Nanostructured graphene Lu, Liqiang

Nanostructured graphene Lu, Liqiang

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Lu, L. (2018). Nanostructured graphene: Forms, synthesis, properties and applications. Rijksuniversiteit Groningen.

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Chapter 2

Dimensions, forms and synthesis of graphene — a critical review

In this Chapter, synthesis routes of graphene and their related foams are reviewed from high dimensions to low dimensions, i.e. from foams to films, and dots.

2.1 Large-area graphene film

Although initially by the mechanical exfoliation of graphene (so-called Scotch Tape-method) from highly oriented pyrolytic graphite one can obtain high quality of graphene, the sizes, thicknesses, and shapes are hard to control. In addition, it was shown in literature that the disorder induced by the substrate during growth of graphene is significant. Thus, the synthesis of large-area and high-quality graphene is still a great challenge and attracts considerable attention.

Currently, for the synthesis of graphene flakes and films the following methods have been developed: (i) mechanical exfoliation of graphite, (ii) liquid exfoliation of graphite, (iii) reduction of graphene oxides, (iv) chemical synthesis, (v) solid-state growth, (vi) epitaxial growth and (vii) chemical vapor deposition (CVD). The current methods have high-temperature as well as low-temperature syntheses.

Chemical vapor deposition, epitaxial growth and solid-state growth are performed at above 1,000 °C, which are the high-temperature synthesis. In contrast, other methods such as exfoliation of graphene, reduction of graphene oxides and chemical synthesis are at or near room temperature belonging to low-temperature synthesis. In this part, some typical methods for synthesis of graphene flakes and films including chemical or physical exfoliation of graphene and graphene oxides, high-temperature CVD and low-temperature growth and their challenges are reviewed.

2.1.1 Exfoliation of graphene and graphene oxide

In terms of the exfoliation method, so far, the original top-down mechanically

exfoliation method is chosen for preparing graphene-based devises because the

high-quality graphene with a low content of defects and high mobility carriers are

easier to be obtained by this method. Nevertheless, this method is neither high throughput nor high-yields.

The other general exfoliation approaches have been altering mechanical exfoliation: chemical efforts to exfoliate and stabilize individual graphene sheets in solution and oxidation, exfoliation of exfoliated graphene oxides from graphite followed with reduction. The solution methods have greater industrial potential in cost and yields in comparison with other mechanically exfoliation and high- temperature methods. The typical solution method is the reduction of graphene oxides developed by Hummers, which dated back to 1958.

Later on various modified Hummers methods were investigated by reducing the defects content, increasing the size of graphene oxides, boost the yields of monolayer graphene oxides.

Typically, the synthesis of graphene oxides contains two steps as shown in Figure 2.1. The first step is oxidation of graphite. During this process, the surface of graphite will be oxidized and functionalized by oxygen groups, resulting in the formation of graphite oxides (GO). The epoxides, alcohols, ketone carbonyls and carboxylic groups interrupt the contiguous lattice of graphene layer of graphite. As a result, the graphite lattice is disrupted and the interlayer spacing increases from 0.335 nm for graphite to more than 0.625 nm for graphite oxides. The well-reacted samples of graphite oxide will have an increased carbon to oxygen atomic ratio lying between 2.1 and 2.9 according to Hummers methods. Hence, GO is a polar material, is very hydrophilic, and readily disperses in water or other polar solvents to form stable colloidal suspensions.

Figure 2.1 Schematic illustration of the chemical oxidization of graphite and exfoliation of graphite oxide to graphene oxide from graphite.

With optimizing the acids and oxidants, the other recent improved Hummers methods not only can make the graphite oxides to be more oxidized, but also to be greener, safer, more facile, and higher yields. The interlayer spacing of GO could reaches above 0.9 nm, indicating much easier for exfoliation.

After the synthesis of graphite oxides by oxidation of nature graphite, the following step is exfoliation of graphite oxides to individual monolayer graphene oxide or few layer graphene oxides by means of water intercalation and ultrasonication or mechanical stirring.

The interlayer van der Waals forces of GO can be overcome by the intercalated

water and oxygen functional groups with the assistance of ultrasonic or mechanical stirring. Thus, the yield of monolayer graphene oxides is related to the oxidation of graphite and post treatments during exfoliation. The degree of oxidation of the graphite can be easily judged by the colour of the GO/water colloidal suspension.

With increasing the oxidation, the product would present from black hue to green and bright yellow.

In contrast to graphene, graphene oxides have very poor conductivity due to the disrupted sp

bonding networks. However, the π-network can be restored by means of reduction reaction and the electrical conductivity can be recovered. This is very important because graphene oxides have better chemistry than graphene owing to its functional groups and good dispersity in water. Graphene oxides become good precursors for synthesis of graphene-based composites. The product of graphene after reduction is named reduced graphene oxide (r-GO). The reduction can occur during the synthesis of composites or via post treatment after GO-based composites. Many reduction methods have been reported, including chemical reduction, thermal reduction and electrochemical reduction, etc.

During reduction, the atomic ratio of oxygen to carbon is reduced. For example, the r-GO reduced by sodium borohydride (r-GO-NaBH

) has a C:O ratio as high as 13.4:1, which is higher than 6.2:1 of r-GO reduced by hydrazine (r-GO-N

H

). Accordingly, the sheet resistances also increase to 59 and 780 k Ω sq

for r-GO-NaBH

and r- GO-N

H

.

Graphene oxides can be taken as precursors to prepare large-area reduced graphene oxide thin films, which can be used for portable electronics, energy storage devices, photovoltaic cells, transparent electrodes and anti-corrosion coatings. Various methods have been developed for preparing GO thin films, involving spraying spinning, dip-coating, rod coating, self-assemble, Langmuir- Blodgett assemble, vacuum filtrating and electrodeposition method of graphene oxides films on arbitrary substrates including metals, plastic, ceramics and others.

After deposition of GO coating, the r-GO film is obtained by reduction of GO. The reduction methods are considered according to the requirements of film and properties of substrates. Pei et al. found by using hydrohalic acids the r-GO films have better flexibility and high electric conductivity than that reduced by N

H

/H

O and NaBH

solutions. A flexible graphene-based transparent conductive film with a sheet resistance of 1.6 kΩ sq

and 85% transparency could be obtained.

To achieve good quality of r-GO film for transparent conductive films, large-area individual graphene oxides flakes of around 10,000 um

were also synthesized and separated using the modified Hummers method by Cheng’s group.

Although the reduced graphene oxides have high scalability, the quality of

graphene is still worse than that obtained by mechanically exfoliation of graphite

because of lots of defects and damages generated during the oxidation and

ultrasonic-assisted exfoliation processes. To avoid damages of graphene layers in

graphite and abandon the reduction processes, another liquid exfoliation of graphite is developed with the assistance of ultrasonication or stirring for direct synthesis of graphene.

