Preparation and characterization of graphene

CHAPTER 3

Preparation and characterization of graphene Contents 3.1 3.2 3.3 3.4 3.5 3.6 3.7 3.8 3.9

Mechanical exfoliation Epitaxial growth Chemical vapor deposition Plasma-enhanced chemical vapor deposition and pulsed laser deposition Wet exfoliation Synthesis of graphene oxide Reduction of graphene oxide Graphene-based composites Analysis and characterization 3.9.1 Optical imaging 3.9.2 Fluorescence quenching 3.9.3 Atomic-force microscopy 3.9.4 Raman spectroscopy 3.9.5 X-ray photoelectron spectroscopy 3.9.6 Transmission electron microscopy 3.10 Summary References Further reading

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3.1 Mechanical exfoliation As follows from the theoretical studies discussed in Chapter 2, the successful utilization of graphene in tribological applications raises some specific requirements not only to the structure of graphene itself but also to the whole “graphene-substrate” system. The critical factors for the graphene structure are the graphene defectiveness, the number of layers, the grain size, the adhesion to the substrate, the stiffness of the substrate, etc. They are defined by a specific method of preparation/synthesis of graphene. Since the first invention of mechanical exfoliation of graphene from pyrolytic graphite in 2004, many approaches have been developed to fabricate Tribology of Graphene https://doi.org/10.1016/B978-0-12-818641-1.00003-4

© 2020 Elsevier Inc. All rights reserved.

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single- and multilayer graphene on different substrate materials. The variety of approaches include techniques such as chemical vapor deposition, liquidphase exfoliation, microcleavage, and oxidization reduction, as illustrated in Fig. 3.1A. The synthesis methods can be divided into two main groups, called top down and bottom up. Top-down approaches are grounded on the exfoliation of graphite. The first historical preparation method of isolated graphene specimens was mechanical exfoliation, proposed by Novoselov and Geim [1]. This method allowed preparing multilayer or single-layer graphene that would be suitable for device fabrication and the study of electronic properties. According to this method, graphene films were made by repeated peeling of small mesas of highly oriented pyrolytic graphite (Fig. 3.1B). This method is also well known as the “Scotch tape method” or microcleavage [2, 3]. It allows preparing films up to 20 μm in size [4]. Despite its initial simplicity, the Scotch tape method was further improved by various researchers. For example, Peng et al. proposed their own modified plan [5]. In their work, a fresh piece of adhesive tape was firmly pressed sticky-side down onto the HOPG surface for about 30 s. Then, the tape was peeled away with thick layers of graphene stuck to it. The part of the tape with graphite layers was refolded on a clean adhesive section of the same piece of tape. After that, the tape was unfolded and the mirrored layer of graphite remained on it. This process was repeated until a large portion of the tape became dark gray. Then, the graphene was transferred to the SiO2/Si substrate by pressing the tape with graphene flakes on it. Zhang et al. developed a micromechanical method to extract fragile graphene samples [6]. Initially, micropillars were fabricated on the HOPG surface using micropatterning, followed by masked anisotropic oxygen plasma etching. Then, individual pillars were removed from the substrate and attached to a micromachined silicon cantilever by a small amount of ultraviolet-sensitive epoxy. In the next step, cantilevers with graphite samples were mounted at the tips of the AFM device for transfer onto an SiO2/Si substrate. AFM tip slid in contact mode against the surface, leading to the shearing of graphene layers onto the substrate. The thickness of the graphene layers was in the range of 10–100 nm, and the lateral size was around 2 μm. These specimens were suitable for electrical experiments. The original Scotch tape method appeared to be extremely simple and effective and allowed the fast growth of graphene-related studies. It has a low entry barrier and does not require a significant investment or complicated

Preparation and characterization of graphene

CVD

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Mechanical exfoliation (research, prototyping)

Quality

(coating, bio, transparent conductive layers, electronics, photonics)

SiC (electronics, RF transistors)

Molecular assembly (nanoelectronics)

Liquid-phase exfoliation (coating, composites, inks, energy storage, bio, transparent conductive layers) Price (for mass production)

(A)

(B) Fig. 3.1 (A) Summary of graphene production methods based on price and quality estimation. (B) The micromechanical cleavage technique (“Scotch tape method”) for producing graphene. Top row: Adhesive tape is used to cleave the top few layers of graphite from a bulk crystal of the material. Bottom left: The tape with graphitic flakes is then pressed against the substrate of choice. Bottom right: Some flakes stay on the substrate, even after removal of the tape. (A: Reprinted by permission from K.S. Novoselov, V.I. Fal’ko, L. Colombo, P.R. Gellert, M.G. Schwab, K. Kim, A roadmap for graphene, Nature 490 (2012) 192, https://doi.org/10.1038/nature11458. Copyright (2012); B: Reprinted from K.S. Novoselov, Nobel lecture: graphene: materials in the flatland, Rev. Mod. Phys. 83 (2011) 837. Copyright (2011) by the American Physical Society.)

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equipment. These factors have helped considerably to broaden the geography of graphene science [3]. This method allowed preparing high-quality graphene specimens with a low number of structural defects [1, 6–8]. Nonetheless, even though this method of mechanical exfoliation has led to numerous discoveries regarding the mechanical properties of graphene, it has two significant drawbacks: low production efficiency and the small size of graphene flakes. Because these limitations are crucial for the tribological applications of graphene, it was used only in some initial tribological studies [7].

3.2 Epitaxial growth Besides mechanical exfoliation and chemical reduction methods, several promising approaches for producing graphene sheets have been reported. In particular, they are chemical vapor deposition (CVD) and epitaxial growth from SiC [9]. Growth from SiC, which is also called SiC sublimation, has been attractive, first of all, for the semiconductor industry because it does not require transfer to another substrate for assembling the final devices [10, 11]. In this process, the sublimation of Si atoms from the substrate occurred during high-temperature heating of the SiC substrates under UHV conditions. As a result, the surface carbon atoms were rearranged into graphene layers. The thickness of the graphene layers depends on the temperature and processing time. The formation of few-layer graphene typically required annealing of the SiC surface at a temperature around 1200°C for several minutes [12]. Even though the synthesis of graphene on SiC substrates looks promising, several substantial disadvantages still limit its application. In particular, it was challenging to control the number of graphene layers. Another uncertainty involved the different epitaxial growth patterns on different SiC polar faces (i.e., Si-face or C-face). Also, the lattice mismatch caused abnormal rotational graphene stacking. The lattice mismatch can lead to delamination between different layers of graphene [11]. Moreover, it is very challenging to grow uniform graphene over a large area due to the occurrence of step bunching on the SiC surface [13]. The issue of nonuniformity was overcome by the gown of graphene on off-axis 3C-SiC(111), which is one of the polytypes of silicon carbide. 3C-SiC substrates could be synthesized in many ways in the form of bulk substrates or coatings [14]. It was demonstrated that the step bunching was eliminated

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during graphene growth due to the synergistic effect of periodic SiC step edges as graphene nucleation sites and the unique thermal decomposition energy of 3C-SiC steps [13]. Besides, several studies were focused on the optimization of synthesis conditions for conventional SiC substrates. For instance, high-quality epitaxial graphene was produced by flash annealing of 6H-SiC in a lead (Pb) atmosphere [15]. It was found that the three top bilayers of SiC are decomposed due to fast heating to
1400°C and cooling. In this case, the sublimation of Si atoms from SiC was accelerated by the Pb atmosphere.

