Induced Growth from a Ag Gas on Cu(111)

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C: Surfaces, Interfaces, Porous Materials, and Catalysis

Induced Growth from a Ag Gas on Cu(111) Carsten Sprodowski, and Karina Morgenstern J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.9b00478 • Publication Date (Web): 18 Mar 2019 Downloaded from http://pubs.acs.org on March 21, 2019

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The Journal of Physical Chemistry

Induced Growth from a Ag Gas on Cu(111) C. Sprodowski† and K. Morgenstern∗,‡ †Leibniz Universität Hannover, Institut f¨ ur Festk¨ orperphysik, Appelstr. 2, D-30167 Hannover, Germany ‡Ruhr-Universität Bochum, Lehrstuhl f¨ ur physikalische Chemie I, Universit¨ atsstr. 150, D-44801 Bochum, Germany E-mail: [email protected]

Abstract We investigate induced growth of a Ag layer on a Cu(111) surface by variable low temperature scanning tunneling microscopy between 100 and 140 K at submonolayer coverage. Without any interference by the scanning process, the Ag atoms from a twodimensional gas on the Cu(111) surface. Imaging the surface at elevated voltage induces nucleation and growth of one-dimensional Ag stripes of monolayer height eventually filling the surface of the imaged area completely. The stripes consist of rods of atoms with a preferential length of (1.88 ± 0.01) nm corresponding to approx. seven or eight Ag atoms on eight to nine Cu hollow sites. At a ratio of approximately 1:3 rods of double length are the second most observed species. The rods stack in the h112i √ directions at the 3 distance of Cu(111). Though all equivalent three surface directions are observed, their abundance is not equally distributed, such that the rod direction aligned with the main scanning directions predominates. At slow growth rates, it is possible to create a striped pattern with one surface direction, only.

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Introduction Although the scanning tunneling microscope (STM) was initially invented to image surfaces at atomic resolution, it was soon discovered that the STM tip, due to its close proximity to the surface atoms, influences and sometimes modifies the surface. 1 Despite being disadvantageous for faithful imaging such tip-substrate interactions were used deliberately to induce modifications of the surface, 2–4 e.g. by moving single atoms, 2,3 molecules 4 or defects 5 in a very precise way. 1 Direct interaction with the tip has been used to induce the motion of adatoms and molecules, 6–8 vibrational excitation to reversibly attach atoms to the tip, 9 and the large field in the STM tunnel junction, of the order to 108 V/m, to alter the surface structures, in particular molecules, their orientation, 10,11 and their conformation. 12–14 With respect to molecular switches, the different manipulation capabilities, i.e. field, chemical, current, were extensively explored not only on metals, 12,15 but also semi-conductors 16–18 and on insulating layers. 19 The creation of such and other artificial nanostructures by tip-sample interaction increased the understanding of surface physics and chemistry. 1 Several approaches also aimed at altering the surface at a larger scale. Quite early the reconstruction of the Au(111) surface was altered through formation of a surface hole by the tip. 20 On the same surface, needles were extracted from the step edges by scanning at some ten nA currents, i.e. in close proximity to the surface. 21 On a hydrogenated silicon surface, the nanoscale electron beam from the STM tip was used for desorption of H atoms, named nanolithography. 22,23 This method was also used to remove the native silicon dioxide from a Si surface 24 and to activate it by pre-adsorbed precursor metal-containing molecules, 25 to pattern a self-assembled monolayer by cross-linking, 26 and to polymerize bucky balls 27 in selected regions only. Recently, we demonstrated that positive voltage pulses lead to the nucleation of molecular clusters from a surface gas of phtalocyanine molecules on Ag(100). 28 Following the same approach, voltage pulses were used to form a supramolecular network of terarylene molecules on Cu(111). 29 Unexplored remains scanning aided growth, which might 2


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lead to specific nanoscale structures. In this article, we show that a few percent of a monolayer of Ag on Cu(111) form a mobile gas on Cu(111) in the temperature range between 100 K and 140 K, which only nucleates in the scanned area. The structures formed consist of short rods of atoms stacked in parallel, but without a strict alignment. Three different domains follow the main surface directions. Thereby, the two directions close to the slow and fast scanning direction are clearly preferred over the surface direction running at approximately 45◦ with respect to the scanning directions. We are thus able to form structures with only one growth direction.