The liquid solution can be N-methyl-2-pyrrolidone, dimethylformamide, dimethyl sulfoxide, N, N-dimethylacetamide, g-butyrolactone, 1,3-dimethyl-2- imidazolidinone and even ethanol-H

O mixture for green and sustainable chemistry. Adjusting the net energetic cost, which is dependent on the balance of graphene and solvent surface energies, the interlayer van der Waals forces can be overcome by the surface energy, peeling off the graphene layer from the graphite and graphene well dispersing in solvents. It also demonstrates a solvent characterized by surface tensions in the region of 40–50 mJ m

is good for exfoliation. For using N-methyl-2-pyrrolidone, the efficiency of exfoliation can go to 92.4 wt.%. This method has industrially scalable potential because of its simplicity and flexibility even possible in production by our kitchen by using a juicer. The organic solvents such as NMP could be reused after separating the graphene. However, this method also has many disadvantages. First, the yields of graphene (~0.01 mg mL

) and monolayer graphene ~1 wt.% are very low. Second, the use of organic solvents increases the costs. The separated graphene is not stable and can re-stack to multilayer or thin graphite aggregates. Third, the graphene resources limit the flake size of graphene. The size of exfoliated graphene IS normally even smaller than that by chemical exfoliation. To achieve millimetre and even larger size of graphene is still a big issue.

2.1.2 Chemical vapor deposition growth of graphene

In contrast, chemical vapor deposition provides a versatile and promising strategy for producing large-area up to centimeters or meters, high-quality, uniform, continuous, and single crystal graphene with controllable size and thicknesses. The CVD growth of graphene normally requires transition metals as catalysts. The interactions of monolayer thin graphitic carbon on surfaces of transitional metals such as Ni and Cu was initially a research topic and was studied before the mechanical exfoliation of graphite.

The formation of pyrolytic graphite by the decomposition and crystallization of hydrocarbons (gases and solid) normally requires a temperature as high as above 1,500 °C to form high crystallinity of graphitic carbon. The introduction of metal, a catalyst, enables these processes with low activation energy at lower temperatures.

Thus, the metal catalysts are critical to the CVD process for dissociation of hydrocarbons, graphitic carbon nucleation, crystal growth, connection and merging of adjacent domains.

Except, the CVD growth conditions such as growing temperature, pressure, gases,

cooling time also influence the size, shape, quality and electronic properties of as-

grown graphene. In this section, the CVD growth processes, kinetics and

mechanisms, influences of catalysts surfaces and CVD growth conditions, strategies

to achieve large-area graphene, and future challenges are reviewed.

Currently, various kinds of metals have been reported for the catalytic growth of graphene, including Cu, Ni, Pt, Co, Fe, Ge, Ru, Ir, Pd and so forth. According to the solubility of carbon in metal, these metals can be classified as low carbon solubility catalysts such as Cu, and high carbon solubility catalysts such as Ni. Due to the different solubility and carbon-metal interactions, the formation mechanisms of graphene are different.

(1) CVD growth of graphene on Cu

Previous research has demonstrated the weak interaction between π network of graphene and electronic structure of copper, which has a closed d-shell electronic configuration. Thus, the linear band structure of free-standing graphene can be preserved on Cu. The use of copper as a catalyst for graphene growth has been developed quite fast since graphene became a “star” on the horizon of materials science. Copper has many advantages for catalyzing graphene growth owing to its good catalysis, low cost, availability of large-area foil and facile manufacturing. Figure 2.2a schematically shows a typical experimental setup of CVD growth of graphene on Cu.

Similar to other CVD processes, precursors with as-deposited atoms are introduced from one end, following with deposition of as- deposited films or coatings onto target bodies or substrates sited in the heating zone. For CVD growth of graphene, hydrocarbon such as methane is commonly used for providing carbon atoms. After CVD growth, normally a single-crystal graphene with large-area up to millimeter size as shown in Figure 2.2b can be obtained on Cu.

Regarding the CVD growth process, typically it involves the following steps as shown in Figure 2.2c: (1) mass transport normally occurs with a steady state gas flow, which has a boundary layer with thickness of δ. During the mass transport, the hydrocarbon species diffuse through the boundary layer and reach the surface of copper. (2) In the following these species are adsorbed on the surface. (3) Then carbon species catalytically decompose to provide active C species that are adsorbed on the Cu surface. These active carbon species have surface diffusion that thermally activated and driven by the concentration gradient of C atoms; (4) after reaching a concentration active carbon species nucleate into graphene domains.

Other carbon species attach to the edges, and are incorporated in the domains that

propagate to large-area domains. (5) The by-products such as hydrogen,

hydrocarbon species, part of carbon atoms desorb and (6) diffuse away from the

surface even together with some catalyst atoms (evaporation of Cu atoms from the

Cu surface also occurs at 1,000−1,080 °C as the melting temperature of Cu is

1,085 °C. It is more severe in vacuum CVD).

Figure 2.2 (a) Schematic illustration of the CVD setup, (b) a graphene domain with hundreds of microns grown on Cu by CVD method,

(c) the general CVD process of graphene growth on Cu catalyst.

(i) Influences of growth conditions

The overall process includes mass transport-controlled zone from the bulk of gas flow through a boundary layer to the surface of Cu, and a surface reaction- controlled zone on the surface of catalyst of Cu as shown in Figure 2.2c. The flux of reactant species to the substrate through the boundary layer F

is given as

= ℎ ( − ) 2.1

where ℎ is the mass transfer coefficient of the gas, is the concentration of gas in bulk of flow and is the concentration of carbon species on the Cu surface.

The flux of reactant species consumed by the surface reaction F

is given as

= 2.2

where is the surface reaction constant.

