**Surface Functionalization of Graphene with Polymers for Enhanced Properties**

Wenge Zheng, Bin Shen and Wentao Zhai

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/50490

## **1. Introduction**

[68] Zhuang, X. D., Chen, Y., Liu, G., Li, P. P., Zhu, C. X., Kang, E. T., Neoh, K. G., Zhang, B., Zhu, J. H., & Li, Y. X. (2010). Conjugated-polymer-functionalized graphene oxide: Synthesis and nonvolatile rewritable memory effect. *Advanced Materials*, 22(15),

[69] Zhuge, F., Dai, W., He, C. L., Wang, A. Y., Liu, Y. W., Li, M., Wu, Y. H., Cui, P., & Li, R. W. (2010). Nonvolatile resistive switching memory based on amorphous carbon.

[70] Zhuge, F., Hu, B. L., He, C. L., Zhou, X. F., Liu, Z. P., & Li, R. W. (2011). Mechanism of nonvolatile resistive switching in graphene oxide thin films. *Carbon*, 49(12),

[71] Zhuge, F., Li, R. W., He, C. L., Liu, Z. P., & Zhou, X. F. (2011, Mar.) Non-volatile resis‐ tive switching in graphene oxide thin films. *Physics and Applications of Graphene-Ex‐*

1731-1735, 1476-1122.

206 New Progress on Graphene Research

3796-3802, 0008-6223.

*Applied Physics Letters*, 96(16), 163505, 0003-6951.

*periments*, 421-438, Print, 978-953-307-217-3, Croatia, InTech.

Graphene, a single-atom-thick sheet of hexagonally arrayed sp2 bonded carbon atoms, has been under the spotlight owning to its intriguing and unparalleled physical properties [1]. Because of its novel properties, such as exceptional thermal conductivity, [2] high Young's modulus, [3] and high electrical conductivity,[4] graphene has been highlighted in fabricat‐ ing various micro-electrical devices, batteries, supercapacitors, and composites [5, 7]. Espe‐ cially, integration of graphene and its derivations into polymer has been highlighted, from the point views of both the spectacular improvement in mechanical, electrical properties, and the low cost of graphite [8, 9]. Control of the size, shape and surface chemistry of the reinforcement materials is essential in the development of materials that can be used to pro‐ duce devices, sensors and actuators based on the modulation of functional properties. The maximum improvements in final properties can be achieved when graphene is homogene‐ ously dispersed in the matrix and the external load is efficiently transferred through strong filler/polymer interfacial interactions, extensively reported in the case of other nanofillers. However, the large surface area of graphene and strong van der Waals force among them result in severe aggregation in the composites matrix. Furthermore, the carbon atoms on the graphene are chemically stable because of the aromatic nature of the bond. As a result, the reinforcing graphene are inert and can interact with the surrounding matrix mainly through van der Waals interactions, unable to provide an efficient load transfer across the graphene/ matrix interface. To obtain satisfied performance of the final graphene/polymer composites, the issues of the strong interfacial adhesion between graphene–matrix and well dispersion of graphene should be addressed.

To date, the mixing of graphene and functionalized graphene with polymers covers the most of the published studies, and the direct modification of graphene with polymers is a

© 2013 Zheng et al.; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2013 Zheng et al.; licensee InTech. This is a paper distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

somewhat less explored approach. However, in many cases, to achieve stable dispersions of graphene and adequate control of the microstructure of the nanocomposites, non-covalent or covalent functionalization of graphene with polymers may be necessary. The non-cova‐ lent functionalization, which relies on the van der Waals force, electrostatic interaction or ππ stacking [10, 12], is easier to carry out without altering the chemical structure of the graphene sheets, and provides effective means to tailor the electronic/optical property and solubility of the nanosheets [13]. The covalent functionalization of graphene derivatives is mainly based on the reaction between the functional groups of the molecules and the oxy‐ genated groups on graphene oxide (GO) or reduced GO (r-GO) surfaces [14, 15], such as ep‐ oxides and hydroxyls on their basal planes and carboxyls on the edges [16]. Compared with non-covalent functionalization, the covalent functionalization of graphene-based sheets holds versatile possibility due to the rich surface chemistry of GO/r-GO. However, it should be pointed out that the non-covalent or covalent attachment of graphene to polymer chains can improve some properties, but may be negative for others, especially those related to the movement of electrons or phonons. Although the functionalization of graphene with poly‐ mers is generally attempted with a view to conferring to the polymer new or improved properties, the polymer may also prevent the aggregation of the graphene sheets, where the graphene-polymer size ratio and molecular weight play important roles. For general bibliog‐ raphy on typical graphene-based nanocomposites, the reader can consult several mono‐ graphs, reviews, and feature articles that summarize the state of the art of the field.

(water, alcohol, and other protic solvents) assisted by mechanical exfoliation, such as ultra‐ sonication and/or stirring, forming colloidal suspensions of "graphene oxide". The chemical‐ ly reduced graphene oxide is produced by chemical reduction of the exfoliated graphene oxide sheets using hydrazine, [22, 24] dimethylhydrazine, [25] sodium borohydride fol‐ lowed by hydrazine, [26] hydroquinone, [ 27] vitamin C, [28] etc. However, the hazardous nature and cost of the chemicals used in reduction may limit its application. The most prom‐ ising methods for large scale production of graphene is the thermal exfoliation and reduc‐ tion of GO. Thermally reduced graphene oxide can be produced by rapid heating of dry GO under inert gas and high temperature [29, 31]. Heating GO in an inert environment at 1050°C for 30 s leads to reduction and exfoliation of GO, producing low-bulk-density TRG sheets, which are highly wrinkled [32]. In the thermal process, the epoxy and hydroxyl sites of GO decompose to produce gases like H2O and CO2, yielding pressures that exceed van der Waals forces holding the graphene sheets together, causing the occurrence of exfoliation.

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It is desirable that stronger bonds are usually formed between the graphene and the poly‐ mers by covalent functionalization of graphene with polymers. However, it is usually diffi‐ cult to realize because ideal graphene lacks functional groups that can be conjugated with. In some cases, when the graphene sheets were exfiolated from GO, incomplete reduction process leaves oxygen-containing functionalities that are then available for further function‐ alizations. Other covalent functionalization strategies typically involve further disruption of the conjugation of the graphene sheets. Although covalent functionalization of graphene will compromise some of its natural conductivity, this method is still valuable in some cases when graphene's other properties are desirable. More details of graphene functionalization

Until now, "grafting to" and "grafting from" techniques have been developed to graft the polymer chains onto the graphene surface. The "grafting from" method relies on the immo‐ bilization of initiators at the surface of graphene, followed by in situ surface-initiated poly‐ merization to generate tethered polymer chains. A number of studies of polymer-

Among the types of "grafting from" polymerization [33, 46], ATRP is the most widely used, and represents the majority of the studies reported. ATRP is almost certainly chosen because it offers the advantages of radical polymerization, that is, a fast initiation process and the development of a dynamic equilibrium between dormant and growing radicals [47]. In ad‐ dition, a wide range of monomers can be polymerized by ATRP with controlled chain length. Moreover, block copolymers can be prepared by ATRP because of the living radical

functionalized graphene by the "grafting from" method have been reported.

**3. Covalent functionalization of graphene with polymers**

via covalent bonds will be discussed below.

**3.1. Functionalizations via "grafting from" method**

*3.1.1. Atom transfer radical polymerization (ATRP)*

The objective of the present work is to provide a broad overview on the methods developed to non-covalently or covalently bind graphene to polymers. The covalent linking of poly‐ meric chains to graphene is at its initial stages and there is significant room for the develop‐ ment of new and improved strategies.

#### **2. The precursor of functionalized graphene**

As we know, GO is the main precursor for the functionalization of graphene with poly‐ mers. It is because that there are multiple oxygen-containing functionalities, such as hydrox‐ yl, epoxy and carboxyl groups on GO sheets [16]. GO is usually produced using different variations of the Staudenmaier [17] or Hummers [18] method in which graphite is oxi‐ dized using strong oxidants such as KMnO4, KClO3, and NaNO2 in the presence of nitric acid or its mixture with sulfuric [19, 20]. For more details about GO, we refer the reader to the extensive review of GO preparation, structure, and reactivity by Dreyer et al and Zhu et al. [19, 20].

Furthermore, the reduction of GO will remove most, but not all, of the oxygen-containing functionalities such as hydroxyl, carboxylic acid and epoxy groups. Therefore, some func‐ tionalization reactions are based on the reduced GO. Generally, GO can be exfoliated using a variety of methods, most commonly by solvent-based exfoliation and reduction in appro‐ priate media or thermal exfoliation and reduction [16, 21] In the former route, the hydrophil‐ ic nature and increased interlayer spacing of GO facilitates direct exfoliation into solvents (water, alcohol, and other protic solvents) assisted by mechanical exfoliation, such as ultra‐ sonication and/or stirring, forming colloidal suspensions of "graphene oxide". The chemical‐ ly reduced graphene oxide is produced by chemical reduction of the exfoliated graphene oxide sheets using hydrazine, [22, 24] dimethylhydrazine, [25] sodium borohydride fol‐ lowed by hydrazine, [26] hydroquinone, [ 27] vitamin C, [28] etc. However, the hazardous nature and cost of the chemicals used in reduction may limit its application. The most prom‐ ising methods for large scale production of graphene is the thermal exfoliation and reduc‐ tion of GO. Thermally reduced graphene oxide can be produced by rapid heating of dry GO under inert gas and high temperature [29, 31]. Heating GO in an inert environment at 1050°C for 30 s leads to reduction and exfoliation of GO, producing low-bulk-density TRG sheets, which are highly wrinkled [32]. In the thermal process, the epoxy and hydroxyl sites of GO decompose to produce gases like H2O and CO2, yielding pressures that exceed van der Waals forces holding the graphene sheets together, causing the occurrence of exfoliation.

## **3. Covalent functionalization of graphene with polymers**

It is desirable that stronger bonds are usually formed between the graphene and the poly‐ mers by covalent functionalization of graphene with polymers. However, it is usually diffi‐ cult to realize because ideal graphene lacks functional groups that can be conjugated with. In some cases, when the graphene sheets were exfiolated from GO, incomplete reduction process leaves oxygen-containing functionalities that are then available for further function‐ alizations. Other covalent functionalization strategies typically involve further disruption of the conjugation of the graphene sheets. Although covalent functionalization of graphene will compromise some of its natural conductivity, this method is still valuable in some cases when graphene's other properties are desirable. More details of graphene functionalization via covalent bonds will be discussed below.

#### **3.1. Functionalizations via "grafting from" method**

somewhat less explored approach. However, in many cases, to achieve stable dispersions of graphene and adequate control of the microstructure of the nanocomposites, non-covalent or covalent functionalization of graphene with polymers may be necessary. The non-cova‐ lent functionalization, which relies on the van der Waals force, electrostatic interaction or ππ stacking [10, 12], is easier to carry out without altering the chemical structure of the graphene sheets, and provides effective means to tailor the electronic/optical property and solubility of the nanosheets [13]. The covalent functionalization of graphene derivatives is mainly based on the reaction between the functional groups of the molecules and the oxy‐ genated groups on graphene oxide (GO) or reduced GO (r-GO) surfaces [14, 15], such as ep‐ oxides and hydroxyls on their basal planes and carboxyls on the edges [16]. Compared with non-covalent functionalization, the covalent functionalization of graphene-based sheets holds versatile possibility due to the rich surface chemistry of GO/r-GO. However, it should be pointed out that the non-covalent or covalent attachment of graphene to polymer chains can improve some properties, but may be negative for others, especially those related to the movement of electrons or phonons. Although the functionalization of graphene with poly‐ mers is generally attempted with a view to conferring to the polymer new or improved properties, the polymer may also prevent the aggregation of the graphene sheets, where the graphene-polymer size ratio and molecular weight play important roles. For general bibliog‐ raphy on typical graphene-based nanocomposites, the reader can consult several mono‐

graphs, reviews, and feature articles that summarize the state of the art of the field.

ment of new and improved strategies.

208 New Progress on Graphene Research

et al. [19, 20].

**2. The precursor of functionalized graphene**

The objective of the present work is to provide a broad overview on the methods developed to non-covalently or covalently bind graphene to polymers. The covalent linking of poly‐ meric chains to graphene is at its initial stages and there is significant room for the develop‐

As we know, GO is the main precursor for the functionalization of graphene with poly‐ mers. It is because that there are multiple oxygen-containing functionalities, such as hydrox‐ yl, epoxy and carboxyl groups on GO sheets [16]. GO is usually produced using different variations of the Staudenmaier [17] or Hummers [18] method in which graphite is oxi‐ dized using strong oxidants such as KMnO4, KClO3, and NaNO2 in the presence of nitric acid or its mixture with sulfuric [19, 20]. For more details about GO, we refer the reader to the extensive review of GO preparation, structure, and reactivity by Dreyer et al and Zhu

Furthermore, the reduction of GO will remove most, but not all, of the oxygen-containing functionalities such as hydroxyl, carboxylic acid and epoxy groups. Therefore, some func‐ tionalization reactions are based on the reduced GO. Generally, GO can be exfoliated using a variety of methods, most commonly by solvent-based exfoliation and reduction in appro‐ priate media or thermal exfoliation and reduction [16, 21] In the former route, the hydrophil‐ ic nature and increased interlayer spacing of GO facilitates direct exfoliation into solvents Until now, "grafting to" and "grafting from" techniques have been developed to graft the polymer chains onto the graphene surface. The "grafting from" method relies on the immo‐ bilization of initiators at the surface of graphene, followed by in situ surface-initiated poly‐ merization to generate tethered polymer chains. A number of studies of polymerfunctionalized graphene by the "grafting from" method have been reported.

#### *3.1.1. Atom transfer radical polymerization (ATRP)*

Among the types of "grafting from" polymerization [33, 46], ATRP is the most widely used, and represents the majority of the studies reported. ATRP is almost certainly chosen because it offers the advantages of radical polymerization, that is, a fast initiation process and the development of a dynamic equilibrium between dormant and growing radicals [47]. In ad‐ dition, a wide range of monomers can be polymerized by ATRP with controlled chain length. Moreover, block copolymers can be prepared by ATRP because of the living radical process. Furthermore, ATRP is probably the most practical technique for preparation of functional polymers because the terminal alkyl halide can be converted to a wide variety of functionalities by using conventional organic synthetic procedures.

cal mechanism as that for carbon nanotubes (CNTs). The relaxation of the polymer chains covalently bonded to the r-GO surface was strongly confined, particularly for segments in close proximity to the r-GO surface. This confinement effect could enhance the thermal conductivity of the polymer nanocomposites. The significant increases in thermal conduc‐ tivity were observed for only 2.0 wt% functionalized r-GO in PS composites. Also, the resulting PS nanocomposites with 0.9 wt% functionalized r-GO revealed around 70% and

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211

57% increases in tensile strength and Young's modulus.

**Figure 2.** Synthetic routes for achieving controllable functionalization of graphene. 28

Furthermore, Gonçalves et al. [40] developed the use of poly(methyl methacrylate) (PMMA) grafted from carboxylic groups in GO as a reinforcement filler. Here the BMPB initiators were immobilized by two esterification reactions: the carboxylic groups of GO were esteri‐ fied with ethylene glycol, followed by reacting with BMPB using the same type of reaction as Lee et al. [38]. In this case, the polydispersity of the grafted PMMA, removed from the GO by hydrolysis, was found to be close to unity, once more suggesting a well-controlled process irrespective of the under estimated molecular weights. This PMMA-functionalized GO showed a good solubility in organic solvents such as chloroform and could be used as reinforcement filler in the preparation of PMMA composite films. Due to the strong interfa‐ cial interactions between the PMMA-functionalized GO and PMMA matrix caused by the presence of short PMMA chains covalently bonded to GO, an efficient load transfer from the GO to the matrix was formed, thus improving the mechanical properties of their nanocom‐ posites, which were more stable and tougher than pure PMMA and its nanocomposites with unmodified GO (Figure 3). For example, addition of 1 wt% PMMA-functionalized GO clear‐ ly led to a significant improvement of the elongation at break, yielding a much more ductile

**Figure 1.** Synthesis of surface-functionalized GO via attachment of an ATRP initiator (a-bromoisobutyryl bromide) fol‐ lowed by polymerization of styrene, butyl acrylate, or methyl methacrylate [38].

Lee et al. [38] have reported a new method for attaching polymer brushes to GO sheet using surface-initiated ATRP. The hydroxyl groups present on the surface of GO were first func‐ tionalized with a wellknown ATRP initiator (a-bromoisobutyryl bromide), and then poly‐ mers of styrene, butyl acrylate, or methyl methacrylate were grown directly via a surfaceinitiated polymerization (SIP) (Figure 1). The authors studied the case of polystyrene (PS) in detail and presented two main conclusions. First, they suggested that the polymer chain length can be tunable by changing the ratio of monomer and initiator modified GO. Sec‐ ond, they reported that the monomer loading can vary the molecular weight of the graft‐ ed PS, which was obtained by gel permeation chromatography (GPC) after detaching by saponification, and the polydispersity was low, which suggested the polymerization pro‐ ceeds in a controlled manner. Furthermore, the PS-functionalized GO was shown to sig‐ nificantly increase the solubility in N,N-dimethylformamide (DMF), toluene, chloroform, and dichloromethane, improving the processing potential of these materials for applica‐ tions in polymer composites.

Fang et al. [43, 44] demonstrated the ability to systematically tune the grafting density and chain length of PS covalently bonded to graphene sheets by combining diazonium addi‐ tion and ATRP. After reduction, r-GO was functionalized with 2-(4-aminophenyl) etha‐ nol, reacted with 2-bromo-2-methylpropi-onyl bromide (BMPB), and subsequently the polymerization of styrene was carried out (Figure 2). Their results showed that the poly‐ dispersity of the high grafting density sample was more uniform than that of the low grafting density, which was attributed to the the degree of functionalization of r-GO sheets with the initiator because the diazonium coupling to graphene follows an identical radi‐ cal mechanism as that for carbon nanotubes (CNTs). The relaxation of the polymer chains covalently bonded to the r-GO surface was strongly confined, particularly for segments in close proximity to the r-GO surface. This confinement effect could enhance the thermal conductivity of the polymer nanocomposites. The significant increases in thermal conduc‐ tivity were observed for only 2.0 wt% functionalized r-GO in PS composites. Also, the resulting PS nanocomposites with 0.9 wt% functionalized r-GO revealed around 70% and 57% increases in tensile strength and Young's modulus.

process. Furthermore, ATRP is probably the most practical technique for preparation of functional polymers because the terminal alkyl halide can be converted to a wide variety of

**Figure 1.** Synthesis of surface-functionalized GO via attachment of an ATRP initiator (a-bromoisobutyryl bromide) fol‐

Lee et al. [38] have reported a new method for attaching polymer brushes to GO sheet using surface-initiated ATRP. The hydroxyl groups present on the surface of GO were first func‐ tionalized with a wellknown ATRP initiator (a-bromoisobutyryl bromide), and then poly‐ mers of styrene, butyl acrylate, or methyl methacrylate were grown directly via a surfaceinitiated polymerization (SIP) (Figure 1). The authors studied the case of polystyrene (PS) in detail and presented two main conclusions. First, they suggested that the polymer chain length can be tunable by changing the ratio of monomer and initiator modified GO. Sec‐ ond, they reported that the monomer loading can vary the molecular weight of the graft‐ ed PS, which was obtained by gel permeation chromatography (GPC) after detaching by saponification, and the polydispersity was low, which suggested the polymerization pro‐ ceeds in a controlled manner. Furthermore, the PS-functionalized GO was shown to sig‐ nificantly increase the solubility in N,N-dimethylformamide (DMF), toluene, chloroform, and dichloromethane, improving the processing potential of these materials for applica‐

Fang et al. [43, 44] demonstrated the ability to systematically tune the grafting density and chain length of PS covalently bonded to graphene sheets by combining diazonium addi‐ tion and ATRP. After reduction, r-GO was functionalized with 2-(4-aminophenyl) etha‐ nol, reacted with 2-bromo-2-methylpropi-onyl bromide (BMPB), and subsequently the polymerization of styrene was carried out (Figure 2). Their results showed that the poly‐ dispersity of the high grafting density sample was more uniform than that of the low grafting density, which was attributed to the the degree of functionalization of r-GO sheets with the initiator because the diazonium coupling to graphene follows an identical radi‐

functionalities by using conventional organic synthetic procedures.

210 New Progress on Graphene Research

lowed by polymerization of styrene, butyl acrylate, or methyl methacrylate [38].

tions in polymer composites.

