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Thickness-dependent electronic structure and interfacial behaviors of iron on faceted MgO(1 1 1) films
Chemical Physics Letters 551 (2012) 92–95
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Chemical Physics Letters journal homepage: www.elsevier.com/locate/cplett
Thickness-dependent electronic structure and interfacial behaviors of iron on faceted MgO(1 1 1) films Mingshan Xue a,b,⇑, Qinlin Guo b,⇑ a b
School of Materials Science and Engineering, Nanchang Hangkong University, Nanchang 330063, China Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, P.O. Box 603, Beijing 100190, China
a r t i c l e
i n f o
Article history: Received 27 June 2012 In final form 13 September 2012 Available online 21 September 2012
a b s t r a c t Based on the vital effect of the interfacial behaviors on tunneling magnetoresistance in Fe/MgO/Fe junctions, the thickness-dependent electronic structure of iron on MgO(1 1 1) films with {1 0 0} facets was investigated. The results illustrated that the chemical interaction at the interface of Fe and MgO films was rather weak. Instead, a upward band bending of MgO with 1.2 eV was obviously observed with the increase of Fe thickness. The particle size effect was responsible for these shifts of core levels owing to the three-dimensional growth of Fe. These results were further testified by the data from the lattice vibration in Fe–MgO system. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction The interactions at metal-oxide interfaces have a considerable effect on the related optical, chemical, electronic and magnetic properties [1,2]. In recent years, much effort has been made to understand these interactions for the application of such materials in sensor, semiconductor, microelectronics, catalysis, magnetic random memory as well as spintronics [3,4]. For example, related studies have indicated that the chemical bonding at the Fe/MgO interface is of much disadvantage to the tunneling magnetoresistance (TMR) in the Fe/MgO/Fe magnetic tunnel junctions [1,5,6]. The oxide layers at the interface obviously impact the tunneling current between two ferromagnetic layers, resulting in the sharp decrease of the TMR value. In the Fe/MgO catalysts, the distribution of iron oxides at the interface also strongly affects the H2S removal ability for the catalytic wet oxidation of H2S [7]. These factors stimulate an increasing interest towards the investigation of Fe–MgO interface. However, the studies between metals and bulk oxide crystals are often confined because of the surface charging problem and the difficulty of sample heating and cooling originating from the insulator character of bulk oxides [8]. Instead, ordered oxide thin films grown epitaxially on single-crystal metal substrates (such as Mo, W, Pt substrates) have been widely used [9,10]. As a low cost material with excellent mechanical, thermal and electronic properties, ordered MgO films grown on metal substrates are widely used as a template to investigate the interactions between metals and oxide surfaces, such as Ni/MgO, Ag/MgO and Zn/MgO [5,9,11,12]. One hand is in that the material with a high ⇑ Corresponding authors. Fax: +86 791 86453210. E-mail addresses: [email protected] (M. Xue), [email protected] (Q. Guo). 0009-2614/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.cplett.2012.09.034
stoichiometry is easily prepared due to the high oxygen affinity of magnesium. Another hand is in that it has a simple rock-salt structure, easily forming ordered films on single-crystal metal substrates. When Fe and MgO layers are fabricated into Fe/MgO/Fe magnetic tunnel junctions, these epitaxial devices exhibit large TMR ratios (e.g., 318% at 10 K) though they are still considerably lower than the values predicted theoretically [13]. The discrepancy between experimental and theoretical results is mainly attributed to the interfacial structures, such as the relative position of the atoms, interfacial oxidation, strain, and structural asymmetry of the interfaces [2–8,13]. In order to further clarify the detailed factors affecting the TMR, the evolution of surface/interface electronic structure at the initial growth stage of Fe on MgO surface should be taken into consideration. In previous works [2–9], the neutral (1 0 0) face of MgO in Fe–MgO system was widely used to investigate the interfacial behaviors of Fe and MgO. In spite of their efforts, whether or not a chemical reaction exists at the Fe–MgO interface is still controversial [2,6,14]. In this Letter, we try to construct a new system of Fe–MgO(1 1 1) with {100} facets, which is chemically more active than general MgO(1 0 0) surface. As a typical polar surface, MgO(1 1 1) films consisting of alternative stacking of Mg layers and O layers have a strong chemical activity due to a low coordination of surface atoms. Because of the so-called ‘polar instability’, the (1 1 1) surface of MgO is easily diverged into {100} facets (including (1 0 0), (0 1 0) and (0 0 1) faces) [15,16]. As a result, asformed surface is characteristic of (1 1 1) face covered by {100} facets, but it has a slightly higher activity than pure (1 0 0) surface of MgO films owing to the exposure of part atoms from (1 1 1) face in the valley region between {100} facets. Such a MgO surface may provide a stronger evidence for the absence of chemical bonding at the Fe–MgO (1 0 0) interface if there is no reactive layer observed
M. Xue, Q. Guo / Chemical Physics Letters 551 (2012) 92–95