Experimental Section STM measurements are performed with a fast scanning STM under ultra-high vacuum (UHV) conditions (2 · 10−10 mbar). The Cu(111) sample was prepared by cycles of Ar+ sputtering (1.3 keV, 3 to 5 · 10−5 mbar, 8 to 15 µA, 10 to 30 min) and annealing up to 970 K (10 to 45 min). Ag is deposited by resistively heating a short Ag wire attached for this purpose to a tungsten filament. The surface is held between 117 K and 150 K during deposition. The surface is exposed to the atomic beam for 53 to 72 s at slightly different fluxes. The fluxes are monitored by a quartz balance and identical coverages are ensured by always depositing an amount that equals a change of frequency of ∆ν = 10 Hz leading to a coverage of the order of 5% of a monolayer, apart from Figs. 3 and 8, where only half of the amount is deposited. During metal deposition the chamber pressure stayed below 3 · 10−10 mbar. The low deposition temperature makes alloy formation unlikely as Cu atoms detach neither from step edges nor from the terrace below 200 K. 30 Changes to the surface are induced and monitored in drift-corrected movies, where the same spot of the surface is imaged continuously. Scanning is thereby from bottom to top of the image. The STM is operated in constant-current mode. However, the program delivers

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for each image the average of the real currents at all pixels. Thus, the current value given varies slightly between subsequent images of a movie. We give the range of currents in the figure captions. All images of induced growth that are analyzed quantitatively are recorded with 256x256 pixel at an image size of 172 nm x 172 nm. The size per pixel is thus 0.45 nm2 . For continuous imaging, the interaction time per pixel (or nm2 ) is directly proportional to the recording time regardless of the scanning speed. We thus plot the decay curves of the remaining free surfaces versus the real time as reference.

Results and discussion

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Figure 1: Pristine Cu(111) surface, 0.84 nA, 782 mV, 120 K; inset: atomic resolution with surface directions marked, 0.52 nA, -33 mV.

The surface preparation leads to a surface with broad terraces of several 10 nm (Fig. 1). This image does not change after deposition of Cu in submonolayer coverage (not shown). At tunneling parameters below ± 2 V, there is no indication of the deposited Ag implying that the Ag atoms are highly mobile and do not nucleate at the measurement temperatures between 100 K and 140 K. This phenomenon is observed for a range of deposition temperatures 4


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Figure 2: Induced growth for Ag on Cu(111): (a,b) first and last STM images of a movie consisting of 101 images recorded at (137 ± 3) K; indicated times in min:s, 23.7 s/image, dwelling time per pixel: 0.36 ms; pixel size: 0.45 nm2 ; deposition at (148 ± 12) K for 53 s, 0.28 ... 0.29 nA, -2.1 V (c) amount of uncovered terrace area A vs. time t with fit of B + C · exp(− τt ) yielding B = −(43 ± 86) nm2 , C = (15472 ± 129) nm2 , τ = (622 ± 14)/s and fit for final decay < 3000 nm2 yielding B = (164 ± 49) nm2 , C = (25531 ± 2376) nm2 , τ = (547 ± 25)/s; inset: on half-logarithmic scale. However, imaging at larger voltages leads to nucleation and growth of a Ag layer. In the first image some islands are nucleated on the terrace and at the step edges (Fig. 2a). We image the same region around 100 times till the terrace is filled completely. For a terrace of 20.000 nm2 , the filling takes 100 times scanning the surface or a bit more than 40 min (Fig. 2b). The decay of the non-covered area is exponential which is obvious in the half-logarithmic plot (Fig. 2c, inset). Note that the decay exponent τ , obtained by fitting B + C · exp(−t/τ ) to the data, is around 10% smaller, if only the second half of the decay is fitted (Fig. 2c, yellow line). For better understanding the induced growth, we now first analyze the initial growth, then the formed structures and finally the time evolution at close to closure of the layer. Though such growth phenomena were observed between 1.2 V and -3 V, we here concentrate on the growth induced at -2.1 V. Scanning at gradually increased image size in the middle of a terrace facilitates to follow nucleation of a single nucleus on the middle of a terrace and its growth without interference 5



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Figure 3: Initial nucleation and growth on terrace: Images from a series at indicated times in min:s; rectangle marks region of previous image; total number of images: 21; pixel size: 0.2, 0.45, and 0.8 nm2 for two images each; image acquisition times: 27.3 .. 40.1 s; (102 ± 2) K; 0.28 - 0.29 nA, -2.1 V, deposition for 30 s at T =(103 ± 8) K; cross marks the Cuh110i directions, as deduced from the inset in Fig. 1. from nuclei at the step edges or elsewhere on the terrace (Fig. 3). The single nucleus is already branched after the first scan (Fig. 3a). The branches follow mostly the h112i directions, i.e. [11¯2], [1¯21], [¯211], [¯1¯12], [¯12¯1], and [2¯1¯1], of the surface lattice at 30o from the closed-packed h110i directions. The three branches lengthen upon the next scan (Fig. 3b) and branch further (Fig. 3c). Some of the novel branches follow the h110i directions of Cu(111). The branches also broaden in width (Fig. 3d). A further nucleus at the top left of the initial island joined it at the next image. Fig. 3d includes a part of a step edge, where another island nucleates, grows, and joins the island (Fig. 3e). Further branches either follow likewise the h112i or the h110i directions of the surface. These are the two main growth directions observed in general. On the last image, independent nucleation on several places is observed (Fig. 3f). Higher resolution reveals that the branches consists of stripes, often two or three in parallel (Fig. 4a). During further growth and filling of the terrace the original structure largely remains, i.e. the structures are stable on the time scale of several minutes (Fig.

