鐵超薄膜在鉑(111)上之表面結構與磁性研究

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(1)Chapter 1. Introduction. The perpendicular magnetic anisotropy (PMA) has been acknowledged for the 21st century digital data recording technology [1]. Ferromagnetic thin films and alloys are very promising candidates for high-density magneto-optical recording materials and hard disk drives because they have the property of PMA [2–5]. Fig. 1-1 illustrates the reading/writing process in longitudinal and perpendicular recording [6]. In perpendicular recording, the magnetization of the data bits is aligned vertically, instead of lying in the disc’s plane as it does in longitudinal recording, perpendicular to the disk. The data bits in this manner will take fewer platters, and smaller grains are harder to reverse the magnetic orientation, resulting in smaller physical bits that are still stable at room temperature. So they can be placed closer together on the platter, thus increasing the recording density of perpendicular recording media. Detailed information about the perpendicular recording technology can be obtained from the publications by S. Khizroev and D. Litvinov [7,8].. Figure 1-1: Longitudinal recording diagram (top) and perpendicular recording diagram (bottom) [6].. 3.

(2) There is a known thermal-related magnetization instability [9], which will limit the efforts toward increasing the areal density in ultrahigh magnetic information storage media. The stability factor defined as KuV/kBT must be greater than 50 to avoid thermal instabilities (where Ku: magnetic anisotropy constant, V: magnetic grain volume, kB: Boltzmann constant, T: temperature) [10]. Consequently, many researchers have been focused their studies on the media with high coercivity and high magnetic anisotropy [11-13]. Lyberatos estimated the increase in areal density of heat-assisted magnetic recording can reach as high as ten times when the writing temperature is close to the Curie temperature (TC) [14]. Therefore, TC of a media for ultrahigh density magnetic recording media must be at a temperature not too much higher than room temperature [9,15]. P. Gambardella et al. found that single Co atom and nanoparticles onto Pt(111) have a giant magnetic anisotropy due to strong spin-orbit coupling induced by the Pt substrate [16]. T. Onoue et al. pointed out that the strong PMA and high coercivity property of CoNi/Pt multiplayer are suitable for the formation of small and stable magnetic bits [17]. J. Feng et al. found that acicular shape grains formed more easily in Al–BaM ferrite films than in simple BaM ones, and the Al–BaM ferrite films with high coercivity and large squareness ratio may be applicable as perpendicular magnetic recording layers with low noise level [18]. Recently ultrathin films, multilayers, and alloys of Fe on transition metal surface have been investigated intensively due to their potential of ultrahigh density near 1 Tbit/in.2 as media for perpendicular magnetic recording [19-21]. For example, Araya-Pochet et al. used surface magneto-optical Kerr effect (SMOKE) to study Fe/Ag(100) [22]. They observed that the direction of the easy axis of the magnetization is out-of-plane when dFe < 2 ML, and it changes to the in-plane direction when dFe > 2 ML. Qiu et al. found that the Kerr signal initially increases linearly as a function of the Fe film thickness and is independent of the intervening Ag layers for Fe(110)/Ag(111) single-crystal superlattices [23]. Abid et al. studied the relation between the magnetization and temperature of Fe/Pt multilayers; they found that the thickness of Pt is the key factor to determine the direction of the easy axis of the magnetization [24]. Lim et. 4.

(3) al. observed that the out-of-plane coercivity increases with thickness but reduces drastically when Fe–Pt layer is thicker than 20 nm on CrRu(200) [25]. Sato et al. found that the PMA of Fe layers in (Co, Fe)/Pt is induced by the interlayer interaction between Fe and Co layers via polarization of Pt [26]. Fe-Pt alloy films with high in-plane coercivity can be obtained by annealing the films in the temperature range between 400. ℃ and 650 ℃ [27]. Especially, the L10 phase of FePt (fct-phase) with a high magnetic anisotropy energy constant (Ku ~7 × 107 erg/cm3) has been identified as a candidate material for ultrahigh density recording media [28]. But at room temperature, the fcc-FePt (face-centered-cubic [fcc] structure) does not exhibit a high Ku and the high anisotropic fct-FePt ordered structure (face-centered-tetragonal [fct] structure) is formed after a high temperature post annealing over 600 °C [29]. The high fabrication temperature is not favorable for fct-FePt application and production. Therefore, many studies are developed to lower down the order–disorder transition temperature of the high-anisotropy FePt L10 films [30-33]. In the theoretical study, Paudyal et al. predicted that the average magnetic dipole moment per atom decreases when the composition of Pt increases for the Fe–Pt alloy [34]. The structural determination, the formation of the surface alloy and the magneto-optical properties of ultrathin Fe films on Pt(111) are studied in this dissertation. The fundamental concept and the experimental techniques are introduced in Chapter 2 and Chapter 3, respectively. All experiments were carried out in situ in two ultrahigh vacuum (UHV) chambers which conditions are controlled in the same circumstances for the two cases. The growth mode, structure, and magnetic properties of Fe/Pt (111) at room temperature and low temperature are described in Chapter 4-1.1 to 4-1.3. The annealing effects on surface structure and magnetic properties of ultrathin films Fe/Pt (111) are discussed in Chapter 4-1.4 and 4-1.5. During the study, the phenomenon of a spin-reorientation transition (SRT) induced by Ag overlayers is observed. The increasing of the interface anisotropy is investigated, and the enhancement of PMA in the Ag/Fe/Pt (111) system is discussed in Chapter 4-2. At the end of this dissertation, we make a conclusion of the relation between structure and magnetism of Fe/Pt (111) in Chapter 5.. 5.

(4) Chapter 2 2-1. Fundamental notions. Why is ultrahigh vacuum necessary ? Most of the experiments for magnetic ultrathin films are performed in situ in an. ultrahigh vacuum (UHV) chamber. Because the mean free paths (λ) of free particles such as electrons or ions will be too small for experimental techniques at pressures above about 10-4 Torr. The mean free path of free particles is λ=. kT. (2-1). 2Pσ. P - pressure [ N m-2 ] k - Boltzmann constant ( = 1.38 × 10-23 J K-1 ) T - temperature [ K ] σ- collision cross section [ m2 ]. where :. As an example, the mean free pass of Ar particle is 4.7 cm at 300 K and 10-4 Torr. If the path of Ar is reduced to 10-6 Torr, then λ is 4.7 m. So UHV (<10-9 Torr) is more than enough to get a suitable environment for the electrons (or ions) emitted from the specimen to reach the analyzer. An atomically clean surface is prepared for study, and such surfaces to be maintained in a contamination-free state for the duration of the experiment of surface properties. Even under UHV conditions, a clean surface would become rapidly covered with contaminants from the residual gas atmosphere [35]. For a given set of conditions (P, T etc.) the number n& s of particles striking a surface of 1 m2 per second is readily calculated using a combination of the ideas of statistical physics, the ideal gas equation and the Maxwell-Boltzmann gas velocity distribution.. n&. s. =. 1 nv 4. [ molecules m-2 sec-1 ]. (2-2). where: n = molecular gas density [molecules m-3] v. = average molecular speed [m s-1]. The molecular gas density is given by the ideal gas equation and the mean molecular speed is obtained from the Maxwell-Boltzmann distribution of gas velocities by integration, yielding n = N/V = P/RT. (R: ideal gas constant = k•NA ; NA: Avogadaros' Number) (2-3). 6.

(5) and. ( m: molecular mass [kg] ).. v = 8kT / mπ. (2-4). By combining Eqs. (2-3), (2-4) with Eq. (2-2) can get the Hertz-Knudsen formula for the incident flux. n&. s. =. P 2 π m kT. [ molecules m-2 sec-1 ].. (2-5). Assuming the average molecule weight of gas is 28, one can get v = 5.17 × 102 m/s at T = 300 K and a monolayer (ML) capacity is about 1018 particles/m2. The sticking coefficient (s) is the probability (s = 0 - 1) that an impinging molecule remains adsorbed after striking the surface. Substituting the above values to Eq. (2-5) can yield -1 6 n& s = 10 × s × P [ML mbar sec ].. (2-6). This means that it needs approximately only 1 to 10 seconds for building-up an ML of gas contamination at a pressure of 10-6 Torr. If we want to have enough time to study the surface without the influence of contamination, we need to decrease the pressure of the chamber to at least 10−10 mbar. That is an important reason why ultrahigh vacuum is needed in studying surface properties.. 2-2. Growth mode and surface structure. 2-2.1 Growth modes In general, thin film growth can be simply classified into three-different growth modes [36]. A schematic representation of these three growth modes is shown in Fig. 2-1. Three different modes are described as follow.. Figure 2-1: The schematic plots of three kinds of growth mode. (a) Frank-van der Merwe (FM:. layer-by-layer),. (b). Stranski-Krastanov. Volmer-Weber (VW: island formation). 7. (SK:. layer-plus-island),. and. (c).

