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Step-Doubling at Vicinal Ni(111) Surfaces Investigated With a Curved Crystal Maxim Ilyn, J. Enrique Ortega, Ana Magaña, Andrew Leigh Walter, Jorge Lobo-Checa, Dimas G Oteyza, and Frederik Schiller J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b11254 • Publication Date (Web): 25 Jan 2017 Downloaded from http://pubs.acs.org on January 26, 2017

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The Journal of Physical Chemistry

Step-doubling at Vicinal Ni(111) Surfaces Investigated with a Curved Crystal Max Ilyn,†,‡ Ana Magaña,¶ Andrew Leigh Walter,§ Jorge Lobo-Checa,∥,⊥ Dimas G. de Oteyza,‡,# Frederik Schiller,† and J. Enrique Ortega∗,†,¶,‡ Centro de F´ısica de Materiales CSIC/UPV-EHU-Materials Physics Center, Manuel Lardizabal 5, 20018-San Sebastian, Spain, Donostia International Physics Centre, Paseo Manuel de Lardizabal 4, 20018-San Sebastian, Spain, Departamento F´ısica Aplicada I, Universidad del Pa´ıs Vasco, 20018-San Sebastian, Spain, Brookhaven National Laboratory, Photon Sciences Directorate, NSLS II, Upton, New York, 11973, USA, Instituto de Ciencia de Materiales de Aragón (ICMA), CSIC-Universidad de Zaragoza, E-50009 Zaragoza, Spain, Departamento de F´ısica de la Materia Condensada, Universidad de Zaragoza, E-50009 Zaragoza, Spain, and Ikerbasque, Basque Foundation for Science, 48011 Bilbao, Spain E-mail: [email protected]

∗

To whom correspondence should be addressed Centro de F´ısica de Materiales CSIC/UPV-EHU-Materials Physics Center, Manuel Lardizabal 5, 20018San Sebastian, Spain ‡ Donostia International Physics Centre, Paseo Manuel de Lardizabal 4, 20018-San Sebastian, Spain ¶ Departamento F´ısica Aplicada I, Universidad del Pa´ıs Vasco, 20018-San Sebastian, Spain § Brookhaven National Laboratory, Photon Sciences Directorate, NSLS II, Upton, New York, 11973, USA ∥ Instituto de Ciencia de Materiales de Aragón (ICMA), CSIC-Universidad de Zaragoza, E-50009 Zaragoza, Spain ⊥ Departamento de F´ısica de la Materia Condensada, Universidad de Zaragoza, E-50009 Zaragoza, Spain # Ikerbasque, Basque Foundation for Science, 48011 Bilbao, Spain †

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Abstract Vicinal surfaces may undergo structural transformations as a function of temperature or in the presence of adsorbates. Step-doubling, in which monatomic steps pair up forming double-atom high staircases, is the simplest example. Here we investigate the case of Ni(111) using a curved crystal surface, which allows us to explore the occurrence of step-doubling as a function of temperature and vicinal plane (miscut α and step type). We find a striking A-type ({100}-like microfacets) versus B-type ({111}-like) asymmetry towards step-doubling. The terrace-width distribution analysis performed from Scanning Tunneling Microscopy data points to elastic step interactions overcoming entropic effects at very small miscut α in A-type vicinals, as compared to B-type steps. For A-type vicinals, we elaborate the temperature/miscut phase diagram, on which we establish a critical miscut αc = 9.3◦ for step-doubling to take place.

Introduction Due to their lower atomic coordination, step atoms are active sites in surface chemical reactions. 1 This makes stepped (or vicinal) surfaces model systems, e.g., to investigate heterogeneous catalysis at the atomic scale. 2–4 In real catalysts the structure of the surface frequently changes during reaction conditions, and it is of fundamental importance to understand the nature of such structural transformations. Under ultra high vacuum (UHV) conditions, vicinal surfaces can also be disrupted, i.e., the canonical monatomic step ladder is often unstable at certain temperatures or upon adsorption, which may induce step bunching or faceting. Atomic reconstructions of high symmetry terraces trigger phase separation of reconstructed and unreconstructed terraces, which occurs under critical temperatures, e.g., on Si, 5 within certain miscut ranges, or in Au, 6 or induced by adsorbates, e.g., on Si 7,8 or Rh. 9 During overlayer growth structural transformations are driven by lattice matching, as shown for a variety of compounds grown on Cu vicinals, 10–13 which induces faceting into lattice-matched Cu vicinal planes. 2 ACS Paragon Plus Environment