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Figure 4: Induced growth for Ag on Cu(111): (a-c) three snapshots of a movie at indicated times in min:s, 0.27...0.41 nA, 2.1 V, (115 ± 1) K; deposition for 75 s at (125 ± 10) K; image acquisition time: 21.3 s; total number of images: 39; dwelling time per pixel: 0.33 ms 4b,c). Even higher resolution reveals that the stripe structure consists of rods of rather uniform length along the closed-packed h110i directions (Fig. 5a). The stripes are separated by (0.36 ± 0.03) nm in h112i directions of Cu(111), corresponding to the next-nearest neighbor, √ or 3, distance of Cu(111) (Fig. 5b, inset). A neighboring rod is placed thus at the closest possible distance with equivalent adsorption sites within the rods. A certain deviation from a rigid stacking, i.e. a waviness of the stripes, suggests a rather weak interaction between the rods. This allows a continuation of the stripes around obstacles, e.g. other stripes (e.g. Fig. 5a, 5th stripe from top). From roughly 4800 analyzed rods, the majority of 56% are close to the fast scanning direction, 32% are close to the slow scanning direction and the minority, at 12%, run at 45◦ to the two scanning directions. This large deviation from an equal distribution underlines that the structures are tip-induced. The length distribution reveals two maxima (Fig. 5b). The first one, at (1.88 ± 0.01) nm, corresponds to 6.5 lattice constants of aAg = 0.289 nm of Ag(111) and the second one, at (3.65 ± 0.03) nm, to 12.6 aAg . In terms of aCu , the first maximum corresponds to 7.4 aCu , the second one to 14.3 aCu . The first maximum thus corresponds to rods that consist of seven or eight Ag atoms on top of eight or nine hollow sites of the surface (Fig. 5b, inset); the second one to 13 to 14 Ag atoms on 14 or 15 Cu hollow sites. The total area of the first 7



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Figure 5: High resolution of stripe pattern: (a) fully covered surface, 0.39 nA, 2.1 V, 133 K (b) Length histogram of rods; inset: schematics of a stripe consisting of five rods; Cu atoms in orange, Ag atoms in grey; (c) surface above monolayer coverage, 0.35 nA, -2.1 V, 153 K (d) Imaging of low-temperature superstructure at 2.1 V; 0.36 nA. maximum is, at 3515 nm2 , three times larger than the one of the second maximum, at 1191 nm2 . The ratio between the rods of the two legths is thus 3:1 in favor of the shorter rods. The larger Ag keeps its lattice spacing along the rods and the length of the rods is determined by this match. We speculate that nucleation of of these rods without scanning is not hindered by repulsive interactions, as for instance for the S/Au(111) system, 31,32 but simply by kinetics, i.e. the binding energy is not sufficient to stabilize the structures in the investigated temperature range down to 100 K. From thi, we can estimate the binding energy to be smaller than 0.3 eV. Further scanning leads not only to a lengthening of the rods but 8


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also second layer nucleation (Fig. 5c). Note that individual rods show different apparent heights at 2.1 V (Fig. 5d). As this effect is tip dependent, we attribute it to an electronic effect.

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Figure 6: Induced growth of Ag on Cu(111), deposition for 72 s at (128 ± 15) K; Tmeas = (115 ± 1) K: (a-d) STM images at indicated times in min:s, total number of images: 48; image acquisition time: 18.75 s; dwelling time per pixel: 0.29 ms; pixel size: 0.45 nm2 ; I = 0.23 to 0.28 nA, V = -2.1 V (e,f) uncovered terrace area A vs. time t; red line is fit of exponential decay B + C · exp(−t/τ ) yielding B = (5167 ± 718) nm2 , C = (21210 ± 614) nm2 , τ = (652 ± 44)/s; blue lines are apparent linear fits to half-log plot yielding slopes of 1.27 · 10−3 and 6.3 · 10−3 , respectively; black line in (c) is fit from Fig. 2c. To further confirm that the structure formation in Fig. 2 is tip-induced, we investigated the filling of a terrace at the 10 K lower temperature than above in Fig. 6. Several small stripes nucleate at the step edges and in the middle of terrace (Fig. 6a). During further imaging at the same parameters, the stripes grow and branch, till they meet other stripes, creating a network with vacancies of different sizes (Fig. 6b). These vacancies fill gradually (Fig. 6c) till the whole scanned area is covered with a uniform Ag layer (Fig. 6d).

