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Large-scale chemical vapor deposition of graphene on polycrystalline nickel films: Effect of annealing conditions
Thin Solid Films 690 (2019) 137565
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Large-scale chemical vapor deposition of graphene on polycrystalline nickel films: Effect of annealing conditions
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Fatima Akhtara, , Jaroslaw Dabrowskia, Marco Liskera, Peter Zaumseila, Sebastian Schulzea, Alex Jouvrayb, Piotr Cabanc, Andreas Maia,d, Christian Wengera,e, Mindaugas Lukosiusa IHP – Leibniz-Institut für Innovative Mikroelektronik, Im Technologiepark 25, 15236 Frankfurt (Oder), Germany Aixtron Ltd., Anderson Road, Swavesey, CB24 4FQ Cambridge, UK c Institute of Electronic Materials Technology, Wolczynska 133, 01-919 Warsaw, Poland d Technical University of Applied Science Wildau, Hochschulring 1, 15745 Wildau, Germany e Brandenburg Medical School Theodor Fontane, 16816 Neuruppin, Germany a
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A R T I C LE I N FO
A B S T R A C T
Keywords: Polycrystalline nickel Surface pretreatment Grain growth Grain boundary grooves Chemical vapor deposition Graphene
In the present study, 8-in. silicon substrates, covered with thin (200 nm) polycrystalline nickel films have been employed for the growth of graphene by chemical vapor deposition. In order to control the uniformity and coverage of the graphene, thin nickel layers were used due to their less deep grain boundary grooves and ability to store less carbon in comparison with thick nickel films (> 500 nm). The preferential sites for the growth of multilayer graphene were influenced by the surface pretreatment of the polycrystalline nickel films at 1025 °C under different ambient conditions (hydrogen and vacuum). Significant differences in the surface morphologies were observed for the annealed nickel films. The growth of larger grains up to ~6 μm for the films annealed in hydrogen could be attributed to hydrogen interstitials. On the other hand, grains up to ~3 μm were extracted for the films annealed in vacuum. Graphene was grown after exposing the annealed Ni films to ethylene at 925 °C. The lower range (42–106 cm−1) of full width at half maxima of the 2D band as determined by Raman spectroscopy was obtained for the films annealed in hydrogen as compared to the ones annealed in vacuum (51–128 cm−1), indicating that the thickness uniformity of graphene was strongly influenced by the surface modifications of nickel films.
1. Introduction Due to its properties, graphene has appeared as a potential candidate in various electronic devices. However, one of the biggest challenge in its integration into microelectronics is the lack of high-quality and wafer-scale synthesis of the graphene. Typically, chemical vapor deposition (CVD) is used for the growth of graphene on metal catalysts [1–9]. Indeed, copper (Cu) catalyst has been employed for large-area graphene growth [9,10], however, a large lattice mismatch between graphene and Cu [11], the rotated graphene domains [12] and the presence of metallic contaminants after a complicated transfer step can degrade the quality of graphene [13]. In contrast to Cu, the growth of graphene on Nickel (Ni) occurs by a segregation phenomenon, this is attributed to the high carbon solubility (~2.03 at. % at 1000 °C) in Ni [14,15]. Due to the excellent lattice match between graphene and Ni, single crystal Ni (111) was reported to be one of the most suitable substrate for the growth of monolayer graphene [14,16–18].
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Nevertheless, high costs and limited sizes are the main drawbacks of single crystal substrates. At this point, polycrystalline Ni films on SiO2/ Si (001) are attractive for the growth of graphene due to their largescale availability and compatibility with Si microelectronics. However, the polycrystallinity of Ni makes the graphene growth a complicated process. Grain boundaries (GBs) in polycrystalline Ni are the most preferential sites for the growth of multilayer graphene, whereas Ni (111) grains favor the growth of monolayer graphene. Therefore, in order to crystallize Ni along the most preferred (111) orientations and to reduce the number of GBs, typically high temperatures (~1100C) [7] and thick Ni films (500 nm) have been used [6–8,15,18]. Thick Ni films were usually employed in order to avoid the dewetting phenomenon at higher annealing temperatures, however, thick Ni films suffer from several drawbacks, for instance, the presence of deeper GBs grooves, which could increase the surface roughness and could be the cause of higher coverage of multilayer graphene [18]. Another disadvantage of thick Ni films is related to their higher carbon storage efficiency that
Corresponding author. E-mail address: [email protected] (F. Akhtar).
https://doi.org/10.1016/j.tsf.2019.137565 Received 18 December 2018; Received in revised form 15 August 2019; Accepted 11 September 2019 Available online 12 September 2019 0040-6090/ © 2019 Elsevier B.V. All rights reserved.
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Renishaw Raman spectrometer with 514 nm excitation wavelength of the laser and beam spot size ≈ 1 μm. X-ray diffraction (XRD; SmartLab diffractometer from Rigaku) was used to determine the crystal structure using CuKα (λ = 1.54186 Å) radiation. The XRD measurements were performed in a parallel beam configuration. Surface morphology of Ni films was investigated by scanning electron microscopy (SEM; Gemini 2, Zeiss) at 1.5 kV. The grains distribution in Ni films was studied by FEI Nova 600 Nanolab equipped with an electron backscattering diffraction (EBSD) detector at 30 keV with 9.5 nA beam current. Atomic force microscopy (AFM; Dimension 5000 SPM System with NanoScope IV Controller from Digital Instruments) in tapping mode was used to monitor the surface roughness of Ni substrates. The elemental composition of as-deposited and annealed samples was determined by x-ray photoelectron spectroscopy (XPS) using a monochromatic MgKα (1253.6 eV) source under 7 × 10−8 Pa with a hemispherical analyzer (PHOIBOS-100). The beam energy and beam current density used for XPS experiments were 300 W and 15 KV, respectively. Moreover, the binding energy was calibrated to gold (Au) 4f line set to 83.95 eV. XPS spectra were analyzed with the CasaXPS software. A Shirley back ground was used in the curve-fitting procedure along with a asymmetric line shapes (LF) for Ni2p spectra. However, Gauss-Lorentzian (GL) line shapes were used to fit the C1s and O1s spectra.