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The simplest structural change that a vicinal surface may undergo is the formation of twoatom high steps. The appearance of step-doubling in semiconductors was deeply discussed in the past, due to its key influence in epitaxial growth. 14,15 But it is also of fundamental importance in catalytic reactions on transition metal surfaces, as shown for the CO oxidation of vicinal Pt(111), where double steps are more active than single steps. 16 In fcc surfaces vicinal to the (111) direction step-doubling is frequently linked to the presence of adsorbates or impurities that prefer to attach to microfacets with distinct orientation. 17–24 In the case of vicinal Ni(111), step-doubling is triggered by minute amounts of adsorbates, such as O2 , 22,23 or by temperature-reversible migration of bulk C impurities. 20 More recently, spontaneous step doubling has been claimed in the clean Ni(775) surface at ∼600 K, 24 although such observation has been later contested. 25 Step-doubling at a clean vicinal surface is energetically favourable when certain conditions are met, 14,26 however the existence of a canonical, temperature-dependent phase transition is more questionable. 15,26 Here we investigate stepdoubling in vicinal Ni(111) as a function of both vicinal angle (density and type of steps) and temperature using a curved Ni crystal. Low Energy Electron Diffraction (LEED) and Scanning Tunneling Microcopy (STM) experiments show the existence of massive, periodic step-doubling in A-type steps, with a miscut angle above a critical threshold. We demonstrate the potential of the curved surface approach to provide the key insights into this fundamental phenomenon.

Results The Ni(111) curved crystal, in short c-Ni(111), is polished with cylindrical shape around the (111) plane (Fig. 1 a), leading to vicinal surfaces with arrays of steps along the [1-10] direction, which separate (111)-oriented terraces. The cylindrical section angle α ∼ 27.5◦ defines the total miscut around the (111) surface, allowing one to examine both A-type ({100}-like microfacets) and B-type ({111}-like microfacets) step arrays at the left and the



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right sides of the sample, respectively. The reduced size (11×11 mm2 ) and large radius of curvature (23.2 mm) of the sample facilitates fast and homogeneous processing in vacuum, while providing easy access to the distinct vicinal directions with the coarse positioning of the STM tip. More importantly, the shallow curvature allows electron beam scanning in LEED, in order to address specific surface orientations. In fact, the ∼ 300µm spot size spreads over a ∆α ∼0.7◦ arc on the curved surface, which permits terrace-width d-dependent LEED analysis of Ni(111) vicinal surfaces. As sketched in Fig. 1 (a), in LEED experiments the sample is moved in front of the (fixed) electron beam to probe the different vicinal planes, defined by their miscut angle α with respect to the (111) surface. The LEED patterns in Figs. 1 (b) and (c) correspond to α= -12◦ and α= 13◦ in the B- and the A-side of the crystal, respectively, acquired at the same temperature T =300 K. The electron energy was set to Ep = 92 eV,leading to out-of-phase scattering conditions for monatomic step arrays, 27 in which the spots of the main hexagon appear neatly split. The size of the splitting with respect to the separation of the main (111) diffraction spots immediately gives the step spacing 2π/d in atomic row units 2π/a⊥ , (a⊥ = 2.16 ˚ A being the row spacing perpendicular to the steps). Despite having nearly the same miscut, one can immediately notice a drastic difference between patterns acquired at opposite sides of the crystal. Instead of the simple splitting of the main spot, in the A side we observe a single bright spot accompanied by weak satellites, which indicate a larger step array periodicity at the A side (d=19 ˚ A) compared to the B side (d=9.5 ˚ A). This is due to the presence of arrays of two-atom high steps 2h (h=2.03 ˚ A) in the former, which leads to the doubling (2d) of the average terrace size (proved in the STM experiments below), and hence to the transition to in-phase interference conditions [see the Supplementary Information (SI)]. In order to test step-doubling over the full range of miscuts we shift the sample in front of the LEED system in 0.5 mm steps, which roughly corresponds to ∆α=1.25◦ variations. During such sample scan, the (111) surface direction is kept parallel to the incident electron


