Experimental study of In thin films on Pd(111) and alloy formation(Article)

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In-Pd and similar systems: top-layer substitutional alloys

Top-layer substitution alloying has been observed for In/Pd(111) at very low coverage. Using perturbed γ-γ-angular correlation (PAC) spectroscopy, Hunger et al. found 10−4 ML of 111In probe atoms deposited on Pd(111) at 80 K occupied mostly substitutional step sites upon annealing to 350 K and mostly substitutional terrace sites at 380 K. The authors claim at 380 K, the axial symmetry from substitutional terrace site occupancy is consistent with In atoms replacing Pd atoms in the top layer. The axial symmetry was extrapolated from the electric field gradient signal so a z-axis tilt could be assigned to a likely adsorption configuration at each given site. This site occupancy was deduced from a simple model inclusive of above terrace, step edge, step edge kink, substitutional step edge, and substitutional terrace In atom positionings. After annealing slightly past 500 K, only signals corresponding to In in substitutional terrace sites were detected [62].
To our knowledge, the above study of In/Pd(111) is the only one performed at coverages √ √ ≤ 1 ML but no substitutional ordered structures, such as Pd(111) ( 3 × 3 )R30◦−In have been reported. The existence of a Pd(111) (√ √ )R30◦−In substitutional 2D surface alloy seems plausible considering the substitutional behavior already observed at very low coverages [62] coupled with surface energies of other systems that exhibit this superstructure. For example, the surface energies of Cu/Sn, Cu/Sb, Pt/Sn, Ni/Sn, Pd/Sb are 1.85/.71, 1.85/.68, 2.55/.71, 2.45/.71, 2.1/.68 compared to 2.1/.69 [J/m2] for Pd/In [101]. This phase can be formed after annealing from over 0.33 ML ( 1ML) of tin deposited on Cu(111), Ni(111), and Pt(111) [6, 116]. Also, annealing a higher coverage √ √ of ∼3.6 ML for Sb/Pd(111), resulted in a ( 3 × 3 )R30◦ LEED pattern. Structural analysis of the closest candidate, Sb/Pd(111), revealed that substitutional Sb in the first √ √ layer formed the structure [19]. These authors also observed a ( 3× 3 )R30◦phase after depositing ∼ 3.6 ML of In on Pd(111) and annealing to around 530 K, but difficulties induced by indium diffusion seemed to have prevented any publications of quantitative findings [19, 117].
Unlike their In-Pd counterpart, Zn-Pd and Ga-Pd thin films on Pd single crystals have been studied from a perspective of IMCs in catalysis. However, structural information on top-layer substitutional alloys is vague in comparison to the more chemically interesting alloys and IMC formations of several layers. Annealing Zn films on Pd(111) or Pd(110) over 600 K results in a Pd-Zn monolayer. The former has been described as a Pd adlayer being substitutional in nature and having a (1×1) LEED pattern [164]. The monolayer for Zn/Pd(110) was identified as a Zn:Pd 1:1 NSIP having a p(2×1) LEED pattern [2, 150]. In both cases, the surfaces showed a low corrugation with a slight Pd-up/Zn-down buckling. Stadlmayr et al. suggested this was from increasing Pd in surface and second layers pulling Zn atoms inward (for Zn/Pd(111)) [151]. More recently, the authors observed similar buckling on Ga/Pd(111) and Ga/Pd(110) films. For Ga/Pd(110), as for Zn/Pd(110), only Pd-up/Zn-down buckling was observed. Annealing both Zn (Ga) films on Pd(111) to around 600 K showed a Zn(Ga)-up/Pd-down to Pd-up/Zn(Ga)-down corrugation change. The Pd-up/Ga-down buckling found on Ga/Pd(111) was not directly associated with a top layer alloy as opposed to Zn/Pd(111) findings. Instead, the authors describe this as a GaPd2 film [149].