3.3 Chemical vapor deposition While top-down methods are grounded on the exfoliation of graphite, bottom-up approaches use carbon molecules as building blocks. Usually, these molecules are obtained from external sources such as carbonaceous gases. CVD is related to bottom-up methods. It is still one of the most promising approaches for the large-scale production of graphene [7]. The successful synthesis of few-layer graphene films using thermal CVD was first reported by Somani et al. [15]. Films were grown on Ni foils using camphor pyrolysis with two horizontal furnaces. Camphor is evaporated in the first furnace at 180°C and pyrolyzed in the second furnace at 700–850°C with argon as the carrier gas. Ni sheets were kept on the alumina boat in the center of the second furnace. This method allowed covering relatively large areas (2  2 cm2 in the first experiments). This study opened the possibility of performing controlled and large-area synthesis of graphene [16]. Despite the first success, CVD suffered several severe downsides and unresolved issues such as the need for minimization of graphene folding and control of the number of layers. However, since the initial report, much progress has been achieved, allowing the deposition of a controlled number of graphene layers on various metal substrates. Moreover, it was demonstrated that graphene can be relatively easily transferred to other substrates after CVD synthesis. The transfer technique can provide high-quality graphene layers without additional mechanical or chemical treatments [16]. For example, Bae et al. demonstrated a roll-to-roll production of graphene films using the CVD method [17]. Initially, graphene was grown on 30-in. copper foils. Then, the copper foil was covered with a polymer layer to support the graphene. In the next step, the copper foil was etched, and the graphene layers were released and transferred onto a target substrate. A schematic of such a process is shown in Fig. 3.2.

Cu etchant

(A)

(C)

Graphene on target

(D)

Fig. 3.2 (A) Schematic of the roll-based production of graphene films grown on a copper foil. The process includes the adhesion of polymer supports, copper etching (rinsing), and dry transfer printing on a target substrate. Wet chemical doping can be carried out using a setup similar to that used for etching. (B) Copper foil wrapping around a 7.5-in. quartz tube to be inserted into an 8-in. quartz reactor. The lower image shows the stage in which the copper foil reacts with CH4 and H2 gases at high temperatures. (C) Roll-to-roll transfer of graphene films from a thermal release tape to a PET film at 120°C. (D) A transparent ultra large-area graphene film transferred on a 35-in. PET sheet. (Reprinted by permission from S. Bae, H. Kim, Y. Lee, X. Xu, J.S. Park, Y. Zheng, et al., Roll-to-roll production of 30-inch graphene films for transparent electrodes, Nat. Nanotechnol. 5 (2010) 574, https://doi.org/10.1038/nnano.2010.132. Copyright (2012).)

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Released polymer support

Target substrate

Graphene on Cu foil

(B)

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Graphene on polymer support

Polymer support

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Depending on the substrate material used in the CVD process, the formation of graphene films may occur through two different mechanisms— reinfusion or surface activation on the catalyst. In the case of metals having intermediate or high solubility of carbon (>0.1 atomic%), such as Co and Ni, the synthesis occurred by reinfusion and involved two stages. The first stage is the diffusion of carbon atoms into the thin metal film during the high-temperature CVD process. Carbonaceous gaseous species such as CH4 decomposed on the hot surface, providing the source of carbon for the first stage. The diffusion was followed by the second stage, which is the precipitation of carbon from the bulk to the surface of the metal during cooling [18]. The surface activation occurred at high temperatures (900–1100°C) on the surfaces of thin metal films or foils. This mechanism has been primarily attributed to the limited solubility of carbon in some metals. In particular, it is realized in metals having very low carbon solubility (<0.001 atomic%) such as Cu [19]. In this case, decomposition of the carbonaceous precursor on the hot surface and accumulation of carbon atoms led to nucleation of the graphene lattice. Fig. 3.3 illustrates the typical processes that occur during the surface-activation process [20], which mainly happens through a fourstep process as follows [21]: 1. Decomposition of methane from CxHy on a Cu catalyst while Cu is exposed to a methane/hydrogen atmosphere. The Cu surface may be undersaturated, saturated, or supersaturated with CxHy species. It depends on the process parameters such as temperature, methane flow, and the partial pressures of methane and hydrogen. 2. Nuclei formation due to local supersaturation of Cu with CxHy. The nuclei formation is possible only in the case of saturation or supersaturation of the Cu substrate. 3. Nuclei grow to form graphene islands on the Cu surface. 4. The Cu surface is fully covered by graphene after the conjunction of the graphene islands. There is an example of the classical growth process and conditions: 1. Loading the Cu substrate into the furnace, evacuation down to
1 Pa, backfilling to atmospheric pressure with H2 or H2/Ar mixture, heating to 1000°C, and maintaining an H2 flow at 50 sccm for 1–2 h. 2. Introduction of 35 sccm of CH4 for the desired growth time (5–20 min). 3. Slow cooling of the furnace to room temperature in the hydrogen atmosphere [19, 22].

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Fig. 3.3 Schematic representation of fundamental processes during epitaxy: hydrocarbon molecules E are deposited on the surface, undergo decomposition via a series of dehydrogenation reactions giving rise to various CxHy species shown as Ed, and H atoms. The new species are all able to diffuse across the surface. Smaller carbon species M and D form and diffuse on the surface, aggregating into larger clusters C. H atoms of the original molecule migrate on the surface and form H2 molecules that evaporate from the surface. Finally, some of the species such as M and D, or even their bigger clusters C, may attach to the island G at its edge. Other processes (not shown) are also possible based on, e.g., diffusion of atoms along the edge of an island, nucleation of the second and higher layers on the islands, downward movement of atoms adsorbed on top of islands to lower layers, and the breakup or dissolution of islands. (Reprinted from H. Tetlow, J. Posthuma de Boer, I.J. Ford, D.D. Vvedensky, J. Coraux, L. Kantorovich, Growth of epitaxial graphene: theory and experiment, Phys. Rep. 542 (2014) 195, https://doi.org/10.1016/j.physrep. 2014.03.003. Copyright (2014), with permission from Elsevier.)

The mechanism of graphene growth on Cu was found to be self-regulated. The self-regulation here means that increasing deposition time led to the saturation of the growth process. In the early stages, the formation of nonuniform distribution graphene flakes having a various number of layers and a high concentration of defects usually occurred. While the growth time increased, the defect concentration reduced, and the distribution of graphene thickness became uniform [22]. In other words, the maximum number of graphene layers is not significantly affected by the growth time, but increasing the growth time improves the spatial uniformity and the structure of the coating. One of the essential deposition parameters of CVD is the processing pressure. In this regard, atmospheric-pressure chemical vapor deposition (APCVD) and low-pressure chemical vapor deposition (LPCVD) should be distinguished. In general, APCVD and LPCVD require different temperatures for optimal graphene growth. A comprehensive experimental