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beam at Ep = 92 eV. If the height of steps is constant, the out-of-phase scattering conditions are fulfilled all across the whole curved surface (see SI), allowing one to observe the homogeneous d-dependent spot-splitting upon surface scanning. Fig. 1 (d) is made with line cuts made across the (1,0) spot [white boxes in Figs. 1 (b) and (c)] in all the LEED patterns across the curved crystal. For the B side of the crystal, the intensity plot reflects the linear variation of the spot spacing expected for vicinal surfaces made of monatomic step arrays with smoothly varying step density, i.e., surfaces with negligible step bunching or faceting. At the A side, a new, intense spot, labeled D, emerges between the split pair (labeled S) for miscuts larger than αc ∼9.3◦ . Therefore, the terrace-size doubling occurs for miscut angles larger than αc . Note that the D spot does not shift in the wave vector scale, indicating in-phase electron interference. This switch of the scattering conditions under the same geometry and beam energy is in fact expected upon step-doubling (see SI). By means of STM we analyze the step structure in the c-Ni(111) crystal at a local scale. Figs. 2 (a-d) show characteristic images taken at the A-side of the crystal, away (α=1.9◦ and α=3.1◦ ), right below (α=7.3◦ ), and above (α=10.1◦ ) the critical miscut αc for step doubling. For the B-side of the crystals, the STM experiments are discussed in the SI. Along some selected profiles perpendicular to the step arrays, we mark the presence of individual steps of double (D) and triple (T ) atomic height on top of (b-c) panels, whereas in panels (c-d) we also mark single steps (S ). In Fig 2. (e) we represent the statistical probability (counts) of a given terrace width d through the corresponding terrace-width distribution (TWD) histogram, which is obtained from the automated processing of STM images, as explained elsewhere. 28 Note that due to the occurrence of both single and double steps, we consider the local step height variation to determine the miscut α in each STM image. From the computed α value we determine the ”effective” mean terrace width < d >= h/ sin α marked with solid red lines in the histograms. < d >, which would be the average terrace width if all steps were monatomic, helps to rationalize STM histograms. At very small miscuts only monatomic steps are seen [Fig. 2 (a)]. Its corresponding TWD in Fig. 2



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(e) exhibits a skewed shape that peaks slightly below < d >=6 nm, indicating that nearly all steps are monatomic. At a somewhat higher miscut in Fig. 2 (b), triple steps appear randomly. The TWD maximum appears close to < d >=3.8 nm, but new features emerge near 2 < d > and 3 < d >, which are marked with dashed red lines. Approaching the onset of step-doubling [Fig. 2 (c)], double steps dominate over single- and triple-atom steps. The combined effect leads to a TWD maximum that is already shifted near 2 < d >=3.2 nm. Above the transition, more than 80% of the sample exhibits double-steps [Fig. 2 (e)], with single steps appearing randomly. The TWD exhibits a clear maximum near 2 < d >=2.3 nm, with a minor peak slightly below < d >. In conclusion, STM results in Fig. 2 agree with the LEED observations of Fig. 1, adding interesting structural information around αc , particularly the progressive substitution of single steps by double steps, rather than a sharp single-to-double step transition. 15,26 The lineshape of the TWD of a vicinal surface reflects the balance of different types of step-step interactions, mostly the so-called entropic interaction, related to the (effectively) forbidden crossing of two steps moving in one-dimension, and the elastic repulsion, which builds up as steps approach each other at high miscuts. 29–31 In this context, we have recently shown that tracking the TWD over a curved surface allows one to readily visualize the switch over from the entropic to the elastic regime, enabling a more accurate way to investigate such universal transition. 32 The TWD analysis of the STM data for the c-Ni(111) crystal is presented in Fig. 3 as an intensity plot. The probability data are normalized to the maximum value in each individual histogram, and the horizontal axis expressed as ”normalized” terrace width d/ < d >. The ticks mark the maximum probability in each of the histograms and the solid red line is a linear fit to the histogram maxima as a function of the miscut α (see SI for a larger scale). The vertical white line marks the d/ < d >=0.94 point, i.e., the peak maximum for a Wigner-Surmise line of order ϱ = 4, which adequately describes the TWD statistics with both entropic and elastic interaction effects. 31 As in the case of c-Pt(111), 32 the TWD maximum steadily shifts towards d/ < d >=1 as the step density increases, which