In-Pd and relevant systems: multilayer alloys and IMCs

After depositing 4 ML of Ga on Pd(111) and annealing between ∼500 to 800 K, a bulk-truncated GaPd2 structure was observed by LEED. Moreover, the authors found annealing 1 ML to similar temperatures resulted in a weak c(4×2) pattern and alluded the structure could also be construed as GaPd2 [149]. These findings suggest that, unlike Zn/Pd(111) films of various thicknesses annealed to similar temperatures, a multilayer surface IMC (GaPd2) existed. This is not surprising since the cohesive energy of Ga in Ga-Pd is more than double that of Zn in Zn-Pd [149]; so much higher annealing temper-atures would be required to drive sub-surface Ga to the top layer prior to desorption.
A two-layer thick ZnPd film on Pd(111) was stable at lower annealing temperatures ( 550 K) in comparison to GaPd2/Pd(111) and revealed a (2×2) LEED pattern. The (2×2) LEED pattern resulted from 3-rotational domains of p(2×1) ZnPd. This was clearly observed and further understood by STM. Weirum et al. found growth of p(2×1) ZnPd bi-layer islands after dosing submonolayer to ∼2 ML of Zn on Pd(111) at room temperature. Above 2 ML, the LEED pattern transformed to a (1×1) phase indicative of the onset of bulk Zn growth over the ZnPd bilayer. The same authors also note that thermal stability of ZnPd/Pd(111) decreased with decreasing film thickness [164].
The results from the most recent In-Pd surface study suggest perpendicular inhomogene-ity in multilayer intermetallic compounds/alloys formed after depositing 4 MLE of In on a polycrystalline Pd film and annealing. After annealing to 363 K, InxPd1−x, with x = 0.63 in top layers and x = 0.51 in sub-surface layers formed, whereas In depletion upon annealing above 623 K leveled-out the concentration gradient (x = 0.19) [128]. The latter potentially compares well with multi-layer InPd3 films found in another In-Pd study [59].
A disordered fcc crystal structure with a lattice constant of 4.0028 ˚A is known to form on sputtered InPd3 thin films [59]. Such a configuration, with theoretical (111) atomic spacings of 2.8304 ˚A is close to Pd(111) bulk registry (2.75 ˚A). A multilayer of similar composition was found after annealing 4-5 ML of Sn layers on Pt(111). It was described as a stacking of fcc Pt3Sn(111) having a (2 × 2) periodicity with the bulk [48]. Bulk Pt3Sn(111) also shows a (2 × 2) phase, but can be reconstructed with a top-layer having (√ × √ )R30◦ ordering resulting from a slight depletion of tin in sublayers. This had the same composition as the Pt(111) (√ × √ )R30◦−Sn bidimensional surface alloy 3 3 [3, 6].
Likewise, the possibility of surface reconstructions on multilayer In-Pd alloys cannot be √ √ overlooked. This may explain the formation of the In/Pd(111) ( 3 × 3 )R30◦ phase reported by Carazzolle and Pancotti et al. [19, 117]. It is possible that In/Pd(111) could form both multilayer and top layer (√ ×√ )R30◦ structures, analogous to (2×2) phases 3 3 of Pt(111)-Sn [48].
Room temperature inter-diffusion of Pd and In has been studied in several amorphous/poly-crystalline multi-layer films, showing some consistency with Pd(100)-In alloys formed upon annealing In/Pd(100) surface films at low temperature ( 400 K). For example, (a) 5 nm and (b) 10 nm of Pd were deposited on an In/Pd covered SiO2 substrate in two separate experiments. The fixed In/Pd layers were respectively 50/100 nm thick.
AES depth profiling data showed 1:1 InPd at ∼35 to 45 nm (a) and within the top 5 nm (b). Before Pd deposition however, nearly symmetric gradients of Pd diffusion into the In layers and In diffusion into the Pd layers was observed at the In/Pd interface region of 50 nm [172]. The authors indicate that the alloy formed in the region above 50 nm was probably In3Pd [172, 173]. In3Pd (tetragonal lattice with a = 3.25 ˚A, c = 4.94 ˚A) [93, 132] was also observed using PAC spectroscopy by placing local 111In probes on a Pd(100) surface preceding 100 ML In deposition at liquid N2 temperature and annealing to 350 K; this was explained as an amorphous In3Pd interface layer [37]. A separate study using the same experimental conditions and method, showed no In-Pd inter-diffusion when annealing to 340 K, however inter-diffusion and formation of several Pd/In compounds occurred when annealing to 360 K. After annealing to 390 K, only In3Pd was detected, whereas further annealing to 450 K, for an extended period, showed a mixture of InPd2 and In3Pd2 [38]. More recently, for a Pd/In2O3 catalyst, In3Pd2 was formed during reduction at 673 K, and overlapped with In3Pd at T > 673 K [93] forming a bcc lattice with a = 9.433 ˚A [53, 93].
This bcc phase had the same lattice constant as In7Pd3; the most recently added inter-metallic compound to the In-Pd phase diagram (Ref. 1.3). In fact, it has been deemed as the corrected interpretation of In3Pd. This probably better explains some mixed or overlapped structures previously observed, as discussed in the last paragraph.