**Figure 2.** Synthetic routes for achieving controllable functionalization of graphene. 28

Furthermore, Gonçalves et al. [40] developed the use of poly(methyl methacrylate) (PMMA) grafted from carboxylic groups in GO as a reinforcement filler. Here the BMPB initiators were immobilized by two esterification reactions: the carboxylic groups of GO were esteri‐ fied with ethylene glycol, followed by reacting with BMPB using the same type of reaction as Lee et al. [38]. In this case, the polydispersity of the grafted PMMA, removed from the GO by hydrolysis, was found to be close to unity, once more suggesting a well-controlled process irrespective of the under estimated molecular weights. This PMMA-functionalized GO showed a good solubility in organic solvents such as chloroform and could be used as reinforcement filler in the preparation of PMMA composite films. Due to the strong interfa‐ cial interactions between the PMMA-functionalized GO and PMMA matrix caused by the presence of short PMMA chains covalently bonded to GO, an efficient load transfer from the GO to the matrix was formed, thus improving the mechanical properties of their nanocom‐ posites, which were more stable and tougher than pure PMMA and its nanocomposites with unmodified GO (Figure 3). For example, addition of 1 wt% PMMA-functionalized GO clear‐ ly led to a significant improvement of the elongation at break, yielding a much more ductile and tougher material. In addition, the presence of PMMA-functionalized GO also stabilized the nanocomposites increasing the onset of thermal decomposition by around 50 °C.

**Figure 3.** Load-displacement nanoidentation (left) and stress-strain curves (right) of films of PMMA and its nanocom‐ posites with PMMA-functionalized graphene. 40

Yang et al. [42] also took advantage of the carboxylic groups to graft poly-(2-dimethylami‐ noethyl methacrylate) (PDMAEMA) onto GO sheets. Here, the BMPB initiators were attach‐ ed onto GO sheets by two steps involving amidation reactions. This functionalization of GO with PDMAEMA not only enhanced the solubility in acidic aqueous solutions (pH = 1), but also in short chain alcohols. Moreover, this solubility allowed this functionalized-GO to be mixed with spherical particles of poly(ethylene glycol dimethacrylate-co-methacrylic acid) to generate decorated GO sheets.

**Figure 4.** The improved thermal and mechanical propertied of PU/GO nanocomposites [33].

the nanocomposites.

Etmimi et al. [35] investigated the preparation of PS/GO nanocomposites via RAFT mediat‐ ed mini-emulsion polymerization. In this process, dodecyl isobutyric acid trithiocarbonate (DIBTC) RAFT agent was attached to the hydroxyl groups of GO through an esterification reaction (Figure 5). The resultant RAFT-grafted GO was used for the preparation of PS/GO nanocomposites in miniemulsion polymerization. The stable miniemulsions were obtained by sonicating RAFT-grafted GO in styrene monomer in the presence of a surfactant, fol‐ lowed by polymerizing using AIBN as the initiator to yield encapsulated PS-GO nanocom‐ posites. The molecule weight and polydispersity of PS in the nanocomposites depended on the amount of RAFT-grafted GO in the system, in accordance with the features of the RAFT polymerization method. The thermal stability of the obtained PS/GO nanocomposites was improved, which may be attributed to the intercalation of PS into the lamellae of graphite. Furthermore, the increased RAFT-grafted GO significantly resulted in the improved mechan‐ ical properties of the nanocomposites. The storage and loss modulus of the nanocompo‐ sites were higher than those of the standard PS when the GO loadings reached 3 and 6%, respectively. Oppositely, as RAFT-grafted GO content increased, the Tg values of the sam‐ ple decreased. This was attributed to the change in the molecule weight of the PS chains in

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#### *3.1.2. Other polymerization methods apart from ATRP*

Besides the ATRP method, polycondensation, [33] ring opening polymerization [34], rever‐ sible addition-fragmentation chain transfer (RAFT) mediated mini-emulsion polymeriza‐ tion [35], direct electrophilic substitution [36], and Ziegler–Natta polymerization [37] have also been used to functionalize the graphene sheets with various polymer chains, which will be discussed below.

Wang et al. [33] functionalized GO sheets with polyurethane (PU) by using the polyconden‐ sation method. Here GO sheets were reacted with 4,4'-diphenylmethane diisocyanate fol‐ lowed by polycondensation of poly(tetramethylene glycol) and ethylene glycol. The presence of PU chains linked to GO remarkably improved the dispersion of GO in PU ma‐ trix, which was confirmed by the morphological study, and make it compatible with pure PU forming strong interfacial interactions that provide an enhanced load transfer between the matrix and the GO sheets thus improving their mechanical properties, as well as their thermal properties. With the incorporation of 2.0 wt% PU-functionalized GO, the tensile strength and storage modulus of the PU nanocomposites increased by 239% and 202%, re‐ spectively (Figure 4). Furthermore, the nanocomposites displayed high electrical conductivi‐ ty, and improved thermal stability of PU was also achieved (Figure 4).

Surface Functionalization of Graphene with Polymers for Enhanced Properties http://dx.doi.org/10.5772/50490 213

**Figure 4.** The improved thermal and mechanical propertied of PU/GO nanocomposites [33].

and tougher material. In addition, the presence of PMMA-functionalized GO also stabilized

**Figure 3.** Load-displacement nanoidentation (left) and stress-strain curves (right) of films of PMMA and its nanocom‐

Yang et al. [42] also took advantage of the carboxylic groups to graft poly-(2-dimethylami‐ noethyl methacrylate) (PDMAEMA) onto GO sheets. Here, the BMPB initiators were attach‐ ed onto GO sheets by two steps involving amidation reactions. This functionalization of GO with PDMAEMA not only enhanced the solubility in acidic aqueous solutions (pH = 1), but also in short chain alcohols. Moreover, this solubility allowed this functionalized-GO to be mixed with spherical particles of poly(ethylene glycol dimethacrylate-co-methacrylic acid)

Besides the ATRP method, polycondensation, [33] ring opening polymerization [34], rever‐ sible addition-fragmentation chain transfer (RAFT) mediated mini-emulsion polymeriza‐ tion [35], direct electrophilic substitution [36], and Ziegler–Natta polymerization [37] have also been used to functionalize the graphene sheets with various polymer chains, which

Wang et al. [33] functionalized GO sheets with polyurethane (PU) by using the polyconden‐ sation method. Here GO sheets were reacted with 4,4'-diphenylmethane diisocyanate fol‐ lowed by polycondensation of poly(tetramethylene glycol) and ethylene glycol. The presence of PU chains linked to GO remarkably improved the dispersion of GO in PU ma‐ trix, which was confirmed by the morphological study, and make it compatible with pure PU forming strong interfacial interactions that provide an enhanced load transfer between the matrix and the GO sheets thus improving their mechanical properties, as well as their thermal properties. With the incorporation of 2.0 wt% PU-functionalized GO, the tensile strength and storage modulus of the PU nanocomposites increased by 239% and 202%, re‐ spectively (Figure 4). Furthermore, the nanocomposites displayed high electrical conductivi‐

ty, and improved thermal stability of PU was also achieved (Figure 4).

posites with PMMA-functionalized graphene. 40

212 New Progress on Graphene Research

to generate decorated GO sheets.

will be discussed below.

*3.1.2. Other polymerization methods apart from ATRP*

the nanocomposites increasing the onset of thermal decomposition by around 50 °C.

Etmimi et al. [35] investigated the preparation of PS/GO nanocomposites via RAFT mediat‐ ed mini-emulsion polymerization. In this process, dodecyl isobutyric acid trithiocarbonate (DIBTC) RAFT agent was attached to the hydroxyl groups of GO through an esterification reaction (Figure 5). The resultant RAFT-grafted GO was used for the preparation of PS/GO nanocomposites in miniemulsion polymerization. The stable miniemulsions were obtained by sonicating RAFT-grafted GO in styrene monomer in the presence of a surfactant, fol‐ lowed by polymerizing using AIBN as the initiator to yield encapsulated PS-GO nanocom‐ posites. The molecule weight and polydispersity of PS in the nanocomposites depended on the amount of RAFT-grafted GO in the system, in accordance with the features of the RAFT polymerization method. The thermal stability of the obtained PS/GO nanocomposites was improved, which may be attributed to the intercalation of PS into the lamellae of graphite. Furthermore, the increased RAFT-grafted GO significantly resulted in the improved mechan‐ ical properties of the nanocomposites. The storage and loss modulus of the nanocompo‐ sites were higher than those of the standard PS when the GO loadings reached 3 and 6%, respectively. Oppositely, as RAFT-grafted GO content increased, the Tg values of the sam‐ ple decreased. This was attributed to the change in the molecule weight of the PS chains in the nanocomposites.

wt %, σc was measured at 0.3 S m−1. We believe that this must originate from a side reaction

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Ultrasound has found important applications in a diverse range of materials and chemi‐ cal syntheses [48, 51]. Both the physical and chemical effects of ultrasound arise from acous‐ tic cavitation: the formation, growth, and collapse of bubbles in liquids irradiated with high intensity ultrasound [48 - 50, 52]. Localized hot spots with temperature of ~5000K and pressures of hundreds of bars are generated during the bubble collapse within liquid, which can induce some chemical reactions that can't take place under normal conditions. Xu et al. [53] reported a convenient single-step sonication-induced approach for the preparation of polymer functionalized graphenes from graphite flakes and a reactive monomer, styrene. In this work, they showed that by choosing a reactive medium as the solvent, the com‐ bined mechanochemical effects of high intensity ultrasound can, in a single step, readily induce exfoliation of graphite to produce functionalized graphenes. Ultrasonic irradiation of graphite in styrene results in the mechanochemical exfoliation of graphite flakes to singlelayer and few-layer graphene sheets combined with functionalization of the graphene with PS chains (Figure 7). The PS chains are formed from sonochemically initiated radical poly‐ merization of styrene. They also tested a variety of other solvents, including toluene, ethyl‐ benzene, 1-dodecene, and 4-vinylpyridine to prepare functionalized graphenes. Only the easily polymerizable reactants containing vinyl groups, styrene and 4-vinylpyridine, lead to stable functionalized graphene. Such functionalized graphene have good stability and solubility in common organic solvents and have great potential for graphene-based compo‐

Moreover, direct photografting reactions of vinyl monomers came into focus for the prepa‐ ration of stable polymer brushes. Two approaches in particular, the sequential "living" pho‐ topolymerization and the self-initiated photografting and photopolymerization (SIPGP), attracted the attention of numerous research groups because of the facile preparation and broad applicability. Steenackers et al. [54] showed that PS chains could covalently bound to graphene by the UV-induced polymerization of styrene (Figure 8). Photopolymerization oc‐ curs at existing defect sites and that there is no detectable disruption of the basal plane con‐ jugation of graphene. This method thus offers a route to define graphene functionality without degrading its electronic properties. Furthermore, photopolymerization with styrene results in self-organized intercalative growth and exfoliation of few layer graphene sheets. Under these reaction conditions, a range of other vinyl monomers exhibits no reactivity with graphene. However, the authors demonstrate an alternative route by which the surface reac‐ tivity can be precisely tuned, and these monomers can be locally grafted via electron-beam-

involving the reduction of GO sheets that occurs in one of the synthetic steps.

*3.1.3. Irradiation-induced polymerization*

site materials.

induced carbon deposition on the graphene surface.

**Figure 5.** The overall synthesis route for the preparation of RAFT immobilized GO nanosheets [35].

**Figure 6.** Fabrication of PP/GO nanocomposites by in situ Ziegler-Natta polymerization [37].

Huang et al. [37] reports the first example of preparation of polypropylene/GO (PP/GO) nanocomposites via in situ Ziegler−Natta polymerization. As illustrated in Figure 6, a Mg/Ti catalyst species was immobilized onto GO sheets by reacting with the surface functional groups including −OH and −COOH. Subsequent propylene polymerization led to the in situ formation of PP matrix, which was accompanied by the nanoscale exfoliation of GO. Inde‐ pendent of the opposing nature of the polymer and GO, a good dispersion of GO sheets in PP matrix was observed, which was verified by morphological examination through TEM and SEM observation. Furthermore, high electrical conductivity was discovered with thus prepared PP/GO nanocomposites, this being the only paper reporting conductive materials prepared by grafting a polymer from graphene sheets. For example, at a GO loading of 4.9 wt %, σc was measured at 0.3 S m−1. We believe that this must originate from a side reaction involving the reduction of GO sheets that occurs in one of the synthetic steps.

#### *3.1.3. Irradiation-induced polymerization*

**Figure 5.** The overall synthesis route for the preparation of RAFT immobilized GO nanosheets [35].

214 New Progress on Graphene Research

**Figure 6.** Fabrication of PP/GO nanocomposites by in situ Ziegler-Natta polymerization [37].

Huang et al. [37] reports the first example of preparation of polypropylene/GO (PP/GO) nanocomposites via in situ Ziegler−Natta polymerization. As illustrated in Figure 6, a Mg/Ti catalyst species was immobilized onto GO sheets by reacting with the surface functional groups including −OH and −COOH. Subsequent propylene polymerization led to the in situ formation of PP matrix, which was accompanied by the nanoscale exfoliation of GO. Inde‐ pendent of the opposing nature of the polymer and GO, a good dispersion of GO sheets in PP matrix was observed, which was verified by morphological examination through TEM and SEM observation. Furthermore, high electrical conductivity was discovered with thus prepared PP/GO nanocomposites, this being the only paper reporting conductive materials prepared by grafting a polymer from graphene sheets. For example, at a GO loading of 4.9 Ultrasound has found important applications in a diverse range of materials and chemi‐ cal syntheses [48, 51]. Both the physical and chemical effects of ultrasound arise from acous‐ tic cavitation: the formation, growth, and collapse of bubbles in liquids irradiated with high intensity ultrasound [48 - 50, 52]. Localized hot spots with temperature of ~5000K and pressures of hundreds of bars are generated during the bubble collapse within liquid, which can induce some chemical reactions that can't take place under normal conditions. Xu et al. [53] reported a convenient single-step sonication-induced approach for the preparation of polymer functionalized graphenes from graphite flakes and a reactive monomer, styrene. In this work, they showed that by choosing a reactive medium as the solvent, the com‐ bined mechanochemical effects of high intensity ultrasound can, in a single step, readily induce exfoliation of graphite to produce functionalized graphenes. Ultrasonic irradiation of graphite in styrene results in the mechanochemical exfoliation of graphite flakes to singlelayer and few-layer graphene sheets combined with functionalization of the graphene with PS chains (Figure 7). The PS chains are formed from sonochemically initiated radical poly‐ merization of styrene. They also tested a variety of other solvents, including toluene, ethyl‐ benzene, 1-dodecene, and 4-vinylpyridine to prepare functionalized graphenes. Only the easily polymerizable reactants containing vinyl groups, styrene and 4-vinylpyridine, lead to stable functionalized graphene. Such functionalized graphene have good stability and solubility in common organic solvents and have great potential for graphene-based compo‐ site materials.

Moreover, direct photografting reactions of vinyl monomers came into focus for the prepa‐ ration of stable polymer brushes. Two approaches in particular, the sequential "living" pho‐ topolymerization and the self-initiated photografting and photopolymerization (SIPGP), attracted the attention of numerous research groups because of the facile preparation and broad applicability. Steenackers et al. [54] showed that PS chains could covalently bound to graphene by the UV-induced polymerization of styrene (Figure 8). Photopolymerization oc‐ curs at existing defect sites and that there is no detectable disruption of the basal plane con‐ jugation of graphene. This method thus offers a route to define graphene functionality without degrading its electronic properties. Furthermore, photopolymerization with styrene results in self-organized intercalative growth and exfoliation of few layer graphene sheets. Under these reaction conditions, a range of other vinyl monomers exhibits no reactivity with graphene. However, the authors demonstrate an alternative route by which the surface reac‐ tivity can be precisely tuned, and these monomers can be locally grafted via electron-beaminduced carbon deposition on the graphene surface.

γ-ray irradiation-induced graft polymerization (Figure 9). Due to the full coverage of PVAc chains and solvated layer formation on GO sheets surface, which weakens the interlami‐ nar attraction of GO sheets, PVAc-functionalized GO was well dispersed in common organ‐ ic solvents, and the dispersions obtained were extremely stable at room temperature without

Surface Functionalization of Graphene with Polymers for Enhanced Properties

Furthermore, Lee at al. [56] have developed a method to selectively fluorinate graphene by irradiating fluoropolymer-covered graphene with a laser (Figure 10). Here the sp<sup>2</sup>

brized graphene would react with the active fluorine radicals, which was produced by photon-induced decomposition of the fluoropolymer under laser-irradiation, and form C-F bonds. However, this reaction only occurred in the laser-irradiated region. The kinetics of C–F bond formation is dependent on both the laser power and fluoropolymer thickness. Furthermore, the resistance of the graphene dramatically increased due to the fluorina‐ tion, while the basic skeletal structure of the carbon bonding network is maintained. This is an efficient method for isolating graphene devices because the laser irradiation on fluoro‐ polymer-covered graphene process produces fluorinated graphene with highly insulating

As commented previously, the "grafting from" method relies on the immobilization of ini‐ tiators at the surface of graphene, followed by in situ surface-initiated polymerization to generate tethered polymer chains. However, this may not be possible in certain cases, where the covalent linkage between the presynthesized polymer and graphene emerges as the only alternative. In order to expand the type of polymers that can be bound to graphene, the cate‐ gory of "grafting-to" method can be employed to achieve this purpose. The "grafting to" technique involves the bonding of preformed end functionalized polymer chains to the sur‐ face of graphene. Therefore, the prepared graphene require adequate functional groups, which could react with specific polymers. Or the polymer has functionalities capable of re‐ acting with either graphene or its chemically broader cousin, GO. In the following part, we will summarize the type of reactions and the families of polymers that have been grafted to

**Figure 10.** The scheme showing a mechanism for fluorination under laser-irradiation.56

**3.2. Functionalizations via "grafting to" method**


217

http://dx.doi.org/10.5772/50490

any aggregation.

properties in a single step.

the graphene.

**Figure 7.** Experimental setup of the one-step mechanochemical process for exfoliation of graphite and sonochemical functionalization of graphene [53].

**Figure 8.** Patterned polymer brush layers on CVD-grown single layer graphene are prepared by UV illumination through a mask in bulk styrene. Surface photopolymerization occurs selectively in illuminated regions of the material [54].

**Figure 9.** The preparation of PVAc grafted GO by γ-ray irradiation-induced graft polymerization [55].

γ-ray radiation-induced graft polymerization has many advantages, including being a singlestep chemical reaction, needing no additives or catalysts, being conducted at room temper‐ ature, cost-effective, and so on. Above all, it is versatile for vinyl monomers that undergo free radiation polymerization, and production can be easily scaled-up. Zhang et al. [55] reported a facile approach to functionalize GO sheets with poly(vinyl acetate) (PVAc) by γ-ray irradiation-induced graft polymerization (Figure 9). Due to the full coverage of PVAc chains and solvated layer formation on GO sheets surface, which weakens the interlami‐ nar attraction of GO sheets, PVAc-functionalized GO was well dispersed in common organ‐ ic solvents, and the dispersions obtained were extremely stable at room temperature without any aggregation.

Furthermore, Lee at al. [56] have developed a method to selectively fluorinate graphene by irradiating fluoropolymer-covered graphene with a laser (Figure 10). Here the sp<sup>2</sup> -hy‐ brized graphene would react with the active fluorine radicals, which was produced by photon-induced decomposition of the fluoropolymer under laser-irradiation, and form C-F bonds. However, this reaction only occurred in the laser-irradiated region. The kinetics of C–F bond formation is dependent on both the laser power and fluoropolymer thickness. Furthermore, the resistance of the graphene dramatically increased due to the fluorina‐ tion, while the basic skeletal structure of the carbon bonding network is maintained. This is an efficient method for isolating graphene devices because the laser irradiation on fluoro‐ polymer-covered graphene process produces fluorinated graphene with highly insulating properties in a single step.

**Figure 10.** The scheme showing a mechanism for fluorination under laser-irradiation.56

#### **3.2. Functionalizations via "grafting to" method**

**Figure 7.** Experimental setup of the one-step mechanochemical process for exfoliation of graphite and sonochemical

**Figure 8.** Patterned polymer brush layers on CVD-grown single layer graphene are prepared by UV illumination through a mask in bulk styrene. Surface photopolymerization occurs selectively in illuminated regions of the material [54].

**Figure 9.** The preparation of PVAc grafted GO by γ-ray irradiation-induced graft polymerization [55].

γ-ray radiation-induced graft polymerization has many advantages, including being a singlestep chemical reaction, needing no additives or catalysts, being conducted at room temper‐ ature, cost-effective, and so on. Above all, it is versatile for vinyl monomers that undergo free radiation polymerization, and production can be easily scaled-up. Zhang et al. [55] reported a facile approach to functionalize GO sheets with poly(vinyl acetate) (PVAc) by

functionalization of graphene [53].

