3C-SiC Hybrid Nanolaminate - ACS Applied Materials

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Graphene/3C-SiC Hybrid Nanolaminate Hao Zhuang, Bing Yang, Steffen Heuser, Nan Huang, Haiyuan Fu, and Xin Jiang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.5b09794 • Publication Date (Web): 09 Dec 2015 Downloaded from http://pubs.acs.org on December 10, 2015

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ACS Applied Materials & Interfaces

Graphene/3C-SiC Hybrid Nanolaminate Hao Zhuang1,#, Bing Yang2,#, Steffen Heuser1, Nan Huang2, Haiyuan Fu1, Xin Jiang1,2,*

1

Institute of Materials Engineering, University of Siegen, Paul-Bonatz-Str. 9-11, 57076 Siegen, Germany

2

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016 China

#

Both authors equally contribute to the work.

Tel: +49 271 740 2966; Fax: +49 271 740 2442 E-mail: [email protected] (Xin Jiang) 1

ACS Paragon Plus Environment


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Abstract In this work, we demonstrate a one-step approach to create graphene/3C-SiC nanolaminate structure using microwave plasma chemical vapor deposition technique. Layer-by-layer arrangement of thin 3C-SiC layers and graphene sheets is obtained with the thicknesses of the individual 3C-SiC layers and graphene sheets being 5-10 nm and 2-5 nm, respectively. An intimate contact between 3C-SiC and the graphene sheets is achieved and the nanolaminate film shows a high room temperature conductivity of 96.1 S/cm. A dedicated structural analysis of the nanolaminates by means of high-resolution transmission electron microscopy (HRTEM) reveals that the growth of the nanolaminates follows an iterative process: preferential graphene nucleation around the planar defects at the central region of the SiC layer, leading to the “splitting” of the SiC layer; and the thickening of the SiC layer after being “split”. A growth mechanism based on both kinetics and thermodynamics is proposed. Following the proposed the mechanism, it is possible to control the layer thickness of the graphene/3C-SiC hybrid nanolaminate by manipulating the carbon concentration in the gas phase, which is further experimentally verified. The high electrical conductivity, large surface area porous structure, feasible integration on different substrates (metal: Mo, semiconductor: Si and 2H-SiC, insulator: diamond) of the graphene/3C-SiC hybrid nanolaminate as well as other unprecedented advantages of the nanolaminate structure make it very promising for applications in mechanical, energy and sensor related areas.

Keywords: Silicon carbide, graphene, laminate, chemical vapor deposition, interface 2


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Introduction Graphene, a two-dimensional nanostructure of sp2-hybridized carbon, has drawn intensive research attention because of its extraordinary electrical, chemical, thermal and mechanical properties.1-2 Even though its applications have been continuously explored over time,3-7 its zero band gap limits its usage in certain electrical and optical related areas. To overcome this shortcoming, integration of graphene with inorganic semiconductors has become popular way. In the course of time, various graphene-based composites, i.e., graphene/Si,8 graphene/SiC,9-11 graphene/ZnO,12 graphene/TiO2,13 etc., have been developed, making a significant advance. Nevertheless, the performance of a composite is known to be highly dependent on the arrangement of its components. For example, an intimate contact between graphene and semiconductor can improve the electron transfer efficiency, which effectively improves the device performance.14 Therefore, tailoring its structure poses an important concern during the fabrication of graphene-based composites,15-16 which, however, still remains a big challenge nowadays. As a special class of composite structures the laminated architecture, a layerby-layer stacking of different components, has been reported to produce a composite with very different characteristics from any individual components. Especially when its individual layer thickness reaches down to the nanometer range (smaller than the characteristic length scale that defines the physical property), so-called hybrid nanolaminate, novel electrical, optical and mechanical properties arise.17-23 Motivated by the ambition to combine the inherent physical and chemical properties of graphene with the unprecedented advantages of nanolaminate structure, several graphene-based hybrid nanolaminates have been recently produced.24-29 Based on different compositing materials, such an endeavor creates new groups of materials that have

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shown outstanding performances in a wide range of applications, i.e., DNA sensor,24 mechanical strengthening,25-26 lithium battery,27-28 and radiation damage resistor.29 In contrast to the tempting rewards, however, the fabrication methods for the graphenebased hybrid nanolaminates are quite limited. Till now, only physical transferring of graphene24-25 and wet chemical approaches27-28 have been reported. Regardless of the complications involved in these methods, the limited fabrication approaches also constrain the types of the hybrid nanolaminates that are currently available. To overcome such a dilemma, in this contribution, we aim at developing a new and simple approach for the fabrication graphene-based hybrid nanolaminates. Among the various semiconductors we choose SiC as the compositing material. This is because it possesses high physical and chemical stabilities, wide electrochemical working potential and suitable band gap for visible light adsorption (2.4 eV for 3C-SiC).30-31 By combining graphene with SiC, the SiC/graphene composites have been demonstrated to show promising applications as sensors, photocatalysis and catalyst supports.9-10 Tailoring the structure of the graphene/SiC composite into a hybrid nanolaminate structure can, from one hand, be a boost for the above applications, because the larger interfacial area between SiC and graphene leads to a higher electron transfer efficiency. From the other hand, it could even evoke new, yet unforeseen, phenomena based on this composite system, which possibly widens its applications. Motivated by the above, we demonstrate in the present study a novel one-step approach to fabricate graphene/SiC hybrid nanolaminates by microwave plasma chemical vapor deposition (MWCVD) technique. An individual layer thickness less than 10 nm has been successfully achieved. The approach is found to be compatible with various substrates, i.e., metal (Mo), semiconductor (Si and 2HSiC), and insulator (polycrystalline diamond), facilitating the universal integration of