(6) (a) Frank-van der Merwe growth mode (FM growth mode): The growth mode is also called layer-by-layer growth. The interaction between substrate and layer atoms is stronger than that between neighboring layer atoms. The subsequent film favors growing after the previous deposited layer formed a complete film. (b) Stranski-Krastanov growth mode (SK growth mode): The thin film grows the first complete layer and changes to the 3-dimensional (3-D) island growth. A certain lattice mismatch between substrate and deposited atoms may not be able to be continued into the bulk of the epitaxial crystal. (c) Volmer-Weber growth mode (VW growth mode): When the interactions between neighboring adatoms are stronger than those of the adatom with the surface, leading to the formation of 3-D adatom clusters or islands.. 2-2.2 Monitoring growth modes There are two analytical techniques as Auger electron spectroscopy (AES) and low-energy electron diffraction (LEED) for us to study the growth mode and to calibrate the deposition rate of adsorbate. AES and LEED data are obtained in situ during film growth in UHV. (1) AES uptake curve AES intensity I can be written as I = I0 • e-h/λ, where h is film thickness, λis the escape depth of Auger electrons characterizing their mean free path in the solid. Assume that the film grows in a layer-by-layer mode for the first n monolayers, the intensity of AES signal at deposition time t can be expressed as Eq. (2-7). [35,37]. {. }. I (t ) = t R I 0 e − n d / λ e − d / λ + I 0 e − nd / λ. (2-7). where : I(t) and I0 - Auger intensities at deposition time t and at the initial stage R - the deposition rate of the adsorbate d – the interlayer distance of the adsorbate. It is easy to see that I(t) is proportional to the deposition time t from Eq. (2-7). As the film finishing a complete layer, the AES uptake curve will change its slope. Hence, the AES uptake curve has a sequence of linear segments for layer-by-layer growth with breakpoints upon completion of a monolayer [38].. 8.

(7) (2) LEED (0,0) beam oscillation The oscillation of LEED (0,0) beam intensity has been used to measure the number of layers of a MBE growth [39,40]. The diagram of LEED diffraction pattern is a result of superposition of the wavefunction of the diffraction electron beam. A sharp and bright diffraction spot results from a well organized surface. By this way, the time evolution of LEED (0,0) beam intensity has local maximum when the adsorbate form a complete layer. Alternatively, its intensity will decrease with the growth of the film without appearance of local maximum. During the measurement of LEED (0,0) beam intensity in our lab, the kinetic energy of the incident electron beam must be chosen carefully to be 60 eV. At this energy level, LEED diffraction pattern should be very surface sensitive. Another way of expressing this requirement is that the out-of-phase condition of diffraction is more sensitive to surface step density on the (0,0) beam [41]. The phase in our experimental conditions is Φ (0, 0) = h∆K⊥ = 2hKcosθ ≈ 5π for 60 eV electron energy and incident angle θ = 5°, where K⊥ is the momentum transferred perpendicular to the surface, and hFe = 0.162 nm, hNi = 0.203 nm and hCo = 2.04 nm are the step height of the adsorbed Ni adatom, Ni adatom and Co adatom, respectively. Fig. 2-2 shows the relation between AES and LEED (0,0) beam intensity to the three kinds of growth mode.. Figure 2-2: The relation between AES and LEED (0,0) beam intensity to the three kinds of growth mode. 9.

(8) 2-2.3 surface free energy and lattice misfit Surface free energy and lattice misfit are usually used to interpret the growth mode of ultrathin metal film on metal surface. (1) Surface free energy: The occurrence of the various growth mode can be made in terms of surface or interface energy γ, i.e., the characteristic free energy (per unit area) to create an additional piece of surface or interface. The growth mode of a film on a particular substrate must follow a universal law which is to minimize the total energy of the system. The energy difference of surface free energy and interface energy can be use to discuss this problem. The excess of free energy per unit area of film/substrate interface is defined as interface energy. Thus, the energy difference ∆γ can be written as Eq. (2-8). ∆γ = γf + γi − γs. (2-8). where γf and γs are the surface free of the adsorbate and the substrate. γi is the interface energy of the adsorbate/substrate interface. If ∆γ < 0, this means that the system can decrease the total energy by decreasing the area of substrate and increasing the areas of film and interface. Hence, the film favors layer-by-layer growth mode. Alternatively, it favors 3-D island growth if ∆γ > 0. Some selected data of our research metals are reported in the following table 2-1. Table 2-1: Surface free energy [42] 2. γ (J/m ). Fe 2.475. Ag 1.30. Pt 2.475. (2) Lattice misfit: A second important parameter determining epitaxial growth is the lattice misfit (lattice mismatch) η. The lattice mismatch parameter η of A atom epitaxial growing on B single crystal is defined as Eq. (2-9), where aA and aB are the atomic nearest neighbor distance (atomic diameter) of bulk A and bulk B.. η=. a A − aB aB. (2-9). If the lattice mismatch between the lattice parameters is not too large, minimizing the. 10.

(9) total energy leads to a situation whereby, below a critical thickness, the misfit can be accommodated by introducing a tensile strain in one layer and a compressive strain in the other such that ultimately the two materials A and B adopt the same in-plane lattice parameter. This regime is called the coherent regime; the lateral planes are in full lattice-registry. Misfits for our research metals on homosymmetric epitaxial substrates are given in the following table 2-2. The deposited film faces a tensile strain if η < 0, on the other hand, it faces a compressive strain if η > 0. Table 2-2: The lattice mismatch parameter η at %. η=. Adsorbate atoms. a A − aB aB. Substrate. aB. Fe aA=2.482 Å. Ag aA=2.89 Å. Pt aA=2.774 Å. Fe. 2.482 Å. 0. 16.44. 11.76. Ag. 2.89 Å. -14.11. 0. -4.01. Pt. 2.774 Å. -10.53. 4.18. 0. 2-3 Basic theories of magnetic ultrathin films When a magnetic field, H, is applied to a material, the response of the material is called magnetic induction B. An easy way to express the relation between B and H is the B - H curve. The ratio of these two quantities, B and H, is the permeability µ (µ = B/H). Magnetization curve of a material is the evolution of magnetic induction inside the material versus external applied field. The magnetic properties of a material are characterized by magnetization M and the way in which M varies with H. Therefore, the relation between B, M, and H can be expressed as B = µ0(H +M). The permeability of free space is µ0 = 4π × 10−7 N/A2. The properties of a material are defined not only by the magnetization, or the magnetic induction, but by the way in which these quantities vary with the applied magnetic field. The ratio of M to H is called the susceptibility: χ = M/H. From the value of susceptibility χ, we can divide the magnetic material into four categories: (1) diamagnetism (χ ≈ -10-5); (2) paramagnetism (χ ≈ 10-3 ~ 10-5, the direction of the magnetization is parallel to the external magnetic field); (3) antiferromagnetism (χ. 11.

(10) ≈ 10-3 ~ 10-5, the direction of the magnetization is antiparallel to the external magnetic field), and (4) ferromagnetism (χ » 1). Now, we just discuss the ferromagnetism.. 2-3.1. Magnetization of ferromagnetism materials. The ferromagnetic material (FM) always has strong dipolar interaction with the neighboring atoms, and this behavior induces the special magnetic properties. Its spontaneous magnetization is one of the most fundamental quantities in magnetism and presents a magnetization much larger than other materials. The property of ferromagnetism is due to the spin and the Pauli exclusion principle. Because of the Pauli principle, two electrons with the same spin state cannot lie at the same "quantum state", and thus feel an effective additional repulsion that lowers their electrostatic energy. This difference in energy is called the exchange energy and induces nearby electrons to align in the same direction. In addition, the dipoles of ferromagnetism tend to align spontaneously without any applied magnetic field. At this place, the magnetization is called spontaneous magnetization. Only atoms with partially filled shells can experience a net magnetic moment in the absence of an external field.. MS. MR. HC. Figure 2-3: The hysteresis loop of a magnetic material showing the variation of M with changing applied field H.. 12.