in the present Fe–MgO (1 1 1) system. In our experiments, such a MgO (1 1 1) surface with {100} facets is easily prepared on Mo(1 1 0) substrate under ultrahigh vacuum (UHV) condition [12]. X-ray photoelectron spectroscopy (XPS) is used to obtain the first indication of the electronic structure and chemical states of Fe and MgO films. High resolution electron energy loss spectroscopy (HREELS) and low-energy electron diffraction (LEED) are used to monitor their lattice vibration and surface structure. The experiments were performed in a UHV chamber which had a base pressure of less than 5 10 10 mbar. The chamber was equipped with LEED, XPS and HREELS. The XPS analyses were carried out using a Mg Ka X-ray source (hv = 1253.6 eV) with a pass energy of 50 eV. The binding energy (BE) was calibrated with respect to the pure bulk Au 4f7/2 (BE = 84.0 eV) and Ag 3d5/2 (BE = 368.3 eV) lines, and the accuracy on the core-level displacement was better than 0.15 eV. In HREELS measurements, the primary electron incident energies were set to 4.9 eV and the typical resolution was 12–13 meV obtained by the full width at the half maximum (FWHM) height of the elastic peak from the MgO(1 1 1) films. The magnesia films were epitaxially grown on a single-crystal Mo(1 1 0) substrate. The substrate surface was treated by annealing at about 1200 K in 10 7 mbar O2 environment, followed by a subsequent flash to 1500 K until no impurity was detected by XPS. The magnesium and iron sources were made of a pure magnesium ribbon (purity >99.9%) and iron wire (purity >99.99%) wrapped tightly around a tungsten wire, respectively. The two sources were thoroughly degassed by thermal treatment before preparation. The deposition rates of Mg and Fe were about 0.15 and 0.08 monolayer (ML)/min by means of XPS, calibrated via the relative intensity of Mg 2p3/2 and Fe 2p3/2 to Mo 3d5/2 lines as a function of deposition time, respectively. For example, in XPS measurements, to monitor the growth rate of Fe, the growth of Fe was checked by measuring the Fe 2p3/2 to Mo 3d5/2 intensity ratio as a function of deposition time, as shown in Figure 1. A straight line was fitted to the data for the initial growth of Fe on Mo(1 1 0) substrate. There was a clear breakpoint for the curve after growing Fe for 12 min, corresponding to 1 ML Fe film. It suggested that the growth of Fe on Mo(1 1 0) surface followed a Stranski–Krastanow (SK) mode: Fe initially grew as two-dimensional patches at submonolayer coverages, and then gradually formed three-dimensional islands on top of the patches. During growth, the deposition rate of Fe was kept constant by means of adjusting the output power of the constant-current direct supply. The MgO thin films with 15 ML were
Figure 1. Statistical XPS intensity ratio of Fe 2p3/2 to Mo 3d5/2 core levels as a function of thickness of Fe on Mo(1 1 0) substrate. The insets give the LEED pattern of MgO(1 1 1) films with {1 0 0} facets, Ep = 72 eV, and the structural schematic of the interface of Fe and MgO, respectively.
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prepared by evaporating Mg in an ambience of 1 10 6 mbar O2 at 673 K, followed by annealing at 800 K in 1 10 7 mbar O2. Then as-prepared MgO(1 1 1) films were used as the support to deposit Fe at room temperature (RT). All the experimental data were in situ measured at RT. The MgO(1 1 1) films with {100} facets were grown on Mo(1 1 0) substrate, which had been described in a previous study [12]. Besides the core-level lines originating from Mg and O, there are no other peaks observed by XPS such as these from C or Mo, indicating a clean MgO surface. The corresponding BE values of Mg 2p and O 1s lines are 50.8 and 531.2 eV, respectively, being in agreement with previous data of MgO films. The inset in Figure 1 gives a typical LEED pattern from the faceted MgO(1 1 1) films. The LEED pattern consists of two sets: the inner one is of hexagonal symmetry and the outer one is split and diffused in non-radical directions with the slight change of the primary energy [17,18]. It is similar to a cleaved bulk MgO (1 1 1) crystal surface, which is easily covered by a series of trigonal pyramids featured by three sets of {1 0 0} facets (consist of (1 0 0), (0 1 0) and (0 0 1) planes) inclined by 54.7° to the original surface normal caused by the polar instability of MgO (1 1 1) surface [17–19]. Furthermore, the formation of (1 0 0) facets on (1 1 1) face is in qualitative agreement with theoretic analysis of the stability of ionic crystal surfaces based on Madelung potential and lattice dynamics calculation [19,20]. Figure 2 gives the XP spectra of Fe 2p core levels as a function of Fe thickness. For Fe, XPS measurements can identify its different chemical states by observing the BE of Fe 2p core levels. Many studies have indicated that the BEs of Fe 2p3/2 core level for Fe0, Fe2+ and Fe3+ states are 707, 710.7 and 711 eV, respectively [21,22]. For initial 0.08 ML Fe, the corresponding BE value of Fe 2p3/2 line is 708.2 eV, as shown in Figure 2. The peak position is far from the BE values of Fe2+ or Fe3+ states (more than 710 eV), implying an absence of interfacial bonding between Fe and MgO(1 1 1) surface. With the increase of Fe coverage, the peak shifts gradually towards a lower BE until the value is 707 eV. The peak at 707 eV means a metallic Fe0 state because the Fe film is formed after growing several ML Fe. However, no obvious LEED pattern was detected from the corresponding Fe film, indicating a disordered Fe film. It may originate from the three-dimensional growth of Fe on MgO(1 1 1) surface, i.e., the small Fe clusters/particles are formed at the initial
Figure 2. XP spectra of Fe 2p core levels with the increase of Fe thickness on MgO(1 1 1) films. The vertical lines give the peak positions for initial and final thickness of Fe.