At steady state, F

=F

, and we get the total flux F

= 2.3

Thus, the growth of graphene is related with the surface reaction and mass

transport. Typically, there are three situations: (1) when ≫ ℎ , the growth of

graphene is controlled by mass transport. This often happens at high temperatures

under atmospheric pressure CVD (APCVD) and at a slow flow rate. (2) When ≪

ℎ , the growth of graphene is controlled by surface reaction. Density functional theory (DFT) calculations reveal that the H-terminated graphene edge on Cu is more stable than the bare graphene edge on Cu. The diffusing C species on Cu are mainly CH

(0 < x < 4), which are generated on the basis of endothermic decomposition. Thus, the attachment of carbon species on domain edges and incorporation into graphene lattice is rate-limited at low temperatures, high flow rate, low pressure or ultrahigh vacuum conditions. (3) When ~ℎ , the growth of graphene is sensitive to many influences such as temperature, pressure, flow rate, position of substrates and so forth. It is found by controlling the pressure, flow rate, pressure and time, the growth rate, growth coverage, domain density and size of graphene can be tuned easily. It should be pointed out that the larger-area graphene CVD films are continuous but are polycrystalline. The properties of as- obtained macroscopic film depend on tuning the domain size, connectivity, and the domain-boundary structure including related defects. During growth, the domain density can influence the domain size as many of the domains are in different orientations and generate many grain boundaries. High domain density can typically lead to smaller domains. The domain density depends on the nucleation density, which is related with the surface reaction. At a condition of high temperature of ~1,000 °C, low pressure and low flow rate of hydrocarbon, is much smaller than ℎ and the surface reaction controls the rate. As shown in Figure 2.3, low pressure and low flow rate will reduce the nucleation density and growth rate of graphene, but as a result, the domain size increases. In turn, under high temperature, high pressure and high flow rate, the growth rate increases but the domain size decreases upon increasing the domain density. In a short growth time, normally the coverage and connectivity of graphene are higher for under later conditions than that under former conditions. Thus, to obtain large-area single- crystal continuous graphene films, adjusting the growth conditions is very significant.

Another interesting growth behaviour of graphene is the formation of various

domain shapes under different conditions. There are four-types of shape found,

which are four-lobed, hexagon, between hexagon with dendritic edges and square,

as shown in Figure 2.4.

It has been observed that normally under vacuum

conditions, the shape is four-lobed or hexagon with dendritic edges. An argument

about the formation of the four-lobed island, which is expected to be single-domain,

is related to the Cu(100) square lattice. However, the low-energy electron

microscopy (LEEM) study exhibits a polycrystalline structure with multiple

domains. Under ambient conditions, the shape mostly is hexagon but sometimes

square. The hexagon is often observed on polycrystalline Cu foil. Raman analysis

reveals that the hexagon domain has zigzag edges, which are known to be formed

by metal-assisted anisotropic etching of graphene and hydrogen plasma etching.

Thus, another possible explanation for the shape formation of domain is related

with the Cu evaporation. The CVD process at high temperatures close to the

Figure 2.3 SEM images of graphene partially grown on Cu under different growth conditions: T (°C) / J

(sccm) / P

(mTorr): (a) 985 / 35/ 460, (b) 1035 / 35 / 460, (c) 1035 / 7 / 460, (d) 1035 / 7 / 160. Scale bars are 10 µm.

Figure 2.4 (a) LEEM image of four-lobed island,

SEM images of graphene (b) hexagonal

domains,

(c) square domains,

(d) hexagonal domains with fractal edges,

and (e)

flower-shaped domains with fractal edges.

melting point of Cu (1,085 °C) may result in the vaporization of Cu, as confirmed by the observation of Cu film deposited on the end of tube. The as-vaporized Cu atoms could disturb the growth of faceted graphene domains. The influence of evaporated Cu on shape formation is more different in low pressure CVD, owing to faster Cu evaporation. Another argument is that the graphene domain shape can be influenced by kinetics either limited by the incorporation of the constituent species at the growth fronts or by the diffusion of constituent species.

(ii) Influences of catalysts

Another strategy for producing large-area single-crystal continuous graphene films is by engineering the surface of copper catalyst. Except the influence of temperature and consumed hydrocarbon species, the nucleation of graphene is also related with the surface of catalyst such as grain boundaries, surface roughness, local defects, impurities and crystallographic orientation. It has been reported that the grain boundaries, impurities and rough surface can serve as active sites and stimulate the formation of graphene nuclei. These sites may lead to strong binding to adsorbates like carbon species for graphene nucleation. Thus, a rough and highly polycrystalline catalyst surface can induce the formation of high-density graphene nuclei, and can generate many grain boundaries and small sized graphene domains.

Electrochemical polishing of Cu substrates is reported as an effective way in getting flat surface and reducing the density of the graphene nuclei.

Except that, controlling oxygen is becoming an alternative and efficient approach to decrease the graphene nucleation density by passivating the active sites.

By using Cu catalyst with oxygen atomic density of ~0.01%, the domain density can be reduced to ~0.9 mm

from 2×10

mm

for oxygen free catalyst (O~10

atomic%). Oxygen can also lower the graphene domain growth and alter the growth kinetics from edge-attachment-limited to diffusion-limited. By control of surface O content, centimeter-scale single-crystal graphene domains with high quality can be achieved.

Moreover, Stefano et al. compared the graphene growth on Cu(111) and preoxidized Cu(111) and found that the graphene grown on oxidized Cu preserves its intrinsic properties while graphene grown on Cu(111) has weak interaction and doping.

At low carbon chemical potential, graphene grown on Cu is preferably single-

crystal and single-layer because a monolayer coverage can prevent the transport of

precursors to the catalyst surface, and prohibit the formation of a second layer,

which is so-called the self-limiting mechanism. However, by probing the graphene

grown on surfaces with various crystal facets after APCVD under a high carbon

species pressure, it is also found that the underlying crystallographic orientation of

Cu can influence the number of layers of graphene as shown in Figure 2.5. All the

graphene grown on Cu has few defects. But the graphene grown on Cu(100) has the

lowest intensity ratio of I

/I

(I

and I

are the Raman intensity of 2D band and G

band), suggesting few-layer graphene. In contrast, the highest I

/I

observed on

Cu(111) indicates the high quality monolayer graphene. The graphene films grown

on Cu(100) surface as well as highly faceted regions like Cu(10,7,6) have lower

I

/I

. This could be due to the limited surface diffusion of carbon. Another argument is that the additional graphene layers could nucleate between the initial graphene and catalyst, probably caused by the diffusion of precursors through defects of the initial graphene layer.

Figure 2.5 (a) Optical image of Cu surface crystal facets identified by EBSD and annealing twins (dotted) present. (b) Raman spectra taken at the selected colored shapes of (a). (c) Raman spatial map of graphene monolayer intensity ratio I

/I

for the region in (a) (Raman pixel size is 7.5 μm at 633 nm excitation). (d) graphene defect intensity ratio I

/I

for the same region (I

is the Raman intensity of D band).

(iii) Atomic connection of adjacent domains

Except enlargement of the domain size, an alternative method for synthesis of

large-area single-crystal graphene film is by connecting the adjacent graphene

domains together. This merging process requires identical oriented graphene

domains, as a misorientation between two graphene domains can induce graphene

boundaries as shown in Figure 2.6 a and b. For that, epitaxial growth can ideally

connect two identical graphene domains into a seamless graphene crystal. Previous

STM, LEEM and LEED studies revealed that graphene on Cu(111) could have a

slight rotation while on Cu(100) it gives a variety of rotation angles as shown in

Figure 2.6 c-f. However, Stefano et al. observed that graphene on Cu(111) also

presented various different orientation angles.