(11) As the external magnetic field is large enough, then FM will become a single domain. At this time the magnetization is the so-called saturation magnetization Ms. Thus, the magnetization M for a ferromagnetic material is a function of the external magnetic field H. The M-H curves of ferromagnetism form the as known hysteresis loop as shown in Fig. 2-3. MS, MR, and HC are the saturation magnetization, remnant magnetization (magnetization at zero external field), and coercive force (field strength needed to demagnetize the material), respectively. The saturated magnetization is temperature dependence. As the temperature increases, thermal fluctuation decreases the tendency for dipoles to align. When the temperature rises beyond a certain temperature, thermal fluctuation will overcome the exchange energy and the spontaneous magnetization drops to zero. The temperature is defined as Curie temperature (TC). As T near TC, TC is determined from the phenomenological power-law fits to the data, M S (T ) = M S (0)(1 −. T β ) TC. (2-10). where MS (T) is the magnetization proportional to the Kerr intensity and it is the saturation magnetization at temperature T, MS (0) is that of zero absolute temperature and β is called critical point exponent which is dependent on theoretical models [43]. It has been illustrated theoretically using several different model, such as β = 0.125 for 2D Ising model, β = 0.23 for 2D XY model and β = 0.368 for 3D Heisenberg model, respectively [44,45]. Ferromagnetism favors to form many magnetic domains to decrease the magnetostatic energy raised from the spontaneous magnetization. A magnetic domain is a region that all the magnetic dipoles inside it align to a same direction. However, the directions of two neighboring magnetic domains are not the same. The transition between two domains, where the magnetic dipoles rotate, is called a Domain wall. The thickness of a domain wall is on the atomic scale. Those domains whose magnetic dipoles inside them have the same direction with applied magnetic field will expand their area by domain wall motion. Hysteresis phenomenon is a characteristic of this motion.. 13.

(12) 2-3.2. Magnetic anisortropy in ultrathin metallic films. The ferromagnetic single crystals will display ‘easy’ and ‘hard’ directions of the magnetization; i.e. the energy required to magnetize a crystal depends on the direction of the applied field relative to the crystal axes. The preferred magnetic moment orientation in ultrathin magnetic films can be quite different from the factors that account for the easy-axis alignment along a symmetry direction of a bulk material, and the strength can also be markedly different. The preference for a magnetization to lie in a particular direction is called magnetic anisotropy. The magnetic anisotropy is also an important role to discuss the hysteresis loops of the ferromagnetic film. It is very sensitive to the local environment, and subtle changes of the atomic spacing could cause significant effects. It describes the circumstance that the energy of a system changes with a rotation of the magnetization [46]. By varying the thicknesses of the individual layers and choosing appropriate materials, it appeared possible to tailor the magnetic anisotropy. The most important in this respect is the change of the easy axis of the magnetization from the in-plane orientation to the direction perpendicular to the plane. This phenomenon is usually referred to as perpendicular magnetic anisotropy (PMA) and is particularly important for information storage and retrieval applications. The energy required to change the direction of magnetization is called the magnetic anisotropy energy (MAE). If we ignore the higher order terms, it can be written as E = Keff · sin2θ. (2-11). where θ is the angle between magnetization M and the normal to the surface of the sample, as shown in Fig. 2-4. A system prefers a out-of-plane easy axis if its Keff > 0, while it favors a in-plane easy axis if Keff < 0. The effective magnetic anisotropy Keff can be expressed by a volume contribution KV and a interfaces contribution KS. The two contributions approximately obeyed the relation: K eff = KV +. 2K S t. (2-12). where t is the thickness of the magnetic thin film [47]. The factor 2 in the interface. 14.

(13) contribution is corresponding to the two same identical magnetic/nonmagnetic interfaces of multilayers. For example, each period of (Co/Pt)n multilayers has two Co/Pt interfaces and one t thickness of Co. The interface term shows a 1/t response. This is represents the difference between the anisotropy of the interface atoms with respect to the inner or bulk atoms. den Broeder et al. is the first group who used Eq. 2-12 in the determination of KV and KS [48]. The two contributions can be obtained by a plot of the product Keff ·t versus t as shown in Fig. 2-5. The vertical axis intercept equals twice the interface anisotropy, whereas the slope gives the volume contribution. The volume and interface contributions listed in Table 2-3 are useful for us to discuss the direction of magnetization easy axis.. Figure 2.4: Schematic diagram of the angle θ between the magnetization M and the surface normal of the film.. Figure 2-5: MAE times the individual Co layer thickness versus the individual Co layer thickness of Co/Pd multilayers. The vertical axis intercept equals twice the interface anisotropy, whereas the slope gives the volume contribution. Data are taken from den Broeder et al. (1991) [49].. 15.

(14) Table 2-3: Magnetic anisotropy value KVFe K SFe −UHV K SPt − Fe K SFe − Ag. Reference. -2.8 MJ/m3. [50]. 0.89 mJ/m2 0.42 mJ/m2 1.45 mJ/m2. [51] [50] [52]. In Fig. 2-5, the negative slope indicates a negative volume anisotropy KV, favoring in-plane magnetization; while the positive KS, favoring perpendicular magnetization. Below a certain thickness t⊥, the surface term dominates, and it results in a perpendicularly magnetized system. If KV and KS are in different signs, Keff decrease as the thickness t increases, then it will change its sign as t increases above t⊥ , at this time the easy axis of the magnetization will change from perpendicular to in-plane, this phenomenon is the so-called spin reorientation transition (SRT). Recently, the SRT of the magnetization in ultrathin ferromagnetic films has attracted tremendous attention. It is known that the SRT marks the magnetization switching with film thickness, temperature, non-magnetic overlayer, alloy composition of thin films [53-57]. An intimate relationship behaved between the representation of SRT and the Keff. The magnetic surface anisotropy is resulted from symmetry-breaking around surface atoms, it is a straightforward consequence of the fact that surface atoms are located in a different environment than the bulk ones. The low symmetry behavior raises a strong magnetic anisotropy which contributes to the surface term KS. The two main sources of the magnetic anisotropy are the magnetic dipolar interaction and the spin–orbit interaction. The spin-orbit interaction is the primary source of the magnetocrystalline anisotropy. If the magnitude of the dipolar interface contribution is of minor importance, the spin-orbit coupling appears to be dominant. Consequently, the spin-orbit interaction is more important for the magnetocrystalline anisotropy in Fe, Ni and Co. The effects of this anisotropy on the magnetization curves are large different along distinct axis of crystal. For the ultrathin film, the reduced symmetry at the surface should result in magnetic anisotropy at the surface differing. 16.

(15) strongly from the bulk, this is the magnetrocrystalline surface anisotropy, and it may be responsible for the perpendicular magnetization. The shape anisotropy, which is associated with the demagnetizing field Hd of the ellipsoidal ferromagnetic specimen, is ascribed to the dipolar interaction. For the rough surface the magnetic dipolar interaction become strong, and Hd is different along distinct axis except for cubic symmetry, thus shape alone is a source of magnetic anisotropy. The Magnetic dipolar anisotropy is raised from magnetostatic energy. The magnetostatic energy results from the magnetic poles on the surface. When a magnetized material has magnetic flux lines flow through its surface with normal component, there exist free standing poles on this surface. It results in an internal magnetic field, demagnetizing field, exists inside the material. Thus, the magnetic field really passes through the material is the sum of external applied field H and demagnetizing field Hd. The strength of Hd is related to the shape of the sample. For a thin film, the magnetostatic energy density can be expressed as Eq. (2-13). Ed = −. µ0 1 = µ 0 M S2 cos 2 θ 2V 2. (2-13). Here, θ is the angle between M and surface normal. Because the thickness of magnetic film does not affected Ed, thus, the magnetostatic energy contributes only to volume anisotropy KV . To minimize the magnetostatic energy, the magnetic dipolar anisotropy favors an in-plane easy axis (θ = 0º). The magneto-elastic anisotropy is the strain- induced anisotropy, which also play an important role for ultrathin ferromagnetic films. The energy per unit volume associated with this effect, for an elastically, isotropic medium with isotropic magnetostriction, can be written as Eme = 1.5 ⋅ λσ sin 2 θ = K me sin 2 θ , where λ and σ are the magnetostriction constant and stress, respectively. θ is the angle between the magnetization and stress directions [58]. The magnetostriction result in magnetic anisotropy can be expressed as anisotropy constant Kme. For positive, the easy magnetic direction will be along a direction of tensile stress, or perpendicular to a compressive stress. When the parameters are constant this contribution can be identified with a volume contribution KV.. 17.

(16) Chapter 3 Experiment Experiments are performed in the stainless steel ultrahigh vacuum (UHV) chamber. There are two ultrahigh vacuum (UHV) chamber in our lab as shown in Fig. 3-1 and Fig. 3-2. The background pressures of the both UHV chambers are 5 × 10−10 Torr. The basic analyzing tools are Auger electron spectroscopy (AES), low-energy electron diffraction (LEED), ultraviolet photoelectron spectroscopy (UPS), and magneto-optical Kerr Effect (MOKE). One of the UHV chamber is used to study the surface structure and the other is focused on the evolution of magnetic properties. We use these tools to study surface structure and magnetic properties of ultrathin Fe films on Pt(111) surface. In this chapter, the sample preparation, instruments of our UHV chamber and experimental techniques are discussed.. 3-1 Sample preparation The substrate surface of Pt(111) (miscut angle < 0.5º) is prepared using sputtering-annealing cycles. To remove the residual carbon, the sample is heated to 800 K in oxygen at a pressure of 1×10−7 Torr for about 5 minutes before the sputtering process. The bombardment energy and ionic current of Ar ions used in the sputtering are 2 keV and approximately 2 – 4 µA, respectively. After the sputtering procedure, the substrate is annealed at a high temperature of 1000 K for one hour, and then it cools slowly to room temperature. The chemical impurities on the surface are checked by AES. The cleaning cycles are repeated until a sharp p(1×1) threefold LEED pattern is clearly observed. Subsequently, Fe atoms are deposited on the Pt(111) surface by molecular beam epitaxial (MBE). Fe atoms are evaporated from homemade coils as shown in Fig. 3-3 (1 mm in diameter) with high purity (99.995%) to a well prepared Pt(111) surface. Throughout the experiments, the deposition rates are calibrated by peak-to-peak intensities in the plot of Auger signal versus deposition time and the oscillation of LEED (0,0) beam intensity for out-of-plane electron.. 18.

