Study on growth mechanisms and electronic structures of polar NiO(1 1 1) ultrathin films using iron oxide buffer layers

Journal of Alloys and Compounds 598 (2014) 224–229

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom

Study on growth mechanisms and electronic structures of polar NiO(1 1 1) ultrathin films using iron oxide buffer layers 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, PO Box 603, Beijing 100190, China

a r t i c l e

i n f o

Article history: Received 22 September 2013 Received in revised form 6 February 2014 Accepted 6 February 2014 Available online 15 February 2014 Keywords: NiO(1 1 1) Buffer layers Polar materials

a b s t r a c t Polar surfaces of metal oxides have attracted considerable attention in fundamental science and technological applications because of their peculiar stabilization mechanisms and unusual adsorption and catalytic properties. In this study, various ordered iron oxide films, including FeO(1 1 1), Fe3O4(1 1 1) and Fe2O3(0 0 0 1), were successfully used as the buffer layers for the growth of polar NiO(1 1 1) films with several nanometer thickness. The results indicated that the iron oxide buffer layers reduced the lattice mismatch between NiO(1 1 1) and Mo(1 1 0) substrate, and also decreased the interfacial energy by means of the interfacial reaction between NiO and iron oxide surfaces, benefiting the initial nucleation and the growth of polar NiO(1 1 1) films. This study will be essential for understanding the physical and chemical properties of polar surfaces. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction In surface science, the stability of polar surfaces has long been one of the puzzles. However, the underlying applied demands and development prospects, e.g., from heterogeneous catalysis and digital storage to spintronics, arouse the studies on the growth and stability mechanisms of polar materials [1–3]. According to classical electrostatic criteria, many bulk terminated oxide surfaces, such as NiO(1 1 1), MgO(1 1 1), and ZnO(0 0 0 1), consist of alternate stacking of cation layers and oxygen layers and have a net dipole moment (infinite surface energy) normal to the surface, resulting in the so-called ‘‘polar instability’’. In the past decade, inspite of having some theoretical studies on the polar surfaces/ interfaces using density function theory [1,4,5], the related experimental information about polar surfaces is lacking since it is difficult to prepare such an ideal polar surface owing to the polar instability. As one of typical rock-salt oxides, nickel oxide (NiO) as a promising transition-metal oxide has attracted considerable attention because of its extensively important applications as catalysts, lithiumion batteries, supercapacitors, gas sensor and especially magnetic devices [6–8]. For example, the NiO(1 1 1) plane may ⇑ Corresponding authors at: Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, PO Box 603, Beijing 100190, China (M. Xue). Tel./fax: +86 791 86453210. E-mail addresses: [email protected] (M. Xue), [email protected] (Q. Guo). http://dx.doi.org/10.1016/j.jallcom.2014.02.023 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.

perform exchange coupling of ferromagnetic films in giant magnetoresistive sensors due to its anti-ferromagnetic structure with a Neel temperature of 250 °C [6]. For a better understanding of their properties, it is of significance and necessity to obtain a high-quality single crystal or well-ordered film of NiO(1 1 1). However, the (1 1 1) plane of NiO consists of alternate stacking of Ni layers and O layers along [1 1 1] direction, being a typical polar surface [9]. To stabilize the polar surface, the surface readily undergoes faceting, reconstruction, hydroxylation, or adsorption of contaminations [10–13]. For instance, it was reported that Ni vacancies induced p(2  2) NiO(1 1 1) [11]. These processes stabilizing the polar surfaces affect their physical and chemical properties (such as new surface states) and are disadvantageous to reveal the inherent properties of the polar materials [10]. Undoubtedly, for these hard-preparing metal oxides, to optimize the growth methods is vital to realize the facile synthesis of well-ordered nanofilms and to develop their electronic, chemical and catalytic performance. Meanwhile, to the best of our knowledge, the magnetic coupling between ferromagnetic and antiferromagnetic materials (such as NiO/Fe3O4 system) as well as the related exchange-coupled mechanism at the interface are far from clear [6,14]. Therefore, a detailed knowledge of the atomic structure of ultrathin NiO(1 1 1) films is necessary for understanding their electronic, chemical, catalytic and magnetic properties. To prepare high-quality NiO(1 1 1) films on different substrates, e.g., Si, Au, Al2O3 and Fe3O4, various growth methods including molecular beam epitaxy, pulse laser deposition, metal organic