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the hybrid nanolaminate. The marriage of the outstanding physical and chemical properties of graphene and 3C-SiC with the unique features of the nanolaminate structure could open the door to a plethora of applications; among them the applications as sensors and photocatalysis are readily foreseen.9-10

Experimental Section Single-crystalline (100) Si wafers were firstly utilized as substrates for the growth of the graphene/SiC hybrid nanolaminate. The reactive gas featured a mixture of H2, methane and Si(CH3)4 [tetramethylsilane (TMS)] with a total flow rate of 400 sccm (standard cubic centimeter per minute at STP). The methane flow rate was fixed at 9 sccm and the concentration of TMS in the gas phase was 345 ppm. During the deposition, the Si substrate was firstly pretreated in pure H2 plasma for 5 min at a microwave power of 700 W, a gas pressure of 20 Torr, and a substrate temperature of ~950 °C. The goal of this pretreatment was to remove the thin natural SiO2 surface layer for epitaxial growth of SiC. Subsequently, CH4 and TMS were introduced to the gas phase in order to initiate the growth of SiC and graphene structures. The total growth duration was 4 hours. During growth, the microwave power, gas pressure and substrate temperature were kept constant. To check the compatibility of the deposition process on other substrates, single crystalline (0001) 2H-SiC wafer, polycrystalline diamond and Mo were also employed as the substrates for deposition. A scanning electron microscope (SEM, Zeiss Ultra 55) was used to obtain the plane and cross-sectional microstructure of the samples. Investigation by transmission electron microscopy (TEM) was carried out on Philips CM20 and FEI G2 F20 to acquire cross-sectional information of the samples as well as the detailed information of individual structures. Micro Raman scattering studies were carried out to understand the structural order of graphene based on the typical phonon lines in the 5



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obtained spectra. The 532 nm line of a Nd:YVO4 diode-pumped solid-state laser was used as the excitation source.

Results and Discussion Structure and properties Figure 1a depicts the overall scanning electron microscopic (SEM) observation of the graphene/3C-SiC hybrid nanolaminates. More detailed observation of the surface and cross-sectional morphology is shown in Figure 1b-d. All the nanolaminates are arranged in clusters (Figure 1b and c). A layer-by-layer stacking of thin layers with their thicknesses in the nanometer range is observed in every cluster, confirming the formation of nanolaminate structure in the present study. The laminated arrangement can be also seen from the cross-sectional view, which is clearly denoted by the arrow in Figure 1d. The amount of layers in the nanolaminates differs, leading to a certain disparity in the width of the nanolaminates. Most of the nanolaminates show thicknesses ranging from ~100 to ~300 nm. Nevertheless, several very thin nanolaminates (tens of nanometers) are also observable in Figure 1c. Their lengths are not uniform either, which varies from ~0.3 µm to ~1 µm. Their height is determined from Figure 1d to be ~2.5 µm. In addition to the above geometrical characteristics, the nanolaminates show certain preferential alignment. A large number of nanolaminates lie in perpendicular to one another; some of them are marked in yellow in Figure 1b for a better illustration. Such a geometrical arrangement leads to the formation of nanometer-sized pores in the film, which in turn makes the nanolaminate film flexible. As shown in Figure 1e, the freestanding nanolaminate film can be bent under external force without breaking. This feature facilitates the transportation of the nanolaminate film onto curved substrates for applications.

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Figure 1. (a) – (d) SEM images of the typical graphene/3C-SiC hybrid nanolaminate grown on the Si substrate. (a) Overall observation; (b) and (c) surface view; (d) crosssectional view, the arrow indicates the observation of laminated structure; (e) an optical photo of a flexible freestanding graphene/3C-SiC hybrid nanolaminate film; (f) Raman spectra.