For epitaxial growth of graphene,

it requires well-defined, highly crystalline catalyst surfaces with a minimum lattice

mismatch to graphene.

Figure 2.6 (a) Schematic illustration of grain-boundary formation (marked in red) by a misorientation between different graphene domains, and two identical domains merging into a larger monocrystalline graphene in the area highlighted in green. (b) SEM image of two domains with a misorientation.

(c) Orientations of graphene domains grown on Cu(100) (d) and Cu(111) lattices. (e) DF-LEEM images of graphene grown on Cu(100) surface (f) and graphene on Cu(111) surface. The colours represent the orientation of graphene domains.

(2) CVD growth of graphene on Ni

Ni is another most widely studied catalyst for the synthesis of graphene. In comparison with Cu, there are many differences on the interaction between graphene and Ni, and graphene growth mechanisms on Ni. Graphene on Ni has a strong interaction owning to the open d-shell structure of Ni.

The overlap between the graphene π and Ni d valence band states completely devastates the linear graphene band dispersion around the Fermi level. Graphene growth on Ni(111) also has a very small lattice mismatch (lattice constant of 2.46 Å for graphene and 2.49 Å for Ni(111)), which are ideally suitable for the epitaxial growth with minimum additional strain induced in graphene.

This facilitates the growth of graphene on the Ni(111) surface. Owing to the high solubility of carbon in Ni, a growth mechanism involving dissolution and segregation of carbon into/out of Ni catalyst leads to controlling of not only monolayer but also multilayer graphene.

Controlling the layer number of graphene on Ni becomes more versatile but also more complicated because of carbon solubility varies with the temperature. What is more, graphene domain shape looks random. Hence, it is often argued that for monolayer graphene CVD a catalyst with low carbon solubility is essential and that for multilayer graphene on high carbon solubility metals adjusting growth via precipitation upon cooling is important.

(i) Kinetic control

A typical CVD process for deposition of graphene on Ni consists of the

following procedures: pre-treatment of a Ni catalyst at an elevated temperature,

introduction of hydrocarbon at a constant high temperature, and cooling down.

During cooling, second and more graphene layers could form after a monolayer graphene on Ni catalyst. For a considerable thick Ni catalyst without having saturated carbon, the graphene formation can be mainly on the growth kinetics of graphene on Ni catalyst without the segregation of carbon onto surface of Ni during cooling, leading to a monolayer or few-layer graphene grown on the surface of Ni.

Thus, a strategy for growing monolayer or few-layer graphene can be achieved by balancing the carbon flux from precursor impingement and dissociation (J

) that is related to carbon diffusion into Ni (J

) and graphene growth (J

), as shown in Figure 2.7.

The enhanced domain size and less thickness of graphene grown on thicker Ni catalyst demonstrate the significant influence of Ni thickness on graphene growth.

Figure 2.7 (a) A schematic illustration of the balance between the carbon flux (J

) from precursor, carbon diffusion into the catalyst (J

), and graphene formation (J

). Higher precursor pressures can lead to a local carbon supersaturation Δc near the catalyst surface which results in graphene monolayer formation. The JD remains limited as indicated by the carbon concentration depth profile in red. (b-d) SEM images of graphene grown on various thicknesses of Ni (b) 550 nm, (c) 25 μm and (d) 250 μm. (b) shows inhomogeneous multilayer graphene formation, while (c and d) show uniform MLG coverage across different Ni grains.

(ii) Crystallographic effect

The readily available Ni catalyst is polycrystalline Ni that is from commercial rolling product or deposited by sputtering. Compared with single-crystal Ni, poly- Ni has many and different grain size, grain orientations, grain boundary and other defects or impurities. Zhang et al. studied the growth behaviour of on single-crystal Ni and poly-Ni.

It is found monolayer/bilayer graphene is preferentially formed on the single crystal surface, whereas polycrystalline Ni catalyses a higher percentage of multilayer graphene, which is attributed to the high density of step edges in grain boundaries that can serve as nucleation sites for multilayer growth.

Thus graphene grown on poly-Ni by CVD normally contains many grain boundaries,

inhomogeneous and multilayer features, whereas fewer nucleation sites on Ni(111)

leads to large-area and few layer graphene. To reduce the grains boundaries and

increase the grain size of Ni, thermal annealing in Ar/H

atmosphere at 900-

1,000 °C before introducing hydrocarbon is efficient, leading to enlarged and

thinner graphene crystals. Besides, deposition of poly-Ni on flat surfaces such as MgO substrate followed with annealing can improve the quality of Ni catalyst.

In terms of the crystal orientation of Ni on graphene formation, Kozlova et al.

used electron backscatter diffraction (EBSD), Raman and energy dispersive X-ray spectroscopy investigated the growth behaviour of graphene on different grains of poly-Ni.

As shown in Figure 2.8, for a short time CVD, the grains with Ni(111) orientation are well covered, whereas non-continuous graphene coverage was observed on grains oriented close to Ni(001) and especially on higher index surfaces. For a long time CVD, almost all Ni grains with various orientations could be covered with graphene, but interestingly thinner graphene observed on grains oriented close to Ni(111) and thinnest graphene is observed on ideal Ni(111). In addition, the shapes of graphene grown on Ni grains of various crystallographic orientations are different. Graphene growth on grains close to Ni(001) appears long dendritic branches oriented along the steps, seeming an extended part to these faces from neighbouring grains. The observed difference in graphene formation on the (001) oriented grains with (111) oriented grains could be due to a larger binding energy of the carbon atoms to (001) surface than (111) surface as each carbon atom interacts with four Ni atoms. Dendritic edges are also observed on graphene grown on Cu. It is still not clear the formation of dendritic branches. Thus a possible explanation could be the dendritic branches are caused by fast etching of graphene on (001) oriented grains before or during cooling, because the strong Ni–C interaction would cause a repulsive interaction within the C–C interaction and cause dissolution at the edge of graphene.

Figure 2.8 (a) SEM image and (b) EBSD orientation map taken from the same area showing the dependence of coverage on grain orientation after 10 min of methane exposure at 1,000 °C and fast cooling.

(iii) Reducing nucleation density

To synthesize large-area graphene films on poly-Ni, another strategy is

reducing nucleation density. However, normally on polycrystalline Ni catalyst

surface the density of graphene nuclei is very high, which results in small size of

graphene domains and inhomogeneous multilayer graphene films as shown in

Figure 2.9 A,C and E. Applying alloyed catalyst such as Au-Ni (as shown in Figure

2.9 B,D and F) is reported an efficient approach to reduce the nucleation density of graphene domains because Au decoration could suppress the edges steps of bulk poly-Ni, which stimulate isothermal graphene nucleation.