M. Xue, Q. Guo / Journal of Alloys and Compounds 598 (2014) 224–229

chemical vapor deposition and magnetron sputtering, have been employed [10,11,13,14]. However, in many cases, the quality of metal oxide films is not obviously enhanced due to the lattice mismatch, the existence of oxygen vacancies as well as the difficulty in creating effective nucleation centers [15]. For example, in previous studies [10,16], the {1 0 0} facets readily appeared for the growth of MgO{1 1 1} or NiO(1 1 1) films on Mo(1 1 0) substrate. In this work, we present the epitaxial growth of polar NiO(1 1 1) films on Mo(1 1 0) substrate using the buffer layers of ordered iron oxide (including FeO(1 1 1), Fe3O4(1 1 1) and Fe2O3(0 0 0 1)). These films are in situ characterized by various surface analytical techniques including X-ray photoelectron spectroscopy (XPS), ultraviolet photoelectron spectroscopy (UPS), Auger electron spectroscopy (AES), high-resolution electron energy loss spectroscopy (HREELS), and low-energy electron diffraction (LEED). The experimental results suggest that the interfacial chemical reaction between NiO and iron oxide surfaces plays a positive role in the initial nucleation and growth of NiO(1 1 1) films. 2. Experimental section The preparation and characterization of the samples were performed in two ultrahigh vacuum (UHV) systems (ESCALAB-5 (VG SCIENTIFIC) and ELS-22 (LEYBOLD-HERAUS GMBH)) described in detailed elsewhere [2,16]. The ESCALAB-5 system with a base pressure of 8  1010 mbar equipping with optical facilities for LEED, dual-anode X-ray sources (Mg and Al) and a He I (hv = 21.2 eV) source for XPS and UPS, respectively. The ELS-22 chamber with a base pressure of 1  1010 mbar is for AES, HREELS and LEED studies. In two chambers, various metal and gas sources were installed. A Mo(1 1 0) single crystal was used as the substrate for the growth of various iron oxide films. It could be either resistively heated to 1300 K or annealed to temperature >2000 K with an electron beam heating assembly at the backside of the sample. A C-type thermocouple (5%Re–W/26%Re–W) was spot welded to the edge of the Mo(1 1 0) substrate for temperature measurements. The substrate surface was treated at 1200 K in 107 mbar oxygen to remove the surface contamination (mainly carbon), followed by a subsequent flash to 1800 K without oxygen until obtaining a clean surface monitored by XPS/AES and LEED. Fe and Ni sources were made of a pure iron wire (purity >99.95%) and a Ni wire (purity >99.994%) wrapped tightly around a tungsten wire, respectively. Before growth, the two sources were thoroughly degassed by thermal treatment. The deposition rates of Fe and Ni were about 0.15 and 0.1 monolayer (ML)/min, respectively, calibrated via the intensity ratios of Fe 2p3/2 to Mo 3d5/2 and Zn 2p3/2 to Mo 3d5/2 lines as a function of deposition time by XPS in ESCALAB-5 chamber, or measured by a quartz crystal oscillator in ELS-22 chamber. The preparation of iron oxides was tried by evaporating Fe at 5  1085  106 mbar O2 pressure at the temperature range of 300–1100 K, or by oxidizing pre-depositing Fe(1 1 0) films at 1  1071  106 mbar O2 pressure at the temperature range of 300–900 K. We emphasized that the selective growth of iron oxide films tightly related to the oxygen partial pressure and substrate temperature. All films had an average thickness of 6–8 ML. Then, these iron oxide films were used as a template to prepare NiO films by depositing Ni in 106 mbar O2 at room temperature (RT). The Mg Ka X-ray source (hv = 1253.6 eV) with a pass energy of 50 eV was used in XPS measurements. 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 of measured BE is better than 0.15 eV. The ultraviolet photon of hv = 21.2 eV was excited by high pure He (99.99%) in UPS measurements. A primary electron beam energy of 3 keV was used in AES measurements. In the HREELS, the data were recorded at the specular direction (60° with respect to the surface normal), and the primary energy of the electron beam is set to 4.9 eV and the full width at the half maximum (FWHM) height of the elastic peak from the substrate gives a typical revolution of 9–10 meV.

