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Accepted Manuscript
Graphene-Analogous Low-Dimensional Materials
Qing Tang, Zhen Zhou
PII: S0079-6425(13)00037-6
DOI: http://dx.doi.org/10.1016/j.pmatsci.2013.04.003
Reference: JPMS 320
To appear in: Progress in Materials Science
Please cite this article as: Tang, Q., Zhou, Z., Graphene-Analogous Low-Dimensional Materials, Progress in
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Graphene-Analogous Low-Dimensional Materials
Qing Tang, Zhen Zhou*
Tianjin Key Laboratory of Metal and Molecule Based Material Chemistry, Key Laboratory of Advanced Energy
Materials Chemistry (Ministry of Education), Computational Centre for Molecular Science, Institute of New
Energy Material Chemistry, Nankai University, Tianjin 300071, P. R. China
ABSTRACT
Graphene, an atomic monolayer of carbon atoms in a honeycomb lattice realized in 2004, has
rapidly risen as the hottest star in materials science due to its exceptional properties. The explosive
studies on graphene have sparked new interests towards graphene-analogous materials. Now many
graphene-analogous materials have been fabricated from a large variety of layer and non-layer
materials. Also, many graphene-analogous materials have been designed from the computational
side. Though overshadowed by the rising graphene to some degree, graphene-analogous materials
have exceptional properties associated with low dimensionality and edge states, and bring new
breakthrough to nanomaterials science as well. In this review, we summarize the recent progress
on graphene-analogous low-dimensional materials (2D nanosheets and 1D nanoribbons) from both
experimental and computational side, and emphasis is placed on structure, properties, preparation,
and potential applications of graphene-analogous materials as well as the comparison with
graphene. The reviewed materials include strictly graphene-like planar materials (experimentally
available h-BN, silicene, and BC3 as well as computationally predicted SiC, SiC2, B, and B2C),
non-planar materials (metal dichalcogenides, metal oxides and hydroxides, graphitic-phase of ZnO,
MXene), metal coordination polymers, and organic covalent polymers. This comprehensive
review might provide a directional guide for the bright future of this emerging area.
* Corresponding author. Tel.: +86 22 23503623; fax: +86 22 23498941.
Email address: [email protected] (Z.Z.)
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Contents 1. Introduction...................................................................................................................................3 2. Experimental synthesis, characterization, and theoretical methods ..............................................6
2.1 Micromechanical cleavage .................................................................................................6 2.2 Chemical exfoliation ...........................................................................................................7 2.3 Chemical vapor deposition ...............................................................................................13 2.4 Surface-assisted epitaxial growth .....................................................................................18 2.5 Synthesis of 1D nanoribbons.............................................................................................20
2.5.1 Unzipping BN nanotubes to produce BN nanoribbons ..........................................20 2.5.2 Growth of ultranarrow MoS2 and WS2 nanoribbons inside carbon nanotubes .....22
2.6 Other synthetic routes .......................................................................................................23 2.7 Characterization ...............................................................................................................25 2.8 Theoretical methods ..........................................................................................................27
3. Planar graphene analogues..........................................................................................................30 3.1 “White graphene”: BN nanosheets and nanoribbons.......................................................31
3.1.1 Comparison between BN and C .............................................................................31 3.1.2 Electronic and magnetic properties of BN nanosheets and nanoribbons ..............33 3.1.3 Band-gap modifications of BN nanosheets and nanoribbons ................................36 3.1.4 BN/graphene hybrid structures ..............................................................................46 3.1.5 Potential applications ............................................................................................55
3.2 Silicene ..............................................................................................................................59 3.2.1 Synthesis of silicene and functionalized silicene ....................................................60 3.2.2 Theoretical investigations of silicene .....................................................................62 3.2.3 Potential applications of silicene ...........................................................................69
3.3 BC3 honeycomb sheets ......................................................................................................70 4. Hypothetical planar graphene analogues ....................................................................................73
4.1 SiC Silagraphene...............................................................................................................73 4.2 SiC2 Silagraphene .............................................................................................................76 4.3 Boron sheets ......................................................................................................................78 4.4 B2C sheet ...........................................................................................................................83
5. Non-planar materials...................................................................................................................85 5.1 Metal dichalcogenides ......................................................................................................85
5.1.1 MoS2 and WS2 ........................................................................................................85 5.1.2 Other layered metal dichalcogenides.....................................................................98
5.2 Layered oxide and hydroxide nanosheets .........................................................................99 5.3 Graphitic-like phase of ZnO............................................................................................104 5.4 MXenes............................................................................................................................108
6. Two-dimensional coordination and covalent organic polymers ...............................................112 6.1 Coordination polymers ...................................................................................................112 6.2 Covalent organic polymers .............................................................................................116
7. Conclusion and prospective ......................................................................................................119 Acknowledgements.......................................................................................................................121 References.....................................................................................................................................121
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1. Introduction
Graphene, a carbon honeycomb with only one-atom thickness isolated from
graphite in 2004, is the first example of true two-dimensional (2D) single-layer
atomic crystal [1]. Despite its short history, graphene shows tremendous attraction to
researchers from different fields and has risen as the most exciting star in materials
science during the past several years [ 2 ]. Its exceptional properties, such as
half-integer quantum Hall effect, ambipolar electric field effect, extremely high
carrier mobility, high thermal conductivity, high specific surface area, and the highest
strength ever mearsured, provide a fertile ground for the possible implementation of
graphene in nanodevices for a large variety of applications, and a lot of recent reviews
have been directed towards its synthesis, properties, and functionalized applications
[3-6].
As a matter of fact, the intensive studies on graphene have sparked new
discoveries towards graphene-analogous materials, comprising single layer or few
layers with compositions other than carbon [7-9]. Initially, following the same
micromechanical exfoliation methodology originally used in graphite, Novoselov et al.
[10] successfully isolated individual single layers from a large variety of layered
materials including h-BN, dichalcogenides (such as MoS2 and NbSe2) and complex
oxides (such as Ba2Sr2CaCu2Ox). Since then, many 2D materials including BC3 [11],
silicene [12], MXene (transition metal carbides and nitrides, such as Ti3C2 [13] and
Ti2C [14]), and coordination polymers (such as [Cu2Br(IN)2]n [15] and polymeric
Fe-phthalocyanine [16]), have been prepared and characterized, and new fabrication
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methods, including liquid-phase exfoliation [17], electrochemical exfoliation [18],
chemical vapor deposition (CVD), and polycondensation reaction, have been
accordingly established. In addition, unique experimental skills were developed for
preparing 1D ribbons, e.g., BN nanoribbons (NRs) were produced by unzipping BN
nanotubes through plasma etching [19] or potassium intercalation [20]. Ultra-narrow
MoS2 and WS2 nanoribbons encapsulated into single- and double-walled carbon
nanotubes were synthesized via chemical reactions inside carbon nanotubes [21,22].
These novel materials have fantastic properties and promising applications. For
example, BN nanosheets are highly insulating, and exhibit superb chemical, thermal,
and oxidation stability, and possess a mechanical strength and thermal conductivity
comparable to those of graphene. These excellent properties aid the uses of layered
BN as thermal radiators, ultraviolet-light laser and emitter devices, as
thermoconductive fillers in polymer or ceramic composites, and as dielectric substrate
for graphene-based electronics [23]. Similarly, MoS2 nanosheets were explored as
high-temperature solid lubricants, nanoelectronics, electrode materials for Li-ion
batteries, and catalysts. MoS2 monolayer with a direct band gap of 1.8 eV serves as a
semiconducting rival of graphene, particularly considering that the zero-band-gap
character of the latter poses a severe drawback for its applicability in logic devices.
Actually, a MoS2-based transistor with a room-temperature mobility of more than 200
cm2/(Vs) and current on/off ratios of up to 1×108 has been fabricated, which is
demonstrated to be comparable to thin Si films or graphene nanoribbons (GNRs).
At the computational side, many graphene-like materials have been explored and
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designed, and fascinating properties distinctive from graphene are predicted even in
advance of experiments. For example, in contrast to graphene, all BN monolayers
have closed-shell singlet ground states, and those with long zigzag edges have slightly
larger band gaps [24]. For BN nanoribbons, the electronic properties depend heavily
on the edge states. The ground states of both fully bare BN nanoribbons and the ones
with a bare N edge and an H-terminated B edge are half metallic. The alignment of
spin at the bare B edge is antiferromagnetic (anti-FM), while that at the bare N edge is
ferromagnetic (FM) for both the fully bare and half-bare zigzag-edged BN
nanoribbons [25]. Hydrogenation can further precisely modulate the electronic and
magnetic properties of BN nanoribbons by controlling hydrogenation ratios [26].
Zigzag SiC nanoribbons are magnetic metals, whose spin polarization originates from
the unpaired electrons localized on the ribbon edges. Interestingly, the zigzag SiC
nanoribbons narrower than ~4 nm present half-metallic behavior without the aid of
external field or chemical modification [27]. Ahead of experimental realization, Li et
al. [28] investigated the stability, magnetic and electronic properties of MoS2 sinlge
layer and nanoribbons with either zigzag or armchair-terminated edges. Zigzag
nanoribbons show ferromagnetic and metallic behavior, while armchair nanoribbons
are nonmagnetic (NM) and semiconducting. Such materials came into truth in 2010
[21]. By computations, planar tetracoordinate silicon (ptSi)-containing SiC2
silagraphene was designed [29], where each silicon atom is bonded with four carbon
atoms in a pure plane, representing the first anti-van't Hoff/Lebel species in the
Si-containing extended system. All these ptSi-containing nanomaterials are metallic.
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Though overshadowed by the rising graphene to some degree,
graphene-analogous materials have exceptional properties, and bring new
breakthrough to nanomaterials science. Very recently, several good reviews have been
published on graphene-like materials, including BN, transition metal dichalcogenides,
and silicene [12,23,30]. Aside from these, many other 2D materials have also gained
recent excitement, such as Si-C binary phase, B and B-C phase, layered V2O5,
MXenes, wurtzite materials (such as ZnO), and coordination or covalent polymers. In
this review, we will summarize recent progress in all these graphene-analogous
materials from both experimental and computational side, and emphasis is placed on
structure, properties, preparation, and potential applications. It is highly expected that
this more comprehensive review might be useful for both experimental and theoretical
peers and provide a directional guide for the bright future of this emerging new area
of research.
2. Experimental synthesis, characterization, and theoretical methods
In this section we will first highlight the recent progress in general synthesis and
characterization of experimentally achieved 2D materials, and some special methods
for individual materials will be introduced in their corresponding section.
Subsequently, we will provide a brief summary on theoretical methods used in
structure modeling and property prediction of graphene-analogous materials.
2.1 Micromechanical cleavage
Micromechanical cleavage, as originally used in peeling off graphene from
graphite, can be extended to other layered materials with weak van der Waals (vdW)
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forces between layers. This method requires repeatedly peeling layered materials and
followed transferring the peeled sample on top of a surface. The micromechanical
cleavage was firstly applied to isolation of h-BN, MoS2, NbSe2, and Ba2Sr2CaCu2Ox
[10]. The resulting 2D sheets are stable under ambient conditions, exhibit high crystal
quality, and are continuous on a macroscopic scale. After that, several groups have
reported synthesis of BN [31,32], MoS2, NbSe2, WSe2 [33], GaS, GaSe [34,35],
Bi2Se3 [36], and Bi2Te3 [37] nanosheets from their layered phases by using this
method, with the obtained thickness ranging from one to ten atomic layers.
Micromechanical cleavage has proven an easy and fast way of obtaining highly
crystalline atomically thin nanosheets. Nevertheless, this method also produces a large
quantity of thicker sheets, and the thinner or monolayer ones only reside in a very
minor proportion; thus this method is not scalable to mass production for potential
engineering applications. Consequently, although mechanical cleavage has shown
some success in isolating thin BN and MoS2 nanosheets, this method is practically
less explored due to its extremely low yield.
2.2 Chemical exfoliation
The critical drawback for micromechanical exfoliation is the large thickness of
obtained majority products. As an alternative method, chemically derived exfoliations,
such as liquid-phase exfoliation, and ion-intercalation induced exfoliation, have been
demonstrated to effectively isolate single layer and few layers from those thicker
structures in large quantities. These chemically derived routes, also regarded as wet
methods, have been widely adopted for synthesizing 2D materials with lateral sizes up
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to several micrometers.
Liquid-phase exfoliation is regarded as a dispersion/exfoliation method, which
consists of sonicating the layered bulk materials in polar solvents, surfactant or
reaction reagents, and then exfoliating the resultant dispersions into separated thin
layers with assistance of centrifugation. The strong affinity between solvent and host
materials weakens the interlayer interactions and thus facilitates the isolation of thin
sheets upon sonication. A proper solvent should have a surface energy that matches
the energy required to overcome the vdW forces of bulk materials, and be able to
form stable dispersion with the host materials against reaggregation. The quality of
yielded materials depends heavily on solution-processing parameters such as
sonication time and centrifugation rate. Synthesis of mono- and few-layered h-BN
nanosheets from single crystalline h-BN via a chemical solution derived method was
first accomplished by Han et al. [38] in 2008. Later, Zhi et al. [39] exfoliated
large-scale BN nanosheets from BN powers dispersed in a strong polar solvent,
N,N-dimethylformamide (DMF). However, the isolated BN nanosheets suffer from
relatively low yields, corresponding to 0.01-0.03 mg/mL after 10 h sonication, which
is mainly due to the weak interactions between BN layers and the solvent molecules.
Using the liquid-phase exfoliation method, Coleman et al. [17] produced single-
and few-layered 2D nanosheets (BN, MoS2, and WS2) dispersed in various organic
solvents (Fig. 1), which were further fabricated into macroscopic films by vacuum
filtration. In their experiments, about 25-30 solvents were examined to check their
effectiveness on each material, and the promising solvents should have a surface
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tension close to 40 mJ/m2, such as N-methyl-pyrrolidone (NMP) and isopropanol
(IPA). This provides useful information for exploring new solvents and solvent blends.
Under their optimized solvent conditions, the lateral sizes of the resulting nanosheets
are 50 to 1000 nm for MoS2 and WS2 and 100 to 5000 nm for BN, with the generated
concentration as high as 0.3 mg/ml for MoS2 (in NMP), 0.15 mg/ml for WS2 (in
NMP), and 0.06 mg/ml for BN (IPA). Transmission electron microscopy (TEM)
images confirmed that the exfoliated nanosheets exhibited well-preserved hexagonal
symmetry.
Fig. 1. (a) Photographs of MoS2 (in NMP), WS2 (in NMP), and BN (in IPA). (b)-(d)
High-resolution TEM (HRTEM) images of BN, MoS2, and WS2 monolayers. (e)
Photograph of BN, MoS2, and WS2 (thickness ~50 µm) films fabricated via vacuum
filtration. Reproduced from Ref. [17]. Copyright © 2011, American Association for
the Advancement of Science.
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Additionally, other polar solvents, such as deionized water [ 40 ],
N-methyl-2-pyrrolidone [41], methanesulfonic acid (MSA) [42], 1,2-dichloroethane
[43], cyclohexylpyrrolidone (CHP) [44], a mixture of ethanol and water [45], sodium
cholate/water solution [46], and formamide [47,48], have been used as dispersion
media of layered materials (Table 1).
Table 1
Overview of adopted solvents and reagents for liquid-phase exfoliation of 2D
materials.
Solvents/Reagents Exfoliated materials (precursor) Ref.
N,N-dimethylformamide (DMF) BN [39]
N-methyl-2-pyrrolidone BN [41]
methanesulfonic acid (MSA) BN [42]
1,2-dichloroethane BN [43]
1,2-dichloroethane/polyphenylenevinylene BN [38]
molten hydroxides BN [49]
cyclohexylpyrrolidone (CHP) WS2, MoS2, MoSe2, MoTe2 [44]
N-methyl-pyrrolidone (NMP) MoS2, WS2, WSe2 [17,50,51]
isopropanol (IPA) BN [17]
deionized water BN [40]
N-vinyl-2-pyrrolidone MoS2 [52]
ethanol/water MoS2, WS2, BN [45]
sodium cholate solution MoS2 [46]
formamide VS2, V2O5 [47,48]
octadecylamine (ODA) BN [53]
amine-terminated polyethylene glycol BN [53]
N2H4, H2O2, HNO3/H2SO4, oleum BN [55]
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Besides the use of polar solvents, liquid-phase exfoliation based on chemical
reactions is another facile route to isolate layered materials. This route, as a matter of
fact, is a combination of, firstly, a chemical modification of a layered material, and
later an ultrasonic treatment in organic solvent. As a first attempt, Lin et al. [53,54]
reported the covalent functionalization of h-BN by amine molecules through Lewis
acid-base reaction, in which the amine species forms covalent complexes with the
electron-deficient B atoms in h-BN surface. This complexation facilitates the
intercalation of functional molecules into the h-BN layered structures and triggers the
exfoliation of functionalized h-BN into solubilized thin nanosheets (3-20 layers) as a
result. In a later report, Nazarov et al. [55] reported the BN functionalizaiton with
inorganic reagents by heating h-BN power with hydrazine (N2H4), H2O2,
HNO3/H2SO4, and oleum. Further ultra-sonication in water or DMF leads to
exfoliation of the functionalized h-BN to form stable colloidal solutions of few-layer
BN nanosheets with lateral dimensions below 1 µm. Very recently, Sainsbury et al.
[ 56 ] have reported the surface covalent functionalization of exfoliated h-BN
nanosheets via nitriene addition reaction, and the generated product demonstrates
enhanced dispersability in organic solvents and facilitates the chemical
compatibilization of h-BN nanosheets within polymer matrices. Besides BN, this
method has also been extensively utilized to realize nanosheets of metal oxides and
hydroxides [57,58].
Another strategy used for exfoliation of layered materials is the ion-intercalation
method, which has been used a lot for exfoliating transition-metal dichalcogenides.
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This method usually includes two basic steps: Li intercalation into the layered space
of bulk materials, and the subsequent exfoliation by immersing the Li-intercalated
compounds in water by ultrasonication. The reaction between the incorporated Li and
water forms LiOH and H2, which expands the interlayer space to facilitate the
exfoliation. At early attempts, single-layer MoS2, WS2, and their related
dichalcogenides were prepared by using butyl lithium as the intercalation reagent [59].
However, the major difficulty in this standard intercalation technique arises from the
low Li concentration, and more rigid conditions, such as elevated reaction
temperature (~100oC) and longer reaction time (~3 days), are generally required to
increase the Li content.
Zeng et al. [18] demonstrated a facile route for high-yield production of
single-layer dichalcogenides (MoS2, WS2, TiS2, TaS2, and ZrS2) through a
controllable electrochemical lithiation process (Fig. 2), which can be easily performed
within 6 hours at room temperature. The electrochemical lithiation was performed in a
battery test system, in which the layered materials were prepared as the cathode, and
the lithium foil was used as the anode. This method has the advantage that the degree
of lithium intercalation is effectively monitored by the discharge process. Subsequent
ultrasonication in water or ethanol yields large-amount single-layer nanosheets, and
the resulting MoS2 monolayers even achieve a high yield of 92%. The developed
electrochemical lithiation can be applied to prepare other 2D materials, like h-BN and
metal selenides or tellurides (such as NbSe2, WSe2, Sb2Se3, and Bi2Te3) [60].
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Fig. 2. Scheme of electrochemical lithiation process. Reproduced from Ref. [18].
Copyright © 2011, Wiley-VCH.
A noteworthy drawback for the Li-intercalation method is that Li intercalation
results in a striking structural phase transformation [61]. Taken MoS2 as an example,
the Mo coordination changes from trigonal prismatic (2H) to octahedral (1T) after Li
insertion. The phase transformation process is detrimental since it leads to formation
of metastable MoS2 structures, and degrade their electronic behaviors and device
functionalizations.
The chemically derived solution processing, relying upon strong polar solvents,
reactive reagents, or ion intercalation, are versatile and up-scalable. The fabricated 2D
materials can be redispersed in common organic solvents, and can be used to deposit
in various environments and substrates not available for mechanical cleavage methods.
This method opens up a whole new range of potential large-scale preparations and
applications of 2D materials for nanodevices, composites, or liquid phase chemistry.
2.3 Chemical vapor deposition
In close resemblance to growth of graphene by thermal decomposition of
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hydrocarbons on substrates, production of other 2D materials can also be
accomplished by high-temperature chemical reactions of molecular precursors on a
surface. This bottom-up self-assembly method is also known as chemical vapor
deposition (CVD). The surface not only serves as a template but also plays the role of
a catalyst to assist the epitaxial growth of solid films. The advantageous aspect for
this surface-assisted method is that clean single layer or few layers can grow without
accompanying thicker flakes.