(17) Figure 3-1: UHV Chamber for studying surface structure.. 19.

(18) Figure 3-2: UHV Chamber for studying magnetic properties.. Figure 3-3: Home-made deposition sources of Fe.. 20.

(19) 3-2. Analyzing instrument. 3-2.1 Auger electron spectroscopy (AES) AES is a standard technique in surface and interface physics. Its sensitivity is about 1% of a monolayer. The information accompanied with Auger analysis comes from the top 2-10 atomic layers. It is very suitable for studying the chemical composition of solid surfaces. Auger process is a secondary electron excitation process which is schematic as follows. Fig. 3-4 is an example of the KLL AES process. A primary electron beam with energy 3-5 keV is used to excite the electron of K-shell. And then, an electron in the L1-shell falls to fill the hole in the K-shell and the energy emitted in this process is simultaneously transferred to a second electron to overcome the binding energy. The remainder energy is retained by this emitted Auger electron as kinetic energy. In the Auger process illustrated, the final state is a doubly-ionized atom with core holes in the L1 and L2,3 shells. But, we still can make a rough estimate of the kinetic energy of the Auger electron from the binding energies of the various levels involved. For the particular example discussed in Fig. 3-4, (3-1). K E = E L1 − E K − E L2 ,3. where EK, EL1, and EL2,3 are the binding energy of K, L1, and L2,3 levels, respectively. It is easy to see that the KE is dependent only on the energy levels of the observed element, and it is independent on the energy of the incident electron beam. Thus, AES is sensitive to all elements except hydrogen and helium.. Figure 3-4: Scheme of the Auger electron emission process.. 21.

(20) In the Auger electron spectra, Auger peaks are superimposed on a rather large continuous background. This background can be removed by differentiating the energy distribution function N(E). Thus the conventional Auger spectrum is the function dN(E)/dE. The peak-to-peak magnitude of an Auger peak in a differentiated spectrum is directly related to the surface concentration of the element which produces the Auger electrons. This can be simply proved by an energy distribution function with Gaussian distribution. For a Gaussian function F ( x) = Ae − x S = A∫. ∞. −∞. 2. e − x dx = A π .. 2. , the area under F is given by: (3-2). 2. The differentiation of F(x) is dF(x)/dx= − 2 Axe − x . Its maximum and minimum are A. 2 e. and − A. 2 , e. respectively. Thus the peak-to-peak height L equals to 2 A. 2 . e. Comparing L and S, L is proportional to S. It is proved that the peak-to-peak magnitude of an Auger peak in a differentiated spectrum is proportional to the surface concentration of the corresponding element. Fig. 3-5 (a) shows the distribution of Auger number N(E) vs energy and Fig. 3-5 (b) shows the same spectrum in such a differentiated form.. Figure 3-5: (a) Auger spectrum of Pd metal - generated using a 2.5 keV electron beam. (b) Differential spectrum of Pd metal - generated using a 2.5 keV electron beam [59].. 22.

(21) 3-2.2 Ultraviolet Photoelectron Spectroscopy (UPS) UPS is a technique for measuring the energy spectrum of electrons emitted during the absorption of ultraviolet radiation. A Helium lamp emitting at 21.2 eV (He I radiation) is used to be the source of UPS. Ultraviolet light incidents to the sample, the low photon energy in UPS means that deep core electron levels cannot be excited and only photoelectrons emitted from the valence band or shallow core levels are detected. Fig. 3-6 illustrates the various electronic energy levels involved in an UPS experiment. The electron absorbs the energy of the incident ultraviolet radiation hν, emitted from the core levels, and reaches the Fermi surface EF. Before reaching the vacuum level, it faces an energy barrier so called the workfunction e φ of the sample. The kinetic energy of the electron which is emitted from the sample and reaches the vacuum can be expressed E = hν − Eb − e φ. (3-3). where Eb is the energy barrier between the core level and the Fermi level. When an UPS experiment is operated, the sample is grounded simultaneously with the spectrometer. Therefore, the Fermi level EF of sample and spectrometer coincide. The maximum kinetic energy measured by spectrometer is Ek (max) = hν − e φ sp. (3-4). in which Eb = 0, where e φ sp is the workfunction of the spectrometer. The cut off kinetic energy measured by the spectrometer is dominated by the secondary electrons which were produced by the photoelectrons. The minimum kinetic energy Ek(min) measured by the spectrometer occurres when the kinetic energy of the secondary electron at the sample surface is zero, thus, Ek(min) can be expressed as Ek(min) = e φ − e φ sp.. (3-5). So, the workfunction of the sample can be obtained by subtracting Eq. (3-5) from Eq. 3-4. (3-6) e φ = hν − (Ek(max) − Ek(min)) = hν − ∆E Due to the analyzer is less sensitive to the electrons of low kinetic energy, one can place the sample at a negative potential of several volts with respect to the analyzer. In this way, the whole spectrum will be shifted towards a value of higher EK but let the ∆E remains the same.. 23.

(22) Figure 3-6: The schematic plot of the energy levels involved in an photoemission from solid of UPS experiment.. 3-2.3. Concentric hemispherical analyzer (CHA). The electron energy analyzer for both AES and UPS is the CLAM-2 hemispherical analyzer of VG MICROTECH. The basic form of a CHA is shown in Fig. 3-7. Two concentric hemispheres of radius R1 (inner) and R2 (outer) are respectively at electric potentials −V1 and −V2 (V2 > V1). An equipotential surface (−V0) between the. 24.

(23) hemispheres has radius R0. Thus, the potential V0 can expressed as Eq. (3-7). V0 = (V1R1 + V2R2 ) / 2R0. (3-7). At first, the electrons must pass through a retard voltage VR which is between the electric lens and the entrance of the hemisphere, after which they pass through the tunnel between the concentric hemisphere plats If the electrons have kinetic energy of E = VR +Ep +W (W is the work function of CHA) in which Ep = eV (0) are injected tangentially to the surface of R0, they will move approximately in a circular orbits along the R0 surface. The base resolution ∆Eb of CHA is not only related to passable kinetic energy Ep of the electron, but also related to the widths of the entrance and exit. It is convenient to make the widths of the entrance (w1) and the exit (w2) to be equal (w1 = w2 = w). From a detail derivation [60], the relation between ∆Eb, Ep, W, and R0 can be expressed as Eq. (3-8). ∆Eb / Ep = 0.63w / R0. (3-8). Since R0 and w are related to the dimension of CHA, thus we can only control the value of Ep ( i.e. adjust V1, and V2) to specify the base resolution. By varying VR, V1, and V2, one can analyze the energy distribution of the electrons which are emitted from the sample surface.. Figure 3-7: The structure of concentric hemispherical analyzer (CHA).. 25.

(24) 3-2.4. Low-energy electron diffraction (LEED). LEED utilizes the wave nature of electrons. Electrons, with energies between 30 and 600 eV are shot in the form of a thin beam onto a surface, as shown in Fig. 3-8. The electrons will be partially reflected on the layers of a crystalline material, just like X-rays. The periodic arrangement of the atoms in the crystalline material determines the direction towards which electrons are diffracted. These back-diffracted electrons are made visible on a phosphor screen. The matter wavelength of an electron can be expressed by the de Broglie’s relation λ = h/p =. h 2me E. =. 150.4 Å E (eV ). (3-9). where E is the kinetic energy of electron in eV. The wavelength of electrons (1−3 Å) is comparable with atomic spacing. This is the necessary condition for diffraction effects which is caused by atomic structure. In order to get a LEED pattern, the sample itself must be a well-ordered surface structure. The structure of experiment is shown in Fig. 3-8. The electrons scatter from the sample surface and pass through a retarded voltage. The retarded voltage Ur is used to select only the electrons which elastically scatter from the surface. Those electrons passed though Ur are then accelerated to strike a fluorescent screen, this cause the phosphor to glow. The revealing pattern of dots is the diffraction pattern. Some of the simple surface structures can be reached by analyzing the diffraction pattern as shown in Fig. 3-9.. Figure 3-8: Schematic structure of RFA for LEED apparatus. 26.

