J. Phys. Chem. C 2009, 113, 7541–7547
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Self-Ordering Behavior on Patterned Model Surfaces Gregory Grochola,* Ian K. Snook, and Salvy P. Russo Department of Applied Sciences, RMIT UniVersity, GPO Box 2476V, Melbourne, Victoria 3001, Australia ReceiVed: December 10, 2008; ReVised Manuscript ReceiVed: February 9, 2009
We use computer simulation to explore phenomena which could be used to grow arbitrary arrangements of nanostructures on metal surfaces via the vacuum deposition process. The method involves a three-stage process where, in the first “pattern formation stage”, a desired pattern of adatom islands is created on a base metal substrate, followed by a second “fixing stage” where an immiscible metallic species is vacuum deposited on the substrate at elevated temperatures to fill vacancy islands, and a final “growth stage” involving the further vacuum deposition of specific metal or nonmetal atomic species. In this work, we use molecular dynamic (MD) simulations to study the self-ordering behavior in the “fixing” and “growth” stage of assembly. We use the model immiscible systems Pt/Ag and Pt/Au with Pt as the base substrate. It was observed in the “fixing” stage that immiscible species deposited at higher temperatures would readily move off substrate adatom islands effectively “flooding” the vacancy areas around the islands, eventually forming a structurally stable, atomically smooth surface of alternating atomic species. In the “growth” stage, substrate species deposited at lower temperatures on such patterned surfaces were observed to preferentially move off surface regions covered by the immiscible species and onto the original adatom island areas, sometimes aggregating into nanostructures several atomic layers high. I. Introduction The process of building nanostructures using “bottom-up” synthesis techniques,1 where atoms spontaneously self-order or are made to order into desired nanostructures, relies critically on phenomena which transports atoms from one region of a substrate to another whereupon preferred aggregation can take place. The identification and control of such phenomena is paramount to progress in the field of fabrication by selfassembly. Stages of the formation process can take place under thermodynamic or kinetic conditions with the resulting structures depending greatly on the chosen environmental conditions, with kinetic growth conditions producing a greater variety of structures. Under kinetic conditions, knowledge of processes at the molecular level becomes very important. This is because almost all self-assembled nanostructures are formed one atom or one molecule at a time. More often than not, a single molecular level phenomenon can govern the morphology of the nanostructures formed. In this regard, molecular dynamics (MD) is an invaluable tool, in that using MD we can identify and explore key and potentially very useful molecular level processes as well as more collective dynamic behavior. An example of this was recently given in a work which simulated the growth of nanorods in an isotropic fluid with the use of model surfactants where collective dynamic behavior of a class of surfactants with specific properties was observed to be responsible for the anisotropic growth of the nanoparticle.2 In this work, it was shown how simple surfactant bonding characteristics can be used to induce nanorod growth. In another example of kinetically dominated nanoparticle growth, it was observed how a simple surface reconstruction mechanism3 transferred adatoms from the (111) surface facet of a decahedral (Dh) nanoparticle to its side (100) faces resulting in the formation of “pancake” Dh nanoparticles which are known to be thermodynamically unstable. * Corresponding author. E-mail:
[email protected].