The application of CVD technique in generating one-atom-thick BN films was
traced back to almost a decade ago. Generally, well-ordered single and double layer
BN domains were synthesized by thermal decomposition of borazine (BN)3H6 or
B-trichloroborazine (ClBNH)3 on catalytic metal surfaces such as Ni, Pd, Pt, Cu, Ag,
Fe, Ir, and Rh [62-64]. During these CVD processes, the rigid growth conditions, such
as high temperature (> 700oC) and ultrahigh vacuum chambers, are required. For the
deposited h-BN, B atom favors to site near a face-centered cubic (fcc) or hexagonal
close-packed (hcp) hollow site, and N atom prefers to locate close to the on-top
position of surface metal atoms [65].
These early studies on CVD growth only demonstrated the formation of BN
nanomeshs with relatively small areas, and the thickness, in most cases, is limited to
one layer. Methods for preparing h-BN layers with large areas and varying thickness
have been proposed in later reports. Recently, through atmospheric pressure (AP) or
low pressure (LP) thermal catalytic CVD methods, polycrystalline Ni film and Cu foil
have proved excellent catalytic substrates for growth of large-area thin BN nanosheets
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(one to five layers) with borazine [ 66 ], amineborane [ 67 - 70 ], or decaborane
(diborane)/ammonia mixture [71,72] as the precursor. Besides metal substrates, BN
nanosheets with thickness lower than 5 nm can be directly fabricated on Si/SiO2
substrates by thermal CVD using B, MgO, and FeO powers as precursors under NH3
gas flow [73], or by catalyst-free microwave plasma (MP) reaction of BF3-H2-N2
mixtures [74]. In further advance, Liu et al. [75] reported the use of graphene as a
supported substrate for depositing BN nanosheets via a two-step CVD process (Fig. 3).
During the first step, graphene is deposited onto Cu foil by the chemical
decomposition of n-hexane at 950 C. During the second step, h-BN layers form on the
generated graphene/Cu foils through thermal decomposition of ammonia borane
(NH3-BH3) at 1000 C. The CVD-produced BN nanosheets can also be realized by
using other precursors (see Table 2). Besides BN, growth by CVD is a promising
approach to produce B-C-N hybrid materials [76].
Fig. 3. Schematic for preparation of graphene/BN stacked film. (a) Cu foil substrate.
(b) Graphene is grown on Cu. (c) Graphene/Cu foil is loaded for growth of h-BN film
on top. (d) and (e) are photographs of graphene (purple), graphene/h-BN film (blue),
and SiO2 (light purple). Reproduced from Ref. [75]. Copyright © 2011, American
Chemical Society.
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Moreover, CVD methods have been used to grow atomically thin sheets of MoS2
on insulating substrates. For example, Zhan et al. [77] synthesized large-area MoS2
atomic layers (one to three layers) on SiO2/Si substrates by CVD using Mo and S as
reagents. In this procedure, the Mo thin films are firstly deposited onto SiO2
substrates by e-beam evaporation, and the pre-deposited Mo layers then react with the
introduced S vapors at 750 C to form MoS2 sheets (Fig. 4). As another alternative
routine, Lee et al. [78] fabricated MoS2 thin films directly deposited on SiO2/Si
substrates by CVD from MoO3 and S precursors under 650 C. Besides, CVD growth
of MoS2 was realized by using CVD-grown graphene on Cu foil as the template,
which gives rise to single crystalline hexagonal MoS2 flakes with a lateral size of
several micrometers [79].
Fig. 4. Schematic for growth of MoS2 on SiO2 substrate. Reproduced from Ref. [77].
Copyright © 2012, Wiley-VCH.
As disclosed by the above examples, surface-assisted CVD growth of thin layers
allows the patterning and large-scale growth of 2D materials, and the yielded
structures can be free of thicker nanosheets. However, this method relies upon the
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interaction between substrate surface and the deposited films, and a very harsh
environment, such as high reaction temperature, is often required, which always poses
significant experimental difficulty. The representative examples are summarized in
Table 2 for 2D materials achieved by CVD techniques. Compared with the simple and
easily-handled chemical exfoliation methods, the CVD strategy is highly costly and
thus might limit its wide applications.
Table 2
Representative 2D materials produced by CVD method.
2D materials Precursor Method T ( C) Ref.
BN borazine APCVD 400 (1000) [66]
BN decaborane/ammonia thermal CVD 1000 [71]
BN diborane/ammonia LPCVD 1025 [72]
BN B, MgO, FeO, NH3 thermal CVD 900-1200 [73]
BN BF3/H2/N2 MPCVD 800 [74]
BN BCl3/NH3/H2/N2 thermal CVD 1000 [80]
BN boron oxide/melamine thermal CVD 1100-1300 [81]
BN (MeO)3B oxidation-nitrification
CVD 627 [82]
hybrid h-BNC methane/amineborane thermal CVD 900-1000 [76]
MoS2 Mo/S vapor thermal CVD 750 [77]
MoS2 MoO3/S thermal CVD 650 [78]
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2.4 Surface-assisted epitaxial growth
Surface-assisted epitaxial growth can be regarded as a modification of CVD
method, in which the substrate surface serves as a seed crystal other than a template or
a catalyst. This method is alternatively considered as molecular beam epitaxy (MBE)
growth. The epitaxial growth has been successfully applied to fabricate
one-atom-thick Si sheets (silicene), and the un-reactive metal Ag with six-fold surface
symmetry provides a promising substrate to facilitate growth of hexagonal silicene or
Si nanoribbons [83].
On Ag(110) surface, in situ deposition of Si sources under untrahigh vaccum
conditions produces self-aligned Si NRs with a honeycomb graphene-like structure
[84-86]. For example, De Padova et al. [85] produced high aspect ratio Si NRs on
Ag(110) substrate, with several nanometers in lengths, 1.6 nm in width and only 0.2
nm in height. In a later work, De Padova et al. [86] grew perfectly straight and aligned
multilayer Si NRs with pyramidal cross section.
On Ag(111) surface, Lalmi et al. [87] synthesized a highly ordered Si monolayer
with a ( 3232 × )R30 superstructure. However, the Si-Si distance (0.19 ± 0.01 nm)
determined by Lalmi et al. is much shorter than the expected value of 0.21 nm as
computed from the ( 3232 × )R30 model. This makes the structure of silicene on
Ag(111) speculative. Lay et al. [88] have recently showed that the discrepancy
concerning the Si-Si separation in Lalmi et al.’s scanning tunneling microscopy (STM)
results can be explained by their misinterpretation of pure Ag(111) surface as the
strained silicene layer.
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Fig. 5. (a) STM image of 2D Si layer on Ag(111)-(1×1). (b) Density functional theory
(DFT) results for silicene on Ag(111): top and side view of the fully relaxed atomic
geometries of the model for silicene on Ag(111) surface. Reproduced from Ref. [89].
Copyright © 2012, American Physical Society.
After that, five other groups independently reported experimental evidences of
ordered silicene phases on the same surface. For example, Vogt et al. [89] reported
epitaxial growth of silicene on Ag(111) substrate by directly depositing evaporated Si
onto Ag samples at temperatures between 220-260 C. The STM topograph (Fig. 5a)
shows a Si adlayer covering Ag(111) surface terraces with a honeycomb-like
appearance. Structurally, the Si atoms are located either above Ag atoms or between
Ag atoms, and the silicene layer forms a uniform (4×4) symmetry with respect to the
unrestructured (1×1) Ag(111) surface. At variance with the planar skeleton of
graphene, silicene prefers to be structurally low-corrugated. The in-plane Si-Si
distance determined from the STM image is about 0.22 nm (±0.01 nm), in agreement
with theoretical prediction of 0.232 nm (Fig. 5b). Similarly, Chiappe et al. [90], Feng
et al. [91], Lin et al. [92], and Jamgotchian et al. [93] also produced silicene
monolayer on Ag(111) substrate. According to their results, the silicene sheets are
rather sensitive to deposition temperature, and the resulting silicene domains present
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different orientations ((4×4), ( 1313 × ), ( 77 × ), ( 3232 × )) with respect to the
(1×1) Ag substrate on varying temperatures.
Besides the Ag substrate, silicene can also be fabricated on Ir(111) surface [94].
Moreover, Fleurence et al. [95] showed that silicene can form through surface
segregation on ZrB2 (0001) thin films grown on Si(111) wafers.
2.5 Synthesis of 1D nanoribbons
1D nanoribbons have properties distinctive from their 2D nanosheets due to edge
states. Growth of high-quality nanoribbons with controllable widths and smooth edges
is important for many technological applications. However, compared with numerous
routes established for fabricating graphene nanoribbons, the experimental
availabilities for facile synthesis of other nanoribbons are comparatively rare. Besides
the availability of silicene nanoribbons, BN, MoS2 and WS2 nanoribbons have been
successfully fabricated by unique experimental techniques.
2.5.1 Unzipping BN nanotubes to produce BN nanoribbons
In analogy to the successful unzipping of carbon nanotubes to fabricate graphene
nanoribbons, BN nanoribbons have been successfully produced by unzipping of BN
nanotubes. Similarly, multiwalled BN nanotubes were unzipped to form nanoribbons
by selective plasma etching or by potassium vapor intercalation.
Initially, Zeng et al. [19] fabricated BN nanoribbons based on Ar plasma etching
of BN nanotubes embedded within a polymer (Fig. 6a). In this process, since side and
bottom walls of the nanotubes were initially imbedded in the poly(methyl
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methacrylate) (PMMA) film, only the top part of the tube-shell was exclusively
etched and gradually removed, resulting in nanotube unzipping through cutting top
walls. The produced BN nanoribbons mainly displayed N-terminated zigzag edges
and surface vacancy defects, with the ribbon’s widths as narrow as 15 nm and the
lengths ranging from several hundred nanometers to several microns.
Erickson et al. [20] developed another synthetic method to produce BN
nanoribbons by the K-intercalation-induced longitudinal splitting of BN nanotubes
(Fig. 6b). Mechanically, K intercalation between the walls induces bond strains
circumferentially around the sidewalls of the tubes, which facilitates bond breakage
along the longitudinal direction of outer walls. This approach results in formation of
narrow (20-50 nm), long (at least 1 µm in length), thin (usually between 2 and 10
layers) and high crystalline BN nanoribbons.
Very recently, Li et al.[96] have reported an in situ unzipping of BN nanotubes
to produce BN nanoribbons, where the unzipping arises from intercalation of BN
nanotubes by Li-NH3 species formed during nanotube synthesis.
Fig. 6. (a) Schematic of the unwrapping process of BN nanotubes induced by plasma
etching and atomic force microscopy (AFM) image of single-layered BN nanoribbons.
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Reproduced from Ref. [19]. Copyright © 2010, American Chemical Society. (b)
Schematic of the potassium-intercalation-induced splitting process of BN nanotubes
and TEM micrograph of a few-layer BN nanoribbon splitting off its parent ribbons.
Reproduced from Ref. [20]. Copyright © 2011, American Chemical Society.
2.5.2 Growth of ultranarrow MoS2 and WS2 nanoribbons inside carbon nanotubes
One method currently developed for producing ultra narrow MoS2 and WS2
nanoribbons is the direct chemical growth inside carbon nanotubes (Fig. 7a). This
method was developed by Wang et al. [21], and the carbon nanotubes behave as the
role of templates that confine the growth of nanoribbons along 1D direction. In this
synthetic method, H3PMo12O40 was firstly encapsulated into carbon nanotubes, and
the filled carbon nanotubes were then heated under H2S/H2 to generate MoS2.
Depending on the diameter of carbon nanotubes, single-layered and bi- or tri-layered
MoS2 nanoribbons can be obtained. Fig. 7b shows the HRTEM image of a
single-layer MoS2 nanoribbon encapsulated in a single-walled carbon nanotube. The
MoS2 nanoribbons clearly exhibit a zigzag-typed edge with four zigzag chains along
the ribbon width (Fig. 7c is the schematic model). Through similar procedure, ultra
narrow WS2 nanoribbons with smooth zigzag edges and 1-3 layered thickness were
fabricated inside carbon nanotubes using H3PW12O40 as the starting material [22].
This method can be potentially applied to yield other transition-metal chalcogenide
nanoribbons.
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Fig. 7. (a) Model of MoS2 nanoribbon inside carbon nanotube. (b) HRTEM image of
a single-layer MoS2 nanoribbon encapsulated in a single-walled carbon nanotube. (c)
Model of the MoS2 nanoribbon in (b). Reproduced from Ref. [21]. Copyright © 2010,
American Chemical Society.
2.6 Other synthetic routes
The above reviewed methods are the typically adopted methods for synthesizing
2D materials, especially for BN and dichalcogenides. Besides, many other routes have
been developed for fabricating thin sheets of BN, metal dichalcogenides and other 2D
materials (see Table 3). For example, BN nanosheets can be realized by ex-situ ion
etching [97], surface segregation on catalytic metals [98], and high-energy electron
beam irradiation inside TEM [99,100]. Wang et al. introduced a “chemical blowing”
method to synthesize BN and C-containing BN nanosheets (one to few layers) in mass
production [101]. This method relies on making large bubbles with thin B–N–H (or
B–N–C–H) polymer walls by releasing H2 from a precursor ammonia borane, in
which the produced H2 bubbles promote thinning process of BN nanosheets.
Moreover, MoS2 nanosheets can be alternatively achieved by laser-thinning technique
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[ 102 ], thermolysis of ammonium thiomolybdate ((NH4)2MoS4) [ 103 ], thermal
evaporation technique [104], and hydrothermal synthesis [105], etc. Thin nanosheets
of other metal dichalcogenides/chalcogenides, such as VS2 [106], MoSe2, WSe2 [51],
TiS2 [107], Bi2Se3 [108], and Bi2Te3 [109], have also been fabricated through various
experimental skills.
Table 3
Other synthetic routes for synthesizing BN and metal dichalcogenides or
chalcogenides.
2D materials Synthesis method Ref.
BN ex-situ ion etching [97]
BN surface segregation from bulk Fe-Cr-Ni alloy doped with B and N [98]
BN chemical reaction of boric acid with urea at 900oC [110]
BCN chemical reaction of activated charcoal with boric acid and urea at
900oC [111]
BN high-energy electron beam irradiation [99,100]
BN “chemical blowing” method using ammonia borane as precursor [101,112]
MoS2 laser-thinning of multilayered MoS2 flakes [102]
MoS2 two-step thermolysis of (NH4)2MoS4 precursor [103]
MoS2 thermal evaporation of MoO3 with S powers [104]
MoS2 annealing of ball-milled MoO3 and S powers at 600oC [113]
MoS2 solvothermal synthesis from MoO3 and thioacetamide [114]
WS2 sulfidation of W18O49 nanorods [115]
MoS2 hydrothermal reaction between MoO3 and KSCN [105]
MoSe2 hydrothermal reaction between molybdic acid and Se [51]
MoS2, WS2 chemical reaction of molybdic (tungstic) acid and thiourea at 773K [105]
MoSe2, WSe2 chemical reaction of molybdic (tungstic) acid and Se at 773K [51]
VS2 chemical exfoliation of VS2
.NH3 precursor formed by hydrothermal reaction of NH4VO3 and thioacetamide
[106]
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TiS2 chemical reaction of TiCl4 and oleylamine/S mixtures at 300oC [107]
Bi2Se3 molecular beam epitaxy (MBE) growth [108]
Bi2Te3 surface-assisted chemical vapor transport (CVT) technique [109]
In addition, the unique fabrications for graphitic-like ZnO and MXene layers, as
well as coordination synthesis for metal organic frameworks (MOFs) and
polymerization reaction for assembling covalent organic frameworks (COFs) will be
reviewed in their corresponding sections later.
2.7 Characterization
The layer thickness of 2D materials is determined mainly by optical microscopy
imaging, AFM, and Raman spectroscopy [116]. A facile observation of the obtained
2D sheets can be realized by optical microscopy, and further combination with AFM
technique offers a fast estimation of the thickness distribution. Additionally, since
vibrational spectrum shows significant thickness dependence, Raman spectroscopy is
also widely used to determine the thickness and to examine the changes in material
properties with thickness (Fig. 8). For example, the AFM images in Fig. 8a and 8d
illustrate the identification of regions with different thicknesses (one to four layers)
for extracted BN and MoS2 nanosheets deposited on oxidized Si wafer. With regard to
Raman spectroscopy, the Raman intensity and the peak shift were employed for
characterizing layer numbers. The characteristic BN peak centered at ~1366 cm-1 in
bulk h-BN originates from the E2g mode of B-N vibration within h-BN layers, which
would shift to higher or lower frequency under compressive or tensile stress. The
monolayer h-BN exhibits a blue shift of up to 4 cm-1, which is explained by hardening
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of the E2g mode because of the slightly shorter B-N bonds in BN monolayer. In
bilayer h-BN, the strain effect becomes dominated, and causes the red shift by 1-2
cm-1. The peak becomes stronger as the thickness increases, ascribed to a low yield of
monolayer BN compared with thicker ones (Fig. 8b). Similarly, for the exfoliated
MoS2 layers, the in-plane E2g1 and the out-of-plane A1g vibration modes near 400 cm-1
are sensitive to film thickness; the former shows red shift and the latter blue shift as
thickness increases, and finally the Raman frequencies converge to bulk values when
the films are thicker than four layers (Fig. 8f).
Fig. 8. (a) Optical microscope of thin BN nanosheets. The insets show AFM images.
(b) Raman spectroscopy of thin BN nanosheets. Reproduced from Ref. [32].
Copyright © 2011, Wiley-VCH. (c) Optical micrograph of thin MoS2 films. (d) AFM
image of MoS2. (e) Model of MoS2 bilayer structure. (f) Raman spectra of thin and
bulk MoS2 films. (g) Atomic displacements of bulk MoS2 crystal. Reproduced from
Ref. [33]. Copyright © 2010, American Chemical Society.
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2.8 Theoretical methods
The advancement of parallel computational powers, especially development of
supercomputer facilities in recent decades, has rapidly paced the theoretical
investigations on various materials ranging from bulk to low-dimensional
nanostructures. The significance of theoretical studies is certainly self-evident, which
can not only provide scientific understanding to elucidate experimentally observed
phenomena, but also offer an important tool for conducting materials design and
property prediction. The fundamental theory for materials modeling and computations
at the atomic scale is based on ab initio or quantum mechanics theory.
Density functional theory (DFT) is currently the most commonly employed
quantum mechanics method, which has evolved into a powerful tool for computing
electronic ground-state properties of a large number of nanomaterials. The entire field
of DFT method relies on the theorem that the ground-state energy of a many-electron
system is a unique and variational functional of the electron density, and this
conceptual proposal is implemented in a mathematical form to solve the Kohn-Sham
(KS) equations.
Within the DFT frameworks, the choices for computational levels and basis sets
vary over a large range. Typically, spin-polarized or non-polarized computations can
be performed with the exchange correlation functional based on a local density
approximation (LDA) considering Perdew-Wang correlational (PWC) [ 117 ], a
generalized gradient approximation (GGA) considering Perdew-Wang (PW91) [118]
or Perdew-Burke-Eruzerhof (PBE) functionals [119], or hybrid functionals (such as
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HSE [120]) taking into account the non-local exchange correlations. The basis sets
widely employed include double numerical basis sets plus p-polarization functions
(DNP) or d-polarization functions (DND) [121,122], double-zeta plus polarization
atomic-orbital basis sets (DZP) [ 123 , 124 ], and plane-wave basis sets [ 125 ].
Particularly, the plane-wave DFT computations are more regularly adopted for
studying nanomaterials [126]. Moreover, to reduce the computational burden caused
by core electrons, psuedopotentials based on the frozen core approximations are
always employed, which consist of effective core potentials [127], DFT-based
semi-core pseudopotentials (DSPP) [128], Troullier-Martins (TM) norm-conserving
pseudopotentials [129], Vanderbilt ultrasoft pseudopotentials (USPP) [130], and
projector-augmented-wave (PAW) pseudopotentials [126].
The computational programs that have been used mainly include VASP [125],
CASTEP [131], PWSCF [132], SIESTA [124], ABINIT [133], and DMol3 [122].
Among them, VASP, CASTEP, PWSCF, and ABINIT are plane-wave codes using
plane-wave basis set and norm-conserving or ultrasoft pseudopotentials, or PAW
method, while DMol3 and SIESTA are all-electron codes using numerical
atomic-orbital basis sets and all-electron or semicore or effective-core
pseudopotentials (for DMol3) and standard norm-conserving pseudopotential (for
SIESTA).
