Article pubs.acs.org/JPCC
Insights on Defect-Mediated Heterogeneous Nucleation of Graphene on Copper Priyadarshini Ghosh,† Shishir Kumar,‡ Gopalakrishnan Ramalingam,§ Vidya Kochat,∥ Madhavan Radhakrishnan,⊥ Sukanya Dhar,‡ Satyam Suwas,⊥ Arindam Ghosh,∥ N. Ravishankar,† and Srinivasan Raghavan*,‡ †
Materials Research Centre, ‡Centre for Nano Science and Engineering, ∥Department of Physics, and ⊥Department of Materials Engineering, Indian Institute of Science, Bangalore 560012, India § Department of Materials Science and Engineering, University of Virginia, Charlottesville, Virginia 22903, United States S Supporting Information *
ABSTRACT: The grain size of monolayer large area graphene is key to its performance. Microstructural design for the desired grain size requires a fundamental understanding of graphene nucleation and growth. The two levers that can be used to control these aspects are the defect density, whose population can be controlled by annealing, and the gas-phase supersaturation for activation of nucleation at the defect sites. We observe that defects on copper surface, namely dislocations, grain boundaries, triple points, and rolling marks, initiate nucleation of graphene. We show that among these defects dislocations are the most potent nucleation sites, as they get activated at lowest supersaturation. As an illustration, we tailor the defect density and supersaturation to change the domain size of graphene from 100 μm2. Growth data reported in the literature has been summarized on a supersaturation plot, and a regime for defect-dominated growth has been identified. In this growth regime, we demonstrate the spatial control over nucleation at intentionally introduced defects, paving the way for patterned growth of graphene. Our results provide a unified framework for understanding the role of defects in graphene nucleation and can be used as a guideline for controlled growth of graphene.
1. INTRODUCTION Since the observation of an ambipolar electric field effect1 in exfoliated micron-sized flakes, graphene has come a long way from a scientific novelty to being explored for device commercialization. A big reason for this advance is the current ability to grow layers of graphene over large areas, on metallic surfaces, by chemical vapor deposition (CVD) using hydrocarbon−hydrogen mixtures.2−8 The polycrystallinity9 of CVD graphene dictates that its properties are controlled by the grain or domain size and grain orientation as in any other polycrystalline material. For many of the properties of current interest such as electronic mobility10−14 and thermal conductivity,15,16 large grain sizes are desirable. On the other hand, defects and grain boundaries can have special properties,17−23 and have also been shown to be chemically more active.24 Therefore, a decrease in grain size or increase in grain boundary length per unit area would benefit applications such as sensors25,26 and electrochemical applications like batteries.27,28 There is thus a clear need for the ability to incorporate the desired length of grain boundary length per unit area, or in other words, engineer the microstructure of graphene. The grain boundary length per unit area of a graphene film growing on a substrate is determined by the nucleation density (number of nuclei per unit area) and the growth rate of the nuclei so formed. A decrease in nucleation density and an © XXXX American Chemical Society
increase in growth rate would result in larger grain sizes and smaller grain boundary density. Nucleation and growth are readily influenced by defects on growth surfaces. It has been documented frequently for growth of graphene on Cu foils,3 which are presently the preferred growth substrates.2−5,29−43 The Cu substrates used for graphene growth are generally coldrolled polycrystalline foils. Upon heating to the CVD graphene growth temperatures of about 1000 °C, the foils develop a predominantly cubic (100) texture.44 The (100) surface of Cu so formed is defective and contains high angle grain boundaries, triple junctions, rolling marks, and low angle grain boundaries or arrays of dislocation lines. In order to reduce these defect densities, increase Cu grain size, and reduce surface roughness, the foils have been subjected to various annealing and polishing treatments.4,38 These efforts have focused on trying to increase the domain size of graphene, for good measure, as electronic applications that require high mobility seem to be of imminent importance. Toward this end, graphene flake sizes as large as 4.5 mm2 have been recently demonstrated on Cu foil annealed at 1077 °C for 7 h with mobilities approaching that of exfoliated graphene.4 Received: October 20, 2014 Revised: December 26, 2014
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DOI: 10.1021/jp510556t J. Phys. Chem. C XXXX, XXX, XXX−XXX
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The Journal of Physical Chemistry C
Figure 1. SEM images showing nucleation and growth of graphene islands on Cu foils annealed for 5 min, 2 h, and 5 h (indicated in bottom left corner of the images) in rows 1, 2, and 3, respectively. Graphene growth time is indicated in the top right corner of the images. Graphene-covered area of the Cu surface appears darker than the bare one. Decrease in density and increase in growth rate of islands with increase in pregrowth annealing time is clearly seen. White arrows point to examples (discussion in text) of islands formed within a Cu grain. Many of these islands also have a white dot at their center. The Cu grain boundaries can be identified (see image (d)) as the lines decorated by the white dots or Cu precipitates. Scale bar in all images: 10 μm.