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stage of Fe deposition (as shown in Figure 1), being different from the case of Fe/Mo system. The particle size effect causes the shift of Fe core levels towards higher BE, resulting in a total shift of 1.2 eV from initial to final growth. Figure 3 shows the XP spectra of Mg KLL, 2p and Fe 3p core levels with the increase of Fe thickness. For Mg 2p and Fe 3p core levels, their peaks are very close so that it is hard to distinguish the shifts of their peaks. However, for Mg KLL spectra, a shift of 1.2 eV towards lower BE is obviously observed with the increase of Fe coverage from 0 to 5 ML. Similarly, the phenomena are observed in the XP spectra of O 1s lines (not shown). The shifts of core-level lines are attributed to the band bending at the MgO surface [23,24]. Due to the work-function difference of Fe and MgO following the Schottky–Mott theory, the bands bend upward towards the surface with a subsequent Fe deposition, resulting in the continuous decrease of BEs of Mg and O core levels [25]. Moreover, the same shift direction and magnitude of the Mg and O core levels support the interpretation of as-observed shifts of Fe 2p core levels as the pLetter size effect rather than chemical interaction at the interface. To completely exclude the existence of iron oxides at the Fe/ MgO interface, we further investigated the oxidation of 1.5 ML Fe films by XPS, as shown in Figure 4. For 1.5 ML Fe film (curve a), the BE of Fe 2p3/2 line is at 707 eV. After only exposing to 10 L (Langmuir, 1 L = 1.33 10 6 mbar s) oxygen at RT, the BE value is shifted to 710.7 eV, which is typically characteristic of iron oxides (i.e., Fe3O4). When the film is further oxidized at 10 6 mbar oxygen at 700 K for 20 min, the peak has almost no shift in BE, but has a better symmetry in the shape of the peak. The result reveals that Fe atoms are easily oxidized in an oxygen ambience, and most of 1.5 ML iron atoms deposited have been oxidized by 10 L oxygen [26]. The inset in Figure 4 shows the corresponding XP spectra of O 1s lines. Compared with the peak position before oxidation (curve a), the O 1s line shifts 0.4 eV towards lower BE, which may be associated with the difference of peak positions of O 1s lines in MgO and iron oxides [27]. To be noticed, before and after oxidation, the shift directions of O 1s and Fe 2p lines are opposite, being different from these during growing Fe films. Thus, combined with the shift of O, Mg and Fe core levels, the result further
Figure 3. XP spectra of Mg KLL, 2p and Fe 3p core levels with the increase of Fe thickness on MgO(1 1 1) films. The curves (a)–(h) correspond to the thickness of 0.08, 0.2, 0.4, 0.7, 1.0, 1.8, 3.0, 5.0 mL Fe.
