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Structural origin of perpendicular magnetic anisotropy in Ni–W thin films
Journal of Magnetism and Magnetic Materials 231 (2001) 85–93
Structural origin of perpendicular magnetic anisotropy in Ni–W thin films Nicolae Sulitanu* Department of Solid State Physics, Faculty of Physics, Al.I. Cuza University, Iasi 6600, Romania Received 12 October 2000; received in revised form 28 November 2000
Abstract The magnetic properties and microstructure of electrodeposited Ni–W thin films (0–11.7 at% W in composition) were studied. The film structures were divided into three regions: an FCC nanocrystalline phase (0–2 at% W), a transition region from FCC nanocrystalline to amorphous phase (2–7 at% W), and an amorphous phase (>7 at% W). In the transition region, (4–5 at% W) films with perpendicular magnetic anisotropy (PMA) were found. The saturation magnetization, magnetic anisotropy field, perpendicular magnetic anisotropy and perpendicular coercivity for a typical Ni–W film (4.5 at% W) were 420 kA/m, 451 kA/m, 230 kJ/m and 113 kA/m, respectively. The microstructure of Ni–W films with PMA is composed of isolated columnar crystalline grains (27–36 nm) with the FCC phase surrounded by the Ni–W amorphous phase. The appearance of the interface between the magnetic core of Ni crystalline grains and the Ni–W non-magnetic boundary layer seems to be the driving mechanism for the appearance of PMA. The origin of PMA in Ni–W films is mainly attributed to the magnetoelastic anisotropy associated with in-plane internal stress and positive magnetostriction. The secondary source of PMA is believed to be the magnetocrystalline anisotropy of h1 1 1i columnar grains and its shape magnetic anisotropy. It is concluded that Ni–W electrodeposited films (4–5 at% W) may be applicable for perpendicular magnetic recording media. # 2001 Elsevier Science B.V. All rights reserved. PACS: 61–68; 71–76 Keywords: Microstructure; Magnetic anisotropy; Ferromagnetic films
1. Introduction The demand for high-density recording systems presents a challenge to the physicist, materials scientist and for structural analysis techniques. In the last few decades, the progress of magnetic disc or tape data storage towards higher areal densities has been facilitated by the development of *Tel.: +40-32-20-1173; fax: +40-32-20-1150. E-mail address: [email protected] (N. Sulitanu).
magnetoresistive thin-film recording heads. The present technology uses thin film media which have in-plane magnetization and record information in the longitudinal mode [1–4], but in the future, perpendicular recording may be able to follow the demand for an increased recording density [1,3,5]. Generally, a perpendicular recording medium requires a magnetically uniaxial film in which a strong, perpendicular anisotropy K?; outweighs the demagnetization energy Kd ¼ m0 Ms2 =2 (where Ms is saturation magnetization and m0 is
0304-8853/01/$ - see front matter # 2001 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 0 1 ) 0 0 0 4 1 - 5
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vacuum magnetic permeability), namely K? > Kd and supports magnetization normal to the film. Several types of material in film form have been investigated, including particulate barium ferrite, sputter-deposited crystalline Co-transition metal alloys, amorphous rare earth-transition metal alloy films, Nd–Fe–B alloys, etc [1,3,5]. In this study, we report on the perpendicular magnetic anisotropy (PMA) in electrodeposited Ni–W thin films. These films have in-plane as well as perpendicular anisotropy depending on their composition, thickness and conditions of preparation [6]. The transition between in-plane and outof-plane magnetization can be studied with these films and offers the possibility of understanding the ‘‘real’’ origin of PMA. Assuming a homogeneously film magnetized to saturation with an applied field H in the normal plane, the total energy per unit sample volume is given by the following equation:
the applied field [8]. In accordance with the above, quantitative characterization of the film properties can be obtained from the quality factor Q¼
K? : Kd
ð4Þ
Here, j and y denote the angles subtended by the saturation magnetization and field, respectively, and the film normal. Ku ¼ K? m0 Ms2 =2 signifies the effective (apparent) uniaxial PMA constant including the shape (demagnetizing) energy influence. Therefore, the intrinsic (correct) PMA constant can be written as