2. EXPERIMENTAL SECTION 2.1. Graphene Synthesis. A He leak tight CVD reactor setup was used for graphene synthesis. A 1 in. horizontal quartz tube inserted in a tube furnace and pumped with a rotary pump served as the reaction chamber (Figure S1a, Supporting Information). Pressure was measured using a 0−1000 Torr Baratron from MKS instruments. All the gases used in our experiments, H2, N2, and CH4, were of ultrahigh purity grade (99.999%). In addition, H2 and N2 were passed through a point of use getter purifier (SAES Pure Gas, Inc.), and 99.8% pure 25 μm thick copper foils obtained from Alfa Aesar were used as substrates for the synthesis. After the substrate was loaded in the reaction chamber, the system was pumped down to a base pressure of ∼1 Torr, and following that, the system was flushed with 250 sccm of H2 for 10 min. A three zone furnace was used to ramp the system to 1000 °C, with a ramping rate of 25 °C min−1, along with a reduced 50 sccm flow of H2. At 1000 °C, the substrates were subjected to annealing for different time intervals. Near 1000 °C, the methane flow was established in an auxiliary line directly connected to the pump for stabilization of the flow. Once the flow in auxiliary line was stabilized, the switching valve was activated to connect the methane line to the inlet (Figure S1a, Supporting Information). The growth of graphene was carried out with the injection of CH4 diluted with H2. The default flow ratio was CH4:H2:::5:1000 sccm. To study kinetics, the ratio was changed as mentioned in the text. The total default pressure during growth was ∼18 Torr. Following growth, the reactor was cooled to room temperature under a 1000 sccm flow of H2. The position of the Cu foil inside the quartz tube does not have any observable effect on the growth. We tested with the positions shown in Figure S1a (Supporting Information), just inside the beginning and at the end of the heating zones. 2.2. EBSD Mapping of Cu Substrate. Orientation image mapping was carried out in a scanning electron microscope (SEM) equipped with EBSD facility. Scans were taken on the
In this paper, we show that this defect structure, particularly dislocations that intersect the Cu surface, can in fact be beneficial. They can be exploited to control the nucleation density in a particular low-supersaturation window. Annealing of the Cu foils can be used to control the density of different types of defects on the Cu surface on one hand. On the other hand, nucleation at different defect sites on the surface of Cu can be triggered at different degrees of vapor-phase supersaturation. In addition to the control that can be gained on the nucleation density using these two parameters, degree of annealing and supersaturation, the growth rate of the nucleated islands changes with the degree of annealing. These three parameters have been used to obtain large-area CVD graphene with different kinds of microstructures. Dislocations, in particular, are shown to be the most helpful defects with nucleation happening at the lowest supersaturation at regions with the highest dislocation densities, rolling marks, or plastic deformation zones. This is then followed by nucleation at slightly higher supersaturation at isolated cores within grains and eventually at all locations on the Cu surface. The last of these methods mentioned above, only supersaturation control with no defect influence, has been used by Kim et al. in a detailed study on the nucleation and growth of graphene on Cu subjected to various treatments.3 In yet another report, the annealing of Cu foils has been studied in detail by Robinson et al.44 The aim therein was to obtain (111) textured Cu. The research reported here differs from these two previous reports in that we quantitatively correlate the nucleation density of graphene with the number of dislocations on the growth surface, thereby adding an extra lever for growth control. We also summarize current literature data in a supersaturation map and identify the window that can be exploited for the above approach of defect con-trolled nucleation and microstructural design. To the best of our knowledge, this is the first time that such a correlation has been established. B
DOI: 10.1021/jp510556t J. Phys. Chem. C XXXX, XXX, XXX−XXX
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The Journal of Physical Chemistry C
Figure 2. Statistical summary of Figure 1 showing effect of annealing Cu foils for 5 min, 2 h, and 5 h prior to growth: (a) ∼100-fold reduction in saturation island density (≤graphene nucleation density) and 20-fold increase in estimated grain size with Cu foil annealing time; (b) no significant increase in time taken for complete graphene coverage with annealing time in spite of the 100-fold nucleation density reduction. This implies an increase in growth velocity as shown by the plot of measured linear growth rates of island edges. Lines show trends and are not fits to data.