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comparison of these methods and an investigation of the pressure effect on the structure of graphene grown on Cu foils was performed by Ullah et al. [23]. It was demonstrated that reducing the pressure led to the growth of more uniform graphene with a smaller number of layers (up to three). Graphene deposited under the atmospheric pressure had a larger number of layers and significant uniformity. These data are in good agreement with the results of other studies [22]. Thus, using APCVD seems to be preferable for the synthesis of high-quality graphene. Besides the processing time and pressure, the quality of graphene is strongly affected by the orientation of grains in a metallic substrate. For example, it was demonstrated that the graphene was weaker on a Cu(100) substrate than on a Cu(111) [24]. Being grown on Cu(111), graphene formed a high-quality and microscopically uniform sheet. Graphene grown on Cu(100) under the same conditions was not able to create the uniform sheet and displayed uncovered step edges. Because increasing the graphene sheet/grain size and minimization of the uncovered edges are the critical requirements for high tribological performance, controlling the structure of the copper substrate became crucial. Managing the copper structure could be performed in several ways, including using liquid Cu substrates and managing the Cu grain orientation through preoxidation or zone annealing. Employing a liquid Cu surface eliminates the grain boundaries in solid polycrystalline Cu, resulting in a uniform nucleation distribution and low graphene nucleation density; it also enables the self-assembly of graphene flakes into compact and ordered structures [25]. As a result, the growth of large-area uniform monolayer graphene was achieved. Later, the synthesis of single-crystalline multilayer graphene domains on liquid copper by APCVD was also demonstrated [26]. In this case, the domains had an umbrella-like shape with diameters up to 0.6 mm. It was shown that incensing the hydrogen flow reduced the nucleation density of graphene, leading to the formation of larger domains. Moreover, due to the immersion of a grown graphene flake into liquid copper, the growth of the uniform flakes was provided because carbon atoms could reach each of the graphene layers. In general, the graphene growth mechanism on a metal catalyst depended on several factors, including crystal structure, lattice parameters, the limit of the carbon solubility in the metal, and thermodynamic parameters such as the temperature and pressure of the system [16]. It involved a series of complex interactions between the gas atmosphere and the Cu substrate. Various structural irregularities and impurities of a Cu surface resulted

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in heterogeneous nucleation of graphene, and its uncontrolled growth [27]. Another strategy for controllable nucleation of graphene could be plasma treatment of the Cu foils. The air plasma treatment allowed removing impurities from the Cu surface, leading to reducing the active sites for single-crystal graphene nucleation [28]. Moreover, the plasma-induced formation of CuO nanoparticles allowed reducing the nucleation density of graphene and promoting rapid growth. Following this approach, the synthesis of hexagonshaped single-crystal low-defective graphene with a size of
5 mm was demonstrated by Cheng et al. [28]. In 2019, a synthesis of large-area single-layer graphene on a Cu substrate using a refined cooking palm oil was demonstrated [29]. For this purpose, a spray injector-assisted CVD was used. The deposition was performed under atmospheric pressure in an Ar-N2-H2 atmosphere. In general, the process diagram was quite similar to the conventional CVD process described above. After annealing in hydrogen, the small amount of oil (
0.05 mL) was injected into the processing chamber. Then the growth was performed at the set temperature of 900–1000°C for 10 min. After the growing stage, the specimen was cooled in a vacuum. It was found that a 1 cm substrateto-nozzle distance and a substrate temperature of 950°C led to the growth of large-area single-layer graphene with coverage up to 97% of the measured area size of 6400 μm2. However, graphene grown under these conditions had a relatively high concentration of structural defects. Using a flashcooling technique was suggested by the authors. Instead, this method could be further improved by combination with one of the substrate pretreatment methods discussed above, such as copper peroxidation or plasma treatment. Another synthesis method of aligned large-area single-crystal graphene flakes was proposed by Reckinger et al. [30]. This method was based on several effects caused by residual oxygen on the Cu surface. A slight oxidation of the Cu surface led to pinning of the grain boundaries and freezing of the thermal recrystallization. Subsequent reduction of Cu in an H2 atmosphere led to the preferential growth of centimeter-sized copper (111) grains through the mechanism of abnormal grain growth. Besides the formation of preferable (111) orientation of Cu grains, the preoxidation of Cu significantly reduced the nucleation density of graphene. The temperature-time diagram shown in Fig. 3.4 summarizes the so-called “standard” (small, misaligned domains) and “novel” (macroscopic, aligned domains) graphene growth conditions, with the corresponding argon and hydrogen flows [30]. The positive effect of thermal preoxidation of Cu foils before CVD growth on reducing the nucleation density was later confirmed and theoretically explained by Liang et al. [27].

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Fig. 3.4 (A) Temperature-time diagram summarizing the different steps of the standard (in blue) and aligned (in red) graphene growth conditions with the corresponding argon and hydrogen flows. Scanning electron microscopy pictures of Cu foils after (B) S#1:H2/ S#2:H2 (highlighted in blue) or (C) S#1:H2/S#2:H2 (highlighted in red). (Reproduced from N. Reckinger, X. Tang, F. Joucken, L. Lajaunie, R. Arenal, E. Dubois, et al., Oxidation-assisted graphene heteroepitaxy on copper foil, Nanoscale 8 (2016) 18751, https://doi.org/10. 1039/C6NR02936A, with permission from The Royal Society of Chemistry.)

Finally, the synthesis of meter-sized single-crystal graphene on copper foil was presented in 2017 by Xu et al. [31]. This result was achieved by epitaxial growth of graphene islands on the Cu(111) surface. First, the copper foil was annealed to form the meter-sized Cu single crystal. Then, graphene was grown by the modified CVD method. Initially, the formation of graphene islands occurred. Using of the Cu single-crystalline substrate led to the uniform crystallographic orientation of the graphene islands. Further coalescence led to the creation of a graphene film with high single crystallinity. Besides Ni and Cu, the CVD growth of graphene was demonstrated for various 3d-5d transition metals. For instance, graphene layers were grown epitaxially on the (0001) faces of ruthenium (Ru) crystals under UHV at high temperatures. In this case, sparse nucleation of graphene allowed the growth of macroscopic single-crystalline domains [32, 33]. Producing graphene layers by CVD was also demonstrated on Rh, Ir, Re, and Pt. [20, 34]

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The synthesis of large-area graphene films on Mo films by APCVD was reported in 2012 [35]. Mo foils were preliminarily annealed in an Ar/H2 atmosphere for removing surface oxides. It was demonstrated that dissolution and segregation governed growth. The cooling rate was the key parameter defining the thickness of graphene films. By decreasing the cooling rate from 10 to 1.5°C/s, the number of graphene layers increased from one to three. Besides bulk Mo foils, single-crystal epitaxial Mo(110) grown on Al2O3 was also used for graphene synthesis. It was proved that removing Mo grain boundaries would significantly improve the quality of graphene on Mo because the grain size of the substrate limits the continuity. In 2014, the wafer-scale CVD growth of graphene on sputtered thin films of molybdenum was proposed [36]. The high melting point of Mo, along with the smooth surface and low thermal expansion, created potentially favorable conditions to produce high-quality, large-area, wrinkle-free graphene. Using very durable Mo instead of Cu opened new possibilities for tribological applications. Due to issues associated with transferring CVD-grown graphene on target substrates, significant efforts were directed to the increasing number of materials available for direct growth by CVD on target substrates such as germanium single-crystals, SiO2 thin films, sapphire, quartz, and glasses [37]. In 2015, a new synthesis approach of epitaxial graphene on transition metal oxide was demonstrated [38]. Initially, C60 was deposited on Ti/SiC. Next, Ti metal and C60 were converted into TiC by carbon reduction. Finally, by O2 intercalation of the graphene/TiC/SiC structure, the final graphene/ TiO2/SiC heterostructure was obtained. Depending on the substrate materials, metal catalyst layers were used to activate graphene growth. Then, the metallic films were etched out, and the graphene layer remained on the substrate. Metal-free direct synthesis of graphene on silica glass substrates by semiatmospheric pressure CVD with methane as the carbon precursor was demonstrated by Barbosa et al. [37]. This method allowed avoiding graphene contamination by a catalyst and minimizing the number of defects induced during chemical etching. Nonetheless, the metal-free CVD growth of graphene often suffers from poor structural uniformity and a slow growth rate caused by the negligible catalytic activity of dielectric substrates. A water-assisted CVD process was developed by Wei et al. to overcome these issues [39]. This method allowed the rapid growth of monolayer graphene film on SiO2/Si substrates without using metal catalysts or ultrahigh temperatures. It was demonstrated that the

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presence of a small amount of water enabled the preferential formation of high-quality monolayer graphene. This phenomenon was attributed to effective reducing of the growth kinetic barrier. Besides, water accelerated the release of oxygen from the SiO2 substrate. The large-scale, metal-free synthesis of graphene on a 200 mm Si wafer was demonstrated by Lukosius et al. [40]. Initially, the 2 μm Ge layer was deposited onto the Si substrate. Then, the CVD process was carried out at a deposition temperature of 885°C and a pressure of 700 mbar. CH4 was used as a source of carbon and Ar/H2 mixture as the carrier gas. The optimized deposition time was found to be 60 min for these conditions. After deposition, graphene sheets with sizes up to 2  2 cm2 were successfully transferred onto another substrate.