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can be interpreted as a transition from a purely entropic to a purely elastic interaction regime in step lattices, as discussed below. From Fig. 3 it is qualitatively clear that this entropic-to-elastic transition happens at a smaller miscut in the A-step side of the crystal. In fact, assuming the linear shift of the TWD maximum, the horizontal dotted line marks the miscut at which the probability maximum in each side reach d/ < d >=0.94. Taking d/ < d >=0.94 as the critical switch-over point from the entropic to the elastic regime, such transition occurs at α ∼3.8◦ in the A-step side of the crystal, and beyond α=8◦ in the B-step side (see also SI). Therefore, the critical step-doubling miscut αc ∼9.3◦ in the A-step side corresponds to step arrays submitted to a regime of rather strong elastic interactions, compared to the entropic repulsion effects. The temperature dependence of step-doubling is tested in Fig. 4 at different miscuts above αc =9.3◦ . In Figs. 4 (a) and (b) we compare the temperature dependent splitting of the (1,0) LEED spot for B- and A-type vicinals at similar α values. In the B side [Fig. 4 (a)], the double S spot is kept throughout the whole temperature range, but the splitting increases by ∼10% above 450 K. This effect can be explained by the increasing contribution of the entropic term in the surface free energy, leading to an apparent reduction in the step spacing measured in LEED (which is smaller than < d >, 33 ). The behavior is notably different for the A side, with the extra spot D emerging below a critical temperature Tc ∼550 K. For α=15.5◦ and α=10.5◦ in the A side [Figs. 4 (c) and 4 (d)], both above αc , Tc remains visually close to 550 K. Both αc (Fig. 1) and Tc (Fig. 4) are more accurately determined from the intensity analysis of individual line scans across the (1,0) spot shown in Figs. 5 (a) and (b), respectively. For αc , we have analyzed the intensity decay as a function of miscut observed in the rightside S peak of Fig. 1 (d). As indicated by the dashed lines, the S intensity dramatically drops between α=8◦ and α=10.5◦ , such that the critical αc =9.3◦ value is defined. In reality, the S peak also belongs to the double step lattice (second diffraction order), and would not strictly be the right ”order parameter” to track the step-doubling transition. 26 However, the



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emergence of the D peak as the miscut increases in Fig. 1 (d) is counterweighted by the strong decrease of the (1,0) spot intensity away from the (111) surface direction (see SI). Thus, there is a large uncertainty in determining αc from the α-dependent D intensity curve. This does not happen when measuring the Tc -dependence of D at fixed α, as shown in Fig. 5 (b). All curves show a sharp drop starting at ∼450 K, with a complete peak quenching at 550-600 K. Finally, we summarize the step-doubling process in vicinal Ni(111) with the T − α phase diagram of Fig. 5 (c). Data points correspond to the intensity analysis of the (1,0) spot of Figs. 5 (a) and (b), assuming the midpoint of the linear drop marked with dashed lines in each panel. To mark the boundary region between single and double steps, we draw a solid line that smoothly tracks the data. The panel spans the whole T − α range here explored for A type vicinals only. In the α axis we reach a maximum value between two well-defined symmetry directions, namely (335), featuring three-atom-row wide terraces, and (112), with two-atom-row wide terraces.