This quality factor, Q, is no longer an absolute measure of whether the magnetization in the remanent state lies in-plane or perpendicular to it. Therefore, perpendicular magnetization in a film can occur for Q > 1, that is when Hk > Ms . The PMA constant, K? , contains all first-order (intrinsic) anisotropy contributions. The shape (stray field) anisotropy is a result of the longrange magnetic dipolar interaction that senses the outer boundaries of the sample and does not depend on the film thickness. It is of particular importance in thin films, and favors an in-plane preferential orientation for the magnetization usually observed. The stray field energy density, Kd , must be taken into account to decide when the easy magnetization direction is out-of-plane. For a pure Ni film electrodeposited on polycrystalline copper substrate, Ms ¼ 4:83 105 A m1 and Kd ¼ m0 Ms2 =2 ¼ 1:465 105 A m1. From a phenomenological point of view, the expression of K? can be separated in two major contributions, that of the magnetocrystalline anisotropy, Kmc and that of the magnetoelastic anisotropy, Kme :
K? ¼ Ku 12m0 Ms2 :
K? ¼ Kmc þ Kme :
E ¼ Ku cos2 y m0 Ms H cosðj yÞ:
ð1Þ
ð2Þ
Among the most accurate ways of determining the effective PMA, Ku , is the torque magnetometer method in which the magnetization of the film is always saturated along the direction of the applied field [7]. Referring to Eq. (2) one might define an intrinsic PMA field Hk? ¼
2K? ¼ Hk M s : m 0 Ms
ð3Þ
Here, the first term signifies the effective PMA field and the second term signifies the anisotropy demagnetizing field (Hd ¼ N ? Ms ¼ Ms ). The determination of the anisotropy field, Hk , is very often done by measuring the in-plane hysteresis loops (in-plane is the hard axis of PMA) but it can be more accurately determined by analysis of the rotational hysteresis energy loss as a function of
ð5Þ
The microscopic origin of the magnetocrystalline anisotropy is the spin–orbit interaction which, in turn, is influenced by the type of crystalline lattice. Therefore, the magnetocrystalline energy depends on the orientation of the magnetization vector relative to the crystalline axes. The lowered symmetry of the polycrystalline films strongly modifies the magnetocrystalline anisotropy as compared to the bulk. In principle, the exchange and the dipolar interactions could also contribute to Kme . The exchange interaction, however, cannot give rise to anisotropy since it is proportional to the scalar product of the spin vectors and is therefore independent of the angle between the spins and the crystalline axes. Moreover, for cubic crystalline materials, it can be shown from symmetry arguments that the sum of the dipole–
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dipole energies cancels [9, p. 249]. Consequently, the spin–orbit interaction will be primarily responsible for the magnetocrystalline anisotropy in FCC Ni. For pure Ni films (100–500 nm) electrodeposited on a polycrystalline copper substrate the mean value of the Kmc is equal to K1 ¼ 5:4 103 J m3. Strain in a ferromagnetic film changes the magnetocrystalline anisotropy and may thereby alter the direction of magnetization. Recently, Sanders et al. [10] established that the magnetoelastic coupling in strained thin films give rise to out-of-plane magnetic anisotropy. When a thin film is subject to in-plain strain an out-of-plane strain frequently appears which is often assumed to be given by the continuum elasticity as a Poisson-type reaction of the in-plane strain. The magnetoelastic energy contribution to the magnetic anisotropy of an elastically isotropic medium with isotropic magnetostriction, is expressed by the magnetoelastic anisotropy constant 3 Kme ¼ ls: 2
ð6Þ
Here, s is the stress that is related to the strain, e, via the elastic modulus, E, by the following equation: s ¼ Ee. The magnetostriction constant l depends on the orientation and can be positive or negative. For pure Ni films l1 1 1 ¼ 2:4 105 and l1 0 0 ¼ 4:6 105 . The magnetoelastic coupling coefficients in thin films can be two to three order of magnitude larger than the magnetocrystalline anisotropy constants [10,11] so that even small strains in the sub-percent range are capable of considerably modifying the magnetic anisotropy. In epitaxial Ni thin films (tens of nanometers in thickness), a strain of about 102 would be sufficient to explain the K? values of (3– 5.4) 104 J m3. In this article, we shall try to establish the contribution to the Ni–W film PMA of the two types of the magnetic anisotropy namely magnetocrystalline and magnetoelastic. Usually, high-density recording media with low noise characteristics are thin films composed of small crystalline grains and each grain is more or less magnetically isolated [2,5]. This means that not only structural but also magnetic properties are important. Therefore, in developing ideal recording media, it is important to clarify the