Note that, although DFT computations give powerful predictions for a wide
range of materials and properties, there are still some situations in which the standard
DFT is not physically accurate. For example, both LDA and GGA functionals are
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well known to underestimate band gaps of semiconductors, with errors commonly
larger than 1 eV compared with experimental values. In particular, the self-interaction
error associated with the classical DFT often yields a qualitatively incorrect prediction
for the band gaps and magnetic properties of strongly correlated systems with robustly
localized electrons. Moreover, the standard DFT is only restricted to evaluate the
ground-state energy and properties, and can not deal with computations of electronic
excited states. Also, the standard DFT computations can not well describe long-range
weak interactions, in which LDA tends to overestimate the binding strength, while
GGA tends to behave opposite.
To overcome these inherent limitations, significant corrections have been
introduced to improve the accuracy of DFT computations. Taken as examples,
applying GW approximation to the electron self-energy is successful in accurately
predicting the excited-state properties and improving on the Kohn-Sham band
structure of semiconductors [134]. The screened exchange hybrid density functional,
HSE, has also been successful in predicting the band gaps of semiconductors, and it
appears to provide a good starting point for perturbative GW corrections. Besides, the
DFT+U method [135], which adds a model Hubbard-U term to the simple DFT
functional, provides a qualitatively correct treatment of the electronic structures of
strongly correlated materials. Moreover, the DFT+D method [136], which introduces
dispersion corrections to the conventional Kohn-Sham DFT energy, has been
developed to remedy drawbacks of DFT regarding the dispersion forces.
In the last paragraph of this section, we will briefly sum up the structural
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modeling of 2D nanosheets and 1D nanoribbons. The 2D sheets under investigation
can be constructed directly from layered bulk materials, or are assembled based on
experimental measurements. Depending on the cutting directions, two types of
nanoribbons with either zigzag or armchair shaped edges can be generated. Following
the conventional notation of graphene nanoribbons, the graphene-analogous
nanoribbons are also specified by the ribbon parameter Nz or Na, which is defined as
the number of parallel zigzag chains or the number of dimer lines along the width
directions of zigzag or armchair ribbons, respectively. Taking BN as an example, Fig.
9 clearly illustrates geometric structures of 2D BN monolayer together with its cutted
zigzag and armchair BN nanoribbons with 9 zigzag chains (9-zigzag BNNR) and 15
dimer lines (15-armchair BNNR), respectively. Other types of graphene-analogous
materials are modeled in a similar way.
Fig. 9. BN monolayer (a) and its derived 9-zigzag BNNR and 15-armchair BNNR (b).
3. Planar graphene analogues
In the following sections, we will devote our attention to atomic structures,
material properties, and potential applications of various 2D and 1D materials beyond
graphene and graphene nanoribbons, which are divided into planar graphene
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analogues (Section 3), hypothetical planar materials that have not been experimentally
realized (Section 4), non-planar materials (Section 5), as well as coordination and
covalent organic polymers (Section 6).
3.1 “White graphene”: BN nanosheets and nanoribbons
3.1.1 Comparison between BN and C
BN nanomaterials deserve special attention since they are carbon’s isoelectronic
analogues. In fact, BN exhibits various crystalline forms in many ways analogous to
carbon, including diamond-like cubic BN, graphite-like h-BN, wurtzite-like BN,
onion-like fullerenes, and BN nanotubes. Within these polymorphs, h-BN is
thermodynamically the most stable and softest form, and has attracted increasing
attention. Bulk h-BN has a similar layered structure as graphite with the exception
that the basal planes in h-BN are positioned directly on top of each other, with the
electron-deficient B atoms in one layer lying over and above the electron-rich N
atoms in adjacent layers (Fig. 10). For graphite, however, the adjacent layers are
stacked offset, and alternating C atoms lie above and beneath the hexagon centers.
Both bulk h-BN and graphite exhibit very similar lattice constants and interlayer
distances (lattice constants: a = 2.456 Å, c = 6.696 Å for graphite, and a = 2.504 Å, c
= 6.661 Å for bulk h-BN; interlayer distances: 3.33-3.35 Å for graphite, and 3.30-3.33
Å for h-BN). Due to their close structural similarity, monolayer h-BN can thus be
regarded as a structural analogue of graphene where the C-C bonds are substituted by
alternating B-N pairs. In analogy to the common use of “white graphite” for bulk
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h-BN, the monolayer h-BN is depicted as “white graphene” [19].
Fig. 10. Graphitic structures for carbon and h-BN.
In spite of their structural resemblance, h-BN nanosystems show strikingly
distinct properties from carbon counterparts. The difference in electronegativity
between boron and nitrogen induces ionicity, which is responsible for narrowing of
sp2 derived π bands and the corresponding loss of conductivity, causing h-BN, in
contrast to semimetallic graphite, to be highly insulating. As a matter of fact, h-BN
based materials, either 3D bulk, 2D sheets, or 1D nanotubes, are all electrically
insulating with wide band gaps of 5-6 eV.
On the other hand, h-BN systems also demonstrate advantageous properties. For
example, BN materials are thermally and chemically more stable than carbon
counterparts, and possess a superb thermal conductivity, extraordinary mechanical
robustness, excellent resistance to oxidation (until temperature over 800 C), and good
optical properties. The high thermal stability of h-BN is given by the strong covalent
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network in the plane and the strong polarizability of individual layers. Note that due
to their oxidation resistance, thin h-BN nanosheets can not be prepared by the
oxidation/exfoliation methods as widely adopted for large-quantity fabrication of
graphene.
3.1.2 Electronic and magnetic properties of BN nanosheets and nanoribbons
The isolated BN monolayer inherits the insulating characteristic of bulk h-BN,
and exhibits an indirect band gap (4.3 eV for LDA and 6.0 eV for GW correction)
[137]; whereas the experimentally measured band gap is 5.97 eV [138], larger than
that of the bulk h-BN (5.2-5.4 eV). For 1D BNNRs, since quantum size and symmetry
effects as well as edge effects all emerge, the electronic and magnetic properties of
BNNRs differ a lot from those of BN sheets, which are closely associated with edge
geometries and depend critically on how the edges are passivated, especially for
zigzag edges [139].
Similar to GNRs, spin-polarized magnetic behaviors were only found in
zigzag-typed BNNRs [140]. Armchair-typed BNNRs, owing to strong coupling of the
dangling bonds of dimeric N and B atoms at the same edge, are ground-state
non-magnetic and exhibit semiconducting behaviors with large indirect gaps [141].
The lowest-energy magnetic configuration of bare zigzag BNNRs corresponds to a
ferromagnetic (FM) spin arrangement at the N edge and an anti-FM spin arrangement
at the B edge (corresponding to (+ , ++) magnetic configuration, Fig. 11). This
behavior is in contrast with that found in zigzag GNRs, where both zigzag C edges
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interact via anti-FM spin arrangement mediated by the carbon backbone. Note that
HSE functional predicts all the five spin configurations of bare zigzag BNNRs to be
semiconducting, whereas semilocal PBE functional predicts a half-metallic behavior
for them.
Fig. 11. Five possible magnetic configurations for bare zigzag BN nanoribbons.
Reproduced from Ref. [140]. Copyright © 2008, American Chemical Society.
The bare zigzag edges of BNNRs are less stable than the armchair ones due to
the existence of un-paired edge dangling bonds, and also have less stability than the
reconstructed edge made of 5-7 membered rings [142], which, however, are stabilized
by passivating the dangling edges with H atoms. When only one edge is passivated,
the zigzag BNNRs with passivated B edge and bare N edge are ferromagnetic half
metals [143]. Conversely, the zigzag BNNRs with bare B edge and passivated N edge
are antiferromagnetic p-type semiconductors with an indirect band gap [25,144].
When both edges are hydrogen passivated, the fully passivated armchair and
zigzag BNNRs are all non-magnetic semiconductors with direct and indirect band
gaps, respectively, and their band gaps are strongly affected by external electric field
(Fig. 12) [145]. Band gaps of armchair BNNRs exhibit a family oscillation behavior
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as width increases, and then converge to a value about 0.02 eV smaller than the BN
monolayer (4.53 eV) due to the presence of weak edge states. Under a transverse
electric field, the band gaps of armchair BNNRs are reduced with increased field
strength regardless of the field direction. In contrast to armchair BNNRs, band gaps of
zigzag BNNRs monotonically decrease with widths and converge to a value about 0.7
eV smaller than the BN monolayer due to the presence of strong edge states.
Furthermore, the polarized edge states make zigzag BNNRs exhibit an asymmetric
response to the electric field.
Fig. 12. (a) LDA band gaps of BNNRs versus widths, and dashed lines indicate the
band gap of BN monolayer. (b) Band gaps of 36-armchair BNNR (filled red circles)
and 84-armchair BNNR (empty blue squares) versus field strength. (c) Band gaps of
27-zigzag BNNR (filled red circles) and 46-zigzag BNNR (empty blue squares)
versus field strength. Reproduced from Ref. [145]. Copyright © 2008, American
Chemical Society.
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In the above considered BNNRs, the edge B (N) atoms are of sp2 type. Another
attractive case is the bihydrogenated edge under hydrogen rich environment, where
the fully saturated edge B (N), bonded to two hydrogen atoms, is of sp3 type [146].
Specifically, zigzag BNNRs composed of bihydrogenated B edge and
monohydrogenated or bihydrogenated N edge are all ferromagnetic metals. Armchair
BNNRs, however, are robust non-magnetic semiconductors regardless of hydrogen
contents.
3.1.3 Band-gap modifications of BN nanosheets and nanoribbons
As mentioned in the previous section, perfect BN nanosheets always have wide
band gaps, which are far beyond the gap values (< 3 eV) desired for most of the
current electronic and optical devices, thus become a severe obstacle in processing
BN-based electronics. The well-adopted strategies, such as defects, doping, as well as
surface and edge modifications, can effectively modulate the electronic properties of
BN monolayer and nanoribbons and enhance their device applicability.
3.1.3.1 Defects
The structural and electromagnetic properties of BN nanosheets/ribbons with
unperturbed high perfection of atomic lattices are outstanding; however, structural
defects, in the forms of intrinsic point defects and extrinsic substitution or doping,
might be present during growth or processing. Deviations from perfections, however,
can be useful in some applications since they make it possible to tailor local properties
of BN lattice and to achieve new functionalities. In this section, we will present
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experimental and theoretical surveys on structural defects of BN nanosheets/ribbons,
and discuss influence of these defects on the properties of BN.
Point defects. The commonly observed point defects in experimentally
synthesized BN sheets are triangle vacancy defects with missing B and N atoms.
These vacancies with different sizes have been detected by HRTEM and can in
principle be created in a controlled manner by electron beam irradiation onto BN
sheets (Fig. 13). Jin et al. [99] and Meyer et al. [100] have fabricated the defective
single layer h-BN nanosheets at the beam energy of 120 kV and 80 kV inside the
TEM. Mono-atomic and multi-atomic vacancies, especially B monovacancy and
N-terminated triangle multivacancies, dominated in these BN sheets [147]. The
triangle-shaped vacancies are created mainly due to knock-on effect, which is
mechanically a quasielastic collision between incident electron and nuclei of the
atoms belonging to the h-BN specimens. Since the defects are created under a
high-energy electron beam, the formation and evolution of the triangle-shaped
multivacancies are governed mainly by kinetic factors such as knock-on displacement
rates or probabilities. The preference of B vacancy indicates that boron atoms are
more easily removed than nitrogen atoms.
Theoretical computations showed that structural stabilities of triangle vacancies
depend strongly on the environmental conditions [148-150]. Under nitrogen-rich or
electron-rich conditions, triangle vacancies with N-terminated edges (such as VB and
VB+3N) are more stable. In contrast, under boron-rich conditions, B-terminated edges
(such as VN and VN+3B) are energetically more favorable. The BN diatomic vacancy
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VBN is stable under nitrogen-rich and neutral conditions.
Fig. 13. (a) HRTEM showing lattice defects in h-BN. (b) Models for atomic defects in
h-BN. Reproduced from Ref. [99]. Copyright © 2009, American Physical Society.
The point defects observed in BN sheets are markedly different from those in
graphene. In graphene, an energetically favorable defect is a five-membered ring
(5MR), pentagon. As a consequence, topological defects like reconstructed
pentagon-heptagon or other odd-numbered rings are readily generated in graphene
under electron beam irradiation, especially at edges. However, no reconstructed
vacancy (5MRs) or Stone-Wales (SW) defect (pentagon-heptagon defect) [151] was
observed in h-BN sheets. This is because formation of homonuclear N-N bonds or
B-B bonds is highly unfavorable, making the local reconstruction to form pentagon or
SW defects extremely difficult. As a comparison, Fig. 14 shows the HRTEM images
of BN and graphene membrances forming under electron beam irradiation at 80 keV.
It is clearly seen that a number of triangle-shaped monovacancies are observable,
which do not exhibit any structural reconstruction. While for graphene, three types of
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defect structures are observed: a carbon monovacancy without reconstruction, a
reconstructed monovacancy forming a pentagon-nonagon (5-9) pair, and a
reconstructed pentagon-heptagon (5-7) SW defect. Notably, the reconstructed BN
divacancy (VBN, a Schottky defect pair) has been observed in BN nanotubes, but no
stable BN divacancy was identified in planar BN nanosheets.
Fig. 14. HRTEM images of (a) h-BN and (b) graphene membrane. Reproduced from
Ref. [100]. Copyright © 2009, American Chemical Society.
The electronic and magnetic properties of BN sheets [148,152] and ribbons [153]
with point defects have been investigated theoretically. Depending on the sizes and
edge terminations of created vacancies, BN monolayer should display substantial
magnetism in proportional to the number of unpaired electrons of the N/B atoms
along the zigzag edges. Especially, a single B vacancy in BN monolayer realizes a
half-metallic feature. Vacancy defects also induce spontaneous magnetization and
manipulate electronic properties of BN nanoribbons. The formation of vacancy occurs
preferentially at the vicinity of B edge than other sites. The effect of B or N vacancy
on zigzag BNNRs shows remarkable dependence on defect sites and concentrations,
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and some special band structures such as spin-polarized semiconductor, half metal,
and spin-gapless semiconductor can be achieved. As a further addition of point
defects, line defects are also present [154]. The defect lines are located at the BN
domain boundaries which have different orientations. The existence of extended line
defects presents a new pathway to tune the properties of BN nanoribbons [155].
3.1.3.2 Substitution and adsorption doping
Foreign atoms (molecules) can be incorporated into the BN lattice or embedded
in the vacancy sites as substitutional impurities, or be adsorbed on the BN surface as
adsorption dopants [156]. Theoretical studies dealing with functionalization of BN
layer predicted that adsorption or substitution doping by magnetic metal impurities
(such as transition metals, V, Cr, Fe, Co, Cu, and Mn [157], and noble metals, Au, Ag,
Pt, and Pd [158]), and non-magnetic impurities (such as Be, B, C, N, O, Al, and Si
[ 159 ]) induce magnetism due to emergence of spin-polarized impurity states.
Localized impurity states occurring in the band gap yields interesting properties for
electronic, magnetic and optical applications.
Carbon doping, due to its atomic similarity with B and N, has attracted a lot of
attention. C-substitution for non-edge boron or nitrogen atom in BNNRs induces
spontaneous magnetization, which is independent on the site of substitution or type of
BNNRs [ 160 ]. Post-synthesis of substitutional C doping of h-BN sheets was
implemented via in situ electron beam irradiation inside an energy-filtering 300 kV
HRTEM (Fig. 15a) [161]. Using essentially the same technique, Krivanek et al. [162]
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reported the substitution of both carbon (C substituting for B and N) and oxygen (O
substituting for N) atoms into the BN honeycomb network at electron energies of 60
keV (Fig. 15b). The substitutional C doping transforms BN nanostructures from
electrical insulators to conductors. The doping mechanism was proposed to rely on
the knockout ejections of B and N atoms and subsequent healing of vacancies with
supplied C atoms. Using DFT static and dynamic computations, Berseneva et al. [163]
explained that the C-substitution process is governed not only by the response of such
systems to irradiation, but also by the energetics of the atomic configurations, and it
costs less energy to substitute B atoms than N, especially when the system is
positively charged.
Fig. 15. (a) Schematic for electron-beam-induced substitutional C doping in a
honeycomb BN lattice. Reproduced from Ref. [161]. Copyright © 2011, American
Chemical Society. (b) DFT simulation of a BN layer containing the observed
substitutional impurities. Red, B; yellow, C; green, N; blue, O. Reproduced from Ref.
[162]. Copyright © 2010, Nature Publishing Group.
Drilling holes with an atomic-size electron beam and filling them with atoms
(such as C or other elements) was a promising method for constructing hybrid
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structures with atomic precision and for tuning electronic properties of diverse B-C-N
structures in a full compositional range between pure BN and C systems.
Surface doping by organic molecules through non-covalent interaction is a
simple and effective method to tune electronic structure of BN nanosystems. Because
of interfacial charge transfer, surface adsorption with strong electron acceptor (tetra
cyanoquinodimethane, TCNQ) or electron donor (tetrathiafulvalene, TTF) molecule
can significantly reduce the wide band gap of pristine BN nanosheets/ribbons and
result in a p- or n-type semiconductor, respectively [164]. The origins for band
modifications are attributed to the introduction of new impurity states near the top
valence band or the bottom conduction band of pristine BN nanosheets/ribbons,
contributed by the highest occupied molecular orbital (HOMO) or lowest unoccupied
molecular orbital (LUMO) levels of the adsorbed molecules.
3.1.3.3 Surface modifications of BN nanosheets
The surface of BN nanosheets can be chemically decorated by various groups
such as hydrogen, fluorine and oxygen, and the hybridized states of B or N atoms in
these cases change from sp2 into sp3. Nevertheless, there is some difference in surface
functionalization between BN and graphene. Structurally, hydrogenated graphene
prefers chair-like configuration (graphane). For hydrogenated BN sheets, however,
the chair configuration is less favorable than boat or stirrup configurations [165]. Note
that the chair-like structure can be stabilized by surface modifications with groups of
large atomic radius such as F, Cl, OH, CH3, and NH2 [ 166 ]. Interestingly,
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hydrogenation opens a band gap in graphene, but it reduces the band gap of BN sheets
[167]. At the DFT-GGA level, the hydrogenated BN monolayer has a direct band gap
of 3.33 eV. The semihydrogenated BN sheets can form by applying electric fields to
the fully hydrogenated ones or by single-side hydrogenation on the substrate. The
semihydrogenation on BN sheet is slightly different from graphene because of the
heteroatomic inequivalence of B and N atoms. Boron-semihydrogenated BN sheet is
more stable and behaves as a ferromagnetic half-metal [167].
Due to its strong electron affinity, F atom has a strong site preference to bond
with B, but a weak bonding with N. The magnetism and band structures in fluorinated
BN can be controlled by applying strain [168]. Promoted by the selectivity of
fluorination with B atoms, surface fluorination of few-layered BN sheets produces
stable F-terminated cubic BN nanofilms [169], which exhibit controllable band gap by
altering thickness or applying electric fields.
Moreover, in analogy to oxidation of graphene to form graphene oxide, the
surface oxidation of BN sheet generates the BN derivative, BN oxide, which tends to
form an O domain or O chain on the BN sheet, and introduces unoccupied impurity
states leading to band gap reduction of BN sheet [170].
Besides BN sheet, hydrogenation also plays a crucial role in engineering
electromagnetic properties of BNNRs, as indicated by Chen et al. [26]. The fully
hydrogenated armchair BNNRs act as non-magnetic semiconductors, while the fully
hydrogenated zigzag BNNRs behave as ferromagnetic metals. The partially
hydrogenated zigzag BNNRs exhibit diverse properties, where as hydrogenation
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coverage increases, a semiconductor→half-metal→metal transition occurs,
accompanied by a nonmagnetic→magnetic transformation (Fig. 16). For all the
hydrogenated zigzag BNNRs, the bands around the Fermi level are dominated by the
edge B/N atoms with sp3 hybridization and the related H atoms, and increasing
hydrogenation ratio causes the highest valence bands and the lowest conduction bands
to shift towards the Fermi level and cross it eventually. Thus, controlling the
hydrogenation ratio can precisely modulate the electronic and magnetic properties of
zigzag BNNRs, which endows BN nanomaterials many potential applications in the
novel integrated functional nano-devices.
Fig. 16. Structural and electronic properties of partially hydrogenated 8-zigzag
BNNRs with different ratios of LH to L0: (a) 1:4, (b) 2:3, (c) 3:2, and (d) 4:1. (e)
Structure of fully hydrogenated zigzag BNNR. Parts f and g show zooms on the
region about the Fermi level of band structure and density of states (DOS) of (d).
Reproduced from Ref. [26]. Copyright © 2010, American Chemical Society.
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It is worth noting that although completely and partially hydrogenated or
fluorinated graphene has been successfully fabricated, the surface-hydrogenated or
-fluorinated BN sheets are rarely available. So far, only the partially hydrogenated
few-layered BN membranes have been experimentally achieved by hydrogen plasma
treatment [171]. Moreover, the surface oxidation under strong oxidizing reagents is a
common reaction of graphene, but the oxidation reaction of BN counterpart becomes
extremely difficult due to its chemical inertness. With these arisen challenges, the
realizations of hydride-, fluoride- and oxide-functionalized BN derivatives are
expected to find their unique technical solutions in future.