Discussion Understanding step-doubling at vicinal surfaces, and particularly the nature of the phase transition, 15,26 requires a systematic research of crystals, as a function of temperature and miscut. Ni surfaces, where step-doubling has been claimed in a variety of cases, 20–24,34,35 are good candidates. The first important question that arises is whether step-doubling in Ni surfaces is inherent to the pristine crystal or it is triggered by impurities. This is difficult to be answered, since small concentrations, e.g., of bulk C or vacuum adsorbates are difficult to avoid. 21 In the present case, the proper estimate is limited by the experimental environment and the detection capabilities to )max = 0.94. A shift in the TWD peak can only be explained when step interactions deviate from the ∼ 1/d2 dipole dependence. Moreover, if step-step interactions change from attractive, to purely entropic and to purely repulsive elastic, one expects the TWD to evolve from a WS-0 function (peak at (d/ < d >)max = 0.70), to the WS-2 (peak at (d/ < d >)max = 0.88), and to a gaussian function ((d/ < d >)max = 1). 31 The TWD



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shift in Fig. 3 strongly suggests that repulsive elastic interaction terms beyond the 1/d2 dipole contribution enter into play in monatomic step arrays (α < αc ), leading to a sizeable increase in the ”effective” elastic interaction (elastic/entropic) when approaching αc . We may use the TWD peak position to qualitatively evaluate the effective elastic interaction at the single-to-double step transition. One should note that step-doubling in the A-step side occurs in the presence of a strongly dominant elastic repulsion between monatomic steps, since the TWD at the transition onset (αc ) is significantly shifted above the canonical WS-4 position ((d/ < d >)max > 0.94). In contrast, for B steps the TWD does not shift over (d/ < d >)max = 0.94 in the miscut range examined in this work. In fact, in the B-side the step-doubling transition is not observed to occur, despite the minor density of double steps detected with STM at all miscuts (see SI). These could be induced by impurities, although the important fact is that they appear randomly distributed, and do not trigger the massive bunching observed in the A-step side of the sample. In summary, the step-doubling transition in vicinal Ni(111) has been examined using a curved crystal. For A-type steps we find a critical miscut and a temperature dependence for a complete single-to-double step transition, not observed in B-type steps. The statistical analysis of STM images shows a characteristic TWD shift in monatopic step arrays, which points to a miscut-dependent increase of the elastic/entropic interaction ratio. Such ratio is maximum at the transition onset in A-type step arrays, and it is significantly lower in B-type lattices. This results reveals the important role of the elastic interactions to trigger the transition from single to double step arrays in vicinal surfaces.

Methods Curved sample processing in Ultra High Vacuum The curved Ni(111) surface (Bihurcrystal Ltd., Spain) is obtained by mechanical erosion of a flat Ni(111) crystal, followed by mechanical polishing down to 0.25 µm grinding. In 10 ACS Paragon Plus Environment

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vacuum, the sample was extensively treated by alternating sputtering (1 keV Ar ion beam) and thermal flashing (800 ◦ C). To remove any possible C segregation, the annealing cycle was periodically performed in the presence of O2 (1×10−8 mbar, 10 minutes). Sputtering was carried out at 45

◦

incidence and parallel to the surface steps.

Low Energy Electron Diffraction Low Energy Electron Diffraction Experiments were performed with a standard three-grid setup (Omicron). Laterally-resolved scans across the curved surface were achieved by reducing the size of the incident electron beam (5◦ incidence from [111] crystal direction) to about 300 µm. Such measuring geometry minimized the spread of the beam over the curved surface, while reducing the curved surface scan to a simple sample translation. We estimate a probing miscut range on the sample ∆α < 0.7◦ , which falls slightly above the surface orientation accuracy commonly achieved in flat crystal samples. 6 We recorded individual LEED patterns at 25 different positions (∆α < 1.5◦ steps).