Figure 4. XP spectra of Fe 2p core levels for: (a) 1.5 mL Fe grown on MgO(1 1 1) surface; (b) exposing it to 10 L oxygen; and (c) continuing to oxidize it at 5 10 7 mbar oxygen at 700 K for 20 min. The inset gives the XP spectra of O 1s lines corresponding to curves (a)–(c).
indicates that the chemical reaction at the interface of Fe and MgO is absent. According to the previous documents, the detailed microscopic mechanism for the upward band bending is complex, but it can be mainly attributed to four factors: interfacial bonding states, adatom-induced states, metal-induced gap states (MIGS) or defect states [11,28]. The first one is excluded in the present case because of the absence of interfacial bonding. At the initial growth stage, although the Fe atoms form clusters/particles, the hybridization between O atoms and absorbed Fe atoms causes the formation of discrete interface states (the adatom-induced states) at the valence band region. These states prompts the Fermi level away from the valence band, making the band bend upward. With the increase of Fe thickness, the metallic states become stronger and stronger and the MIGS are decayed exponentially within the band gap of MgO, bringing a further bending of the band. When Fe films are characteristic of metallic states, the MIGS reach saturation, resulting in the stabilization of Fermi level, i.e., all the core levels stop shifting. To be emphasized, it is difficult to avoid the effect of defect states on the interfacial electronic structures since there are always various defects such as vacancies and interstitial atoms in such-prepared oxide films [29]. The defect states may cause the quantitative change of core level shifts, but do not directly affect the direction of core level shifts [9]. Conclusively, the adatominduced states, MIGS and defect states are responsible for the interfacial electronic structure at the initial stage of Fe deposition. As a result, these factors causes a upward band bending of MgO, resulting in a total shift of 1.2 eV of the Fermi level, as shown in Figure 5. Similarly, the BE shifts and band bending of Ag and Au on oxides have been reported [5,9,30]. In HREELS measurements, the energy losses of surface optical phonons are easily monitored owing to the long-range electronic fields associated with dipole-active excitation in metal oxides with a high crystal quality [31]. In Figure 6a, the HREEL spectrum for a clean MgO film was shown. The main loss peak at 81 meV as well as its multiple loss peaks at 161 and 240 meV are attributed to the excitation of a long wavelength surface optical phonon (as known as Fuchs–Kliewer mode) [32]. With the increase of Fe thickness, the main loss peak at 81 meV gradually becomes weak in intensity, but no obvious shift of the main loss peak or additional peaks (such as the lattice vibration of Fe–O) are observed, being indicative of a weak interaction at the Fe/MgO interface. To be worth noting, the
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ing a formation of iron oxide after oxidizing Fe films [34]. In contrast, before exposing the Fe film to oxygen ambience the absence of the shoulder peak in Figure 6b–d implies that there is no interfacial layer of iron oxides formed at the Fe/MgO interface or the interfacial bonding is rather weak. These results are in good agreement with these obtained from XPS data. In summary, the initial growth, interfacial bonding and electronic structure of iron on faceted MgO(1 1 1) surface have been investigated by various surface analysis techniques. The results indicated a three-dimensional growth of Fe on faceted MgO(1 1 1) surface. The core-level shift of Fe at the initial stage of Fe deposition was mainly attributed to the particle size effect, but not originated from the chemical bonding at the interface. The results will be constructive for the understanding of Fe/MgO interface as well as the development of related magnetic devices. Acknowledgment Figure 5. Schematic energy diagram at the band gap region of MgO as a function of Fe thickness.
We gratefully acknowledge the financial support of this Letter by the Natural Science Foundation of China (Grant No. 21103084), the Natural Science Foundation of Hangkong (Grant No. 2010ZE56012), the Knowledge Innovation Program of the Chinese Academy of Sciences (Grant No. KJCX2-EW-W09), ‘973’ Program of China (2012CB921700), the Natural Science Foundation of Department of Education of Jiangxi Province (Grant No. GJJ11173), and the Natural Science Foundation of Department of Science and Technology of Jiangxi Province (Grant No. 20122BAB202013). References
Figure 6. HREEL spectra of: (a) clean MgO(1 1 1) films; (b–d) 0.5, 1.5 and 3 mL Fe on MgO(1 1 1) surface; (e) corresponding films after oxidizing 3 mL Fe at 5 10 7 mbar oxygen at 700 K for 20 min; and (f) a clean Fe3O4(1 1 1) films grown on Mo(1 1 0) substrate for comparison.