3.1.3.4 Edge modifications of BN nanoribbons
The edges are typically more reactive than the surfaces, and edge modifications
can effectively control the electronic properties of BNNRs. Hydrogen atoms are the
typical modifiers for BN edges, and are usually used as the reference for comparison
with other edge modifiers. Chemical functionalization of zigzag BNNRs edges by Cl,
OH, and NO2 results in a narrowed band gap, while decorating with F atom leads to
an enlarged gap [172]. For these edge decorations, B edges are more favorable than N
edges for adsorption. Especially, edge functionalizations by F and OH give higher
stability compared with the hydrogen passivated counterparts, suggesting that these
modifications have great potentials to be realized experimentally. Interestingly, both
O-edge and S-edge terminated zigzag BNNRs become metallic [173]. The O-edge
terminated armchair BNNRs, however, remain semiconducting [174]. The chemically
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modified armchair BNNRs by different transition metals can be semiconductors,
half-metals or metals with diverse magnetic ground states depending on the types of
metals [175]. These interesting band-gap modulations by edge decorations may be
exploited for nanoelectronic applications.
3.1.4 BN/graphene hybrid structures
Stimulated by the intensive studies of graphene and BN, the hybrid analogues of
graphene and h-BN, containing boron, nitrogen and carbon, have also attracted
intensive interests. Because of their commensurate structural parameters and distinct
electronic properties, graphene and BN are considered as good candidates for
fabricating B/N/C materials offering new functionalities. Other than forming solid
solution of B, N, and C, which has been studied extensively, it is necessary to explore
the possibility of making a graphene/BN composite, where the two phases coexist
separately, in a layered structure or in a same plane. Such a novel composite is our
main focus here.
A layered graphene/BN superlattice can be upheld together by vdW interactions,
such as placing graphene on single layer BN to form a bilayer system [176] and on
two or more BN layers to construct a multilayer heterostructure [177], or designing
3D superlattice with alternate stacking of graphene and BN monolayer [178]. Other
heterostructures such as monolayer [179] and bilayer [180] graphene sandwiched
between two BN layers or monolayer BN sandwiched between two graphene layers
[181] have also been constructed.
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For graphene/BN bilayer, the most stable configuration has one C atom on top of
a B atom while the other C atom sites above the center of the BN hexagon ring, and a
small gap (53 meV) is opened around the K-point in graphene resulting from
symmetry-breaking of the two carbon sublattices [176] (Fig. 17). The band gap of
graphene/BN bilayer or multilayer structure can be modulated by tuning the interlayer
spacing and stacking pattern [182], hydrogenating [183], introducing strain [184], or
applying an electric field in the direction perpendicular to the hybrid layers [185].
Note that according to Xu et al.’s report [186], the response of electronic properties of
graphene/BN heterostructurs on the applied external pressure might lead to the
potential realization of atomic-scale pressure sensors.
Fig. 17. Model (a) and band structure (b) of single layer graphene on h-BN.
Reproduced from Ref. [176]. Copyright © 2007, American Physical Society.
Experimentally, many efforts were devoted to synthesize graphene/BN hybrid
heterolayers [187]. For example, Gannett et al. [188] and Kim et al. [189] separately
grew graphene on Cu foil through CVD technique, and then transferred the fabricated
graphene onto the exfoliated BN underlayer. Graphene has also been grown by CVD
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onto BN monolayer on Ru(0001) [190] and Ni(111) [191] surfaces. Alternatively,
Bresnehan et al. [192] firstly grew large-scale BN layer on Cu foil by a catalytic
thermal CVD, and followed by the transfer of BN to quasi-freestanding epitaxial
graphene grown on SiC wafer. Besides the use of metal catalysts, Son et al. [193]
reported a direct catalyst-free growth of flat graphene pads by atmospheric CVD
process on top of mechanically exfoliated BN substrate. Lin et al. [194] introduced a
hydrogen flame synthesis method to successfully grow graphene on BN flakes using
PMMA as carbon sources. Very recently, Haigh et al. [195] have reported a
step-by-step preparation of graphene/BN multilayer heterostructures in which the
mono- and bilayer graphene were individually interlaid between atomically thin h-BN
crystals. As a representative case, Fig. 18a sketches the hybrid devices with two
graphene layers (dark grey) intercalated into thin BN layers (blue) deposited on top of
an oxidized silicon wafer (violet and grey). Fig. 18b shows the optical microscopy of
fabricated heterostructure, and the graphene Hall bar is highlighted with thin black
lines. The scanning electron microscopy (SEM) images of a cross-sectional specimen
extracted from the device are presented in Fig. 18c and d, and the red circles (known
as “bubbles”) indicate the trapped chemical species, mostly the surface-adsorbed
hydrocarbons. These adsorbates are detrimental to graphene performance, and can be
further eliminated by thermal annealing to obtain clean and atomically flat interface
[196]. The high-resolution scanning transmission electron microscopy (STEM) image
in Fig. 18e provides a useful tool to count the atomic layer sequences, and clearly, this
unique device is alternatively stacked with one graphene layer and four BN layers
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(Fig. 18f).
(a) (b) (c)
(d)
(e) (f)
Fig. 18. (a) Schematic illustration that two graphene monolayers are interlaid with
h-BN crystals. (b) Optical image, (c) SEM micrograph, and (d) low-magnification
SEM image of the device. (e) High-resolution bright-field aberration-corrected STEM
image of the graphene-h-BN heterostructure. (f) Schematic of the atomic layer
sequence for this device. Reproduced from Ref. [195]. Copyright © 2012, Nature
Publishing Group.
Actually, because of the good lattice match and the atomically smooth surface
free of dangling bonds and charge traps, h-BN serves as an excellent back-gate
dielectric layer for graphene-based field-effect transistors [ 197 , 198 ], electron
tunneling devices [199], and is used as an appealing supporting-substrate to enhance
carrier mobilities of graphene electronics that have better quality than on SiO2 or
other high-k dielectrics substrates such as HfO2 and Al2O3 [200 - 202 ]. As a
comparison, Fig. 19 presents a side-by-side STM topography of typical graphene
corrugations on BN and SiO2 substrate [203], and they obviously reveal a reduced
roughness of graphene/BN over large areas than its graphene/SiO2 counterpart.
Moreover, the charge density inhomogenety in graphene/BN system is also
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significantly reduced as compared with graphene/SiO2 system, which is clearly
observed from the charge density maps in Fig. 19c and d. Note that placing graphene
onto BN can actually yield improved device performance, yet contrary to the
theoretical prediction; the existence of gap-opening for graphene/BN system has not
been detected by spectroscopy measurements, which is mainly because of the stacking
misalignment of the graphene/BN lattices [204].
Fig. 19. Comparing topography and charge density for graphene/BN and
graphene/SiO2. Reproduced from Ref. [203]. Copyright © 2011, American Chemical
Society.
In another way, a small fraction of BN region can be substituted with C atoms or
vice versa to form hybridized single layer superlattice, and there has been some
success in preparing BxCyNz compositions. For example, based on CVD growth, Ci et
al. [76] fabricated hybrid BNC atomic sheets containing randomly and separately
distributed h-BN and C phases, as shown by the atomic model in Fig. 20. In their
experiment, methane and amine borane (NH3-BH3) were adopted as the precursors for
C and BN, respectively. By controlling the relative proportion of C and BN sources
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during the CVD process, the atomic ratio of B, N, and C of BNC hybrids can be
flexibly tuned over a wide scale ranging from pure BN to pure graphene, and the
electrical properties can be accordingly controlled from insulators to conductors by
increasing the C concentration. Lin et al. [205] developed a method to convert
graphene oxide (GO) into BCN nanosheets through substitutional doping by reacting
GO with B2O3 and ammonia at 900-1100 C. Based on a sequential CVD growth,
Sutter et al. [206] prepared monolayer graphene-BN heterostructures on Ru(0001)
substrate by exposing the metal surface to high-purity ethylene and borazine,
respectively. Pakdel et al. [207] developed a non-catalytic CVD method for growing
boron nitride-carbon (BN-C) phase-separated composite nanosheet coatings on
Si/SiO2 substrates. A chemical blowing method which relies on making large bubbles
within thin B-N-C-H polymer walls was also used to produce Cx-BN nanosheets after
high-temperature annealing and collapse of the polymer bubbles [101]. Moreover,
Raidongia et al. [111] synthesized a graphene-like material with composition of BCN
by reacting high-surface-area activated charcoal with a mixture of boric acid and urea
at 900 C. This material mainly consists of two to three layers, showing a relatively
high surface area of 2911 m2g-1 and exhibiting a high propensity for CO2 adsorption
and hydrogen uptake. The BCN material exhibitd femtosecond third-order nonlinear
optical susceptibility and hot carrier dynamics, making BCN an attractive candidate
for untrafast optical applications [208].
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Fig. 20. Atomic model of the h-BNC film showing hybridized h-BN and graphene
domains. Reproduced from Ref. [76]. Copyright © 2010, Nature Publishing Group.
Theoretically, hybrid graphene/BN single-layer sheets were predicted to display
tunable properties depending on the arrangement of the separated graphene or BN
regions and the compositional ratio of carbon to BN in the structure [ 218 ].
Structurally, graphene has a smaller interatomic distance (1.42 Å) than BN (1.45 Å),
and the incorporation of graphene into BN lattice would induce a tensile stress after
relaxation, which ensures the overall planarity of the C-BN composite system.
Because of quantum confinement, the intrinsically semiconducting graphene islands
embedded in the BN matrix gain the character of quantum dots [219]. Typically, a
hybrid BCN honeycomb sheet with triangle-shaped graphene quantum-dot shows
flat-band ferromagnetism, while the magnetism vanishes when the embedded
graphene adopts a hexagonal-shape quantum dot [220]. The magnetism mainly
originates from π electrons of the embedded graphene, especially located at the C
atoms along the zigzag edges of the graphene flakes. Essentially, the substitution of
one B atom by one C atom will bring one more electron into the system, while the
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substitution of one N atom by one C atom will introduce one more hole, and the net
effect of the hybrid system becomes hole-injected or electron-injected if the
embedding border is connected with C-N (in N-rich condition) or C-B (in B-rich
condition) bonds. In both cases, the band gap of the hybrid architecture reduces with
increasing the embedded graphene size. On the other hand, the BN quantum dots
embedded into graphene lattices would lead to gap opening of graphene regardless of
sizes and geometries of the interior BN domains [221,222].
Structural hybridization by substituting graphene C-C pairs with isoelectronic
BN pairs also achieves significant tuning on the electronic transmission of 1D
graphene/BN nanoribbons [223,224]. Note that half metallicity can be realized in
graphene nanoribbons by applying an external electric field, yet a carefully designed
graphene-BN material such as zigzag graphene/BN nanoribbons with dihydrogenated
B edge and the zigzag graphene nanoroads embedded in BN can eliminate the need of
electric field to possess intrinsic half metallicity [225,226]. For BN-embedded zigzag
GNRs (Fig. 21) [227], a gradual replacement of the zigzag C-C chains in the middle
part by zigzag B-N chains transforms the system finally to the zigzag BN nanoribbons,
and the electronic structures vary accordingly with the doping concentration. At a
high doping concentration, the hybrid nanoribbons with terminated polyacene C
chains at the edge and all the substituted B-N chains in the internal part act as
half-metallic antiferromagnets for all the widths. Application of electric field on
zigzag BN nanoribbons with polyacene edges shows that the half-metallic behavior
sustains over sufficiently large electric field strength. The Lewis acid character of
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boron is identified to be responsible for the charge transfer from adjacent C atoms to
the B atoms, resulting in an interface potential gradient analogous to the effect of
external electric field and invoking half metallicity. Furthermore, at a precise B-C-N
composition, the BC2N sheet is a direct-gap semiconductor with a band gap of 1.6 eV
which is intermediate between graphene and BN, and the derived BC2N nanoribbons
can be semiconducting, metallic, or half metallic, which is width- and edge-dependant
[228].
Fig. 21. A hydrogen passivated 8-ZGNR without (upper) and with (lower) BN doping.
Doping concentration increases gradually along the directions of solid arrows.
Electric field was applied along the y-axis direction. Reproduced from Ref. [227].
Copyright © 2009, American Physical Society.
Note that the C-BN hybridized phase not only shows great potential for band-gap
engineering, but also demonstrates promising applications in hydrogen storage since
the interface between graphene and BN exhibits high affinity for gas adsorption,
especially at the zigzag-type interface [111,229]. Beyond this, Wang et al. [230] have
recently showed that the BCN graphene exhibits efficient metal-free electrocatalytic
activity for oxygen reduction reaction (ORR) in fuel cells. Lei et al. [231] revealed
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that the BCN nanosheets showed promising storage performance in lithium ion
batteries, and exhibited a stable capacity of ~100 mA h g-1 at 2 A g-1 for 5000 cycles.
3.1.5 Potential applications
The unique and exceptional properties with regard to the mechanical, electronic,
thermal, and chemical aspects favor BN nanosheets (BNNS) for use in a wide range
of electronic and composite applications. For example, h-BN nanosheets are widely
used as ultraviolet-light laser devices [81,232], field emitters [233], semiconductor
diodes [234], and insulating thermoconductive nanofillers in polymer or ceramic
composites [39,235]. For example, ultrathin BN nanosheets protruded from Si3N4
nanowires [233] or BN fibers [236] as well as porous BN nanospheres [237] display
excellent field emission performance with electron emission property comparable to
that of carbon nanotubes. BN layers also widely act as good insulating substrates for
graphene-based electronics, as discussed in the previous section.
On the other hand, due to its attractive combination of electrical insulation, high
thermal conductivity, and optical transparency, layered h-BN holds great promise in
applications as polymeric composite fillers. According to the experimental results, the
incorporation of exfoliated BN sheets into PMMA [39] and
poly[2,2'-(p-oxydiphenylene)-5,5'-bibenzimidazole (OPBI) [42] matrix showed
greatly improved mechanical and thermal properties as compared with the neat
polymers. As an example, Fig. 22 presented a comparative study on the thermal
mechanical analysis (TMA) and mechanical strength test of blank PMMA polymer
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and its BN composite (with 0.3 wt% BN addition). By comparing with the parent
PMMA, the embedment of BN nanosheets in PMMA results in a remarkable
reduction of the coefficient of thermal expansion (CTE) and leads to enhanced elastic
modulus (22% increase) and strength (11% increase). Besides, Song et al. [238] have
recently showed that nanocomposites containing exfoliated BN nanosheets and
poly(vinyl alcohol) or epoxy resin polymers possess superior thermal transport
performance. Furthermore, the BN layers could be used as nanofillers in heat transfer
fluids. Taha-Tijerina et al. [239] have recently reported the synthesis of a stable
nanofluid with BN fillers in mineral oil (MO), a commonly used heat-transfer fluid in
transformers. This novel BN/MO nanofluid has a high thermal conductivity which
tends to increase with h-BN filler concentration, and only 0.1wt% addition of h-BN in
MO can achieve a significantly enhanced thermal conductivity of around 76% at 373
K.
Fig. 22. (a) Comparative TMA and (b,c) mechanical strength tests on blank PMMA
and its BN nanosheet composite. The inset in (a) shows CTE data before and after
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glass-transition temperature. Reproduced from Ref. [39]. Copyright © 2009,
Wiley-VCH.
Moreover, functionalized BN nanosheets also demonstrated great potential to
fabricate polymer composites. In an experimental work, Liu et al. [240] recorded that
BN nanosheets functionalized by P(S-b-MMA) modifier serve as a better nanofiller to
improve the tensile strength of PMMA and polystyrene (PS) matrix than the
unmodified BN filler. Quiet recently, Sainsbury et al. [241] described a novel
functionalization strategy to yield hydroxyl-functionalized BN nanosheets. This
involves a two-step procedure: initially covalently grafting alkoxy groups to the B
atoms in h-BN lattice via oxygen radical attack, and the subsequent hydrolytic
defunctionalization of the alkoxy groups to generate surface OH-functionalized BN
nanosheets (OH-BNNSs) (Fig. 23a). Note that the bare BNNSs are typically
superhydrophobic, while surface modification by introducing hydrophilic OH
functional groups renders the nanosheets water-soluble. Fig. 23b presents photographs
of water, pristine BNNSs, and OH-BNNSs (from left to the right) under an irradiation
of a 532 nm green laser, which clearly shows the enhanced solubility and scattering
behavior of the OH-BNNSs relative to the pristine BNNSs. When used as polymer
nanofillers, the OH-BNNSs significantly reinforce the polyvinyl alcohol (PVA) in its
mechanical strength (Fig. 23c and d) over the pristine BNNSs, since the presence of
surface OH groups provide better chemical compatibilization and dispersability with
the host matrix.
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Fig. 23. (a) Reaction scheme indicating the two-step procedure for oxygen radical
functionalization of BN nanosheets. (b) Photographs of water, pristine BNNSs and
OH-BNNSs in water. (c) Stress-strain curves for PVA reference and BNNS:PVA,
OH-BNNS:PVA composites. Inset shows the transparent samples on glass
microscope slides. (d) Young’s modulus, ultimate tensile strength, strain to break, and
toughness for the PVA reference and composites. Reproduced from Ref. [241].
Copyright © 2012, American Chemical Society.
In addition to the applications mentioned above, BN nanosheets are also
promising as an alternative for graphene to plane electronics where lithography
technique can be conveniently employed, as superhydrophobic surface coating, as
nanoscale support for metal and metal oxide catalysts. For example, the BN nanomesh
deposited onto transition metal surfaces is thermally very stable and serves as a
template for surface self-assembly of well-ordered organic molecules [242], water
molecules [243], rare gas atoms [244], transition metal atoms [245], or metallic
nanoparticles [246].
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These substrate-supported molecular or atomic arrays can be used to produce
functional supramolecular BN nanostructures, which demonstrate great potential in
catalytic and sensing applications. For instance, a theoretical study by Gao et al. [247]
suggested that Au supported on BN surface exhibits enhanced adsorption and
catalytic activation for O2 molecule. Experimentally, Wang et al. [248] showed that
Au- and Pt-nanoparticles loaded onto BN nanosheets function as efficient catalysts
towards CO oxidation. Wang et al. [249] also showed that Au nanoparticle-loaded
h-BN nanocrystals can be used for oxidation of benzyl alcohol to benzaldehyde. Apart
from the catalytic applications, Lin et al. [250] reported the fabrication of Ag
nanoparticle-decorated BN nanosheets, which were further transferred onto quartz
substrates to fabricate reusable and oxidation-resistent surface enhanced Raman
spectroscopy (SERS) sensor devices.
Note that the 1D BN nanoribbons are theoretically predicted to have significant
potentials in the area of nano-electronics and spintronics; however, their realistic
applications remain currently less explored by experimentalists. This challenge is
possibly due to the large experimental difficulty in producing high-quality BNNRs
with uniform and smooth edges.
3.2 Silicene
Silicene, a one-atom-thick silicon sheet arranged in a 2D honeycomb lattice, is a
new allotrope of silicon, similar to graphene. Due to their structural resemblance,
single-layer silicene is also alternatively considered as a Si-based 2D counterpart of
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graphene. However, unlike graphene, silicene favors to adopt a low-corrugated
honeycomb configuration (about 0.44 Å buckling is present) [251] rather than a planar
one, forming a mixed sp2-sp3 like hybridized state [252]. This difference stems from
the fact that for C atoms, the fully sp2-hybridized state is more stable than the
sp3-hybridized state, while the case for Si is reverse. This could also be otherwise
understood by the instability of Si atoms to form strong Si-Si π bonds to stabilize the
planar three-fold coordination. Due to its unique bonding behavior, bulk Si can not
form a layered phase like graphite, yet one exception is that the layered Si atoms can
survive in alkaline-earth-metal silicides (such as CaSi2), which are structurally
composed of alternative layer-by-layer packing of hexagonal alkaline-earth-metal
layers and corrugated Si(111) atomic layers.
3.2.1 Synthesis of silicene and functionalized silicene
The possible existence of silicene has been theoretically conjectured since 1994
and has been just synthesized in recent years [253]. Note that silicene is a meta-stable
structure, which is 1.56 eV/atom less stable than bulk Si. This indicates that growth of
silicene necessitates the use of deposition substrates. Actually, although free-standing
silicene has not been realized up to now, experimental evidence for the existence of
epitaxial silicene does exist, which has been discussed in detail in the previous
Section 2.4.