3.4 Plasma-enhanced chemical vapor deposition and pulsed laser deposition Plasma-enhanced chemical vapor deposition (PECVD) was a further development of the CVD method. Using plasma assistance allowed decreasing the deposition temperature to
350°C. The simultaneous study showed that the substrate does not play any role in the dissociation of the hydrocarbon precursor during PECVD [41]. Thus, this method could be applied to a wide variety of substrate materials. The effect of the substrate, in this case, is just limited to different adsorption and the dissociation of the hydrocarbon precursor. Thus, various substrates should exhibit different growth rates, and the optimal process temperature also could be different. For the first time, a few layers of graphene were produced by a radio frequency PECVD from a CH4/H2 gas mixture and a substrate temperature of
700°C [42]. A wide variety of substrate materials was successfully used, including quartz, silicon, nickel, platinum, germanium, titanium, tungsten, stainless steel, tantalum, and molybdenum. Before graphene synthesis, substrates were preheated in hydrogen plasma to the target temperature. The synthesis was performed under a reduced pressure of 40 Torr. The growth time was in the range of 1–3000 s. Opposite to the conventional CVD, PECVD is not self-regulated. Increasing the processing time led to the sequenced creation of the smooth layer on the surface, the formation of the curling crack edges, and the growth on the free-standing graphene on these edges. Depending on the substrate material, graphene demonstrated different adhesion to the substrate due to the formation of carbides in some cases.

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Later, significant effort was directed to understand the graphene growth mechanism and to optimize experimental conditions to control the thickness of the graphene film. As a result, successful graphene synthesis was demonstrated on quartz, silicon, and platinum substrates [43]. It was found that the initial growth stages were dependent on the substrate material. In particular, a base multilayer graphene layer with a large concentration of irregular cracks was formed in the Pt substrate. Then, the growth direction was changed from parallel to perpendicular on the crack edges. Finally, the formation of free-standing graphene flakes occurred. It happened similarly, as was observed by Malesevic et al. [42]. In the case of the quartz substrate, a monolayer of the amorphous carbon layer was firmed first. Then, it was covered by the monolayer of graphene. Finally, the formation of vertical flakes occurred. As for the Si substrate, the SiC layer was formed first, and it was covered with graphene during the next stage, followed by the formation of vertical graphene flakes. Zhang et al. demonstrated for the first time that large-area graphene single crystals could be grown on h-BN using PECVD [44]. In this case, the synthesis was performed at 500 °C. Initially, h-BN flakes were mechanically exfoliated and transferred onto SiO2/Si substrates using the Scotch tape method. Then, graphene was grown on the BN surface. This approach is interesting because it could open new ways for direct synthesis of graphene on various substrates. Later, the deposition of high-quality graphene at a reduced temperature of 350°C was demonstrated [23]. In this case, high-temperature annealing of Cu substrates, which was usually required for removing oxide layers, was not necessary because oxides were forcefully removed by plasma at the same temperature. In general, it is worth noting that the PECVD growth mechanism involves a balance between graphite deposition and etching caused by hydrogen and argon ions [16]. Thus, the number of defects in the graphene could be increased, leading to a negative influence on the mechanical and tribological properties of coatings [22]. An exciting option of synthesizing graphene by the CVD is the possibility of substitutional doping of graphene. It is performed by introducing other gases such as N2 or NH3 during the process [7]. PECVD was also found to be an efficient graphene doping method, and it has been applied to the direct synthesis of doped graphene by various research groups. Pulsed laser deposition (PLD) is another prominent method for the synthesis of graphene coatings [45]. It allows growing few-layer graphene sheets on different metal substrates (Ni, Cu, Co, and Fe). The quality of graphene

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deposited by PLD is close to that obtained by CVD, but the deposition temperature can be decreased to 300 K. This method also could be combined with other deposition techniques. Such a combination may open new ways of synthesis for graphene and its derivatives. For example, a fresh deposition approach of N-doped few-layer graphene synthesized by femtosecond pulsed laser ablation from a solid-state nitrogen carbide a-C:N source was proposed by Maddi et al. [46]. A femtosecond laser beam was focused on a graphite target to form a carbon-based plasma expanding in an atmosphere of gaseous nitrogen. The 10 nm thick a-C:N film was next covered by a nickel film (150 nm) by vacuum thermal evaporation. The a-C:N/Ni sandwich is then heated to 780°C in a vacuum followed by natural vacuum cooling, thereby forming few-layer N-doped graphene on the top surface.

3.5 Wet exfoliation Besides using continuous graphene films, there is another approach based on the utilization of single- and few-layer graphene flakes with a size of several micrometers. These flakes also could be used for producing coatings or as fillers/reinforcements for various composite materials. In the case of coatings, the overlapping of chips allowed producing pseudocontinuous films. The graphene flakes could be prepared using different methods such as chemical exfoliation, which is based on the penetration of chemical reactants between graphene sheets of graphite and distortion of weak Van der Waals bonds. The use of graphite, graphite intercalation compounds, or expandable graphite allows the production of dispersions of high-quality graphene sheets [47]. For example, a colloidal suspension of unoxidized graphene sheets in organic solvents such as N-methylpyrrolidone (NMP) was obtained by sonication of graphite powder [48]. Electrochemical treatment of graphite was also used to generate a colloidal suspension of graphene flakes [49]. The method of graphene preparation using pyrene-based molecules was proposed by An et al. [50]. This method allowed an effective exfoliation of single-, few-, and multilayer graphene flakes into stable aqueous dispersions. In this case, the pyrene-based molecules acted as molecular wedges, according to the MD simulation results discussed in Section 2.4. This method allowed obtaining a high yield of graphene in water without residual graphite flakes in the dispersion. Electrochemical exfoliation of graphene sheets into aqueous solutions of different inorganic salts ((NH4)2SO4, Na2SO4, K2SO4, etc.) was proposed

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by Parvez et al. [51]. It was demonstrated that exfoliation in these electrolytes allowed producing graphene with a large lateral size (up to 44 μm), a high yield (>85%, 3 layers), and a low oxidation degree. The suspensions prepared by various exfoliation methods could be further used for deposition of graphene films by multiple techniques such as vacuum filtration, drop casting, spray coating, spin coating, or electrospray coating [52, 53].