3. Results and discussion 3.1. Growth and electronic structures of ordered iron oxide films In order to identify the as-prepared buffer layers, it is of much necessity to know about the characters of various iron oxides in structure and chemical states since Fe and O can form a number of phases with different stoichiometries and crystal symmetries. Table 1 gives some parameters of iron and its oxides. The single crystal of iron is body-centered cubic. Rock-salt FeO has the same structure with NaCl crystal and along h1 1 1i direction, the oxygen

225

sublattice consists of a close packed face-centered cubic (fcc) structure (O–O distance: 3.04 Å) with interstitial sites occupied by Fe2+ species. Magnetite Fe3O4 is characteristic of the inverse spinel structure and along h1 1 1i direction the oxygen ions form a close-packed fcc sublattice (O–O distance: 2.97 Å) with interstitial arrangement of Fe2+ and Fe3+ species. There are two layers: one is tetrahedrally occupied by Fe3+ cations, while the other is octohedrally occupied by equal Fe2+ and Fe3+ species. Hematite a-Fe2O3 has the corundum structure with hexagonal symmetry and along h0 0 0 1i direction the oxygen ions arrange into a hexagonal packed structure (O–O distance: 2.90 Å) with two layer Fe3+ cations in the interstitial sites. Due to electrostatic interaction and spin orbit coupling between Fe 2p core holes and unpaired 3d electrons [17,18], the BEs and peak shape of Fe core-level lines in different iron oxides are obviously distinguishable by XPS, as shown in Fig. 1. For example, to distinguish the Fe2+ or Fe3+ states, XPS can detect the presence of satellite peaks of Fe 2p core levels because of the shake-up process originating from the transfer of an electron from a 3d orbital to an empty 4s orbital during excitation of Fe 2p core-level photoelectron, which has been reported in our previous work [19]. By depositing Fe at the oxygen pressure of 1  107 mbar at 650 K, followed by annealing at 850 K without O2, the corresponding BE and satellite peak of Fe 2p3/2 line are 710 and 716 eV in XPS spectra (Fig. 1), respectively, being in well agreement with Fe2+ state [20]. Rock-salt FeO(1 1 1) surface has a O–O distance of 3.04 Å. The Mo(1 1 0) substrate has lattice constants with a0 = 3.15 Å and b0 = 2.73 Å, a distorted hexagonal structure. The lattice mismatch in two dimensional surface is 4.3% and 10.9% along a0 and b0 direction, corresponding to tensile and compressive stress, respectively. Due to the impact of interfacial strain, the tensile strain along a0 direction fitly reduces the compressive strain along b0 direction to release the strain energy, and the growth of FeO films along [1 1 1] direction recovers to a symmetric hexagonal structure from the distorted hexagonal structure of Mo(1 1 0) [15,21]. The growth mechanism is most likely similar to the case of MgO(1 1 1)/Mo(1 1 0) system [2]. Previous experimental results indicated that Fe2O3 films were hard fabricated under the low oxygen pressure in UHV condition [22,23]. According to theoretical Fe–O phase diagram, it is possible to prepare Fe2O3 films under UHV condition [24]. In our experiments, we found that it was difficult to obtain Fe2O3 films by co-depositing Fe at O2 ambient with various pressure (<5  105 mbar) at substrate temperature from 300 to 1000 K. We tried to investigate the situation of oxidizing Fe films at different temperature. After depositing 3–5 nm Fe(1 1 0) films on Mo(1 1 0) substrate, they were oxidized at 106 mbar O2 at a certain substrate temperature. After oxidization at 600 K, we found that the Fe 2p3/2 line and its weak satellite peak were at 711.1 and 719 eV (Fig. 1), respectively. Compared with the reference data, the results suggested the appearance of pure Fe3+ states [25,26]. Although Fe2O3 has four polymorphs with a-, b-, c- and e-Fe2O3, only the a-Fe2O3 crystal is most likely stable under UHV condition. Based on the bulk structure of a-Fe2O3, the (0 0 0 1) face is of hexagonal structure with a unit cell of 5.036 Å and has a O–O distance of 2.90 Å, being similar to Mo(1 1 0) surface in lattice match. The result suggests that the iron oxide films after annealing at 600 K corresponds to Fe2O3 films [19]. Among iron oxides, the Fe3O4 is the most stable one in structure. From the chemical point of view, every cell of Fe3O4 can be considered as [FeO]1[Fe2O3]1. By co-deposition of Fe at 106 mbar O2, Fe3O4 films are easily obtained. The corresponding XP spectrum of Fe 2p core levels is shown in Fig. 1. The BE of Fe 2p3/2 line is 710.3 eV, being consistent with the previous results [19,24,25]. In addition, the FWHM of Fe 2p3/2 peak is obviously wider than these from FeO or Fe2O3 due to the coexistence of Fe2+ and Fe3+