Figure 2.9 SEM images of graphene grown on Ni (A, C, E) and Ni-Au alloy catalysts (B, D, F).

(3) Challenges

For most applications, the quality such as size, crystallinity, layer thickness of graphene is very crucial. The CVD synthetic technique consists of catalytic growth of graphene on metal substrates (Cu or Ni) and post-growth transfer of the graphene film to a substrate of interest. For growth process, the as-obtained graphene is generally polycrystalline and contains a significant amount of domain boundaries that limit intrinsic physical properties of graphene. Considering influences of graphene domain boundaries on the electrical, mechanical, and thermal properties, the synthesis of “single-crystalline graphene” free from domain boundaries is an important and challenging issue.

Most of the previous works are mainly on grown transferable graphene on metal supports by CVD approach. The post-transfer process of graphene to other substrates especially insulating substrates BN, SiO

, glass or plastic are normally complicated, including etching the metallic support, pressing, heating. During the transfer process, there are many problems limiting this method to industrial use, due to not only the increased cost, but also the degraded quality owing to introduction of wrinkles, cracks, and contamination into the graphene. Therefore, it is strategically important to develop a transfer-free, scalable growth process for depositing high-quality graphene on dielectric surfaces without post-growth treatments.

Researchers developed transfer-free graphene on metallic and non-metallic

substrates. Mark P. Levendorf et al. reported a method preparing transfer-free

graphene by grown graphene on an evaporated copper film in CVD conditions.

Pan et al. prepared transfer-free graphene on SiO

insulator wafers by rapidly thermal treatment sputtered carbon and metal layers at temperatures as high as 1,200 °C.

Shin and co-workers using pyrolysis of aromatic carbon materials and self-assemble process synthesized few-layer graphene on SiO

substrates avoided the transfer process.

The above results were all performed at relatively high temperatures (normally ~1,000 °C), which is necessary for dissociation of hydrocarbon. However, the high temperature limits the choice of substrates and adds additional transferring of graphene to other thermal unstable materials such as polymeric materials. Developing low-temperature growth of graphene directly on thermal unstable materials is very important.

2.1.3 Low-temperature growth of graphene

Low-temperature synthesis provides a strategy for overcoming the high- temperature problems. Herein low temperature means temperatures below 800 °C, but mostly below 600 °C. With reducing the temperature from ~1,000 °C, commonly the graphene quality decreases fast in particular below 800 °C. The crystallinity of graphene reduces and defect content increases, leading to a decrease of the carrier transport. At low temperatures, the dissolution of carbon in Ni also dramatically reduces. Thus, the growth mechanism of graphene at low temperatures may different from that at high temperatures. Addou et al. studied the graphene growth at 550-650 °C by using LEEM and Auger electron spectroscopy (AES) microscope, and found a self-limiting monolayer graphene growth mechanism on Ni(111).

Yang performed the graphene growth on Cu at 500 °C under APCVD, and reported a low sheet resistance of ~54 kΩ sq

due to the high defects and low crystallinity.

It is known that at a low temperature such as 600 °C the dissociation of hydrocarbon such as methane is difficult because of its high activation energy for decomposition. To address the problem, plasma-enhanced CVD (PECVD) has been developed to reduce the activation energy and assist the dissociation of hydrocarbon at low temperatures. Kim et al. developed the microwave plasma chemical vapor deposition (MPCVD) to reduce the synthesis temperature from 750 °C down to 450 °C.

With decreasing the temperature, the defect content largely increases. The as-synthesized graphene at 450 °C presents a sheet resistance of 1,855 Ω sq

, which is lower than that of the graphene catalysed on Cu at 500 °C under APCVD. Lee et al. using PECVD synthesised large-area formation of continuous transparent graphene layers on glass at 450 °C. The as-measured sheet resistance is locally as low as 500 Ω sq

.

Surface wave plasma chemical vapor deposition is also employed for synthesis of graphene-based transparent conductive films at a low-temperature range of 300–400 °C. The sheet resistance of the films reaches ~1.0 kΩ sq

.

Except PECVD techniques, alternative carbon sources with low activation

energy for decomposition are used. Li et al. investigated the graphene growth using

benzene as the hydrocarbon source on Cu, and found good-quality monolayer graphene domains with sizes of ~2-3 μm can be achieved at temperatures as low as 300 °C as shown in Figure 2.10.

First principle calculations give a 1.47 eV of activation energy of dehydrogenation of benzene on Cu(111), which is much lower than that of methane (1.77 eV). In addition, benzene is suggested to have a lower nucleation barrier for graphene compared with methane, allowing graphene growth at such low temperature. In compared with benzene, using another solid source of polymethylmethacrylate (PMMA) required a higher temperature of ~400 °C and as-generated graphene has higher defects content. Different polycyclic aromatic hydrocarbons will induce different quality of graphene.

Figure 2.10 Graphene synthesized by using benzene: (a) Raman spectra of graphene synthesized at 500 and 300 °C, respectively; (b and c) SEM images of graphene grown at 500 and 300 300 °C, respectively. Scale bars are 2 μm.

It is found that a higher quality of monolayer graphene flakes can be achieved at a growth temperature as low as 550 °C on Cu catalyst by using coronene than those by using pentacene or rubrene, because of avoiding the formation of pentagons in lattice structures.

Besides using PMMA, Zhu et al. studied using other solid carbon sources including poly(vinylpyrrolidone) (PVP), polystyrene (PS), and polyethyl glycerol (PEG) for synthesis of graphene on stainless steel at 500 °C.

Raman analysis as shown in Figure 2.11 indicate the defect contents of graphene obtained from polymers has a sequence of PS>PMMA≈PEG>PVP.

Unfortunately, no discussion on the difference is given. Nonetheless when comparing the chemical formula of these polymers, we can conclude that non- carbon atoms such as hydrogen, oxygen, and nitrogen varies, lead to different types of C−H, C−O or C=O, and C−N bonds. These bonds and heteroatoms could increase the activation energy of dissociation and graphene nucleation.

Another advantage of using low temperature solid-state growth is facilitating heteroatom-doped graphene films. Under thermal treatment, the formation of C−C and C−dopant bonds to is in competition to achieve the minimum internal energy.

C−C has higher bond energies of 346 kJ mol

than that of C−N (305 kJ mol

), and

C−S (272 kJ mol

), resulting in C−C bond formation at high temperatures. Thus,

the dopant incorporation favour low temperature conditions. Zhang et al. used

pentachloropyridine as sources synthesized N-doped graphene on Cu at 230 °C

under vacuum conditions, and tetrabromothiophene synthesized S-doped graphene on Cu at 300 °C under vacuum condition. The as-synthesized N-doped graphene- based FETs delivered a highly n-type electrical property with electron mobility of 80.1−302.7 cm

V

s

, whereas S-doped graphene-based devices display electron mobility of 2.6−17.1 cm

V

s

.