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Here, for a terrace of around 1.6 ·104 nm2 the filling takes around 50 times scanning the surface within 1000 s (Fig. 6e). This is faster than the case shown in Fig. 2. Nonetheless, the filling of the terrace, again fitted by B + C · exp(t/τ ), yields a very similar rate rate of τ = (652 ± 44)/s as above. The different filling time results from different preexponential factors approx. 15000 vs. 21000 nm2 . At the approx. 20 K lower measurement temperature than in Fig. 2, the filling rate does thus not decrease as expected for a thermally activated process. As compared to the filling of a terrace of similar size shown in Fig. 2, the filling is even faster; needing only half of the time. This suggests that the structures are only metastable and in dynamic equilibrium with the Ag gas phase with a higher detachment rate of the atoms from the growing structures at a higher temperature. This leads to an increase of the filling time with temperature. More clearly than in Fig. 2, the half log plot yields two exponential decays yielding decay exponents of (738 ± 236)/s and (37.7 ± 3.1)/s, respectively, crossing at 720 s or 3560 nm2 (Fig. 6f). (a)

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Figure 7: Final filling of terrace: (a-c) STM images from a movie recorded at (136 ± 1) K at indicated times in min:s, scanning time per image: 22 s; dwelling time per pixel: 0.34 ms; pixel size: 0.45 nm2 ; 24 images per movie; 0.28 .. 0.38 nA, -2.1 V (d) Uncovered terrace area A vs. time t with fit of B + C · exp(t/τ ) yielding τ = (434 ± 201)/s. To understand why the decay rate is smaller towards the end of the decay, we now analyze in more detail the final decay for the already branched structure (Fig. 7a). The vacancies decrease in size (Fig. 7b) till they are all disappeared (Fig. 7c). Here, fitting yields τ = (434 ± 201)/s (Fig. 7d). Also this exponent is smaller than the overall exponents for the two other cases, but again different from the final decay exponent of the other two. This 10


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suggests some variance in the final decay, which we attribute to the local structural variation in the remaining vacancies. These vacancies are different both, in size and in shape (Fig. 6c, Fig. 7a,b).

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Figure 8: Filling of individual vacancies: (a) STM image at t=0 s; islands are numbered; image acquisition time: 44 s, dwelling time per pixel: 0.7 ms; pixel size: 0.45 nm2 ; 0.32 nA, -2.1 V, T = (125 ± 3) K, deposition at (109 ± 9) K for 30 s ∆ν = 5 Hz; 44.4 s/ image; 64 images; dwelling time per pixel: 0.67 ms (b) Area A of vacancy island marked in (a) vs. time t (c) same as (b) on half-logarithmic scale (d) magnificaton of (c). To demonstrate our point, we plot the area of the vacancies in Fig. 8a individually (Fig. 8b-d). Note that some of the vacancies pinch off, for instance vacancy number 2 at its top and vacancy number 4 at the bottom, leading to small vacancies that are numbered 5 and 7 (Fig. 8a,b). These vacancies are rather small and follow the slow decay of such small islands (cf. to vacancy number 1). On a half-logarithmic plot the final decay for vacancies smaller than 500 nm2 is linear, i.e. A ∝log(t), (Fig. 8c) down to an island size of a few ten nm2 , at most, where the decay decelerates (Fig. 8d). The latter corresponds to a size, where rods of full length, can no longer be fitted into the remaining holes. Both deviations from a global filling rate are thus traced back to local structural variation.