226

M. Xue, Q. Guo / Journal of Alloys and Compounds 598 (2014) 224–229

Table 1 Structural and chemical parameters of iron and its oxides. Parameter

Type

Crystal structure Lattice constants of oxygen sublattice (Å) BE of Fe 2p3/2 (eV) BE of satellite peak of Fe 2p3/2 (eV)

Fe

FeO

Fe3O4

Fe2O3

Body cubic 707 + 0.2 -

Rock-salt 3.04 709.8 + 0.2 716.0

Inverse spinel 2.97 710.3 + 0.2 -

Corundum 2.90 710.9 + 0.2 719.0

Sat.

Sat.

Fe

3+ Fe state

Intensity (a. u.)

Intensity (a. u.)

O 2p

Fe2O3

FeO Fe3O4

Fe3O4

Sat.

2+ Fe state

Sat. Fe2O3

735 730 725 720 715 710 705 Binding Energy (eV) Fig. 1. XPS spectra of Fe 2p core levels in pure Fe(1 1 0), FeO(1 1 1), Fe3O4(1 1 1), and Fe2O3(0 0 0 1) films grown on Mo(1 1 0) substrate.

states in Fe3O4 [27]. There is no satellite peak observed. The characteristic satellite peaks of Fe2+ and Fe3+ 2p3/2 lines are at 716 eV and 719 eV, respectively. When the chemical states of Fe2+ and Fe3+ appear in a compound, the overlap of satellite peaks of Fe2+ and Fe3+ states makes the region from 716 to 719 eV become flat, resulting in no observable peaks. During the experimental process, it is also found that Fe3O4 films can be obtained by the transformation from FeO or Fe2O3 films. When FeO films are further exposed to 106 mbar O2 at 650 K, part of Fe2+ states are oxidized to Fe3+ states, causing the formation of Fe3O4 films [28]. The increasing O2 pressure benefits the oxidization from Fe2+ to Fe3+ states, while the suitable temperature makes Fe and O atoms obtain enough energy and results in the transfer and rearrangement of Fe and O atoms, tending to form Fe3O4 films with a well thermodynamic stability. The similar transformation of Fe3O4 films from FeO films formed on Pt(1 1 1) substrate by layer-by-layer growth techniques was also reported [24,28]. As for the transformation from Fe2O3 films to Fe3O4 films, it is mainly attributed to the reduction of part of Fe3+ ions at annealing temperature of more than 650 K. Fig. 2 shows the UPS spectra of various iron oxide films. Just as most of oxides, the O 2p states are main responsible for the intense emission from 4 eV to 7 eV [29], while the features below 3.5 eV originate from the contribution of Fe 3d emissions. For FeO films, Fe2+ states correspond to 3d6 electron configuration, and the final state spin configuration 3d5 out of Fe2+ 3d6 initial state is responsible for the emission peak within 1.0 eV [30,31]. For Fe2O3 films, Fe3+ states characterized 3d5 electron configuration, and the ultraviolet photon excites the emission of 3d4 from 3d5 electron configuration. As a result, there are no obvious peaks within 1.0 eV. Instead, A peak at 3.5 eV appears and associates with the evolution

10

8

6 4 2 Binding Energy (eV)

FeO

0

Fig. 2. UPS spectra of FeO(1 1 1), Fe3O4(1 1 1) and Fe2O3(0 0 0 1) films, hv = 21.2 eV. The characteristic peaks of Fe2+ and Fe3+ originating from Fe 3d5L and 3d6L states are marked.