A great deal of work has been done on substrate deposition with the aim of creating patterned nanostructures through selfordering behavior. Ag deposition on a Pt(111) substrate can be seen as a prototypical immiscible system. Adatom island formation for the Ag/Pt(111) system as well as for a number of other systems has also been studied extensively as a kinetic process; see, for example, the review by Brune4 or the work by Fichthorn and Scheffler.5 Gambardella et al.6 looked at the formation of Ag nanowires utilizing steps on the Pt(997) system as a template, finding that under certain conditions (at 400 K) smooth monolayer Ag strips attach adjacent to the Pt step edges. Ro¨der et al.7 showed that Ag atoms deposited on a Pt substrate with steps will not mix with Pt below 620 K; however, above 620 K Ag, atoms were found to intermix with Pt step edges forming 2D nanoclusters of Ag within the Pt surface structure near the step edge. Conversely, it was found that 2D Pt clusters also can form within Ag surface regions. Deposition of Ag on a patterned substrate consisting of two layers of Ag on top of a Pt substrate has also been performed.8,9 In this work, the top monolayer of Ag atoms formed an ordered strain relief array of altering fcc and hcp domains. Interestingly, it was shown that Ag adatoms formed self-organized clusters on fcc regions. The ordering behavior was based on preferential adsorption of adatoms on fcc domains and a high domain barrier in going from the fcc to hcp regions. However, we are not aware of any attempt to subsequently deposit Pt atoms onto such patterned substrates after a monolayer of Ag has been deposited. Such a deposition technique has been used to “build up” structures in other systems, typically to build CaF2 nanostructures such as nanowires and rows of nanodots on a Si substrates after it have been covered by a monolayer of CaF1.10 Hence, it would be very interesting to study the formation of Pt structures on surfaces patterned with Ag regions. The ultimate aim of the present work is to identify and explore phenomena which can be used to transport vacuum-deposited metal atoms (or potentially other nanostructure building blocks
10.1021/jp810866f CCC: $40.75 2009 American Chemical Society Published on Web 04/14/2009
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such as, for example, other nonmetal atoms or small nanoparticles) on a metal substrate to various neighboring regions through preferential binding. The model systems used here were the immiscible Ag/Pt and Au/Pt systems; however, since these systems behaved in similar ways we believe the findings here are applicable to other metallic immiscible systems. In this work, we investigated two stages of the proposed growth process and show both stages to be theoretically feasible. The “growth” stage of the process involves the deposition of the desired building blocks for the nanostructures which in this case consisted of Au, Ag, or Pt metal atoms. The previous “fixing” stage involved the construction of an atomically smooth prepatterned substrate and is also shown to be feasible; however, our main results are concerned with the self-ordering properties of deposited atoms on the constructed substrate and not how to construct such a substrate. In this regard, the method does depend on the formation of adatom island patterns, which we did not look at in this work. However, such patterned substrates can be constructed as thin film growth has been extensively studied and well understood. In the next section, we outline the MD methods and the interatomic potential used. In our discussion section, we detail and identify the various processes and their origin from an interatomic perspective. II. Methods Two types of MD simulations were performed in this work. The first involved standard MD where substrate atoms were allow to move under the influence of their neighbors, while the second involved a static array of substrate atoms upon which moving atoms were deposited. Experimentally, deposition times are much slower than that which can be reproduce in molecular dynamics simulations. Static lattice simulations were used as they allowed for at least an order of magnitude slower deposition times and/or larger simulation cells to be used. The details of these simulations are given below. A. Full Molecular Dynamics Simulations. 1. Cell Setup and Deposition. Clean and adatom island patterned Pt(111) and Pt(100) substrate slabs 5-6 layers thick were set up and allowed to equilibrate in an NVT ensemble. Lattice constants for the Pt system were obtained in separate bulk NPT simulation and used for the slab simulations to obtain near-zero pressure. Periodic boundary conditions were used at the ends of the slab. Vapor atoms were introduced into the simulation cell above the substrate and were given random coordinates and a velocity directed toward the surface. Two types of depositions were performed. The first which we term the “fixing” stage involved the deposition of Ag or Au atoms on substrates containing adatom or vacancy islands. This was done to explore the formation of the immiscible species monolayer on adatom patterned surfaces. The second set of simulations (of main interest in this work) which we term the “growth” stage involved setting up ideal prepatterned Pt/Ag and Pt/Au surfaces followed by the deposition of Pt atoms onto these surface. 2. Molecular Dynamics Details. We used a Broughton thermostat11 to generate an NVT phase space in between insertion events and the Verlet algorithm to integrate Newton’s equations of motion, with an MD time step of 2.0 fs. Deposition times are given in the figure captions for each system. B. Hybrid Static Lattice and Molecular Dynamic Simulations. 1. Cell Setup and Deposition. First, prepatterned Pt/Ag and Pt/Au substrates were set up in the same manner as above, making sure the correct zero pressure substrate lattice constants