5:1000 sccm. At this very large H2 flow the effect of oxygen as reported by Choubak et al. is unlikely to have an effect as the oxygen potentials in our system are very low as described in the Supporting Information, S3.43 The white specks/particles seen in the images are composed of Cu and commonly observed in CVD graphene.3 They condense from the vapor phase during post-growth cool down and play no role in nucleation of graphene at the growth temperature (Supporting Information, section 2). The statistical analysis of the effect of pregrowth annealing time on surface coverage of graphene is summarized in Figure 2a,b in terms of saturation island density (obtained from island density vs growth time plot Supporting Information, section 4, Figure S4), grain size in the coalesced layer (estimated from maximum island density), island velocity and “100%” coverage time. It is clear from Figure 2 that there is a reduction in the saturation density of graphene islands with an increase in the annealing time of the Cu foils. For foils annealed for 2 h or longer at 1000 °C, the saturation island density observed is 2 orders of magnitude lower than samples annealed for 5 min (Figure 2a). Since some of the nuclei might have coalesced, the actual nucleation density would be higher. From the images, it is evident that the difference between observed island density and actual nucleation density would be highest for the 5 min annealed samples and lower for the remaining two. Thus, the 100-fold reduction in island density observed for the annealed case is a conservative estimate of the reduction in nucleation density. In the remainder of this paper, to avoid confusion, only the terms nuclei and nucleation density will be used. No significant change to the nucleation density was observed on annealing for longer than 2 h. Thus, from the point of view of nucleation density reduction alone, this could be construed as an upper limit for annealing of Cu foils. AFM analysis (Supporting Information, section 5) shows that the surface roughness of the Cu foils used does not change substantially on annealing and therefore can be discarded as a possible cause for the reduction in nucleation density observed.45 In addition, XPS (Supporting Information, section 6), shows only Cu and Cu2+ peaks with the latter decreasing with annealing time. Given that Cu is the active catalyst, the reduction in nucleation density cannot be explained by this feature either. No attempts were made to determine the exact time between 5 min and 2 h at which the nucleation density would plateau off as it was not deemed necessary for the defect-nucleation density correlations to be established in this paper. Assuming hexagonal grain shape at coalescence, the reduction in nucleation density would result in a 4-fold decrease in grain length per cm2 from 166 to 40 m.