from 3d5 to 3d4 states [32]. However, for Fe3O4 films, the 3d emission near EF consists of two final state spin configuration 3d5 out of Fe2+ states and one quintet final states spin configuration 3d4 out of Fe3+ states [31,32]. As a result, there are two peaks near EF:0.7 eV and 2.7 eV separated by the quartet-sextet exchange interaction energy [33]. In addition, there are two additional phenomena: the intensity of the characteristic peak of Fe2+ states in Fe3O4(1 1 1) films is as strong as 1/3 of that in FeO(1 1 1) films, being in accord with the fact that there are 1/3 cations of Fe2+ states every cell of Fe3O4; the peak at 2.7 eV is obviously observed in Fe3O4 films, while it is not obvious in Fe2O3 films. It is most likely due to the overlap of O 2p and Fe 3d states, resulting in the strong peak at 3.5 eV. In a word, the main characteristics distinguishing Fe2+ from Fe3+ states are the presence of a Fe2+ 3d6-derived peak within 1.0 eV and a Fe3+ 3d5-derived peak at 2.7 eV near EF. Fig. 3 shows the HREELS spectra for three kinds of iron oxide films. In HREELS the electron energy losses are representative of the main Fuchs–Kliewer photon features, which is dominated by long-range electric fields associated with dipole-active excitations. For FeO(1 1 1) films, the main loss peak at 62 meV as well as the multiple loss peak observed at 124 meV characters the lattice vibration of Fe2+–O. Combining with the corresponding hexagonal LEED pattern in the inset in Fig. 3, these data suggest the formation of ordered FeO(1 1 1) films. For Fe3O4 film, there are two main loss peaks at 50.5 and 72 meV. Similarly, together with the typical LEED pattern they suggests the formation of ordered Fe3O4(1 1 1) films. Interestingly, compared with Fe3O4(1 1 1) films, the peak at 50.5 meV in Fe2O3(0 0 0 1) films has no shift but the peak at 72 meV shifts to 77.5 meV. Obviously, the shift and split of these loss peaks are strongly associated with the change of iron oxides in chemical states and crystal structures.

227

M. Xue, Q. Guo / Journal of Alloys and Compounds 598 (2014) 224–229

Fig. 3. HREELS spectra of (a) FeO(1 1 1), (b) Fe3O4(1 1 1) and (c) Fe2O3(0 0 0 1) films, Ep = 4.9 eV. The insets show the corresponding LEED patterns of FeO(1 1 1) with Ep = 54 eV, Fe3O4(1 1 1) with Ep = 56 eV, O–Fe2O3(0 0 0 1) surface with Ep = 60 eV.

Fig. 5. HREELS spectrum of NiO(1 1 1) films grown on Fe3O4(1 1 1) films, Ep = 4.9 eV. The inset shows the corresponding LEED pattern of NiO(1 1 1) films with Ep = 65 eV.

As-prepared iron oxide thin films were used as the buffer layers for the growth of NiO films on Mo(1 1 0) substrate by evaporating Ni in 106 mbar oxygen at RT. Fig. 4 shows the XPS spectrum of Ni 2p core levels for nickle oxide grown on Fe3O4(1 1 1) buffer layers. The value of BE is 853.8 eV for Ni 2p3/2 core level, corresponding to Ni2+ state [25]. The position and shape of Ni 2p and O 1s peaks are consistent with that of NiO reported elsewhere [10], i.e. the formation of NiO for as-prepared thin films. Fig. 5 shows the HREELS result for as-prepared NiO film. The main loss peak at 67 meV as well as the multiple loss peaks at 134 and 199 meV characters the surface optical phonons (Fuchs– Kliewer mode) from the lattice vibration of NiO single crystal, which is in well agreement with that from single crystals of NiO [29]. The inset in Fig. 5 demonstrates the corresponding LEED pattern with the hexagonal spots, indicating the epitaxial growth of NiO(1 1 1) films on Fe3O4(1 1 1) buffer layers. Fig. 6 shows the UPS spectrum of corresponding NiO films. In the valence region, the large peak at 5.1 eV is from the non-bonding O 2p emission, and the peak at 2.4 eV represents the screened Ni 3d emission, which is well consistent with previous data [10,29]. Therefore, the UPS result together with XPS, HREELS and LEED results confirm the formation of ordered NiO(1 1 1) films on

Intensity (a. u.)

3.2. Growth of NiO(1 1 1) films

12

10

8

6

4

2

0

Binding Energy (eV) Fig. 6. UPS spectrum of NiO(1 1 1) films grown on Fe3O4(1 1 1) films.