main loss peak remains existing after depositing 3 ML Fe (Figure 6d). Considering the high sensitivity of HREELS to surface species with 1-2 ML coverage, the result strongly suggests that the surface of MgO is not completely covered by 3 ML Fe, i.e., the growth of Fe follows a three-dimensional growth mode. However, when the 3 ML Fe film is oxidized at 5 10 7 mbar oxygen at 700 K for 20 min (Figure 6e), the intensity of the main loss peak at 81 meV is obviously increased. The result indicates that part of surface area of MgO occupied by Fe is exposed again, i.e., iron and its oxides are agglomerated at 700 K during oxidation. Temperature-induced agglomeration has been widely observed [33]. To be noticed, compared curve e with the HREEL spectrum (curve f) of a pure Fe3O4(1 1 1) film grown on Mo(1 1 0) substrate, a shoulder peak at 55 meV at the left of the main loss peak is observed. It is associated with the Fe–O stretching mode, indicat-
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Contents lists available at SciVerse ScienceDirect
Chemical Physics Letters journal homepage: www.elsevier.com/locate/cplett
Thickness-dependent electronic structure and interfacial behaviors of iron on faceted MgO(1 1 1) films Mingshan Xue a,b,⇑, Qinlin Guo b,⇑ a b
School of Materials Science and Engineering, Nanchang Hangkong University, Nanchang 330063, China Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, P.O. Box 603, Beijing 100190, China
a r t i c l e
i n f o
Article history: Received 27 June 2012 In final form 13 September 2012 Available online 21 September 2012
a b s t r a c t Based on the vital effect of the interfacial behaviors on tunneling magnetoresistance in Fe/MgO/Fe junctions, the thickness-dependent electronic structure of iron on MgO(1 1 1) films with {1 0 0} facets was investigated. The results illustrated that the chemical interaction at the interface of Fe and MgO films was rather weak. Instead, a upward band bending of MgO with 1.2 eV was obviously observed with the increase of Fe thickness. The particle size effect was responsible for these shifts of core levels owing to the three-dimensional growth of Fe. These results were further testified by the data from the lattice vibration in Fe–MgO system. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction The interactions at metal-oxide interfaces have a considerable effect on the related optical, chemical, electronic and magnetic properties [1,2]. In recent years, much effort has been made to understand these interactions for the application of such materials in sensor, semiconductor, microelectronics, catalysis, magnetic random memory as well as spintronics [3,4]. For example, related studies have indicated that the chemical bonding at the Fe/MgO interface is of much disadvantage to the tunneling magnetoresistance (TMR) in the Fe/MgO/Fe magnetic tunnel junctions [1,5,6]. The oxide layers at the interface obviously impact the tunneling current between two ferromagnetic layers, resulting in the sharp decrease of the TMR value. In the Fe/MgO catalysts, the distribution of iron oxides at the interface also strongly affects the H2S removal ability for the catalytic wet oxidation of H2S [7]. These factors stimulate an increasing interest towards the investigation of Fe–MgO interface. However, the studies between metals and bulk oxide crystals are often confined because of the surface charging problem and the difficulty of sample heating and cooling originating from the insulator character of bulk oxides [8]. Instead, ordered oxide thin films grown epitaxially on single-crystal metal substrates (such as Mo, W, Pt substrates) have been widely used [9,10]. As a low cost material with excellent mechanical, thermal and electronic properties, ordered MgO films grown on metal substrates are widely used as a template to investigate the interactions between metals and oxide surfaces, such as Ni/MgO, Ag/MgO and Zn/MgO [5,9,11,12]. One hand is in that the material with a high ⇑ Corresponding authors. Fax: +86 791 86453210. E-mail addresses: [email protected] (M. Xue), [email protected] (Q. Guo). 0009-2614/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.cplett.2012.09.034
stoichiometry is easily prepared due to the high oxygen affinity of magnesium. Another hand is in that it has a simple rock-salt structure, easily forming ordered films on single-crystal metal substrates. When Fe and MgO layers are fabricated into Fe/MgO/Fe magnetic tunnel junctions, these epitaxial devices exhibit large TMR ratios (e.g., 318% at 10 K) though they are still considerably lower than the values predicted theoretically [13]. The discrepancy between experimental and theoretical results is mainly attributed to the interfacial structures, such as the relative position of the atoms, interfacial oxidation, strain, and structural asymmetry of the interfaces [2–8,13]. In order to further clarify the detailed factors affecting the TMR, the evolution of surface/interface electronic structure at the initial growth stage of Fe on MgO surface should be taken into consideration. In previous works [2–9], the neutral (1 0 0) face of MgO in Fe–MgO system was widely used to investigate the interfacial behaviors of Fe and MgO. In spite of their efforts, whether or not a chemical reaction exists at the Fe–MgO interface is still controversial [2,6,14]. In this Letter, we try to construct a new system of Fe–MgO(1 1 1) with {100} facets, which is chemically more active than general MgO(1 0 0) surface. As a typical polar surface, MgO(1 1 1) films consisting of alternative stacking of Mg layers and O layers have a strong chemical activity due to a low coordination of surface atoms. Because of the so-called ‘polar instability’, the (1 1 1) surface of MgO is easily diverged into {100} facets (including (1 0 0), (0 1 0) and (0 0 1) faces) [15,16]. As a result, asformed surface is characteristic of (1 1 1) face covered by {100} facets, but it has a slightly higher activity than pure (1 0 0) surface of MgO films owing to the exposure of part atoms from (1 1 1) face in the valley region between {100} facets. Such a MgO surface may provide a stronger evidence for the absence of chemical bonding at the Fe–MgO (1 0 0) interface if there is no reactive layer observed