normal plane for the as-received and the annealed samples. All of the samples for EBSD were electropolished in a Struer’s electropolishing setup with Struer’s D2 solution containing phosphoric acid, ethanol, distilled water, and 1-propanol. TSL Orientation Image Map (OIM, version 5.2) software was used to collect and analyze the EBSD data. In OIM micrographs, neighboring pixels separated by a minimum misorientation of 2° (grain tolerance angle) were defined as grain boundaries, and grains were defined to possess a minimum of two data point (pixels) in multiple rows with neighboring points of misorientation within the limits of tolerance angle defined earlier. 2.3. Etch Pit Density Measurement. The Cu samples annealed at 1000 °C were etched at room temperature with an 8 M HNO3 solution. The etchant solution was made with highpurity distilled water. The samples were cleaned with glacial acetic acid prior to each etching experiment. Etching was carried out for various time periods. After etching, samples were cleaned properly with high purity DI water and then analyzed using both optical microscopy and scanning electron microscopy. Pit density was obtained by image analysis in “ImageJ” software. 2.4. Nucleation Density Measurement. Graphene appears darker compared to the Cu surface in scanning electron micrograph. Hence, graphene nuclei can be distinguished, and their density can be measured from SEM of the partially covered Cu surface. The nuclei density was calculated in the same manner as pit density. For each data point, nucleation density was calculated from ∼20−25 random locations from the 1 cm2 partially covered graphene. Nucleation density for each growth condition was averaged over at least three trial runs. 2.5. Growth Rate Measurement. The linear growth rate of individual graphene grain was determined from SEM images taken at two different times in the linear portion of the JMAK plot as V = (r2 − r1)/(t2 − t1). V, r, and t are velocity, radius, and time, respectively.
3. RESULT AND DISCUSSION 3.1. Effect of Pregrowth Annealing on Nucleation Density and Growth Rates. We start with the observations depicted in Figure 1. The SEM images show, qualitatively, that with an increase in the annealing time, 5 min, 2 h, and 5 h, of Cu foils prior to graphene growth, there is a reduction in the island density and an increase in the growth rate of the nucleated domains. Unless otherwise stated and as discussed later, the CH4:H2 flow ratio used for graphene growth was C
DOI: 10.1021/jp510556t J. Phys. Chem. C XXXX, XXX, XXX−XXX
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nucleation at these relatively higher energy defects, no such preference is observed for twin boundaries having lower energies as shown in Figure 3c. As shown in Figure 1d,e,g,h, the Cu deposits also preferentially nucleate at these defect sites during post-growth cool down. In Figure 1d, they decorate the grain boundaries of the Cu foil. In addition to the effect of flow rates during cooling discussed in Supporting Information (section 2), the density of Cu precipitation decreased with the increase in graphene coverage (Supporting Information, Figure S8), implying that they did not like to nucleate on graphene covered areas. In many instances, following almost “100% graphene coverage” the Cu precipitates formed only along the grain boundaries of the islands in the graphene layer as seen in Figure 3d. In contrast to the above observation of nucleation of graphene and the Cu particles at defect sites, one apparently anomalous feature was the observation of graphene islands formed within a Cu foil grain (see Figure 1d, for instance). These islands apparently therefore have no preference for any of the aforementioned high energy defects. In most, if not all, cases, there are a few Cu deposits within a graphene island and in many cases almost at the center as if the graphene island nucleated at that spot. This presence of such Cu deposits points to an enhanced preference, energetically speaking, for a particular location within a Cu foil grain by both the graphene nuclei as well as the Cu island. Thus, these apparently anomalous nuclei also probably lie on a defect within the grain such as a dislocation core. These observations of preferred nucleation thus point to the fact that not all points on the annealed surface of Cu are energetically equivalent from the point of view of nucleation. However, in order to exploit this preference for nucleation and the change in nucleation densities and growth rates on annealing for graphene microstructure design, the correlation between annealing and the surface structure of Cu needs to be first established. This is the topic of the next section. 3.3. EBSD and Etch Pit Characterization of Surface Defect Structure Evolution on Annealing. The orientation
As expected, the change in density of grain boundary has a direct effect on electrical properties of CVD graphene. We observed a 10-fold decrease in resistance of graphene ribbons of fixed dimensions, obtained from 5 min and 2 h annealed Cu (Supporting Information, section 7). 3.2. Preferential Nucleation at High Energy Defect Sites. A closer look across the spectrum represented by Figure 1 indicated that nucleation did have a preference for defects as shown in Figure 3 and reported in the literature.4,46−49 We see
Figure 3. Selection of images showing preference of graphene nucleation for different surface defect sites: (a) nucleation at grain boundaries and triple junctions (red) on 3 h annealed Cu sample; (b) nucleation on rolling marks on a 5 min annealed Cu substrate; (c) no nucleation observed at low energy twin boundaries in 5 min annealed Cu substrate (yellow); (d) polycrystalline graphene layers on 2 h annealed Cu whose boundaries are decorated with Cu (graphene grain boundary indicated by blue arrow). Scale bar in all images: 2 μm.