216 New Progress on Graphene Research

As commented previously, the "grafting from" method relies on the immobilization of ini‐ tiators at the surface of graphene, followed by in situ surface-initiated polymerization to generate tethered polymer chains. However, this may not be possible in certain cases, where the covalent linkage between the presynthesized polymer and graphene emerges as the only alternative. In order to expand the type of polymers that can be bound to graphene, the cate‐ gory of "grafting-to" method can be employed to achieve this purpose. The "grafting to" technique involves the bonding of preformed end functionalized polymer chains to the sur‐ face of graphene. Therefore, the prepared graphene require adequate functional groups, which could react with specific polymers. Or the polymer has functionalities capable of re‐ acting with either graphene or its chemically broader cousin, GO. In the following part, we will summarize the type of reactions and the families of polymers that have been grafted to the graphene.

#### *3.2.1. Esterification/amidation reactions*

Esterification/amidation reactions between carboxylic groups in GO and hydroxyl or amine groups in the polymer have been widely investigated [14, 15, 44, 57, 60]. In this respect, poly(vinyl alcohol) (PVA) was covalently bonded to GO [14, 57] and r-GO [14] by using a typical catalytic system for esterification (Figure 11). After functionalization with PVA chains, the solution processability of graphene was significantly improved. And the degree of functionalization was shown to be low, probably due to steric hindrance caused by the huge volume of GO. However, due to the presence of the huge graphene sheets, significant changes in the crystalline properties as well as in the tacticity of the polymer were observed. The originally semicrystalline PVA became completely amorphous, and the Tg increased by 35 °C after bonded to GO sheets. The decrease in crystallinity was attributed to the intercala‐ tion of PVA chains between the graphene sheets as well as the formation of "secondary" bonds, for example, hydrogen bonding that breaks intra- and interchain bonds. Finally, it has been demonstrated that the reaction is favoured at specific conformations at the isotactic sequences where the hydroxyl groups are more exposed (lower internal steric hindrance) than in the syndiotactic counterpart.

alized GO. The existence of these secondary bonds can lead to some additional ordering that

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219

**Figure 12.** Comparison of (a) storage modulus and (b) tan δ curves for neat PVC (square), PVC functionalized CNTs

Furthermore, conjugated polymer-functionalized graphene materials have also been pre‐ pared by esterification/amidation reactions [60, 63]. In these cases, the ends of the conjugat‐ ed polymers were bonded to the functional groups on the graphene sheets. As a result, the solubility of the obtained functionalized graphene was significantly improved in common solvents, enabling device preparation by solution processing. Thus, GO functionalized with both triphenylamine-based polyazomethine-modified GO (TPAPAM-GO) and poly(3-hex‐ ylthiophene) modified GO (P3HT-GO) (Figure 13) can be incorporated into specific devices by simple spin coating to obtain composites that exhibit non-volatile memory effect as well as higher power conversion efficiency for solar cells, demonstrated in the cases of TPAPAM-

**Figure 13.** Synthesis procedure for chemical grafting of CH2OH-terminated P3HT chains onto graphene, which in‐ volves the SOCl2 treatment of GO (step 1) and the esterification reaction between acyl-chloride functionalized GO and

(triangle), PVC functionalized GO (solid circle), and PVC functionalized r-GO (open circle) [59].

GO and P3HT-GO, respectively.

MeOH-terminated P3HT (step 2) [47].

alters the segmental mobility and consequently the final properties.

**Figure 11.** Schematic Illustration of the Esterification of GO with PVA [14].

A similar strategy has been approached to functionalize r-GO with poly(vinyl chloride) (PVC) [59]. In this step, the susceptible groups in PVC chains could react with the functional groups on r-GO sheets by esterification, which was provided by a nucleophilic substitution reaction [61, 62]. Furthermore, several methodologies to prepare r-GO/PVC nanocomposites and the optimum conditions have been established. The covalent attachment of r-GO to ap‐ propriately functionalized PVC is the only effective method to produce nanocomposites with improved thermal and mechanical properties (Figure 12). The absolute values of the mechanical and thermal properties of PVC-functionalized GO nanocomposites are higher than those for a similar system using MWNTs as reinforcement because of the higher aspect ratio of the r-GO sheets with respect to the MWNTs. The introduction of r-GO also increases the Tg of the composites, reflecting the changes in the mobility of the PVC chains. However, due to the lower strength of the "secondary bonds", such as halogen bonding and hydrogen bonding for PVC and PVA respectively, the changes in Tg for PVC were much lower than those reported for PVA, which had a similar degree of functionalization with PVC-function‐ alized GO. The existence of these secondary bonds can lead to some additional ordering that alters the segmental mobility and consequently the final properties.

*3.2.1. Esterification/amidation reactions*

218 New Progress on Graphene Research

than in the syndiotactic counterpart.

**Figure 11.** Schematic Illustration of the Esterification of GO with PVA [14].

Esterification/amidation reactions between carboxylic groups in GO and hydroxyl or amine groups in the polymer have been widely investigated [14, 15, 44, 57, 60]. In this respect, poly(vinyl alcohol) (PVA) was covalently bonded to GO [14, 57] and r-GO [14] by using a typical catalytic system for esterification (Figure 11). After functionalization with PVA chains, the solution processability of graphene was significantly improved. And the degree of functionalization was shown to be low, probably due to steric hindrance caused by the huge volume of GO. However, due to the presence of the huge graphene sheets, significant changes in the crystalline properties as well as in the tacticity of the polymer were observed. The originally semicrystalline PVA became completely amorphous, and the Tg increased by 35 °C after bonded to GO sheets. The decrease in crystallinity was attributed to the intercala‐ tion of PVA chains between the graphene sheets as well as the formation of "secondary" bonds, for example, hydrogen bonding that breaks intra- and interchain bonds. Finally, it has been demonstrated that the reaction is favoured at specific conformations at the isotactic sequences where the hydroxyl groups are more exposed (lower internal steric hindrance)

A similar strategy has been approached to functionalize r-GO with poly(vinyl chloride) (PVC) [59]. In this step, the susceptible groups in PVC chains could react with the functional groups on r-GO sheets by esterification, which was provided by a nucleophilic substitution reaction [61, 62]. Furthermore, several methodologies to prepare r-GO/PVC nanocomposites and the optimum conditions have been established. The covalent attachment of r-GO to ap‐ propriately functionalized PVC is the only effective method to produce nanocomposites with improved thermal and mechanical properties (Figure 12). The absolute values of the mechanical and thermal properties of PVC-functionalized GO nanocomposites are higher than those for a similar system using MWNTs as reinforcement because of the higher aspect ratio of the r-GO sheets with respect to the MWNTs. The introduction of r-GO also increases the Tg of the composites, reflecting the changes in the mobility of the PVC chains. However, due to the lower strength of the "secondary bonds", such as halogen bonding and hydrogen bonding for PVC and PVA respectively, the changes in Tg for PVC were much lower than those reported for PVA, which had a similar degree of functionalization with PVC-function‐

**Figure 12.** Comparison of (a) storage modulus and (b) tan δ curves for neat PVC (square), PVC functionalized CNTs (triangle), PVC functionalized GO (solid circle), and PVC functionalized r-GO (open circle) [59].

Furthermore, conjugated polymer-functionalized graphene materials have also been pre‐ pared by esterification/amidation reactions [60, 63]. In these cases, the ends of the conjugat‐ ed polymers were bonded to the functional groups on the graphene sheets. As a result, the solubility of the obtained functionalized graphene was significantly improved in common solvents, enabling device preparation by solution processing. Thus, GO functionalized with both triphenylamine-based polyazomethine-modified GO (TPAPAM-GO) and poly(3-hex‐ ylthiophene) modified GO (P3HT-GO) (Figure 13) can be incorporated into specific devices by simple spin coating to obtain composites that exhibit non-volatile memory effect as well as higher power conversion efficiency for solar cells, demonstrated in the cases of TPAPAM-GO and P3HT-GO, respectively.

**Figure 13.** Synthesis procedure for chemical grafting of CH2OH-terminated P3HT chains onto graphene, which in‐ volves the SOCl2 treatment of GO (step 1) and the esterification reaction between acyl-chloride functionalized GO and MeOH-terminated P3HT (step 2) [47].

#### *3.2.2. Nitrene cycloaddition*

Nitrene chemistry, an approach used to functionalize graphene with single molecules [64, 65], has also been extended to polymers [66, 67].Nitrene chemistry is a versatile tool that al‐ lows the functionalization of graphene with a pool of functionalities, potentiating graphene solubility, and dispersion in a wide variety polymeric matrices. By using of this technology, He and Gao [67] have reported a general and versatile approach to graft of polymers onto graphene sheets. In their experiment, a wide range of immobilized functional groups were used to graft specific polymers from it (Figure 14). Though the cycloadditions of nitrene rad‐ icals and thermal reduction of GO occurs simultaneously, the conductivity of graphene di‐ minished after functionalization, values of around 300-700 S/m for PS and poly(ethylene glycol) (PEG)-functioanlzied graphene were obtained. This was mainly due to the high amounts of graphene in the final product, which is reasonable because although the poly‐ mer is linked to graphene by the ends, the molecular weight of the polymers is low, making the mass percentage of graphene high.

irradiation of graphene can cause a considerable amount of defects on the graphene surface, which might produce reactive sites in situ as a result of the high temperature and pressure during bubble collapse. Moreover, the original defects on pristine graphene can also be easi‐ ly destroyed and produce reactive sites during ultrasonic irradiation. The PVA chain radi‐ cals produced by sonochemical degradation of the PVA solution can react easily with graphene, because of the reactive sites formed on the graphene surface, and readily func‐ tionalized them via "grafting to" method. The content of PVA on graphene was estimated to be ~35%. The f-G could be well dispersed in the PVA matrix by a simple solution mixing and casting procedure. Due to the effective load transfer between f-G and PVA matrix, the mechanical properties of the f-G/PVA films were significantly improved. Compared with the p-G/PVA films, a 12.6% increase in tensile strength and a 15.6% improvement of Young's modulus were achieved by addition of only 0.3 wt% f-G. Moreover, this simple ultrasonica‐

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tion technique could enable us to functionalize graphene with other polymers.

**Figure 15.** Schematic illustration of the sonochemical preparation process of PVA-functionalized graphene [72].

part of the main chain that can be directly related to changes in the final properties.

**4. Non-covalent functionalization of graphene with polymers**

As a summary, "grafting-to" methods are highly versatile since they take advantage of the chemistry of GO that can be appropriately modified with a wide variety of functional groups providing a capacity for reaction with almost any type of polymers. In addition, "grafting-to" method allows the selection of the location of graphene, that is, at the end or as

As shown in the aforementioned examples, the covalent functionalization of polymers on graphene-based sheets holds versatile possibility due to the rich surface chemistry of GO/r-GO. Nevertheless, the non-covalent functionalization, which almost relies on hydrogen bonding or π–π stacking, is easier to carry out without altering the chemical structure of the capped r-GO sheets, and provides effective means to tailor the electronic/optical property and solubility of the nanosheets. The first example of non-covalent functionalization of r-GO sheets was demonstrated by the in situ reduction of GO with hydrazine in the presence of

**Figure 14.** General strategy for the preparation of functionalized graphene sheets by nitrene chemistry and the fur‐ ther chemical modifications [67].

Besides the methods mentioned above, other "grafting to" approaches have also been inves‐ tigated, such as the opening of maleic rings in maleic acid (MA) grafted polyethylene by amine functionalized graphene [68], nucleophilic epoxy-ring opening in GO by amine groups in biocompatible poly-l-lysine [69], atom transfer nitroxide radical coupling (ATNRP) of PNIPAM and 2,2,6,6-tetramethylpiperidine-1-oxyl-modified graphene [70], and simultaneous reduction of GO and radical grafting of PMMA by phase transfer [71].

#### *3.2.3. Irradiation-induced radical grafting*

Shen et al. [72] have reported PVA-functionalized graphene (f-G) could be prepared by ul‐ trasonication of pristine graphene (p-G) in a PVA aqueous solution (Figure 15). Ultrasonic irradiation of graphene can cause a considerable amount of defects on the graphene surface, which might produce reactive sites in situ as a result of the high temperature and pressure during bubble collapse. Moreover, the original defects on pristine graphene can also be easi‐ ly destroyed and produce reactive sites during ultrasonic irradiation. The PVA chain radi‐ cals produced by sonochemical degradation of the PVA solution can react easily with graphene, because of the reactive sites formed on the graphene surface, and readily func‐ tionalized them via "grafting to" method. The content of PVA on graphene was estimated to be ~35%. The f-G could be well dispersed in the PVA matrix by a simple solution mixing and casting procedure. Due to the effective load transfer between f-G and PVA matrix, the mechanical properties of the f-G/PVA films were significantly improved. Compared with the p-G/PVA films, a 12.6% increase in tensile strength and a 15.6% improvement of Young's modulus were achieved by addition of only 0.3 wt% f-G. Moreover, this simple ultrasonica‐ tion technique could enable us to functionalize graphene with other polymers.

*3.2.2. Nitrene cycloaddition*

220 New Progress on Graphene Research

the mass percentage of graphene high.

ther chemical modifications [67].

*3.2.3. Irradiation-induced radical grafting*

Nitrene chemistry, an approach used to functionalize graphene with single molecules [64, 65], has also been extended to polymers [66, 67].Nitrene chemistry is a versatile tool that al‐ lows the functionalization of graphene with a pool of functionalities, potentiating graphene solubility, and dispersion in a wide variety polymeric matrices. By using of this technology, He and Gao [67] have reported a general and versatile approach to graft of polymers onto graphene sheets. In their experiment, a wide range of immobilized functional groups were used to graft specific polymers from it (Figure 14). Though the cycloadditions of nitrene rad‐ icals and thermal reduction of GO occurs simultaneously, the conductivity of graphene di‐ minished after functionalization, values of around 300-700 S/m for PS and poly(ethylene glycol) (PEG)-functioanlzied graphene were obtained. This was mainly due to the high amounts of graphene in the final product, which is reasonable because although the poly‐ mer is linked to graphene by the ends, the molecular weight of the polymers is low, making

**Figure 14.** General strategy for the preparation of functionalized graphene sheets by nitrene chemistry and the fur‐

Besides the methods mentioned above, other "grafting to" approaches have also been inves‐ tigated, such as the opening of maleic rings in maleic acid (MA) grafted polyethylene by amine functionalized graphene [68], nucleophilic epoxy-ring opening in GO by amine groups in biocompatible poly-l-lysine [69], atom transfer nitroxide radical coupling (ATNRP) of PNIPAM and 2,2,6,6-tetramethylpiperidine-1-oxyl-modified graphene [70], and

Shen et al. [72] have reported PVA-functionalized graphene (f-G) could be prepared by ul‐ trasonication of pristine graphene (p-G) in a PVA aqueous solution (Figure 15). Ultrasonic

simultaneous reduction of GO and radical grafting of PMMA by phase transfer [71].

**Figure 15.** Schematic illustration of the sonochemical preparation process of PVA-functionalized graphene [72].

As a summary, "grafting-to" methods are highly versatile since they take advantage of the chemistry of GO that can be appropriately modified with a wide variety of functional groups providing a capacity for reaction with almost any type of polymers. In addition, "grafting-to" method allows the selection of the location of graphene, that is, at the end or as part of the main chain that can be directly related to changes in the final properties.

#### **4. Non-covalent functionalization of graphene with polymers**

As shown in the aforementioned examples, the covalent functionalization of polymers on graphene-based sheets holds versatile possibility due to the rich surface chemistry of GO/r-GO. Nevertheless, the non-covalent functionalization, which almost relies on hydrogen bonding or π–π stacking, is easier to carry out without altering the chemical structure of the capped r-GO sheets, and provides effective means to tailor the electronic/optical property and solubility of the nanosheets. The first example of non-covalent functionalization of r-GO sheets was demonstrated by the in situ reduction of GO with hydrazine in the presence of poly(sodium 4-styrenesulfonate) (PSS) [12], in which the hydrophobic backbone of PSS sta‐ bilizes the r-GO, and the hydrophilic sulfonate side groups maintains a good dispersion of the hybrid nanosheets in water.

Liu et al. synthesized thermoresponsive graphene-polymer nanocomposites. They first took advantage of RAFT polymerization to synthesize a well-defined thermoresponsive pyrene terminated poly(N-isopropylacrylamide) (PNIPAAm), followed by attachment onto the bas‐ al plane of graphene sheets via π-π stacking interactions (Figure 16) [81]. The lower critical solution temperature (LCST) of pyrene-terminated PNIPAAm was measured to be 33 ° C. However, after the pyrene-functional polymer functionalized with graphene sheets, the re‐ sultant graphene composites were also thermoresponsive in aqueous solutions, but with a

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Similarly, Liu et al. also prepared pH sensitive graphene-polymer composites by functionali‐ zation of graphene with a pyrene-terminated positive charged polymer, poly(2-N,N'-(di‐ methyl amino ethyl acrylate)) (PDMAEA), and a negatively charged polymer, polyacrylic acid (PAA) [80]. During the process, a pyrene-terminated RAFT agent was used to prepare the pyrene-terminated PDMAEA and PAA. When manipulating the pH of the graphene −composite suspensions, phase transfer between the aqueous and organic phases was ob‐ served. Self-assembly of the two oppositely charged graphene-polymer composites afforded layer-by-layer (LbL) structures as evidenced by high-resolution scanning electron microsco‐ py (SEM) and quartz crystal microbalance (QCM) measurements (Figure 17). In addition to RAFT mechanism, π-orbital rich polymers have also been synthesized by using of ATRP method for functionalization of r-GO to afford the fluorescent and water-soluble graphene

**Figure 17.** Synthesis of pH sensitive pyrene-polymer composites via π-π stacking interactions for the self-assembly of

Moreover, conjugated polyelectrolytes with various functionalities have been used to modi‐ fy r-GO nanosheets [83 - 85], in the hope to achieve good solubility in different kinds of sol‐

lower LCST of 24 °C.

composites via π-π stacking interactions [82].

functionalized graphene into layered structures [80].

*4.1.2. Conjugated polyelectrolytes*

#### **4.1. Functionalizations via π–π stacking interactions**

π–π stacking interactions usually occur between two relatively large non-polar aromatic rings having overlapping π orbitals. They can be comparable to covalent attachment in strength and hence provide more stable alternatives to the weaker hydrogen bonding, elec‐ trostatic bonding and coordination bonding strategies. Furthermore, π–π stacking function‐ alization does not disrupt the conjugation of the graphene sheets, and hence preserves the electronic properties of graphene.

#### *4.1.1. Polymers with pyrene end-groups*

In order to functionalize graphene with polymers via π–π stacking, one strategy is for the polymer chains to be synthesized with pyrene moieties as the termini of the polymer chains. RAFT polymerization can be a useful tool to achieve this aim. Polymers with pyrene endgroups have been made using RAFT mechanism in several recent papers [73, 80].

**Figure 16.** A schematic depicting the synthesis of pyrene-terminated PNIPAAm using a pyrene-functional RAFT agent and the subsequent attachment of the polymer to graphene [ 81].

Liu et al. synthesized thermoresponsive graphene-polymer nanocomposites. They first took advantage of RAFT polymerization to synthesize a well-defined thermoresponsive pyrene terminated poly(N-isopropylacrylamide) (PNIPAAm), followed by attachment onto the bas‐ al plane of graphene sheets via π-π stacking interactions (Figure 16) [81]. The lower critical solution temperature (LCST) of pyrene-terminated PNIPAAm was measured to be 33 ° C. However, after the pyrene-functional polymer functionalized with graphene sheets, the re‐ sultant graphene composites were also thermoresponsive in aqueous solutions, but with a lower LCST of 24 °C.

Similarly, Liu et al. also prepared pH sensitive graphene-polymer composites by functionali‐ zation of graphene with a pyrene-terminated positive charged polymer, poly(2-N,N'-(di‐ methyl amino ethyl acrylate)) (PDMAEA), and a negatively charged polymer, polyacrylic acid (PAA) [80]. During the process, a pyrene-terminated RAFT agent was used to prepare the pyrene-terminated PDMAEA and PAA. When manipulating the pH of the graphene −composite suspensions, phase transfer between the aqueous and organic phases was ob‐ served. Self-assembly of the two oppositely charged graphene-polymer composites afforded layer-by-layer (LbL) structures as evidenced by high-resolution scanning electron microsco‐ py (SEM) and quartz crystal microbalance (QCM) measurements (Figure 17). In addition to RAFT mechanism, π-orbital rich polymers have also been synthesized by using of ATRP method for functionalization of r-GO to afford the fluorescent and water-soluble graphene composites via π-π stacking interactions [82].

**Figure 17.** Synthesis of pH sensitive pyrene-polymer composites via π-π stacking interactions for the self-assembly of functionalized graphene into layered structures [80].

