.doi.org/10.1111/j.1551-2916.2009.02941.x. [124] Berner S, Corso M, Widmer R, Groening O, Laskowski R, Blaha P, et al. Boron nitride nanomesh: functionality from a corrugated monolayer. Angew Chem Int Ed 2007;46:5115–9. http://dx.doi.org/ 10.1002/anie.200700234. [125] Gong Y, Shi G, Zhang Z, Zhou W, Jung J, Gao W, et al. Direct chemical conversion of graphene to boron- and nitrogen- and carbon-containing atomic layers. Nat Commun 2014;5. http://dx.doi. org/10.1038/ncomms4193.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 831

[126] Gorbachev RV, Riaz I, Nair RR, Jalil R, Britnell L, Belle BD, et al. Hunting for monolayer Boron nitride: optical and raman signatures. Small 2011;7:465–8. http://dx.doi.org/10.1002/smll. 201001628. [127] Preobrajenski AB, Vinogradov AS, Ma˚rtensson N. Monolayer of h-BN chemisorbed on Cu(111) and Ni(111): the role of the transition metal 3d states. Surf Sci 2005;582:21–30. http://dx.doi.org/10. 1016/j.susc.2005.02.047. [128] Nagashima A, Tejima N, Gamou Y, Kawai T, Oshima C. Electronic structure of monolayer hexagonal Boron nitride physisorbed on metal surfaces. Phys Rev Lett 1995;75:3918–21. http://dx.doi. org/10.1103/PhysRevLett.75.3918. [129] Paszkowicz W, Pelka JB, Knapp M, Szyszko T, Podsiadlo S. Lattice parameters and anisotropic thermal expansion of hexagonal boron nitride in the 10–297.5 K temperature range. Appl Phys A 2002;75:431–5. http://dx.doi.org/10.1007/s003390100999. [130] Burat O, Bouchier D, Stambouli V, Gautherin G. Characterization and growth mechanisms of boron nitride films synthesized by ion-beam-assisted deposition. J Appl Phys 1990;68:2780–90. http://dx.doi.org/10.1063/1.347173. [131] Novoselov KS, Jiang D, Schedin F, Booth TJ, Khotkevich VV, Morozov SV, et al. Two-dimensional atomic crystals. Proc Natl Acad Sci USA 2005;102:10451–3. [132] Kaasbjerg K, Thygesen KS, Jacobsen KW. Phonon-limited mobility in n-type single-layer MoS2 from first principles. Phys Rev B 2012;85:115317. [133] Ayari A, Cobas E, Ogundadegbe O, Fuhrer MS. Realization and electrical characterization of ultrathin crystals of layered transition-metal dichalcogenides. J Appl Phys 2007;101. 014507–014507. [134] Mak KF, Lee C, Hone J, Shan J, Heinz TF. Atomically thin MoS2: a new Direct-Gap semiconductor. Phys Rev Lett 2010;105:136805. [135] Radisavljevic B, Radenovic A, Brivio J, Giacometti V, Kis A. Single-layer MoS2 transistors. Nat Nano 2011;6:147–50. http://dx.doi.org/10.1038/nnano.2010.279. [136] Wang QH, Kalantar-Zadeh K, Kis A, Coleman JN, Strano MS. Electronics and optoelectronics of two-dimensional transition metal dichalcogenides. Nat Nano 2012;7:699–712. http://dx.doi.org/ 10.1038/nnano.2012.193. [137] Perkins FK, Friedman AL, Cobas E, Campbell PM, Jernigan GG, Jonker BT. Chemical vapor sensing with monolayer MoS2. Nano Lett 2013;13:668–73. http://dx.doi.org/10.1021/nl3043079. [138] CLAUSS FJ. Chapter 4-Molybdenum disulfide. In: CLAUSS FJ, editor. Solid Lubr. Self-Lubr. Solids. Academic Press; 1972. p. 75–112. [139] Chianelli RR, Siadati MH, De la Rosa MP, Berhault G, Wilcoxon JP, Bearden R, et al. Catalytic properties of single layers of transition metal sulfide catalytic materials. Catal Rev 2006;48:1–41. http://dx.doi.org/10.1080/01614940500439776. [140] Mak KF, He K, Shan J, Heinz TF. Control of valley polarization in monolayer MoS2 by optical helicity. Nat Nano 2012;7:494–8. http://dx.doi.org/10.1038/nnano.2012.96. [141] Chhowalla M, Shin HS, Eda G, Li L-J, Loh KP, Zhang H. The chemistry of two-dimensional layered transition metal dichalcogenide nanosheets. Nat Chem 2013;5:263–75. http://dx.doi.org/10.1038/ nchem.1589. [142] Wilson JA, Yoffe AD. The transition metal dichalcogenides discussion and interpretation of the observed optical, electrical and structural properties. Adv Phys 1969;18:193–335. http://dx.doi.org/ 10.1080/00018736900101307. [143] Yoffe AD. Layer compounds. Annu. Rev Mater Sci 1973;3:147–70. http://dx.doi.org/10.1146/ annurev.ms.03.080173.001051. [144] Doran NJ. Electronic structure and band theory of transition metal dichalcogenides. Physica BþC 1980;99:227–37.