instances of preferential nucleation at grain boundaries and triple junctions (Figure 3a) and aligned nuclei along rolling marks (Figure 3b). In comparison, to the preferential
Figure 4. Kernel average misorientation maps for the annealed samples: (a) as-received, (b) 5 min, (c) 2 h, (d) 5 h. The sharp reduction in pixel to pixel misorientation within grains between the 2 and 5 h annealed samples is clearly seen. Most misorientation after 2 h of annealing is confined to the grain boundaries of Cu. Figures S8 and S9 of the Supporting Information have more information on texture and grain size, which increases sharply between the as-received and 5 min annealed sample. Scale bar in (a) and (b): 100 μm. Scale bar in (c) and (d): 200 μm. D
DOI: 10.1021/jp510556t J. Phys. Chem. C XXXX, XXX, XXX−XXX
Article
The Journal of Physical Chemistry C
These results in the previous two sections thus show that both nucleation density and Cu surface defect density decrease with annealing times. We now show that there is a quantitative correlation between the two. Also, as shown in Figure S11 (Supporting Information), this observed preference is a kinetic and not a thermodynamic one. 3.4. Quantitative Correlation of Graphene Nucleation Density and Etch Pit Defect Density. This qualitative correlation between defects and nucleation is cemented further through the quantitative correlation between the nucleation density and etch pit density plotted in Figure 6a for the three samples annealed for 5 min, 2 h, and 5 h and grown at a CH4:H2 ratio of 5:1000. The drop in island density on annealing from 90 × 106 to 5 × 106 cm−2 is replicated by the drop in etch pit density from 30 × 106 to 5 × 106 cm−2. The ratio of the nucleation density to the etch pit density is plotted in Figure 6b. It is seen that in all three samples nucleation density is almost the same as the etch pit density, and their ratio falls in between 1 and 3. Even the departure from 1 of this ratio is in large part due to an underestimation of the etch pit density as discussed previously, especially for the 5 min annealed sample in which the dislocation cores are not as well separated. These data show that under these conditions graphene nucleates predominantly at the defects. Thus, the grains that constitute the final monolayer have their origins predominantly in nuclei that formed at defects. This apparently simple observation is a strong and significant correlation between the density of graphene islands and the defect concentration of the Cu foils on which they are grown. Subsequent analysis of the data in the Discussion on the variation in nucleation density with temperature as observed in these studies versus those published in the literature will bolster this thesis further. We now return to our previous observation of islands at the center of graphene grains that had a Cu dot sitting at the center. The preceding quantitative correlation clearly shows that the graphene islands that nucleate within the Cu grain do so at dislocations. The presence of the bright dot at the center of the islands is then due to Cu precipitation at these cores as well. Dislocations are the most potent nucleation sites for graphene growth. In order to verify this aspect, a 2 h annealed Cu foil, which according to EBSD data (Supporting Information, Figures S9 and S10) and TEM data (Supporting Information, section 11), would have been relatively free of rolling related deformation was indented to introduce an array of local islands of plastic deformation and hence dislocations. This sample subjected to graphene growth with 1 sccm of CH4. As can be seen from Figure 7, under this condition, graphene is nucleated only around the indent edges−plastic deformation zones−and nowhere else. This is a clear demonstration of the site selectivity possible for more controlled growth using the methods described in this paper. 3.5. Discussion. 3.5.1. Nucleation at Defects. The first question that arises is whether this site selectivity is a thermodynamic one or a kinetic one. If it were thermodynamic, it would mean that the energy barrier for nucleation at different sites becomes finite at supersaturations or CH4 partial pressures greater than a certain critical value. At supersaturations lower than this value nucleation can never be triggered at that site. As a reference, the minimum partial pressure required to nucleate bulk graphite from a CH4:H2 gas mixture can be easily
image map (OIM) of the Cu foils annealed for various times is presented in the Supporting Information (Figure S9). As can be seen with reference to the legend in Figure S9, only grain boundaries, predominantly low angle (