Figure 1f shows the Raman spectra of the nanolaminates, from which the phase information can be derived. The typical features of graphite including the G band and 2D band are observed with higher intensity of the G band than that of the 2D band, indicating the presence of multilayer graphene. Other peaks like D band and D+D’ band are also observable because of the existence of defects or disordering in graphene. In addition to the features related to graphene, a small and broad peak locating at ~800 cm-1 also presents. This peak can be assigned to the transverse optical (TO) phonon of SiC,32 confirming the presence of SiC. The low intensity of the SiC Raman peak is because of the very low Raman efficiency of SiC in comparison to that of carbon,33 but it doesn’t imply an extremely low amount of SiC in the structure. Similar phenomenon has been observed while co-depositing SiC structures with diamond.34-35

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Figure 2 shows plane-view transmission electron microscopy (TEM) observation of the nanolaminates. In accordance with the SEM observations in Figure 1, the low magnification TEM image in Figure 2a shows the porous characteristic of the nanolaminate film. Inset of Figure 2a depicts the corresponding selective area electron diffraction (SAED) pattern. Diffraction rings are clearly observable indicating the overall structure is randomly oriented. Such a phenomenon is actually in good accordance with the SEM observation: even though the nanolaminates are locally perpendicular to one another, these perpendicular nanolaminates are randomly distributed in the film, showing a random orientation of the overall structure. Most of the diffraction rings can be assigned to the 3C-SiC crystal planes. A weak ring is also observed close to the center of the pattern (denoted as “G”). It is assigned to the graphene plane with an inter-planar spacing of ~0.34 nm. Its low intensity can be explained by the small amount of graphene in the structures and the damage in its crystallinity while preparing the TEM sample (induced by Ar ion irradiation during ion milling).

Figure 2. TEM images of the graphene/3C-SiC hybrid nanolaminates from the surface view. (a) Overall morphology, inset shows the corresponding SAED pattern; (b) and

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(c) higher magnification TEM image of one individual nanolaminate; (d) HRTEM image of nanolaminate, the inset shows the corresponding FFT pattern; (e) IFFT reconstruction of (d), the red color indicates 3C-SiC and the blue color indicates the graphene sheets; (f) HRTEM image of another hybrid nanolaminate.

Higher magnification TEM images regarding one nanolaminate are shown in Figure 2b – d. From Figure 2b and c, it is obvious that two kinds of layers with different thicknesses exist: the thicker one is 4-10 nm (refer to the statistical results in the following) and the thinner one is 2-5 nm in thickness, confirming the formation of hybrid nanolaminates structure in the present study. Figure 2d shows the high resolution TEM (HRTEM) image of the layers. The corresponding fast Fourier transformation (FFT) pattern in the inset confirms again the co-existence of graphene and 3C-SiC. Nevertheless, the pattern indexed to graphene is not sharp but slightly spreads out because of the disordering in the graphene sheets. Figure 2e depicts the inverse FFT (IFFT) re-construction of Figure 2d from the FFT pattern. The distribution of graphene and 3C-SiC can be distinguished by the different colors: red for 3C-SiC and blue for graphene. By comparing Figure 2e with Figure 2d, it can be concluded that the thicker layers in Figure 2b and c are 3C-SiC and the thinner ones are graphene. Figure 2f shows a more detailed HRTEM observation at the graphene/3C-SiC interface. The basal plane of the SiC layer is observed to be the (111) 3C-SiC plane with an interplanar spacing of 0.25 nm. The graphene sheets are directly adjacent to the 3C-SiC crystallite with its basal plane parallel to the (111) plane of 3C-SiC, showing an intimate contact between graphene and 3C-SiC. Besides Si substrate, the current deposition approach can be also applied to directly deposit the nanolaminates on other substrates. Figure S1 in Supporting

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Information presents the SEM images of nanolaminate films deposited on different technologically important substrates such as metal (Mo), semiconductor (2H-SiC) and insulator (diamond). Similar structures are obtained on all the substrates, indicating the feasibility in the easy integration of the nanolaminates. In addition to the above, the nanolaminate film also shows high room temperature conductivity of 96.1 S/cm, as measured by four-point-probe technique. It is noteworthy that the electrical conductivity of the nanolaminate film is dependent on the thickness of the film. For the film thickness of 0.5, 1 and 4 µm, the conductivities are 27.5, 60.1 and 94.3 S/cm, respectively, with the saturation of the conductivity achieved at a film thickness of ~2.5 µm. Such a thickness dependent conductivity change is due to the different structures between the root and the surface. As will be discussed in the following, the root contains more SiC that has lower electrical conductivity, which leads to the lower overall electrical conductivity when the nanolaminate film is thin. When the thickness of the film increases, the amount of graphene along the growth direction increases, which causes the increase in the electrical conductivity. When the film reaches a certain thickness, equilibrium is reached and the amount of graphene keeps unchanged with increasing thickness. As a result, the conductivity of the film becomes constant.

Table

1

compares

the

conductivities

of

several

different

graphene/semiconductor composites fabricated using different strategies. Among them, the graphene/3C-SiC nanolaminate film possesses the highest conductivity, endowing it with the ability to be applied as electrodes. The high electrical conductivity, large surface area porous structure, intermediate band-gap of SiC, feasible integration on different substrates as well as the good chemical and mechanical stability of SiC and graphene make this nanolaminate structure very promising for a wide range of applications. Among them, the applications that have

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been already demonstrated to be beneficial using the combination of graphene and SiC, i.e., sensors9 and photocatalysis10, are foreseen.

Table 1. The electrical conductivity of the graphene/semiconductor composites fabricated using different methods.