Grochola et al. were used. Full MD simulations were used to obtain averaged relaxed atomic positions of the substrate lattice at the relevant temperatures. These positions were then used in the static lattice simulations. Static lattice atoms interacted with the deposited atoms using the full EAM interatomic potential, but were not allowed to move during the simulation. Two complications naturally arise as a result of a static substrate, the first of which involves the temperature of the deposited atoms. For strongly interacting metallic systems, a large amount of kinetic energy is released as a result of the deposition. Normally, deposited atoms quickly transfer this kinetic energy into the substrate; however, a static lattice has no such transfer mechanism, resulting in very hot adatoms with an unrealistically high diffusion rate. We attempted to overcome this problem through the use an Andersen stochastic thermostat to regulate the temperature of the deposited atoms upon contact with the substrate but found the thermostat to limit surface diffusion as it is known to do.12 This was found to be inconsistent with the full MD simulations or with real deposition where the temperature of the deposited atom and the local substrate atoms take a finite time to equilibrate after an impingement event. For our simulations we chose to use the Boughton velocity scaling thermostat for the movable atoms coupled with a single event velocity scaling procedure to bring newly deposited atoms to the substrate temperature within a designated time, specifically 1 ps. The procedure involved scaling the velocity of the deposited atom by a factor of 0.9 at every time step for up to 1 ps as soon as the impinging atom comes within intermolecular potential range of the substrate up to the point at which the atom’s temperature is at the desired input temperature. The time of 1 ps was chosen by analyzing the instantaneous temperature of deposited atoms using full MD simulations directly after initial contact with the substrate where it was found that 1 ps was a sufficient length of time to bring the instantaneous temperature of a deposited atom to the desired input temperature. Once the deposited atom is at this temperature its velocity is not altered again except through standard Boughton temperature scaling. Such a scaling procedure ensured that deposited atoms did not have an artificially high diffusion rate directly after attachment. Observation of simulations confirmed that single adatoms stayed on the static lattice for approximately the same amount of time as was observed for full MD simulations. While this scaling procedure may seem crude, it was sufficient for our purposes and should not have an impact on the diffusion characteristics of the system. This is because adatom diffusion on the (100) surface was only activated after adatoms aggregated into dimers, trimers, and larger clusters at some later stage far from the initial impingement event (see discussion section below). A second complication for the static lattice simulations is the lower entropic effects experienced by the deposited atoms. For example, surface interfacial free energies between the aggregated nanostructures and the substrate are higher for a static substrate than they are for a dynamic lattice. This resulted in nanostructures which did not resemble those observed for the full MD simulations. To take this effect into account, we chose to scale all lattice interactions with the deposited atoms to mimic entropic effects. We scaled the interactions until the nanostructures behaved in the same manner as the full MD simulations. A scaling factor of 80% was found to produce structures very similar to full MD simulations, although a value between 75 and 85% produced similar results, with the higher scaling values naturally producing slightly higher, more compact nanostructures.
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TABLE 1: Calculated and Fitted Data for the Alloy Systems AgPt and AuPta AgPt (AB) (eV/atom) AuPt (AB) (eV/atom) HoF A3B (L12) HoF AB (L10) HoF AB3 (L12) DLHoS B in host A DLHoS A in host B Pt scaling factor used
0.11 (0.10) 0.14 (0.18) 0.10 (0.06) 0.41 0.33 1.7
0.068 (0.08) 0.093 (0.11) 0.072 (0.06) 0.26 0.27 1.45
a Shown are the actual heats of formation (HoF) and dilute limit heats of solution (DLHoS) for our alloy AgPt and AuPt systems. The first values are the actual model values while the values in parentheses are the experimental values taken from Takizawa et al.15
C. Interatomic Potential. We refitted the cross-interaction potentials using the Johnson EAM potentials13,14 for individual species and the Johnson’s pair potential scheme. For the cross potentials we fitted the heats of formation for alloy structures and dilute limit heats of solution. With the Johnson scheme only a single fitting value, namely the relative scaling of the two electron density functions, need be fitted to obtain the cross pair potential if the EAM potentials for two species are known. As a first step, all potentials are normalized in the standard way, i.e., such that the equilibrium electron embedding density for each atom in an fcc lattice at zero pressure is 1.0 and the gradient of the embedding energy function for an embedding density of 1.0 is equal to 0.0. The procedure then involved assuming an electron density scaling factor for one of the species and calculating the heats of formation of the A3B and AB3 alloys with an L12 crystalline structure and an AB alloy with an L10 structure, and the dilute limit heats of solutions for species A in a metal B and vice versa. The electron density scaling factor alters the alloying strength of the elements. If the above heats of formation did not match experimental15 and theoretical values,16 then a new electron density scaling factor was used until a good fit was obtained. The dilute limit heats of solution were monitored to make sure they were consistent with what other EAM potentials predict. A complete set of heats of formation and dilute limit heats of solutions, obtained in this work, is given in Table 1 for the AuPt and AgPt system. III. Results and Discussion A. (111) Surface Behavior. 