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relationships between structural and magnetic properties. Although many studies on the relationship between structural properties and PMA have been reported, the correlation between them does not seem to be sufficiently clear [1,5,12–15]. Up to now, a number of different sources of the structural anisotropy underlying the magnetic anisotropy have been postulated, ranging from stress, to a columnar microstructure, to a chemical or topological short-range order which becomes oriented perpendicular to the film due to some generally unspecified deposition process [5]. However, none of the above mechanisms have provided an entirely satisfactory explanation for the special properties of the Se films. The present study is focused on the correlation between microstructural and magnetic properties of electrodeposited Ni– W thin films and emphasises that the presence of the 5d-transition element, W, is crucial to the microstructure displayed by this kind of material. On the other hand, the W atoms present in the Ni matrix induces a grained structure because W atoms tend to segregate to the grain boundaries, W being known to have a high surface free energy [16]. Tungsten impurity atoms dissolved in a ferromagnetic Ni crystalline matrix change the spin–orbit coupling in the Ni–W magnetic system since the spin–orbit coupling is quite pronounced for these heavy atoms. On the other hand, the Ni ferromagnetic host atoms can induce by hybridisation a sizeable magnetic moment on the W impurity atoms, which are normally non-magnetic [17]. Therefore, the intrinsic origin of PMA in electrodeposited Ni–W thin films is determined by the changes in the electronic and local atomic structures induced by the tailored specific microstructures depending on the W content.
2. Experiment The Ni–W films (320 nm in thickness) were prepared by electrodeposition from aqueous plating baths containing sodium tungstate [18]. The films were deposited on copper substrates in discoid form. Deposits of different composition (0–11.7 at% W) were obtained by varying the sodium tungstate content of the fresh plating
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solution. The composition of Ni–W films was carried out by electron probe microanalysis (EPMA). The structural state of samples was defined by X-ray diffraction (XRD) in MoKa radiation. The grain morphology and grain size were directly determined from scanning electron microscopy (SEM) micrographs and from the broadening of XRD peaks using the Scherrer peaks broadening formula [19]. The magnetization and effective perpendicular anisotropy constant of the samples were determined with an automated torque-magnetometer (ATQM) in a maximum external field of 480 kA/m. Magnetization curves (hysteresis loops) were measured using a vibrating sample magnetometer (VSM). A magnetic field up to 500 kA/m was applied parallel or perpendicular to the film surface to measure both in-plane and perpendicular magnetization curves.
Fig. 1. Tungsten content dependence of the saturation magnetization Ms and of the anisotropy field Hk .
3. Results and discussion Fig. 1 presents the values of Ms and Hk for Ni– W films 320 nm thick depending on the W content. The saturation magnetization Ms slightly and monotonically decreases until 460 kA/m when the W content increases in the range below 2 at%. Ms begins to decrease more rapidly and linearly at W contents higher than 2 at% and is 175 kA/m for 11.7 at% W (films with higher W content are under investigation). Such a rapid decrease of Ms cannot be explained by simple dilution of the Ni matrix. The fast decrease suggests that W co-deposition enhances the formation of the less crystallised phase mainly composed of Ni and W. The anisotropy field, Hk; increases with W content increasing and reaches a maximum value of 455 kA/m for 4.5 at% W. In the whole region of W contents between 4–5 at%, one has Hk > Ms and, therefore, these films clearly exhibit PMA. Fig. 2 shows XRD diagrams of the films with various W contents (0, 2, 4.5 and 8 at%). The XRD diagrams clarified that the Ni–W electrodeposited films with W content below 2 at% are composed of FCC Ni crystalline grains whose [1 1 1] axes are preferentially oriented in the
Fig. 2. X-ray diffraction diagrams of films with thickness D of 320 nm and various W contents.