3.6 Synthesis of graphene oxide The chemical conversion of graphite to graphene oxide was developed as an affordable production method of graphene-based single sheets in considerable quantities. GO can be produced from graphite using various oxidants such as concentrated nitric or sulfuric acids or potassium permanganate. The formation of GO from graphite includes the three following independent steps. The first step is the conversion of graphite into a graphite intercalation compound (GIC). The second step is the conversion of the GIC into oxidized graphite, which was defined as pristine graphite oxide (PGO). This step involves diffusion of the oxidizing agent into the preoccupied graphite galleries. This rate-determining step makes the entire process diffusivecontrolled. The third step is the conversion of PGO into conventional GO after exposure to water. This step involves the hydrolysis of covalent sulfates and the loss of all interlayer bonding [54]. In the case of the classical Hummers’ method [55], graphite oxidation was achieved by the harsh treatment of one equal weight of graphite powders in a concentrated H2SO4 solution containing three equal weights of KMnO4 and 0.5 equal weight of NaNO3. The Hummers’ method has significant advantages over previous techniques: (1) the reaction can be completed within a few hours; (2) KMnO4 was used to improve the reaction safety, avoiding the evolution of explosive ClO2; and (3) the use of NaNO3 instead of fuming HNO3 eliminates the formation of acid fog [56]. It was demonstrated that the appropriate oxidation time of graphite depended on the size and shape of the primary particles [57]. While the oxidation of the large agglomerations was completed in several days, the oxidation of flake powder was achieved within just 2 h. Thus, the oxidation time could be decreased by reducing the particle size (e.g., by grinding). This means that the small and flakey particles are more suitable for large-scale GO preparation. The oxidation degree of GO produced by modified Hummers’ methods could be controlled by adjusting the parameters of the process, such as the temperature and duration of different steps [58].

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Initially, the Hummers’ method received intense attention due to its high efficiency and satisfying reaction safety. However, it still has several significant drawbacks: (1) the oxidation procedure releases toxic gases such as NO2 and N2O4; and (2) the residual Na+ and NO3 ions are difficult to remove from the wastewater formed from the processes of synthesizing and purifying GO [56]. Thus, significant efforts were directed to developed ecofriendly methods of GO synthesis. Dimiev and Tour [54] found that excluding the NaNO3, increasing the amount of KMnO4, and performing the reaction in a 9:1 mixture of H2SO4/H3PO4 allowed improving the efficiency of the oxidation process. This improved method provides a more massive amount of hydrophilic oxidized graphene material as compared to the classical Hummers’ method or Hummers’ method with additional KMnO4. Chen et al. [56] also proposed an improved Hummers’ method without using NaNO3 for the synthesis of GO. This enhanced method eliminates the generation of toxic gases and simplifies the procedure of purifying waste liquid that allowed decreasing the cost of GO synthesis. Further improvement was achieved by Zaaba et al. [59]. In this method, the GO also was synthesized without NaNO3 and an ice bath, allowing the process to be carried out at room temperature. The size and number of layers are essential parameters that define the further applicability of GO. Typically, it allows producing GO flakes with a scale in the range from several micrometers to several hundred micrometers. Improvements of the Hummers’ method allowed producing single-layer graphene oxide with a high yield of 171  4% [60]. The solubility and stability of GO was significantly affected by the reduction process. Solutions of GO in NMP, ethylene glycol, and water presented significant long-term stability with solubility values reaching
8.7 μg/mL for NMP. Also, the dispersion behavior of GO changed after its reduction, introducing better interaction with solvents such as o-DCB (
9 μg/mL) and CN (8.1 μg/ mL) [61].

3.7 Reduction of graphene oxide GO could be further converted to rGO by chemical reduction using various reducing agents, including hydrazine and sodium borohydrate [47, 62]. Hydrazine was the most efficient reductant, but due to its toxic nature, there is a high demand to use green reductants for RGO synthesis. Thus, many efforts were directed to design green and ecofriendly processes of GO reduction [63]. Green reducing agents include organic acids, plant extracts, sugars,

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proteins, amino acids, and even microorganisms. For instance, the green reduction of GO using two different phytoextracts of Mangifera indica L. (dry mango leaf ) and Solanum tuberosum L. (potato) was demonstrated [64]. It was observed that the polyphenols in the plant extract act both as reducing and stabilizing agents. Thus, the resultant graphene possesses good solubility and stability in an aqueous medium. Nevertheless, among all the green reductants, ascorbic acid-AA (vitamin C) has achieved primary interest as it has proven to be an excellent alternative in synthesizing RGO, challenging the toxic yet potent reductants such as hydrazine [63]. Besides chemical methods, GO could be reduced physically, for instance, using UV irradiation or thermal reduction. For example, reducing GO carried out by UV irradiation of GO dispersion in the presence of N,Ndimethylformamide (DMF) was demonstrated by Wu et al. [65]. It was found that a stable dispersion of rGO was produced by UV irradiation of a GO solution in the presence of DMF. The rapid and mild thermal reduction of GO with the assistance of microwaves in a mixed solution of N,Ndimethylacetamide and water was demonstrated by Chen et al. [66]. The whole variety of graphene synthesis methods was summarized by Raccichini et al. [67] in the diagram shown in Fig. 3.5. Selection of the process plays a crucial role in determining the properties of the final product (Table 3.1). For example, the use of methods such as mechanical exfoliation, epitaxial growth on SiC, and bottom-up approaches are still mostly limited to fundamental research due to limited scalability and high production cost. Nonetheless, significant progress was achieved in the development of the bottom-up methods, and they look most suitable for realization in the ideal graphene-based tribological system described in Section 2.7.

3.8 Graphene-based composites Graphene-based composites are a broad class of materials utilizing graphene reinforcement and categorized by a composite matrix. In general, the composite matrix may be polymer, metal, or ceramic. A manufacturing process is mostly defined by the type of matrix. The fabrication of polymer-based composites is relatively simple because it does not require high temperature and pressures [68]. As a result, degradation of the reinforcement can be readily avoided. The graphene should not form aggregates and must be well dispersed to enhance the interfacial interaction with the matrix to maximize the advantage of graphene as an effective reinforcement for high-strength polymer composites. The most common method for preparing graphene-based

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Fig. 3.5 Schematic of the most common graphene production methods. Each method has been evaluated in terms of graphene quality (G), cost aspect (C; a low value corresponds to the high cost of production), scalability (S), purity (P), and yield (Y) of the overall production process. (Reprinted by permission from R. Raccichini, A. Varzi, S. Passerini, B. Scrosati, The role of graphene for electrochemical energy storage, Nat. Mater. 14 (2014) 271, https://doi.org/1038/nmat4170. Copyright (2014).) Table 3.1 Comparison of methods and properties of graphene. Method

Mechanical exfoliation (cleavage) Epitaxy CVD, intact CVD, transferred PECVD Chemical exfoliation, oxidation, reduction a

Dimension, up to

Adhesion

20 μm

Weak/moderate


200 mm
1 m
1 m
200 mm
100 mm

High High Weak Moderate/higha Weak

Depending on the material of substrate.

Durability a

High High Moderate/higha Low Moderate Low

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polymer composites is to mix both components into a solvent and evaporate it with or without vacuum conditions to form a composite. The most efficient distribution of components during the first step is achieved by bath or tip sonication [68–72]. Besides, polymer-based composites could be manufactured by in situ polymerization and melt mixing. The melt mixing technique uses a high temperature and shear forces to disperse the reinforcement phase in the polymer matrix. The high temperature liquefies the polymer phase and allows easy dispersion or intercalation of GO and reduced graphene sheets. In situ polymerization starts with the dispersion of GO or RGO in the monomer followed by its polymerization. Graphene-reinforced metal and ceramic composites are mostly produced by powder metallurgy techniques. The initial powder mixing is a critical step in ensuring the homogenous dispersion of graphene particles. Graphene should be uniformly dispersed in matrices without agglomeration to take advantage of the high surface area and nanostructure. A common problem with mixing graphene is the prevention of particle agglomeration, so the techniques of mixing ceramic/metal matrix powder and graphene powder require high energy to overcome the high surface energy of graphene that causes the agglomeration of graphene particles [73]. After preparation of a graphene-content mixture, it could be sintered using high temperatures and pressures.