1. Surface-Embedded Immiscible Atoms. The first self-ordering phenomena observed involved the behavior of deposited Ag atoms on a substrate with Ag atoms embedded in the first layer of a Pt substrate. Ag atoms were observed to move away from the embedded Ag surface and bind preferentially to the Pt surface. This resulted in a flux of Ag atoms moving away from Ag regions toward Pt regions. Such deposited atoms were seen to be effectively confined by the strips and aggregated into elongated monolayer islands in the regions in between the Ag strips shown in Figures 1 and 2. This effect was observed for both the Ag/Pt system and the Au/Pt system, but it tended to be more efficient for the Ag/Pt system. Experimentally, the same confinement effects have been observed by Gambardella et al.6 resulting in the formation of Ag nanowires on Pt(997) substrate. We believe the preference which deposited atoms Au or Ag have to the Pt surface regions stem from two effects. First, the surface energy of the system can be lowered by covering the high surface energy Pt surface with low surface energy Au or Ag species; see, for example, Tsong et al.17 Second, through a metallic bonding argument, namely, we can note that the Johnson EAM potential function predicts Pt atoms to have a
Figure 1. Elongated Ag adatom monolayer formations formed via the vacuum deposition of Ag on a substrate containing two 5 atom thick Ag strips running in a [110] direction embedded in the first layer of a Pt surface. The simulation was carried out at 473 K, using a deposition time of 0.4 ns.
Figure 2. Elongated Ag adatom monolayer formations formed via the vacuum deposition of Ag on a substrate containing two 8 atom thick Ag strips running in a [112] direction embedded in the first layer of a Pt surface. The simulation was carried out at 473 K, using a deposition time of 0.4 ns.
higher relative bonding electron density function,18 at any particular radius r as compared to both the Ag and Au atoms; see Figure 3. Furthermore, the Pt system has a smaller lattice constant than both Ag and Au. Hence, Ag (or Au) atoms embedded in a Pt substrate experience a relatively high electron density environment, due to neighboring Pt atoms, and hence an embedding energy which is nearer to the minima. From these facts, Ag adatoms sitting on an Ag-covered substrate will not recover as much embedding energy for both itself and the Ag atoms underneath than if it was sitting on a pure Pt surface, and hence the preference for Pt regions. The same applied to Au, although the effect is weaker due to the smaller difference in electron density functions. 2. Internally Embedded Immiscible Atoms. Next we detail another surface ordering effect arising from the embedding of Ag atoms in the second layer of the substrate. This time the opposite to that discussed in section 1 takes place; namely, deposited Ag atoms are attracted to regions containing the Ag atoms in the second layer, forming elongated monolayer islands; see Figure 4. Again, the explanation of this interesting behavior comes from the differences in electron density functions. Ag atoms in the second layer create a lack of electron embedding density for the top layer of Pt atoms. These Pt atoms benefit more from the extra electron density supplied by the deposited Ag atoms
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Figure 3. Comparison of the Johnson EAM atomic electron density function for the three elements Ag, Au, and Pt. Note that Pt has a greater electron density function at any particular atomic radius r than Au and Ag, leading to the self-ordering effect observed in this work when Ag or Au is embedded in a Pt substrate.
Figure 4. Elongated Ag adatom monolayer formations formed via the vacuum deposition of Ag on a substrate containing two 3 atom thick Ag strips running in a [110] direction embedded in the second layer of a Pt surface. The simulation was carried out at 673 K, using a deposition time of 0.4 ns.
than other Pt substrate atoms which have Pt atoms embedded underneath. Again, the same effect is observed for the Pt/Au system. We can combine both phenomena into a more complicated substrate structure, by alternatively embedded Ag atoms in the first layer and the second layer. This resulted in a very efficient movement of deposited atoms toward regions with Pt atoms in the first layer and Ag atoms in the second layer. As an example we show in Figure 5 a substrate with a square mesh of Au strips embedded in the first layer and a square patch of Au atoms embedded in the second layer (not visible) situated in the center of the Au first layer strips. The structure again confines Au atoms which tended to move toward the square Pt regions with Au atoms embedded in the second layer aggregating into monolayer adatom islands. Figures 6a-c and 7a-c show the effect that deposition temperature had on the adatom island formation on various patterned surfaces. The figures show deposition temperatures of 473, 673, and 873 K. Pt atoms were also deposited on these surfaces with similar ordering behavior, although the binding of Pt atoms was much stronger to Pt regions, and hence the movement of deposited atoms away from the Ag embedded atoms was not as efficient as with Ag or Au deposited atoms, with Pt monolayer islands sometimes remaining in energetically unfavorable positions,
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Figure 5. Ag adatom monolayer formations formed via the vacuum deposition of Ag on a substrate containing 3 and 4 atom thick Ag strips running in an [110] and [112] direction, respectively, embedded in the first layer of a Pt surface. The substrate also has a 16 atom Ag square embedded in the second layer of the substrate at the center of each square region (not visible). The simulation was carried out at 673 K, using a deposition time of 0.4 ns.