At present, the silicene structures are obtained mainly by the surface-assisted
epitaxial growth, such as the observed silicene nanoribbons on Ag(110) [84-86] and
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silicene sheets on Ag(111) [88-93], Ir (111) [94], or ZrB2(0001) surface [95]. Unlike
the high oxidation reactivity of typically sp3-bonded Si surface, De Padova et al. [254]
discovered that the 1D grafting of silicene nanoribbons grown on Ag(110) surface are
room-temperature oxidation-resistant, which compares favorably with graphene.
Dávila et al. [255] have recently showed that isolated or self-aligned Si nanoribbons
on Ag(110) can be easily attacked by atomic or molecular hydrogen adsorption.
Note that even prior to the synthesis of silicene, 2D functionalized Si nanosheets
have been prepared by Nakano et al. using solution-phase methods. For example,
layered siloxene [Si6H3(OH)3], which is composed of corrugated Si layers with Si
atoms being terminated by H and OH groups, alternatively, was exfoliated by using
sodium dodecyl sulfate as dispersant [256]. To prepare oxygen-free modified Si
nanosheets, they performed the exfoliation of layered polysilane (Si6H6) as a result of
reaction with n-decylamine, and the resulting Si nanosheets are covalently covered
with organic amine groups [257]. In particular, they also synthesized oxygen-free,
phenyl-modified organosilicon nanosheet (Si6H4Ph2) by reacting the layered Si6H6
with PhMgBr [258] (Fig. 24). The yielding Si6H4Ph2 exhibits a completely flat plane
surface, with a measured sheet thickness of 1.11 nm and a periodic distance of 0.94
(0.98) nm of surface phenyl groups, in consistency with the thickness (0.98 nm) and
periodicity (1.0 nm) calculated on the basis of its atomic architecture. Very recently,
Nakano et al. have also synthesized alkyl-modified Si nanosheets through
hydrosilylation of layered Si6H6 [259]. Structurally, these modified Si nanosheets are
the functionalized derivatives of 2D silicene structure, which can be regarded as a
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new class of 2D functional Si materials. Note that unlike the bare silicene sheet, the
functionalized Si nanosheets can stably exist as a free-standing form, and become
soluble in organic solvents. Thus they are likely to be deposited onto various
substrates for characterization and application.
Fig. 24. (a) AFM image of [Si6H4Ph2]. (b) Line profile along the black line in (a). (c)
Model of [Si6H4Ph2]. (d) AFM image of the surface of [Si6H4Ph2]. (e) Line profile
along the black line in (d). (f) Top view of the model structure for [Si6H4Ph2].
Reproduced from Ref. [258]. Copyright © 2010, American Chemical Society.
3.2.2 Theoretical investigations of silicene
The free-standing silicene sheet and its nanoribbons have been widely
investigated through first-principles computations, and their electronic properties bear
great resemblance to those of graphene.
2D silicene. Structually, the buckled Si hexagonal sheet is optimized to have a
lattice constant of a = 3.86 Å, and the nearest Si-Si distance is relaxed to be 2.28 Å.
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Compared with the C-C distance (1.42 Å) in graphene, such large Si-Si distance leads
to greatly weakened π-π overlaping, and thus the planar Si structure can not be
stabilized. Due to its structural corrugation, silicene has a lower D3h symmetry instead
of the D6h symmetry of graphene (Fig. 25a). Despite this, they share essentially
similar electronic structures and consequently analogous physico-chemical properties.
Electrically, silicene is a gapless semimetal, and its charge carrier behaves like a
massless Dirac fermion due to the crossing of π and π* bands at the Fermi level,
similar to graphene (Fig. 25b). Note that although the free-standing silicene is
expected to have a zero gap, a minigap can be opened in epitaxial silicene, resulting
from the symmetry-breaking induced by the interaction with the Ag substrate [260].
The existence of Dirac ferumions for silicene on Ag(111) has been observed by
angular-resolved photoelectron spectroscopy (ARPES) [89] and scanning tunneling
spectroscopy (STS) [261], and an Fermi velocity of vF = 1.3 × 106 ms−1 is obtained,
which is comparable to that found for graphene. Besides, the graphene-like Dirac
cones have also been observed by ARPES in monolayer and multilayer silicene
nanoribbons grown on Ag(110) [85,86] .
Fig. 25. (a) Top and side view of silicene. (b-d) Band structures of silicene around Ef
at three different vertical electric fields. Reproduced from Ref. [263]. Copyright ©
2011, American Chemical Society.
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Because of their similarity in electronic properties, all graphene expectations,
such as high-speed nanoelectronic devices based on ballistic transport, can be
consequently transferred to silicene due to its high compatibility with the current
Si-based semiconductor industries. Moreover, since Si atoms have greater intrinsic
spin-orbit coupling strength than C atoms, silicene with topologically nontrivial
electronic structure is theoretically predicted to have a spin-orbit band gap of 1.55
meV that is much larger than that of graphene, and exhibit quantum spin Hall effect in
an accessible temperature regime [262]. This property makes silicene particularly
interesting for applications as spin Hall effect devices.
Silicene is a versatile material, because not only it owns the fascinating
properties of graphene, but also the structural puckering leads to a more reactive
surface, making band-gap-tuning of silicene much easier, which will be discussed in
the following part.
The electronic structures of silicene can be tuned by external electric field. For
example, Ni et al. [263] revealed that the band gap opens in buckled silicene by
applying a vertical electric field, and the size of the band gap increases linearly with
electric field strength. Specifically, the opened band gap for silicene is 0.08 eV under
E = 0.51 V/Å and becomes doubled under E = 1.03 V/Å (at the GGA/DNP level) (Fig.
25c and d). It is also remarkable that a topological quantum phase transition is
induced in silicene by changing the electric field [264]. The main effect of electric
field is to break the symmetry in A and B sublattices of silicene’s honeycomd and
hence leads to gap opening in the band structure. It should be recalled that by
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imposing an external electric field graphene still remains its zero-gap semimetallic
character since its two sublattices remain equivalent. As a result, the biased
monolayer graphene did not function effectively as a field effect transistor (FET). For
comparison, the effective band-gap response of monolayer silicene as a function of
external electric field makes it a more suitable candidate for FET devices.
In addition, silicene can be alternatively modified by surface functionalization.
For example, hydrogenation opens a band gap in silicene, and the fully hydrogenated
silicene (also known as polysilane or silicane) with a composition of SiH is a
wide-band-gap semiconductor, with its band gap ranging between 2 (DFT-LDA level)
and 4 eV (GW approximation) [265]. Similar to graphene, the fully hydrogenated
silicene also prefers the chair-like configuration, but exhibits an indirect band gap.
Interestingly, surface decoration by halogen atoms (F, Cl, or Br) changes the indirect
gap of chair-like polysilane into the direct one, and the resulting gap value (1~2 eV) is
predicted to be smaller than the hydrogenated counterpart [266]. For the partially
hydrogenated silicene, ferromagnetism is introduced by attaching hydrogen atoms on
one side of the sheet (semihydrogenated silicene) [267]. The semihydrogenated
silicene is identified as a magnetic semiconductor with an estimated gap of 0.94 eV
(DFT-GGA level, Fig. 26), and is corrected to be 1.74 eV with HSE06 functional.
Osborn et al. [ 268 ] also studied the partial hydrogenation of silicene with
hydrogenation ratio ranging between 3.1% and 100%, and showed that by specifically
patterning the silicene with H domains, a metal-semiconductor-insulator functionality
can be obtained.
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Fig. 26. Top (a) and side (b) view of semihydrogenated silicene, and its band structure
(c). Reproduced from Ref. [267]. Copyright © 2012, American Chemical Society.
Moreover, surface adsorption of metal atoms provides a variety of ways to
functionalize and modify the electronic properties of silicene. Lin et al. [269] revealed
that metal atoms, such as alkali and alkali-earth metals (Li, Na, K, or Ca), groups III
and IV metals (Al, Ga, In, or Sn), and transition metals (Ti, Fe, Co, Ni, Pd, Pt, or Au),
form strong interactions with silicene, and the formed complexes exhibit rich
electronic properties with the introduced s or d electron states lying above or below
the Fermi level. The strong binding of metal adatoms to silicene makes it potentially
useful in gas storage, superconductivity and catalysis. Osborn et al. [270] additionally
studied the Li chemisorption on silicene, and showed that the completely lithiated
silicene (composition of SiLi) is a stable and semiconducting material with a band gap
of 0.37 eV.
1D silicene nanoribbons. Like graphene, the applications of silicene in future
nanoelectronics would need a band gap. In our above discussion, the electric field
engineering only provides an external control, and the surface functionalization would
induce the change of hybridized state into sp3. Another effective method to achieve
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gap opening in silicene is to reduce the dimensionality into 1D nanoribbons.
Due to the presence of quantum and edge effects, silicene nanoribbons exhibit
intriguing properties distinct from those of silicene, which are electrically width and
edge dependent [251]. Because of the instability of edge dangling bonds, the bare
armchair or zigzag edges undergo a sharp (2×1) reconstruction (Fig. 27), while the
reconstruction disappears when the edges are terminated by H atoms. The armchair Si
nanoribbons, with either bare or H-terminated edges, are all non-magnetic
semiconductors. Similar to armchair GNRs, their semiconducting gaps exhibit an
oscillatory behavior, in which the band gap is smaller for n = 3p + 2 than for n = 3p or
3p +1 (n is an integer) (Fig. 27 b and c). Zigzag Si nanoribbons, on the other hand, are
metallic or semiconducting depending on whether their edges are saturated or not.
Particularly, the bare zigzag ribbons have a non-magnetic and metallic ground state
(Fig. 27d). After H-saturation, the antiferromagnetic (opposite spin on the two
different edges) semiconducting state becomes most stable (Fig. 27e), and the band
gap of zigzag Si ribbons decreases monotonically as the width increases (Fig. 27f),
sharing great similarity with zigzag GNRs [271]. Like the fully hydrogenated GNRs,
the fully hydrogenated armchair- or zigzag-silicene nanoribbons are all non-magnetic
semiconductors [272].
The electromagnetic behaviors of zigzag Si nanoribbons can be modified by
external electric field or magnetic field. Especially, when a transverse electric field is
applied, the zigzag Si ribbons become half-metals [271], and a spin-polarized current
can be generated [273]. By applying a vertical magnetic field, Xu et al. [274] found
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that zigzag Si nanoribbons can switch between anti-FM and FM coupled states, and a
large magnetoresistance (MR) is obtained due to the large current difference between
the semiconducting anti-FM state and the metallic FM state. This suggests potential
technological applications of silicene in spin-valve devices [275].
Fig. 27. (a) Ideal and relaxed bare 10-armchair Si nanoribbon. Band gaps as a
function of width for bare (b) and H-terminated (c) Si armchair nanoribbons. (d) Bare
Si zigzag nanoribbons showing two different (2×1) reconstruction indicated by ‘‘1’’
and ‘‘2’’ and the band structure corresponding to ‘‘1’’. Reproduced from Ref. [251].
Copyright © 2009, American Physical Society. (e) Spin density distribution of
H-terminated 6-zigzag Si nanoribbon at anti-FM state and its band structure. (f) Band
gaps of H-terminated zigzag ribbons vs the number Nz of ribbon width. Reproduced
from Ref. [271]. Copyright © 2009, American Institute of Physics.
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3.2.3 Potential applications of silicene
As an important structural analogue of graphene, silicene possesses extraordinary
electronic properties, and its reactive surface can be modified by many methods.
Besides, it is advantageous that silicene can be easily integrated into Si-based
semiconductor industries. Although the current progress on silicene is still at its
infancy, silicene is expected to be developed for many technical prospects [276]. For
example, due to its high electron mobility, an intrinsic opening of spin-orbital band
gap, and tunable electronic structures, silicene can be used in nanoelectronics,
spintronics, and photonic devices. Thermal transport computations by Pan et al. [277]
suggested that armchair silicene nanoribbons may be used as high performance
thermoelectric materials. Hu et al. [278] showed that silicene with divacancy defects
exhibits high selectivy for H2 over H2O, N2, CO, CO2, and CH4, and can be
potentially used as hydrogen purification membrane. Moreover, since Si
nanomaterials have been largely explored for Li-ion battery anodes due to their high
energy density, silicene sheets should also be a promising anode candidate for Li-ion
batteries, because of its large specific surface area allowing for occupation of more Li
ions and contributing to a higher energy density.
Note that unlike graphene which can form a free-standing sheet, silicene has only
been found to grow on metal substrates. However, to promote silicene into the next
pace for nanoelectronic infrastructure, it is of great importance to grow silicene on
insulator substrates, and these synthesis challenges would encourage more theoretical
and experimental studies in the future.
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3.3 BC3 honeycomb sheets
BC3 is a new graphite-like material with a layered hexagonal structure, which
can be viewed as the substitutional boron doping of graphite. Bulk BC3 was
successfully synthesized by reacting benzene and boron trichloride at about 800°C
[279]. Recently, the uniform BC3 sheets have been prepared in an epitaxial way onto
NbB2(0001) substrate via C substitution in the B honeycomb layer of metal diboride
[280], and the phonon dispersion curves of BC3 sheets have been determined [281].
Compared with pure carbon materials, BC3 has the promising technological benefits
such as high oxidation resistance and superconductivity [282].
Fig. 28. Atomic arrangement of BC3. Reproduced from Ref. [283]. Copyright © 2012,
American Chemical Society.
The most stable ordered structure at the BC3 composition is composed of carbon
hexagons and orderly distributed boron atoms, where six carbons construct a hexagon
and each boron atom is linked to three separated hexagons. Although boron is
electron-deficient as compared with carbon, the charge in BC3 sheet is actually
transferred from boron to carbon. Note that in contrast with graphene where the π
electrons are delocalized over entire carbon lattice, the BC3 lattice displays the
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aromatic carbon rings with six π-electrons located over each carbon hexagon, and are
separated by the anti-aromatic hexagons without π-electrons consisting of carbon and
boron atoms (Fig. 28) [283].
A number of studies have examined the electronic properties, defects, doping,
and surface functionalization of BC3 sheet. BC3 is a non-magnetic semiconductor with
an indirect LDA band gap of 0.54 eV. By inducing defects (vacancy, or antisites), the
electronic and magnetic properties of BC3 sheet can be tailored remarkably [284]. Lin
et al. [285] showed that by substituting the B or C atoms in BC3 lattice with transition
metals (Fe, Co, or Ni), a rich variety of properties ranging from magnetic
semiconductor, nonmagnetic semiconductor, to magnetic metal can emerge. Note that
generation of O-defects in BC3 sheet has been detected by angle-resolved x-ray
absorption near edge structure (XANES) after the sample is subject to Ar+ ion
bombardment and subsequent exposure to air [286]. The O-defects are formed by
firstly creating C vacancies and then substituting the vacancy sites with O atoms.
Moreover, it is possible to tailor electronic structures of BC3 by hydrogenation.
Through step-by-step hydrogenation, a semiconductor→metal transition appears in
hydrogenated BC3 sheet [ 287 ]. Ding et al. [ 288 ] further demonstrated that
half-metallicity is achieved in semihydrogenated BC3.
BC3 nanoribbons also show a rich variety of electronic properties. According to
Ding et al.’s report [289], armchair BC3 nanoribbons are all semiconductors, while
zigzag nanoribbons can be semiconductors or metals depending on the edge
geometries. Specifically, the zigzag nanoribbons with BCBC terminations are
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semiconductors, and those with BCCC or CCCC edge terminations are metallic.
Additionally, Dutta et al. [290] showed that for the case of armchair BC3 nanoribbons,
removal of the passivating hydrogen atoms from the edge B atoms leads to a higher
stability and makes the narrow ribbons become metallic, and then a transition to
semiconductor occurs as the ribbon width increases. Because of the rich electronic
properties, BC3 and its derivatives are potential low-dimensional materials in
nanoelectronics.
The intriguing BC3 material demonstrates promising applications in hydrogen
storage and Li-ion batteries. For example, Zhang and Alavi [291] as well as Sha et al.
[292] predicted that bulk BC3 can be used as potential hydrogen storage candidate.
Compared with the large energy barrier for H2 diffusion into crystallite graphite, bulk
BC3 exhibits higher reactivity towards H2 and leads to reduction of the activation
energy for H2 migration. Within the BC3 lattice, the intercalated H2 molecule
undergoes dissociative chemisorption to form C-H bonds, and the adsorption strength
and H migration barrier are modest, making layered BC3 suitable for reversible
hydrogen storage under near-ambient conditions. Except for bulk BC3, Yang et al.
[293,294] showed that metal atoms (Li, or Ca) can be strongly adsorbed on BC3
monolayer without clustering, and Li-doped and Ca-doped BC3 sheets are the
favorable candidates for hydrogen storage. Besides, BC3 can also be promising as an
anode material for Li-ion batteries. Based on Kuzubov et al.’s computations [295], the
layered BC3 sheets are predicted to be the Li+ intercalation hosts with the Li ion
capacity (corresponding to Li2BC3 stoichiometry) surpassing that of graphite.
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4. Hypothetical planar graphene analogues
4.1 SiC Silagraphene
For Si-C phases, a typical material with considerable interests is SiC. Si and C in
bulk SiC favor sp3 hybridization, and SiC naturally occurs in the form of cubic,
hexagonal and rhombohedral structures. Recently, a planar phase of SiC with
sp2-hybridized feature resembling graphene was theoretically predicted to have high
structural stability, and the electronic properties of this graphene-like material have
been studied [27,296].
The graphene-like SiC sheet consists of alternating Si and C atoms with each Si
atom having three C atoms as its nearest neighbors and vice versa, and the Si-C bond
length is optimized to be 1.79 Å. Due to the Si-C ionicity, monolayer SiC is a
semiconductor with a direct gap of about 2.55 eV (Fig. 29a), which is enlarged to be
3.7 eV after GW correction [297]. The SiC nanoribbons with H-passivated edges
possess interesting behaviors. The armchair SiC nanoribbons are non-magnetic
semiconductors, and the resulting direct band gaps exhibit sawtooth-like oscillation
features (Fig. 29c). The zigzag SiC nanoribbons are magnetic metals except for
thinner ribbons with small direct band gaps [27,296]. The spin-polarized magnetism
comes from the edge Si and C atoms (Fig. 29e). However, the two zigzag edges are
not antiferromagnetically coupled as in zigzag GNRs, but are ferrimagnetically
coupled, where the ferromagnetic chains at the two edges have opposite orientations
of magnetic moments but of different values, and thus the net magnetic moments for
the zigzag systems are not zero. This is ascribed to the unequal values of local
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magnetic moments at the edge C and Si atoms. Interestingly, Sun et al. [27] reported
that zigzag SiC nanoribbons narrower than 4 nm present half-metallic behaviors (Fig.
29d). Based on HSE functional computations, Ding et al. [298] have recently revealed
that asymmetric H-terminations can break the magnetic degeneracy in zigzag SiC
nanoribbons and favor the ferromagnetic state, and meantime, the nanoribbons with
bare Si edges or two-hydrogen saturated Si (C) edges are converted into half metals.
Fig. 29. (a) Geometry and band structure of SiC sheet. (b) Structures of armchair and
zigzag SiC nanoribbons with width W. (c) Band gap of armchair SiC nanoribbons as a
function of ribbon width W. (d) Band structure of zigzag SiC nanoribbons with W=7.
Dotted red lines are spin-up bands, and solid blue lines are spin-down bands. The
projected band structure of 2D SiC sheet is shown by shaded areas. (e) Spin densities
for the W=7 zigzag SiC nanoribbon. Reproduced from Ref. [27]. Copyright © 2008,
American Institute of Physics.
The electronic structures and magnetization of SiC sheets and nanoribbons can
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be further manipulated by other strategies such as transverse electric field [299],
surface modification [300], strains [301], defects [302], and doping [303,304]. For
example, Lou et al. [299] found that the narrow zigzag SiC nanoribbons are turned
into ferromagnetic metals from ferrimagnetic semiconductors when an external
transverse electric field is applied. Substitutional impurities, adatom adsorption and
vacancy defects have been demonstrated to yield controllable functionalization of SiC
honeycomb structure [302]. Moreover, carrier doping, such as hole, electron, or
chemical doping (B or N doping), is used to manipulate magnetism of narrow zigzag
SiC nanoribbons [303,304]. Like graphene, BN or silicene, hydrogen is also an
important surface modifier for SiC. The fully hydrogenated SiC sheet in its stable
chair-like structure is a non-magnetic semiconductor with a direct band gap of 3.84
eV (DFT-GGA) [305], enlarged to be 5.3 eV at the GW approximation [297]. For the
semihydrogenated case, Xu et al. [300] reported the induced FM and anti-FM
properties in SiC when C and Si atoms are hydrogenated, respectively. Surface
hydrogenation also plays a crucial role in tuning the electronic properties of SiC
nanoribbons. Guan et al. [ 306 ] reported that by carefully controlling the
hydrogenation ratio, a transition of the anti-FM semiconductor→FM metal→anti-FM
half metal→NM semiconductor is achieved in zigzag SiC nanoribbons.