3.9 Analysis and characterization 3.9.1 Optical imaging Different methods are used for the visualization of single- and few-layer graphene such as optical microscopy, scanning electron microscopy (SEM), high-resolution transmission electron microscopy (TEM), and their combinations. Initially, the optical microscope was primarily used because it is a cheap and nondestructive method. However, handling optical microscopy requires graphene layers laying on a thin dielectric layer such as SiO2 or Si3N4 [74]. High contrast between the graphene and the substrate can be obtained by choosing the appropriate optical properties and thickness of the dielectric layer. In this case, the visual contrast is based on the FabryPerot interference between the Si substrate and the dielectric surface layer. Another essential factor that defined contrast is the wavelength of the incident light. For instance, Blake et al. [75] used different narrowband filters to detect graphene sheets on SiO2 layers having different thicknesses. The authors show that graphene visibility strongly depends on both the light

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Fig. 3.6 Graphene crystallites on 300 nm SiO2 imaged with white light (A), green light, and another graphene sample on 200 nm SiO2 imaged with white light (C). Single-layer graphene is clearly visible on the left image (A), but even three layers are indiscernible on the right (C). Image sizes are 25  25 μm2. Top and bottom panels show the same flakes as in (A) and (C), respectively, but are illuminated through various narrow bandpass filters with a bandwidth of
10 nm. The flakes were chosen to contain areas of different thickness so that one can see changes in graphene’s visibility with the increasing numbers of layers. The trace in (B) shows step-like changes in contrast for one, two, and three layers (trace averaged over 10-pixel lines). This proves that the contrast can also be used as a quantitative tool for defining the number of graphene layers on a given substrate. The measured thickness of the GO sheet here is
1 nm. (Reprinted with permission from P. Blake, E.W. Hill, A.H. Castro Neto, K.S. Novoselov, D. Jiang, R. Yang, et al., Making graphene visible, Appl. Phys. Lett. 91 (2007) 063124, https://doi.org/10.1063/1.2768624. Copyright (2007) American Chemical Society.)

wavelength and the thickness of the SiO2 layer. It was found that by using monochromatic illumination, graphene can be isolated for any SiO2 thickness. Nevertheless, 300 nm and, mainly,
100 nm were observed to be most suitable for the visual detection of graphene (Fig. 3.6).

3.9.2 Fluorescence quenching It was reported in 2010 that graphene-based sheets could be made highly visible under a fluorescence microscope by quenching the emission from a dye coating [76]. As was mentioned in Section 3.4, optical imaging for graphene requires specific substrates. In contrast, the fluorescence quenching

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mechanism does not have such a limitation. In this case, graphene, reduced graphene oxide, or even graphene oxide sheets deposited on arbitrary substrates can be visualized with good contrast for layer counting. Also, this method could be used for direct observation of suspended graphene flakes in solution. The imaging mechanism involves quenching the emissions from dye-coated graphene, GO, and RGO. The contrast originated due to the chemical interaction between the dye and the graphene. The charge was transferred from the dye molecule to GO, which causes the quenching of fluorescence [76, 77]. The contrast depended on the material of the substrate and varied from 0.07 to 0.78 for glass and SiO2 correspondingly, which is enough for clear visualization [78]. Correlation between AFM and FQM images is shown in Fig. 3.7, where the FQM image is compared to the AFM image of the same area. Besides, this technique can be used for

(A)

(B)

2 µm

3 2 1 0

50 µm

0

1

2

3

4

X axis (µm)

5

6

Height (nm)

(C)

–1

Fig. 3.7 (A) Fluorescence image of a GO monolayer deposited on T4-functionalized, 300 nm thick SiOx. (B) AFM image of the area indicated in (A), showing the GO sheet partially folded over itself (z range 30 nm). (C) Height profile is taken across the black line in (B). The measured thickness of the GO sheet here is
1 nm. (Reprinted with permission from E. Treossi, M. Melucci, A. Liscio, M. Gazzano, V. Palermo, High-contrast visualization of graphene oxide on dye-sensitized glass, quartz, and silicon by fluorescence quenching, J. Am. Chem. Soc. 131 (2009) 15576, https://doi.org/10.1021/ja9055382. Copyright (2009) American Chemical Society.)

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visualization of the GO/RGO film on plastic substrates [79]. However, this technique is based on dye addition on the graphene surface. Even though the dye coating can be conveniently removed afterward by rinsing without disrupting the graphene layers, further use of the same sample might be limited due to the attachment of unwanted functional groups.

3.9.3 Atomic-force microscopy This technique is commonly used to determine the graphene layer thickness at the nanometer scale. The typical cross-sectional profile of mono- and fewlayer graphene is illustrated in Fig. 3.7. The vertical distance between the top surfaces of graphene layers is called “step high” (Fig. 3.7C). The sum of the step heights of all layers gives the total thickness of the graphene film. The step height of monolayer graphene equals its thickness. Various authors reported the thickness of the monolayer graphene in the range of 0.5–1.7 nm, depending on specimens and measurement conditions [80, 81]. Gupta et al. [82] demonstrated that the thickness of multilayer graphene linearly increased with the constant increment of 0.35 nm (Fig. 3.8). In their study, the thickness of the monolayer graphene was 0.86 nm. The 0.33 nm difference between the 0.35 increment and the thickness of the monolayer graphene was called the “AFM offset.” It was attributed to the different attraction forces between

Fig. 3.8 Effective nGL film height versus assigned n. The straight line is a least-square fit to the data. The apparent thickness of a graphene layer is t ¼ 0.35  0.01 nm, and the AFM offset parameter is t0 ¼ 0.33 nm (see text for discussion). (Reprinted with permission from A. Gupta, G. Chen, P. Joshi, S. Tadigadapa, P.C. Eklund, Raman scattering from highfrequency phonons in supported n-graphene layer films, Nano Lett. 6 (2006) 2667, https://doi.org/10.1021/nl061420a. Copyright (2006) American Chemical Society.)

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the tip and graphene compared to the tip and SiO2. Thus, the variation in the thickness of the monolayer graphene reported by various authors could be due to a slight difference of adhesion between tips, graphene, and substrates. The comprehensive comparison of measurement results of the thickness of monolayer graphene was performed by Shearer et al. [81]. It was demonstrated that using advanced AFM modes such as the PeakForce tapping mode allowed reducing the error in measuring the first layer from 0.3–1.3 nm to 0.1–0.3 nm. It was also concluded that the pressure applied to the AFM tip was the critical parameter for the accurate measurement of graphene because it mitigated the effect of the absorbate layer between a substrate and a graphene layer. Also, that observation meant that more precise results could be achieved in the dry or vacuum conditions. Despite the fact that the linear thickness increase of multilayer graphene with an increasing number of layers is intuitive, nonlinear behavior was also observed by some authors. For instance, Ptak et al. [83] found that in their experiment, the step high of exfoliated graphene transferred onto an SiO2 substrate reduced with increasing the number of layers. The thickness of monolayer graphene on SiO2 is around 0.8 nm. The step high reduced it to 0.5 nm for the second layer and reached
0.34 nm for the fifth layer. Thus, the step high of the fifth layer is equivalent to the interlayer distance of bulk graphite. The nonlinear behavior, in this case, could be attributed to the effect of water absorption from the atmosphere. Different AFM modes allowed the study of the electrical, mechanical, frictional, magnetic, and even elastic properties of graphene [16]. It is an indispensable instrument for the mechanical and tribological characteristics of graphene. For instance, AFM was used to measure the strength and Young’s modulus of graphene [79] as well as for simulation of its tensile and compressive properties [84]. Because of the extremely high lateral and vertical force resolutions of AFM, it is an excellent tool for investigating nanotribological phenomena [4]. Using an AFM device in the friction force microscopy (FFM) mode allows measuring the friction coefficient under various loads. After a friction test, wear can be immediately evaluated using a topography measurement mode. However, due to technical limitations, it is challenging to assess large areas of graphene by AFM. First, the maximum scanning area of most commercial devices is in the range from 50  50 to 100  100 μm2. Second, scanning large areas requires a compromise between image quality (resolution) and scanning time. Moreover, AFM imaging is mostly limited to topographic contrast. Thus, graphene oxide and the graphene layers cannot be distinguished in the normal direction.