#### *4.1.2. Conjugated polyelectrolytes*

poly(sodium 4-styrenesulfonate) (PSS) [12], in which the hydrophobic backbone of PSS sta‐ bilizes the r-GO, and the hydrophilic sulfonate side groups maintains a good dispersion of

π–π stacking interactions usually occur between two relatively large non-polar aromatic rings having overlapping π orbitals. They can be comparable to covalent attachment in strength and hence provide more stable alternatives to the weaker hydrogen bonding, elec‐ trostatic bonding and coordination bonding strategies. Furthermore, π–π stacking function‐ alization does not disrupt the conjugation of the graphene sheets, and hence preserves the

In order to functionalize graphene with polymers via π–π stacking, one strategy is for the polymer chains to be synthesized with pyrene moieties as the termini of the polymer chains. RAFT polymerization can be a useful tool to achieve this aim. Polymers with pyrene end-

**Figure 16.** A schematic depicting the synthesis of pyrene-terminated PNIPAAm using a pyrene-functional RAFT agent

and the subsequent attachment of the polymer to graphene [ 81].

groups have been made using RAFT mechanism in several recent papers [73, 80].

the hybrid nanosheets in water.

222 New Progress on Graphene Research

electronic properties of graphene.

*4.1.1. Polymers with pyrene end-groups*

**4.1. Functionalizations via π–π stacking interactions**

Moreover, conjugated polyelectrolytes with various functionalities have been used to modi‐ fy r-GO nanosheets [83 - 85], in the hope to achieve good solubility in different kinds of sol‐ vents, and at the same time acquire added optoelectronic properties. Qi et al. has specially designed an amphiphilic coil–rod–coil conjugated triblock copolymer (PEG-OPE, chemical structure shown in Figure 18A) to improve the solubility of graphene-polymer nanocompo‐ sites in both high and low polar solvents [33, 83]. In the proposed configuration, the conju‐ gated rigid-rod backbone of PEG-OPE can bind to the basal plane of the r-GO via the π-π stacking interaction (Figure 18B), whereas the lipophilic side chains and two hydrophilic coils of the backbone form an amphiphilic outer-layer surrounding the r-GO sheet. As a re‐ sult, the obtained r-GO sheets with a uniformly coated polymer layer (Figure 18C) are solu‐ ble in both organic low polar (such as toluene and chloroform) and water-miscible high polar solvents (such as water and ethanol).

In another study, Qi et al. demonstrated the preparation of highly soluble r-GO hybrid ma‐ terial (PFVSO3-r-GO) by taking advantage of strong π–π interactions between the anionic CPE and r-GO (Figure 19) [84]. The resulting CPE-functionalized r-GO (PFVSO3-r-GO) shows excellent solubility and stability in a variety of polar solvents, including water, etha‐ nol, methanol, dimethyl sulfoxide, and dimethyl formamide. The morphology of PFVSO3-r-GO is studied, which reveal a sandwich-like nanostructure. Within this nanostructure, the backbones of PFVSO3 stack onto the basal plane of r-GO sheets via strong π–π interactions, while the charged hydrophilic side chains of PFVSO3 prevent the rGO sheets from aggregat‐ ing via electrostatic and steric repulsions, thus leading to the solubility and stability of PFVSO3-rGO in polar solvents. Furthermore, the presence of PFVSO3 within r-GO induces photoinduced charge transfer and p-doping of r-GO. As a result, the electrical conductivity of PFVSO3-r-GO is not only much better than that of GO, but also than that of the unfunc‐

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225

High temperature and strong shear forces are usually involved during the melt blending process, which tends to fracture the nanoparticle aggregates, and endow polymer chains with the ability to diffuse into the gaps of the nanoparticle interlayer. Furthermore, as suggested by theoretical and experimental studies [86, 87] chemical or physical interac‐ tions can be formed between the fillers and the polymer components. Zhang et al [88] found the melt blending led to enhanced interactions between PS and CNTs, which was indicat‐ ed by increased amount of PS linked to CNTs and therefore dramatically increased solubil‐ ity of CNTs in some solvents. Taking advantage of this method, Zhou et al [89] obtained PS-coated CNTs through simple melt mixing of PS with CNTs. Furthermore, Lu et al [90] studied the styrene-butadiene-styrene tri-block copolymer (SBS)/CNTs composite, and their results showed that there were interactions between CNTs and SBS occurred during melt mixing, leading to an improvement of the mechanical properties of SBS/CNTs composites, as well as the homogeneous dispersion of CNTs in SBS. The mechanism of melt blending on these enhanced interactions was mainly attributed to the formation of π-π stacking between the aromatic system of π-electrons of PS and the π-electrons system of CNTs during

Melt Blending can also graft PS chains onto the surface of graphene sheets via π-π interac‐ tions. The interaction between graphene and PS was significantly enhanced by melt blend‐ ing, which led to an increased amount of PS-functional graphene (PSFG) exhibiting good solubility in some solvents [10]. The mechanism for the in-situ formation of π-π stacking was addressed, as illustrated in Figure 20. It was proposed that the strong shear action ap‐ plied by extruder could stretch the PS chains and endow the polymer chains with possibility to diffuse into the interlayer gap of graphene sheets. Moreover, the PS chains could be push‐ ed towards to graphene sheets to form the π-π stacking under high shear forces. The UV-vis absorption spectroscopy of PSFG presented an obvious red shift, suggested the presence of

tionalized r-GO.

melt blending [10, 88].

π-π stacking between PS and graphene.

*4.1.3. π–π stacking induced by melt blending*

**Figure 18.** A) Chemical structure of PEG-OPE. (B) Schematic illustration of fabrication of PEG-OPE stabilized r-GO sheets. (C) Tapping-mode AFM image and cross-sectional analysis of PEG-OPE-r-GO on mica [67].

**Figure 19.** A) Chemical structure of the newly designed PFVSO3. B) Schematic illustration of the synthesis of PFVSO3 stabilized r-GO in H2O: step 1, oxidative treatment of graphite (gray-black) yields single-layer GO sheets (brown); step 2, chemical reduction of GO with hydrazine in the presence of PFVSO3 produces a stable aqueous suspension of PFVSO3-functionalized r-GO sheets (PFVSO3-r-GO). C) Photograph of aqueous dispersions of GO (i), r-GO (ii), PFVSO3-r-GO (iii), and PFVSO3 (iv) [84].

In another study, Qi et al. demonstrated the preparation of highly soluble r-GO hybrid ma‐ terial (PFVSO3-r-GO) by taking advantage of strong π–π interactions between the anionic CPE and r-GO (Figure 19) [84]. The resulting CPE-functionalized r-GO (PFVSO3-r-GO) shows excellent solubility and stability in a variety of polar solvents, including water, etha‐ nol, methanol, dimethyl sulfoxide, and dimethyl formamide. The morphology of PFVSO3-r-GO is studied, which reveal a sandwich-like nanostructure. Within this nanostructure, the backbones of PFVSO3 stack onto the basal plane of r-GO sheets via strong π–π interactions, while the charged hydrophilic side chains of PFVSO3 prevent the rGO sheets from aggregat‐ ing via electrostatic and steric repulsions, thus leading to the solubility and stability of PFVSO3-rGO in polar solvents. Furthermore, the presence of PFVSO3 within r-GO induces photoinduced charge transfer and p-doping of r-GO. As a result, the electrical conductivity of PFVSO3-r-GO is not only much better than that of GO, but also than that of the unfunc‐ tionalized r-GO.

#### *4.1.3. π–π stacking induced by melt blending*

vents, and at the same time acquire added optoelectronic properties. Qi et al. has specially designed an amphiphilic coil–rod–coil conjugated triblock copolymer (PEG-OPE, chemical structure shown in Figure 18A) to improve the solubility of graphene-polymer nanocompo‐ sites in both high and low polar solvents [33, 83]. In the proposed configuration, the conju‐ gated rigid-rod backbone of PEG-OPE can bind to the basal plane of the r-GO via the π-π stacking interaction (Figure 18B), whereas the lipophilic side chains and two hydrophilic coils of the backbone form an amphiphilic outer-layer surrounding the r-GO sheet. As a re‐ sult, the obtained r-GO sheets with a uniformly coated polymer layer (Figure 18C) are solu‐ ble in both organic low polar (such as toluene and chloroform) and water-miscible high

**Figure 18.** A) Chemical structure of PEG-OPE. (B) Schematic illustration of fabrication of PEG-OPE stabilized r-GO

**Figure 19.** A) Chemical structure of the newly designed PFVSO3. B) Schematic illustration of the synthesis of PFVSO3 stabilized r-GO in H2O: step 1, oxidative treatment of graphite (gray-black) yields single-layer GO sheets (brown); step 2, chemical reduction of GO with hydrazine in the presence of PFVSO3 produces a stable aqueous suspension of PFVSO3-functionalized r-GO sheets (PFVSO3-r-GO). C) Photograph of aqueous dispersions of GO (i), r-GO (ii), PFVSO3-r-

sheets. (C) Tapping-mode AFM image and cross-sectional analysis of PEG-OPE-r-GO on mica [67].

polar solvents (such as water and ethanol).

224 New Progress on Graphene Research

GO (iii), and PFVSO3 (iv) [84].

High temperature and strong shear forces are usually involved during the melt blending process, which tends to fracture the nanoparticle aggregates, and endow polymer chains with the ability to diffuse into the gaps of the nanoparticle interlayer. Furthermore, as suggested by theoretical and experimental studies [86, 87] chemical or physical interac‐ tions can be formed between the fillers and the polymer components. Zhang et al [88] found the melt blending led to enhanced interactions between PS and CNTs, which was indicat‐ ed by increased amount of PS linked to CNTs and therefore dramatically increased solubil‐ ity of CNTs in some solvents. Taking advantage of this method, Zhou et al [89] obtained PS-coated CNTs through simple melt mixing of PS with CNTs. Furthermore, Lu et al [90] studied the styrene-butadiene-styrene tri-block copolymer (SBS)/CNTs composite, and their results showed that there were interactions between CNTs and SBS occurred during melt mixing, leading to an improvement of the mechanical properties of SBS/CNTs composites, as well as the homogeneous dispersion of CNTs in SBS. The mechanism of melt blending on these enhanced interactions was mainly attributed to the formation of π-π stacking between the aromatic system of π-electrons of PS and the π-electrons system of CNTs during melt blending [10, 88].

Melt Blending can also graft PS chains onto the surface of graphene sheets via π-π interac‐ tions. The interaction between graphene and PS was significantly enhanced by melt blend‐ ing, which led to an increased amount of PS-functional graphene (PSFG) exhibiting good solubility in some solvents [10]. The mechanism for the in-situ formation of π-π stacking was addressed, as illustrated in Figure 20. It was proposed that the strong shear action ap‐ plied by extruder could stretch the PS chains and endow the polymer chains with possibility to diffuse into the interlayer gap of graphene sheets. Moreover, the PS chains could be push‐ ed towards to graphene sheets to form the π-π stacking under high shear forces. The UV-vis absorption spectroscopy of PSFG presented an obvious red shift, suggested the presence of π-π stacking between PS and graphene.

fabricated a bienzyme biosensing system for the detection of maltose by successive LbL as‐ sembly of functionalized graphene, GOD, and glucoamylase (GA), which showed great

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227

**Figure 21.** Schematic Illustration of the Strategies for layer-by-layer assembly of graphene multilayer films for enzyme-

**Figure 22.** Schematic illustration of the electron-withdrawing from graphene by PDDA to facilitate the ORR process

promise in highly efficient sensors and advanced biosensing systems.

based glucose and maltose biosensing [95].

[97].

**Figure 20.** Schematic for the forming of π-π stacking in the process of melt blending [10].

#### **4.2. Functionalizations via hydrogen bonding**

Hydrogen bonding is very common in the biological world. An individual hydrogen bond is not very strong (2–8 kcal/mol); however, multiple hydrogen bonds will afford strong inter‐ actions as seen in DNA hybridization. Liang et al. prepared PVA nanocomposites with gra‐ phene by dispersing GO sheets into PVA matrix at molecular level [91]. The authors considered that the increased tensile strength and Young's modulus of the PVA/GO compo‐ site films were caused by the strong hydrogen bonding interactions between the residual oxygen-containing groups of GO sheets, such as epoxides and hydroxyls on their basal planes and carboxyls on the edges, and hydroxyl groups of the PVA chains. Polymer/ graphene nanocomposites with other hydrophilic polymers, epoxy, poly(acrylonitrile) and polyaniline exhibited extraordinary high increase in modulus or glass transition tempera‐ ture, attributed to hydrogen bonding interactions [92 - 94].

## **5. Applications of functionalized graphene**

The use of graphene or functionalized graphene materials usually exploits properties such as the large surface area or high electrical conductivity. Currently, the applications of func‐ tionalized graphene are focused on clean energy devices and electronic devices, as well as on sensors, medical devices and catalysis. Two examples will be discussed below.

Non-covalent chemical modification between functionalized pyrene and graphene was adopted to achieve patterned arrays of glucose oxidase (GOD) for potential applications in glucose sensors, cell sensors and tissue engineering [96]. In research by Zeng et al. as illus‐ trated in Figure 21, [95] chemically reduced GO was functionalized by pyrene-grafted poly(acrylic acid) (PAA) in aqueous solution on the basis of π-π stacking as well as van der Waals interactions. Then PAA-functionalized graphene (PAA-graphene) was LbL assem‐ bled with poly(ethyleneimine) (PEI). Graphene multilayer films facilitated the electron transfer, enhancing the electrochemical reactivity of H2O2. On the basis of this property, they fabricated a bienzyme biosensing system for the detection of maltose by successive LbL as‐ sembly of functionalized graphene, GOD, and glucoamylase (GA), which showed great promise in highly efficient sensors and advanced biosensing systems.

**Figure 20.** Schematic for the forming of π-π stacking in the process of melt blending [10].

Hydrogen bonding is very common in the biological world. An individual hydrogen bond is not very strong (2–8 kcal/mol); however, multiple hydrogen bonds will afford strong inter‐ actions as seen in DNA hybridization. Liang et al. prepared PVA nanocomposites with gra‐ phene by dispersing GO sheets into PVA matrix at molecular level [91]. The authors considered that the increased tensile strength and Young's modulus of the PVA/GO compo‐ site films were caused by the strong hydrogen bonding interactions between the residual oxygen-containing groups of GO sheets, such as epoxides and hydroxyls on their basal planes and carboxyls on the edges, and hydroxyl groups of the PVA chains. Polymer/ graphene nanocomposites with other hydrophilic polymers, epoxy, poly(acrylonitrile) and polyaniline exhibited extraordinary high increase in modulus or glass transition tempera‐

The use of graphene or functionalized graphene materials usually exploits properties such as the large surface area or high electrical conductivity. Currently, the applications of func‐ tionalized graphene are focused on clean energy devices and electronic devices, as well as

Non-covalent chemical modification between functionalized pyrene and graphene was adopted to achieve patterned arrays of glucose oxidase (GOD) for potential applications in glucose sensors, cell sensors and tissue engineering [96]. In research by Zeng et al. as illus‐ trated in Figure 21, [95] chemically reduced GO was functionalized by pyrene-grafted poly(acrylic acid) (PAA) in aqueous solution on the basis of π-π stacking as well as van der Waals interactions. Then PAA-functionalized graphene (PAA-graphene) was LbL assem‐ bled with poly(ethyleneimine) (PEI). Graphene multilayer films facilitated the electron transfer, enhancing the electrochemical reactivity of H2O2. On the basis of this property, they

on sensors, medical devices and catalysis. Two examples will be discussed below.

**4.2. Functionalizations via hydrogen bonding**

226 New Progress on Graphene Research

ture, attributed to hydrogen bonding interactions [92 - 94].

**5. Applications of functionalized graphene**

**Figure 21.** Schematic Illustration of the Strategies for layer-by-layer assembly of graphene multilayer films for enzymebased glucose and maltose biosensing [95].

**Figure 22.** Schematic illustration of the electron-withdrawing from graphene by PDDA to facilitate the ORR process [97].

The planar structure and superb conductivity of graphene also provide an appropriate plat‐ form for novel electrochemical sensors. A metal-free electrocatalyst for the oxygen reduction reaction (ORR) was achieved with graphene sheets functionalized with an electron acceptor, poly(diallyldimethylammonium chloride) (PDDA) (Figure 22). The resultant positively charged graphene composite was demonstrated to show remarkable electrocatalytic activity toward ORR with better fuel selectivity, tolerance to CO poisoning, and long-term stability than that of the commercially available Pt/C electrode [97]. The observed ORR electrocata‐ lytic activity induced by the intermolecular charge-transfer provides a general approach to various carbon-based metal-free ORR catalysts for oxygen reduction.

out without altering the chemical structure of the capped rGO sheets, and provides effective means to tailor the electronic/optical property and solubility of the nanosheets. π-π stacking interactions usually occur between two relatively large non-polar aromatic rings having overlapping π orbitals. In order to functionalize graphene with polymers via π-π stacking, one strategy is for the polymer chains to be synthesized with pyrene moieties as the termini of the polymer chains. Moreover, conjugated polyelectrolytes with various functionalities have been used to functionalize graphene via π-π stacking. The hydrogen bonding interac‐ tions between the residual oxygen-containing groups of graphene sheets and hydroxyl groups of the hydrophilic polymer chains, such as PVA chains, have also been used to func‐ tionalize the graphene in order to obtain increased tensile strength and Young's modulus.

Surface Functionalization of Graphene with Polymers for Enhanced Properties

http://dx.doi.org/10.5772/50490

229

Ningbo Key Lab of Polymer Materials, Ningbo Institute of Material Technology and Engi‐

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**Author details**

, Bin Shen and Wentao Zhai

\*Address all correspondence to: wgzheng@nimte.ac.cn

neering, Chinese Academy of Sciences, China

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Wenge Zheng\*

**References**

#### **6. Conclusion**

Generally speaking, functionalization of graphene with polymers can be achieved via ei‐ ther covalent or non-covalent interactions. Furthermore, the functionalization of graphene is always based on the graphene from previously prepared GO, which has multiple oxygencontaining functionalities, such as hydroxyl and carboxyl groups. In the covalent functional‐ izations, "grafting to" and "grafting from" techniques have been developed to graft the polymer chains onto the graphene surface. The "grafting from" method relies on the immo‐ bilization of initiators at the surface of graphene, followed by in situ surface-initiated poly‐ merization to generate tethered polymer chains. The "grafting to" technique involves the bonding of preformed end functionalized polymer chains to the surface of graphene. Com‐ paring both methods through the examples described above, it seems that "grafting-to" method allows the covalent bonding of a wider variety of polymers to graphene. The poly‐ mers "grafted from" graphene are those produced principally by some type of radical poly‐ merization, such as ATRP and RAFT. However, the most relevant polymers grafted from graphene, such as PS and PMMA have also been attached to graphene by the "grafting-to" method. In principle, for polymers obtained by ATRP the "grafting-from" method might be the most appropriate, but this depends on the features desired in the final composite, be‐ cause through "grafting to" method graphene forms part of the bulk polymer whereas in the "grafting from" it is limited to a terminal group. Moreover, except for a few exceptions, functional polymers and polymers synthesized by condensation reactions are principally bound to graphene by the grafting-to approach. Regarding to use PVA and PVC to function‐ alize graphene, we select the "grafting to" method due to the experimental procedures em‐ ployed to obtain these polymers. In fact, it is well known that PVA is not prepared by polymerization of the corresponding monomer, and the majority of PVC is synthesized by free-radical polymerization through suspension or bulk processes. Finally, radiation-in‐ duced graft polymerization has been utilized in the "grafting to" and "grafting from" tech‐ niques, due to its advantages, including being a single-step chemical reaction, needing no additives or catalysts, being conducted at room temperature, cost-effective, and so on.

The covalent functionalization of polymers on graphene-based sheets holds versatile possi‐ bility due to the rich surface chemistry of GO/r-GO. Nevertheless, the non-covalent func‐ tionalization, which almost relies on hydrogen bonding or π-π stacking, is easier to carry out without altering the chemical structure of the capped rGO sheets, and provides effective means to tailor the electronic/optical property and solubility of the nanosheets. π-π stacking interactions usually occur between two relatively large non-polar aromatic rings having overlapping π orbitals. In order to functionalize graphene with polymers via π-π stacking, one strategy is for the polymer chains to be synthesized with pyrene moieties as the termini of the polymer chains. Moreover, conjugated polyelectrolytes with various functionalities have been used to functionalize graphene via π-π stacking. The hydrogen bonding interac‐ tions between the residual oxygen-containing groups of graphene sheets and hydroxyl groups of the hydrophilic polymer chains, such as PVA chains, have also been used to func‐ tionalize the graphene in order to obtain increased tensile strength and Young's modulus.