Here SiC only represents one example of the graphene-like binary compounds of
group-IV elements. Many other binary phases, such as GeC, SnGe, SiGe, SnSi, and
SnC, have been theoretically studied on their structural, electronic and mechanical
properties [307]. Among their optimized monolayer structures, GeC and SnC form a
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planar honeycomb structure like SiC, whereas SnGe, SiGe, and SnSi prefer a
low-buckled geometry for stabilization like silicene, with the buckled distance in the
range of 0.55~0.73 Å. Electrically, all these binary phases are predicted to be
semiconductors. Depending on the structural composition, their LDA band gaps vary
over a wide scale from 0.02 eV (SiGe) to 2.09 eV (GeC) [307].
4.2 SiC2 Silagraphene
The Si atoms in the above planar SiC sheet are three-fold coordinated. Motivated
by the intensive studies on planar tetracoordinate C (ptC) (known as the anti-van't
Hoff/Lebel compound), planar tetracoordinate Si (ptSi) chemistry has also invoked
tremendous interests for scientists [308]. From a theoretical perspective, Li et al. [29]
have recently predicted an interesting phase of SiC2 silagraphene where ptSi can be
stably embedded into periodic 2D graphitic network.
SiC2 sheet is designed based on a planar SiC4 molecule containing a ptSi (Fig.
30a). SiC4 has C2v symmetry, and is a local minimum at the potential energy surface.
Using planar SiC4 as the building unit, 2D infinite SiC2 sheet in an ideally planar
arrangement is accordingly constructed (Fig. 30b). In SiC2 silagraphene, each Si atom
is bonded with four C atoms to form a ptSi moiety, while each C atom is bonded with
two ptSi atoms and one C atom. The Si-C bond (1.916 Å) is slightly longer than the
typical Si-C single bond (1.87-1.91 Å), while the C-C bond (1.332 Å) is characterized
as a typical C=C double bond. Structurally, one C=C bond and its four neighboring Si
atoms resembles the structure of ethylene. In a reasonable sense, the
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planarity-preferred ethylene-like skeletons act as a driving force to yield a perfectly
planar SiC2 network. The SiC2 silagraphene is a stable phase, as a local minimum at
the potential energy surface, having a large binding energy (6.04 eV/atom) which is
almost in coincidence with that of the SiC silagraphene, and possessing high kinetic
(up to 800 K) and thermodynamic stabilities.
Fig. 30. (a) B3LYP/6-311+ G* optimized structure of SiC4 molecule; the distances
are in angstroms. The number in bracket is computed from MP2 procedure. (b) Top
and side views of SiC2 silagraphene. (c) Band structure and partial DOS of SiC2
silagraphene. Reproduced from Ref. [29]. Copyright © 2010, American Chemical
Society.
The electronic properties of SiC2 related nanomaterials are in stark contrast to
their carbon and SiC analogues. All the ptSi-containing nanomaterials, ranging from
SiC2 silagraphene sheet to nanoribbons, are robustly metallic. Analyzing partial DOS
of SiC2 silagraphene reveals that the states at the Fermi level are mainly contributed
by C-2p and Si-3p states, and the contribution of C-2p is more pronounced than Si-3p,
thus the characteristic metallicity of SiC2 is governed by the π-type C=C bonds (Fig.
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30c).
4.3 Boron sheets
Boron, having one electron fewer than carbon, has always been fascinating the
research community due to its richness of bonding chemistry. The natural crystalline
structure of boron is the α-rhombohedral solid, where the boron atoms form
symmetric 12-atom icosahedral clusters in a crystalline state. Though boron has only
three valence electrons, it can still undergo sp2 hybridization like carbon. The
existence of planar or quasiplanar boron clusters in the 10- to 15-atom range has been
evidenced in recent experiments [309]. However, no evidence of 2D planar structure
exists in its crystals since they are mainly built from B12 icosahedra, and as a
consequence, searching for a planar boron sheet with high structural stability remains
an intriguing but challenging task.
As a matter of fact, researchers have theoretically proposed several models of 2D
boron sheets. In contrast to carbon, the graphene-like hexagonal B monolayer is
unstable and tends to be distorted. This is mainly due to the fact that the
electron-deficient characters of boron tend to promote boron atoms to interact with
each other in forming multicenter bonds (usually forming three-center bonds with
B-B-B structural units). Herein, it shoud be mentioned that although the free-standing
graphene-like boron sheet is unlikely to exist, boron can actually form hexagonal
layered structures in AB2-type materials (such as MgB2, AlB2, BeB2, and TiB2), in
which boron atoms carry a negative charge of 1 and thus acquire configuration
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similar to that of carbon. Similarly, theoretical studies made by Zhang et al. [310]
indicated that the hexagonal geometries can be stabilized on metal surfaces (Mg, Al,
Ti, Au, and Ag). Considering the instability of the standalone hexagonal phase, Evans
et al. [311], Cabria et al. [312], and Kunstmann et al. [313] later independently
proposed that boron sheet constructed by a planar triangular lattice with B7 building
units (filled hexagons with interior boron atoms at the center) is more stable than the
hexagonal boron sheet. However, the six-fold coordination of boron atoms in the
triangular phase is not compatible with the p-orbital symmetries, and the flat plane is
also energetically unstable with respect to buckling which breaks the triangular
symmetry (constructed as hexagonal pyramidal B7 units). Hence in the cases of both
hexagonal and triangular sheets, distortions provide the stability of 2D boron sheet.
Later, after Szwacki et al. [314] reported an unusually stable and highly
symmetric B80 fullerene in 2007, Tang et al. [315] immediately proposed a novel
boron monolayer comprised of a hybrid of triangular and hexagonal motifs, known as
α sheet. The α-sheet has been predicted to be energetically more stable than previous
triangular or hexagonal boron sheet (Fig. 31) [316]. This sheet is ground-state planar
and preserves the symmetry of the triangular lattice, and the binding energy is ~93%
of that of bulk α-rhombohedral solid. In α boron sheet, about 2/3 of the hexagons are
occupied by additional boron atoms, whereby 3/4 of the boron atoms have five nearest
neighbors and 1/4 of the boron atoms have six nearest neighbors, and the hollow
hexagons are evenly and symmetrically distributed within the lattices. The stability of
α-sheet is explained through balance between two-center bonding in hexagonal
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regions and three-center bonding in triangular regions. Specifically, the hexagonal
sheet is an electron-deficient system with part of the in-plane sp2-bonding states
unoccupied, thus acts as an electron-acceptor. While the triangular sheet has a surplus
of electrons in anti-bonding states, thus acts as an electron-donor. After forming the
mixed hexagon-triangle lattice, the electrons provided by the triangular regions
occupy those empty bonding states, thereby markedly enhancing the structural
stability.
Fig. 31. Models of various 2D B sheets. The stability order is: α-sheet > β-sheet >
buckled triangular sheet > perfect triangular sheet > distorted hexagonal sheet >
perfect hexagonal sheet. Reproduced from Ref. [316]. Copyright © 2008, American
Chemical Society.
As further additions to α-sheet, other B sheets based on hybrid hexagonal and
triangular motifs have been proposed (Fig. 32), and some of them are comparable or
even more stable than the known α-sheet. For example, Tang et al. [317] revealed that
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for single-layered B sheets made of hexagons and triangles, their ground-state
configurations may be either buckled or flat depending on the hexagon-to-triangle
ratio, or the “hexagon hole density” η. They showed that flat B sheets are located at
the range of 1/9 < η < 1/5, while otherwise the sheets stay buckled. Note that the
α-sheet has a η value of 1/9. According to Lau et al.’s opinion [316], the α-sheet is
more stable than the planar β-sheet (Fig. 31), which has η value of 1/8. Zope et al.
[318] designed a planar snub B sheet with a hexagon hole density of 1/7, and the
triangular regions in snub sheet form extended zigzag stripes. The proposed snub B
sheet is about 0.02 eV/atom less stable than the α-sheet. Besides, Özdoğan et al. [319]
proposed a γ-sheet, which is decorated similarly to the α-sheet but with parallel
hexagonal holes, yet the α-sheet is still comparatively more stable. Furthermore,
Penev et al. [320] discovered another two ground-state structures of planar B layers
that are as stable as the α-sheet but have higher hexagon density (1/8 and 2/15). More
recently, Yu et al. [321] have predicted a meta-stable B sheet (struc-1/8 B sheet),
which is 0.01 eV/atom less stable than α-sheet or the 1/8- and 2/15-B sheet.
Especially, Wu et al. [322] found that compared with the originally planar α-sheet, the
slightly buckled α-sheet (named α'-sheet) becomes energetically more stable;
moreover, their search yields two highly stable B monolayers, the planar α1-sheet and
β1-sheet, respectively, both with a hole density of η = 1/8. The α1-sheet and β1-sheet
are predicted to possess greater cohesive energies (by more than 60 meV/atom) than
the α-sheet or the 1/8- and 2/15-B layers. These proposed B monolayers with
comparable stabilities (particularly α-sheet, 1/8-sheet, 2/15-sheet, buckled α'-sheet,
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α1-sheet, and β1-sheet) might constitute the leading candidates for the low-energy
configurations of 2D B sheets.
Fig. 32. Various structures of B monolayers based on hybrid hexagon-triangle motifs:
snub B sheet (Reproduced from Ref. [318]. Copyright © 2010, Elsevier.), γ-B sheet
(Reproduced from Ref. [319]. Copyright © 2010, American Chemical Society), 1/8-
or 2/15-B sheet (Reproduced from Ref. [320]. Copyright © 2012, American Chemical
Society.), struc-1/8 B sheet (Reproduced from Ref. [321]. Copyright © 2012,
American Chemical Society.), as well as α , α1, and β1-B sheet (Reproduced from Ref.
[322]. Copyright © 2012, American Chemical Society.). The red and yellow balls in
α -B sheet denote boron atoms moving outward or inward from the plane.
In terms of the electronic structure, nearly all the 2D boron monolayers are
rigorously metallic, with the exception of α- and α'-sheets. It is noteworthy that
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although α- and α'-sheets are predicted to be metallic based on PBE functionals, the
hybrid PBE0 functionals suggest that both of them are semiconducting with an
indirect gap of 1.4 and 1.1 eV, respectively [322].
The electronic properties of α-B nanoribbons and potential applications of α-B
sheet have been studied. The α-B nanoribbons with line or zigzag edges are
electrically metallic; however, the presence of semiconducting behaviors is observed
in the two-hydrogen saturated armchair and zigzag boron nanoribbons [ 323 ].
Moreover, due to the unique geometry, the hexagonal regions of the α-B sheet can
interact strongly with alkali metal atoms (Li, Na, or K), and the doped α-B sheet is
predicted to be a promising hydrogen storage material [324].
4.4 B2C sheet
B2C sheet is a new B-C phase, proposed by Wu et al. [325], where each C atom
is bonded with four B atoms, forming tetra-coordination C moiety. The B2C sheet is
composed of closely-packed hexagons and rhombi, and its structural design is
essentially the 2D extended network of a predicted ptC molecule CB4 (Fig. 33a). The
B2C sheet has a C2-rotation symmetry and two reflection symmetries with respect to
the vector a and b, respectively. The optimized lattice constants are a = 2.558 Å and b
= 3.453 Å, and the B-C and B-B bond length is 1.557 and 1.685 Å, respectively (Fig.
33b). Note that not forming a completely planar structure, B2C sheet is slightly
corrugated, where B-layer and C-layer are separated by a tiny distance of ~0.085 Å
(Fig. 33c). From the deformation and total electronic density (Fig. 33d and e), the C
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atoms form multi-center electron-deficient covalent bonds with four B atoms, and
meantime, the two-center B-B bonds form between two neighboring motifs.
Fig. 33. (a) CB4 molecule (left) and a C2V-CB4 motif (right). (b) Top and (c) side view
of B2C monolayer. (d) Deformation electronic density of B2C. Blue and yellow region
refers to electron rich and deficient region, respectively. (e) Total electronic density
projected on B2C surface. Reproduced from Ref. [325]. Copyright © 2009, American
Chemical Society.
B2C has many appealing properties. For example, the B2C network is predicted
to have excellent mechanical and chemical stability due to strong covalent bonds. The
elastic strength of B2C is found to lie between those of graphene and boron monolayer.
Electronically, B2C sheet and derived nanoribbons are predicted to be uniformly
metallic. Based on first-principles lattice dynamics and electron-phonon coupling
computations, Dai et al. [326] showed that B2C sheet is a 2D phonon-mediated
superconductor with a relatively high transition temperature (19.2 K). Similar to BC3,
B2C sheet can also hold metal atoms strongly, and the Li-doped [327] B2C sheet is a
promising system to storage hydrogen.
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5. Non-planar materials
5.1 Metal dichalcogenides
Transition metal dichalcogenides (TMDs), in a general formula MX2 (M = Mo,
W, V, Nb, Ta, Ti, Zr, Hf, and X = S, Se, Te), represent an intriguing family of
materials with prospects for a broad range of unique properties and applications [328].
TMD materials have layered structures consisting of stacks of X-M-X sandwiches
held together by vdW force. Depending on the different compositions, TMDs can
have hexagonal or rhombohedral symmetry, and the central metal atoms can form
octahedral or trigonal prismatic coordination configuration. Due to their weak
interlayer interactions, isolated sheets of TMDs can be cleaved along the layer plane
similar to the case of graphite, and these mono- or few-layer structures are regarded as
non-planar graphene analogues.
Among the large group of layered TMDs, MoS2 and WS2 are the two most
widely studied materials in recent years, and hence in the subsequent section, we
mainly summarize recent advances in experimental fabrications, fundamental
properties, and potential applications of MoS2 and WS2 thin nanosheets/ribbons.
5.1.1 MoS2 and WS2
MoS2 and WS2 exhibit similar structural features. They are both characterized as
hexagonal layered configurations in which the atoms in the layer are bonded with
strong covalent bonding, while the layers are packed together with weak interlayer
forces like graphite and h-BN (Fig. 34). Within MoS2 or WS2 monolayer, each Mo or
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W atom is coordinated to six S atoms, and each S atom is coordinated to three Mo or
W atoms, forming a trigonal prismatic coordination configuration.
Fig. 34. Layered structure of MoS2, and its single layer has a thickness of 6.5 Å.
Reproduced from Ref. [328]. Copyright © 2011, Nature Publishing Group.
Experimentally, thin nanosheets of MoS2 and WS2 can be prepared through
many techniques. For example, because of the weakly-bonded layers, single and few
layers of MoS2 and WS2 can be extracted by using top-down techniques such as
micromechanical cleavage, Li-ion intercalation and exfoliation, and liquid-phase
exfoliation (see Sections 2.1 and 2.2). Surface ripples are observed in mechanically
exfoliated single-layer MoS2, like the case of graphene, which reflects similar stability
mechanism of both 2D materials and can explain the degradation of electron mobility
of MoS2 due to exfoliation [329]. Alternatively, using bottom-up technologies, such
as CVD growth on insulating substrates, hydrothermal reaction at high temperature,
and thermal evaporation technique, high-quality MoS2 and WS2 flakes with hundreds
of nanometers to a few microns in size can be generated (see Sections 2.3 and 2.6).
These developed methods also hold great promise for preparing atomically thin films
of other TMDs (such as VS2, TiS2, NbSe2, WSe2, MoSe2, and MoTe2), and other types
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of metal chalcogenides (such as GaS, GaSe, Bi2Se3, and Bi2Te3), which has been
discussed in Sections 2.1-2.6.
5.1.1.1 Electronic properties
MoS2 and WS2 sheets. Unlike the semimetallic characteristic of graphene, both
MoS2 [330] and WS2 [331] sheets have intrinsic band gaps, and might serve as an
important complement to graphene in the field of semiconducting applications.
Theoretical studies on the lattice dynamics [332], electronic structures [333], and
dielectric properties [334] of MoS2 have been reported.
The electronic properties of MoS2 are highly associated with their thickness
owing to perpendicular quantum confinement effect (Fig. 35). Bulk MoS2 is an
indirect-band-gap semiconductor (~1.2 eV), and recent first-principles computations
have revealed that intercalation of K atoms into interstitial sites of bulk MoS2 can lead
to a metallic characteristic as a result of electron donation from K 4s orbital to the
conduction band of MoS2 [335]. Bilayer MoS2 is also an indirect semiconductor with
a gap of ~1.6 eV. However, single-layer MoS2 is a direct-band-gap semiconductor
with a larger gap (~1.9 eV) [336,337] and shows a strong photoluminesence absent in
the bulk form [338,339]. The confinement induced indirect-to-direct gap transition in
the single-layer limit also appears in WS2 [331] and other TMDs such as MoSe2,
MoTe2 [340], and WSe2 [341]. The indirect-to-direct gap crossover can also be
realized in few-layered MoSe2 by thermal excitement [342].
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Fig. 35. (a) Band structures of bulk and monolayer MoS2. (b) Band structures of bulk
and monolayer WS2. Reproduced from Ref. [337]. Copyright © 2011, American
Physical Society.
The presence of a direct band gap in single-layer MoS2 makes it promising for
integration of MoS2-based microelectronics into photoelectronic devices which allow
for efficient electron-phonon energy conversion. More importantly, the intrinsically
moderate band gap of single-layer MoS2 in the visible frequency range overcomes the
drawback of either zero band gap of graphene or wide band gap of h-BN, implying its
prospective semiconducting applications in complementing graphene or BN as
next-generation nanoelectronics and photoelectronics.
Also, due to its unique electronic structure, MoS2 monolayer has great potential
for valleytronics and spintronics applications. MoS2 monolayer is a two-valley
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direct-gap semiconductor where the conduction band and valence band edges are both
located at the K point. Note that the two inequivalent valleys are separated in the
Brillouin zone by a large momentum, the valley index is thus robust against
intervalley scattering, and consequently can be used as a potential information carrier.
Monolayer MoS2 is known to lack inversion symmetry, which can lead to valley Hall
effect and valley-dependent optical selection rules for interband transitions at the
k-points. In addition, compared with the nearly vanishing spin-orbit-coupling in
graphene, MoS2 has a stronger spin-orbit-coupling interaction originated from the d
orbitals of the heavy metal atoms, which results in a strong spin-orbital-induced band
splitting. The inversion symmetry breaking along with the strong spin-orbit
interaction leads to coupled spin and valley physics in MoS2 monolayer [343].
Actually, experimental evidence for the valley polarization of MoS2 monolayer has
been achieved by photoexciting the samples with circularly polarized light, as has
been recently reported by three experimental groups [344-346], and this might lead to
development of potential “valleytronics” devices based on MoS2 monolayer.
Moreover, due to strong in-plane Mo-S covalent bonds, MoS2 nanosheets have
good mechanical strengths. The high Young’s modulus obtained for suspended
ultrathin MoS2 nanosheets with thickness of 5-25 layers (E = 0.33 ± 0.07 TPa) are
comparable to that of graphene oxide [347], and the in-plane stiffness (180 ± 60 Nm-1)
and effective Young’s modulus (270 ± 100 GPa) of single-layer MoS2 are estimated to
be comparable to those of steel [348]. These excellent mechanical properties make
MoS2 nanosheets attractive as reinforcing elements in composites or for fabrication of
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flexible electronic devices.
Note that the electronic structures of MoS2 sheets can not only be tuned by
varying the number of layers, but also can be modulated by application of mechanical
strain, electric field, surface adsorption, and defects. For example, Scalise et al. [349]
showed that the band gaps of monolayer and bilayer MoS2 decrease gradually upon
application of bi-axial strain, and finally change to be metallic. Similarly,
semiconductor-to-metal transition also appears for monolayer and bilayer MoSe2,
MoTe2, WS2, and WSe2 under tensile and shear strains [350], or bilayer MoS2, MoSe2,
MoTe2, and WS2 under perpendicular electric field [351]. For the adatom adsorption,
He et al. [352] showed that H, B, C, N, O, and F atoms chemically adsorb on MoS2
monolayer, and these atoms significantly modify the electronic structures into n-type
or p-type semiconductors and induce magnetism in MoS2. The MoS2 sheets can also
obtain local magnetism through adsorption of 3d transition metals (such as Co, Cr, Fe,
Mn, Mo, Sc, Ti, and V), as well as Si and Ge [353].