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Bulk IMC and alloy surfaces

Bulk alloy and IMC surface preparation under UHV often involves sputtering – annealing cycles to remove contaminates and obtain flat, reconstructed surfaces suitable for char-acterization by various experimental methods. During ion bombardment at the surface (sputtering), one of the elements can be removed from the surface at a faster rate than the other (preferential sputtering). Often, the element having less mass is ejected. Ac-cording to Bardi, the light element irreversibly depleted after extended ion bombardment for Pt80Co20(100) [6, 7], Pt80Fe20(111) [6, 11], and NiAl(100) [6, 105]. In NiAl(111), Al resegregated to the surface above ∼400 ◦C after being depleted from sputtering. Further annealing between 900 ◦C and 1000 ◦C resulted in a faint (√ × √ )R30◦ LEED 3 3 pattern [108]. A (√ × √ )R30◦ phase was also observed for Pt3Sn, but only after ion bombardment depleted subsurface tin [3, 6]. Furthermore, it was suggested that sputtering-induced loss of titanium of the selvedge region for Pt3Ti(100) was responsi-ble for a pure, rather than a mixed surface termination [6]; the bulk terminated planes of Pt3Ti(100), an AuCu3 – type ordered alloy, alternate between 1:1 Pt:Ti (mixed termina-tion) and 100 % Pt (pure termination) [4]. Excess Pt in Pt3Ti was expected to lead to Pt segregation [6, 148] and indeed was the case for Pt-enriched Pt3Ti(111) [6, 24, 120]. The preferentially sputtered species can be the heavier element when its bonding is weaker in the solid. Recently, Zn depletion was noted after sputtering an Al57Pd30Zn13 complex metallic alloy [177]. This being different than findings in Al-rich ternary quasicrystals where Al, the lightest element, was often preferentially sputtered [87, 177].
The notion of altered layer(s) from ion bombardment/segregation is a known phe-nomenon and can be attributed to athermal and thermal processes. At or below RT, the athermal processes of displacement mixing and radiation-induced segregation cause the preferentially sputtered species to reach the surface. On the other hand, radiation-enhanced segregation and Gibbsian segregation are thermal processes. Gibbsian segre-gation increases the likelihood of the ejected element also being the segregated species [168]. Although not stated by the authors, this may explain the resegregation of Al while annealing NiAl(111) [108]. It seems that annealing a preferentially sputtered sur-face first establishes a local equilibrium between the altered surface and bulk regions. After annealing to ∼60 – 70 % of the melting temperature, bulk diffusion occurs and equilibrium with the bulk is established [168]. Here, equilibrium segregation can be ob-served and interface defects vanish. For example, Pt25Ni75(111) STM showed a network of dark lines appeared from misfit dislocations between altered and bulk regions. They disappeared when reaching the annealing temperature of bulk-surface thermodynamic equilibrium [163, 168].
Relevant findings of preferential sputtering and/or surface segregation have not been reported for the M-Pd (M=Ga,Zn,In) IMC/alloy surfaces of recent interest. Sputtering bulk GaPd(111) and GaPd(111) single crystal surfaces did not result in ejection of a particular element [124–126]. For Zn-Pd and In-Pd bulk surfaces, to our knowledge, UHV surface preparation details regarding these phenomena are non-existent.