### **Author details**

The planar structure and superb conductivity of graphene also provide an appropriate plat‐ form for novel electrochemical sensors. A metal-free electrocatalyst for the oxygen reduction reaction (ORR) was achieved with graphene sheets functionalized with an electron acceptor, poly(diallyldimethylammonium chloride) (PDDA) (Figure 22). The resultant positively charged graphene composite was demonstrated to show remarkable electrocatalytic activity toward ORR with better fuel selectivity, tolerance to CO poisoning, and long-term stability than that of the commercially available Pt/C electrode [97]. The observed ORR electrocata‐ lytic activity induced by the intermolecular charge-transfer provides a general approach to

Generally speaking, functionalization of graphene with polymers can be achieved via ei‐ ther covalent or non-covalent interactions. Furthermore, the functionalization of graphene is always based on the graphene from previously prepared GO, which has multiple oxygencontaining functionalities, such as hydroxyl and carboxyl groups. In the covalent functional‐ izations, "grafting to" and "grafting from" techniques have been developed to graft the polymer chains onto the graphene surface. The "grafting from" method relies on the immo‐ bilization of initiators at the surface of graphene, followed by in situ surface-initiated poly‐ merization to generate tethered polymer chains. The "grafting to" technique involves the bonding of preformed end functionalized polymer chains to the surface of graphene. Com‐ paring both methods through the examples described above, it seems that "grafting-to" method allows the covalent bonding of a wider variety of polymers to graphene. The poly‐ mers "grafted from" graphene are those produced principally by some type of radical poly‐ merization, such as ATRP and RAFT. However, the most relevant polymers grafted from graphene, such as PS and PMMA have also been attached to graphene by the "grafting-to" method. In principle, for polymers obtained by ATRP the "grafting-from" method might be the most appropriate, but this depends on the features desired in the final composite, be‐ cause through "grafting to" method graphene forms part of the bulk polymer whereas in the "grafting from" it is limited to a terminal group. Moreover, except for a few exceptions, functional polymers and polymers synthesized by condensation reactions are principally bound to graphene by the grafting-to approach. Regarding to use PVA and PVC to function‐ alize graphene, we select the "grafting to" method due to the experimental procedures em‐ ployed to obtain these polymers. In fact, it is well known that PVA is not prepared by polymerization of the corresponding monomer, and the majority of PVC is synthesized by free-radical polymerization through suspension or bulk processes. Finally, radiation-in‐ duced graft polymerization has been utilized in the "grafting to" and "grafting from" tech‐ niques, due to its advantages, including being a single-step chemical reaction, needing no additives or catalysts, being conducted at room temperature, cost-effective, and so on.

The covalent functionalization of polymers on graphene-based sheets holds versatile possi‐ bility due to the rich surface chemistry of GO/r-GO. Nevertheless, the non-covalent func‐ tionalization, which almost relies on hydrogen bonding or π-π stacking, is easier to carry

various carbon-based metal-free ORR catalysts for oxygen reduction.

**6. Conclusion**

228 New Progress on Graphene Research

Wenge Zheng\* , Bin Shen and Wentao Zhai

\*Address all correspondence to: wgzheng@nimte.ac.cn

Ningbo Key Lab of Polymer Materials, Ningbo Institute of Material Technology and Engi‐ neering, Chinese Academy of Sciences, China

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**Chapter 9**

**Graphene Nanowalls**

http://dx.doi.org/10.5772/51528

**1. Introduction**

Mineo Hiramatsu, Hiroki Kondo and Masaru Hori

Graphene, hexagonal arrangement of carbon atoms forming one-atom thick planar sheet, is a promising material for future electronic applications due to their high electrical conductiv‐ ity as well as chemical and physical stability [1]. Planar graphene films with respect to the substrate have been synthesized using various methods including mechanical exfoliation from highly oriented pyrolytic graphite, chemical exfoliation from bulk graphite, thermal decomposition of carbon-terminated silicon carbide, and chemical vapor deposition (CVD)

On the other hand, plasma-enhanced CVD (PECVD) is among the early methods to synthe‐ size of vertically standing few layer graphenes or graphene nanowalls (GNWs) [7–12]. GNWs can be described as self-assembled, vertically standing, few-layered graphene sheet nanostructures, which are also called as carbon nanowalls, carbon nanosheets, and carbon nanoflakes. As illustrated in Fig. 1, the sheets form a self-supported network of wall struc‐ tures with thicknesses ranging from a few nanometers to a few tens of nanometers. GNWs have a high density of atomic scale graphitic edges that are potential sites for electron field emission, which might lead to the application in flat panel displays and light sources [13,14].

> © 2013 Hiramatsu et al.; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

© 2013 Hiramatsu et al.; licensee InTech. This is a paper distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Additional information is available at the end of the chapter

on metals such as nickel and copper substrates [2–6].

**Figure 1.** Schematic illustration of graphene nanowalls


## **Graphene Nanowalls**

Mineo Hiramatsu, Hiroki Kondo and Masaru Hori

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/51528

## **1. Introduction**

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Graphene, hexagonal arrangement of carbon atoms forming one-atom thick planar sheet, is a promising material for future electronic applications due to their high electrical conductiv‐ ity as well as chemical and physical stability [1]. Planar graphene films with respect to the substrate have been synthesized using various methods including mechanical exfoliation from highly oriented pyrolytic graphite, chemical exfoliation from bulk graphite, thermal decomposition of carbon-terminated silicon carbide, and chemical vapor deposition (CVD) on metals such as nickel and copper substrates [2–6].

**Figure 1.** Schematic illustration of graphene nanowalls

On the other hand, plasma-enhanced CVD (PECVD) is among the early methods to synthe‐ size of vertically standing few layer graphenes or graphene nanowalls (GNWs) [7–12]. GNWs can be described as self-assembled, vertically standing, few-layered graphene sheet nanostructures, which are also called as carbon nanowalls, carbon nanosheets, and carbon nanoflakes. As illustrated in Fig. 1, the sheets form a self-supported network of wall struc‐ tures with thicknesses ranging from a few nanometers to a few tens of nanometers. GNWs have a high density of atomic scale graphitic edges that are potential sites for electron field emission, which might lead to the application in flat panel displays and light sources [13,14].

© 2013 Hiramatsu et al.; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2013 Hiramatsu et al.; licensee InTech. This is a paper distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

The large surface area of GNWs is useful as templates for the fabrication of other types of nanostructured materials, electrodes for energy storage devices and biosensors [15–20]. These graphene-based materials are applied in the field of electrochemistry including elec‐ trode for fuel cell and electrochemical sensors. In these applications, graphene-based materi‐ als are often decorated with metal nanoparticles and other materials.

nar geometry, RF power is inductively coupled into the process chamber with a planar-coil antenna through a quartz (fused silica) window, and plasma is generated in the chamber. Plasma density of ICP discharge is on the order of 1012 cm-3. Figure 2 shows a schematic of ICP reactor with planar geometry used for the growth of GNWs. The ICP reactor was 16 cm in diameter and 30 cm in height. A one-turn coil antenna with a diameter of 10 cm was set on a quartz window at the top of reactor. Si or SiO2-coated Si substrates were set on the mid‐

Graphene Nanowalls

237

http://dx.doi.org/10.5772/51528

**Figure 2.** Schematic of inductively coupled plasma reactor with planar geometry used for the growth of GNWs.

**Figure 3.** SEM images of the GNW films grown by ICP-CVD employing a mixture of CH4 and Ar for (a–b) 5 min, (c–d) 15

Figures 3(a) –3(f) show scanning electron microscopy (SEM) images of GNW films grown by ICP-CVD employing a mixture of CH4 and Ar for different growth periods. After the nuclea‐ tion stage of GNWs, growth of less-aligned, isolated carbon sheets with a semicircular shape standing on the substrate is confirmed as shown in Figs. 3(a) and 3(b). As the growth period increased, density of isolated nanosheets increased and those standing almost vertically on the substrate continued preferably to spread faster. Then, spreading nanosheets met one an‐

min, and (e–f) 30 min. Scale bar: 1 µm [21].

dle of the substrate holder at 10 cm below the quartz window.

In this chapter, basic properties of GNWs and their prospective applications are described. First of all, synthesis and characterization of GNWs are outlined. PECVD is becoming one of the most promising techniques for the production of carbon materials including diamond, aligned carbon nanotube films and GNWs, due to its feasibility and potentiality for largearea production with reasonable growth rates at relatively low temperatures. In the present study, GNW growth using inductively coupled plasma (ICP) enhanced CVD is featured, since the ICP CVD system has advantages of simple design and scalability to large area growth. The growth mechanism, characterization of GNWs, and several decoration techni‐ ques of GNW surface are described.

Due to the large surface area of GNWs, we can expect a variety of electrochemical applications using GNWs such as batteries, capacitors, and sensors. To these ends, GNWs are often decorat‐ ed with nanoparticles or films. In the latter half of this chapter, application of GNWs as elec‐ trode for fuel cell is described. GNWs were grown on the carbon fiber paper. Then metal organic chemical fluid deposition (MOCFD) using supercritical fluid (SCF) was applied to form platinum (Pt) nanoparticles on the surface of GNWs. Using this method, highly dis‐ persed Pt nanoparticles of approximately 2 nm in diameter were formed on the surface of GNWs grown on the carbon fiber paper. Furthermore, the application as a biosensor using GNWs is described. As another example, the electrocatalytic activity of GNWs for determin‐ ing dopamine, ascorbic acid and uric acid in phosphate buffer solution was investigated. The ability of GNWs as a platform to create graphene-based hybrid materials is demonstrated.

#### **2. Growth of graphene nanowalls and their growth mechanism**

#### **2.1. Growth of graphene nanowalls using inductively coupled plasma CVD**

Synthesis techniques for GNWs and related vertical graphene structures are similar to those used for diamond films and carbon nanotubes (CNTs). In general, a mixture of hydrocarbon and hydrogen or argon gases, typically CH4 and H2, is used as source gases for the synthesis of GNWs. Unlike the CNT growth, GNWs can be fabricated on a variety of substrates, in‐ cluding Si, SiO2, Al2O3, Ni, Ti, and stainless steel, at substrate temperatures of 500–750 ˚C without the use of catalysts [8]. To date, GNWs have been grown using various PECVD methods employing microwave plasma, inductively coupled plasma (ICP), capacitively cou‐ pled plasma (CCP) with H radical injection, very high frequency (VHF) plasma with H radi‐ cal injection, electron beam excited plasma, and DC plasma [7–12,19].

Radio frequency (RF: 13.56 MHz) ICP is one of high-density plasmas, and has been used to etch several materials including Si, SiOx, SiNx, and metal films in the LSI fabrication process. The ICP is operated at relatively low pressures below 100 mTorr (13.3 Pa). In the case of pla‐ nar geometry, RF power is inductively coupled into the process chamber with a planar-coil antenna through a quartz (fused silica) window, and plasma is generated in the chamber. Plasma density of ICP discharge is on the order of 1012 cm-3. Figure 2 shows a schematic of ICP reactor with planar geometry used for the growth of GNWs. The ICP reactor was 16 cm in diameter and 30 cm in height. A one-turn coil antenna with a diameter of 10 cm was set on a quartz window at the top of reactor. Si or SiO2-coated Si substrates were set on the mid‐ dle of the substrate holder at 10 cm below the quartz window.

The large surface area of GNWs is useful as templates for the fabrication of other types of nanostructured materials, electrodes for energy storage devices and biosensors [15–20]. These graphene-based materials are applied in the field of electrochemistry including elec‐ trode for fuel cell and electrochemical sensors. In these applications, graphene-based materi‐

In this chapter, basic properties of GNWs and their prospective applications are described. First of all, synthesis and characterization of GNWs are outlined. PECVD is becoming one of the most promising techniques for the production of carbon materials including diamond, aligned carbon nanotube films and GNWs, due to its feasibility and potentiality for largearea production with reasonable growth rates at relatively low temperatures. In the present study, GNW growth using inductively coupled plasma (ICP) enhanced CVD is featured, since the ICP CVD system has advantages of simple design and scalability to large area growth. The growth mechanism, characterization of GNWs, and several decoration techni‐

Due to the large surface area of GNWs, we can expect a variety of electrochemical applications using GNWs such as batteries, capacitors, and sensors. To these ends, GNWs are often decorat‐ ed with nanoparticles or films. In the latter half of this chapter, application of GNWs as elec‐ trode for fuel cell is described. GNWs were grown on the carbon fiber paper. Then metal organic chemical fluid deposition (MOCFD) using supercritical fluid (SCF) was applied to form platinum (Pt) nanoparticles on the surface of GNWs. Using this method, highly dis‐ persed Pt nanoparticles of approximately 2 nm in diameter were formed on the surface of GNWs grown on the carbon fiber paper. Furthermore, the application as a biosensor using GNWs is described. As another example, the electrocatalytic activity of GNWs for determin‐ ing dopamine, ascorbic acid and uric acid in phosphate buffer solution was investigated. The ability of GNWs as a platform to create graphene-based hybrid materials is demonstrated.

**2. Growth of graphene nanowalls and their growth mechanism**

**2.1. Growth of graphene nanowalls using inductively coupled plasma CVD**

cal injection, electron beam excited plasma, and DC plasma [7–12,19].

Synthesis techniques for GNWs and related vertical graphene structures are similar to those used for diamond films and carbon nanotubes (CNTs). In general, a mixture of hydrocarbon and hydrogen or argon gases, typically CH4 and H2, is used as source gases for the synthesis of GNWs. Unlike the CNT growth, GNWs can be fabricated on a variety of substrates, in‐ cluding Si, SiO2, Al2O3, Ni, Ti, and stainless steel, at substrate temperatures of 500–750 ˚C without the use of catalysts [8]. To date, GNWs have been grown using various PECVD methods employing microwave plasma, inductively coupled plasma (ICP), capacitively cou‐ pled plasma (CCP) with H radical injection, very high frequency (VHF) plasma with H radi‐

Radio frequency (RF: 13.56 MHz) ICP is one of high-density plasmas, and has been used to etch several materials including Si, SiOx, SiNx, and metal films in the LSI fabrication process. The ICP is operated at relatively low pressures below 100 mTorr (13.3 Pa). In the case of pla‐

als are often decorated with metal nanoparticles and other materials.

ques of GNW surface are described.

236 New Progress on Graphene Research

**Figure 2.** Schematic of inductively coupled plasma reactor with planar geometry used for the growth of GNWs.

**Figure 3.** SEM images of the GNW films grown by ICP-CVD employing a mixture of CH4 and Ar for (a–b) 5 min, (c–d) 15 min, and (e–f) 30 min. Scale bar: 1 µm [21].

Figures 3(a) –3(f) show scanning electron microscopy (SEM) images of GNW films grown by ICP-CVD employing a mixture of CH4 and Ar for different growth periods. After the nuclea‐ tion stage of GNWs, growth of less-aligned, isolated carbon sheets with a semicircular shape standing on the substrate is confirmed as shown in Figs. 3(a) and 3(b). As the growth period increased, density of isolated nanosheets increased and those standing almost vertically on the substrate continued preferably to spread faster. Then, spreading nanosheets met one an‐ other; eventually resulting in the formation of linked nanowalls as shown in Figs. 3(c) and 3(d). During the early growth stage after the nucleation up to the steady-state growth, the growth for the inclined smaller nanowalls was terminated, while the vertical nanowalls preferentially continued to grow. Therefore, with the increase of growth period in the early stage, the spacing between nanowalls at their top increased gradually, and then became al‐ most saturated, resulting in the formation of two-dimensional carbon sheets standing verti‐ cally on the substrate with high aspect ratio. As shown in Figs. 3(e) and 3(f), with the further increase of growth period during the steady-state growth, the height of vertical aligned GNWs increased, while the thickness of nanowalls and the spacing between nanowalls be‐ came almost saturated with keeping the morphology of GNWs.

electron microscopy (TEM), in-plane synchrotron X-ray diffraction, and Raman spectrosco‐

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239

http://dx.doi.org/10.5772/51528

Figure 5(a) shows a typical SEM image of GNW film, indicating the vertical growth of the two-dimensional carbon sheets with honeycomb structure on the substrate. Actually, the morphology and structure of GNW film depend on the source gases, pressure, process tem‐ perature, as well as the type of plasma used for the growth of GNWs. In addition to the ver‐ tically standing maze-like structure, isolated very thin nanosheets, less aligned petal-like,

Figure 5(b) shows a low-magnification TEM image of a piece of typical GNW of a microme‐ ter-high planar nanosheet structure with a relatively smooth surface. The GNW is composed of nano-domains of a few tens of nanometers in size. The high-resolution TEM image of the GNW shown in Fig. 5(c) reveals the graphene layers, which indicates the graphitized struc‐ ture of the GNWs. The spacing between neighboring graphene layers was measured as ap‐

**Figure 5.** a) Typical SEM image of GNW film, (b) low-resolution TEM image of GNW on a microgrid, and (c) high-reso‐

**Figure 6.** SR X-ray diffraction pattern of GNW film measured at beam line BL13XU of SPring-8 [22].

lution TEM image of GNW, showing graphene layers at the fold of GNWs.

highly branched type, and a kind of porous film have been fabricated so far.

py using a 514.5 nm line of argon laser.

proximately 0.34 nm.

Growth rate curve for the GNWs fabricated by ICP-CVD employing CH4/Ar system was ob‐ tained by measuring the height of the nanowalls for differing period of growth (0–120 min). Figure 4 shows the average height of GNW films as a function of growth period. As shown in Fig. 4, the height of GNWs almost linearly increased with the increase of growth period in the range from 10 to 120 min, while it took approximately 5 min for the nucleation. The growth rate of GNWs in the steady-state condition was constant at approximately 60 nm/min.

**Figure 4.** Wall height of GNWs as a function of growth period. The growth rate data were obtained from the samples grown for different period on Si substrates by ICP-CVD employing CH4/Ar system [21].

#### **2.2. Characterization of graphene nanowalls**

As was illustrated in Fig. 1, GNWs can be described as graphite sheet nanostructures with edges that are composed of stacks of graphene sheets standing almost vertically on the sub‐ strate. The sheets form a self-supported network of wall structures with thicknesses in the range from a few nanometers to a few tens of nanometers, and with a high aspect ratio. In this section, typical GNWs grown using PECVD are characterized using SEM, transmission electron microscopy (TEM), in-plane synchrotron X-ray diffraction, and Raman spectrosco‐ py using a 514.5 nm line of argon laser.

other; eventually resulting in the formation of linked nanowalls as shown in Figs. 3(c) and 3(d). During the early growth stage after the nucleation up to the steady-state growth, the growth for the inclined smaller nanowalls was terminated, while the vertical nanowalls preferentially continued to grow. Therefore, with the increase of growth period in the early stage, the spacing between nanowalls at their top increased gradually, and then became al‐ most saturated, resulting in the formation of two-dimensional carbon sheets standing verti‐ cally on the substrate with high aspect ratio. As shown in Figs. 3(e) and 3(f), with the further increase of growth period during the steady-state growth, the height of vertical aligned GNWs increased, while the thickness of nanowalls and the spacing between nanowalls be‐

Growth rate curve for the GNWs fabricated by ICP-CVD employing CH4/Ar system was ob‐ tained by measuring the height of the nanowalls for differing period of growth (0–120 min). Figure 4 shows the average height of GNW films as a function of growth period. As shown in Fig. 4, the height of GNWs almost linearly increased with the increase of growth period in the range from 10 to 120 min, while it took approximately 5 min for the nucleation. The growth rate

**Figure 4.** Wall height of GNWs as a function of growth period. The growth rate data were obtained from the samples

As was illustrated in Fig. 1, GNWs can be described as graphite sheet nanostructures with edges that are composed of stacks of graphene sheets standing almost vertically on the sub‐ strate. The sheets form a self-supported network of wall structures with thicknesses in the range from a few nanometers to a few tens of nanometers, and with a high aspect ratio. In this section, typical GNWs grown using PECVD are characterized using SEM, transmission

grown for different period on Si substrates by ICP-CVD employing CH4/Ar system [21].

**2.2. Characterization of graphene nanowalls**

of GNWs in the steady-state condition was constant at approximately 60 nm/min.

came almost saturated with keeping the morphology of GNWs.

238 New Progress on Graphene Research

Figure 5(a) shows a typical SEM image of GNW film, indicating the vertical growth of the two-dimensional carbon sheets with honeycomb structure on the substrate. Actually, the morphology and structure of GNW film depend on the source gases, pressure, process tem‐ perature, as well as the type of plasma used for the growth of GNWs. In addition to the ver‐ tically standing maze-like structure, isolated very thin nanosheets, less aligned petal-like, highly branched type, and a kind of porous film have been fabricated so far.

Figure 5(b) shows a low-magnification TEM image of a piece of typical GNW of a microme‐ ter-high planar nanosheet structure with a relatively smooth surface. The GNW is composed of nano-domains of a few tens of nanometers in size. The high-resolution TEM image of the GNW shown in Fig. 5(c) reveals the graphene layers, which indicates the graphitized struc‐ ture of the GNWs. The spacing between neighboring graphene layers was measured as ap‐ proximately 0.34 nm.