makes the thickness control of the graphene challenging [11,18]. Indeed, as reported in literature, the coverage of mono-bilayer graphene on polycrystalline Ni films varied in the range from 70% to 87%, indicating the complexity of this system [6,11,18,19]. Therefore, the purpose of this work is to employ thin (200 nm) Ni films deposited on 8in. Si wafers and to prepare its surface for the growth of graphene by modifying its grain sizes, orientations, and GBs by different pretreatment. One of the advantages of thin Ni films as compared to the thick films is the presence of less deep GBs grooves, which can play a significant role in reducing the growth of multilayer graphene. 200 nm Ni films were annealed at 1025 °C under different ambient conditions (hydrogen/vacuum), keeping other parameters constant to examine their effect on the surface roughness, grain sizes, and crystallinity of the annealed Ni films. After annealing, these Ni films were exposed to ethylene at 925 °C for the growth of graphene, and the quality of the obtained graphene films was then compared. The findings of this study revealed that annealing in hydrogen led to a substantial increase in the sizes of Ni grains as compared to the annealing in vacuum. This reflects a higher coverage of a thinner graphene on the Ni films annealed in hydrogen, as measured by Raman, indicating the influence of grain sizes on the thickness of graphene. 2. Experimental details
3. Results & discussions 2.1. Deposition, annealing of Ni/SiO2/Si, and growth of graphene The crystal quality and the surface morphology of the as-deposited and the annealed Ni films were investigated by XRD, SEM, AFM, and EBSD. The diffraction patterns were recorded by specular ω–2θ scans in the angular range of 2θ = 40°–55° using for as-deposited and annealed Ni films (hydrogen/vacuum) as shown in Fig. 2a. As-deposited Ni was polycrystalline as the reflections at 44.52°, assigned to Ni (111) orientation, and a weak Ni (200) reflection at 51.87° corresponding to Ni crystallites with (100) orientations, were measured. The Ni (111) orientation has the lowest surface/interface energy with respect to other crystallographic orientation, which is a driving force for their growth. XRD curves of the Ni films annealed in hydrogen (black line) and vacuum (red line) at 1025 °C are also presented in Fig. 2a. A slight shift of Ni (111) and Ni (200) reflections to higher angles is observed after annealing in hydrogen and vacuum. These shifts could be attributed to the different thermal expansion coefficients between Ni (14 × 10−6 K−1) and the underlying Si substrate (4 × 10−6 K−1) [7,20]. A detailed analysis of the Ni (111) reflections was performed by conducting ω scans to determine the mosaic spread in the films appearing due to randomly oriented crystallites. In these experiments, the detector was fixed at the center of Ni (111) diffraction peak, and the sample was tilted. The full width at half maxima (FWHM) of the obtained curve is an objective measure for the mosaicity of the Ni films. The values of FWHM are plotted as a function of annealing temperatures and are depicted in the inset of Fig. 2a. After annealing at 1025 °C, the values of FWHM decreased from 9.63° (as-deposited Ni) to 5.43° (annealed in hydrogen) and to 5.89° (annealed in the vacuum). This reduction indicates that the mosaicity of Ni (111) planes is lower in both annealing conditions. However, the lower values of FWHM of Ni (111) reflection for the films annealed in hydrogen demonstrate that the alignment of Ni grains is enhanced compared to films annealed in vacuum. SEM images taken with an HE-SE2 detector (high-efficiency secondary electron detector; used to investigate surface topography) are shown in Fig. 2b–c for the films annealed in hydrogen and vacuum, respectively. Firstly, inhomogeneity in Ni grain size and shape can be noticed in the SEM images. This inhomogeneity is the result of a grain growth phenomenon in which lateral grain size exceeds the film thickness. During the grain growth mechanism, faster-growing grains dominate the slower ones, and this process continues until all un-favored grains diminish leading to a texture development in thin films [21,22]. Secondly, surface roughness was increased due to the grain
Ni films were deposited from a Ni source by DC magnetron sputtering method on 100 nm thick thermal SiO2 on 8-in. Si (001) substrate. The depositions were carried out at room temperature in ~55 sccm argon. During the deposition process, plasma was ignited at 750 W and the pressure was kept at 0.533 Pa. The growth of graphene on polycrystalline Ni is a complicated process that can be divided into four steps as shown in Fig. 1. In the first step, the Ni grain growth was performed by annealing the Ni films for 10 min in a CVD chamber at 1025 °C (label a). Annealing was performed in different ambient conditions, hydrogen (P = 10 Pa) or vacuum (P = 1.08 × 10−2 Pa). After annealing, the temperature was reduced to 925 °C (label b), where hydrocarbons (C2H4) were let into the chamber and the pressure in the chamber was kept at 10 Pa (label c). After 5 min, the supply of the hydrocarbons was stopped and the system was allowed to cool down. As a result, carbon atoms started to precipitate and after crystallization on the Ni surface formed a graphene sheet (label d). 2.2. Characterization Quality of the deposited graphene layers was monitored by a
Fig. 1. Illustration of the temperature-time profile of graphene growth process by CVD. 2
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Fig. 2. XRD ω–2θ scans of (a) as-deposited Ni film (blue line), Ni films annealed in hydrogen (black line) and vacuum (red line). The intensity is normalized in a logarithmic scale. FWHM of Ni (111) reflection obtained from ω scans as a function of annealing temperature is shown in the inset of Fig. 2a. SEM images of Ni films annealed in (b) hydrogen and (c) vacuum. AFM images of (d) as-deposited and Ni films annealed in (e) hydrogen and (f) vacuum. SEM image of as-deposited Ni film is shown in the inset of Fig. 2d. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
hydrogen. This is in agreement with the literature report where wide and flat Ni (111) nanostructures have been observed due to the dissolution of hydrogen [25]. According to another study, the presence of interstitial hydrogen causes additional stresses in the Ni films that could result in the volume expansion of the Ni grains [20]. On the contrary, Ni films annealed in vacuum were not under the influence of hydrogen and the grain growth was affected by abnormal grain growth mechanism and grain sizes up to 3 μm were recorded. In comparison with the reports in the literature, different grain sizes were obtained for different thicknesses of Ni films. For 500 to 550 nm thick Ni films, the grain sizes varied from 1 to 20 μm [6–8,27,28], whereas even larger grains (≥50 μm) were obtained for Ni foils (25 μm) [29]. Although larger grains were reported for thick Ni, the disadvantages of thick Ni were previously discussed. For the thinner Ni films (≤300 nm), the grain sizes were found to be much smaller as compared to the thicker films. The grain sizes were reported within the range of 1 to 4 μm upon reducing the thickness of the Ni films to 300 nm [30–33]. Literature reports on the grain sizes for thin ≤200 nm Ni films (as used in this work) are rather limited and were in the range from 150 nm to 1 μm [34–36]. The authors in the work [35,36] attribute the smaller grains to the simultaneous growth of Ni grains and carbon diffusion during the annealing step; these smaller grains resulted in inhomogeneous growth of graphene, indicating the importance of preannealed Ni films. In our study, the crystallization of Ni films in different annealing conditions before graphene growth resulted in larger grains up to 6 μm.