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in the present Fe–MgO (1 1 1) system. In our experiments, such a MgO (1 1 1) surface with {100} facets is easily prepared on Mo(1 1 0) substrate under ultrahigh vacuum (UHV) condition [12]. X-ray photoelectron spectroscopy (XPS) is used to obtain the first indication of the electronic structure and chemical states of Fe and MgO films. High resolution electron energy loss spectroscopy (HREELS) and low-energy electron diffraction (LEED) are used to monitor their lattice vibration and surface structure. The experiments were performed in a UHV chamber which had a base pressure of less than 5 10 10 mbar. The chamber was equipped with LEED, XPS and HREELS. The XPS analyses were carried out using a Mg Ka X-ray source (hv = 1253.6 eV) with a pass energy of 50 eV. The binding energy (BE) was calibrated with respect to the pure bulk Au 4f7/2 (BE = 84.0 eV) and Ag 3d5/2 (BE = 368.3 eV) lines, and the accuracy on the core-level displacement was better than 0.15 eV. In HREELS measurements, the primary electron incident energies were set to 4.9 eV and the typical resolution was 12–13 meV obtained by the full width at the half maximum (FWHM) height of the elastic peak from the MgO(1 1 1) films. The magnesia films were epitaxially grown on a single-crystal Mo(1 1 0) substrate. The substrate surface was treated by annealing at about 1200 K in 10 7 mbar O2 environment, followed by a subsequent flash to 1500 K until no impurity was detected by XPS. The magnesium and iron sources were made of a pure magnesium ribbon (purity >99.9%) and iron wire (purity >99.99%) wrapped tightly around a tungsten wire, respectively. The two sources were thoroughly degassed by thermal treatment before preparation. The deposition rates of Mg and Fe were about 0.15 and 0.08 monolayer (ML)/min by means of XPS, calibrated via the relative intensity of Mg 2p3/2 and Fe 2p3/2 to Mo 3d5/2 lines as a function of deposition time, respectively. For example, in XPS measurements, to monitor the growth rate of Fe, the growth of Fe was checked by measuring the Fe 2p3/2 to Mo 3d5/2 intensity ratio as a function of deposition time, as shown in Figure 1. A straight line was fitted to the data for the initial growth of Fe on Mo(1 1 0) substrate. There was a clear breakpoint for the curve after growing Fe for 12 min, corresponding to 1 ML Fe film. It suggested that the growth of Fe on Mo(1 1 0) surface followed a Stranski–Krastanow (SK) mode: Fe initially grew as two-dimensional patches at submonolayer coverages, and then gradually formed three-dimensional islands on top of the patches. During growth, the deposition rate of Fe was kept constant by means of adjusting the output power of the constant-current direct supply. The MgO thin films with 15 ML were
Figure 1. Statistical XPS intensity ratio of Fe 2p3/2 to Mo 3d5/2 core levels as a function of thickness of Fe on Mo(1 1 0) substrate. The insets give the LEED pattern of MgO(1 1 1) films with {1 0 0} facets, Ep = 72 eV, and the structural schematic of the interface of Fe and MgO, respectively.
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prepared by evaporating Mg in an ambience of 1 10 6 mbar O2 at 673 K, followed by annealing at 800 K in 1 10 7 mbar O2. Then as-prepared MgO(1 1 1) films were used as the support to deposit Fe at room temperature (RT). All the experimental data were in situ measured at RT. The MgO(1 1 1) films with {100} facets were grown on Mo(1 1 0) substrate, which had been described in a previous study [12]. Besides the core-level lines originating from Mg and O, there are no other peaks observed by XPS such as these from C or Mo, indicating a clean MgO surface. The corresponding BE values of Mg 2p and O 1s lines are 50.8 and 531.2 eV, respectively, being in agreement with previous data of MgO films. The inset in Figure 1 gives a typical LEED pattern from the faceted MgO(1 1 1) films. The LEED pattern consists of two sets: the inner one is of hexagonal symmetry and the outer one is split and diffused in non-radical directions with the slight change of the primary energy [17,18]. It is similar to a cleaved bulk MgO (1 1 1) crystal surface, which is easily covered by a series of trigonal pyramids featured by three sets of {1 0 0} facets (consist of (1 0 0), (0 1 0) and (0 0 1) planes) inclined by 54.7° to the original surface normal caused by the polar instability of MgO (1 1 1) surface [17–19]. Furthermore, the formation of (1 0 0) facets on (1 1 1) face is in qualitative agreement with theoretic analysis of the stability of ionic crystal surfaces based on Madelung potential and lattice dynamics calculation [19,20]. Figure 2 gives the XP spectra of Fe 2p core levels as a function of Fe thickness. For Fe, XPS measurements can identify its different chemical states by observing the BE of Fe 2p core levels. Many studies have indicated that the BEs of Fe 2p3/2 core level for Fe0, Fe2+ and Fe3+ states are 707, 710.7 and 711 eV, respectively [21,22]. For initial 0.08 ML Fe, the corresponding BE value of Fe 2p3/2 line is 708.2 eV, as shown in Figure 2. The peak position is far from the BE values of Fe2+ or Fe3+ states (more than 710 eV), implying an absence of interfacial bonding between Fe and MgO(1 1 1) surface. With the increase of Fe coverage, the peak shifts gradually towards a lower BE until the value is 707 eV. The peak at 707 eV means a metallic Fe0 state because the Fe film is formed after growing several ML Fe. However, no obvious LEED pattern was detected from the corresponding Fe film, indicating a disordered Fe film. It may originate from the three-dimensional growth of Fe on MgO(1 1 1) surface, i.e., the small Fe clusters/particles are formed at the initial
Figure 2. XP spectra of Fe 2p core levels with the increase of Fe thickness on MgO(1 1 1) films. The vertical lines give the peak positions for initial and final thickness of Fe.