Compositing material

Fabrication Method

Conductivity (S/cm)

Ref.

Fe3O4 nanoparticles

Wet chemistry

0.002

48

Si nanoparticles

Physical mixing in solution

13.1-33.1

8

SnO2 nanoparticles

Wet chemistry

0.023

49

Si3N4 microcrystals

Spark plasma sintering

40

50

CuO nanosheets

Wet chemistry

0.16

51

Al2O3 nanoparticles


59

52

TiO2 nanoparticles

Wet chemistry

0.2

13

ZnO/Zn(OH)2 nanoparticles

Wet chemistry

2.11

12

SiC nanoparticles


1.02

11

SiC-Graphene nanolaminate

Chemical vapor deposition

96.1

This work

Growth behavior After understanding the structure and properties of the graphene/3C-SiC hybrid nanolaminate, now our aim is to study its growth behavior to reveal its growth mechanism. TEM cross-sectional analysis was thus carried out and the results are shown in Figure 3. Figure 3a and b present the overall cross-sectional observation of the nanolaminates. All the nanolaminates are arranged in bunches. They grow from the Si substrate with a thin root. With their growth continuing, new layers emerge and slightly spread to the sides. The HRTEM observation in Figure S2 in Supporting Information shows that the root is composed by 3C-SiC nanosheet, whose basal plane 11



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is the (111) SiC plane, covered by graphene on both sides (graphene/SiC/graphene sandwich-like structure). Majority of the 3C-SiC nanosheets at the root is 9 – 22 nm in thickness (see the statistical results in the following). Since the root possesses a sheet-like morphology, the film appears porous even at the early growth stage.

Figure 3. TEM images of the graphene/3C-SiC hybrid nanolaminate deposited on (001) Si wafer. (a) and (b) the overall cross-sectional images, the region below the dashed line in (b) shows a preferential orientation; (c) high magnification TEM image at the Si/SiC interface, the dashed arrows indicate the existence of planar defects in SiC; (d) SAED of (c), the circles indicate the misorientation in SiC; (e) HRTEM image at the SiC/Si interface showing every five SiC (111) planes match with every four Si (111) planes.

For the whole cross-section, two regions, being separated by the dashed line, are clearly observable in Figure 3b. Above the line, there is no obvious preferential orientation. Below the line, however, most of the sheets lie at an angle of 55º to the substrate surface. This is more clearly observable in the high magnification TEM 12


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image of this region in Figure 3c. Fifty-five degree is a typical angle between the {111} and {001} facets in a face-centered-cubic (FCC) crystal structure. Keeping in mind that the Si substrate is (001) oriented, this implies the 3C-SiC nansheets are epitaxially grown on Si, as its basal plane is the (111) SiC plane. It is further confirmed by the corresponding SAED pattern of Figure 3c and the HRTEM observation at the SiC/Si interface, as shown in Figure 3d and e, respectively. From the HRTEM image in Figure 3e, a good epitaxial matching between 3C-SiC and the Si substrate is clearly observable with every five 3C-SiC (111) planes matching with every four Si (111) planes, confirming the epitaxial growth of 3C-SiC on Si substrate from the atomic scale. The SAED pattern in Figure 3d also demonstrates the epitaxial relationship between 3C-SiC and Si substrate with most of the diffraction points of 3C-SiC matching with those of Si substrate. Nevertheless, misorientation can be also observed, as marked by the circles in Figure 3d, which stems from the misoriented sheets, i.e., bending during the growth (it will be discussed in the following). The epitaxial growth of 3C-SiC on Si is not a surprising phenomenon because of the high similarities in the crystallographic structure between 3C-SiC and Si.36 A similar case has also been observed in our previous studies while using TMS for the growth of 3CSiC film as well as 3C-SiC nanosheets.37-38 The SAED pattern in Figure 3d also depicts weak streaks between the diffraction spots of 3C-SiC, indicating the existence of planar defects (i.e., stacking faults and twins) in 3C-SiC. The planar defects lie parallel to the {111} planes of 3CSiC, which is clearly distinguishable by the contrast differences in the SiC nanosheets in Figure 3c (marked by the dashed arrows). Their formation is a universal phenomenon during the growth of SiC based structures because of their low formation energy.35-38 Especially when 3C-SiC grows epitaxially on Si, it offers a mechanism to