direction perpendicular to the film plane. Although the films with higher W content of 2 at% still present diffraction peaks corresponding to the (1 1 1) planes of FCC Ni grains, they have an amorphous phase volume fraction because the (1 1 1) plane intensity rapidly decreases with an increase of W content above 2 at%. This finding is in accordance with SEM investigations. Fig. 3 shows SEM micrographs for films with two different W contents: 4.5 at% W (Fig. 3a) and 6.5 at% W (Fig. 3b). The film structures were divided into three regions: FCC nanocrystalline phase (0–2 at% W), transition region from FCC nanocrystalline to amorphous phase (2–7 at% W) and amorphous phase (>7 at% W). The crystalline Ni–W films (52 at% W) consists of FCC
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Fig. 3. Scanning electron micrographs of Ni–W films: (a) 2 at% W and (b) 4.5 at% W.
crystalline grains (45–54 nm) and, therefore, we should properly call them nanocrystalline films. The Ni–W films with 2–7 at% W consisted of isolated crystalline grains (12–45 nm) with FCC phase surrounded by an amorphous phase and, therefore, these films will be called nanocrystalline–amorphous films. XRD data for films with more than 7 at% W show only a single broad peak (1 1 1) and no sharp (crystalline) peaks; i.e. consistent with a ‘‘X-ray amorphous’’ structure. This means that the films have a higher degree of disorder, but are still crystalline at least on a nanoscale. Otherwise, SEM investigations show that films with a W content somewhat above this critical value still contain small volume fractions of precipitate with diameters less than 10 nm that cannot be detected by large angle XRD. As a general remark, it was observed from XRD diagrams that diffraction peak corresponding to Ni (1 1 1) planes shifted to the side of the lower diffraction angle 2y with an increase of W content
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above 2 at% (Fig. 2). This means that lattice spacing between the adjacent Ni (1 1 1) planes is expanded due to the large internal stress resulting from the W atom penetration into the Ni matrix. The W is not soluble in the Ni lattice because of the high difference in atom diameters, >18% [20]. The nanocrystalline – amorphous Ni–W films with a (1 1 1) highly favored crystallographic orientation indicate the presence of a strong texture in the films, together with a strong magnetocrystalline anisotropy which is associated with a strong [1 1 1] uniaxial magnetic anisotropy (Fig. 1). Since the films have a grainy structure, there is a close correlation between the crystalline texture and the columnar grain structure [21]. Otherwise, a columnar grain structure was observed in the course of SEM investigation on the Ni–W films crosssection. Actually, this finding is in accordance with Fig. 3a, which shows a plan view W map of the Ni–4.5 at% W film. Tungsten enriched regions can be observed with brighter contrast in this figure. Almost no contrast was observed for the sample with lower W content than 2 at%. Fig. 3a shows that W clear segregated to the Ni grain boundaries and this effect is important for W which has high surface free energy [16,22] (bright contrast around of small region, 36 nm, with weak contrast). The width of the W enriched region was about 7 nm. It seems certain that Surely the crystal lattice in the boundary regions is fairly disordered and consists of very fine grains and/or partially amorphous phase. If so, the W enriched region can effectively isolate grains, and consequently can reduce exchange coupling between neighboring grains. Consequently, the film coercivity increases. Fig. 4 shows the in-plane Hck and, respectively, perpendicular Hc? coercivity as function of W content. In the composition region below 4.5 at% W, Hck increases up to about 48 kA/m. However, Hck is very small (maximum 48 kA/m) comparatively with Hc? , which also increases with an increase of W content and has a sharp peak (113 kA/m) at about 4.5 at% W. The small values for Hck and the similar variations of Hck and Hc? indicate that the magnetization is preponderantly oriented out-of-the-film plane. Beside the isolation of the Ni grains, the coercivity is influenced by the in-plane strain through a magnetostriction
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Fig. 5. Typical M2H curves (in-plane, k and perpendicular, ?) for a typical Ni–W film (4.5 at% W) with perpendicular magnetic anisotropy .