growth, as determined by AFM. The root mean square surface roughness (Rq) was recorded by scanning 5 × 5 μm2 areas of the as-deposited and the annealed Ni films as shown in Fig. 2d–f. The Rq values increased from 1.6 nm (as-deposited Ni) to ~28 nm for the films annealed in hydrogen and ~10 nm for the ones annealed in vacuum. As compared to literature, the roughness of the films annealed in hydrogen was lowered than the reported values of 36.3 nm obtained after 15 min annealing (900 °C) of 500 nm Ni under 600 sccm H2 pressure [18]. Lower surface roughness in this study could be attributed to less deep GBs grooves of thin films as compared to thick films. The grain sizes of the as-deposited Ni films were around 20 nm (inset of Fig. 2d) and they increased to several microns depending on the annealing conditions. Indeed, as extracted from 30 × 80 μm2 EBSD maps in Fig. 3a–b, grains up to ~6.0 μm diameter were measured within the films annealed in hydrogen (Fig. 3c squares) and up to ~3.0 μm were observed within the films annealed in the vacuum (Fig. 3d circles). The grain sizes below 0.6 μm could not be adequately indexed and have been removed [23]. According to EBSD and SEM results, larger grains were recorded for the films annealed in hydrogen in compare to the films annealed in vacuum. The difference in Ni grain sizes as a function of different annealing ambient could be attributed to the presence of hydrogen. Hydrogen is known to incorporate in the interstitial sites in a host metal lattice during annealing [20,24–26]. This may generate stresses in the Ni films and thus a volume expansion of the Ni crystallites. Therefore, we assume that larger Ni grains (6 μm) for the films annealed in hydrogen could be attributed to the stress induced by interstitial 3
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Fig. 3. EBSD analysis: colour map, Orientation maps for Ni films annealed in (a) hydrogen (b) vacuum. Grain distributions measured by EBSD for Ni films annealed at 1025 °C in (c) hydrogen (squares) (d) vacuum (circles) and after graphene growth at 925 °C (triangles).
Fig. 4. XPS analysis: (a) O1s (b) Ni2p and (c) C1s spectra of Ni films as-deposited (black line), annealed at 1025 °C in hydrogen (magenta line) and vacuum (blue line) and after graphene growth at 925 °C (green line). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Fig. 5. EBSD images of Ni films annealed in (a) hydrogen (b) vacuum after graphene growth. SEM images of graphene grown on Ni films annealed in (c) hydrogen (d) vacuum.
Before starting the growth of graphene, the Ni films were investigated by XPS. Fig. 4 summarizes the XPS data that provides complementary information about the chemical state of the surface of the Ni-films. The O1s spectrum of the as-deposited Ni (black line) as presented in Fig. 4a is fitted to two main components at 529.6 eV (FWHM = 1.4 eV) and 531.7 eV (FWHM = 2.0 eV), which are assigned to NiO and Ni2O3, respectively [37]. In Fig. 4b, the Ni2p spectrum of the as-deposited Ni (black line) is composed of Ni° (852.6 eV), Ni+2 (854.1 eV) and Ni+3 (856.1 eV) corresponding to metallic Ni (FWHM = 1.5 eV), NiO (FWHM = 2.0 eV) and Ni2O3 (FWHM = 2.4 eV), respectively. Some satellite (shakeup) peaks at a few eV above the main line are also visible in the spectrum; the satellite peaks of Ni oxides are located at higher binding energy than that of the metallic Ni. They appear as the result of a change in the Coulomb potential of electrons due to the removal of core electron; their orbitals relax and liberate energy which is utilized to shake up an outer electron to a higher d-orbital [38]. Moreover, the peaks in the C1s spectrum (black line) of the as-deposited Ni as demonstrated in Fig. 4c are attributed to CeH (285.0 eV), CeO (286.5 eV), and C]O (288.9 eV). The FWHM of the main line contributions are 1.5 eV, 1.5 eV and 1.7 eV, respectively. These XPS results demonstrated that the as-deposited Ni films were contaminated when they were exposed to air; however, these contaminants have been removed after annealing in hydrogen (magenta lines) and vacuum (blue lines) as shown in Fig. 4a–c. In order to grow graphene, the Ni films annealed in hydrogen and in vacuum were exposed to ethylene (Partial pressure = 5 Pa) and hydrogen (Partial pressure = 5 Pa) for 5 min at 925 °C and XPS analysis was applied again. In Fig. 4c, the evolution of an sp2 bonded carbon signal at ~284.6 eV (FWHM = 1.1 eV) in C1s core level spectrum (green line) implies the formation of graphene. The obtained value of
the binding energy in this work is comparable to the reported values, where authors described that the binding energy of graphene on polycrystalline Ni lies in the range of 284.3 to 284.5 eV [11,27,32,39]. In contrast to the polycrystalline Ni films, the binding energy of graphene on single crystal Ni (111) has been reported to be within the range of 284.8 to 285.0 eV [40,41]. It could be attributed to the stronger interaction between graphene and single crystal Ni (111) as compared to the graphene/polycrystalline Ni system [41]. Furthermore, no oxide signals were recorded in the O1s and Ni2p spectra, as shown in Fig. 4a–b (green lines). It is worth mentioning here that in-situ XPS measurements have been performed in this study (both after annealing the Ni films and after the growth of graphene) due to interconnected XPS and CVD chambers. The grain distributions after graphene growth were again investigated by EBSD and are demonstrated in Fig. 3c–d (triangles). The orientation maps of Ni films collected after the graphene growth are shown in Fig. 5a–b. By comparing the orientation maps (Fig. 3a–b) and the grain distributions in Fig. 3c–d for the Ni films annealed in hydrogen (squares) and vacuum (circles) with the orientation maps (Fig. 5a–b) and grain distributions (Fig. 3c–d triangles) for the films obtained after the graphene growth, one can notice that they remained the same. Indeed, the graphene growth did not influence the grain sizes within the Ni films. In other words, the annealing procedure carried out with a higher thermal budget than the one during graphene growth forced the Ni film morphology to a structural state which was stable under the graphene growth conditions. SEM is a powerful tool to identify different numbers of graphene layers. The surface morphology of graphene was investigated by SEM; images taken with the in-lens detector are shown in Fig. 5c–d. This detector was used since it generates different colour
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Fig. 6. (a) Raman spectra of graphene grown on Ni films annealed in (a) hydrogen (b) vacuum; different colours corresponding to spectra taken from different areas. Raman maps (2D-FWHM) of graphene grown on Ni films annealed in (c) hydrogen (d) vacuum.