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release the stress generated by the lattice mismatch between 3C-SiC and Si.36 In the present study, the existence of these planar defects is also an essential requirement to trigger the 2-dimensional growth of 3C-SiC nansheets by creating sub-steps at the side facet of the nanosheets. Its mechanism has been discussed thoroughly in the previous studies,38 and will be briefly discussed later in the section of “growth mechanism”. In addition to the diffraction pattern of 3C-SiC, a weak diffraction ring denoted as “G” also exists in the SAED pattern. As observed in Figure S2, these graphene sheets form at both sides of the 3C-SiC nanosheets at the root. From Figure 2 and Figure 3, the 3C-SiC layers at the root are observed to be significantly thicker than the individual 3C-SiC layers in the final hybrid nanolaminates. To quantify such differences, a statistical survey of the 3C-SiC layer thicknesses in both regions is done. Figure 4a and b show the histograms of the thickness of individual 3C-SiC layers at the root and in the final hybrid nanolaminate, respectively, as derived from the TEM observations. The histogram in Figure 4a is relatively wide and it reaches a maximum at a thickness of 14-16 nm. For the 3C-SiC layers in the final nanolaminate (Figure 4b), its thickness ranges from 4 nm to 10 nm and the peak lies at 5-6 nm. The thickness distribution is narrower with 90% of the 3C-SiC layers having a thickness of 4-8 nm. Therefore, the 3C-SiC layers at the root are almost twice as thick as those in the final nanolaminates. In contrast, the thickness of the graphene layers at the root doesn’t differ much from that in the final nanolaminate. Moreover, the distribution of the thickness of the graphene layer is also much narrower with 80% of the graphene layers having a thickness of 2-3.5 nm. In this context, the thickness change in the SiC layers represents the key to the growth of the nanolaminate structure.

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Figure 4. (a) and (b) Histograms of the thickness of the individual SiC layer (a) at the root and (b) in the final hybrid nanolaminates, as derived from the TEM images, (c) and (d) HRTEM images of two typical 3C-SiC layers that develop new layers. Insets show the magnified images of the indicated regions. The yellow arrows represent the nucleation and growth of graphene; the red arrows represent the occurrence of bending in the SiC layer; the white arrows denote the planar defects.

To understand how the thin 3C-SiC layers are developed from the thick ones, HRTEM observations of two typical 3C-SiC layers that just started developing thinner layers are presented in Figure 4c and d. The 3C-SiC (denoted as “SiC”) and graphene (denoted as “G”) layers are clearly distinguishable in both figures. The stripe contrast, an indication of planar defects, exists throughout the 3C-SiC layers. Graphene sheets are observed to nucleate around these planar defects at the central region of the growth front of the 3C-SiC layers, as denoted by the yellow arrows in the image. The nucleation of the graphene sheets leads to a slight distortion in the SiC lattice, as marked by the red arrows and shown in the inset of Figure 4c and d. Consequently, Half of the 3C-SiC layer starts to bend away from the other. With the growth process 15



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continuing, two new thin 3C-SiC layers separated by the graphene sheet are developed, forming SiC/graphene/SiC structure as the initial form of the hybrid nanolaminates. It looks as if the SiC layer is “split” into two parts by a “graphene wedge”. Obviously, the spacing between the two 3C-SiC layers constrains the thickness of the graphene sheets in the nanolaminate. Moreover, the bending of the SiC layer certainly leads to a change in the orientation of the SiC layers. It is obvious from Figure 3a and b that bending takes place continuously during the growth process, which results in an overall random orientation of the nanolaminates, as suggested by the SAED pattern in Figure 2a. Nevertheless, locally preferential orientation of the nanolaminates (Figure 1, some nanolaminates lie perpendicular to each other) is still kept due to the influence of initial epitaxial growth. After being “split”, the SiC layers start to grow thicker. This can be easily proven by adding up the thicknesses of all the individual 3C-SiC layers in one final nanolaminate; the sum is significantly larger than the thickness of the initial 3C-SiC layer at the root, indicating the thickening of the SiC layers during growth. Figure 5 shows the HRTEM images depicting the thickening process of two typical SiC layers. Two different ways are presented. In Figure 5a, the SiC layer thickens from 4.7 nm to 7.0 nm through a layer-by-layer expansion of its (111) plane. The red arrows in Figure 5a indicate several sites where the new 3C-SiC atomic planes form on the (111) plane. Around these sites, the graphene sheets just bend because of their good flexibility to follow the shape of the 3C-SiC layer. In Figure 5b, in addition to the layer-by-layer expansion of the (111) SiC plane, thickening by the continuous formation of dislocation is also observed. Two of the dislocations are marked by “T” in Figure 5b. Inserting the extra half-plane of atoms in the 3C-SiC lattice leads to the distortion of the nearby planes of atoms and increase of the 3C-SiC layer thickness. As a result, the

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thickness of the 3C-SiC layer quickly increases from 4.9 nm to 5.9 nm over a length of only ~14.5 nm after being “split”. These two ways of thickening are commonly observed during the growth of 3C-SiC layers in the present study, making the 3C-SiC layers thicker and thicker until they are again “split” into two.

Figure 5. (a) and (b) HRTEM images showing the thickening of the SiC layers during growth. The red arrows in (a) denote the insertion of new layers on the side of the SiC layer.