Fig. 4. In-plane Hck and, respectively, perpendicular Hc? coercivity versus W content for Ni–W films.
mechanism [23]. The qualitative confirmation that the stress really acts on Ni–W films was obtained from the large differences in the character of the in-plane hysteresis loops exhibited by samples with different W content [24]. Samples of pure Ni have tensile stress and exhibit almost square hysteresis loops with low coercivity of approximately 6 kA/m whereas the Ni–W films having high tensile stress show severely sheared rhombic hysteresis loops with high coercivity (15–48 kA/m), and these films have PMA. From Fig. 4 it is evident that the coercivity of the fine grained (12–45 nm) Ni–W films is much larger than the coercivity observed in coarse-grained (50–54 nm) pure Ni films. This would be one reason why Ni–W fine grained films retain a high magnetic anisotropy energy. Fig. 5 shows typical M2H hysteresis loops of 320 nm thick Ni–W film with 4.5 at% W measured along two orthogonal directions of the magnetization of the magnetic field relative to the sample plane. The area enclosed between the parallel and perpendicular loops gives the magnetic anisotropy energy. The in-plane hysteresis loop is a characteristic loop for films with stripe domain structure [25], and the perpendicular hysteresis loop is similar to that of films which exhibit PMA and magnetization reversal by magnetic domain wall displacement [26]. The film has a clear PMA, characterised by K? ¼ 230 kJ/m and Hc? ¼ 113 kA/m. These values of K? and Ms ¼
Fig. 6. Magnetic torque curve (clockwise curve) measured while rotating a 480 kA/m magnetic field in the normal plane to a Ni–W (4.5 at% W) film.
420 kA/m are almost equal to those of Co–Cr thin films [27]. Moreover, the Ni–W films exhibit a low anisotropy dispersion switching field distribution, S * ¼ 4 kA/m that gives rise to the right shoulder in the perpendicular M2H hysteresis loop. The squareness ratio of the in-plane M2H loop was low, S ¼ 0:6, and indicated a weak interaction field between grains [28]. These results indicate that there are regions with strong uniaxial anisotropy in the direction perpendicular to the film plane. Further, the high-field (480 kA/m) torque curves obtained by rotating the magnetic field round of the film plane normal axis were measured to determine the magnitude and direction of the magnetic anisotropy. Fig. 6 shows the torque curve measured in the plane normal to the Ni–4.5 at% W film plane. Since the magnetic easy and hard axes correspond to the direction where the magnetic torque
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becomes zero with negative and, respectively, positive slope, the film has PMA. Furthermore, a detailed examination of the torque curve reveals that the magnetic easy and hard axes are not exactly in respective normal and in-plane directions, and consequently the angle between these axes is not exactly 908. This shows that direction of the hard axis anisotropy is inside the film plane and that of the easy axis anisotropy is tilted from the film normal direction. Since the normal plane, in which the external magnetic field rotates in this experiment is arbitrary, no more details about the dependence of tilted angle on this film can be provided, although this tilt may depend on the deposition conditions. However, the tilting angle for PMA is less than 38. Fig. 7 shows the PMA constant, K? and demagnetizing anisotropy constant Kd , in relation to the W content. In the composition range between 4 and 5 at% W, K? > m0 Ms =2, and has the same variation with Hk (Fig. 1) and Hc? (Fig. 4). This finding is a new confirmation that these films exhibit PMA. We see that K? in Fig. 7 increases steeply when the W content is less than 4.5 at%, while K? decreases steadily as the W content increases above 4.5 at%. This suggests that the origin of magnetic aniso-
Fig. 7. Tungsten content dependence of the effective perpendicular anisotropy K? and the demagnetizing anisotropy Kd .