The preference for growth in the fast scanning direction can be used to create a unidirectional pattern. We scan slower than for the data analyzed above and utilize a step at a slight angle to the scanning direction as nucleation center. From this step edge, we create rods in scanning direction (from bottom to top, Fig. 9a,b). The stripes are close to the slow scanning direction. In this way, we are able to create regions of several 100 nm2 with mostly 11



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Figure 9: Unidirectional growth: (a) 0.42 nA, -2.48 V, deposition 60 s at (118 ± 17) K, 75.95 s/image; 1.16 ms/pixel (b) 0.29 nA, -2.1 V, deposition 30 s at (110 ± 8) K, 86.22 s/image, 1.32 ms/pixel (c) 0.39 nA, 1.06 V, deposition 120 s at (144 ± 24) K, 27.3 s/image, 0.42 ms/pixel (d) 0.31 nA, -2.1 V, deposition at (109 ± 9) K for 30 s, 22 s/image, 0.34 ms/pixel (e) 0.31 nA, -2.1 V, deposition at (109 ± 9) K for 30 s, 33 s/image, 0.5 ms/pixel . a uniform direction (Fig. 9c). Finally, we zoom out after creation of some structures (Fig. 9d,e). As expected for an induced growth, the region next to the scanned region does not exhibit the striped pattern. As discussed in the introduction, STM induced changes can result from direct interaction with the tip, the current or the field in the STM tunnel junction. 1 Direct interaction demands a close tip-object distance, which is achieved by lowering the voltage and increasing the current. As the structures here are induced ABOVE a voltage threshold, chemical interaction is not the origin of the structures. Manipulation by the current results from inelastic tunneling of the electrons, i.e. they deposit their energy into the manipulated object, usually into vibrations. For atoms on a metal surface, possible vibrations are phonons at the surface or external vibrations of the adsorbate against the surface, i.e. frustrated rotation or translation vibrations. Both are in energy in the range of a few ten meV. Thus for direct 12


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excitation, the structures should form at much lower voltage than observed here. Vibrations can also be excited indirectly by electrons deposited into an electronic state, usually a molecular orbital. Such states do not exist for the metal/metal system. Here, the only states of the system would be field emission states that are, however, much higher in energy. 33 This leaves the electric field as the likely reason for the structure formation. The field can either polarize the adsorbate atoms or alter the potential energy surface. In the first case polarization induces an electrostatic interaction with the tip, confining the atoms in its vicinity and thus facilitating nucleation. In the second case, potentials that are deeper may hold the atoms together for long enough, so that other atoms attach and hinder the dissolution of the formed structure. As these two possibilities cannot be discriminated experimentally, we hope to induce theoretical investigations that clarify the process underlying the tip-induced nucleation and growth.

Conclusion We showed that mobile Ag atoms nucleate in the field of the STM tip to form stripes consisting of stacked rods with specific lengths, which reflects the large lattice mismatch between Ag and Cu. The rods line up in the perpendicular direction with some flexibility, leading to many stripes that eventually fill the whole scanned area. Our study thus presents a novel aspect of manipulation by tip-sample interaction, namely scanning-induced growth.

Acknowledgments This work was supported by the Research Training group Confinement-controlled Chemistry, which is funded by the Deutsche Forschungsgemeinschaft (DFG) - GRK2376 / 331085229.

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References (1) Morgenstern, K.; Lorente, N.; Rieder, K.-H. Controlled manipulation of single atoms and small molecules using the scanning tunnelling microscope. phys. stat. sol. b 2013, 250, 1671- 1751. (2) Eigler, D.M.; Schweizer, E.K. Positioning single atoms with a scanning tunnelling microscope. Nature 1990, 344, 524-526. (3) Stroscio, J.A.; Eigler, D.M. Atomic and molecular manipulation with the scanning tunneling microscope. Science 1991, 254, 1319-1326. (4) Meyer, G.; Neu, B.; Rieder, K.-H. Controlled lateral manipulation of single molecules with the scanning tunneling microscope. Appl. Phys. A 1995, 60, 343-345. (5) Wong, D.; Velasco Jr., J.; Ju, L.; Lee, J.; Kahn, S.; Tsai, H.-Z.; Germany, C.; Taniguchi, T.; Wanatbe, K.; Zettl, A.; et al. Characterization and manipulation of individual defects in insulating hexagonal boron nitride using scanning tunneling microscopy. Nat. Nanotechn. 2015, 10, 949-954. (6) Bartels, L.; Meyer, G.; Rieder, K.H. Basic steps of lateral manipulation of single atoms and diatomic clusters with a scanning tunneling microscope tip. Phys. Rev. Lett. 1997, 79, 697700. (7) Böhringer, M.; Morgenstern, K.; Schneider, W.-D.; Berndt, R. Scanning tunneling microscope-induced molecular motion and its effect on the image formation. Surf. Sci. 2000, 4457, 3750. (8) Trembulowicz, A.; Ehrlich, G.; Antczak, G. Surface diffusion of gold on quasi-hexagonal reconstructed Au(100). Phys. Rev. B 2011, 84, 245445. (9) Yang, J.; Nacci, C.; Marinez-Blanco, J.; Kanisawa, K.; Fölsch, S. Vertical manipulation

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Induced Growth from a Ag Gas on Cu(111)

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