(25) Figure 3-9: Schematic LEED images for fcc(111), fcc(110) and fcc(100) surface structure.. 3-2.5. Surface magneto-optical Kerr effect (SMOKE). The magnetic properties of thin film system are studied by means of the SMOKE technique. Magneto-optical (MO) effects in magnetic materials can serve as very sensitive probes of the magnetic order. The discovery of the MO effect in 1876 John Kerr observed the rotation of the polarization plane of linearly polarized light upon reflection from the surface of a magnetic piece of iron [61]. Thus, the dependence of the polarization state of light reflected from ferromagnetic surfaces on the magnetization state of the condensed matter, called magneto-optical Kerr effect (MOKE), has been used for many years to observe magnetic domains at bulk magnetic surfaces. The microscopic origin of the MO effects comes from the spin-orbital interaction. The electric and magnetic fields of the incident light affect the orbital motion of the electron, which in turn is coupled via their spins to the magnetization. After Bader et al. noted its easy use to take hysteresis loops from ultrathin films, the effect, called now the surface magneto-optical Kerr effect (SMOKE), has found wide application in ultrathin film magnetism, down to the monolayer [62]. As shown in Fig. 3-10, when a polarized beam of light propagates in the a magnetic field B inside a magmetic material, the material. 27.

(26) interacts with the right and left circularly polarized components of the beam with different refractive indices. These two circular polarizations are reflected with different reflectivity because the Fresnel reflection coefficients depend on the refractive index. When. there. is. a. phase. difference. exists,. the. reflected. beam. becomes. ellipitically-polarized. Thus, the reflection of polarized light from the surface of a magnetic material carries information of the magnetization of the material. This phenomenon is known as the magneto-optical Kerr effect and is used in magneto-optical recording systems. There are three varieties of operation modes for surface magneto-optical Kerr effect: polar, longitudinal and transverse, as shown in Fig. 3-10. In the polar Kerr effect, the magnetization is perpendicular to the surface of the sample, and consequently lies parallel to the plane of the incidence. In the longitudinal one, the magnetization lies parallel to both the sample surface and the incidence plane. The magnetization of the transverse geometry lies in the surface of the sample and is perpendicular to the incidence plane. In the polar and longitudinal configurations, the MO effect consists of an M-dependent change in elliptical polarization of the reflected beam.. Figure 3-10: Scheme of the three varieties of operation modes for surface magneto-optical Kerr effect. θK is the so-called Kerr rotation angle and it is proportional to the magnetization of the sample.. 28.

(27) In Fig. 3-10, incident light is labeled as ”P-polarized” when its E field is in the plane of incidence. The component of which its E field is perpendicular to the plane of incidence is called ”S-polarized”. The reflected light is elliptical polarized with its principle axis has a Kerr rotation angle angle (θK) with respect to the plane of incidence. From the theoretical derivaiton [23], the Kerr rotation angle θK and the Kerr ellipticity εK are proportional to the magnetization of the sample, and the relation between the components of the electric field, EP and ES, can be written as EP / ES = θK + iεK.. (3-10). In ordered to plot the hysteresis loop, we make the direction of the analyzer has a small angle δ (θK < δ < 1º ) with respect to the S-polarized direction of the incident light. The intensity I of the reflected light passes though the analyzer can be written as Eq. 3-10 in which sinδ ≈ δ and cosδ ≈ 1. I = EP. 2. (δ. 2. + 2δθ K. ). (3-11). Let I0 be the intensity with zero magnetization, the difference in intensity between I and I0 can be written as Eq. (3-11) and θK can be derived to has the form in I = I − I0 = |EP |2 (2 δ θK) = I0 ( 2 θK / δ ) θK = (δ / 2)(∆I / I0) ∝ Magnetization.. (3-12) (3-13). Therefore, one can describe the evolution of magnetization versus applied magnetic field by measuring the S’-polarized component of the reflected light. A schematic figure of the experimental arrangement is shown in Fig. 3-11. We use a green light with its wavelength of 532 nm as the source light of the SMOKE experiment. The light incident onto the sample surface and is reflected by the surface. The intensity of the reflected light is measured by a photodetector. The evolution of the intensity versus the applied magnetic field is plotted by the computer.. 29.

(28) Figure 3-11: Schematic diagram for the SMOKE setup.. 30.

(29) Chapter 4 Results and Discussion 4-1. The Growth Mode and Magnetic Properties of Fe Ultrathin Films on Pt (111). The growth of epitaxial monolayers of thin film, alloy, and multilayer has direct influence on the electronic state and magnetic properties. To study the growth behavior and alloy formation of Fe films on Pt can be beneficial to understand these magnetic properties of ferromagnetic systems. In this section, Auger electron spectroscopy (AES) and low-energy electron diffraction (LEED) are used to study the initial growth of Fe ultrathin films deposited on a Pt(111) surface. Some interesting structural phases are observed. The alloy formation at different coverage is investigated.. 4-1.1. The growth mode of ultrathin Fe Films on Pt(111). AES uptake curve is used to study the growth mode when Fe ultrathin films are grown on Pt(111) at room temperature. Auger signals of Pt NNV 237 eV and Fe LMM 703 eV as a function of the deposition time are shown in Fig. 4-1 (a). The Auger intensities of the two curves are linear relationship with the deposition time in the time intervals of 0–6 minute, 6–12 minute, and 12–18 minute, with the exception of two kinks located at 6 min and 12 min. These curves then change to an exponential tendency after t > 18 minute. The change of the slope on the Auger uptake curve is interpreted as the layer-by-layer growth of a complete atomic layer on a flat surface [63]. This indicates that the growth mode of 0–3 ML Fe/Pt(111) is in layer-by-layer mode, after which it turns to three-dimensional (3-D) island growth. The result shows that the growth of Fe deposits on Pt(111) surface at room temperature is the Stranski-Krastanow (SK) mode with the deposition rate of 6 min/ML.. 31.

(30) Figure 4-1: (a) Auger uptake curves for Pt NNV 237 eV and Fe LMM 703 eV Auger peak intensity versus deposition time. The break points at t = 6, 12 and 18 minutes correspond to 1.0, 2.0 and 3.0 ML of Fe coverage, respectively. (b) The intensity of the specular beam of LEED as a function of deposition time (I-t profile). The beam energy of 72 eV is used. Three peaks locate at the same deposition times that of AES break points. The oscillation of LEED intensity has been used to measure the number of layers of material deposited during molecular beam epitaxy (MBE) growth [39]. The peak height intensity of LEED (0,0) beam versus deposition time is shown in Fig. 4-1 (b). The energy of normally incident electrons is 72 eV. The out-of-phase condition,φ(0,0) = ∆k⊥•h = n π, n: odd number, is taken in our experiment, where ∆k⊥ is the momentum transfer. 32.

(31) perpendicular to the surface, and h ≈ 2 Å is the step height of the adsorbed Fe atoms. Since the specular beam is a direct measurement of the momentum change perpendicular to the surface, it is more sensitive to the surface step formation and can oscillate in a layer-by-layer growth [64]. In Fig. 4-1(b), the maximum of the intensity means that one monolayer grows completely, so that 1 ML, 2 ML and 3 ML Fe thin films occur at 6, 12 and 18 minutes, respectively. The result of LEED (0,0) beam is consistent with that of AES. This indicates that the growth mode of 0-3 ML Fe/Pt(111) is in layer-by-layer mode, after which it turned to 3-D island growth. Before we obtain the conclusion of the growth mode, further analysis is necessary. The number of the Auger electrons emitted from a solid surface decline with the thickness of the films after they pass through. The AES intensity I can be written as I = I 0 e − h / λ , where h is film thickness, λ is mean free path of Auger electron. The AES intensity of Pt substrate is related to not only the film thickness of Fe but also the growth mode of the Fe films. Assume that the Fe film grows in a layer-by-layer mode for the first n monolayers, the intensity of Pt 237 eV Auger signal at deposition time t can be expressed as Eq. 4-1 I t = I 0 e − ( n −1) hFe / λ237 eV { x [t − (n − 1) t1 ML ] e − hFe / λ237 eV + [1 − x(t − ( n − 1) t1ML )] } 144424443 144 42444 3 144 42444 3 (A). (4-1). (C ). ( B). where n is the nth Fe layer at deposition time t; I0 is the Pt 237 eV Auger intensity at t = 0. The deposition rate and the time needed to build up one Fe monolayer are denoted as x and t1 ML , respectively. From Fig. 4-1 (a), t1 ML is equals to 6 minutes. The distance between two adjacent Fe layers, hFe is about 2 Å in the fcc [111] direction. λ237 eV = 5.35 Å is the mean free path of Pt 237 eV Auger electron. The term (A) in equation (1) is the AES intensity when the Pt substrate builds up ( n − 1) ML Fe layers. Term (B) and term (C) in Eq. 4-1 are the percentages of covered and uncovered surfaces by the nth Fe layer at deposition time t, respectively. Eq. 4-1 is correct only under the condition of layer-by-layer growth, i.e. n ≤ m . After substituting t1 ML with 6 minutes, Eq. 4-1 can be rewritten as following. I 0e. It − nh Fe / λ237 eV. = − x [t − 6(n − 1)](e hFe / λ237 eV − 1) + e hFe / λ237 eV. 33. (4-2).