graphene can be grown in both conditions. The FWHM of the 2D band is presented in Fig. 6c–d. The range of the FWHM values lied in the range 42–106 cm−1 for hydrogen annealed Ni films and 51–128 cm−1 for the films annealed in vacuum. The lower values of FWHM indicated the presence of thinner graphene on hydrogen annealed films [18,37,47]. Moreover, the I2D/IG ratio ~2.4 (label 1) in Fig. 6c suggested the growth of mono–bilayer graphene [32,37]. The I2D to IG ratio dropped to ~0.6 (label 2) and ~0.5 (label 3) in Fig. 6c, indicating the presence of few layers (3–5) and multilayer graphene [32,37], respectively. In the case of the films annealed in vacuum, the I2D/IG ratio ~1.7 (label 1), ~0.5 (label 2) and ~0.4 (label 3) in Fig. 6d revealed the growth of thicker graphene. To summarize, Raman analysis demonstrated the growth of thinner graphene on the films annealed in hydrogen as compared to the ones annealed in vacuum. The higher coverage of thinner graphene on the films annealed in hydrogen was attributed to larger Ni grains and reduced grain boundaries. On the contrary, smaller grains and therefore increased GBs for the films annealed in vacuum facilitated the growth of multilayer graphene.
contrasts of a surface according to its work functions. The charge transfer from metal to the layers of graphene is different, which results in different work function values and therefore the colors in the SEM image. The non-uniformity of the graphene coverage is clearly visible in SEM images. The surface of the scanned area was divided into three distinct areas based on their work functions. Herein, the light grey or brighter areas corresponded to the surfaces with low work function or thinner graphene (label 1). Upon increasing the number of grown graphene layers, their area became darker (labels 2 and 3) [32,42,43]. Finally, the quality of graphene was monitored by Raman spectroscopy. A Raman spectrum of graphene consists of three different peaks known as G peak (~1580 cm−1), D peak (~1350 cm−1) and 2D peak (~2700 cm−1) [44,45]. The G peak is the result of in-plane CeC stretching mode and corresponds to doubly degenerate E2g phonons at the Brillouin zone center (Γ point). The D peak or defect mode is the result of a double resonance process and requires a TO phonon and a defect in its activation. Whereas the 2D peak, which is also known as a graphene fingerprint, is the overtone of the D peak and is caused by a second order two-phonon Raman process involving two TO phonons near the K point [44–46]. The I2D/IG and FWHM of the 2D band provide the basis for determining the number of graphene layers. The height of the 2D band peak corresponding to monolayered graphene is larger than the G peak and has a symmetric line shape that can be fitted with a single Lorentzian peak. The height and shape of the 2D band peak changes with increasing thickness of the graphene film, for instance, thicker graphene demonstrated a broadened 2D peak with reduced amplitude [45]. As shown in Fig. 6a–b, Raman spectra were monitored at three different regions depending on the thickness of the grown graphene films. The presence of small defect peaks indicates that a good quality
4. Conclusions Two different pretreatment conditions (hydrogen and vacuum) were employed to anneal 200 nm thick polycrystalline nickel films deposited on 8-in. 100 nm SiO2/Si (001) substrates. Both annealing conditions favored the crystallization of Ni films that resulted in a strong {111} texture; however, omega scans of Ni (111) reflections indicated that the mosaicity of the grains was reduced in hydrogen annealed Ni films. According to the AFM analysis, the surface roughness of Rq ~28 nm was recorded for hydrogen annealed films, whereas for vacuum annealed films it was found to be Rq ~10 nm. EBSD measurements 6
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revealed larger grains for the films annealed in hydrogen. Hydrogen ions incorporate into the interstitial sites of Ni films thus enabling the growth of large Ni grains. As indicated by Raman spectroscopy, the reduced bandwidth of the 2D peak of the graphene grown on the Ni films annealed in hydrogen indicated the relevance of grain size for the growth of thinner graphene. However, the surface of Ni films annealed in hydrogen was rougher. The smaller Ni grains for the films annealed in vacuum resulted in increased number of grain boundaries that facilitated the growth of multilayer graphene.
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Declaration of Competing Interest None. Acknowledgments The authors are pleased to acknowledge the funding support received from the Federal Ministry of Education and Research for Graphica project [grant number 03XP0004A]. The authors would also like to thank Dr. Martin Albrecht for the Electron Backscattering Diffraction measurements. References [1] L. Gao, W. Ren, H. Xu, L. Jin, Z. Wang, T. Ma, L.-P. Ma, Z. Zhang, Q. Fu, L.-M. Peng, X. Bao, H.-M. Cheng, Repeated growth and bubbling transfer of graphene with millimeter-size single-crystal grains using platinum, Nat. Commun. 3 (2010) 699, https://doi.org/10.1038/ncomms1702. [2] C.-M. Seah, S.-P. Chai, A.R. Mohamed, Mechanisms of graphene growth by chemical vapor deposition on transition metals, Carbon 70 (2014) 1–21, https://doi.org/10. 1016/j.carbon.2013.12.073. [3] Y. Zhang, L. Zhang, C. Zhou, Review of chemical vapor deposition of graphene and related applications, Acc. Chem. Res. 46 (2013) 2329–2339, https://doi.org/10. 1021/ar300203n. [4] J. Coraux, A.T. N’Diaye, M. Engler, C. Busse, D. Wall, N. Buckanie, F.M. Heringdorf, R.V. Gastel, B. Poelsema, T. Michely, Growth of graphene on Ir(111), New J. Phys. 11 (2009), https://doi.org/10.1088/1367-2630/11/2/023006 023006. [5] H. Ago, Y. Ito, N. Mizuta, K. Yoshida, B. Hu, C.-M. Orofeo, M. Tsuji, K.-I. Ikeda, S. Mizuno, Epitaxial chemical vapor deposition growth of single-layer graphene over cobalt film crystallized on sapphire, ACS Nano 4 (2010) 7407–7414, https:// doi.org/10.1021/nn102519b. [6] A. Reina, S. Thiele, X. Jia, S. Bhaviripudi, M.S. Dresselhaus, J.A. Schaefer, J. Kong, Growth of large-area single- and bi-layer graphene by controlled carbon precipitation on polycrystalline Ni surfaces, Nano Res. 2 (2009) 509–516, https://doi. org/10.1007/s12274-009-9059-y. [7] S. Thiele, A. Reina, P. Healey, J. Kedzierski, P. Wyatt, P.