Fe3O4(1 1 1) buffer layers. Similarly, the epitaxial growth of NiO(1 1 1) films on FeO(1 1 1) and Fe2O3(0 0 0 1) buffer layers can be achieved. 3.3. Interfacial electronic structures

O 1s

Intensity (a.u.)

NiO

Fe3O4

2p3/2 536 534 532 530 528 526 Binding Energy (eV)

2p1/2

Sat.

Sat.

880 875 870 865 860 855 850 Binding Energy (eV) Fig. 4. XPS spectrum of Ni 2p core levels in NiO(1 1 1) films grown on Fe3O4(1 1 1) films. The inset shows the corresponding XPS spectra of O 1s core levels in NiO(1 1 1) and Fe3O4(1 1 1) films.

NiO(1 1 1) and Fe3O4(1 1 1) are characteristic of rock-salt and inverse spinel structures, respectively. Structurally, it is possible for the epitaxial growth of ordered NiO films on Fe3O4(1 1 1) buffer layers because both NiO(1 1 1) (O–O distance: 2.95 Å) and Fe3O4(1 1 1) surfaces are hexagonal symmetry in two-dimensional (2D) cells and their lattice mismatch in oxygen sublattice is very small (0.6%). In our experiments, at the topmost surface the Fe3O4(1 1 1) buffer layers are most likely O-terminated owing to the existence of oxygen ambience as described in experimental section. In Fe3O4(1 1 1) films the O-terminated layer consisted of a close-packed oxygen matrix benefits to attract cations onto the surface. Thus, a wetting layer of NiO can be formed by means of the absorption of Ni or Ni–O species on the O-terminated layer. Accordingly,   Fe–O–Ni–O–Ni   layers along the surface normal are grown. This is to say, it is vital for the interfacial layer as a seed layer to result in an ordered NiO film. Obviously, if there is a chemical interaction occurring between O-terminated iron oxide layer and Ni at the interface, it will be helpful for the formation of a wetting layer of NiO.

228

M. Xue, Q. Guo / Journal of Alloys and Compounds 598 (2014) 224–229

Fig. 7. XPS spectra of (a) Ni 2p and (b) Fe 2p core levels for Fe3O4(1 1 1)-supported Ni with the coverage of 0, 0.4, 1.0, 2.0, and 3.0 ML (labeled a–e), respectively. (f) Corresponds to the surface after oxidizing 1 ML Ni grown on Fe3O4(1 1 1) films.

In order to detect the chemical behaviors at the interface of NiO and Fe3O4(1 1 1) buffer layers, Ni was deposited step-by-step on Fe3O4(1 1 1) films without oxygen. Fig. 7(a) shows the XPS spectra of Ni 2p core levels with the increase of Ni thickness. At the initial coverage of 0.4 ML Ni, the Ni 2p3/2 line is at about 853.6 eV, being very close to 853.8 eV characterizing Ni2+ state. With the increase of Ni thickness, the peak gradually shifts towards lower BE until it reaches 852.8 eV characterizing metallic Ni (curves c–e in Fig. 7(a)). Fig. 7(b) shows the corresponding XPS spectra of Fe 2p core levels. After depositing 0.4 ML Ni, the BE of Fe 2p3/2 has a slight shift towards lower BE. At the same time, the characteristic satellite peak of Fe2+ state at 716.0 eV seems becoming strong relative to the intensity of Fe 2p3/2 line. These results strongly suggest the reduction of part of Fe3+ states in Fe3O4, i.e., a partial reduction from Fe3+ to Fe2+. Combining the partial oxidation of Ni atoms with partial reduction of Fe3O4 at the interface, it is indicated that there is a chemical reaction at the interface of NiO and Fe3O4(1 1 1) buffer layers, which benefits the initial nucleation and growth of NiO films. Generally, the growth mode of metal-oxide films associates with the surface energies of the substrate (Es) and the film (Ef) as well as their interfacial energy (Ei) [34]. In the equilibrium condition of growth, the total energy of the whole system can be expressed as: DE = Ef + EiEs. If DE < 0, the growth of the films follows a layer-by-layer growth mode, while if DE > 0 the films are grown as 3D structure. In general, most of metal oxides have lower surface energies than metals, so that ordered oxide films can be epitaxially grown on metal substrates. In contrast, the clusters or islands of metals are formed on metal oxide substrates [24]. For example, the Ag clusters on MgO(1 0 0) are formed, while MgO(1 0 0) films on Ag(1 0 0) are formed since EMgO(1 0 0) (0.9 J/m2) is lower than EAg(1 0 0) (1.25 J/m2) [5]. However, in oxide-oxide system DE is dependent on the interfacial interaction between the film and the substrate owing to the similar value of Ef and Es [34,35]. In the present work, the reactive layer at the interface between Ni and iron oxide surface tends to the minimization of Ei (0) [36], benefiting the initial nucleation and growth of NiO(1 1 1) films on iron oxide surface. In addition, it is noticed that there is no signal from iron oxide films detected by XPS when the NiO films are >5 ML thick, suggesting that a layer-by-layer mode