**Figure 5.** a) Typical SEM image of GNW film, (b) low-resolution TEM image of GNW on a microgrid, and (c) high-reso‐ lution TEM image of GNW, showing graphene layers at the fold of GNWs.

**Figure 6.** SR X-ray diffraction pattern of GNW film measured at beam line BL13XU of SPring-8 [22].

The crystallinity of GNWs was analyzed using synchrotron X-ray surface diffraction at graz‐ ing incidence and exit at the beamline BL13XU of SPring-8 [22]. The in-plane diffraction technique measures diffracted beams, which are scattered nearly parallel to the sample sur‐ face and hence measures lattice planes that are perpendicular to the sample surface. The Xray beam was incident on the GNW film sample at grazing angle of 0.3˚ relative to the substrate surface. Figure 6 shows the grazing incidence in-plane X-ray diffraction pattern of GNW film sample. An intense 002 Bragg peak, the plane of which is normal to the substrate, is at 2θ=16.9˚and there are also weak 100/101, 004, and 110 Bragg peaks. The interlayer spac‐ ing d002 was determined from the 002 peak by applying Bragg's law with a wavelength of 0.1003 nm. It was found to be 0.342 nm for all samples, which is slightly larger than that of bulk graphite (0.335 nm).

covered with amorphous carbon layer, which is later confirmed using Raman spectroscopy. In 1 min, the onset of nano-graphene growth was observed in the top view SEM image of depos‐ its formed for 1 min as shown Fig. 8(b). Figures 8(c) and 8(d) show tilted SEM image and crosssectional TEM image of deposits, respectively, formed for 2 min. For the cross-sectional TEM observation, sample surface was coated with the epoxy resin, the deposits embedded in the epoxy resin were peeled off from the substrate, and then the substrate side (interface side) of resin embedding the deposits was coated again with the epoxy resin. In Fig. 8(d), the red line corresponds to the interface to the substrate surface and the red arrow indicates the growth di‐ rection. In 2 min, the formation of isolated graphene sheets was observed on the amorphous carbon layer as shown in Fig. 8(c). The thickness of amorphous carbon layer was estimated to

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**Figure 8.** SEM and TEM images of deposits formed at the nucleation stage. (a) Top view SEM image of substrate sur‐ face after 30 sec. (b) Top view SEM image of deposits formed for 1 min, indicating the commencement of nano-gra‐ phene growth. (c) Tilted SEM image of deposits formed for 2 min. (d) Cross-sectional TEM image of deposits formed for 2 min. Deposits were detached from the Si substrate. Red line in (d) corresponds to the interface to the substrate

Figure 9 shows Raman spectra of deposits formed on SiO2-coated Si substrate in the nuclea‐ tion stage. The peak around 950 cm-1 originates from the Si wafer [27]. The broad peak at 1340 cm-1 in the Raman spectrum at 30 sec indicates that the deposits are amorphous carbon or diamond-like carbon. Namely, during the nucleation period, graphene component was

be 30 nm from the TEM observation shown in Fig. 8(d).

surface and the red arrow indicates the growth direction [21].

scarcely contained in the underlying interface layer shown in Fig. 9.

Raman spectrum for GNWs grown on Si substrate is shown in Fig. 7. Typical Raman spec‐ trum for the GNWs has G band peak at 1590 cm-1 indicating the formation of a graphitized structure and D band peak at 1350 cm-1 corresponding to the disorder-induced phonon mode. The peak intensity of D band is twice as high as that of G band. The G band peak is accompanied by a shoulder peak at 1620 cm-1 (D' band), which is associated with finite-size graphite crystals and graphene edges [23,24]. The 2D band peak at 2690 cm-1 is used to con‐ firm the presence of graphene and it originates from a double resonance process that links phonons to the electronic band structure [25,26]. The strong and sharp D band peak and D' band peak suggest a more nanocrystalline structure and the presence of graphene edges and defects, which are prevalent features of GNWs [8,10,12].

**Figure 7.** Typical Raman spectrum of GNWs measured using the 514.5 nm line of an Ar laser [21]

#### **2.3. Nucleation of vertical nanographenes**

From the temporal dependence of nanowall height shown in Fig. 3, nucleation of GNWs is con‐ sidered to occur before the commencement of steady-state growth. Significant interest exists in clarifying the nucleation mechanism of GNWs at the very early stage. Figure 8(a) shows a top view SEM image of SiO2 substrate surface after 30 sec growth. First the surface of substrate was covered with amorphous carbon layer, which is later confirmed using Raman spectroscopy. In 1 min, the onset of nano-graphene growth was observed in the top view SEM image of depos‐ its formed for 1 min as shown Fig. 8(b). Figures 8(c) and 8(d) show tilted SEM image and crosssectional TEM image of deposits, respectively, formed for 2 min. For the cross-sectional TEM observation, sample surface was coated with the epoxy resin, the deposits embedded in the epoxy resin were peeled off from the substrate, and then the substrate side (interface side) of resin embedding the deposits was coated again with the epoxy resin. In Fig. 8(d), the red line corresponds to the interface to the substrate surface and the red arrow indicates the growth di‐ rection. In 2 min, the formation of isolated graphene sheets was observed on the amorphous carbon layer as shown in Fig. 8(c). The thickness of amorphous carbon layer was estimated to be 30 nm from the TEM observation shown in Fig. 8(d).

The crystallinity of GNWs was analyzed using synchrotron X-ray surface diffraction at graz‐ ing incidence and exit at the beamline BL13XU of SPring-8 [22]. The in-plane diffraction technique measures diffracted beams, which are scattered nearly parallel to the sample sur‐ face and hence measures lattice planes that are perpendicular to the sample surface. The Xray beam was incident on the GNW film sample at grazing angle of 0.3˚ relative to the substrate surface. Figure 6 shows the grazing incidence in-plane X-ray diffraction pattern of GNW film sample. An intense 002 Bragg peak, the plane of which is normal to the substrate, is at 2θ=16.9˚and there are also weak 100/101, 004, and 110 Bragg peaks. The interlayer spac‐ ing d002 was determined from the 002 peak by applying Bragg's law with a wavelength of 0.1003 nm. It was found to be 0.342 nm for all samples, which is slightly larger than that of

Raman spectrum for GNWs grown on Si substrate is shown in Fig. 7. Typical Raman spec‐ trum for the GNWs has G band peak at 1590 cm-1 indicating the formation of a graphitized structure and D band peak at 1350 cm-1 corresponding to the disorder-induced phonon mode. The peak intensity of D band is twice as high as that of G band. The G band peak is accompanied by a shoulder peak at 1620 cm-1 (D' band), which is associated with finite-size graphite crystals and graphene edges [23,24]. The 2D band peak at 2690 cm-1 is used to con‐ firm the presence of graphene and it originates from a double resonance process that links phonons to the electronic band structure [25,26]. The strong and sharp D band peak and D' band peak suggest a more nanocrystalline structure and the presence of graphene edges and

bulk graphite (0.335 nm).

240 New Progress on Graphene Research

defects, which are prevalent features of GNWs [8,10,12].

**2.3. Nucleation of vertical nanographenes**

**Figure 7.** Typical Raman spectrum of GNWs measured using the 514.5 nm line of an Ar laser [21]

From the temporal dependence of nanowall height shown in Fig. 3, nucleation of GNWs is con‐ sidered to occur before the commencement of steady-state growth. Significant interest exists in clarifying the nucleation mechanism of GNWs at the very early stage. Figure 8(a) shows a top view SEM image of SiO2 substrate surface after 30 sec growth. First the surface of substrate was

**Figure 8.** SEM and TEM images of deposits formed at the nucleation stage. (a) Top view SEM image of substrate sur‐ face after 30 sec. (b) Top view SEM image of deposits formed for 1 min, indicating the commencement of nano-gra‐ phene growth. (c) Tilted SEM image of deposits formed for 2 min. (d) Cross-sectional TEM image of deposits formed for 2 min. Deposits were detached from the Si substrate. Red line in (d) corresponds to the interface to the substrate surface and the red arrow indicates the growth direction [21].

Figure 9 shows Raman spectra of deposits formed on SiO2-coated Si substrate in the nuclea‐ tion stage. The peak around 950 cm-1 originates from the Si wafer [27]. The broad peak at 1340 cm-1 in the Raman spectrum at 30 sec indicates that the deposits are amorphous carbon or diamond-like carbon. Namely, during the nucleation period, graphene component was scarcely contained in the underlying interface layer shown in Fig. 9.

So far, several papers have been published on the observation of GNW growth in the early growth stage and the nucleation mechanism for the formation of vertical layered-graphenes using various CVD methods [8,28-35]. The things in common in previous observations are that there is an induction period of 1–5 min before the onset of vertical nano-graphene growth and an interface layer exists between vertical nano-graphenes and the surface of Si and SiO2 substrates. Zhu, et al. [31,32] reported the presence of graphenes parallel to the substrate surface before the onset of vertical nanosheet growth, although neither Raman spectrum nor TEM image of base layer was attached. In their model, these few-layer graphe‐ nes would grow parallel to the substrate surface until a sufficient level of force develops at the grain boundaries to curl the leading edge of the top layers upward. The vertical orienta‐ tion of these sheets would result from the interaction of the plasma electric field with their anisotropic polarizability [32]. In the case of GNW growth in the present work, on the other hand, the interface layer under the CNWs was an amorphous or diamond-like carbon, which is similar to the cases using radical injection PECVD, multi-beam CVD, and DC PECVD [8,28-30,33,34]. Moreover, ion bombardment on the surface will play an important role for nucleation by creating active sites for neutral radical bonding [30]. The existence of amorphous or diamond-like carbon interface layer will enable us to grow GNWs and similar structures on a variety of substrates without catalyst.

dangling bonds, followed by two-dimensional growth and subsequent formation of nano‐ graphene sheets with random orientation. (5) Among the nucleated graphene sheets with random orientations, those standing almost vertically on the substrate continued preferably to grow up faster to vertically standing nanosheets owing to the difference in the growth rates along the strongly bonded planes of graphene sheets expanding and in the weakly bonded stacking direction. Reactive carbon species arriving at the edge of the graphene lay‐ er are easily bonded to the edge, and eventually the graphene layer would expand prefera‐ bly along the direction of radical diffusion, perpendicular to the electrode plane. On the other hand, low-lying inclined graphene sheets were shadowed by the high-grown vertical graphene sheets. As a result, the amounts of reactive carbon species arriving at the low-ly‐ ing inclined graphene sheets decreased, resulting in the termination of growth for the in‐ clined smaller nanowalls, while the vertical nanowalls preferentially continued to grow. As growth period increased, spreading vertical nanowalls met one another, eventually result‐ ing in the formation of linked nanowalls similar to a maze. With further increase of growth period, the spacing between nanowalls at their top increased gradually, and then became al‐ most saturated, resulting in the formation of two-dimensional graphene sheets standing ver‐ tically on the substrate with high aspect ratio. In the steady-state growth condition, the

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height of nanowalls increased almost linearly with keeping their morphology.

Due to the large surface area (high surface-to-volume ratio) of GNWs, we can expect a va‐ riety of electrochemical applications using GNWs such as batteries, capacitors, and gas sen‐

semiconductors, and insulators, by using several techniques including vacuum evapora‐ tion, sputtering, CVD, and plating. Figure 11 shows schematic illustrations of decorated

sors. To these ends, GNWs are decorated with nanoparticles or films of metals,

**Figure 10.** Illustration of the initial growth model of GNWs

**3.1. Decoration of graphene nanowall surface**

**3. Functionalization of graphene nanowalls**

**Figure 9.** Raman spectra of deposits formed in the nucleation stage on SiO2-coated Si substrate. Peak around 950 cm-1 originates from the Si wafer [8].

#### **2.4. Growth mechanism of graphene nanowalls**

Figure 10 depicts the initial growth process of GNWs. A model for the initial growth mecha‐ nism is as follows. (1) In the beginning, hydrocarbon radicals such as CH3 are adsorbed on the substrate, forming a very thin amorphous carbon layer. Ion irradiation induces the for‐ mation of dangling bonds on the growing surface, resulting in the formation of nucleation sites. (2) Adsorbed carbon species are migrating on the surface and condensed to form nano‐ islands with dangling bonds. (3) Ion irradiation would also enhance the adsorption of CHx radicals on the surface. (4) Small, disordered graphene nanosheets are nucleated at these dangling bonds, followed by two-dimensional growth and subsequent formation of nano‐ graphene sheets with random orientation. (5) Among the nucleated graphene sheets with random orientations, those standing almost vertically on the substrate continued preferably to grow up faster to vertically standing nanosheets owing to the difference in the growth rates along the strongly bonded planes of graphene sheets expanding and in the weakly bonded stacking direction. Reactive carbon species arriving at the edge of the graphene lay‐ er are easily bonded to the edge, and eventually the graphene layer would expand prefera‐ bly along the direction of radical diffusion, perpendicular to the electrode plane. On the other hand, low-lying inclined graphene sheets were shadowed by the high-grown vertical graphene sheets. As a result, the amounts of reactive carbon species arriving at the low-ly‐ ing inclined graphene sheets decreased, resulting in the termination of growth for the in‐ clined smaller nanowalls, while the vertical nanowalls preferentially continued to grow. As growth period increased, spreading vertical nanowalls met one another, eventually result‐ ing in the formation of linked nanowalls similar to a maze. With further increase of growth period, the spacing between nanowalls at their top increased gradually, and then became al‐ most saturated, resulting in the formation of two-dimensional graphene sheets standing ver‐ tically on the substrate with high aspect ratio. In the steady-state growth condition, the height of nanowalls increased almost linearly with keeping their morphology.

**Figure 10.** Illustration of the initial growth model of GNWs

So far, several papers have been published on the observation of GNW growth in the early growth stage and the nucleation mechanism for the formation of vertical layered-graphenes using various CVD methods [8,28-35]. The things in common in previous observations are that there is an induction period of 1–5 min before the onset of vertical nano-graphene growth and an interface layer exists between vertical nano-graphenes and the surface of Si and SiO2 substrates. Zhu, et al. [31,32] reported the presence of graphenes parallel to the substrate surface before the onset of vertical nanosheet growth, although neither Raman spectrum nor TEM image of base layer was attached. In their model, these few-layer graphe‐ nes would grow parallel to the substrate surface until a sufficient level of force develops at the grain boundaries to curl the leading edge of the top layers upward. The vertical orienta‐ tion of these sheets would result from the interaction of the plasma electric field with their anisotropic polarizability [32]. In the case of GNW growth in the present work, on the other hand, the interface layer under the CNWs was an amorphous or diamond-like carbon, which is similar to the cases using radical injection PECVD, multi-beam CVD, and DC PECVD [8,28-30,33,34]. Moreover, ion bombardment on the surface will play an important role for nucleation by creating active sites for neutral radical bonding [30]. The existence of amorphous or diamond-like carbon interface layer will enable us to grow GNWs and similar

**Figure 9.** Raman spectra of deposits formed in the nucleation stage on SiO2-coated Si substrate. Peak around 950 cm-1

Figure 10 depicts the initial growth process of GNWs. A model for the initial growth mecha‐ nism is as follows. (1) In the beginning, hydrocarbon radicals such as CH3 are adsorbed on the substrate, forming a very thin amorphous carbon layer. Ion irradiation induces the for‐ mation of dangling bonds on the growing surface, resulting in the formation of nucleation sites. (2) Adsorbed carbon species are migrating on the surface and condensed to form nano‐ islands with dangling bonds. (3) Ion irradiation would also enhance the adsorption of CHx radicals on the surface. (4) Small, disordered graphene nanosheets are nucleated at these

structures on a variety of substrates without catalyst.

**2.4. Growth mechanism of graphene nanowalls**

originates from the Si wafer [8].

242 New Progress on Graphene Research

## **3. Functionalization of graphene nanowalls**

#### **3.1. Decoration of graphene nanowall surface**

Due to the large surface area (high surface-to-volume ratio) of GNWs, we can expect a va‐ riety of electrochemical applications using GNWs such as batteries, capacitors, and gas sen‐ sors. To these ends, GNWs are decorated with nanoparticles or films of metals, semiconductors, and insulators, by using several techniques including vacuum evapora‐ tion, sputtering, CVD, and plating. Figure 11 shows schematic illustrations of decorated

GNWs with different morphologies. GNWs are used as the templates to fabricate other types of nanostructures. The morphology of GNWs decorated with nanoparticles or film de‐ pends on the deposition methods of materials. Conformal deposition (Fig. 11(a)) and gap filling (Fig. 11(b)) will be achieved using metal-organic chemical vapor deposition (MOCVD), sputtering, atomic layer deposition, and plating in liquid phase. In Fig. 11(c), thin film or aggregation of nanoparticles would be formed on the top edges of GNWs. Diamond sur‐ face is modified with several types of surface termination, e.g. C–NH2, C–OH, and C– COOH, and DNA and proteins were immobilized on the surface of diamond and nanodia‐ mond films for bio-sensing application [35–38]. As is the case with the diamond surface, the edges of GNWs can also be modified with similar surface termination, and covalent im‐ mobilization of DNA and proteins on the GNWs will be realized. Previously, Wu et al. used GNWs as templates to fabricate large surface area materials, including conformal dep‐ osition of Au and Cu by electron beam evaporation; conformal deposition of ZnO, TiO2, SiOx, SiNx, and AlOx by atomic layer deposition; conformal deposition of Ni, NiFe and CoN‐ iFe nanoparticles by electrochemical deposition; gap filling with dispersed Fe nanoparti‐ cles by the immersion of GNWs into a mixed solution of Fe particles and isopropanol in an ultrasonic bath; Se deposition on the top of edges of GNWs by electrochemical deposi‐ tion [39–41]. In Fig. 11(d), on the other hand, metal nanoparticles are dispersed on the sur‐ face of GNWs, which is a kind of nanocomposite. This morphology can be moderately achieved by plating, sputtering, and laser ablation.

Because of the unique structure of GNWs with high surface-to-volume ratio, GNWs can be potentially used as catalyst support materials for the electrodes of fuel cells. In this applica‐ tion, it is required to support platinum (Pt) nanoparticles as catalysts on the GNW surface. It is well known that the specific activity of catalysts is strongly related to their size, disper‐ sion, and compatibility with supporting materials. Highly dispersed catalyst nanoparticles with small size and narrow size distribution supported on the surface of carbon nanostruc‐ tures are ideal for high electrocatalyst activity due to their large surface-to-volume ratio. To support metal nanoparticles on the surface of carbon nanostructures, metal compounds in the form of liquids are generally employed. A few papers have been published on the prepa‐ ration of Pt nanoparticles on CNT surfaces by the reduction of Pt salt precursors such as H2PtCl6 in solution [42,43]. However, it is difficult to treat the entire surface of GNWs with a metal compound in a liquid phase, because of the high surface tension of GNWs due to their high aspect ratio with narrow interspaces. On the other hand, in gas phase deposition such as sputtering and CVD, metal nanoparticles are deposited only around the tops of GNWs and tend to easily clump together, resulting in the formation of larger particles or films on

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**Figure 12.** Phase diagram of substance (SCF supercritical fluid, Pc critical pressure, Tc critical temperature) and proper‐

As an alternative approach to support metal nanoparticles on the surfaces of dense, aligned CNTs and GNWs with narrow interspaces, we have demonstrated a method employing metal-organic chemical fluid deposition (MOCFD), where supercritical carbon dioxide (sc-CO2) is used as a solvent of metal-organic compounds [45,46]. The phase diagram of sub‐ stance is shown in Fig. 12, together with the properties of liquid phase, gas phase, and supercritical fluid (SCF). Generally, materials can exist in three phases depending on the temperature and pressure, namely, solid, liquid, and gas. The SCF possesses attractive prop‐ erties of both the gas and the liquid phases. Rapid diffusion and permeation are realized by its gas-like diffusivity and viscosity, while its liquid-like density enables dissolution of a wide range of materials. To produce an SCF phase, the temperature and pressure of the ma‐ terial are required to exceed the critical point. The critical point of sc-CO2 exists at 7.38 MPa

the top of carbon nanostructures [44].

ties of liquid phase, gas phase, and supercritical fluid

**3.2. Supercritical fluids**

**Figure 11.** Schematic illustrations of decorated GNWs with different morphologies: (a) conformal deposition, (b) gap filling, (c) deposition on the top edges of GNWs, and (d) dispersed nanoparticle deposition on the surface of GNWs

Because of the unique structure of GNWs with high surface-to-volume ratio, GNWs can be potentially used as catalyst support materials for the electrodes of fuel cells. In this applica‐ tion, it is required to support platinum (Pt) nanoparticles as catalysts on the GNW surface. It is well known that the specific activity of catalysts is strongly related to their size, disper‐ sion, and compatibility with supporting materials. Highly dispersed catalyst nanoparticles with small size and narrow size distribution supported on the surface of carbon nanostruc‐ tures are ideal for high electrocatalyst activity due to their large surface-to-volume ratio. To support metal nanoparticles on the surface of carbon nanostructures, metal compounds in the form of liquids are generally employed. A few papers have been published on the prepa‐ ration of Pt nanoparticles on CNT surfaces by the reduction of Pt salt precursors such as H2PtCl6 in solution [42,43]. However, it is difficult to treat the entire surface of GNWs with a metal compound in a liquid phase, because of the high surface tension of GNWs due to their high aspect ratio with narrow interspaces. On the other hand, in gas phase deposition such as sputtering and CVD, metal nanoparticles are deposited only around the tops of GNWs and tend to easily clump together, resulting in the formation of larger particles or films on the top of carbon nanostructures [44].