namely on the immiscible surface regions. We can speculate that this self-ordering behavior is general in nature and should apply to other metal species. B. (100) Surface Behavior. For the (100) surface Ag atoms in the first substrate layer also created similar self-ordering behavior as was observed on the (111) surface, namely, deposited atoms tended to move off these regions and onto Pt regions. However, for the (100) surface we mainly deposited Pt atom as these atoms tended to actually build up much more readily than that observed on the (111) surface. Figures 8 and 9 show the result of the deposition process. Exchange events where Ag atoms in the first substrate layer are ejected onto the surface and swapped with Pt atoms were a problem at higher temperatures. Exchanged Pt atoms tended to provide attachment points on the Ag-covered regions for other deposited Pt atoms, eventually resulting in the bridging of nanostructures (see Figure 9). However, it was found that these exchange events could be exponentially suppressed by lowering the deposition temperature. For example, in the simulations shown in Figure 8, a temperature of 400 K resulted in only a single swapping event. Experimentally, as mentioned before, the deposition process takes place much slower, and hence temperatures could be taken much lower than what could be achieved here due to computational time constraints on MD simulations. 1. Diffusion Mechanism on the (100) Surface. Interestingly, Pt atoms tended to diffuse across the Ag-covered surface through a different mechanism than was observed for the (111) surface. Single-atom diffusion on the (100) surface was observed to be very slow as is well-known to be the case;19 however, as Pt atoms formed dimers, trimers, and larger clusters, diffusion activity increased. The trimer cluster tended to be mostly equilateral and hence tended to be unstable on the square (100) surface since an equilateral triangle is incommensurate with the lattice hollow site positions for the (100) surface. Tetramers also tended to form rhombic arrangements with some atoms sitting off hollow lattice sites as opposed to a square arrangement with all four atoms sitting on hollow lattice sites. Movement of atoms to Pt regions tended to take place via a bridging process to Pt regions, whereby a gradual “pulling” of small clusters into Pt surface regions was observed, be they dimers, trimers, or larger clusters, almost akin to cluster coalescence events on the
Nanostructures on Metal Surfaces
Figure 6. (a-c) Effects of different deposition temperatures on the Ag adatom monolayer formations formed via vacuum deposition of Ag on a substrate containing two 3 atom thick Ag strips running in a [110] direction embedded in the first layer of a Pt surface. All figures show identical time points and deposited numbers of Ag atoms. Panel a shows multiple adatom island formations forming at 473 K, while at 673 K in panel b the increased mobility of the deposited adatoms created larger but fewer nucleated adatom islands. In panel c, carried out at 873 K, we see single large aggregation of Ag adatom for each enclosed Pt strip. The simulations were carried out using a deposition time of 0.4 ns.
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Figure 7. (a–c) Effects of different deposition temperatures on the elongated Ag adatom monolayer formations formed via the vacuum deposition of Ag on a substrate containing two 3 atom thick Ag strips running in a [110] direction embedded in the second layer of a Pt surface. All parts show identical time points and deposited numbers of Ag atoms. Panel a shows multiple elongated adatom island formations nucleating at 473 K, while at 673 K in panel b the increased mobility of the deposited adatoms created long near-ideal nanostrip adatom formations. At 873 K, in panel c, we again see single long strips of Ag adatom aggregations attached to regions which have Ag atoms embedded in the second layer; however, at these temperatures the adatoms islands are much more disordered. The simulations were carried out using a deposition time of 0.4 ns.
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Figure 8. Pt nanostructure formed via the vacuum deposition of Pt on a Pt substrate which was covered with a monolayer of Ag atoms except for a single square patch of 6 × 6 Pt atoms embedded in the first layer of the surface (not visible). Note that the figure contains four periodic simulation cells for visual clarity. The simulation was carried out at 400 K, using a deposition time of 1.6 ns.