832

HANDBOOK OF CRYSTAL GROWTH

[145] Bromley RA, Murray RB, Yoffe AD. The band structures of some transition metal dichalcogenides. III. Group VIA: trigonal prism materials. J Phys C Solid State Phys 1972;5:759. [146] Neto AC. Charge density wave, superconductivity, and anomalous metallic behavior in 2D transition metal dichalcogenides. Phys Rev Lett 2001;86:4382. [147] Zeng H, Dai J, Yao W, Xiao D, Cui X. Valley polarization in MoS2 monolayers by optical pumping. Nat Nano 2012;7:490–3. http://dx.doi.org/10.1038/nnano.2012.95. [148] Lee C, Yan H, Brus LE, Heinz TF, Hone J, Ryu S. Anomalous lattice vibrations of single- and fewlayer MoS2. ACS Nano 2010;4:2695–700. http://dx.doi.org/10.1021/nn1003937. [149] Najmaei S, Liu Z, Zhou W, Zou X, Shi G, Lei S, et al. Vapour phase growth and grain boundary structure of molybdenum disulphide atomic layers. Nat Mater 2013;12:754–9. [150] Zhan Y, Liu Z, Najmaei S, Ajayan PM, Lou J. Large-area vapor-phase growth and characterization of MoS2 atomic layers on a SiO2 substrate. Small 2012;8:966–71. http://dx.doi.org/10.1002/smll. 201102654. [151] Lin Y-C, Zhang W, Huang J-K, Liu K-K, Lee Y-H, Liang C-T, et al. Wafer-scale MoS2 thin layers prepared by MoO3 sulfurization. Nanoscale 2012;4:6637–41. http://dx.doi.org/10.1039/ C2NR31833D. [152] Lee Y, Lee J, Bark H, Oh I-K, Ryu GH, Lee Z, et al. Synthesis of wafer-scale uniform molybdenum disulfide films with control over the layer number using a gas phase sulfur precursor. Nanoscale 2014;6:2821–6. http://dx.doi.org/10.1039/C3NR05993F. [153] Lee K, Gatensby R, McEvoy N, Hallam T, Duesberg GS. High-performance sensors based on molybdenum disulfide thin films. Adv Mater 2013;25:6699–702. http://dx.doi.org/10.1002/adma. 201303230. [154] Prasad TP, Diemann E, Mu¨ller A. Thermal decomposition of (NH4)2MoO2S2, (NH4)2MoS4, (NH4)2WO2S2 and (NH4)2WS4. J Inorg Nucl Chem 1973;35:1895–904. http://dx.doi.org/10.1016/ 0022-1902(73)80124-1. [155] Liu K-K, Zhang W, Lee Y-H, Lin Y-C, Chang M-T, Su C-Y, et al. Growth of large-area and highly crystalline MoS2 thin layers on insulating substrates. Nano Lett 2012;12:1538–44. http://dx.doi.org/ 10.1021/nl2043612. [156] Shi Y, Zhou W, Lu A-Y, Fang W, Lee Y-H, Hsu AL, et al. Van der Waals epitaxy of MoS2 layers using graphene as growth templates. Nano Lett 2012;12:2784–91. [157] Mann J, Sun D, Ma Q, Chen J-R, Preciado E, Ohta T, et al. Facile growth of monolayer MoS2 film areas on SiO2. Eur Phys J B 2013;86:1–4. http://dx.doi.org/10.1140/epjb/e2013-31011-y. [158] Ji Q, Zhang Y, Gao T, Zhang Y, Ma D, Liu M, et al. Epitaxial monolayer MoS2 on mica with novel photoluminescence. Nano Lett 2013;13:3870–7. http://dx.doi.org/10.1021/nl401938t. [159] Wu W, De D, Chang S-C, Wang Y, Peng H, Bao J, et al. High mobility and high on/off ratio fieldeffect transistors based on chemical vapor deposited single-crystal MoS2 grains. Appl Phys Lett 2013;102. http://dx.doi.org/10.1063/1.4801861. [160] Ling X, Lee Y-H, Lin Y, Fang W, Yu L, Dresselhaus MS, et al. Role of the seeding promoter in MoS2 growth by chemical vapor deposition. Nano Lett 2014;14:464–72. http://dx.doi.org/10.1021/ nl4033704. [161] van der Zande AM, Huang PY, Chenet DA, Berkelbach TC, You Y, Lee G-H, et al. Grains and grain boundaries in highly crystalline monolayer molybdenum disulphide. Nat Mater 2013;12:554–61. [162] Wang X, Feng H, Wu Y, Jiao L. Controlled synthesis of highly crystalline MoS2 flakes by chemical vapor deposition. J Am Chem Soc 2013;135:5304–7. [163] Lee Y-H, Zhang X-Q, Zhang W, Chang M-T, Lin C-T, Chang K-D, et al. Synthesis of large-area MoS2 atomic layers with chemical vapor deposition. Adv Mater 2012;24:2320–5. http://dx.doi.org/10. 1002/adma.201104798.