**Figure 12.** Phase diagram of substance (SCF supercritical fluid, Pc critical pressure, Tc critical temperature) and proper‐ ties of liquid phase, gas phase, and supercritical fluid

#### **3.2. Supercritical fluids**

GNWs with different morphologies. GNWs are used as the templates to fabricate other types of nanostructures. The morphology of GNWs decorated with nanoparticles or film de‐ pends on the deposition methods of materials. Conformal deposition (Fig. 11(a)) and gap filling (Fig. 11(b)) will be achieved using metal-organic chemical vapor deposition (MOCVD), sputtering, atomic layer deposition, and plating in liquid phase. In Fig. 11(c), thin film or aggregation of nanoparticles would be formed on the top edges of GNWs. Diamond sur‐ face is modified with several types of surface termination, e.g. C–NH2, C–OH, and C– COOH, and DNA and proteins were immobilized on the surface of diamond and nanodia‐ mond films for bio-sensing application [35–38]. As is the case with the diamond surface, the edges of GNWs can also be modified with similar surface termination, and covalent im‐ mobilization of DNA and proteins on the GNWs will be realized. Previously, Wu et al. used GNWs as templates to fabricate large surface area materials, including conformal dep‐ osition of Au and Cu by electron beam evaporation; conformal deposition of ZnO, TiO2, SiOx, SiNx, and AlOx by atomic layer deposition; conformal deposition of Ni, NiFe and CoN‐ iFe nanoparticles by electrochemical deposition; gap filling with dispersed Fe nanoparti‐ cles by the immersion of GNWs into a mixed solution of Fe particles and isopropanol in an ultrasonic bath; Se deposition on the top of edges of GNWs by electrochemical deposi‐ tion [39–41]. In Fig. 11(d), on the other hand, metal nanoparticles are dispersed on the sur‐ face of GNWs, which is a kind of nanocomposite. This morphology can be moderately

**Figure 11.** Schematic illustrations of decorated GNWs with different morphologies: (a) conformal deposition, (b) gap filling, (c) deposition on the top edges of GNWs, and (d) dispersed nanoparticle deposition on the surface of GNWs

achieved by plating, sputtering, and laser ablation.

244 New Progress on Graphene Research

As an alternative approach to support metal nanoparticles on the surfaces of dense, aligned CNTs and GNWs with narrow interspaces, we have demonstrated a method employing metal-organic chemical fluid deposition (MOCFD), where supercritical carbon dioxide (sc-CO2) is used as a solvent of metal-organic compounds [45,46]. The phase diagram of sub‐ stance is shown in Fig. 12, together with the properties of liquid phase, gas phase, and supercritical fluid (SCF). Generally, materials can exist in three phases depending on the temperature and pressure, namely, solid, liquid, and gas. The SCF possesses attractive prop‐ erties of both the gas and the liquid phases. Rapid diffusion and permeation are realized by its gas-like diffusivity and viscosity, while its liquid-like density enables dissolution of a wide range of materials. To produce an SCF phase, the temperature and pressure of the ma‐ terial are required to exceed the critical point. The critical point of sc-CO2 exists at 7.38 MPa (72.8 atm) and 31.1 ˚C. Among SCFs, sc-CO<sup>2</sup> is particularly attractive since it is environmen‐ tally friendly and safe due to its low toxicity, low reactivity and nonflammability.

hydrogen radical injection, which comprises a parallel-plate very high frequency (VHF: 100 MHz), capacitively coupled plasma region, and a hydrogen radical injection source that em‐ ploys a surface-wave-excited microwave (2.45 GHz) H2 plasma [11]. As can be seen from these TEM images, the spatial density of the Pt nanoparticles (particle numbers/area) sup‐ ported on the GNW surface strongly depended on the sample temperature during the SCF-MOCFD, while the average size of the Pt nanoparticles increased from 1.5 to 3 nm with an

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**Figure 13.** Schematic of the supercritical fluid, metal-organic chemical fluid deposition (SCF-MOCFD) system [53]

*Ex situ* X-ray photoelectron spectroscopy (XPS) analysis was conducted to gain an insight in‐ to the state of the platinum for the supported Pt surface fabricated by the SCF-MOCFD. Fig‐ ure 14(d) shows an XPS spectrum of the Pt(*4f*) region of the Pt-supported GNW film after the SCF-MOCFD at a sample temperature of 150˚C. The presence of two prominent sets of Pt (*4f*) peaks, corresponding to the *4f 7/2* and *4f 5/2* orbital states, is further confirmation of platinum being present on the GNW surface. The peak regions in Fig. 14(d) can be fitted with two sets of peaks at 71.4 eV (*4f 7/2*) and 74.6 eV (*4f 5/2*) [54]. These correspond to platinum in the metallic state, indicating that only pure Pt exists without being oxidized on the sur‐

The Pt/C ratio of the GNW film surface was obtained from the ratio of the intensities of the XPS C(*1s*) and Pt(*4f*) peaks. Figure 14(e) shows the relative Pt/C ratio of the surface of the GNW film as a function of the temperature of the GNW sample during the SCF-MOCFD process. As the sample temperature during the SCF-MOCFD increased up to 120�?˚C, the relative Pt/C ratio of the surface of the GNW film increased gradually. By further increasing

the sample temperature above 120˚C, the relative Pt/C ratio increased rapidly.

**3.5. Mechanism of platinum nanoparticle formation by metal-organic chemical fluid**

In the case of Pt deposition by the conventional process including impregnation, depressuri‐ zation, and reduction, PtMe2(cod) was dissolved in sc-CO2 and impregnated into the sup‐

increase of the sample temperature from 120 to 170˚C.

face of the GNWs after the SCF-MOCFD.

**deposition using supercritical carbon dioxide**

In the case of Pt deposition, the SCF using sc-CO<sup>2</sup> was first applied to the preparation of pol‐ ymer-supported Pt nanoparticles using dimethyl (1,5-cyclooctadiene) platinum(II), (PtMe2(cod)) as a precursor [47]. Erkey's group has demonstrated the preparation of Pt nanoparticles on a wide range of materials, including carbon aerogel, carbon black, silica aerogel, alumina, and Nafion [47–52]. In their method, PtMe2(cod) was dissolved in sc-CO<sup>2</sup> and impregnated into the supporting materials, and after depressurization the impregnated PtMe2(cod) molecules were then reduced to metallic Pt nanoparticles by heat treatment or by chemical reduction with hydrogen, resulting in the formation of uniformly dispersed nanoparticles with narrow size distributions. However, it took almost 10 hours to complete this process. In our case, in contrast, the supporting carbon nanostructures such as CNTs and GNWs were selectively heated in the sc-CO2 with Pt precursors during the process. Therefore, at the heated surface of the carbon nanostructures during *in situ* thermal reduc‐ tion under the SCF environment, decomposition of the adsorbed precursor molecules and growth of the particles would occur without reduction process.

#### **3.3. Experimental procedure of metal-organic chemical fluid deposition using supercritical carbon dioxide**

Figure 13 shows the SCF-MOCFD system employing sc-CO2 used for the deposition of Pt nanoparticles on the surface of GNWs. The MOCFD system is composed of two high-pres‐ sure stainless steel vessels with a compressor, heating units, and a reservoir for the metalorganic compound. The preliminary vessel contains a screw agitator. The temperature and pressure in each vessel can be set independently, so that two different supercritical condi‐ tions employing CO2 can be produced in these two vessels. As the precursor, (methylcyclo‐ pentadienyl) trimethyl platinum ((CH3C5H4)Pt(CH3)3: MeCpPtMe3) dissolved in hexane was used. In the preliminary vessel, the precursor was stirred with the sc-CO<sup>2</sup> for about 30 min. In the impregnation vessel, the selective heating of GNW samples during the MOCFD proc‐ ess facilitated selective metal deposition on the surface of the carbon nanostructures. In the preliminary vessel, the pressure and temperature of sc-CO2 were maintained at 11 MPa and 50˚C, respectively, and MeCpPtMe<sup>3</sup> was dissolved in the sc-CO2. In the impregnation vessel, the pressure and temperature of sc-CO2 were maintained at 9 MPa and 70˚C, respectively, and the temperature of GNW samples was controlled in the range of 70–170˚C. The solu‐ tions were mixed and Pt nanoparticles formation was carried out for 30 min; the vessel was then depressurized slowly in 30 min to atmospheric conditions. After the depressurization, additional heat treatment was not carried out in the present work.

#### **3.4. Characterization of platinum nanoparticles formed by metal-organic chemical fluid deposition using supercritical carbon dioxide**

Figures 14(a)–14(c) show TEM images of the Pt-supported GNW surface after the SCF-MOCFD at sample temperatures of 120, 150, and 170˚C, respectively. In this experiment, GNW samples were fabricated on the Si substrate by fluorocarbon (C2F6) PECVD assisted by

hydrogen radical injection, which comprises a parallel-plate very high frequency (VHF: 100 MHz), capacitively coupled plasma region, and a hydrogen radical injection source that em‐ ploys a surface-wave-excited microwave (2.45 GHz) H2 plasma [11]. As can be seen from these TEM images, the spatial density of the Pt nanoparticles (particle numbers/area) sup‐ ported on the GNW surface strongly depended on the sample temperature during the SCF-MOCFD, while the average size of the Pt nanoparticles increased from 1.5 to 3 nm with an increase of the sample temperature from 120 to 170˚C.

(72.8 atm) and 31.1 ˚C. Among SCFs, sc-CO<sup>2</sup> is particularly attractive since it is environmen‐

In the case of Pt deposition, the SCF using sc-CO<sup>2</sup> was first applied to the preparation of pol‐ ymer-supported Pt nanoparticles using dimethyl (1,5-cyclooctadiene) platinum(II), (PtMe2(cod)) as a precursor [47]. Erkey's group has demonstrated the preparation of Pt nanoparticles on a wide range of materials, including carbon aerogel, carbon black, silica aerogel, alumina, and Nafion [47–52]. In their method, PtMe2(cod) was dissolved in sc-CO<sup>2</sup> and impregnated into the supporting materials, and after depressurization the impregnated PtMe2(cod) molecules were then reduced to metallic Pt nanoparticles by heat treatment or by chemical reduction with hydrogen, resulting in the formation of uniformly dispersed nanoparticles with narrow size distributions. However, it took almost 10 hours to complete this process. In our case, in contrast, the supporting carbon nanostructures such as CNTs and GNWs were selectively heated in the sc-CO2 with Pt precursors during the process. Therefore, at the heated surface of the carbon nanostructures during *in situ* thermal reduc‐ tion under the SCF environment, decomposition of the adsorbed precursor molecules and

tally friendly and safe due to its low toxicity, low reactivity and nonflammability.

growth of the particles would occur without reduction process.

additional heat treatment was not carried out in the present work.

**deposition using supercritical carbon dioxide**

**supercritical carbon dioxide**

246 New Progress on Graphene Research

**3.3. Experimental procedure of metal-organic chemical fluid deposition using**

Figure 13 shows the SCF-MOCFD system employing sc-CO2 used for the deposition of Pt nanoparticles on the surface of GNWs. The MOCFD system is composed of two high-pres‐ sure stainless steel vessels with a compressor, heating units, and a reservoir for the metalorganic compound. The preliminary vessel contains a screw agitator. The temperature and pressure in each vessel can be set independently, so that two different supercritical condi‐ tions employing CO2 can be produced in these two vessels. As the precursor, (methylcyclo‐ pentadienyl) trimethyl platinum ((CH3C5H4)Pt(CH3)3: MeCpPtMe3) dissolved in hexane was used. In the preliminary vessel, the precursor was stirred with the sc-CO<sup>2</sup> for about 30 min. In the impregnation vessel, the selective heating of GNW samples during the MOCFD proc‐ ess facilitated selective metal deposition on the surface of the carbon nanostructures. In the preliminary vessel, the pressure and temperature of sc-CO2 were maintained at 11 MPa and 50˚C, respectively, and MeCpPtMe<sup>3</sup> was dissolved in the sc-CO2. In the impregnation vessel, the pressure and temperature of sc-CO2 were maintained at 9 MPa and 70˚C, respectively, and the temperature of GNW samples was controlled in the range of 70–170˚C. The solu‐ tions were mixed and Pt nanoparticles formation was carried out for 30 min; the vessel was then depressurized slowly in 30 min to atmospheric conditions. After the depressurization,

**3.4. Characterization of platinum nanoparticles formed by metal-organic chemical fluid**

Figures 14(a)–14(c) show TEM images of the Pt-supported GNW surface after the SCF-MOCFD at sample temperatures of 120, 150, and 170˚C, respectively. In this experiment, GNW samples were fabricated on the Si substrate by fluorocarbon (C2F6) PECVD assisted by

**Figure 13.** Schematic of the supercritical fluid, metal-organic chemical fluid deposition (SCF-MOCFD) system [53]

*Ex situ* X-ray photoelectron spectroscopy (XPS) analysis was conducted to gain an insight in‐ to the state of the platinum for the supported Pt surface fabricated by the SCF-MOCFD. Fig‐ ure 14(d) shows an XPS spectrum of the Pt(*4f*) region of the Pt-supported GNW film after the SCF-MOCFD at a sample temperature of 150˚C. The presence of two prominent sets of Pt (*4f*) peaks, corresponding to the *4f 7/2* and *4f 5/2* orbital states, is further confirmation of platinum being present on the GNW surface. The peak regions in Fig. 14(d) can be fitted with two sets of peaks at 71.4 eV (*4f 7/2*) and 74.6 eV (*4f 5/2*) [54]. These correspond to platinum in the metallic state, indicating that only pure Pt exists without being oxidized on the sur‐ face of the GNWs after the SCF-MOCFD.

The Pt/C ratio of the GNW film surface was obtained from the ratio of the intensities of the XPS C(*1s*) and Pt(*4f*) peaks. Figure 14(e) shows the relative Pt/C ratio of the surface of the GNW film as a function of the temperature of the GNW sample during the SCF-MOCFD process. As the sample temperature during the SCF-MOCFD increased up to 120�?˚C, the relative Pt/C ratio of the surface of the GNW film increased gradually. By further increasing the sample temperature above 120˚C, the relative Pt/C ratio increased rapidly.

#### **3.5. Mechanism of platinum nanoparticle formation by metal-organic chemical fluid deposition using supercritical carbon dioxide**

In the case of Pt deposition by the conventional process including impregnation, depressuri‐ zation, and reduction, PtMe2(cod) was dissolved in sc-CO2 and impregnated into the sup‐ porting materials, and after depressurization the impregnated PtMe2(cod) molecules were then reduced to metallic Pt nanoparticles by heat treatment or by chemical reduction with hydrogen. It was proposed that the precursor molecules in the sc-CO2 phase are adsorbed on the surface during the impregnation period, and after depressurization these adsorbed molecules are in turn reduced to elemental platinum and the resulting particles at the sur‐ face continue to grow until all the adsorbed precursor molecules are converted to the metal.

[56]. The reaction temperature at the surface would be a significant factor influencing the particle number density and particle size. When the temperature of GNWs is increased, both reduction of metal-organic precursors and surface migration of Pt atoms would be enhanced, which may lead to an increase in the particle number density and particle size. As can be seen from the TEM images in Figs. 14(a)–14(c), the average size of Pt nanopar‐ ticles increased from 1.5 to 3 nm with an increase in the sample temperature from 120

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**Figure 15.** Illustration of the formation model of Pt nanoparticles on the surface of carbon nanostructures using met‐

**4. Fabrication of graphene nanowalls on carbon fiber paper for fuel cell**

Carbon fiber paper (CFP) or carbon fiber cloth, which are composed of an open mesh of car‐ bon fibers, have been used as gas diffusion layer in proton exchange membrane (PEM) fuel cell application [57–59]. PEM fuel cells have been widely recognized as the most promising candidates for future power generating devices in the automotive, distributed power gener‐ ation and portable electronic applications. Waje, et al. demonstrated the preparation of Pt nanoparticles 2–2.5 nm in size on organically functionalized CNTs grown on CFP [60]. Very recently Lisi, et al. demonstrated the growth of GNWs on CFP by hot-filament CVD and in‐ vestigated the microstructure of the GNWs both at the tip and at the fiber–nanowall base interface [61]. Our current interest in GNWs is to use them as catalyst supports for Pt in

al-organic chemical fluid deposition in the supercritical CO2

**application**

to 170˚C, while the Pt particle number density increased drastically.

**Figure 14.** a)–(c) High-resolution TEM images of the surface of the GNWs supporting Pt nanoparticles after the SCF-MOCFD at sample temperatures of 120, 150, and 170�?˚C, respectively. Scale bar: 20 nm. (d) X-ray photoelectron spec‐ troscopy spectrum of the Pt-supported GNW film after SCF-MOCFD. (e) Relative Pt/C ratio of the surface of the GNW film as a function of temperature of the GNW sample during SCF-MOCFD [53]

In our case, in contrast, the supporting carbon nanostructures were selectively heated in the sc-CO2 with Pt precursors during the process. Figure 15 depicts the model of Pt nano‐ particle formation on the surface of carbon nanostructures using metal-organic chemical fluid deposition in the supercritical CO2. At the heated surface of the carbon nanostruc‐ tures during *in situ* thermal reduction under the SCF environment, decomposition of the adsorbed precursor molecules and growth of the particles would occur. GNWs have been reported to consist of nano-domains a few tens of nanometers in size, and individual GNWs were found to have many defects [24,55]. It is suggested that the surface-migrat‐ ing Pt adatoms, produced by the decomposition of MeCpPtMe3 precursors, merge to form Pt clusters from several Pt atoms preferentially at chemically active sites such as defects and grain boundaries on the surface of the carbon nanostructures, resulting in the nucle‐ ation of Pt nanoparticles. It has been reported recently that the density of Pt nanoparti‐ cles formed by SCF-MOCFD method increased with increase in the surface defect density [56]. The reaction temperature at the surface would be a significant factor influencing the particle number density and particle size. When the temperature of GNWs is increased, both reduction of metal-organic precursors and surface migration of Pt atoms would be enhanced, which may lead to an increase in the particle number density and particle size. As can be seen from the TEM images in Figs. 14(a)–14(c), the average size of Pt nanopar‐ ticles increased from 1.5 to 3 nm with an increase in the sample temperature from 120 to 170˚C, while the Pt particle number density increased drastically.

porting materials, and after depressurization the impregnated PtMe2(cod) molecules were then reduced to metallic Pt nanoparticles by heat treatment or by chemical reduction with hydrogen. It was proposed that the precursor molecules in the sc-CO2 phase are adsorbed on the surface during the impregnation period, and after depressurization these adsorbed molecules are in turn reduced to elemental platinum and the resulting particles at the sur‐ face continue to grow until all the adsorbed precursor molecules are converted to the metal.