Figure 9. Pt nanostructure formed via the vacuum deposition of Pt on a Pt substrate which was mostly covered with a monolayer of Ag atoms but contained four equally spaced square patches of 3 × 3 Pt atoms embedded in the first layer of the surface (not visible). Note that the figure contains four periodic simulation cells for visual clarity. The simulation was carried out at 400 K, using a deposition time of 3.2 ns.
surface. This process was not always completely successful in removing Pt adatoms from the Ag-covered surface; sometimes small Pt clusters did not have enough time to move to the Pt areas before they were ultimately stabilized by further impinging Pt atoms on the Ag surface. This resulted in the formation of bridging structures (see Figure 9). It should be noted that the deposition rate used here is 7-8 orders of magnitude faster than that performed experimentally. Hence, given the many orders of a magnitude time difference, we can have confidence that more ideal nanostructures would form experimentally even at lower temperatures, since ideal nanostructures situated on Pt surface regions were more thermodynamically stable due to the above preferential bonding arguments. 2. Static Lattice Simulations. Static lattice simulations were performed to show the types of structures which could potentially be growth for larger systems. Figure 10 shows a hollow square arrangement with a nanoparticle at the center. The
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Figure 10. Shows Pt nanostructures formed via the vacuum deposition of Pt on a static Pt substrate covered in first layer Ag atoms but containing an outline of a square made of first layer Pt substrate atoms with a single square patch made of 3 atom thick Pt atoms again embedded in the first layer of the surface (not visible). Note, the figure above contains four periodic simulation cells for visual clarity. The temperature of the moveable atoms was 400 K, using a deposition time of 0.8 ns.
Figure 11. Shows Pt nanolines formed via the vacuum deposition of Pt on a static Pt substrate containing strips of 4 atom thick Ag atoms embedded in the first layer of the surface (not visible). Note, the figure above contains four periodic simulation cells for visual clarity. The temperature of the moveable atoms was 400 K, using a deposition time of 0.8 ns.
structure was not perfect since there was an unwanted bridging event between the central nanoparticle and the square nanostructured wall. However, again given the many orders of a magnitude slower deposition times which takes place experimentally we can have some confidence that near perfect structures can be obtained in the laboratory. Figure 11 shows strips of Ag atoms and the subsequent build up of channels. C. Formation Process of the Atomically Smooth Monolayer: “Fixing” Stage. In separate work we briefly studied the formation of the patterned slab structures used above. First, we set up Pt adatom islands on a slab substrate and then deposited Ag atoms onto the substrate. At low temperatures Ag atoms rarely moved off the top of adatom islands. However at higher temperatures, namely 600-800 K we observed that impinged Ag atoms would readily move off Pt adatom islands and into the vacancy island areas, completely filling in the large area vacancy. Figures 12 and 13 show some examples of this. It was observed that Ag edge atoms within this vacancy space had very high surface diffusion coefficients. Furthermore as the Ag aggregation nears complete filling of the vacancy space the
Nanostructures on Metal Surfaces
Figure 12. Results of Ag vacuum deposition on a Pt substrate containing vacancy islands. Ag adatoms were observed to readily move off the Pt ledge into the vacancy islands, subsequently filling these islands. Note that the figure contains four periodic simulation cells for visual clarity. The simulation was carried out at 873 K, using a deposition time of 3.2 ns.