-L. Hsu, C. Keast, J. Schaefer, J. Kong, Engineering polycrystalline Ni films to improve thickness uniformity of the chemical-vapor-deposition-grown graphene films, Nanotechnology 21 (2010) 015601, , https://doi.org/10.1088/0957-4484/21/1/ 015601. [8] A. Reina, X. Jia, J. Ho, D. Nezich, H. Son, V. Bulovic, M.S. Dresselhaus, J. Kong, Large area, few-layer graphene films on arbitrary substrates by chemical vapor deposition, Nano Lett. 9 (2009) 30–35, https://doi.org/10.1021/nl801827v. [9] X. Wu, G. Zhong, L. Darsie, H. Sugime, S. Esconjauregui, A.W. Robertson, J. Robertson, Growth of continuous monolayer graphene with millimeter-sized domains using industrially safe conditions, Sci. Rep. 6 (2016) 21152, , https://doi. org/10.1038/srep21152. [10] S. Bae, H. Kim, Y. Lee, X. Xu, J.S. Park, Y. Zheng, J. Balakrishnan, T. Lei, H.R. Kim, Y. Il Song, Y.-J. Kim, K.S. Kim, B. Ozyilmaz, J.-H. Ahn, B.H. Hong, S. Iijimaet, Rollto-roll production of 30-inch graphene films for transparent electrodes, Nat. Nanotechnol. 5 (2010) 574–578, https://doi.org/10.1038/nnano.2010.132. [11] N. Liu, L. Fu, B. Dai, K. Yan, X. Liu, R. Zhao, Y. Zhang, Z. Liu, Universal segregation growth approach to wafer-size graphene from non-noble metals, Nano Lett. 11 (2011) 297–303, https://doi.org/10.1021/nl103962a. [12] C. Jeon, H.-N. Hwang, W.-G. Lee, Y.G. Jung, K.S. Kim, C.-Y. Park, C.-C. Hwang, Rotated domains in chemical vapor deposition-grown monolayer graphene on Cu (111): an angle-resolved photoemission study, Nanoscale 5 (2013) 8210–8214, https://doi.org/10.1039/C3NR01700A. [13] G. Lupina, J. Kitzmann, I. Costina, M. Lukosius, C. Wenger, A. Wolff, S. Vaziri, M. Ostling, I. Pasternak, A. Krajewska, W. Strupinski, S. Kataria, A. Gahoi, M.C. Lemme, G. Ruhl, G. Zoth, O. Luxenhofer, W. Mehr, Residual metallic contamination of transferred chemical vapor deposited graphene, ACS Nano 9 (2015) 4776–4785, https://doi.org/10.1021/acsnano.5b01261. [14] A. Dahal, M. Batzill, Graphene–nickel interfaces: a review, Nanoscale 6 (2014) 2548–2562, https://doi.org/10.1039/c3nr05279f. [15] X. Li, W. Cai, L. Colombo, R.S. Ruoff, Evolution of graphene growth on Ni and Cu by carbon isotope labeling, Nano Lett. 9 (2009) 4268–4272, https://doi.org/10.1021/ nl902515k.
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Large-scale chemical vapor deposition of graphene on polycrystalline nickel films: Effect of annealing conditions
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Fatima Akhtara, , Jaroslaw Dabrowskia, Marco Liskera, Peter Zaumseila, Sebastian Schulzea, Alex Jouvrayb, Piotr Cabanc, Andreas Maia,d, Christian Wengera,e, Mindaugas Lukosiusa IHP – Leibniz-Institut für Innovative Mikroelektronik, Im Technologiepark 25, 15236 Frankfurt (Oder), Germany Aixtron Ltd., Anderson Road, Swavesey, CB24 4FQ Cambridge, UK c Institute of Electronic Materials Technology, Wolczynska 133, 01-919 Warsaw, Poland d Technical University of Applied Science Wildau, Hochschulring 1, 15745 Wildau, Germany e Brandenburg Medical School Theodor Fontane, 16816 Neuruppin, Germany a
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A R T I C LE I N FO
A B S T R A C T
Keywords: Polycrystalline nickel Surface pretreatment Grain growth Grain boundary grooves Chemical vapor deposition Graphene
In the present study, 8-in. silicon substrates, covered with thin (200 nm) polycrystalline nickel films have been employed for the growth of graphene by chemical vapor deposition. In order to control the uniformity and coverage of the graphene, thin nickel layers were used due to their less deep grain boundary grooves and ability to store less carbon in comparison with thick nickel films (> 500 nm). The preferential sites for the growth of multilayer graphene were influenced by the surface pretreatment of the polycrystalline nickel films at 1025 °C under different ambient conditions (hydrogen and vacuum). Significant differences in the surface morphologies were observed for the annealed nickel films. The growth of larger grains up to ~6 μm for the films annealed in hydrogen could be attributed to hydrogen interstitials. On the other hand, grains up to ~3 μm were extracted for the films annealed in vacuum. Graphene was grown after exposing the annealed Ni films to ethylene at 925 °C. The lower range (42–106 cm−1) of full width at half maxima of the 2D band as determined by Raman spectroscopy was obtained for the films annealed in hydrogen as compared to the ones annealed in vacuum (51–128 cm−1), indicating that the thickness uniformity of graphene was strongly influenced by the surface modifications of nickel films.
1. Introduction Due to its properties, graphene has appeared as a potential candidate in various electronic devices. However, one of the biggest challenge in its integration into microelectronics is the lack of high-quality and wafer-scale synthesis of the graphene. Typically, chemical vapor deposition (CVD) is used for the growth of graphene on metal catalysts [1–9]. Indeed, copper (Cu) catalyst has been employed for large-area graphene growth [9,10], however, a large lattice mismatch between graphene and Cu [11], the rotated graphene domains [12] and the presence of metallic contaminants after a complicated transfer step can degrade the quality of graphene [13]. In contrast to Cu, the growth of graphene on Nickel (Ni) occurs by a segregation phenomenon, this is attributed to the high carbon solubility (~2.03 at. % at 1000 °C) in Ni [14,15]. Due to the excellent lattice match between graphene and Ni, single crystal Ni (111) was reported to be one of the most suitable substrate for the growth of monolayer graphene [14,16–18].
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Nevertheless, high costs and limited sizes are the main drawbacks of single crystal substrates. At this point, polycrystalline Ni films on SiO2/ Si (001) are attractive for the growth of graphene due to their largescale availability and compatibility with Si microelectronics. However, the polycrystallinity of Ni makes the graphene growth a complicated process. Grain boundaries (GBs) in polycrystalline Ni are the most preferential sites for the growth of multilayer graphene, whereas Ni (111) grains favor the growth of monolayer graphene. Therefore, in order to crystallize Ni along the most preferred (111) orientations and to reduce the number of GBs, typically high temperatures (~1100C) [7] and thick Ni films (500 nm) have been used [6–8,15,18]. Thick Ni films were usually employed in order to avoid the dewetting phenomenon at higher annealing temperatures, however, thick Ni films suffer from several drawbacks, for instance, the presence of deeper GBs grooves, which could increase the surface roughness and could be the cause of higher coverage of multilayer graphene [18]. Another disadvantage of thick Ni films is related to their higher carbon storage efficiency that
Corresponding author. E-mail address: [email protected] (F. Akhtar).
https://doi.org/10.1016/j.tsf.2019.137565 Received 18 December 2018; Received in revised form 15 August 2019; Accepted 11 September 2019 Available online 12 September 2019 0040-6090/ © 2019 Elsevier B.V. All rights reserved.