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stage of Fe deposition (as shown in Figure 1), being different from the case of Fe/Mo system. The particle size effect causes the shift of Fe core levels towards higher BE, resulting in a total shift of 1.2 eV from initial to final growth. Figure 3 shows the XP spectra of Mg KLL, 2p and Fe 3p core levels with the increase of Fe thickness. For Mg 2p and Fe 3p core levels, their peaks are very close so that it is hard to distinguish the shifts of their peaks. However, for Mg KLL spectra, a shift of 1.2 eV towards lower BE is obviously observed with the increase of Fe coverage from 0 to 5 ML. Similarly, the phenomena are observed in the XP spectra of O 1s lines (not shown). The shifts of core-level lines are attributed to the band bending at the MgO surface [23,24]. Due to the work-function difference of Fe and MgO following the Schottky–Mott theory, the bands bend upward towards the surface with a subsequent Fe deposition, resulting in the continuous decrease of BEs of Mg and O core levels [25]. Moreover, the same shift direction and magnitude of the Mg and O core levels support the interpretation of as-observed shifts of Fe 2p core levels as the pLetter size effect rather than chemical interaction at the interface. To completely exclude the existence of iron oxides at the Fe/ MgO interface, we further investigated the oxidation of 1.5 ML Fe films by XPS, as shown in Figure 4. For 1.5 ML Fe film (curve a), the BE of Fe 2p3/2 line is at 707 eV. After only exposing to 10 L (Langmuir, 1 L = 1.33 10 6 mbar s) oxygen at RT, the BE value is shifted to 710.7 eV, which is typically characteristic of iron oxides (i.e., Fe3O4). When the film is further oxidized at 10 6 mbar oxygen at 700 K for 20 min, the peak has almost no shift in BE, but has a better symmetry in the shape of the peak. The result reveals that Fe atoms are easily oxidized in an oxygen ambience, and most of 1.5 ML iron atoms deposited have been oxidized by 10 L oxygen [26]. The inset in Figure 4 shows the corresponding XP spectra of O 1s lines. Compared with the peak position before oxidation (curve a), the O 1s line shifts 0.4 eV towards lower BE, which may be associated with the difference of peak positions of O 1s lines in MgO and iron oxides [27]. To be noticed, before and after oxidation, the shift directions of O 1s and Fe 2p lines are opposite, being different from these during growing Fe films. Thus, combined with the shift of O, Mg and Fe core levels, the result further
Figure 3. XP spectra of Mg KLL, 2p and Fe 3p core levels with the increase of Fe thickness on MgO(1 1 1) films. The curves (a)–(h) correspond to the thickness of 0.08, 0.2, 0.4, 0.7, 1.0, 1.8, 3.0, 5.0 mL Fe.
Figure 4. XP spectra of Fe 2p core levels for: (a) 1.5 mL Fe grown on MgO(1 1 1) surface; (b) exposing it to 10 L oxygen; and (c) continuing to oxidize it at 5 10 7 mbar oxygen at 700 K for 20 min. The inset gives the XP spectra of O 1s lines corresponding to curves (a)–(c).