Scanning Tunneling Microscopy imaging and analysis STM images have been systematically recorded using a variable temperature STM setup (Omicron). The coarse movement of the STM piezos allows us to access different Ni surface orientations on the curved crystal. To image the curved surface we used a tunneling current of 0.3 nA, and a sample bias of +1V. The analytical process of the STM images has been explained in detail in Ref. 6 In summary, we perform a thorough analysis of individual frames with sizes between 200×200 and 50×50 nm2 , depending on the terrace width, using the WSXM software. 37 The STM analysis is limited to surface areas exhibiting homogeneous step arrays in the µm scale. STM images are then automatically processed, determining terrace width d values in each individual linescan in the STM frame, as well as the stepheight, which is automatically corrected to multiples of the nominal step height h=2.03 ˚ A. From the data, we directly obtain probability histograms, expressed as a function of discrete 11 ACS Paragon Plus Environment


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d values (in atomic row units a⊥ ). Each histogram of Fig. 2 is the average of, at least, three different frames taken at the same coarse position of the tip (1 µm2 ). The program corrects the effect of double steps to render the average miscut α at each d point, from which we determine the magnitude < d >= h/ sin α, defined as the average terrace size if all steps were monatomic. STM experiments were carried out at 300 K.

Supporting Information Supplementary information and data are provided for both LEED (geometry and reciprocal space sketch for the curved c-Ni(111) sample, quantitative LEED splitting variation across the curved sample, determination of the miscut angle scale) and STM experiments (analysis of the B-side of the c-Ni(111) sample, data plot of the terrace width distribution maxima across the c-Ni(111) surface).

Acknowledgement We acknowledge financial support from the Spanish Ministry of Economy (Grant MAT201346593-C6-4-P) and Basque Government (Grant IT621-13).

Contribution M. I., A. M., A. L. W., J. L.-C., D. G. de O., and F. S. performed experiments. M. I. and A. M. analyzed data. M. I. and J. E. O. designed research and discussed experimental data. J. E. O. wrote the paper. All authors thoroughly reviewed the article.

Competing financial interests The authors declare no competing financial interests.


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20. Hamilton, J. C.; Jach, T. Structural phase transitions in nickel at the Curie temperature. Phys. Rev. Lett. 1981, 46, 745–748. 21. Niu, L.; Koleske, D.; Gaspar, D.; King, S.; Sibener, S. Reconstruction kinetics of a stepped metallic surface: step doubling and singling of Ni(977) induced by low oxygen coverages. Surf. Sci. 1996, 356, 144 – 160. 22. Pearl, T. P.; Sibener, S. J. Oxygen driven reconstruction dynamics of Ni(977) measured by time-lapse scanning tunneling microscopy. J. Chem. Phys. 2001, 115, 1916–1927. 23. Pearl, T.; Darling, S.; Niu, L.; Koleske, D.; Gaspar, D.; King, S.; Sibener, S. Influence of oxygen dissolution history on reconstruction behavior of a stepped metal surface. Chem. Phys. Lett. 2002, 364, 284 – 289. 24. Usachev, D. Y.; Dobrotvorskii, A. M.; Varykhalov, A. Y.; Rybkin, A. G.; Adamchuk, V. K. Structural stability of stepped nickel surfaces. Phys. Solid State 2011, 53, 1277–1282. 25. Mom, R. V.; Hahn, C.; Jacobse, L.; Juurlink, L. B. LEED analysis of a nickel cylindrical single crystal. Surf. Sci. 2013, 613, 15 – 20. 26. Einstein, T. L.; Jung, T. M.; Bartelt, N. C.; Williams, E. D.; Rottman, C. Step doubling and related transitions on vicinal surfaces. J. Vac. Sci. Technol. A 1992, 10, 2600–2605. 27. Henzler, M. Atomic steps on single crystals: Experimental methods and properties. Appl. Phys. 1976, 9, 11–17. 28. Miccio, L. A.; Setvin, M.; M¨ uller, M.; Abad´ıa, M.; Piquero, I.; Lobo-Checa, J.; Schiller, F.; Rogero, C.; Schmid, M.; Sánchez-Portal, D. et al. Interplay between steps and oxygen vacancies on curved TiO2(110). Nano Lett. 2016, 16, 2017–2022. 29. Jeong, H.-C.; Williams, E. D. Steps on surfaces: experiment and theory. Surf. Sci. Rep. 1999, 34, 171 – 294. 15 ACS Paragon Plus Environment