Figure 2.11 (a) Raman spectra of graphene grown on stainless steel at 500 °C using PS, PVP, PEG, and PMMA as solid carbon sources,

(b-e) Skeletal formula of PS, PMMA, PEG and PVP.

Figure 2.12 (a-c) SEM micrographs of graphene on Ni with implantation temperature of 600 °C. In (b) the MLG fragments of (c) are delimited with white dashed lines to underline the facets with 120 ° angles.

Although either plasma-enhanced PVD technique or liquid or solid

hydrocarbon resources can significantly reduce the activation energy of

dissociation, the graphene formation still require heating at ~300-400 °C for

decomposition of hydrocarbon, meanwhile the as-synthesized graphene remains

many defects. Consideration of the strong interplay between carbon and Ni atoms,

sputtered amorphous carbon provides possible solutions to lower down the

temperature further. Zheng et al. performed the graphene growth on free poly-Ni

surface by using sputtered C at relatively high temperature 650-950 °C.

The

thickness of graphene layer can be controlled by adjusting the feedstock of carbon

layer. However, the author did not conduct the graphene growth behaviour at lower

temperatures. Gutierrez et al. studied the graphene formation by implanting of

carbon in poly-Ni at 450-600 °C and following with remaining at same

temperatures as shown in Figure 2.12.

Few layer (~4) graphene fragments with

facetted edges with angles equal or multiples of 60 ° are observed, indicating the

hexagonal structure of graphite crystals. The study proves graphene can form on Ni surface through diffusion at low temperatures.

Normally, the as-synthesized graphene from amorphous carbon contains lots of grain boundaries and uncontrolled layer thicknesses. This is caused by the high density of grain boundaries in sputtered poly-Ni and fast carbon diffusion. During temperature ramping, the carbon diffuses and start nucleation to graphene at step edges of poly-Ni in early stages even before reaching the desired temperatures. By using a barrier layer of Al

O

is useful for overcoming the premature carbon dissolution and diffusion during ramping, facilitating the grain expansion of Ni.

Compared with the condition without barrier layer, large area monolayer graphene can be synthesized at 600 °C by using an Al

O

barrier layer as shown in Figure 2.13.

Figure 2.13 Schematic illustration of inhibiting carbon dissolution at early heating stage to improve graphene formation from catalytic transformation of amorphous carbon (black) by introducing a diffusion barrier (green), on the right are optical images of graphene on Ni formed at 600 °C without and with a barrier layer (Al

O

).

Another work for growth of graphene by using diffusion of carbon atoms is

done by Kwak et al., who pasted graphite on poly-Ni and synthesized transfer-free

graphene films on non-metallic substrates, such as SiO

/Si, glass and plastic

substrates, via diffusion-assisted synthesizing method at near room-temperature

(25-260 °C).

The diagrams in Figure 2.14 represent (from left to right) the

elementary steps of the process, which includes dissociation of C–C bonds at C/Ni

interface, diffusion of carbon atoms, followed by heterogeneous nucleation at the

defect sites, and growth of graphene via lateral diffusion of C atoms along

Ni/substrate interface. This opened up new possibilities for preparing graphene

films on non-conducting substrates at close to room temperature, and reduced the

limitations of substrate choice due to the elevated temperatures. However, it still

requires the dissociation of C–C bonds of graphite. The synthesis of large-area and

single-layer graphene is still limited by the nickel grain size, which is enlarged by

high temperature treatment. In addition, the growth mechanism, the function of

Ni/substrate interface, and the size and layer controllability are still not clear.

Overall, synthesis of transfer-free graphene on metal substrates at low- temperatures is a big challenge.

Figure 2.14 Schematic diagrams of graphene growth mechanisms in diffusion-assisted synthesis.

2.2 Three-dimensional porous graphene

Construction of 3D porous graphene requires efforts as normally graphene layer tend to restack due to the strong π-π interactions. Currently, there are two main approaches for the synthesis of 3D porous graphene: one by cross-linking of reduced graphene oxides, another by chemical vapor deposition growth of graphene onto porous template followed with removing templates, which is named template-assisted CVD growth.

2.2.1 Cross-linking of graphene oxides

Graphene oxide, an amphiphilic macromolecule with a hydrophobic basal plane and hydrophilic edges (particular for partially reduced graphene oxides), is a good precursor for construction of 3D porous graphene monoliths.

The oxygen functional groups can be used for cross-linking to increase the connection. The interactions of electrostatic interaction, hydrogen bonding, and covalent bonding are normally employed for cross-linking. The as-built graphene foam normally presents the structure as shown in Figure 2.15a.

The van der Waals adhesion between layers can be overcome by the internal elastic energy. Owing to the use of defective graphene oxides, the final 3D graphene foam usually contains defect- induced holes on 2D units.

Many fabricated graphene monoliths have the fragile problems and small

recoverable deformation before failure. These mechanical problems often occur in

porous carbon materials. It reflects construction of graphene oxides to foam

requires careful processing. For a rational construction of GO into 3D foam, various

kinds of techniques such as hydrothermal and solvothermal, freeze-drying and in

situ chemical reduction have been developed. To increase the elastic modulus,

high-temperature post treatments are also often used. Shi et al. successfully

fabricated mechanically strong graphene aerogels with good electrically conductive

and thermally stability via a hydrothermal process (as shown in Figure 2.15b).

The size of pores is from submicrometer to several micrometers and the walls

consist of stacked thin graphene sheets.

Figure 2.15 (a) Models of 3D porous graphene without interconnected porous structure,

(b) SEM image of a 3D porous graphene by hydrothermal-assisted self-assembly of reduced graphene oxides.

The inherent flexibility of graphene sheets is a crucial property for constructing the 3D macrostructures. It is found that, with increase the C/O ratio, the storage modulus, elasticity, yield stress and electrical conductivity increases.

Maximally, a storage modulus of 450–490 kPa, elastic modulus of 290 kPa and yield stress of ~24 kPa can be achieved and over other crosslinked polymer hydrogels when the C/O atomic ratio increases to 5.3. During hydrothermal, metal/

metal oxides nanoparticle can be generated and loaded onto the 3D graphene foam to form composites.

Hydrothermal-assisted assembly is very facile, but still limited to the large-size production due to the limitation of container. In addition, the graphene sheets are usually randomly linked in a random fashion.