Growth mechanism From the results discussed above it is clear that the graphene/3C-SiC hybrid nanolaminates form via an iterative process: graphene induced “splitting” of the SiC layer and the thickening of the SiC layer after being “split”. The thickening and spitting of 3C-SiC layers control the thickness of 3C-SiC layer and the spacing between two 3C-SiC layers constrains the thickness of the graphene sheets. Now a natural question arises: what kind of mechanism governs the nucleation of graphene on the growth front of the 3C-SiC layer? Understanding this mechanism not only

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offers us the ability to control the thickness of the layers but also sheds us some light in obtaining other graphene-based nanolaminates. To find out the nucleation mechanism, it is necessary to understand the kinetics and thermodynamics of the current growth process. It is generally accepted that nucleation in vapor deposition follows a kinetically controlled process on the surface: adsorption and desorption of adatoms, diffusion of the adatoms, formation and growth of clusters.39 The clusters are formed when the adatoms combine with each other due to the fluctuation of the local adatom concentration and the formed clusters act as the sink for the adatoms in the neighborhood. Such a kinetic process controls the distribution of the nuclei. On the other hand, thermodynamics determines the growth mode if the process does not proceed too far from equilibrium:40 the atoms of the crystal will arrange in such a way that the total surface energy is minimized, i.e., the shape of the crystal is controlled by the thermodynamics during growth. In the present study, the main species contributing to the growth of SiC and graphene are CH3 and SiH3 based on our previous theoretical study.41 When the deposition starts, the carbon species not only react with the Si species on the substrate surface37 but also directly react with the pristine Si substrate.42 Both reactions lead to the formation of 3C-SiC at the present deposition conditions. As a result, a thin (~20 nm) but dense epitaxial 3C-SiC layer firstly forms on the Si substrate as shown in Figure 3c. Owing to the large lattice mismatch between 3C-SiC and Si, this initial epitaxial layer contains a large amount of planar defects. Our previous investigation has shown that, these planar defects accelerate the 2-dimensional growth of 3C-SiC by forming sub-steps, which initializes the growth of 3C-SiC nanosheets.38 Here, thermodynamics controls the basal plane of the 3C-SiC nanosheets to be the (111) plane to lower the total free energy, as the close-packed plane has the lowest surface

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energy.38 Such a 2-dimensional growth of 3C-SiC leads to the exposure of a large area of (111) plane to the reactive plasma. Keeping in mind that the {111} 3C-SiC planes readily catalyze the growth of graphene,43 graphene is formed on both sides and graphene/SiC/graphene sandwich structure is obtained as the root (see Figure S2). When the Si substrate is fully covered by 3C-SiC, it no longer takes part in the reaction with the incoming carbon species. The growth of 3C-SiC is only possible by the combination of carbon adatoms with Si adatoms on the growth front of 3C-SiC. Nevertheless, the amount of carbon reactive species in the gas phase is far higher than that of Si reactive species, because the amount of methane is 65 times larger than that of TMS in the gas phase. This results in a large amount of carbon adatoms on the growth front of 3C-SiC. Of course, they can combine with the Si adatoms for the growth of 3C-SiC. However, because of the low concentration of the Si reactive species, most of the carbon adatoms remain unreacted and diffuse on the surface. These carbon adatoms can be viewed as “supersaturated” carbon adatoms. According to the classical vapor phase growth model, when these carbon adatoms diffuse to the graphene sheets at the sides of the SiC layer, they will be captured by the graphene sheets and contribute to their growth.39 In other words, the graphene layer acts as the sink for the diffusing carbon adatoms in their neighborhood.39 Figure 6a shows the schematic diagram concerning the concentration of these carbon adatoms on the growth front of the 3C-SiC layer. The concentration of the carbon adatoms decreases to zero at the border of the graphene layer; it increases with increasing distance from the graphene layer and reaches its maximum in the center region between two graphene layers.40,

44

Here, the maximum concentration of the carbon adatoms is

dependent on the of the 3C-SiC layer: the thicker 3C-SiC layer, the higher will be the maximum concentration of these carbon adatoms on its surface.44 Once the maximum

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concentration of the carbon adatoms exceeds the critical concentration for the graphene formation, nucleation of graphene starts. The upper cut-off size of 10 nm for the individual SiC layer in the final nanolaminate (see Figure 4b) suggests that the critical concentration is most likely reached when the thickness of the SiC layer reaches 10 nm. It is noteworthy that the growth front of the SiC layer contains a large number of sub-steps created by the planar defects. Because of their lower potential well, these sub-steps are the preferential sites for the nucleation of graphene to take place.45 As a result, graphene nucleates around the planar defects in the center region of the SiC layer, leading to the separation and bending of the SiC layers. Two new thinner 3C-SiC layers are thus formed. The above process is schematically summarized in Figure 6b for a better illustration.

Figure 6. (a) Schematic diagram of “supersaturated” carbon adatom concentration on the growth front of the SiC layer. (b) Schematic illustration of the initial nucleation of 20


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graphene and bending of the SiC layer: (I). The growth front of the SiC layer contains a large amount of sub-steps created by the planar defects. Carbon and Si adatoms are adsorbed on the growth front of the SiC layer. The concentration of carbon adatoms is higher than that of Si adatoms. Some carbon adatoms are possible to combine with Si adatoms to form SiC. However, due to the high concentration of carbon, most of the carbon adatoms are unreacted and diffuse as “supersaturated” carbon adtoms on the surface of the growth front of SiC layer. Once their concentration reaches the critical concentration, graphene starts to nucleate at the sub-steps created by the planar defects at the central region of the SiC layer (indicated by the white arrows). (II). The formation of graphene leads to the distortion in the lattice of the SiC layer. (III). The SiC layer bends and separates with each other because of the lattice distortion. (IV). The growth of SiC and graphene continues and SiC/graphene/SiC structure forms as the initial form of graphene/SiC hybrid nanolaminate.