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30. Joós, B.; Einstein, T. L.; Bartelt, N. C. Distribution of terrace widths on a vicinal surface within the one-dimensional free-fermion model. Phys. Rev. B 1991, 43, 8153–8162. 31. Einstein, T. Using the Wigner–Ibach surmise to analyze terrace-width distributions: history, user’s guide, and advances. Appl. Phys. A 2007, 87, 375–384. 32. Walter, A. L.; Schiller, F.; Corso, M.; Merte, L. R.; Bertram, F.; Lobo-Checa, J.; Shipilin, M.; Gustafson, J.; Lundgren, E.; Brión-R´ıos, A. X. et al. X-ray photoemission analysis of clean and carbon monoxide-chemisorbed platinum(111) stepped surfaces using a curved crystal. Nat. Commun. 2015, 6, 8903. 33. In principle, the peak in the TWD function (dmax ) should correspond to the d value measured in LEED, i.e., the spot spacing in LEED rather corresponds to 2π/dmax and not to 2π/ < d >. If the entropic contribution becomes dominant at high temperatures, the TWD becomes more asymmetric and dmax shifts towards smaller values(< d > stays constant). The same effect is shown in Fig. 3 at fixed temperature, but as a function of miscut, i.e., average spacing < d >. 34. Thapliyal, H. V.; Blakely, J. M. Reconstruction of stepped nickel surfaces. J. Vac. Sci. Technol. 1978, 15, 600–605. 35. Shen, Q.; Chang, J. P.; Navrotski, G.; Blakely, J. M. X-ray diffraction study of the Ni(111) vicinal surface. Phys. Rev. Lett. 1990, 64, 451–454. 36. XPS experiments performed in a different vacuum setup indicate a maximum 0.5% atomic concentration of C impurities within the probing depth of the XPS for C 1s electrons (∼10 atomic layers of Ni, 1200 eV kinetic energy). Such low concentration of impurities is frequent in XPS spectra from other nominally clean metal crystal samples measured at the same chamber, i.e., it could partly be due to a background signal from the detection system.


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37. Horcas, I.; Fernández, R.; Gómez-Rodr´ıguez, J. M.; Colchero, J.; Gómez-Herrero, J.; Baro, A. M. WSXM: A software for scanning probe microscopy and a tool for nanotechnology. Rev. Sci. Instrum. 2007, 78, 013705.



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Figures


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Page 19 of 24

Graphical TOC Entry 2d

15

15

10

10

2h

d

αC = 9.3°

h

5

incoming

2π/d

e-

0

5

A steps

0

B steps

diffracted e-5

B steps

d

A steps

-5

-10

-10

2π/d 0.2

-10 -5 0 5 10 15

Miscut angle α (°)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.4

0.6

0.8

Wave vector (× 2 π/a⊥)




(a)

2π/d

incoming e-

(d) diffracted e15

15

D B steps

10

d

αC = 9,3o

10

A steps 5

-10 -5 0 5 10 15

5


A steps

(b)

α = -12o, B steps

(c)

α = 13o, A steps

0

0

B steps -5

2π/d

-5

-10

S 0.2

0.4

-10

S 0.6


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0.8

Wave vector (× 2 π/a⊥) Figure 1: Low Energy Electron Diffraction (LEED) of the curved Ni(111) surface. (a) Schematic description of the Ni curved crystal and the LEED geometry. The crystal is horizontally moved to scan the indicated miscut range. A and B steps respectively have {100}- and {111}-like microfacets. (b-c) Representative LEED patterns taken at α = −12◦ and α = 13◦ , respectively, with electron energy Ep = 92 eV and the substrate held at 300 K. d Intensity mapping of the (1,0) LEED spot [vertically integrated area indicated in (b) and (c)] as a function of the miscut angle α. B steps show a linear increase of 2π/d splitting (S peaks) from the center of the crystal. The extra spot D that develops in the A-side indicates that step-doubling starts from α = 9.3◦ .