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This issue could be overcome by AFM. In this case, phase imaging facilitated distinguishing between defect-free pristine graphene and the functionalized one. This method is based on the difference in the interaction forces between the AFM tip and the attached functional group. Paredes et al. demonstrated the tapping mode for determining the thickness of graphene sheets [85]. It was found that the unreduced graphene oxide has a thickness of
1.0 nm, and for chemically reduced GO it was 0.6 nm. This difference was attributed to the hydrophilicity of GO.

3.9.4 Raman spectroscopy Since the Raman spectrum of graphene was first measured in 2006 [86], Raman spectroscopy had become one of the most popular techniques for the characterization of disordered and amorphous carbons, fullerenes, nanotubes, diamonds, and carbon chains. Raman spectroscopy works for all graphene samples. Moreover, it can be used for the identification of unwanted byproducts, structural damage, functional groups, and chemical modifications introduced during the preparation, processing, or transferring of graphene. As a result, a Raman spectrum is indispensable for quality control and for comparing samples used by different research groups [86, 87]. The spectra of all carbon-based materials show only a few prominent features, regardless of the final structure. Identification of these features allows the characterization of graphene layers regarding the number of layers, strain, doping concentration, impact of temperature, and presence of defects. The Raman spectrum of graphite and multilayer graphene mainly consists of the set of peaks called D, G, and 2D. They are located around 1350, 1580, and 2700 cm1, correspondingly. A comparison of Raman spectra between pristine and defective graphene is shown in Fig. 3.9 [86]. The G band is associated with the double degenerated E2g phonon mode at the center of the Brillouin zone. This band arises due to the in-plane vibration of the sp2 carbon atoms [16]. The D band appears due to the breathing modes of six-atom rings and requires a defect for its activation. It represents disorder in the atomic arrangement or edge effect of graphene, ripples, and charge puddles [87]. The spectra of the pristine graphene do not show the D peak that confirms the absence of defects. The 2D band is the second-order peak, and it has almost double the frequency of the D band. Because it arises from a double-resonant electronic process, the 2D band is sensitive not only to the vibrational features of graphene but also to the electronic structure. As a result, Bernal (ABA) and Rhombohedral (ABC)

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Fig. 3.9 Raman spectra of pristine (top) and defective (bottom) graphene. The main peaks are labeled. (Reprinted by permission from A.C. Ferrari, D.M. Basko, A.C. Ferrari, D.M. Basko, Raman spectroscopy as a versatile tool for studying the properties of graphene, Nat. Nanotechnol. 8 (2013) 235, https://doi.org/10.1038/nnano.2013.46. Copyright (2013).)

multilayer graphene demonstrated clear differences in the shape and width of the Raman 2D peak [88]. Increasing the temperature led to the redshift of the G peak [89]. In the case of few-layer graphene and graphite, a significant change in the shape and intensity of Raman peaks was observed (Fig. 3.10). The 2D band splits into just two components for bulk graphite and multilayer graphene, and into four components for bilayer graphene. For multilayer graphene, an exact number of layers could be evaluated based on the distance between these components. It increased from
25.4 to 44 cm1 when the number of layers was raised from 3 to 10. The number of graphene layers also effects the relative intensity of the 2D peak and the position of the G peak. Increasing the number of layers from 1 to 6 increased the I(G)/I(2D) intensity ratio from
0.3 to
1 and shifted the G peak position from 1585 down to 1581 cm1 [90]. These findings allowed evaluating the spatial uniformity of the thickness of multilayer graphene by Raman mapping [22, 35]. The intensity ratio between the G band and the disorder-induced D band could be used for quantification of the defect’s density in graphene. In particular, the phenomenological model describing the relation between the ID/IG ratio and the graphene concentration was proposed for monolayer graphene by Lucchese et al. [91]. The ID/IG ratio as the function of the

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Fig. 3.10 (A) Plot of the ratio of the integrated intensities of the G and D0 peak versus the number of stacked layers (average value and standard deviation). (B) G line frequency vs. the number of stacked layers (average value and standard deviation). (C) G peak for HOPG (upper peak), double- (middle peak), and single-layer (lower peak) graphene. The vertical dashed line indicates the reference value for bulk graphite. (D) D0 peaks for an increasing number of graphene layers along with HOPG as a bulk reference. The dashed lines show the Lorentzian peaks used to fit the data; the solid lines are the fitted results. The single peak position for the single-layer graphene is at 2678.8  1.0 cm1. The peak positions of the two innermost peaks for double-layer graphene are 2683.0  1.5 and 2701.8  1.0 cm1. On the left, the value for splitting from the double-layer graphene up to HOPG is presented. All peaks are normalized in amplitude and vertically offset. (Adapted with permission from D. Graf, F. Molitor, K. Ensslin, C. Stampfer, A. Jungen, C. Hierold, L. Wirtz, Spatially resolved Raman spectroscopy of single- and few-layer graphene, Nano Lett. 7 (2007) 238, https://doi.org/10.1021/nl061702a. Copyright (2007) American Chemical Society.)

average distance between defects LD has a well-defined peak shape with the maximum around 3 nm. Reducing the LD from 25 to 6 nm led to a parabolic increase of the ID/IG ratio from 0.5 to
3.5 (right side of the peak). Then, the coalescence of defects started at LD
6 nm, the LD reduced further, and the ID/IG ratio lowered to 0.5 (right side of the peak). The ID/IG results on HOPG were different from those on monolayer graphene, showing an increase and saturation of ID/IG with decreasing LD. Thus, several-layer graphene should exhibit behavior lying between these two cases.

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The zigzag and armchair edges of graphene also could be distinguished based on the evaluation of the intensity of a disorder-induced Raman peak [92]. In particular, the armchair edge induced a sharp D peak while in the case of the zigzag edge, the peak was relatively weak; the intensity of the peak was approximately four times lower. This observation was explained based on the double-resonance theory applied to the one-dimensional defect. Because the D band is sensitive to defects, its relative intensity could be used for evaluation of the quality of the graphene films. It helps to estimate the spatial distribution of defects using Raman mapping. For instance, this approach was used by Won et al. [22] to evaluate graphene defectiveness during different stages of CVD growth on copper foils. In this work, the spatial distribution of I(D)/I(G) was plotted and compared to optical images. In combination with the evaluation of the I(G)/I(2D) ratio, this method allows obtaining all information about the uniformity of graphene films. Doping affected the Raman spectra of graphene. Increasing the charge carrier concentration led to decreasing the relative intensity and shift of the D and 2D peaks [93]. In particular, the ID/IG ratio reduced and the G peak position increased (blueshifted). Besides, the Raman spectra of graphene are sensitive to mechanical stress. Contrary to other factors, the presence of mechanical stress led to shifting the peak positions rather than a variation of their intensities. For instance, the compressive stress of several GPa led to a blueshift of the spectra in comparison to the bulk graphite [94]. Tensile stress led to the redshift of the 2D peak [95]. Under a 6% strain, the peak was shifted from 2682 down to 2676 cm1 . Raman spectroscopy can be used for evaluation of the mechanical properties of graphene. For example, compressive and tensile strain in the graphene layer can be evaluated based on a change in the G and 2D peaks with applied stress. The splitting of the G peak and the redshift was observed with an increase in strain, whereas the 2D peak also redshifted without splitting for small deformations around 0.8% [96]. Ni et al. found the opposite behavior for epitaxial graphene on an SiC substrate [94]. The blueshift of all the Raman bands of the epitaxial graphene occurred due to the compressive residual stress [16]. Raman spectroscopy allowed determining the Gruneisen parameters of suspended graphene sheets under uniaxial and biaxial strain [97]. Also, the effects of the graphene-substrate interaction on the strain and the relation between the mechanical and thermal properties were presented along with characterization of the thermal properties of graphene with Raman spectroscopy. The effect of various factors on the Raman spectra is summarized in Table 3.2.