(32) From Eq. 4-2, it is easy to see that the left-hand side I t /( I 0 × e − nh Fe / λ 237 eV ) is proportional to t − 6( n − 1) when n ≤ m . The evolution of I t /( I 0 × e − nh Fe / λ 237 eV ) versus t − 6( n − 1) is shown in Fig. 4-2, in which n is equa1s to 1, 2, 3, 4, and 5 at time intervals from 0 to 6 minutes, 6 to12 minutes, 12 to 18 minutes,18 to 24 minutes, and 24 to 30 minutes, respectively. The curves are linear and parallel each other when n ≤ 3 only. This proves. that the growth of ultrathin Fe films on Pt(111) is layer-by-layer for the first three monolayer, after which it turns to three-dimensional island growth.. Figure 4-2: The evolution of Pt 237 eV Auger intensity I t /( I 0 × e − nh Fe / λ 237 eV ) versus each t-6(n-1) minutes of deposition time. From the results of AES and LEED, we can conclude that the growth of Fe/Pt(111) is the SK mode. It grows 3 ML in a layer-by-layer mode, and then turns into 3-D island growth. This conclusion is agreed with the result of STM images [65]. The first peak height of the maxima is close to the original intensity of the substrate. It represents that the first layer of Fe is in an ordered state. The growth mechanism of the ultrathin Fe film on a flat Pt(111) at room temperature can be considered from two important parameters: lattice mismatch and surface free fcc energy. From the viewpoint of the lattice mismatch, the lattice constant of fcc Fe, a Fe =. 3.59 Å [66] and of Pt, a Ptfcc = 3.92 Å, The lattice mismatch between Fe and Pt is -10.53%. It implies a considerable epitaxial tensile strain of the first Fe layer. This indicates that the system favors 3-D island growth. The other viewpoint is based on the surface or. 34.

(33) interface energy γ, according to Young’s equation ∆γ = γFe + γint - γPt, where γFe=2.475 J/m2, γPt=2.475 J/m2 and γint are the surface free energy of Fe, Pt and the interface energy, respectively [42]. The γint is small and usually neglected in the bimetal system. The three-dimensional. island. growth. is. favored. for. ∆γ. >. 0,. while. the. monolayer-by-monolayer growth is favored for ∆γ ≤ 0. For ∆γ = γFe - γPt = 0, similar surface energies make the growth of epitaxial Fe films on Pt substrate relatively easy. It favors the monolayer-by-monolayer growth. Due to the competition of surface free energy and lattice mismatch, it is reasonable that the ultrathin Fe films turn to 3-D island growth after growing 3 ML on Pt(111) in the layer-by-layer mode.. 4-1.2. The surface structure of ultrathin Fe Films on Pt(111). The structural evolution of Fe/Pt(111) at different coverage is observed by LEED. The LEED pictures and corresponding schematic patterns are shown in Fig. 4-3. When the coverage of Fe is below 2 ML, the p(1×1) LEED spots are observed as shown in Fig. 4-3 (b). This is a pseudo-(1×1) structure because the LEED spots are slightly diffused if compared to the clean Pt(111) LEED pattern. When the coverage is near 2 ML, The LEED pattern shows a marked change. There appear to be four satellite spots around each substrate spot, with two at the inner circle and two in the outer circle. The same LEED pattern has also been found on the surface of an 2.7 ML thick Fe film on Cu(111) and on the surface of thicker films of Fe on Ru(0001) [67,68]. These additional spots have been interpreted that these patterns are those expected from a uniform distribution of bcc Fe(110) grown onto the hexagonal net of pseudomorphic Fe on Pt(111) in the Kurdjumov-Sachs (KS) orientation [69]. Using the bulk parameters of Fe and Cu, J. Shen et al. have simulated the LEED pattern of the KS orientation which is shown in the upperright of Fig. 4-4 (a). Except one spot in the inner circle of the simulated pattern which is too close to substrate spot, will combine with substrate spot experimentally. So we can find only four satellite spots around each substrate spot. The experimental and the simulated LEED patterns agree well with each other. The six satellite spots reflect the six kinds of atomic relationship between the bcc(110) and fcc(111) structure as shown in the. 35.

(34) bottom picture of Fig. 4-4 (b).. Figure 4-3: LEED patterns for different coverage Fe on Pt(111) are observed at room temperature. Below 2 ML the films have a p(1×1) pattern inherited from the substrate. From 2 ML on, satellite spots start to appear around the substrate spots. In the schematic LEED diagrams, the area of the circle is proportional to the spot size of LEED.. 36.

(35) (b). (a). Figure 4-4: Demonstration of the Kurdjumov-Sachs superstructure. (a) Simulated LEED pattern of Fe bcc(110) on Cu(111) with Kurdjumov-Sachs orientation. (b) Real-space schematic view of the six KS domains on the Cu(111) substrate. The KS orientation is a special case of the one-dimensional matching between bcc (110) and fcc(111) with sides of the rhombic unit meshes of the film and the substrate to be parallel [70]. The structural transition from fcc to bcc is further analyzed by an LEED I-V study of the system. Fig. 4-5 shows the (0,0) beam intensity versus energy curves for films with different thickness. These curves are obtained with the incident angle of the primary beam to be about 5° off the surface normal. When Fe atoms deposit on Pt(111) surface, the peaks have clearly shifted towards higher energies. Such a shift, within the kinematic approximation, indicates that the average interlayer distance of the films becomes smaller than that of the substrate. The interlayer distance d of subsurface region can be determined from the I−V profile of the LEED specular beam. From the Bragg diffraction condition 2dcosθ = nλ = nh / 2meE , one can obtain the useful formula [71, 56]: E = n2 (. h2 ) 8me d 2 cos 2 θ. (4-3). where θ = 5º is the angle between the directions of incident electron beam and surface normal in our experiment, E is the kinetic energy of the incident electrons, me is the electron mass, and h is the Planck constant. The interlayer distance d can be calculated from the slope of the E − n2 plot. The relation between E and n2 is perfectly linear, and. 37.

(36) therefore the interlayer distance can be determined. We have calculated the interlayer distance as a function of the film thickness. The results are shown in Fig. 4-6. Two regions can be clearly distinguished as indicated by the dashed line: in the low thickness region where the films show a p(1×1) pattern the interlayer distance of the films is close to 2.073 Å of the fcc(111) Fe, while in the high thickness region the interlayer distance of Fe films becomes 2.035 Å apparently smaller. This distance indicates that the surface of film is bcc(110).. Figure 4-5: LEED I-V spectrum of the (0,0) beam for the Fe/Pt(111) system. At higher thickness the peaks shift towards higher energies. The dashed lines indicate the peak position of the substrate.. 38.

(37) Figure 4-6. The interlayer distances of x ML Fe on Pt(111) as a function of Fe thickness. The vertical dashed line separates films with different LEED structure. The Fe films have a fcc-like structure below 2 ML, and a bcc structure above. UPS is a powerful surface analysis tool of electronic structural identification. The change of density of state is very sensitive to the surface structure and surface composition. The evolution of the electronic structure in the valence band is interesting during the Fe initial growth. He I (hν= 21.2 eV) ultraviolet resonance radiation is used for the experimental light source. We measure the UP spectra of different thickness of Fe (dFe) grown on Pt(111). The results are shown in Fig. 4-7. The sharp Fermi edge and the d-band peak with a binding energy of 4.2 eV (A peak) are observed on the clean Pt(111) surface. The peak heights of the Fermi edge and the d-band decline, and their peaks broaden gradually as the coverage of Fe increases. The binding energy of the d-band shifts from 4.2 to 3.9 eV as the coverage of Fe increases to 1.5 ML. In order to observe the variations of these UP spectra more clearly, difference spectra (the spectrum of the Fe overlayer minus that of the clean surface of Pt) are plotted in Fig. 4-8. Due to the fact that the peak height or intensity of the UPS spectra is similar to the Auger peak intensity, the magnitude of the received signals is proportional to the intensities from one specific chemical element. At first, both peaks near the Fermi edge will combine to form one peak. 39.