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(a)

(b)

α=1.9°°

T

α=3.1°° T T T D

(c) α=7.3°° D SD SS D D DD SS DS D S DD SS SD

(e)

T

S

T

T

D D D D D D DD SS SS

(d)

DS

α=10.1°°

D S DDSDDDDD

D DSS DD DD S SD SDDDS DDD

1

α = 1.9

0

α = 3.1

0

D

α = 7.3

0

0

α = 10.1

0 0

5

10

d (nm)

0

5

10

0

d (nm)

5

0

d (nm)

5

d (nm)

Figure 2: Scanning Tunneling Microscopy (STM) analysis of the A-side of the curved Ni(111) surface. (a-d) STM images taken at increasing miscuts. Random triple (T ) steps appear as the step density increases from (a) to (b). Double (D) steps become abundant close to the critical miscut in (c), and dominate over single steps S at high miscuts in (d). (e) Terrace width (d) distribution histograms for the respective images shown above. < d > (red line) stands for the ”effective” statistical mean terrace width, which is linked to the indicated local miscut by sin α = h/ < d >, where h is the monatomic step height (see the text). Dashed lines mark the successive 2 < d > and 3 < d > terrace widths. The histogram peaks at < d > in the two left panels and at 2< d > in the two right panels, in agreement with the LEED experiment of Fig. 1.



0

1

2

3

10

Double step

8 6

miscut angle α (deg)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Strongly elastic 4

Asteps

2

1.0

0 0

0.5

0.0

-2

Bsteps -4 -6 -8

Strongly elastic -10 0

1

2

3

d/ Figure 3: Evolution of the terrace width distribution across the curved Ni(111) crystal. The intensity plot is made from individual STM histograms like those shown in Fig. 2 (e), measured at different points across the sample. The left scale refers to the local miscut sin α = h/ < d >, where α is determined from the statistical analysis of each image. The ticks mark d values for the maximum probability in each of the histograms. The solid line is a linear fit of such d-maxima for each side of the crystal (see SI). The horizontal dashed lines mark the cross over from the entropic to the elastic regime at d/ < d >=0.94 (vertical dashed line, see the text).


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(a)

Temperature (K)

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(c)

S

B steps

A steps

α=-12o

α=15,5o

S

(b)

(d)

D

S

A steps

A steps

α=13o

α=10,5o

S

Wave vector (× 2 π/a⊥) Figure 4: Temperature dependence of the (1,0) spot of Fig. 1 for selected sample positions. In (a) and (b) we compare B-type and A-type vicinals, respectively, with α =12◦ and α =13◦ miscuts. S and D stand for the two split spots of the monatomic step array and the central beam in step-doubling, respectively, as in Fig. 1 (d). In (c) and (d) we show the temperature dependence for α =15.5◦ and α =10.5◦ miscuts in the A-step side.



(a) 1.0

(b) 1.0

Miscut angle α

D peak intensity

S peak Intensity

0

0.5

10.5 0 13.0 0 14,25 0 15.5

0.5

0.0

0.0 4

6

8

10

12

14

16

300

400

500

600

700

Temperature (K)

o

Miscut α ( )

(c) 800 (335)

700 Temperature (K)

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(112)

Single steps 600 500 400

Double steps

300 0

5

10 Miscut angle α (º)

15

20

Figure 5: Step-doubling phase-transition diagram for A-type Ni(111) vicinals. (a) S peak intensity evolution from individual (1,0) spot line scans of Fig. 1 (d). Error bars are indicated for a single data point, but are the same for all. (b) D peak intensity evolution from the individual (1,0) spot line scans of Fig. 4. (c) Phase-diagram for step doubling in A-type vicinal Ni(111), with phase-boundary data as determined from the inflection point in the intensity drop region in panels (a) and (b). The bars mark the width of such transition region in each case, as defined by vertical dashed lines in (a) and (b). The solid red line is a smooth guide-line that tracks the data.


Step-doubling at Vicinal Ni(111) Surfaces ... - ACS Publications

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