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Renishaw Raman spectrometer with 514 nm excitation wavelength of the laser and beam spot size ≈ 1 μm. X-ray diffraction (XRD; SmartLab diffractometer from Rigaku) was used to determine the crystal structure using CuKα (λ = 1.54186 Å) radiation. The XRD measurements were performed in a parallel beam configuration. Surface morphology of Ni films was investigated by scanning electron microscopy (SEM; Gemini 2, Zeiss) at 1.5 kV. The grains distribution in Ni films was studied by FEI Nova 600 Nanolab equipped with an electron backscattering diffraction (EBSD) detector at 30 keV with 9.5 nA beam current. Atomic force microscopy (AFM; Dimension 5000 SPM System with NanoScope IV Controller from Digital Instruments) in tapping mode was used to monitor the surface roughness of Ni substrates. The elemental composition of as-deposited and annealed samples was determined by x-ray photoelectron spectroscopy (XPS) using a monochromatic MgKα (1253.6 eV) source under 7 × 10−8 Pa with a hemispherical analyzer (PHOIBOS-100). The beam energy and beam current density used for XPS experiments were 300 W and 15 KV, respectively. Moreover, the binding energy was calibrated to gold (Au) 4f line set to 83.95 eV. XPS spectra were analyzed with the CasaXPS software. A Shirley back ground was used in the curve-fitting procedure along with a asymmetric line shapes (LF) for Ni2p spectra. However, Gauss-Lorentzian (GL) line shapes were used to fit the C1s and O1s spectra.
makes the thickness control of the graphene challenging [11,18]. Indeed, as reported in literature, the coverage of mono-bilayer graphene on polycrystalline Ni films varied in the range from 70% to 87%, indicating the complexity of this system [6,11,18,19]. Therefore, the purpose of this work is to employ thin (200 nm) Ni films deposited on 8in. Si wafers and to prepare its surface for the growth of graphene by modifying its grain sizes, orientations, and GBs by different pretreatment. One of the advantages of thin Ni films as compared to the thick films is the presence of less deep GBs grooves, which can play a significant role in reducing the growth of multilayer graphene. 200 nm Ni films were annealed at 1025 °C under different ambient conditions (hydrogen/vacuum), keeping other parameters constant to examine their effect on the surface roughness, grain sizes, and crystallinity of the annealed Ni films. After annealing, these Ni films were exposed to ethylene at 925 °C for the growth of graphene, and the quality of the obtained graphene films was then compared. The findings of this study revealed that annealing in hydrogen led to a substantial increase in the sizes of Ni grains as compared to the annealing in vacuum. This reflects a higher coverage of a thinner graphene on the Ni films annealed in hydrogen, as measured by Raman, indicating the influence of grain sizes on the thickness of graphene. 2. Experimental details
3. Results & discussions 2.1. Deposition, annealing of Ni/SiO2/Si, and growth of graphene The crystal quality and the surface morphology of the as-deposited and the annealed Ni films were investigated by XRD, SEM, AFM, and EBSD. The diffraction patterns were recorded by specular ω–2θ scans in the angular range of 2θ = 40°–55° using for as-deposited and annealed Ni films (hydrogen/vacuum) as shown in Fig. 2a. As-deposited Ni was polycrystalline as the reflections at 44.52°, assigned to Ni (111) orientation, and a weak Ni (200) reflection at 51.87° corresponding to Ni crystallites with (100) orientations, were measured. The Ni (111) orientation has the lowest surface/interface energy with respect to other crystallographic orientation, which is a driving force for their growth. XRD curves of the Ni films annealed in hydrogen (black line) and vacuum (red line) at 1025 °C are also presented in Fig. 2a. A slight shift of Ni (111) and Ni (200) reflections to higher angles is observed after annealing in hydrogen and vacuum. These shifts could be attributed to the different thermal expansion coefficients between Ni (14 × 10−6 K−1) and the underlying Si substrate (4 × 10−6 K−1) [7,20]. A detailed analysis of the Ni (111) reflections was performed by conducting ω scans to determine the mosaic spread in the films appearing due to randomly oriented crystallites. In these experiments, the detector was fixed at the center of Ni (111) diffraction peak, and the sample was tilted. The full width at half maxima (FWHM) of the obtained curve is an objective measure for the mosaicity of the Ni films. The values of FWHM are plotted as a function of annealing temperatures and are depicted in the inset of Fig. 2a. After annealing at 1025 °C, the values of FWHM decreased from 9.63° (as-deposited Ni) to 5.43° (annealed in hydrogen) and to 5.89° (annealed in the vacuum). This reduction indicates that the mosaicity of Ni (111) planes is lower in both annealing conditions. However, the lower values of FWHM of Ni (111) reflection for the films annealed in hydrogen demonstrate that the alignment of Ni grains is enhanced compared to films annealed in vacuum. SEM images taken with an HE-SE2 detector (high-efficiency secondary electron detector; used to investigate surface topography) are shown in Fig. 2b–c for the films annealed in hydrogen and vacuum, respectively. Firstly, inhomogeneity in Ni grain size and shape can be noticed in the SEM images. This inhomogeneity is the result of a grain growth phenomenon in which lateral grain size exceeds the film thickness. During the grain growth mechanism, faster-growing grains dominate the slower ones, and this process continues until all un-favored grains diminish leading to a texture development in thin films [21,22]. Secondly, surface roughness was increased due to the grain