(38) as soon as Fe atoms deposit on Pt surface. It indicates that the interaction between Pt and Fe atoms is very significant. Secondly, the peak A of binding energy 4.2 eV diminishes and shifts as the Fe coverage increases. Thirdly, the peak B of binding energy 2.9 eV shifts to lower binding energy as the Fe coverage increases. Specifically, peaks A and B will disappear when the Fe coverage increases to 2 ML. We can suppose that the surface structural transition from fcc to bcc arises as the Fe coverage arrives 2 ML. The result of UPS is consistent with that of LEED patterns and LEED I-V observations.. Figure 4-7: The variations of UPS spectra of Fe/Pt(111) correspond to different coverage of Fe (dFe). Peak A is located at the d-band peak of Pt with a binding energy of 4.2 eV and peak B is located at the d-band peak with a binding energy of 2.9 eV.. 40.

(39) Figure 4-8: The corresponding UPS difference spectra correspond to Fig. 4-7. Peak D is the minimum in the d-band of Pt with a binding energy of 4.2 eV. The UPS is useful in studying the valence band near the Fermi level. Since the energy of an electron on a conductor’s surface is affected by the surface’s properties, the work function may change whenever the surface changes. The work function is hard to measure due to the analyzer in our lab is less sensitive to the energy region of secondary electrons. A directed evidence of the change of work function is the electron kinetic energy located at Fermi edge (Fermi energy; E KFermi ) at the UPS spectra. Thus, we discuss the change of E KFermi , this helps us to understand the change of work function during the Fe films growth processes. UPS results near the Fermi edge of Fe/Pt(111) films are shown in Fig. 4-7. The electron kinetic energy located at Fermi edge E KFermi moves to the left (lower electron kinetic energy) when the Fe coverage is 0.5 ML and moves back to the right (higher electron kinetic energy) at dFe = 1 ML, after which it decreases linearly with dFe to a minimum at dFe = 4 ML then moves back to the right at dFe = 5 ML. From the shift of E KFermi of the different Fe coverage on Pt(111) in Fig. 4-7, one can obtain the change of the work function as a function of Fe coverage. The result is shown in Fig. 4-9. It is shown that work function decreases approximately 0.1 eV at dFe = 0.5 ML and. 41.

(40) slightly increase 0.07 eV at dFe = 1 ML. After dFe > 1 ML, the work function change will get more minus until dFe = 4 ML. The evolution of work function change versus the coverage of Fe between 1 ML ≤ dFe ≤ 5 ML is consistent in the mode for work-function changes due to adsorption by metals [72].. Figure 4-9: The evolutions of the change in work function versus different coverage of Fe correspond to Fig. 4-7.. 42.

(41) 4-1.3 Magnetic property of ultrathin Fe Films on Pt(111) We use SMOKE to measure the hysteresis loops in the polar and the longitudinal direction during the initial deposition of Fe on Pt(111). Fig. 4-10 (a) shows the evolutions of both polar and longitudinal Kerr hysteresis loops versus the thickness of Fe films. In Fig. 4-10 (a), all the loops were measured at room temperature and the maximum applied magnetic field we used was 300 Oe. No Kerr signal is observed when the thickness of Fe (dFe) ≤ 1.0 ML, even the sample was cooled down to 165 K. We also measured the Kerr signal with a maximum applied magnetic field of 900 Oe, there still no Kerr signal could be observed. The Kerr hysteresis loops measure at 165 K of 1 ML Fe/Pt(111) are shown in Fig. 4-11. It is possible that the Curie temperature is below 165 K or the coercivity is higher than 900 Oe for 1.0 ML Fe/Pt(111). The other possible reason is that the 1.0 ML Fe/Pt(111) system is not ferromagnetic. Both the polar and the longitudinal Kerr hysteresis loops can be obtained when dFe = 1.5 ML. Since the coercivity (HC) of the polar hysteresis loop is much greater than the in-plane one, the easy axis of the magnetization is in the in-plane direction. Only longitudinal Kerr hysteresis loop can be observed when dFe ≥ 2.0 ML, i.e., the easy axis is in the in-plane direction. The thickness dependences of the saturated polar and longitudinal Kerr intensities for Fe/Pt(111) are shown in Fig. 4-10 (b). The longitudinal Kerr intensities increase with the thickness of Fe films when dFe ≥ 1.5 ML. The evolution of longitudinal coercivitiy vs Fe coverage is shown in Fig. 4-10 (c). The HC of the longitudinal hysteresis loop is very low in this system, it reaches a saturated value of 32 Oe when the thickness of Fe is higher than 2.0 ML.. 43.

(42) Figure 4-10: (a) The polar and longitudinal Kerr hysteresis loops at different coverage of Fe (dFe) thin films on Pt(111) at room temperature; (b) The evolutions of the polar and longitudinal Kerr intensities for Fe/Pt(111) correspond to different coverage of Fe. (c) The evolution of longitudinal coercivity versus Fe coverage. 44.

(43) Figure 4-11: The polar and longitudinal Kerr hysteresis loops of 1 ML Fe/Pt(111) at the different low temperature. The maximum of applied magnetic field is 900 Oe.. Our results have some differences with the Nahm’s report [73]. They did not observe any hysteresis loops when Fe thickness is less than 4.1 ML. At dFe = 4.1 ML, only polar hysteresis loop with a coercivity of 50 Oe can be observed. When dFe > 4.1 ML, the easy axis of the magnetization is in the in-plane direction. The in-plane coercivity is only about 9 Oe for dFe = 5.1 ML. The reason why these differences occurred may be due to the different methods of depositing Fe films. The method they used is e-beam heating while the one we used is thermal evaporation. Our deposition rate is lower than theirs, thus our Fe atoms have enough time to diffuse on the surface to form an ordered arrangement. Indeed, the satellites of LEED pattern surrounding each integer spots were observed when dFe > 1.0 ML. This is the evidence that Fe atoms are in an ordered state on Pt(111) surface. D. Repetto et al. have reported the similar results to our but they detect a spin reorientation transition (SRT) from in-plane to out-of-plane when dFe is reduced below a critical coverage of dcrit ≈ 2.8 ML. Such a SRT is found for many thin-film/substrate systems and is commonly ascribed to the dominant role of the interface anisotropy [74]. Experimentally, Fe films deposited on non-magnetic substrate is a well-studied system which shows a complex correlation between magnetic and structural properties depending on deposition temperature and preparation procedure [74-78]. Now we study 45.

(44) the magnetic property of Fe/Pt(111) which Fe films are grown at 180 K (low temperature; LT). Liquid nitrogen is used to cool the sample continuously and control the temperature at 180 K in the deposition process and the MOKE measurements are taken at LT. The evolutions of both polar and longitudinal Kerr hysteresis loops versus the thickness of Fe films growth on Pt(111) at LT are shown in Fig. 4-12 (a). For Fe coverage below dFe= 0.75 ML no Kerr signals can be obtained in the MOKE experiments at 180 K. For a film growing from monolayer islands this value is typically 0.65 ML [79]. This result is consistent with the assumption that hysteresis appears when the film has passed the percolation threshold. In the LT growth, part of the deposited Fe is present in the second layer therefore the onset of hysteresis to higher coverages. It will be studied by means of LEED in the following section. Only polar Kerr hysteresis loop is detected when the thickness of Fe is 0.75 ML. The results show that the LT-grown Fe films are ferromagnetic for thicknesses greater than 0.75 ML. The longitudinal Kerr hysteresis loop is irregular with increasing Fe thickness up to 1.0 ML. For Fe thicknesses ≥1.3 ML, only longitudinal Kerr hysteresis loop can be detected which the ultrathin Fe films on Pt(111) are present in-plane easy axis. Clearly visible is a spin reorientation transition (SRT) from out-of-plane to in-plane when dFe is increased to a critical coverage of dcrit = 1.3 ML. Such a SRT is found for many thin-film/substrate systems and is commonly ascribed to the dominant role of the interface anisotropy [80-82]. Fig. 4-12 (b) shows the remanence Kerr intensity of Fe films on Pt(111) as a function of thickness for LT growth, giving a result consistent with the finding in Refs. [74, 76]. Fig. 4-12 (c) shows HC of Fe films on Pt(111) as a function of dFe for LT growth. The saturation of longitudinal HC = 70 Oe at LT is larger than 35 Oe at room temperature. HC decreases as the temperature increases. This result is in agreement with that the relation between the coercive force and the temperature of a thin film can be described by the mathematical formula: H C = H C 0 (1 − T. TB. (4-4). ). where HC0 is the coercive force at T = 0 K, and TB is the blocking temperature [83].. 46.