Freezing-assisted assembly of graphene oxides provides an alternative solution to obtain porous graphene foam. Different from hydrothermal method on aspect of microstructures, the later contains a hierarchical cork-like cellular structure.

However, direct freezing of GO dispersions can result in a random arranged porous structure. The secret for cork-like cellular structure is by using ice- segregation-induced self-assembly during freezing. Qiu et al. found when a well- dispersed partially reduced GO dispersion is frozen, partially reduced GO sheets are concentrated at the boundary of ice crystals and then aligned along the growth direction of ice due to the squeezing effect. As a result, a continuous honeycomb- like network can be formed. After thawing the ice, the network retains its connectivity. When the C/O atomic ratio in partially reduced GO was tuned to be around 1.93, a cellular structure can be obtained. The resulting biomimetic graphene-based monoliths exhibit a combination of ultralow density, superelasticity, good electrical conductivity and high efficiency of energy absorption. Besides, cross linkers are also beneficial due to their elasticity and fatigue resistance in resulting aerogels. Gao et al. synthesized lamellar chitosan–

a

graphene oxide scaffolds from a homogeneous mixture via bidirectional freezing and annealing.

The resulting composite showed high compressibility, super- elasticity, and low density. Alternative cross-linkers such as sodium dodecyl sulfate and cellulose have also been reported.

2.2.2 Template-assisted CVD growth

In comparison with cross-linking of graphene oxides, template-assisted CVD growth can produce more pristine graphene materials. Currently the templates for graphene growth contain two types, metallic scaffolds and porous metal oxides/ceramic materials. With respect to the metallic templates, Ni foam is a typical material and commonly used. Chen et al. initially synthesized porous graphene foam by CVD growth of graphene onto commercial Ni foam at 1,000 °C as shown in Figure 2.16.

After synthesis, they used poly(methyl methacrylate) (PMMA) protective coating layer to support the hollow graphene skeleton. The as- synthesized porous graphene exhibits higher crystalline, purity, and better robust compared with cross-linked RGO foam. The BET surface area reached 850 m

g

, and the electrical conductivity could be increased to 1,000 S m

, which is 3–4 orders of magnitude higher than that of hydrothermal graphene foam. Even bending for 10,000 cycles to a radius of 2.5 mm, there was only a small increase of resistance (~2.7%).

Figure 2.16 Models of 3D porous graphene without interconnected porous structure.

In another work, Cu foam was employed for CVD growth of graphene.

However, the as-synthesized graphene foam has poor connectivity and higher defects than that from Ni foam. This could be because of the different growth mechanism of graphene on Ni and Cu. Although the graphene foam prepared by using commercial Ni and Cu foam have good chemical and physical properties, but the products are very fragile and normally require supporting materials such as PMMA. In addition, the pore size is in hundreds of microns, which is much bigger than that of derived from hydrothermal graphene foam.

For many energy storage devices, nanopores are very significant in designing

device. Being the template-directed method, reducing the pore size of metallic

templates can pave the way for obtaining 3D nanoporous graphene (3D NPG).

Therefore, developing nanoporous metallic templates are very important for synthesis of 3D nanoporous graphene.

With respect to synthesis of nanoporous metals, one of the most favored approaches by the porous-metal community is dealloying, also named selective etching. Before dealloying, good candidates preferentially binary solid solutions should be selected. During etching, the less noble component is selectively dissolved whereas the nobler part remains and rearranges to form the bicontinuous porous architecture with interconnected pores and solid phase.

Ito et al. firstly investigated CVD growth of graphene onto dealloyed Ni nanoporous skeletons at

~800 -1,000 °C. The as-prepared nanoporous graphene have pore size of 0.1-2.0 µm and BET surface area above 1,200 m

g

. The free-standing 3D NPG does not require supporting materials, and preserves better physical characteristics than those made by assembly.

Versatilely, heteroatom doped NPG can be also prepared by using this method. However, it is found during CVD growth, that the nanoporous Ni templates have severe structural transformation. The ligament size and pore size occur fast coarsening, typically from ~10 nm to 0.1-2.0 µm, which is due to the high diffusion rate of Ni atoms at high temperatures. This phenomenon occurs more obvious in another research, in which 3D NPG is synthesized via CVD growth onto dealloyed nanoporous Cu templates.

The coarsening of pores and ligaments lead to extended pore size of 3D NPG and makes the synthesis uncontrollable. In fact, there are more problems with the dealloying. Firstly, readily prepared binary or ternary alloys with alloying elements in solid solution are required. Secondly, the etching process is time consuming particularly for bulk alloys due to the resistance of ions volumetric diffusion.

Therefore, it is hard to synthesize large-size nanoporous metals by dealloying.

Thirdly, etching of metals requires strongly acidic or oxidizing and toxic chemicals.

Fourthly, the etched wastes need complicated processes for recycle, resulting in increased costs. In addition, the etching process may introduce impurities from etchants and oxides into the porous structure. Thus, developing a controlled approach for synthesis of 3D NPG to prevent the coarsening of pores and ligaments of metallic templates, and an approach for versatile, controllable, and industrially waste-recycled synthesis of porous metal templates are significant in future.

Except using metallic templates, other metal oxides/ceramics are also used as templates for CVD growth of graphene.

Huang’s team developed high- temperature catalyst-free CVD growth of graphene onto porous SiO

materials.

Shi et al. and Tang employed CaO as templates for CVD growth of graphene at

~1,000 °C.

Besides, other oxides such as MgO are also investigated for

synthesis of nanoporous graphene.

The advantage of using metal oxides/ceramics

is lower costs of precursors, and stable ligaments and pores even at high

temperature. However, compared with the NPG derived from metallic templates,

the nanoporous graphene obtained using oxides/ceramics have poor quality such

as high defects and lower electric conductivity. The CVD growth onto oxides also requires much longer time than those using porous metals.

2.3 Low-dimensional graphene quantum dot

2.3.1 Synthesis

Up to date, the synthesis of GQDs contains many approaches, for example nanolithorgraphy, solvothermal or microwave-assisted oxidation cutting, pyrolysis, electrochemical exfoliation, chemical exfoliation, nanotomy assisted exfoliation, organic synthesis, decomposition of fullerene, etching and so forth. From an aspect of the starting precursors, the synthetic methods can be classified as top-down methods and bottom-up methods.

(1) Top-down methods

Top-down methods normally require readily available graphene, graphene derivatives or carbon materials containing few-layer graphitic carbon. Graphene quantum dots can be directly cut, extracted or exfoliated, and unfolded from the body materials by means of chemical or physical methods. The commonly used feedstock materials are graphene, graphene oxides, carbon nanotube, fullerenes, graphite, soot, coal, carbon fibres, activated carbon and pyrolysis carbon, which consist of graphene-like carbon segments.