Atomic defects such as vacancy, dislocation, and doping also serve as significant
modifications for MoS2. Like the defect formation in graphene and BN, atomic
defects in TMDs can be generated by electron irradiation. Specifically, Liu et al. [354]
observed the W and S monovacancies in monolayer WS2 nanoribbons encapsulated
inside carbon nanotubes at an electron acceleration voltage of 60 kV. The vacancies
were created mainly at the ribbon edges, and the W vacancy was found to induce a
drastic structural deformation. Moreover, Hansen et al. [355] observed the atomic Mo
(S) defects at the edges of MoS2 clusters under an 80 kV electron irradiation. Komsa
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et al. [356] have recently reported the vacancy formation in MoS2 sheets under an 80
keV electron beam. The HRTEM image in Fig. 36a clearly indicates formation of a
number of S vacancies accompanied by the crack formation and lateral shrinkage of
the sheet. The simulated TEM images of single and double S vacancies (Fig. 36 d and
e) are in consistency with the observed vacancy structures in measured TEM images
(Fig. 36f). Atomic vacancies can be further filled by other impurity species, forming
electron-beam-mediated substitutional doping defects.
Fig. 36. (a) HRTEM image of single-layer MoS2. Models and simulated HRTEM
images for (b, d) single and (c, e) double vacancy, respectively. Experimental TEMs
are shown in (f). Reproduced from Ref. [356]. Copyright © 2012, American Physical
Society.
Theoretical studies indicated that creation of monovacancy (Mo or S) or
divacancy (2S or MoS) defect in MoS2 monolayer does not induce any magnetic
moment, while the construction of MoS2 triple vacancies gives rise to significant
magnetic moments [353]. The MoS2 triple vacancy in MoS2 is predicted to be able to
attract H2O molecule and catalyze the dissociation of H2O into its constituent H and O
atoms [357].
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In a very recent study, Zou et al. [358] have investigated the dislocations and
grain boundaries in single layers MoS2 and WS2. Their simulations indicated that the
dislocation cores in MoS2 and WS2 form concave polyhedras (5|7, 6|8, and 4|6
polygons) with dreidel-like shapes. The linear arrays of these dislocations then form
grain boundaries between the mutually tilted crystalline domains. The presence of
grain boundaries introduces midgap electronic states within the band gap of MoS2 or
WS2 and acts as sinks for the carriers.
As for doping defects, most theoretical investigations focused on mixed TMD
phases with either metal or chalcogen mixing. Ivanovskaya et al. [359] studied the
bilateral doping within the mixed MoS2-NbS2 system, and found that substitutional
Nb doping in MoS2 leads to a semiconductor-metal transition, while Mo doping in
NbS2 does not alter the metallic character of NbS2 system. On the other hand, Komsa
et al. [360] studied the single layers of MoS2xSe2(1−x) alloy. The lowest-energy
configurations of MoS2xSe2(1−x) favor having dissimilar atoms in the nearest neighbor
sites in the chalcogen sublattice. The band structures of the alloy systems resemble
their binary MoS2 or MoSe2 constitutes, and their direct gaps (1.65-2.0 eV) can be
tuned between the constituent limits. The mixed TMD materials can be prepared by
CVD growth or be exfoliated from bulk mixed phases, and their tunable properties
should be beneficial for optoelectronic applications.
MoS2 nanoribbons. Recently, MoS2 derived nanoribbons have also fascinated the
research community. Li et al. [28] computationally studied MoS2 nanoribbons with
widths up to ~6 nm (Fig. 37). Armchair MoS2 nanoribbons are non-magnetic and
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semiconducting, and their direct band gaps converge to nearly ~0.56 eV as the width
increases (Fig. 37d). Due to edge effect, the band gaps of armchair MoS2 nanoribbons
are smaller than that of the 2D single-layer MoS2. In comparison, zigzag MoS2
nanoribbons are ferromagnetic metals. The induced magnetism mainly concentrates
on the edge Mo and S atoms (Fig. 37c), and the total magnetic moments increase with
increasing ribbon width or thickness. Note that ferromagnetism has been detected in
flat MoS2 clusters and MoS2 sheets, which might be partly due to the presence of
zigzag edges at the grain boundaries of MoS2 crystals or the presence of lattice defects
[361 ,362]. Zigzag MoS2 nanoribbons are energetically more feasible than the
armchair counterparts with comparable widths, while the case in graphene
nanoribbons is opposite. Many other theoretical investigations have also deepened the
understanding of MoS2 nanoribbons, and similar to MoS2 sheets, the properties of
MoS2 nanoribbons can be modulated by edge functionalization (such as H, and S),
adatom adsorption, and vacancy defects, or by applying strains or electric fields
[363-365].
Fig. 37. Geometries of 8-zigzag (a) and 15-armchair (b) MoS2 nanoribbon. The spin
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density distribution of 8-zigzag MoS2 nanoribbon (c), and the energy gap variation of
armchair nanoribbons as a function of ribbon width (d) (8 ≤ Na ≤ 20, inset).
Reproduced from Ref. [28]. Copyright © 2008, American Chemical Society.
Ultra-narrow MoS2 and WS2 nanoribbons with zigzag-shaped edges have been
synthesized by Wang et al. via the bottom-up chemical growths inside carbon
nanotubes (see Section 2.5.2) [21,22]. For these zigzag-edged MoS2 nanoribbons, the
Mo edge tends to be half-saturated by S atoms, while the S edge prefers to be bare.
5.1.1.2 Potential applications
Transistors. The semiconducting ultra-thin MoS2 layers and other TMDs,
because of their intrinsic advantages such as high structural stability, tunable band
gaps between 1-2 eV which can lead to high on/off ratios, subnanometer thickness,
and high carrier mobility, are attractive for use as channel materiasl in low-power
FETs. The TMD-based FETs generally possess a high on/off ratio of larger than 106,
and the carrier mobility of around 200-500 cm2/(Vs) is comparable to the
single-crystal Si [366]. As the first attempt of integrating the mechanically exfoliated
single-layer MoS2 into FET electronics, Radisavljevic et al. [328] fabricated a
MoS2-based n-type transistors using HfO2 as the gate dielectric (Fig. 38). This device
demonstrates a high room-temperature channel mobility of more than 200 cm2/(Vs)
and current on/off ratios of up to 1×108, comparable to the mobility achieved in thin
Si films or GNRs. A high performance p-type FET based on single layered WSe2 with
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hole-doped source/drain contacts and high-κ top-gate dielectric (ZrO2) has also been
fabricated, which exhibited a hole mobility of ~250 cm2/ (V s), subthreshold swing of
~60 mV/dec, and on/off ratio of >106 [367]. Note that the coating of high-κ dielectric
plays a significant role in improving the carrier mobility of semiconducting channel
materials due to the dielectric engineering.
Fig. 38. Schematic of HfO2-top-gated monolayer MoS2 FET device. Reproduced from
Ref. [328]. Copyright © 2011, Nature Publishing Group.
Many other groups also reported the high performance of FETs made of MoS2
films based on the mechanical exfoliation technique [368-370]. For example, Zhang
et al. [371] fabricated an electric double layer transistor (EDLT, a FET gated by ionic
liquids) using thin MoS2 flakes as channel materials and demonstrated that this device
displays ambipolar transistor operation. Braga et al. [372] realized an ambipolar FET
by coupling exfoliated thin flakes of WS2 with an ionic liquid dielectric. Kim et al.
[373] showed high-mobility and low-power thin-film FETs based on the multilayer
MoS2 crystals. MoS2 FET devices are sensitive to environmental moisture, and
exposure to ambient oxygen or water can drastically reduce the device performances
[374].
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The MoS2/WS2 layers prepared from liquid exfoliation or CVD growths have
also been implemented into FET fabrications, and they exhibit similar electrical
characteristics [52]. For example, Hwang et al. [375] reported the realization of
back-gated FETs made of the chemically exfoliated multilayer WS2, which
demonstrate a high (~105×) on/off current ratio, and show clear photo response to
visible light. Pu et al. [376] fabricated a CVD-grown MoS2 thin-film transistor using
ion gel (gelation of an ionic liquid) as gate dielectrics, and this device is characterized
by a high on/off ratio of 105 and a mobility of 12.5 cm2/(V·s) at a low operating
voltage of 0.68 V.
Moreover, the large in-plane carrier mobility of MoS2 also makes it attractive for
fabrication of other digital devices such as phototransistor, integrated circuits, and
signal amplifier. For instance, Yin et al. [377] fabricated a phototransistor based on
MoS2 single layer, which exhibits better phtoresponsivity (7.5 mA/W) than the
graphene-based device (1 mA/W). Lee et al. [378] fabricated top-gate phtotransistors
based on single-, double-, and triple-layer MoS2 sheets, and their devices with triple
MoS2 layers exhibited excellent photodetection capabilities for red light, while those
with single- or double-layers were useful for green light detection. The first
implementation of single-layer MoS2 into the integrated circuits was reported by
Radisavljevic et al. [379], and this device can be used to operate as inverters. Wang et
al. [380] also fabricated the fully integrated multistage circuits based on bilayer MoS2
transistors, such as a logic inverter, a static random access memory, and a five-stage
ring oscillator.
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Sensors. MoS2 nanosheets are also favored in electronic sensing applications.
The EFTs made from MoS2 sheets have been used to detect adsorption of NO gases
[370]. The sensing mechanism is ascribed to the p-type doping effect of the adsorbed
NO molecules, which changes the electrical resistance of the originally n-type doped
MoS2. Likewise, a flexible thin-film transistor using MoS2 as the active channel and
reduced graphene oxide (rGO) film as the drain and source electrodes is used as a gas
sensor for NO2 detection [ 381 ]. Furthermore, the electrochemically reduced
single-layer MoS2 nanosheets have been demonstrated as sensitive glucose and
dopamine detectors [382].
Energy harversting. MoS2 and WS2 nanosheets can be additionally applied to
energy haversting field. One hand, the laterally confined layered structures of MoS2
and WS2 nanosheets favor their use as active electrode materials in Li-ion batteries.
For example, Seo et al. [115] prepared WS2 nanosheets from W18O49 precursor, and
showed that the electrochemical lithiation capacity of WS2 nanosheets was greatly
improved as compared with the bulk WS2, owing to their enhanced surface area and
easier diffusion accessibility of Li ions. Xiao et al. [383] prepared exfoliated
MoS2/polyethylene oxide (MoS2/PEO) nanocomposites and demonstrated high
capacity (1000 mAh g-1) and good cycle stability for the composite. Moreover, by
combining semiconducting MoS2 with conductive carbon materials (graphene, carbon
nanotubes, or amorphous carbon), better electrochemical performances are expected
[384,385]. Zigzag MoS2 nanoribbons were computationally predicted to effectively
bind the Li ions close to the edges with high Li mobility, and can be alternatively
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developed as cathode materials for Li-ion batteries [386]. Besides, MoS2 thin layers
are also promising cathode materials for rechargeable Mg batteries [114].
On the other hand, the tunable band gaps of MoS2 and WS2 in the visible regions
make them potential candidates for photovoltaic solar cell applications. A hybrid bulk
heterojunction (BHJ) solar photovoltaic cell employing MoS2/TiO2 nanocomposites
and poly 3-hexylthiophene (P3HT) as the active layers is demonstrated to have a
photoconversion efficiency of 1.3% [387]. A Schottky-barrier solar cell made from
MoS2 nanomembranes, indium tin oxide (ITO), and Au with a stacked structure of
ITO-MoS2-Au has been recently demonstrated to possess a photoconversion
efficiency of 1~2% [388].
5.1.2 Other layered metal dichalcogenides
MoS2 and WS2 only represent two examples of the broad family of layered
TMDs. Aside from them, many other metal chalcogenides have also earned
tremendous attention, such as MoSe2 [389], WSe2 [390], VS2, VSe2 [391], NbS2,
NbSe2 [392], SnS2 [393], Bi2Se3 [394,395], and Bi2Te3 [396]. For example, WSe2
[390] would undergo an isostructural semiconductor-semimetal phase transition under
high pressure. By applying a tensile strain, a ferromagnetic behavior can be
theoretically induced in VS(Se)2 [391] and NbS(Se)2 [392] monolayers. Sun et al.
[393] synthesized SnS2 single layers through liquid exfoliation, and these single layers
have superior photocatalytic performances to the bulk phase in achieving effective
visible-light water splitting. Particularly, Bi2Se3 and its related materials such as
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Bi2Te3, and Sb2Te3 were identified as a new class of topological insulators [394],
which were proposed for the realization of dissipation-less interconnects, low-power
electronics and quantum computing devices, as well as for thermoelectric
applications.
5.2 Layered oxide and hydroxide nanosheets
Layered oxides and hydroxides typically have strong interlayer interactions
(mostly electrostatic interactions), and their nanosheets are usually synthesized by
chemical exfoliation of the parent layered crystals in exfoliating solutions. So far, a
wide variety of oxides (such as Cs0.7Ti1.825O4, Ca2Nb3O10, K0.45MnO2, K4Nb6O17,
RbTaO3, KTiNbO5, and Cs6+xW11O36) and hydroxides (such as
M2+1−xM3+
x(OH)2·An−x/n mH2O, where M2+ = Mg2+, Fe2+, Co2+, Ni2+, Zn2+, etc., and
M3+ = Al3+, Fe3+, Co3+; RE(OH)2.5 xH2O·An−0.5/n, where RE = Nd, Sm, Eu, Gd, Tb, Dy,
Ho, Er, Tm, Yb, Lu, Y, etc.) have been delaminated into their charge-bearing 2D
nanosheets (see the recent review [397]). The electrostatic repulsion between the
negatively charged oxide nanosheets or the positively charged hydroxide nanosheets
leads to the formation of highly stable aqueous colloidal solutions.
Generally, the exfoliation of most layered metal oxides is achieved firstly by the
protonation reaction of pristine oxides through acid treatment (such as HCl), and
followed by the intercalation of organic ions into the protonated oxides via the
addition of chemical intercalators. The most commonly used intercalator is
tetrabutylammonium ion (TBA+), although tetramethylammonium ion (TMA+) and
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ethylammonium ion have also been successfully used. The intercalation process of
organic ions (such as TBA+) was mediated via ion-exchange reactions (such as
H+-TBA+), which expands the host structures and decreases the electrostatic
interactions between host layers, assisting the exfoliation of layered materials. In a
similar way, layered metal hydroxides can be exfoliated through insertion of bulky
intercalated agents (such as NO3−, ClO4
−, and dodecyl sulfate (DS) ions) and
subsequent treatment with organic solvents (such as butanol, and formamide) under
stirring, heating, or sonicating conditions. As an example, Fig. 39 presents the
schematic illustration for the exfoliation of DS ion intercalated nickel hydroxide [398].
In this process, exfoliation was induced by intercalation of DS ions into Ni(OH)2
under the reflux of formamide solution. AFM images show homogeneous thickness of
1.1 nm for the exfoliated hexagonal layers, larger than that estimated from
crystallographic thickness, which is likely due to surface absorption of water and DS
ions.
Fig. 39. Schematic for exfoliation of DS ion intercalated nickel hydroxide and AFM
image of the isolated hexagonal layers. Reproduced from Ref. [398]. Copyright ©
2008, American Chemical Society.
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2D oxide and hydroxide nanosheets exhibit diverse properties depending on their
compositions, and demonstrate many promising applications. Most oxide nanosheets
are wide band-gap semiconductors, thereby facilitating their potential use as
semiconducting hosts, photocatalysts, dielectric materials, etc. For example, titanium
oxide nanosheets present excellent photochemical properties [399], and the Co
substituted titanium oxide nanosheets can achieve room-temperature ferromagnetism
[400]. Nanosheets of MnO20.45- [401] and RuO2.1
0.2- [402] undergo electrochemical
redox reactions and are useful for electrochemical energy storage. The Cs4W11O362-
nanosheets exhibited photochromic properties, showing reversible color change upon
UV irradiation [403]. In the case of double hydroxide nanosheets, due to their unique
anion-exchangeability and biocompatibility, they are useful candidates as anion
exchangers to remove toxic anions or as drug and gene delivery nanocarriers for
pharmaceutical and biological applications [404]. Moreover, hydroxide nanosheets
can be used as catalysts or catalyst supporters. For example, the tungstate-exchanged
layered double hydroxide is a promising biomimetic catalyst for oxidative
bromination [405]. In addition, as in the case of oxide nanosheets, hydroxide
nanosheets also exhibit versatile electronic, magnetic, and optical properties, and most
of them have ferromagnetic and half-metallic characteristics. For example, Co2+-Al3+
hydroxide nanosheets act as nanoscale ferromagnetic layers and exhibit an interesting
magneto-optical response in the UV-visible region [406]. Furthermore, nanosheets of
europium hydroxide exhibit characteristic red emission based on the Eu3+
photoactivator, which suggests potential applications of ultra-thin films of this
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material as optical devices [407].
In addition to the layered oxides and hydroxides with strong electrostatic
attractions between the host layers and the cationic (for oxides) or anionic (for
hydroxides) interlayer species (as discussed above), we will consider another
exceptional layerd oxide with weakly bonded layers, V2O5, in the subsequent part.
Fig. 40. Structure model of bulk V2O5.
Bulk V2O5 has an orthorhombic symmetry, and is a typical oxide
semiconductor with a layered structure stacking along the (001) direction,
characterized as strong in-plane V-O covalent bonding and weak vdW coupling
between adjacent layers (Fig. 40) [408]. The individual V2O5(001) monolayer is
strongly corrugated and built up of distorted VO5 square pyramids, which are
alternatively pointed upward and downward. Three inequivalent O atoms exist on
the monolayer: the outermost single-coordinated O atoms (O1) are doubly bonded to
the V atoms, while the other two types are two- and three-coordinated bridging O
atoms (O2 and O3). Bulk V2O5 has a visible band gap of ~2.4 eV, and the (001)
monolayer has similar physical properties and stability as its bulk crystal and can be
in principle produced by cleaving the weak interlayer interaction. As a matter of fact,
the thin nanosheets of V2O5 have been recently fabricated via liquid exfoliation
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technique in polar solvents [48].
Fig. 41. (a) Schematic of exfoliation of layered bulk V2O5 into {001}-oriented
few-layer V2O5 nanosheets. (b-d) Cathode performance of {001}-oriented few-layer
V2O5 nanosheets and bulk V2O5 in Li-ion batteries: (b) initial galvanostatic
charge–discharge voltage profiles at a current density of 59 mA g-1 (0.2 C). Here, 1 C
= 294 mA g-1; (c) cycling performance at 0.2 C; (d) rate capability at various charge
and discharge rates. (e) The Ragone plot of V2O5 nanosheets for power and energy
performance, in comparison with some advanced energy storage and conversion
devices. Reproduced from Ref. [48]. Copyright © 2013, Royal Society of Chemistry.
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The high surface-bulk ratios and the absence of interlayer interactions offer 2D
V2O5 nanosheets unique electronic properties and important functional applications.
According to DFT computations, the single-layer V2O5 sheets and nanoribbons
possess tunable electronic and magnetic properties that are susceptible to surface
hydrogenation [409]. Du et al. [410] pointed out that V2O5 monolayer could
theoretically promote the dehydriding kinetics in magnesium hydride. Recently, Rui
et al. [48] have successfully prepared few-layer V2O5 nanosheets with thickness of
3.1-3.8 nm through direct exfoliation of bulk V2O5 in formamide (Fig. 41). Because
of the short Li diffusion paths provided by the ultrathin thickness, the obtained V2O5
nanosheets demonstrated excellent performances as cathode materials in Li-ion
batteries. As shown in Fig. 41b-d, the ultrathin V2O5 nanosheets display larger
reversible Li capacity, higher Coulombic efficiency, and better rate capability than the
bulk V2O5 electrode. Moreover, the V2O5 nanosheet electrode achieves a higher
energy density (158 W h/kg at a power rate of 20 kW/kg) than the pseudocapacitors
(1–70 W h/kg at a power rate of 1-20 kW/kg) (Fig. 41e), indicating the possible
applications of V2O5 nanosheets as superior electrochemical energy-storage devices
with both high-power and high-energy densities.
5.3 Graphitic-like phase of ZnO
Most of the group II-VI and III-V binary compounds such as ZnO, AlN, and
GaN are naturally crystallized in a stable wurtzite structure with a diamond-like
configuration (Fig. 42 a). However, first-principles computations [411,412] predicted
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that the polar (0001)-oriented ultrathin films of wurtzite materials (such as AlN, BeO,
GaN, ZnO, and ZnS) with only a few atomic layers favor a new form of stable
graphitic structures, where each cation and anion adopt a sp2 hybridized state and are
arranged in planar three-fold coordination instead of the bulk-like tetrahedral
configuration (Fig. 42b,c). This interesting structural transformation is driven by
surface charge transfer mechanism, and the surface polarity is completely removed
after forming the graphitic-like structure. Unlike the metallic characteristic of polar
(0001) wurtzite films, the depolarized graphitic structures are semiconducting instead.
Fig. 42. (a) Wurtzite structure of bulk ZnO. Zn atoms are depicted in light blue, O in
red. Reproduced from Ref. [411]. Copyright © 2005, Royal Society of Chemistry. (b)
Polar wurtzite structure of ZnO(0001) film. (c) Graphitic structure. Reproduced from
Ref. [412]. Copyright © 2006, American Physical Society.