Quantifying in-house XPS data

Quantification of XPS data is usually not a trivial process. Many authors have devoted much of their careers developing better methods, to include refinements in predicting IMFPs, and determination/eradication of instrumental and analytic error [70, 81–85, 104, 141, 142, 154, 155, 159], to improve the reliability of concentration measurements. As shown later, in Eq. 2.20, the normalization to account for 100 % of the surface area sampled depends on the photoionization cross section σ and the IMPF λ(Ek). However, a dependence on an analyzer transmission function should also be considered (this was negligible for the synchrotron data). The transmission function is defined as follows:  ATF = a2 b a2 + R2 (2.16)
where R = (Ek/Ep) and a,b are parameters relevant to the analyzer. For the Sphera analyzer used in our experiments, a = 21.0025 and b = 0.2015 (provided by the manu-facturer). Ek and Ep are electron kinetic energy and analyzer pass energy respectively; normally, Ep was set at 10 eV. Other factors (OF) such as photon flux, angular efficiency based on the angle between detected electrons and X-rays impinging the surface, detection area, and photoelectron detection efficiency [61, 162] also contribute to the total intensity (Ii). Ii is defined as: Ii = ni · σi · λi · AT Fi · OF (2.17).

Table of contents :

1 Introduction 
1.1 Overview
1.2 Intermetallic compounds in heterogeneous catalysis
1.3 Methanol steam reforming (MSR)
1.3.1 The reaction and its importance
1.3.2 Recent interest in M-Pd (M=Zn, Ga, In)
1.4 Other interest in In-Pd intermetallics
1.5 Surface alloys
1.5.1 Intermixing and segregation
1.5.2 In-Pd and similar systems: top-layer substitutional alloys
1.5.3 In-Pd and relevant systems: multilayer alloys and IMCs
1.6 Bulk In-Pd phases
1.6.1 Elemental palladium and indium
1.6.2 Discovery of the phases
1.6.3 In-rich Phases
1.6.4 Pd-rich Phases
1.6.5 InPd
1.7 Bulk IMC and alloy surfaces
1.8 Summary
2 Experimental Methods 
2.1 Experimental Background
2.1.1 Experimental set-up
2.1.2 STM
2.1.3 LEED
2.1.4 XPS
2.1.5 Diamond Beamline I06
2.2 XPS data analysis
2.2.1 μXPS
2.2.2 Inelastic mean free path (IMFP) calculations
2.2.3 Quantifying in-house XPS data
2.2.4 Polycrystalline and bulk InPd surfaces
2.2.5 In/Pd(111) surface IMCs and alloys
3 Surface structures of In-Pd intermetallic compounds 
3.1 Overview
3.2 Part I: Experimental study of In thin films on Pd(111) and alloy formation(Article)
3.3 Part II: Theoretical study of surface energies, alloying and segregationeffects (Article)
4 Bulk InPd IMC surfaces 
4.1 Introduction
4.2 Growth of bulk InPd
4.2.1 InPd single crystals
4.2.2 Polycrystalline InPd
4.3 Single crystal InPd surfaces
4.3.1 Overview
4.3.2 InPd(100)
4.3.3 InPd(110)
4.3.4 InPd(111)
4.4 The polycrystalline surface
4.5 Discussion and Conclusion
5 Conclusion 
A Pendry R-Factor


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