The new graphene nuclei again act as the sink for the carbon adatoms. Therefore, the concentration of carbon adatoms on the two new thinner 3C-SiC layers is reduced to below the critical concentration. Consequently, no further nucleation of graphene occurs on the newly formed 3C-SiC layers. The new 3C-SiC layers and graphene sheets thus grow simultaneously. It is noteworthy that, the newly developed 3C-SiC layers tend to grow thicker, as shown in Figure 5. This is thermodynamically favored. The formation of laminate structure costs extra energy; this energy are “stored” as the interface energy at graphene/SiC interfaces. For a unit volume, the interface energy ∆G can be expressed as: ∆G =

2γ AB

λ

, where γAB is the specific

surface tension of graphene/SiC interface and λ is the spacing between the layers.46 In this context, the thicker the SiC and graphene layers, the fewer interfaces exist in the 21



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unit volume, which leads to fewer interfacial areas and lowers the total system energy. As a result, the SiC and graphene layers tend to grow thicker to reduce the total surface energy of the whole system. However, once the 3C-SiC layers reach a thickness larger than 10 nm, the nucleation of graphene takes place again on the surface because of the increased concentration of carbon adatom, “splitting” SiC layers into two. The spacing between the two SiC layers also constrains the maximum thickness of the graphene layer. The processes of thickening and “splitting” take place in turn, and graphene/3C-SiC hybrid nanolaminates finally form. However, a local fluctuation in the surface energy and the supply of reactive species also exists during the growth process, which is commonly observed during the CVD growth of nanostructures.47 This leads to the formation of certain small-sized thin nanolaminates, which can be also observed in Figure 1c. Nevertheless, the amount of such thin nanolaminates is very small, which has no effect on the conclusions we draw in the present study. The growth mechanism proposed above suggests that the concentration of carbon adatoms plays an important role in controlling the layer thickness of the graphene/3C-SiC hybrid nanolaminate. It is known that the carbon adatom concentration is also dependent on the supplied methane concentration. This allows us to control the layer thickness by manipulating the gas phase composition: decreasing the methane concentration can decrease the concentration of carbon adatoms on 3CSiC; a thicker SiC layer is thus required to achieve the critical carbon adatom concentration for the nucleation of graphene, leading to the thicker individual 3C-SiC layer in the final nanolaminates. We confirmed this by depositing the nanolaminate structure using different methane concentrations. Three different CH4 flow rates are used: 3 sccm, 6 sccm and 15 sccm, as shown in Figure 7. Hybrid nanolaminates were

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obtained in all cases. With decreasing methane concentration the thickness of the 3CSiC layer increases significantly. As shown Figure 7d, majority of SiC layers deposited at CH4 flow rate of 3 sccm has a thickness of 17-32 nm, while majority of the SiC layers deposited at CH4 flow rate of 15 sccm is only 2-5 nm in thickness. In addition to the increased thickness of 3C-SiC layers at low CH4 flow rate, a decrease in the thickness of the graphene layer is also observable. This is easy to understand: the growth of thicker graphene layers requires more carbon species; a lower CH4 concentration thus leads to the thinner graphene layers. The phenomenon shown above not only proves the validity of the proposed growth mechanism, but also offers the possibility to achieve a good control over the layer thickness of the graphene/3CSiC hybrid nanolaminate in the present approach. Since the electrical conductivity of the nanolaminate film is mainly conveyed by the graphene sheets in the structure, the decrease in the graphene content in the nanolaminate should, in principle, result in a decrease in its electrical conductivity. This is confirmed by measuring the electrical conductivities of the above samples, which are 2.8, 21.3 and 101.6 S/cm for the nanolaminate film deposited at methane flow rates of 3, 6 and 15 sccm, respectively, showing an increase in the conductivity with the increasing graphene amount.

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Figure 7. (a) – (c) TEM images of the SiC layer in the hybrid nanolaminates deposited using different methane concentrations while keeping the other parameters constant. (a) CH4 = 3 sccm; (b) CH4 = 6 sccm; (c) CH4 = 15 sccm. (d) Statistical survey of the thickness of the SiC layer in the samples deposited with different methane concentrations.