Ni films were deposited from a Ni source by DC magnetron sputtering method on 100 nm thick thermal SiO2 on 8-in. Si (001) substrate. The depositions were carried out at room temperature in ~55 sccm argon. During the deposition process, plasma was ignited at 750 W and the pressure was kept at 0.533 Pa. The growth of graphene on polycrystalline Ni is a complicated process that can be divided into four steps as shown in Fig. 1. In the first step, the Ni grain growth was performed by annealing the Ni films for 10 min in a CVD chamber at 1025 °C (label a). Annealing was performed in different ambient conditions, hydrogen (P = 10 Pa) or vacuum (P = 1.08 × 10−2 Pa). After annealing, the temperature was reduced to 925 °C (label b), where hydrocarbons (C2H4) were let into the chamber and the pressure in the chamber was kept at 10 Pa (label c). After 5 min, the supply of the hydrocarbons was stopped and the system was allowed to cool down. As a result, carbon atoms started to precipitate and after crystallization on the Ni surface formed a graphene sheet (label d). 2.2. Characterization Quality of the deposited graphene layers was monitored by a
Fig. 1. Illustration of the temperature-time profile of graphene growth process by CVD. 2
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Fig. 2. XRD ω–2θ scans of (a) as-deposited Ni film (blue line), Ni films annealed in hydrogen (black line) and vacuum (red line). The intensity is normalized in a logarithmic scale. FWHM of Ni (111) reflection obtained from ω scans as a function of annealing temperature is shown in the inset of Fig. 2a. SEM images of Ni films annealed in (b) hydrogen and (c) vacuum. AFM images of (d) as-deposited and Ni films annealed in (e) hydrogen and (f) vacuum. SEM image of as-deposited Ni film is shown in the inset of Fig. 2d. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
hydrogen. This is in agreement with the literature report where wide and flat Ni (111) nanostructures have been observed due to the dissolution of hydrogen [25]. According to another study, the presence of interstitial hydrogen causes additional stresses in the Ni films that could result in the volume expansion of the Ni grains [20]. On the contrary, Ni films annealed in vacuum were not under the influence of hydrogen and the grain growth was affected by abnormal grain growth mechanism and grain sizes up to 3 μm were recorded. In comparison with the reports in the literature, different grain sizes were obtained for different thicknesses of Ni films. For 500 to 550 nm thick Ni films, the grain sizes varied from 1 to 20 μm [6–8,27,28], whereas even larger grains (≥50 μm) were obtained for Ni foils (25 μm) [29]. Although larger grains were reported for thick Ni, the disadvantages of thick Ni were previously discussed. For the thinner Ni films (≤300 nm), the grain sizes were found to be much smaller as compared to the thicker films. The grain sizes were reported within the range of 1 to 4 μm upon reducing the thickness of the Ni films to 300 nm [30–33]. Literature reports on the grain sizes for thin ≤200 nm Ni films (as used in this work) are rather limited and were in the range from 150 nm to 1 μm [34–36]. The authors in the work [35,36] attribute the smaller grains to the simultaneous growth of Ni grains and carbon diffusion during the annealing step; these smaller grains resulted in inhomogeneous growth of graphene, indicating the importance of preannealed Ni films. In our study, the crystallization of Ni films in different annealing conditions before graphene growth resulted in larger grains up to 6 μm.
growth, as determined by AFM. The root mean square surface roughness (Rq) was recorded by scanning 5 × 5 μm2 areas of the as-deposited and the annealed Ni films as shown in Fig. 2d–f. The Rq values increased from 1.6 nm (as-deposited Ni) to ~28 nm for the films annealed in hydrogen and ~10 nm for the ones annealed in vacuum. As compared to literature, the roughness of the films annealed in hydrogen was lowered than the reported values of 36.3 nm obtained after 15 min annealing (900 °C) of 500 nm Ni under 600 sccm H2 pressure [18]. Lower surface roughness in this study could be attributed to less deep GBs grooves of thin films as compared to thick films. The grain sizes of the as-deposited Ni films were around 20 nm (inset of Fig. 2d) and they increased to several microns depending on the annealing conditions. Indeed, as extracted from 30 × 80 μm2 EBSD maps in Fig. 3a–b, grains up to ~6.0 μm diameter were measured within the films annealed in hydrogen (Fig. 3c squares) and up to ~3.0 μm were observed within the films annealed in the vacuum (Fig. 3d circles). The grain sizes below 0.6 μm could not be adequately indexed and have been removed [23]. According to EBSD and SEM results, larger grains were recorded for the films annealed in hydrogen in compare to the films annealed in vacuum. The difference in Ni grain sizes as a function of different annealing ambient could be attributed to the presence of hydrogen. Hydrogen is known to incorporate in the interstitial sites in a host metal lattice during annealing [20,24–26]. This may generate stresses in the Ni films and thus a volume expansion of the Ni crystallites. Therefore, we assume that larger Ni grains (6 μm) for the films annealed in hydrogen could be attributed to the stress induced by interstitial 3
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Fig. 3. EBSD analysis: colour map, Orientation maps for Ni films annealed in (a) hydrogen (b) vacuum. Grain distributions measured by EBSD for Ni films annealed at 1025 °C in (c) hydrogen (squares) (d) vacuum (circles) and after graphene growth at 925 °C (triangles).
Fig. 4. XPS analysis: (a) O1s (b) Ni2p and (c) C1s spectra of Ni films as-deposited (black line), annealed at 1025 °C in hydrogen (magenta line) and vacuum (blue line) and after graphene growth at 925 °C (green line). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Fig. 5. EBSD images of Ni films annealed in (a) hydrogen (b) vacuum after graphene growth. SEM images of graphene grown on Ni films annealed in (c) hydrogen (d) vacuum.