J. Phys. Chem. C, Vol. 113, No. 18, 2009 7547 of Ag or Au atoms embedded in the first and/or second layers of the substrate. Metal species were observed to preferentially move off any surface regions where the first layer consisted of immiscible species, namely, Ag or Au and onto Pt areas aggregating into nanostructures several layers high for (100) surface and into monolayer adatom islands for the (111) surface. Conversely, deposited atoms tended to preferentially aggregate on top of surface regions with immiscible atoms embedded in the second layer of the substrate. The effect was found to be due to changes in lattice site absorption energies resulting from the fact that Pt atoms have a higher relative bonding electron density function at any particular atomic radius than both Ag and Au atoms embedded in the substrate. Such atoms deposited on the substrate move toward surface regions which gained the most energy from the electron density of the deposited atoms and move off first layer embedded Ag and Au regions which benefit less from the deposited atom’s electron density. The construction of the ordered alloy base substrate surface was found to be favorable at higher temperatures (600-800 K) using vacuum deposition. Deposited immiscible species were readily observed to move off the top of adatom islands preferentially binding to the outer walls of the island and effectively “flooding” the vacancy area around the islands forming structurally stable atomically smooth surface regions of alternating atomic species. Acknowledgment. This work was supported by the Australian Partnership of Advanced Computing (APAC), Victorian Partnership of Advanced Computing (VPAC), the Australian Research Council (ARC), and the RMIT Platform Technologies Research Institute. References and Notes
Figure 13. Results of Ag vacuum deposition on a Pt substrate containing adatom islands. Ag adatoms were observed to readily move off the Pt adatom islands, flooding the vacancy space and subsequently filling this space completely. The simulation was carried out at 873 K, using a deposition time of 1.2 ns.
remaining vacancies had very high surface diffusion coefficients. Hence, even if the last remaining Ag atom was by chance deposited on a Pt region it would still more likely step down into the Ag layer as it came in contact with the mobile vacancy creating a closed surface. However, as observed here, Ro¨der et al.7 found at temperatures above 620 K some intermixing of Ag into the Pt island takes place at the step edge of the Pt island with Pt atoms becoming displaced (see Figure 12 and 13). Hence, slightly lower temperatures than 620 K may be needed to prevent the pollution of Pt islands with Ag defects. This may affect the movement of Ag adatoms off Pt islands, however, a small amount of Ag atoms which may be trapped on Pt islands (depending on the size of the Pt island) are not expected to interfere with the “growth” stage, namely with subsequent Pt deposition any small number of Ag atoms would simply get integrated into the growing Pt nanostructures. The integration of such Ag atoms was observed here in cases where there were exchange events with Ag atoms been ejected as adatoms. IV. Conclusion We explored the behavior of vacuum-deposited metal atom on certain patterned Ag/Pt and Au/Pt substrates containing areas
(1) Bottom-up Nanofabrication. Supramolecules, Self-Assemblies, and Organized Films; Ariga, K., Singh Nalwa, H., Eds.; American Scientific Publishers: Los Angeles, 2008. (2) Grochola, G.; Snook, I. K.; Russo, S. P. Chem. Phys. 2007, 127, 194707. (3) Grochola, G.; Russo, S. P.; Snook, I. K. J. Chem. Phys. 2007, 127, 224705. (4) Brune, H. Surf. Sci. Rep. 1998, 31, 121. (5) Fichthorn, K. A.; Scheffler, M. Phys. ReV. Lett. 2000, 84, 5371. (6) Gambardella, P.; Blanc, M.; Brune, H.; Kuhnke, K.; Kern, K. Phys. ReV. B 2000, 61, 2254. (7) Ro¨der, H,; Schuster, R.; Brune, H.; Kern, K. Phys. ReV. Lett. 1993, 71, 2086. (8) Brune, H.; Giovannini, M.; Bromann, K.; Kern, K. Nature 1998, 394, 451. (9) Bromann, K.; Giovannini, M.; Brune, H.; Kern, K. Eur. Phys. J. D 1999, 9, 25. (10) Himpsel, F. J.; Kirakosian, A.; Crain, J. N.; Lin, J.-L.; Petrovykh, D. Y. Solid State. Commun. 2001, 117, 149. (11) Broughton, J. Q.; Gilmer, G. H.; Weeks, J. D. J. Chem. Phys. 1981, 75, 5129. (12) Frenkel, D.; Smit, B. Understanding molecular simulation; Academic: New York, 1996. (13) Johnson, R. A. Phys. ReV. B 1989, 39, 12554. (14) Johnson, R. A. Phys. ReV. B 1990, 41, 9717. (15) Takizawa, S.; Terakura, K.; Mohri, T. Phys. ReV. B 1989, 39, 5792. (16) Sluiter, M. H. F.; Colinet, C.; Pasturel, A. Phys. ReV. B 2006, 73, 174204. (17) Tsong, T. T.; Ng, Y. S.; McLane, S. B. J. Chem. Phys. 1980, 73, 1464. (18) Note that the electron density function in the EAM formalism is representative of the electron density available for bonding, and not the total atomic electron density. (19) Ayrault, G.; Ehrlich, G. J. Chem. Phys. 1974, 60, 281.
JP810866F