indicates that the chemical reaction at the interface of Fe and MgO is absent. According to the previous documents, the detailed microscopic mechanism for the upward band bending is complex, but it can be mainly attributed to four factors: interfacial bonding states, adatom-induced states, metal-induced gap states (MIGS) or defect states [11,28]. The first one is excluded in the present case because of the absence of interfacial bonding. At the initial growth stage, although the Fe atoms form clusters/particles, the hybridization between O atoms and absorbed Fe atoms causes the formation of discrete interface states (the adatom-induced states) at the valence band region. These states prompts the Fermi level away from the valence band, making the band bend upward. With the increase of Fe thickness, the metallic states become stronger and stronger and the MIGS are decayed exponentially within the band gap of MgO, bringing a further bending of the band. When Fe films are characteristic of metallic states, the MIGS reach saturation, resulting in the stabilization of Fermi level, i.e., all the core levels stop shifting. To be emphasized, it is difficult to avoid the effect of defect states on the interfacial electronic structures since there are always various defects such as vacancies and interstitial atoms in such-prepared oxide films [29]. The defect states may cause the quantitative change of core level shifts, but do not directly affect the direction of core level shifts [9]. Conclusively, the adatominduced states, MIGS and defect states are responsible for the interfacial electronic structure at the initial stage of Fe deposition. As a result, these factors causes a upward band bending of MgO, resulting in a total shift of 1.2 eV of the Fermi level, as shown in Figure 5. Similarly, the BE shifts and band bending of Ag and Au on oxides have been reported [5,9,30]. In HREELS measurements, the energy losses of surface optical phonons are easily monitored owing to the long-range electronic fields associated with dipole-active excitation in metal oxides with a high crystal quality [31]. In Figure 6a, the HREEL spectrum for a clean MgO film was shown. The main loss peak at 81 meV as well as its multiple loss peaks at 161 and 240 meV are attributed to the excitation of a long wavelength surface optical phonon (as known as Fuchs–Kliewer mode) [32]. With the increase of Fe thickness, the main loss peak at 81 meV gradually becomes weak in intensity, but no obvious shift of the main loss peak or additional peaks (such as the lattice vibration of Fe–O) are observed, being indicative of a weak interaction at the Fe/MgO interface. To be worth noting, the
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ing a formation of iron oxide after oxidizing Fe films [34]. In contrast, before exposing the Fe film to oxygen ambience the absence of the shoulder peak in Figure 6b–d implies that there is no interfacial layer of iron oxides formed at the Fe/MgO interface or the interfacial bonding is rather weak. These results are in good agreement with these obtained from XPS data. In summary, the initial growth, interfacial bonding and electronic structure of iron on faceted MgO(1 1 1) surface have been investigated by various surface analysis techniques. The results indicated a three-dimensional growth of Fe on faceted MgO(1 1 1) surface. The core-level shift of Fe at the initial stage of Fe deposition was mainly attributed to the particle size effect, but not originated from the chemical bonding at the interface. The results will be constructive for the understanding of Fe/MgO interface as well as the development of related magnetic devices. Acknowledgment Figure 5. Schematic energy diagram at the band gap region of MgO as a function of Fe thickness.
We gratefully acknowledge the financial support of this Letter by the Natural Science Foundation of China (Grant No. 21103084), the Natural Science Foundation of Hangkong (Grant No. 2010ZE56012), the Knowledge Innovation Program of the Chinese Academy of Sciences (Grant No. KJCX2-EW-W09), ‘973’ Program of China (2012CB921700), the Natural Science Foundation of Department of Education of Jiangxi Province (Grant No. GJJ11173), and the Natural Science Foundation of Department of Science and Technology of Jiangxi Province (Grant No. 20122BAB202013). References
Figure 6. HREEL spectra of: (a) clean MgO(1 1 1) films; (b–d) 0.5, 1.5 and 3 mL Fe on MgO(1 1 1) surface; (e) corresponding films after oxidizing 3 mL Fe at 5 10 7 mbar oxygen at 700 K for 20 min; and (f) a clean Fe3O4(1 1 1) films grown on Mo(1 1 0) substrate for comparison.
main loss peak remains existing after depositing 3 ML Fe (Figure 6d). Considering the high sensitivity of HREELS to surface species with 1-2 ML coverage, the result strongly suggests that the surface of MgO is not completely covered by 3 ML Fe, i.e., the growth of Fe follows a three-dimensional growth mode. However, when the 3 ML Fe film is oxidized at 5 10 7 mbar oxygen at 700 K for 20 min (Figure 6e), the intensity of the main loss peak at 81 meV is obviously increased. The result indicates that part of surface area of MgO occupied by Fe is exposed again, i.e., iron and its oxides are agglomerated at 700 K during oxidation. Temperature-induced agglomeration has been widely observed [33]. To be noticed, compared curve e with the HREEL spectrum (curve f) of a pure Fe3O4(1 1 1) film grown on Mo(1 1 0) substrate, a shoulder peak at 55 meV at the left of the main loss peak is observed. It is associated with the Fe–O stretching mode, indicat-
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