Before starting the growth of graphene, the Ni films were investigated by XPS. Fig. 4 summarizes the XPS data that provides complementary information about the chemical state of the surface of the Ni-films. The O1s spectrum of the as-deposited Ni (black line) as presented in Fig. 4a is fitted to two main components at 529.6 eV (FWHM = 1.4 eV) and 531.7 eV (FWHM = 2.0 eV), which are assigned to NiO and Ni2O3, respectively [37]. In Fig. 4b, the Ni2p spectrum of the as-deposited Ni (black line) is composed of Ni° (852.6 eV), Ni+2 (854.1 eV) and Ni+3 (856.1 eV) corresponding to metallic Ni (FWHM = 1.5 eV), NiO (FWHM = 2.0 eV) and Ni2O3 (FWHM = 2.4 eV), respectively. Some satellite (shakeup) peaks at a few eV above the main line are also visible in the spectrum; the satellite peaks of Ni oxides are located at higher binding energy than that of the metallic Ni. They appear as the result of a change in the Coulomb potential of electrons due to the removal of core electron; their orbitals relax and liberate energy which is utilized to shake up an outer electron to a higher d-orbital [38]. Moreover, the peaks in the C1s spectrum (black line) of the as-deposited Ni as demonstrated in Fig. 4c are attributed to CeH (285.0 eV), CeO (286.5 eV), and C]O (288.9 eV). The FWHM of the main line contributions are 1.5 eV, 1.5 eV and 1.7 eV, respectively. These XPS results demonstrated that the as-deposited Ni films were contaminated when they were exposed to air; however, these contaminants have been removed after annealing in hydrogen (magenta lines) and vacuum (blue lines) as shown in Fig. 4a–c. In order to grow graphene, the Ni films annealed in hydrogen and in vacuum were exposed to ethylene (Partial pressure = 5 Pa) and hydrogen (Partial pressure = 5 Pa) for 5 min at 925 °C and XPS analysis was applied again. In Fig. 4c, the evolution of an sp2 bonded carbon signal at ~284.6 eV (FWHM = 1.1 eV) in C1s core level spectrum (green line) implies the formation of graphene. The obtained value of
the binding energy in this work is comparable to the reported values, where authors described that the binding energy of graphene on polycrystalline Ni lies in the range of 284.3 to 284.5 eV [11,27,32,39]. In contrast to the polycrystalline Ni films, the binding energy of graphene on single crystal Ni (111) has been reported to be within the range of 284.8 to 285.0 eV [40,41]. It could be attributed to the stronger interaction between graphene and single crystal Ni (111) as compared to the graphene/polycrystalline Ni system [41]. Furthermore, no oxide signals were recorded in the O1s and Ni2p spectra, as shown in Fig. 4a–b (green lines). It is worth mentioning here that in-situ XPS measurements have been performed in this study (both after annealing the Ni films and after the growth of graphene) due to interconnected XPS and CVD chambers. The grain distributions after graphene growth were again investigated by EBSD and are demonstrated in Fig. 3c–d (triangles). The orientation maps of Ni films collected after the graphene growth are shown in Fig. 5a–b. By comparing the orientation maps (Fig. 3a–b) and the grain distributions in Fig. 3c–d for the Ni films annealed in hydrogen (squares) and vacuum (circles) with the orientation maps (Fig. 5a–b) and grain distributions (Fig. 3c–d triangles) for the films obtained after the graphene growth, one can notice that they remained the same. Indeed, the graphene growth did not influence the grain sizes within the Ni films. In other words, the annealing procedure carried out with a higher thermal budget than the one during graphene growth forced the Ni film morphology to a structural state which was stable under the graphene growth conditions. SEM is a powerful tool to identify different numbers of graphene layers. The surface morphology of graphene was investigated by SEM; images taken with the in-lens detector are shown in Fig. 5c–d. This detector was used since it generates different colour
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Fig. 6. (a) Raman spectra of graphene grown on Ni films annealed in (a) hydrogen (b) vacuum; different colours corresponding to spectra taken from different areas. Raman maps (2D-FWHM) of graphene grown on Ni films annealed in (c) hydrogen (d) vacuum.
graphene can be grown in both conditions. The FWHM of the 2D band is presented in Fig. 6c–d. The range of the FWHM values lied in the range 42–106 cm−1 for hydrogen annealed Ni films and 51–128 cm−1 for the films annealed in vacuum. The lower values of FWHM indicated the presence of thinner graphene on hydrogen annealed films [18,37,47]. Moreover, the I2D/IG ratio ~2.4 (label 1) in Fig. 6c suggested the growth of mono–bilayer graphene [32,37]. The I2D to IG ratio dropped to ~0.6 (label 2) and ~0.5 (label 3) in Fig. 6c, indicating the presence of few layers (3–5) and multilayer graphene [32,37], respectively. In the case of the films annealed in vacuum, the I2D/IG ratio ~1.7 (label 1), ~0.5 (label 2) and ~0.4 (label 3) in Fig. 6d revealed the growth of thicker graphene. To summarize, Raman analysis demonstrated the growth of thinner graphene on the films annealed in hydrogen as compared to the ones annealed in vacuum. The higher coverage of thinner graphene on the films annealed in hydrogen was attributed to larger Ni grains and reduced grain boundaries. On the contrary, smaller grains and therefore increased GBs for the films annealed in vacuum facilitated the growth of multilayer graphene.
contrasts of a surface according to its work functions. The charge transfer from metal to the layers of graphene is different, which results in different work function values and therefore the colors in the SEM image. The non-uniformity of the graphene coverage is clearly visible in SEM images. The surface of the scanned area was divided into three distinct areas based on their work functions. Herein, the light grey or brighter areas corresponded to the surfaces with low work function or thinner graphene (label 1). Upon increasing the number of grown graphene layers, their area became darker (labels 2 and 3) [32,42,43]. Finally, the quality of graphene was monitored by Raman spectroscopy. A Raman spectrum of graphene consists of three different peaks known as G peak (~1580 cm−1), D peak (~1350 cm−1) and 2D peak (~2700 cm−1) [44,45]. The G peak is the result of in-plane CeC stretching mode and corresponds to doubly degenerate E2g phonons at the Brillouin zone center (Γ point). The D peak or defect mode is the result of a double resonance process and requires a TO phonon and a defect in its activation. Whereas the 2D peak, which is also known as a graphene fingerprint, is the overtone of the D peak and is caused by a second order two-phonon Raman process involving two TO phonons near the K point [44–46]. The I2D/IG and FWHM of the 2D band provide the basis for determining the number of graphene layers. The height of the 2D band peak corresponding to monolayered graphene is larger than the G peak and has a symmetric line shape that can be fitted with a single Lorentzian peak. The height and shape of the 2D band peak changes with increasing thickness of the graphene film, for instance, thicker graphene demonstrated a broadened 2D peak with reduced amplitude [45]. As shown in Fig. 6a–b, Raman spectra were monitored at three different regions depending on the thickness of the grown graphene films. The presence of small defect peaks indicates that a good quality
4. Conclusions Two different pretreatment conditions (hydrogen and vacuum) were employed to anneal 200 nm thick polycrystalline nickel films deposited on 8-in. 100 nm SiO2/Si (001) substrates. Both annealing conditions favored the crystallization of Ni films that resulted in a strong {111} texture; however, omega scans of Ni (111) reflections indicated that the mosaicity of the grains was reduced in hydrogen annealed Ni films. According to the AFM analysis, the surface roughness of Rq ~28 nm was recorded for hydrogen annealed films, whereas for vacuum annealed films it was found to be Rq ~10 nm. EBSD measurements 6
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revealed larger grains for the films annealed in hydrogen. Hydrogen ions incorporate into the interstitial sites of Ni films thus enabling the growth of large Ni grains. As indicated by Raman spectroscopy, the reduced bandwidth of the 2D peak of the graphene grown on the Ni films annealed in hydrogen indicated the relevance of grain size for the growth of thinner graphene. However, the surface of Ni films annealed in hydrogen was rougher. The smaller Ni grains for the films annealed in vacuum resulted in increased number of grain boundaries that facilitated the growth of multilayer graphene.
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