Table 3.2 Effect of structural factors on Raman spectra of graphene. D peak

G peak

2D peak

ID/IG

IG/I2D

Ref.

Number of layers Defect concentration Doping Temperature Tensile stress Compressive stress

– Increasing – – – –

Redshift Widening/splitting Blueshift Redshift Redshift Blueshift

Widening/blueshift – Blueshift – Redshift Blueshift

– Increasing Reducing – – –

Increasing – – – – –

[90] [91] [93] [89] [95] [94]

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Factors

79

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3.9.5 X-ray photoelectron spectroscopy X-ray photoelectron spectroscopy (XPS), also known as ESCA (electron spectroscopy for chemical analysis), is an ultimate tool for the examination of a surface. It is a quantitative spectroscopic technique that allows the evaluation of elemental composition and the chemical state of elements that exist within a material. First, it is an indispensable method for assessing the oxidation of graphene and its derivatives. In this case, primary attention should be paid to the carbon and oxygen peaks. The most representative carbon peak is C 1s, having a binding energy around 284 eV. The O 1s peak has a binding energy around 532 eV. The presence of chemical bonds between carbon, oxygen, and hydrogen leads to the appearance of satellite peaks. For example, Fig. 3.11 illustrates the variation of carbon and oxygen peaks during the high-temperature reduction of GO in the Ar atmosphere and the comparison to the graphite standard [98]. A change of the oxidation rate leads to a redistribution of intensities of these satellite peaks. Besides, XPS can be used to evaluate doping (functionalization) of graphene [99]. For instance, fluorine induces a significant chemical shift of the C 1s binding energy, allowing the quantification of composition and bonding type [100]. Nitrogen doping of graphene caused a shift of the N 1 s peak. Moreover, two predominant binding conditions of nitrogen in graphene could be distinguished [101]. XPS was used for monitoring the defect formation in graphene film during evaluation of the degradation mechanism under dry sliding [22]. In this case, the creation of the defects and voids in the continuous graphene film during the sliding test led to the oxidation of a substrate, causing the appearance of Cu-O peaks on the XPS spectra.

3.9.6 Transmission electron microscopy Transmission electron microscopy (TEM) is used to visualize nano-sized materials with atomic-scale resolution. Because graphene and GO/RGO are layers of several atoms thick, TEM should be one of the most suitable tools to allow resolving the atomic features of the graphene [16]. However, the use of traditional TEM is limited because the operation at high voltage damages the monolayer, whereas at the low operating voltage, the resolution is insufficient. The development of a new class of TEM devices with aberration correction in combination with a monochromator allowed overcom˚ ing this issue. It was demonstrated that such devices could provide 1 A resolution at an acceleration voltage of only 80 kV [102, 103]. For the first

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Fig. 3.11 Results of C 1 s and O 1 s spectra fitting into C and O chemical groups for XPS spectra recorded for (A) FL-GO, (B) FL-RGO, and (C) graphite standards. (Adapted from L. Stobinski, B. Lesiak, A. Malolepszy, M. Mazurkiewicz, B. Mierzwa, J. Zemek, et al., Graphene oxide and reduced graphene oxide studied by the XRD, TEM, and electron spectroscopy methods, J. Electron Spectrosc. 195 (2014) 145, https://doi.org/10.1016/j.elspec.2014.07. 003. Copyright (2014), with permission from Elsevier.)

time, the direct high-resolution images of the graphene lattice were demonstrated by Mayer et al. [102]. It becomes possible to observe every single carbon atom arranged in a hexagonal fashion. It was shown that the imperfection and topological peculiarities in graphene affected the electronic and mechanical properties, and this can be determined using an aberrationcorrected low-voltage TEM.

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Besides, the atomic lattice is visible in high-angle annular dark-field (HAADF) images acquired in a scanning transmission electron microscope (STEM) [103]. In this case, the electron beam focused onto a monoatomic layer of atoms, such as graphene. Bright contrast corresponds to atoms and dark contrast to the gaps between them. Such images are directly interpretable because they are a direct depiction of the ball-and-stick model of an atomic lattice structure. TEM was successfully used for the direct observation of structural defects in the monolayer graphene membrane [104]. The defects were generated under 300 keV electron-beam irradiation. Various defects such as monovacancy, divacancy, and Stone-Wales defects were observed (Fig. 3.12). It was shown how the transformation occurred step by step by nucleation and growth of low-energy multivacancy structures constructed of rotated hexagons and other polygons. Besides the point defects, a grain boundary is another type of defect that is especially pronounced in two-dimensional materials. Using TEM in combination with the scanning probe not only allowed visualizing these defects, but also demonstrated that the grain boundaries severely weaken the mechanical strength of graphene membranes [105]. It was shown that two crystals were stitched together by a series of pentagons, heptagons, and distorted hexagons. The grain boundary was not straight, and the defects along the boundary were not periodic. The last observation was quite crucial because the aperiodicity was contrasted with many theoretical models [105].

Fig. 3.12 Elementary defects and frequently observed defect transformations under irradiation. Atomic bonds are superimposed on the defected areas in the bottom row. Creation of the defects can be explained by atom ejection and reorganization of bonds via bond rotation. (A) Stone-Wales defect, (B) defect-free graphene, (C) V1(5-9) single vacancy, (D) V2(5-8-5) divacancy, (E) V2(555-777) divacancy, and (F) V2(5555-6-7777) divacancy. Scale bar is 1 nm. (Reprinted with permission from J. Kotakoski, A.V. Krasheninnikov, U. Kaiser, J.C. Meyer, From point defects in graphene to two-dimensional amorphous carbon, Phys. Rev. Lett. 106 (2011) 105505. Copyright (2011) by the American Physical Society.)

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3.10 Summary Since the first exfoliation of graphene flakes by the Scotch tape method, huge progress has been achieved in the development of advanced synthesis methods. Some of these methods are almost “perfect” regarding quality, and they are mostly used for research purposes due to their extreme cost. Other methods are not able to produce such “perfect” graphene, but they are cheap enough to be used in the industry. Thus, there is always a trade off between cost and scalability on the one hand and the quality of graphene on the other. Besides the general requirements to the “price/quality” ratio, tribological applications rise additional demands, such as substrate adhesion and area coverage. From the tribological perspective, the CVD-based methods seem to be the most promising, especially if the transfer to another substrate is not required. Besides a variety of synthesis methods, analytical techniques were developed. Raman spectroscopy is of the most interest because it is an ultimate informative and nondestructive method for graphene analysis. With a combination of in situ techniques, it could be an indispensable method for tribological studies of graphene. Other methods such as XPS and TEM could provide additional information about the chemical state and structure of graphene. AFM is widely used for the evaluation of mechanical and nanotribological properties of graphene, as was discussed in this chapter.

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Further reading [106] K.S. Novoselov, V.I. Fal’ko, L. Colombo, P.R. Gellert, M.G. Schwab, K. Kim, A roadmap for graphene, Nature 490 (2012) 192, https://doi.org/10.1038/ nature11458.