Shortly after the appealing theoretical prediction, the presence of graphitic phase
was confirmed by an epitaxial-growth experiment of ultrathin ZnO films. Tusche et al.
[413] observed two-monolayer-thick non-polar ZnO(0001) films on Ag substrates,
which provided a direct evidence for the presence of planar ZnO sheets (Fig. 43). The
observed depolarization of two-monolayer ZnO film is accompanied by a lateral 1.6%
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expansion of the lattice parameter and a 3% reduced Zn-O bond length within the
sheets. The experimentally observed graphitic ZnO films on Ag substrate can only
persist up to 2-3 layers thickness, and then a transition to the bulk wurtzite takes place,
which is much lower than the theoretical predicted 16 layers. This inconsistency can
be explained by the possible effects such as structure defects, film roughening, and
compressive strain induced by lattice mismatch. Noteworthy, except for graphene-like
ZnO observed in experiments, other types of graphene-like II-VI (ZnS, BeO, or MgO)
and Ш-V (AlN, or GaN) phases still remain hypothetical.
Fig. 43. Models of wurtzite-like ZnO(0001) bilayer structure (a) and the
experimentally determined graphitic-like structure (b). High resolution STM image of
graphite-like ZnO film about 2.2 monolayers in thickness. Reproduced from Ref.
[413]. Copyright © 2007, American Physical Society.
The electronic and magnetic properties of ZnO [414], GaN [415,416], AlN [417],
BeO [418] and MgO [419] in their graphitic hexagonal phases have been intensively
examined, and we mainly discuss the case of ZnO since similar properties are
expected in other types of graphitic phases. The ZnO monolayer is a direct-gap
semiconductor (2.62 eV at the DFT/GGA level, and 3.57 eV under the GW correction
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[420]), and presents ferromagnetic coupling after Co-doping [421]. The creation of
Zn vacancy brings on local magnetic moments, while other types of defects such as O
vacancy, Zn+O double vacancy, and Zn-O antisite defects cause no magnetism for the
system [414]. Doping the O sublattice of ZnO monolayer with C or B atoms results in
a “NM→anti-FM→FM” transition, accompanied by a semiconductor→half metal
transition, and doping by N atoms leads to a p-type semiconductor with the anti-FM
ground state [422]. Moreover, the electronic and magnetic properties of graphene-like
ZnO sheets can be modulated by adatom adsorption of H and F atoms [423,424]. For
example, DFT computations on hydrogenation of graphitic ZnO nanosheets revealed
that the electronic and magnetic properties depend on the adsorption sites of H sites
and the sheet thickness.
For the cutted 1D ribbons of ZnO, the single-layered armchair ZnO nanoribbons
are non-magnetic semiconductors, and their band gaps tend to reduce monotonically
with increasing transverse field strength [425]. The single-layered zigzag ZnO
nanoribbons are metallic and exhibit ferromagnetic responses, which emerge from the
unpaired electrons localized on the oxygen-terminated edges. The edge magnetism of
zigzag ZnO nanoribbons can be modulated by applying an external field [426] (Fig.
44) or edge functionalization with sulfur or thiol groups [427]. In addition, half
metallicity was found in edge-passivated wide ZnO nanoribbons when only Zn edge
is saturated by hydrogen, CH3- or NH2- groups [428]. By applying a proper
axis-strain, one can induce a metal→semiconductor transition in zigzag ZnO
nanoribbons [429]. Interestingly, Botello-Méndez et al. [430] found that in multilayer
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zigzag ZnO nanoribbons, edge atoms on top layer form covalent bonds with edge
atoms on bottom layer. As a result, the magnetism and metallicity disappear in the
even-numbered layer. The metallicity still retains in odd-numbered layer but the
magnetic moment quickly vanishes as the layer number increases to five layers.
Fig. 44. (left) Schematic of a zigzag ZnO nanoribbon under a transverse electric field.
(right) Magnetic moments per supercell of different zigzag ZnO nanoribbons as a
function of applied electric fields. Reproduced from Ref. [426]. Copyright © 2010,
American Chemical Society.
5.4 MXenes
Another family of layered materials, the layered ternary metal carbides, nitrides
or carbonitrides, also known as “MAX” phase, has fascinated the researchers’
interests quiet recently. The MAX phases have a general formula of Mn+1AXn (n=1, 2,
3), where M represents early transition metals (M=Ti, Sr, V, Cr, Ta, Nb, Zr, Mo, or Hf),
A represents main-group sp elements (mostly III A or IV A), and X represents either C
or N as well as both. They correspond to a big family with more than 60 members and
constitute a new class of layered materials combining the properties of both metals
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and ceramics [431].
Structurally, MAX phase can be described as the inter-growth structures with
alternative stacking of hexagonal MX layers and close-packed planar A atomic layers.
As an example, Fig. 45 presents a typical structure of M2AX systems, where the
chemical bonds between A-M elements are weaker than the covalent M-X-M bonds,
which makes it possible to extract A layer from the layered solid. However, different
from graphite with weak vdW interlayer interaction, the MX layers in MAX phase are
held together by A-containing layers via partially ionic bonding, and the consequential
bond strengths are quite strong, making separation of MX layers from MAX phases
cannot be easily realized with mechanical cleavage or direct dispersion and
ultrasonication.
Fig. 45. Typical bulk structure of layered M2AX, where M, A, and X represent
transition metal, A-element, and carbon/nitrogen, respectively.
To isolate elementary MX layers from bulk MAX, one alternative strategy is to
use an effective reactant to selectively etch the interleaved A-containing layers to
decrease the interlayer interactions, but without destroying the layered morphology of
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MX layers. A pioneering work in this line by Gogotsi group is the successful
exfoliation of Ti3C2 nanosheets from Ti3AlC2 by firstly immersing Ti3AlC2 powers
into 50% HF (Ti3AlC2 + 3HF = AlF3 + 3/2H2 + Ti3C2) to extract the Al layer, and
followed by ultra-sonication treatment in methanol (Fig. 46) [432]. TEM images show
that the exfoliated Ti3C2 nanosheets exhibit hexagonal symmetry, with thickness
ranging from single layer and double layers, to multilayers. Such as-designed
procedure by Gogotsi group has been successfully applied to isolate other MX layers,
Ti2C, Ta4C3, TiNbC, (V0.5Cr0.5)3C2, and Ti3CNx (x < 1), starting from their
Al-containing MAX phases [433]. Under the aqueous environment of HF solutions,
the outer surfaces of the exfoliated MX layers are usually chemically-terminated with
F or/and OH functional groups. With the ease of the as-developed method, one can
assume that more and more MX nanosheets could be exfoliated from their bulk MAX
phases.
Fig. 46. (a) Exfoliation process for Ti3AlC2. (b) TEM images of exfoliated MXene
nanosheets. Inset shows selected area electron diffraction (SAED) pattern confirming
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hexagonal symmetry of the planes. (c) TEM images of single- and double-layer
MXene sheets. Reproduced from Ref. [432]. Copyright © 2011, Wiley-VCH.
These isolated 2D carbide and carbonitride nanosheets (Mn+1Xn stoichiometry)
are alternatively termed as “MXene”, as suggested by Gogotsi group, since they have
hexagonal structure and electronic properties comparable to those of graphene.
MXene materials display diversified electronic properties depending on the
appropriate surface treatment. For example, bare Ti3C2 monolayer acts as a magnetic
metal, while its derived Ti3C2F2 and Ti3C2(OH)2 are semiconductors with small band
gaps [434]. Khazaei et al. have recently studied the electronic properties of various
MXene systems, M2C (M = Sc, Ti, V, Cr, Zr, Nb, or Ta) and M2N (M = Ti, Cr, or Zr)
with surface chemically functionalized with F, OH, or O groups [435]. All the bare
MXene monolayers are electrically metallic. After surface functionalization, the
Sc2CF2, Sc2C(OH)2, Sc2CO2, Ti2CO2, Zr2CO2, and Hf2CO2 become semiconductors
with band gaps around 0.25-2.0 eV, and the functionalized Cr2CF2, Cr2C(OH)2,
Cr2NF2, Cr2N(OH)2, and Cr2NO2 become ground-state ferromagnetic.
The metallic or narrow-band-gap semiconducting characteristics endow the
MXene-based materials with intrinsic advantage of good electrical conductivity, and
hence favor their potential application as electrode materials for Li-ion batteries. A
theoretical study on Ti3C2 monolayer shows that Li adsorption forms strong
interaction with Ti3C2 host, and the bare Ti3C2 monolayer exhibits a low barrier (0.07
eV) for Li diffusion and high Li storage capacity (up to Ti3C2Li2 stoichiometry) [434].
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This gives insightful prospect for the application of Ti3C2 nanosheets as an anode
material for Li ion batteries. Surface functionalization by F or OH groups, however,
degrades the Li diffusion and decreases the Li storage capacity, and thus should be
avoided in synthetic experiments. The performance of MXene as an anode material
has also been tested in the experiments. Naguib et al. [436,437] conducted detailed
electrochemical tests to evaluate Ti2C’s performance as an anode material in Li-ion
batteries designed for rapid charging and discharging. The key finding in their study
was that the material retained its charge capacity during 1,000 rapid
charging/discharge cycles.
6. Two-dimensional coordination and covalent organic polymers
Linking molecular building blocks into extended networks by either coordination
bonding or strong covalent bonding is now commonly used in the synthesis of
coordination polymers or covalent organic polymers. These materials, especially their
2D forms, give a novel class of atomically regular and uniform 2D crystals beyond
graphene, and exert great scientific attention for both their synthesis and applications.
6.1 Coordination polymers
Coordination polymers (CPs) are a family of nanoporous materials constructed
by organic linkers and metal centers. Depending on the selective building blocks
(metals and ligands), they exhibit a rich variety of architectures and diverse
physicochemical properties. There are numerous examples of coordination polymers
obtained in crystal phase as lamellar materials. However, 2D coordination networks
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created as ordered monolayers are rarely reported.
Encouragingly, impressive progress has been achieved in isolating
one-atom-thick CP flakes. Amo-Ochoa et al. [15] presented successful exfoliation of
single layer CP extracted from layered crystals of [Cu2Br(IN)2]n (IN = isonicotinato)
using liquid-phase sonication. AFM topography image (Fig. 47b) shows a very dense
and homogenous distribution of [Cu2Br(IN)2]n flakes on highly oriented pyrolitic
graphite (HOPG) substrate. The typical heights (Fig. 47c) of 5 ± 0.15 Å are in good
agreement with the thickness (5.5 Å) expected for single atomic layers. Theoretical
investigations [438] verified that monolayer [Cu2Br(IN)2]n has a strong tendency to
adsorb NO and NO2, which suggests the technological possibility of using
[Cu2Br(IN)2]n CP as molecular sensors.
Fig. 47. (a) Schematics of [Cu2]3+ coordination environment, single layer
[Cu2Br(IN)2]n, and the bulk crystal of [Cu2Br(IN)2]n. (b) AFM topography image of
sonicated [Cu2Br(IN)2]n on HOPG. (c) Height of profile across the green line in (b).
Reproduced from Ref. [15]. Copyright © 2010, Royal Society of Chemistry.
In addition to exfoliating CP nanosheets from layered crystals, a direct
self-assembly is a more straightforward procedure to create 2D coordination polymers,
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and co-deposition of metals and organic ligands on crystalline solid surface offers a
promising approach. Li et al. [ 439 ] reported self-assembly of
5,10,15,20-tetra(4-pyridyl)porphyrin (2HTPyP) with Cu on Ag substrate (Fig. 48a). In
this process, Cu reacts in two ways: for one way, neutral Cu(0) atoms are coordinated
by peripheral pyridyl (py) substituent to form py-Cu-py bridge coordination; for
another way, Cu(0) is coordinated by inner nitrogen atoms of the porphyrin
macrocycle and, at elevated temperatures (450 K), oxidized to Cu(II). This
coordination network represents a mixed-valence polymer consisting of an ordered
array of Cu(II) and Cu(0) centers (Fig. 48b).
Fig. 48. (a) Self-assembly of 2HTPyP and Cu to form a mixed-valence Cu(II) and
Cu(0) network. (b) High-resolution image of 2HTPyP with Cu deposited on Au(111)
surface after 450 K annealing: yellow refers to Cu(II) and green to Cu(0). Reproduced
from Ref. [439]. Copyright © 2012, American Chemical Society.
Moreover, Abel et al. [16] reported formation of polymeric Fe-phthalocyanine
(poly-FePc) single sheet, which was obtained by co-evaporation of Fe and
1,2,4,5-tetracyanobenzene (TCNB) in 1:2 stoichiometry on surfaces such as Ag(111),
Au(111), as well as the insulating NaCl/Ag(100) surfaces. The poly-FePc film has a
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square structure with a measured periodicity of 1.15 ± 0.1 nm (Fig. 49). The
poly-FePc nanosheet is theoretically predicted to be an anti-FM semiconductor,
whereas a FM half-metal can be achieved when the central Fe atoms are replaced by
Mn [440].
Fig. 49. (a) Reaction scheme of poly-FePc. (b,c) STM images of poly-FePc formed on
Ag(111) surface. (d) STM image of poly-FePc formed on insulating NaCl island
deposited on Ag(100). Reproduced from Ref. [16]. Copyright © 2010, American
Chemical Society.
Besides the metal-porphyrin complexes, other types of organometallic sheets
have been recently reported. For example, Bauer et al. [ 441 ] synthesized a
monolayered organometallic sheet based on the reversible complexation of Fe2+ and
hexafunctional terpyridine monomers. Shi et al. [442] reported that on Au(111),
molecular ligands of 1,3,5-tris(pyridyl)benzene (TPyB) can form
metallo-supramolecular self-assembly with either Cu or Fe. Beyond the experimental
reports, Kan et al. [ 443 ] theoretically designed organometallic porous sheets
assembled by transition metals (Mn and V) and benzene molecules, which are ideal
magnetic materials with room-temperature ferromagnetism.
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6.2 Covalent organic polymers
Covalent organic polymers or frameworks (COFs), made entirely from light
elements (H, C, N, B, and O), are another intriguing class of porous materials that
allow for the integration of organic molecular units with strong covalent bonds into
rigid 2D or 3D architectures. Compared with coordination polymers, COFs provide a
great technological benefit concerning the aspect of their higher structural stability,
since the intermolecular covalent bonds are more stable than the coordination bonds.
Fabrications of COFs typically involve solvothermal reactions of
self-condensation of boronic acids or co-condensation with polyols [444], imine bond
formation [445], trimerization of nitriles [446], borosilicate bond (B-O-Si) [447] and
hydrazone bond formation [ 448 ]. In addition, surface-confined polymerization
reactions of 2D COFs have gained extensive interests, such as radical addition [449],
surface condensation reaction [450], and STM tip- and electron beam-induced surface
polymerization [451].
Most of the produced COF crystals, the boronate-ester derived COFs in
particular, are layered structures composed of π-π stacked covalent sheets that exhibit
staggered graphite-like or elipsed BN-like arrangements, and the single layer sheet
represents a promising model for 2D COFs. Recently, the oriented COF multilayer
films (hexagonal COF-5) grown on graphene substrate have been demonstrated by
solvothermal condensation of 1,4-phenylenebis(boronic acid) (PBBA) and
2,3,6,7,10,11-hexahydroxytriphenylene (HHTP), which showed improved crystallity
than the COF powders (Fig. 50) [452]. The COF-5 monolayer is predicted to be a
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semiconductor with a band gap of 4.0 eV (HSE functional), and its electron and hole
carriers are separated into the HHTP and PBBA units, respectively, suggesting the
potential applications of COF-5 in organic electronic and photovoltaic devices [453].
Fig. 50. Solvothermal condensation of HHTP and PBBA in the presence of a
substrate-supported single layer graphene (SLG) surface provides COF-5 as both a
film on the graphene surface, as well as a powder precipitated in the bottom of the
reaction vessel. Reproduced from Ref. [452]. Copyright © 2011, American
Association for the Advancement of Science.
Another recent example of monolayer COFs is a 2D heterotriangulene covalent
polymer. Bieri et al. [454] applied the on-surface synthesis approach to obtain
molecular-thin 2D polymer from tribromo-substituted dimethylmethylene bridged
triphenylamine (DTPA) precursors. The chemical step leading to DTPA films relys on
aryl-aryl homo-coupling between monomers (Fig. 51). By annealing the Ag(111)
surface to temperatures of 400 C, an ultraflat and methyl-cleaved 2D
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heterotriangulene polymer with covalent network formed. First-principles studies
[455] demonstrated that the TDPA porous sheet exhibits robust ferromagnetic
half-metallicity under an external strain, which makes the DTPA sheet an ideal
candidate for a spin-selective conductor.
Fig. 51. (a) Structure of DTPA and a fraction of the covalent network. (b) STM
topograph of an ultraflat, methyl-cleaved covalent network obtained by annealing to
400 C, with the structural model and corresponding STM simulation shown in (c).
Reproduced from Ref. [454]. Copyright © 2011, Royal Society of Chemistry.
Compared with the diverse synthesis routes and versatile structures of
coordination polymers, the construction of COFs has been restricted to only a limited
number of monomers, and the lack of suitable procedures that utilized other molecular
units has hindered further advancement of this emerging field. Besides, COFs are
usually hard to characterize since COFs have no single crystal structures. It is thus
important to extend the limited number of synthetic methods and monomer units
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available and improve the structural crystallity to go forward these emerging
materials.
7. Conclusion and prospective
The explosive studies on graphene have generated great enthusiasm towards the
explorations of graphene-analogous materials, which is a novel field that has received
growing interests at a spectacular pace in recent years. Different from graphene only
composed of carbon, graphene-analogous materials have more structural, bonding and
functional versatility. The distinctive properties of various graphene-analogous
materials make them promising candidates as 2D alternatives of graphene, and an
impressively large number of papers have appeared in this rising area during the past
few years.
The underlying goal of this review is to document and systematize recent
progress in experimental and theoretical exploitations of graphene-like
low-dimensional materials (2D nanosheets and 1D nanoribbons), including planar
graphene-like materials (h-BN, silicene, and BC3), theoretically predicted planar
materials (SiC, SiC2, B, and B2C), non-planar materials (metal dichalcogenides, metal
oxides and hydroxides, and MXene), metal coordination polymers, and organic
covalent polymers. In our overview, we mainly pay attention to the experimental
synthesis, characterization, functional applications, and theoretical predictions on
stability, electronic and magnetic properties of different systems.
Among them, h-BN is particularly unique due to its structural resemblance to
graphene but with totally distinct properties. BN nanosheets/ribbons have been
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synthesized by many techniques, and the structural defects such as vacancies were
observed. Although BN is characterized as a wide-band-gap semiconductor, its
electronic and magnetic properties can be modified to a large scale by defects, doping,
and surface or edge functionalizations. Si has the same valence number as C, yet
single layer Si, silicene, is not planar as graphene bur forms a slightly corrugated
configuration. The Si-C hybrid phase, SiC and SiC2, can form stable planar geometry.
Boron is a fascinating element, and its bonding is featured with an electron-deficient
character. The planar boron-containing materials, including B, BC3, and B2C
nanosheets have been proposed. MoS2 and WS2 nanosheets/ribbons have been
successfully produced, and they have moderate band gaps, suggesting the prospective
applications in nanoelectronics and photoelectronics. 2D nanosheets of layered oxides
and hydroxides can be fabricated via chemical exfoliation, and the resulting
nanosheets always form colloidal suspensions. Particularly, layered V2O5 was
exfoliated into thin nanosheets, which are promising cathode materials for Li-ion
batteries. Moreover, polar (0001) films of wurtzite ZnO with a few atomic layers
would transform into more stable graphitic structures, and the graphene-like ZnO
sheets have been observed in experiment. MAX phase is another family of layered
materials with strongly interacted MX and A atomic layers, and its separated MX
layers (“MXene”) have great potential as anode materials for Li-ion batteries. Besides,
2D coordination polymers and organic covalent polymers with tunable architectures
are another interesting area. They possess unique electromagnetic properties, and can
be additionally applied for gas storage and detection.
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Compared with the intensive research focus on graphene, investigations on
graphene-analogous materials have just come into the very beginning, and many
challenges exist, such as their large-scale fabrications and definitive characterizations
with desired thickness. Besides, many properties of graphene-analogous materials are
not fully explored, such as the practically measured mechanical, electronic, magnetic,
and optical properties, and their possible applications are still waiting for the further
explorations. However, we are highly expected that this emerging rich area has a
bright future and will rapidly grow to be an advancing field, and will generate
exciting properties and amazing applications.
Acknowledgements
This work was supported by NSFC (21073096 and 21273118) in China.
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