most likely occurs for the growth of NiO(1 1 1) films on Fe3O4(1 1 1) buffer layers. Similarly, in our experiments the epitaxial growth of NiO(1 1 1) films on FeO(1 1 1) and Fe2O3(0 0 0 1) buffer layers follows such a growth mode. The NiO(1 1 1) films have the same rock-salt structure as MgO(1 1 1), being polar instability. It is noted that there are no facets or reconstruction observed in LEED pattern of NiO(1 1 1) films, and also no observable impurity in XPS and UPS measurements. Surface defects (e.g., Ni vacancies) or absorption (e.g., hydroxyl) may be responsible for its stability [9,10,37]. Compared with NiO(1 1 1) surface with {1 0 0} facets, the existence of these Ni vacancies induces the decrease of the surface energy, benefiting the nucleation of NiO on Fe3O4(1 1 1) buffer layers. In fact, the relatively big LEED pattern in Fig. 5 suggests the existence of part of surface vacancies. Besides the factor of lattice mismatch, the similar hexagonal structures (Table 1) between iron oxide and NiO(1 1 1) favor the epitaxial growth of NiO(1 1 1) films on iron oxide buffer layers. In addition, it is found that various ordered iron oxide films (including FeO(1 1 1), Fe3O4(1 1 1) and Fe2O3(0 0 0 1)) can be used as the buffer layers to grow NiO(1 1 1) ultrathin films. These results suggest that ordered iron oxide films, without considering the chemical states of iron, are good candidates as buffer layers for the growth of NiO(1 1 1) films because such binary oxide systems is highly correlative. The study of these systems will be helpful to develop novel materials and investigate interesting surface properties.

4. Conclusions Various ordered iron oxide films including FeO(1 1 1), Fe3O4(1 1 1) and Fe2O3(0 0 0 1) have been epitaxially grown on Mo(1 1 0) substrate and were used as the buffer layers for the growth of polar NiO(1 1 1) ultrathin films without {1 0 0} facets. These ordered iron oxide films as the effective buffer layers have two main roles: one is to reduce the lattice mismatch between NiO(1 1 1) and Mo(1 1 0), and another is to decrease the interfacial energy due to the interfacial interaction between Ni and iron oxide surface. The facts suggest that ordered iron oxide films are ideal buffer layers for the growth of NiO(1 1 1) films. The results will

M. Xue, Q. Guo / Journal of Alloys and Compounds 598 (2014) 224–229

be of advantage to further investigate the growth mechanisms and surface properties of polar materials.

[14] [15] [16] [17]

Acknowledgments We gratefully acknowledge the financial support of this work by the Natural Science Foundation of China (Grant Nos. 21273276, 21103084) and National Basic Research Program of China (2012CB921700).

[18] [19] [20] [21] [22] [23] [24] [25]

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]