Conclusions In conclusion, graphene/3C-SiC hybrid nanolaminates with high electrical conductivity are obtained using microwave plasma chemical vapor deposition technique. The thicknesses of the individual layers are 4-10 nm for 3C-SiC and 2-5 nm for graphene. Dedicated HRTEM analysis reveals that the growth of the nanolaminates follows an iterative process: the preferential nucleation of graphene at the planar defects of 3C-SiC, leading to the “splitting” of the 3C-SiC layer; and the thickening of the “split” 3C-SiC layers. Both kinetics and thermodynamics are believed to govern the above processes. Based on the proposed mechanism, controlling the layer thickness in the nanolaminate is predicted to be possible by controlling the carbon concentration in the gas phase, which is further confirmed experimentally. The approach demonstrated in the present study is compatible with a wide range of substrates, such as metal (Mo), semiconductor (Si and 2H-SiC) and insulator (diamond). It not only enables the universal integration of the graphene/3CSiC nanolaminate structure, but also sheds some light for the fabrication of other graphene based nanolaminate structure by means of chemical vapor deposition.

Acknowledgement

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The authors thank Deutsche Forschungsgemeinschaft (DFG JI 22/24-1) for the financial support. B. Yang thanks National Nature Science Foundation of China (Grants No. 51402309) for financial support.

Supporting Information Available: The SEM images of the nanolaminates obtained on other substrates and the TEM image of the root of the nanolaminate are shown in Supporting Information. This material is available free of charge via the Internet at http://pubs.acs.org.

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TOC Figure

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Figure 1. (a) – (d) SEM images of the typical graphene/3C-SiC hybrid nanolaminate grown on the Si substrate. (a) Overall observation; (b) and (c) surface view; (d) cross-sectional view, the arrow indicates the observation of laminated structure; (e) an optical photo of a flexible freestanding graphene/3C-SiC hybrid nanolaminate film; (f) Raman spectra. 67x31mm (300 x 300 DPI)



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Figure 2. TEM images of the graphene/3C-SiC hybrid nanolaminates from the surface view. (a) Overall morphology, inset shows the corresponding SAED pattern; (b) and (c) higher magnification TEM image of one individual nanolaminate; (d) HRTEM image of nanolaminate, the inset shows the corresponding FFT pattern; (e) IFFT re-construction of (d), the red color indicates 3C-SiC and the blue color indicates the graphene sheets; (f) HRTEM image of another hybrid nanolaminate. 88x55mm (300 x 300 DPI)


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Figure 3. TEM images of the graphene/3C-SiC hybrid nanolaminate deposited on (001) Si wafer. (a) and (b) the overall cross-sectional images, the region below the dashed line in (b) shows a preferential orientation; (c) high magnification TEM image at the Si/SiC interface, the dashed arrows indicate the existence of planar defects in SiC; (d) SAED of (c), the circles indicate the misorientation in SiC; (e) HRTEM image at the SiC/Si interface showing every five SiC (111) planes match with every four Si (111) planes. 85x58mm (300 x 300 DPI)



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Figure 4. (a) and (b) Histograms of the thickness of the individual SiC layer (a) at the root and (b) in the final hybrid nanolaminates, as derived from the TEM images, (c) and (d) HRTEM images of two typical 3CSiC layers that develop new layers. Insets show the magnified images of the indicated regions. The yellow arrows represent the nucleation and growth of graphene; the red arrows represent the occurrence of bending in the SiC layer; the white arrows denote the planar defects. 78x64mm (300 x 300 DPI)


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Figure 5. (a) and (b) HRTEM images showing the thickening of the SiC layers during growth. The red arrows in (a) denote the insertion of new layers on the side of the SiC layer. 80x68mm (300 x 300 DPI)



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Figure 6. (a) Schematic diagram of “supersaturated” carbon adatom concentration on the growth front of the SiC layer. (b) Schematic illustration of the initial nucleation of graphene and bending of the SiC layer: (I). The growth front of the SiC layer contains a large amount of sub-steps created by the planar defects. Carbon and Si adatoms are adsorbed on the growth front of the SiC layer. The concentration of carbon adatoms is higher than that of Si adatoms. Some carbon adatoms are possible to combine with Si adatoms to form SiC. However, due to the high concentration of carbon, most of the carbon adatoms are unreacted and diffuse as “supersaturated” carbon adtoms on the surface of the growth front of SiC layer. Once their concentration reaches the critical concentration, graphene starts to nucleate at the sub-steps created by the planar defects at the central region of the SiC layer (indicated by the white arrows). (II). The formation of graphene leads to the distortion in the lattice of the SiC layer. (III). The SiC layer bends and separates with each other because of the lattice distortion. (IV). The growth of SiC and graphene continues and SiC/graphene/SiC structure forms as the initial form of graphene/SiC hybrid nanolaminate. 117x165mm (300 x 300 DPI)


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Figure 7. (a) – (c) TEM images of the SiC layer in the hybrid nanolaminates deposited using different methane concentrations while keeping the other parameters constant. (a) CH4 = 3 sccm; (b) CH4 = 6 sccm; (c) CH4 = 15 sccm. (d) Statistical survey of the thickness of the SiC layer in the samples deposited with different methane concentrations. 77x71mm (300 x 300 DPI)


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TOC 39x17mm (300 x 300 DPI)


3C-SiC Hybrid Nanolaminate - ACS Applied Materials

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