Chapter 19 • Chemical Vapor Deposition of Two-Dimensional Crystals 833

[164] Zhu W, Low T, Lee Y-H, Wang H, Farmer DB, Kong J, et al. Electronic transport and device prospects of monolayer molybdenum disulphide grown by chemical vapour deposition. Nat Commun 2014;5. [165] Yu Y, Li C, Liu Y, Su L, Zhang Y, Cao L. Controlled scalable synthesis of uniform, high-quality monolayer and few-layer MoS2 films. Sci Rep 2013;3. [166] Wu S, Huang C, Aivazian G, Ross JS, Cobden DH, Xu X. Vapor–solid growth of high optical quality MoS2 monolayers with near-unity valley polarization. ACS Nano 2013;7:2768–72. http://dx.doi. org/10.1021/nn4002038. [167] Song I, Park C, Hong M, Baik J, Shin H-J, Choi HC. Patternable large-scale molybdenium disulfide atomic layers grown by gold-assisted chemical vapor deposition. Angew Chem Int Ed 2014;53: 1266–9. http://dx.doi.org/10.1002/anie.201309474. [168] Elı´as AL, Perea-Lo´pez N, Castro-Beltra´n A, Berkdemir A, Lv R, Feng S, et al. Controlled synthesis and transfer of large-area WS2 sheets: from single layer to few layers. ACS Nano 2013;7:5235–42. http://dx.doi.org/10.1021/nn400971k. [169] Song J-G, Park J, Lee W, Choi T, Jung H, Lee CW, et al. Layer-controlled, wafer-scale, and conformal synthesis of tungsten disulfide nanosheets using atomic layer deposition. ACS Nano 2013;7:11333–40. http://dx.doi.org/10.1021/nn405194e. [170] Gatensby R, McEvoy N, Lee K, Hallam T, Berner NC, Rezvani E, et al. Controlled synthesis of transition metal dichalcogenide thin films for electronic applications. Appl Surf Sci 2014;297: 139–46. http://dx.doi.org/10.1016/j.apsusc.2014.01.103. [171] Lee Y-H, Yu L, Wang H, Fang W, Ling X, Shi Y, et al. Synthesis and Transfer of single-layer transition metal disulfides on diverse surfaces. Nano Lett 2013;13:1852–7. http://dx.doi.org/10.1021/ nl400687n. [172] Zhang Y, Zhang Y, Ji Q, Ju J, Yuan H, Shi J, et al. Controlled growth of high-quality monolayer WS2 layers on sapphire and imaging its grain boundary. ACS Nano 2013;7:8963–71. http://dx.doi.org/ 10.1021/nn403454e. [173] Lu X, Utama MIB, Lin J, Gong X, Zhang J, Zhao Y, et al. Large-area synthesis of monolayer and few-layer MoSe2 films on SiO2 substrates. Nano Lett 2014;14:2419–25. http://dx.doi.org/10.1021/nl5000906. [174] Shaw J, Zhou H, Chen Y, Weiss N, Liu Y, Huang Y, et al. Chemical vapor deposition growth of monolayer MoSe2 nanosheets. Nano Res 2014;7:1–7. http://dx.doi.org/10.1007/s12274-014-0417-z. [175] Wang X, Gong Y, Shi G, Chow WL, Keyshar K, Ye G, et al. Chemical vapor deposition growth of crystalline monolayer MoSe2. ACS Nano 2014;8:5125–31. http://dx.doi.org/10.1021/nn501175k. [176] Huang J-K, Pu J, Hsu C-L, Chiu M-H, Juang Z-Y, Chang Y-H, et al. Large-area synthesis of highly crystalline WSe2 monolayers and device applications. ACS Nano 2013;8:923–30. http://dx.doi.org/ 10.1021/nn405719x. [177] Beechem TE, et al. ]. ACS Nano 2014;8:1655–63. [178] Chae SH, et al. Nat.Mater 2013;12:403–9. [179] Biedermann LB, et al. New J. Phys. 2010;12:125016. [180] Lock EH, et al. Nano Lett. 2012;12:102–7.