J. Goniakowski, F. Finocchi, C. Noguera, Rep. Prog. Phys. 71 (2008) 016501. M.S. Xue, Q.L. Guo, J. Chem. Phys. 127 (2007) 054705. C. Noguera, J. Phys.: Condens. Matter 12 (2000) R367. T. Franz, J. Zabloudil, F. Mittendorfer, L. Gragnaniello, G. Parteder, F. Allegretti, S. Surnev, F.P. Netzer, J. Phys. Chem. Lett. 3 (2012) 186. J. Goniakowski, C. Noguera, Phys. Rev. B 60 (1999) 16120. M. Pilard, O. Ersen, S. Cherifi, B. Carvello, L. Roiban, B. Muller, F. Scheurer, L. Ranno, C. Boeglin, Phys. Rev. B 76 (2007) 214436. G. Barcaro, A. Fortunelli, G. Granozzi, Phys. Chem. Chem. Phys. 10 (2008) 1876. G. Peng, L.R. Merte, J. Knudsen, R.T. Vang, E. Laegsgaard, F. Besenbacher, M. Mavrikakis, J. Phys. Chem. C 114 (2010) 21579. W.B. Zhang, B.Y. Tang, J. Chem. Phys. 128 (2008) 124703. S. Wang, S. Liu, J. Guo, K. Wu, Q. Guo, Surf. Sci. 606 (2012) 378. A. Barbier, C. Mocuta, H. Kuhlenbeck, K.F. Peters, B. Richter, G. Renaud, Phys. Rev. Lett. 84 (2000) 2897. A. Bouzoubaa, D. Costa, B. Diawara, N. Audiffren, P. Marcus, Corros. Sci. 52 (2010) 2643. R. Yamauchi, G. Tan, D. Shiojiri, K. Koyama, S. Kaneko, A. Matsuda, M. Yoshimoto, Japan, J. Appl. Phys. 51 (2012). 06FF16.

[26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37]

229

P. Stoyanov, A. Gottschalk, D.M. Lind, J. Appl. Phys. 811 (1997) 5010. M.S. Xue, Q. Guo, K. Wu, J. Guo, J. Cryst. Growth 311 (2009) 3918. M.S. Xue, Q. Guo, K. Wu, J. Guo, Langmuir 24 (2008) 8760. A.P. Grosvenor, B.A. Kobe, M.C. Biesinger, N.S. Mclntyre, Surf. Interface Anal. 36 (2004) 1564. L. Yin, I. Adler, T. Tsang, L.J. Matienzo, S.O. Grim, Chem. Phys. Lett. 24 (1974) 81. M.S. Xue, S. Wang, K. Wu, J. Guo, Q. Guo, Langmuir 27 (2011) 11. J.S. Corneille, J.-W. He, D.W. Goodman, Surf. Sci. 338 (1995) 211. M.S. Xue, Q. Guo, K. Wu, J. Guo, J. Chem. Phys. 129 (2008) 234707. W. Weiss, Surf. Sci. 377–379 (1997) 943. K.J. Kim, D.W. Moon, S.K. Lee, K.-H. Jung, Thin Solid Films 360 (2000) 118. W. Weiss, W. Ranke, Prog. Surf. Sci. 70 (2002) 1. J.F. Moulder, W.F. Stickle, P.E. Sobol, K.D. Bomben, in: J. Chastain (Ed.), Handbook of X-ray Photoelectron Spectroscopy, Perkin–Elmer, Eden prairie, MN, 1992. S. Gota, E. Guiot, M. Henriot, M. Gautier-Soyer, Phys. Rev. B 60 (1999) 14387. V.K. Lazarov, S.A. Chambers, M. Gajdardziska-Josifovska, Phys. Rev. Lett. 90 (2003) 216108. W. Ranke, M. Ritter, W. Weiss, Phys. Rev. B 60 (1999) 1527. V.E. Henrich, P.A. Cox, The Surface Science of Metal Oxides, Academic, Cambridge University, 1994. Yu.S. Dedkov, M. Fonin, D.V. Vyalikh, J.O. Hauch, S.L. Molodtsov, U. Rudiger, G. Guntherodt, Phys. Rev. B 70 (2004) 073405. Y.Q. Cai, M. Ritter, W. Weiss, A.M. Bradshaw, Phys. Rev. B 58 (1998) 5043. R. Zalecki, A. Kolodziejczyk, A. Kozlowski, J. Korecki, N. Spiridis, Z. Kakol, Phys. Stat. Sol. (b) 243 (2006) 103. R. Zalecki, A. Kolodziejczyk, W. Tokarz, A. Kozlowski, Z. Kakol, Acta Phys. Pol. B 34 (2003) 1527. C.T. Campbell, Surf. Sci. Rep. 27 (1997) 1. Sh.K. Shaikhutdinov, R. Meyer, D. Lahav, M. Baumer, T. Kluner, H.-J. Freund, Phys. Rev. Lett. 91 (2003) 76102. Y. Gao, Y.J. Kim, S.A. Chambers, G. Bai, J. Vac. Sci. Technol. A 15 (1997) 332. J. Xing, M. Takeguchi, M. Tanaka, Y. Nakayama, Surf. Interface Anal. 44 (2012) 1483.