248 New Progress on Graphene Research

**Figure 14.** a)–(c) High-resolution TEM images of the surface of the GNWs supporting Pt nanoparticles after the SCF-MOCFD at sample temperatures of 120, 150, and 170�?˚C, respectively. Scale bar: 20 nm. (d) X-ray photoelectron spec‐ troscopy spectrum of the Pt-supported GNW film after SCF-MOCFD. (e) Relative Pt/C ratio of the surface of the GNW

In our case, in contrast, the supporting carbon nanostructures were selectively heated in the sc-CO2 with Pt precursors during the process. Figure 15 depicts the model of Pt nano‐ particle formation on the surface of carbon nanostructures using metal-organic chemical fluid deposition in the supercritical CO2. At the heated surface of the carbon nanostruc‐ tures during *in situ* thermal reduction under the SCF environment, decomposition of the adsorbed precursor molecules and growth of the particles would occur. GNWs have been reported to consist of nano-domains a few tens of nanometers in size, and individual GNWs were found to have many defects [24,55]. It is suggested that the surface-migrat‐ ing Pt adatoms, produced by the decomposition of MeCpPtMe3 precursors, merge to form Pt clusters from several Pt atoms preferentially at chemically active sites such as defects and grain boundaries on the surface of the carbon nanostructures, resulting in the nucle‐ ation of Pt nanoparticles. It has been reported recently that the density of Pt nanoparti‐ cles formed by SCF-MOCFD method increased with increase in the surface defect density

film as a function of temperature of the GNW sample during SCF-MOCFD [53]

**Figure 15.** Illustration of the formation model of Pt nanoparticles on the surface of carbon nanostructures using met‐ al-organic chemical fluid deposition in the supercritical CO2

## **4. Fabrication of graphene nanowalls on carbon fiber paper for fuel cell application**

Carbon fiber paper (CFP) or carbon fiber cloth, which are composed of an open mesh of car‐ bon fibers, have been used as gas diffusion layer in proton exchange membrane (PEM) fuel cell application [57–59]. PEM fuel cells have been widely recognized as the most promising candidates for future power generating devices in the automotive, distributed power gener‐ ation and portable electronic applications. Waje, et al. demonstrated the preparation of Pt nanoparticles 2–2.5 nm in size on organically functionalized CNTs grown on CFP [60]. Very recently Lisi, et al. demonstrated the growth of GNWs on CFP by hot-filament CVD and in‐ vestigated the microstructure of the GNWs both at the tip and at the fiber–nanowall base interface [61]. Our current interest in GNWs is to use them as catalyst supports for Pt in PEM fuel cells. In this section, in order to demonstrate the usefulness of GNWs in the fuel cell application, GNWs were directly grown on the CFP using PECVD. Subsequently, Pt nanoparticles were formed on the surface of GNWs using the SCF-MOCFD method. This configuration ensures that all the supported Pt nanoparticles are in electrical contact with the external electrical circuit. Such a design would improve Pt utilization and potentially de‐ crease Pt usage.

was carried out for 30 min. Compared to the SEM images of typical as-grown GNW films without the SCF treatment, no change in the surface morphology was observed after the SCF treatment. It was confirmed that the unique nanostructure of the GNWs was main‐ tained, even after being exposed to the high-pressure fluid. Figure 17(a) shows SEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min. It was found that dispersed Pt nanoparticles of approximately 2 nm in diameter were support‐ ed on the surfaces of GNWs grown on CFP. The area density of Pt nanoparticles on GNW surface was approximately 3×1012 cm-2. Figure 17(b) shows TEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min. The GNWs consist of nano-domains and individual GNWs were found to have many defects. As shown in Fig. 17(b), Pt nanoparticles were formed preferentially at the domain boundaries on the surface

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**Figure 17.** a) SEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min. (b)

A test cell unit using Pt-supported GNW/CFP electrode was constructed experimentally. A schematic of a single PEM fuel cell with active surface area of 1×1 cm2 is shown in Fig. 18. The membrane electrode assembly (MEA) consists of a Nafion 115 membrane in combina‐ tion with Pt- supported GNWs on CFPs. At present, the voltage-current curve for the test PEM fuel cell exhibited unexpectedly poor performance due to the high ohmic resistance. Very recently, Shin, et al. demonstrated the preparation of Pt nanoparticles with an average diameter of 3.5 nm on GNWs by a solution-reduction method, and investigated the electro‐ catalytic activity of powdered Pt/ GNWs peeled off from the substrate [63]. They suggested that the domain structure of GNWs is useful as catalytic support for fuel cells, although they did not demonstrate an actual PEM fuel cell using Pt/GNWs. For the evaluation of the Ptsupported GNWs on CFPs as catalyst layer for the practical use in our case, it is necessary to realize optimum operating conditions including compression, temperature, pressure, and

TEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min

**4.3. Fuel cell unit using Pt-supported GNW/CFP electrode**

of the GNWs.

ionomer loading.

#### **4.1. Growth of graphene nanowalls on carbon fiber paper**

Commercially available CFP (Engineered Fibers Technology, Spectracarb 2050A porous car‐ bon-carbon paper, 200 μm thick) was decorated with GNWs using RF-ICP employing Ar/CH4 mixture. The growth experiments were carried out on CFP and Si substrates for 30 min at RF power of 500 W, total gas pressure of 20 mTorr, substrate (CFP) temperature of 720 °C, and flow rates of Ar and CH4 of 100 and 50 sccm, respectively.

**Figure 16.** a), (b) SEM images of carbon fiber used in this study. (c)–(f) SEM images of GNWs grown on carbon fiber by ICP-CVD for 30 min at different magnifications [62].

Figures 16(a) and 16(b) show SEM images of carbon fiber used in this study. SEM images of GNWs grown on carbon fiber by ICP-CVD for 30 min at different magnifications are shown in Figs. 16(c)–16(f), indicating that GNWs were successfully grown on the CFP using ICP-CVD. As shown in Fig. 16(e), GNWs were grown almost vertically on the surface of carbon fibers forming paper structure. The height of GNWs grown on the CFP at the surface facing the plas‐ ma was about 1.5 μm. Interestingly the GNW growth proceeds in a conformal manner all the way around each carbon fiber into the CFP to a depth of a few tens of micrometers.

#### **4.2. Pt nanoparticle formation on GNW-decorated carbon fiber paper**

Pt nanoparticles were prepared on the GNWs grown on the CFP by the SCF-MOCFD meth‐ od. The pressure and temperature of sc-CO2 were 10 MPa and 130 ˚C, respectively, and the temperature of GNW-decorated CFP was maintained at 180 °C. Pt nanoparticle formation was carried out for 30 min. Compared to the SEM images of typical as-grown GNW films without the SCF treatment, no change in the surface morphology was observed after the SCF treatment. It was confirmed that the unique nanostructure of the GNWs was main‐ tained, even after being exposed to the high-pressure fluid. Figure 17(a) shows SEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min. It was found that dispersed Pt nanoparticles of approximately 2 nm in diameter were support‐ ed on the surfaces of GNWs grown on CFP. The area density of Pt nanoparticles on GNW surface was approximately 3×1012 cm-2. Figure 17(b) shows TEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min. The GNWs consist of nano-domains and individual GNWs were found to have many defects. As shown in Fig. 17(b), Pt nanoparticles were formed preferentially at the domain boundaries on the surface of the GNWs.

**Figure 17.** a) SEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min. (b) TEM image of the surface of the GNW supporting Pt nanoparticles after the SCF-MOCFD for 30 min

#### **4.3. Fuel cell unit using Pt-supported GNW/CFP electrode**

PEM fuel cells. In this section, in order to demonstrate the usefulness of GNWs in the fuel cell application, GNWs were directly grown on the CFP using PECVD. Subsequently, Pt nanoparticles were formed on the surface of GNWs using the SCF-MOCFD method. This configuration ensures that all the supported Pt nanoparticles are in electrical contact with the external electrical circuit. Such a design would improve Pt utilization and potentially de‐

Commercially available CFP (Engineered Fibers Technology, Spectracarb 2050A porous car‐ bon-carbon paper, 200 μm thick) was decorated with GNWs using RF-ICP employing Ar/CH4 mixture. The growth experiments were carried out on CFP and Si substrates for 30 min at RF power of 500 W, total gas pressure of 20 mTorr, substrate (CFP) temperature of

**Figure 16.** a), (b) SEM images of carbon fiber used in this study. (c)–(f) SEM images of GNWs grown on carbon fiber by

Figures 16(a) and 16(b) show SEM images of carbon fiber used in this study. SEM images of GNWs grown on carbon fiber by ICP-CVD for 30 min at different magnifications are shown in Figs. 16(c)–16(f), indicating that GNWs were successfully grown on the CFP using ICP-CVD. As shown in Fig. 16(e), GNWs were grown almost vertically on the surface of carbon fibers forming paper structure. The height of GNWs grown on the CFP at the surface facing the plas‐ ma was about 1.5 μm. Interestingly the GNW growth proceeds in a conformal manner all the

Pt nanoparticles were prepared on the GNWs grown on the CFP by the SCF-MOCFD meth‐ od. The pressure and temperature of sc-CO2 were 10 MPa and 130 ˚C, respectively, and the temperature of GNW-decorated CFP was maintained at 180 °C. Pt nanoparticle formation

way around each carbon fiber into the CFP to a depth of a few tens of micrometers.

**4.2. Pt nanoparticle formation on GNW-decorated carbon fiber paper**

**4.1. Growth of graphene nanowalls on carbon fiber paper**

ICP-CVD for 30 min at different magnifications [62].

720 °C, and flow rates of Ar and CH4 of 100 and 50 sccm, respectively.

crease Pt usage.

250 New Progress on Graphene Research

A test cell unit using Pt-supported GNW/CFP electrode was constructed experimentally. A schematic of a single PEM fuel cell with active surface area of 1×1 cm2 is shown in Fig. 18. The membrane electrode assembly (MEA) consists of a Nafion 115 membrane in combina‐ tion with Pt- supported GNWs on CFPs. At present, the voltage-current curve for the test PEM fuel cell exhibited unexpectedly poor performance due to the high ohmic resistance. Very recently, Shin, et al. demonstrated the preparation of Pt nanoparticles with an average diameter of 3.5 nm on GNWs by a solution-reduction method, and investigated the electro‐ catalytic activity of powdered Pt/ GNWs peeled off from the substrate [63]. They suggested that the domain structure of GNWs is useful as catalytic support for fuel cells, although they did not demonstrate an actual PEM fuel cell using Pt/GNWs. For the evaluation of the Ptsupported GNWs on CFPs as catalyst layer for the practical use in our case, it is necessary to realize optimum operating conditions including compression, temperature, pressure, and ionomer loading.

at a scan rate of 100 mV/s. As shown in Fig. 19, a work potential window nearly 3 V was obtained for the GNW electrode, which was comparable to that for B-doped diamond elec‐ trode. The oxygen evolution for the GNW electrodes occurs at about 1 V and the hydrogen evolution at –2 V. Reported potential windows for glassy carbon and highly oriented pyro‐

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**Figure 19.** Cyclic voltammograms for the GNWs and boron-doped diamond electrodes in HNO3 (0.2 Mol/l); at 100

**Figure 20.** Cyclic voltammograms of GNW electrode in the solution of PBS with DA, AA, and UA, at 100 mV/s scan rate

Electrochemical activity of GNW electrode has been investigated by cyclic voltammetry measurements in an aqueous solution of ferrocyanide and a faster electron transfer between the electrolyte and the nanowall (or nanosheet) surface has been demonstrated [17–20]. Fur‐ thermore, biosensing with GNWs is a promising application. Dopamine (DA), ascorbic acid (AA), and uric acid (UA) are compounds of great biomedical interest, which all are essential biomolecules in our body fluids. Figure 20 shows cyclic voltammogram curves obtained from GNW electrode in the phosphate buffer solution (PBS) with DA, AA, and UA, at 100 mV/s scan rate. At present, researches on the sensing of biological molecules became popu‐ lar. Shang, et al. demonstrated the excellent electrocatalytic activity of multilayer graphene nanoflake films for simultaneously determining DA, AA and UA in PBS [17]. The GNWbased electrochemical platform, which possesses large surface area with edges and electro‐

lytic graphite electrodes are 1.5 and 2.0 V, respectively [67].

mV/s scan rate

**Figure 18.** A schematic of a test single proton exchange membrane (PEM) fuel cell. The membrane electrode assembly (MEA) consists of a Nafion 115 membrane in combination with Pt-loaded GNWs on CFPs [62].

#### **5. Toward the application to electrochemical sensors and biosensors**

There has been an increase in research into the biological applications of CVD diamond. Di‐ amond has attractive characteristics for some biological applications, such as its wide poten‐ tial window, chemical–physical stability, and biocompatibility. DNA and proteins were immobilized on the surface of diamond and nanodiamond films for bio-sensing application [35–38]. Covalent modification of diamond surfaces with molecular monolayers serves as a starting point for linking biomolecules such as DNA and proteins to surfaces. In these cases, diamond surface is modified with several types of surface termination, e.g., C–NH2, C–OH, and C–COOH. It is considered that the surfaces and edges of GNWs can also be modified with similar surface termination. Therefore, covalent immobilization of DNA and proteins on the GNWs will be realized. In the near future, GNWs will be used as a stable, highly se‐ lective platform in subsequent surface hybridization processes.

Recently, graphene has proved to be an excellent nanomaterial for applications in electro‐ chemistry. Graphene-based materials with large surface area are useful as electrodes for electrochemical sensors and biosensors [62–64]. Especially, GNWs and related carbon nano‐ structures can be one of the best electrode materials to investigate basic electrochemical phe‐ nomena, due to their edge defects and graphene structures.

It is interesting to investigate the electrochemical properties of GNWs as electrochemical electrodes. In order to evaluate the potential window of GNW electrode, electrochemical measurements were carried out with a conventional three-electrode arrangement controlled by a commercial potentiostat, with a GNW (boron (B)-doped diamond for comparison) working electrode, a Pt coil counter electrode and an Ag/AgCl reference electrode. The back face of the substrate and the electrical contact were protected from the electrolyte solution by insulating epoxy adhesive, so that only an active GNW electrode area was exposed to the electrolyte solution. Figure 19 shows the cyclic voltammograms for the GNW and B-doped diamond electrodes in HNO3 (0.2 mol/l) obtained in the potential range between –5 and 5 V at a scan rate of 100 mV/s. As shown in Fig. 19, a work potential window nearly 3 V was obtained for the GNW electrode, which was comparable to that for B-doped diamond elec‐ trode. The oxygen evolution for the GNW electrodes occurs at about 1 V and the hydrogen evolution at –2 V. Reported potential windows for glassy carbon and highly oriented pyro‐ lytic graphite electrodes are 1.5 and 2.0 V, respectively [67].

**Figure 18.** A schematic of a test single proton exchange membrane (PEM) fuel cell. The membrane electrode assembly

There has been an increase in research into the biological applications of CVD diamond. Di‐ amond has attractive characteristics for some biological applications, such as its wide poten‐ tial window, chemical–physical stability, and biocompatibility. DNA and proteins were immobilized on the surface of diamond and nanodiamond films for bio-sensing application [35–38]. Covalent modification of diamond surfaces with molecular monolayers serves as a starting point for linking biomolecules such as DNA and proteins to surfaces. In these cases, diamond surface is modified with several types of surface termination, e.g., C–NH2, C–OH, and C–COOH. It is considered that the surfaces and edges of GNWs can also be modified with similar surface termination. Therefore, covalent immobilization of DNA and proteins on the GNWs will be realized. In the near future, GNWs will be used as a stable, highly se‐

Recently, graphene has proved to be an excellent nanomaterial for applications in electro‐ chemistry. Graphene-based materials with large surface area are useful as electrodes for electrochemical sensors and biosensors [62–64]. Especially, GNWs and related carbon nano‐ structures can be one of the best electrode materials to investigate basic electrochemical phe‐

It is interesting to investigate the electrochemical properties of GNWs as electrochemical electrodes. In order to evaluate the potential window of GNW electrode, electrochemical measurements were carried out with a conventional three-electrode arrangement controlled by a commercial potentiostat, with a GNW (boron (B)-doped diamond for comparison) working electrode, a Pt coil counter electrode and an Ag/AgCl reference electrode. The back face of the substrate and the electrical contact were protected from the electrolyte solution by insulating epoxy adhesive, so that only an active GNW electrode area was exposed to the electrolyte solution. Figure 19 shows the cyclic voltammograms for the GNW and B-doped diamond electrodes in HNO3 (0.2 mol/l) obtained in the potential range between –5 and 5 V

**5. Toward the application to electrochemical sensors and biosensors**

(MEA) consists of a Nafion 115 membrane in combination with Pt-loaded GNWs on CFPs [62].

252 New Progress on Graphene Research

lective platform in subsequent surface hybridization processes.

nomena, due to their edge defects and graphene structures.

**Figure 19.** Cyclic voltammograms for the GNWs and boron-doped diamond electrodes in HNO3 (0.2 Mol/l); at 100 mV/s scan rate

**Figure 20.** Cyclic voltammograms of GNW electrode in the solution of PBS with DA, AA, and UA, at 100 mV/s scan rate

Electrochemical activity of GNW electrode has been investigated by cyclic voltammetry measurements in an aqueous solution of ferrocyanide and a faster electron transfer between the electrolyte and the nanowall (or nanosheet) surface has been demonstrated [17–20]. Fur‐ thermore, biosensing with GNWs is a promising application. Dopamine (DA), ascorbic acid (AA), and uric acid (UA) are compounds of great biomedical interest, which all are essential biomolecules in our body fluids. Figure 20 shows cyclic voltammogram curves obtained from GNW electrode in the phosphate buffer solution (PBS) with DA, AA, and UA, at 100 mV/s scan rate. At present, researches on the sensing of biological molecules became popu‐ lar. Shang, et al. demonstrated the excellent electrocatalytic activity of multilayer graphene nanoflake films for simultaneously determining DA, AA and UA in PBS [17]. The GNWbased electrochemical platform, which possesses large surface area with edges and electro‐ chemical activity, offers great promise for providing a new class of nanostructured electrodes for biosensing and energy-conversion applications.

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#### **6. Conclusion**

Self-organized graphite sheet nanostructures composed of graphene have been studied in‐ tensively. Graphene nanowalls and related sheet nanostructures are layered graphenes with open boundaries. The sheets form a self-supported network of wall structures with thick‐ nesses in the range from a few nanometers to a few tens of nanometers, and with a high as‐ pect ratio. The large surface area and sharp edges of graphene nanowalls could prove useful for a number of different applications.

Using graphene nanowalls as templates would be the most promising and important appli‐ cation. Graphene nanowalls can be used as templates for fabricating a variety of nanostruc‐ tured materials based on the surface morphology of the graphene nanowalls and nanocomposites of carbon and nanoparticles of other materials. These structures could prove useful in batteries, sensors, solar cells, electrodes, and biomedical devices. For this purpose, it is necessary to establish decorating methods of graphene nanowall surface with a variety of materials. Furthermore, it is important to evaluate electrochemical characteris‐ tics of nanocomposites of carbon and other materials systematically.

In order to demonstrate the usefulness of graphene nanowalls in the fuel cell application, graphene nanowalls were directly grown on the carbon fiber paper using the inductively coupled plasma-enhanced chemical vapor deposition method. Subsequently, highly dis‐ persed Pt nanoparticles 2 nm in size were formed on the surface of graphene nanowalls us‐ ing metal-organic chemical deposition employing supercritical CO2. This configuration ensures that all the supported Pt nanoparticles are in electrical contact with the external elec‐ trical circuit. Such a design would improve Pt utilization and potentially decrease Pt usage. Pt-supported graphene nanowalls grown on the carbon fiber paper will be well suited to the application for the electrodes of fuel cells.

Furthermore, the application as a biosensor using GNWs was briefly described. The GNWbased electrochemical platform offers great promise for providing a new class of nanostruc‐ tured electrodes for electrochemical sensing, biosensing and energy-conversion applications.

#### **Author details**

Mineo Hiramatsu1\*, Hiroki Kondo2 and Masaru Hori2

1 Meijo University, Japan

2 Nagoya University, Japan

#### **References**

chemical activity, offers great promise for providing a new class of nanostructured

Self-organized graphite sheet nanostructures composed of graphene have been studied in‐ tensively. Graphene nanowalls and related sheet nanostructures are layered graphenes with open boundaries. The sheets form a self-supported network of wall structures with thick‐ nesses in the range from a few nanometers to a few tens of nanometers, and with a high as‐ pect ratio. The large surface area and sharp edges of graphene nanowalls could prove useful

Using graphene nanowalls as templates would be the most promising and important appli‐ cation. Graphene nanowalls can be used as templates for fabricating a variety of nanostruc‐ tured materials based on the surface morphology of the graphene nanowalls and nanocomposites of carbon and nanoparticles of other materials. These structures could prove useful in batteries, sensors, solar cells, electrodes, and biomedical devices. For this purpose, it is necessary to establish decorating methods of graphene nanowall surface with a variety of materials. Furthermore, it is important to evaluate electrochemical characteris‐

In order to demonstrate the usefulness of graphene nanowalls in the fuel cell application, graphene nanowalls were directly grown on the carbon fiber paper using the inductively coupled plasma-enhanced chemical vapor deposition method. Subsequently, highly dis‐ persed Pt nanoparticles 2 nm in size were formed on the surface of graphene nanowalls us‐ ing metal-organic chemical deposition employing supercritical CO2. This configuration ensures that all the supported Pt nanoparticles are in electrical contact with the external elec‐ trical circuit. Such a design would improve Pt utilization and potentially decrease Pt usage. Pt-supported graphene nanowalls grown on the carbon fiber paper will be well suited to the

Furthermore, the application as a biosensor using GNWs was briefly described. The GNWbased electrochemical platform offers great promise for providing a new class of nanostruc‐ tured electrodes for electrochemical sensing, biosensing and energy-conversion applications.

and Masaru Hori2

electrodes for biosensing and energy-conversion applications.

tics of nanocomposites of carbon and other materials systematically.

**6. Conclusion**

254 New Progress on Graphene Research

for a number of different applications.

application for the electrodes of fuel cells.

Mineo Hiramatsu1\*, Hiroki Kondo2

1 Meijo University, Japan

2 Nagoya University, Japan

**Author details**


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