(45) Figure 4-12: (a) The polar and longitudinal Kerr hysteresis loops at different coverage of Fe thin films on Pt(111) at 180 K; (b) The evolutions of the polar and longitudinal Kerr intensities for Fe/Pt(111) correspond to different coverage of Fe (dFe) at LT. (c) The evolution of the polar and longitudinal coercivity versus Fe coverage at LT. Fig. 4-13 shows LEED I-V spectrum of (0,0) beam intensity versus energy curves for different thickness Fe films deposited on Pt(111) at 180 K. We have calculated the interlayer distance as a function of the film thickness from Eq. 4-3. The results are shown in Fig. 4-14. For LT-deposited films, the LEED I-V results show a linear increase of the interlayer distance, starting from about 1 ML Fe, and a value 15% larger than the substrate interplanar distance is observed for a 2 ML thick Fe film. The interlayer distance shows a pronounced behavior for the Fe films grown at room temperature. We have detected the LEED I-V spectrum of (0,0) beam intensity for dFe > 2 ML, but the. 47.

(46) relative variation of the (0,0) beam intensity versus energy is desultory. The results indicate that Fe atoms deposited on Pt(111) at 180 K is 3-D islands which have irregular, ramified shape and grow in height with increasing Fe coverage. Compared with RT growth, a high island density in Fe films on Pt(111) substrates is observed, which is caused by the reduced mobility of Fe adatoms at low temperature. This structural evolution resembles the evolution of the lateral spacing with increasing Fe thickness, shows in Fig. 4-14, and it is evidence of a structural transformation within the Fe films during the film growth for a low temperature. The MOKE and LEED I-V results, we suggest that the structure of the Fe film changes from fcc-like to a bcc-like starting from 1 ML for low temperature growth. The similar structural transformation starts from 2 ML for the Fe films grown on Pt(111) at room temperature. The phenomenon also observes for the growth of Fe on Cu(100) at low temperature [84].. Figure 4-13: LEED I-V spectrum of the (0,0) beam curves for the clean surface of the Pt(111) substrate, and for different thickness Fe films on Pt(111) deposited at 180 K, respectively.. 48.

(47) Figure 4-14: The Vertical interlayer distances versus thickness of Fe films deposited on Pt(111) at 180 K. The value of horizontal dashed line is 2.27 Å for the interlayer distance of the Pt(111). Compared with RT growth, a high island density in Fe films on Pt(111) substrates is observed, which is caused by the reduced mobility of Fe adatoms at low temperature. The polar magnetization shows a strong temperature dependence in the investigated temperature range. In RT grown Fe/Pt(111) gives the magnetoelastic anisotropy due to the larger strain. It is a contribution to the in-plane anisotropy. It is often assumed that a different. surface. roughness. can. modify. the. surface. magnetoelastic. and. magnetocrystalline anisotropy energy. A judge of the roughness-induced contribution to the perpendicular anisotropy can be obtained considering theoretical models by Bruno [85]. For LT grown Fe films deposited on flat Pt(111) surface have much rougher films surface than RT grown films then the easy axis is in the out-of-plane direction. The 1 ML Fe/Pt(111) sample has been grown at substrate temperature of 193 K and slowly warms up at different annealing temperature for 15 min. To measure the polar and longitudinal Kerr hysteresis loops at each annealing temperature. Only polar Kerr signal can be observed during the annealing process. Fig. 4-15 shows the evolution of polar and longitudinal Kerr hysteresis loops versus sample temperature for 1 ML Fe/Pt(111) grown at 193 K. The Kerr intensity and coercivity in out-of-plane gradually decrease to zero. 49.

(48) from 193 K to 300 K. The result can confirm that a different surface roughness may modify the surface magnetocrystalline and magnetostatic anisotropy.. Figure 4-15: The evolution of polar and longitudinal Kerr hysteresis loops versus sample temperature (TS) for 1 ML Fe/Pt(111) grown at 193 K. (b)Temperature dependence of Kerr intensity and coercivity of 1 ML Fe/Pt(111) at 193 K.. 50.

(49) 4-1.4. Annealing effects on surface structure and magnetic properties of ultrathin films 1.2 ML Fe on Pt(111). We use AES to study the changes of the surface compositions during the annealing processes of 1.2 ML Fe/Pt(111). AES is a useful technique to study the diffusing dynamic during the heating process. The heating rate is about 15 ºC per minute during the heating processes. Each temperature of measuring AES data is held for 20 minutes in the annealing process. The annealing for 20 minutes is enough for the sample to reach the thermal equilibrium. We investigate the alloy formation of Fe-Pt by Auger signals of Pt NNV 237 eV and Fe LMM 703 eV. The inelastic mean free paths of Pt 237 eV and Fe 703 eV Auger signals are 5.35 Å and 12.53 Å, respectively [86]. Fig. 4-16 shows the evolution of Fe 703 eV and Pt 237 eV AES intensities versus annealing temperature (Ta) for 1.2 ML Fe/Pt(111). From 300 K to 770 K, Fe 703 eV signal do not has significant changes until the temperature is higher than 770 K. But Pt 237 eV signal increases about 1.5 times when 520 K ≤ Ta ≤ 620 K then do not vary until Ta = 770 K. These results indicate that the mixing of Fe and Pt atoms occurs at 520 K ≤ Ta ≤ 620 K in the interface, without diffusing into the bulk of Pt. As the temperature increases, the Fe and Pt signals reach equilibrium when 620 K ≤ Ta ≤ 720 K. It indicates that Fe atoms diffuse into the substrate to form a Fe–Pt interface alloy starting at 620 K near the interface. The surface Fe concentration, is nearly constant and indicates the formation of an surface alloy that has a constant Fe composition over a temperature range of 620–720 K. Owing to the sensitivity of Auger electron; we did not observe significant changes for the Fe 703 eV Auger signals between 620 K and 720 K. The result is in agreement with observations made by Jerdev et al. [87]. They used low energy ion scattering (LEIS) to determine the surface Fe coverage that the Fe coverage in this temperature range was 0.5 ML. This structure can correlate with a γ2 phase in the bulk Pt–Fe phase diagram which has 55% Pt and 45% Fe [88]. Simultaneously, the surface structure and magnetic property of 1.2 ML Fe/Pt(111) will occurs significant variations as the formation of Fe-Pt alloy. The relation between the magnetic and structural properties of the Fe-Pt alloy will be studied as follow. The Pt signal begins to increases significantly when Ta > 770 K, and meantime Fe 237. 51.

(50) eV signal drops obviously. These changes of signals indicate that a great deal of Fe atoms diffuse deeply into the Pt substrate. As the annealing temperature exceeds 950 K, all most Fe atoms diffuse into the substrate and the Pt atoms occupied almost all surface sites.. Figure 4-16: Pt 237 eV and Fe 703 eV Auger intensity versus annealing temperature for 1.2 ML Fe/Pt(111). The scale of Fe is 2 times of Pt. The Auger-signal ratio, R= I(Fe703)/I(Pt237), versus the annealing temperature is shown in Fig 4-17. It is clearly to view that the thermal motion of Fe and Pt atoms can be partitioned into four sections by Ta= 520 K, 620 K, 770 K and 830 K, respectively. The result agrees with the analysis of Fig. 4-16. This change just corresponds to our LEED measurement as shown with the insert patterns in Fig. 4-17. Each pattern is recorded at room temperature after annealing at the high temperature for 20 minutes. The kinetic energy of the incident electron beam is 64 eV. For 300 K ≤ Ta < 520 K, the LEED pattern shows a blurred six-fold symmetry (1×1) main spots. This indicates that the surface structure is a major fcc-like phase, which is coherent with the Pt substrate. Because the lattice mismatch (8.4%) is considerable between Fe and Pt that the surface of incommensurate structure is not flat but becomes corrugated. When Ta is between 520 K and 720 K, the blurred six-fold symmetry (1×1) main spots becomes sharp and bright that. 52.

(51) it is same as the LEED pattern of the well prepared Pt(111) surface. From the result of Fig 4-16, Fe atoms have been mixing and diffusing into Pt to form a Fe–Pt alloy near the interface at this temperature. The sharp (1×1) pattern is due to the ordered fcc FePt surface alloy formed on the surface layer. At temperatures higher than 770 K, the R decreases linearly with Ta for a great deal Fe atoms deep diffusion into the Pt bulk. The LEED patterns have a clear (2×2) diffracted spots as shown in Fig. 4-17. The (2×2) LEED pattern is also reported for the Pt80Fe20(111) bulk alloy surface and explained by the existence of two kinds of Pt atoms in the topmost layer [89]. The changes of the surface compositions in the surface layer are agreed with the AES signals vary in Fig. 4-16 at 800 K < Ta ≤ 870 K.. Figure 4-17: The evolutions of the ratio of AES intensity R=I(Fe703)/I(Pt237) versus annealing temperature (Ta) for 1.2 ML Fe/Pt(111) system. All the LEED patterns are taken at room temperature after annealing at each temperature for 20 minutes. The LEED patterns are measured at E = 64 eV and two annealing temperature range are also shown.. 53.