Various top-down approaches are reviewed as following.

Electron-beam lithography was early developed to carve graphene flakes into graphene quantum dots or graphene nanoribbons.

The high-quality graphene obtained by mechanical exfoliation was transferred to desired substrates, and then patterned by electron-beam lithography into required geometries. The dimension of the carved GQDs can have the size of 20-100 nm, but minimum to 10 nm due to the resolution limitation of the state-of-the-art lithography. The as-carved GQDs of only a few nanometers in width remain conductive and reveal a confinement gap of up to 0.5 eV, illustrating the graphene-based molecular-scale electronics. Mohanty developed diamond-edge-induced nanoscale-cutting of graphite followed with exfoliation for producing nanostructured graphene quantum dots with size down to 15 nm.

The as-synthesized GQDs have low content of defects (I

/I

of 0.22–0.28) and roughness < 1 nm. The lithography and mechanical cutting method have advantages in producing or patterning the defined dimensions of GQDs, but the production yield is too low and the processing requires much care and sophisticatedly skills. The minimum size of GQDs also has processing limitations.

Another interesting work done by Lu et al. is unfolding C

spheres on ruthenium

catalyst into GQDs. Ru has strong interplay with C

induced by surface vacancies

in the Ru single crystal. At elevated temperatures, C

clusters dissociate on the

terrace and unfold to graphene quantum dots. The-as-obtained GQDs have size of

2.7-10 nm and various shapes. The triangular-shaped, parallelogram-shaped,

trapezoid-shaped and hexagon-shaped (with lateral dimension 5‒10 nm) GQDs exhibit an energy gap of 0.8, 0.6, 0.4 and 0.25 eV, respectively.

In contrast, chemical-based methods provide higher production yields and smaller sizes. A typical approach is hydrothermal or solvothermal assisted oxidative cutting of graphene oxides. Graphene oxides have oxygen-rich functional groups (C=O,–COOH, C–OH and C–O–C) at the edges and on the basal planes so that GO has good solubility in water. Under weak alkaline (pH= 8) conditions at 200 °C, micron-sized graphene oxides flakes can be cut into GQDs.

AFM results present that the diameters of GQDs are in the range of 5–13 nm and the thickness is in 0.5-2.0 nm, corresponding to 1~6 layers of graphene (shown in Figure 2.17).

The mechanism of the cutting process is suggested that the underlying C–C bonds of linear epoxy groups rupture to form an epoxy chain, where epoxy pairs are further oxidized to energetically favourable carbonyl pairs. By controlling the alkaline (pH > 12) conditions and using monolayer GO sheets as precursors, the size of GQDs can be cut down to 1.5–5 nm and the thickness is only in few layers.

Figure 2.17 (a) Mechanism of the hydrothermal cutting of graphene oxides into GQDs: a mixed epoxy chain consisting of epoxy and carbonyl pair groups (left) ruptured under the hydrothermal treatment and led to a complete cut (right).

(b) An AFM image of the GQDs synthesized by the solvothermal route and their height distributions.

In addition to cutting, chemical exfoliation can be used for precursors with few oxygen functional groups and high crystallinity. Peng et al. chemically exfoliated GQDs from carbon fibres in H

SO

/HNO

mixtures at 80-120 °C.

The GQDs have different sizes and emission colors excited at 318, 331, and 429 nm, respectively.

With decreasing the reaction temperature, the size increases whereas the band gap

decreases from 3.9 eV to 2.89 eV. Other carbon materials such as coal, activated

carbon, carbon black, pyrolyzed self-assembled hexa-peri-hexabenzocoronene were

also performed for synthesizing GQDs chemical exfoliation in highly oxidative acid

systems.

Lin et al. developed a method of potassium–intercalation assisted

exfoliation of graphite.

They firstly synthesized potassium–graphite intercalation

compounds, followed with reaction between compounds with ethanol and then exfoliated the graphite by means of the generated hydrogen gas. The violent combustion breaks the graphene layers into GQDs with the assistance of ultrasonication. AFM results indicate that the as-synthesized GQDs have an average size of around 20 nm and are 1-2 layers thick. Although the above chemical exfoliations are effective, the reaction conditions are very extreme and toxic. The as-synthesised GQDs contain rich defects and functional groups. In a further study, a greener way is developed by exfoliation of graphite nanoparticles in water- ethanol system for producing ~4 nm GQDs with pure sp

carbon crystalline structure.

However, the yield is very low. Thus, developing a fast, low-cost, large- scale and environmentally friendly method for producing defects controllable GQDs is a big challenge.

(2) Bottom-up methods

Theoretically, graphene quantum dot can be synthesized by rigid chemical synthesis of polycyclic aromatic hydrocarbon molecules. Müllen’s group conducted extensive researches on the atomically precise fabrication of nanographenes through synthetic organic chemistry.

A typical synthetic protocol is as following, using 1,4-diiodo-2,3,5,6-tetraphenylbenzene to synthesize hexaphenylbenzene, followed with twofold lithiation and reaction with 2-isopropoxy- 4,4,5,5- tetramethyl-[1,3,2] dioxaborolane, led to the bis-boronic ester, which reacted with diiodobenzene by Suzuki-Miyaura polymerization and generate polyphenylenes.

Finally, the polyphenylenes converted to graphene nanoribbons on basis of intramolecular Scholl reaction.

During synthesis, the aggregation of graphene moieties can affect to size expansion. To overcome this problem, Yan et al. used multiple 2′,4′,6′-triakyl phenyl groups covalently attached to the edges of the graphene moieties to stabilize the resultant graphene. The face-to-face interaction between the graphene dots is moderated, leading to stable and large graphene quantum dots.

Up-to-date, the synthesis of nanographenes can be categorised as non-conventional methods, structures incorporating seven- or eight-membered rings, selective heteroatom doping, and direct edge functionalization. Recent engineering of robust topological quantum phases on atomically precise graphene nanoribbons present the significance of chemical synthesis of nanographenes.

In comparison with organic synthesis, pyrolysis of carbohydrates can more

facilely generate GQDs with various size and photoluminescence properties. The

pyrolysis reaction can occur in gases or in a liquid system. For example, Dong

synthesized GQDs of 15 nm in width and 0.5–2.0 nm in thickness by pyrolysis of

citric acid.

The as-synthesized GQDs exhibit blue emission. Tang et al. reported

that the pyrolysis of sugar using microwave-assisted hydrothermal method can

generate GQDs with lateral size of ~3.4 nm.

Other carbohydrates or biomass such

as L-glutamic acid, coffee grounds were also investigated as resources for synthesis

of GQDs.

It is suggested that most of the carbohydrates, which contain C, H,

and O, can be used to prepare GQDs under heating or hydrothermal conditions.