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tropy for each composition region is different and this behavior is reflecting different structural states of the Ni–W films. The nanocrystalline phase predominates in films with less than 4.5 at% W while the amorphous phase begin to predominate the nanocrystalline phase in films with more than 4.5 at% W. Moreover, films with PMA (4–5 at% W) consisted of isolated crystalline grains (27– 36 nm) with the FCC phase surrounded by the amorphous phase (Fig. 3a). If we suppose that Ni–W films with PMA, having thickness D, consist of cylindrical grains with total diameter d, the contribution of the shape anisotropy of columnar crystalline grains, each surrounded by a nonmagnetic W-rich segregated layer with width a at the grain boundary, can be expressed as [29] Kdg ¼ a
D d m0 Ms2 : dD 2
ð7Þ
The Ms reduction can be attributed to the ferromagnetism reduction in the grain boundary layer in consequence of not only local structural variation but also from an enhanced W solubility. We can assume that Ms0 for pure Ni film and Ms for Ni–W film are related through the following expression: aS Ms ¼ Ms0 1 : ð8Þ V Here, S and V represent the mean grain surface and, the mean grain volume respectively. For Ni–W films with PMA the width of nonmagnetic surface layer should be approximately 1–1.5 nm, and the Kdg contribution to PMA is negligible. This finding is in agreement with the theoretical estimate of Snyder and Kryder [30] for sputtered Co–Cr columnar films. Therefore, if shape columnar grain anisotropy is not responsible for PMA, it would have to be associated with a more complicated mechanism like stress along the grain boundary and magnetostriction, or perhaps ferromagnetic–antiferromagnetic coupling. It is more likely that PMA would be related to magnetocrystalline anisotropy, to over-all film stresses and magnetostriction or possibly related to pair ordering. Generally, it is known that a Ni single crystal has a cubic magnetocrystalline anisotropy
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(Kmc ¼ 1:41 103 J m3) which is averaged out in polycrystalline Ni with completely random grain orientation. However, nanocrystalline-amorphous Ni–W films with high crystallographic (1 1 1) texture have magnetocrystalline anisotropy and this is confirmed by the specific in-plane M–H hysteresis loop (rhombic character, Fig. 5). Ni–W films with PMA are perfectly (1 1 1) textured, the film normal is the easy axis and the maximum value of magnetocrystalline anisotropy field is 455 kA/m. Therefore, the magnetocrystalline anisotropy of the small crystalline grains has a welldefined alignment axis i.e. the magnetocrystalline anisotropy in Ni–W films (4–5 at% W) can be considered to be uniaxial in nature. Our results on the correlation between the magnetocrystalline anisotropy and structural anisotropy in Ni–W films (4–5 at% W) support the theoretical results of Hua et al. [31] and Garcia-Otero et al. [32] regarding equilibrium micromagnetic configurations in a cubic Ni particle system. However, it is believed that the observed PMA in Ni–W films is not caused mainly by the magnetocrystalline anisotropy of Ni small grain even if its value sensibly changes in Ni–W films. The W grain boundary segregation could cause strong elastic distortions of the Ni crystalline lattice near the grain boundary accompanied by an exchange coupling between neighboring grains. Therefore, the small Ms of Ni–W films (compared with pure Ni films) results in domination of the straininduced anisotropy over the magnetocrystalline and microshape anisotropies. For nanocrystallineamorphous Ni–W films (2–7 at% W), both Hc? = Hck and S? =Sk ratios increase with increasing W content (Figs. 4 and 5). On the other hand, the diffraction peak corresponding to Ni (1 1 1) planes is shifted to the side of the lower diffraction angle 2y with an increase of W content (Fig. 2). This finding can be understood as follows: during Ni and W alloying, the W atoms induce an in-plane tensile stress (s > 0) acting parallel to the film surface. This stress results in a stress-induced magnetic anisotropy. The magnitude of the Ni magnetostriction constant, l, decreases with dilution, reflecting in a decrease of Ms (Fig. 1). It is known that l of polycrystalline Ni film can change its sign from negative to
positive when the Ni lattice is doped in low concentration with a metallic non-magnetic impurity like, for example, 2.7 at% Cr [9, p. 284]. A similar behavior also should occur in Ni–W alloy films. Therefore, in Ni–W thin films the strains can both change the magnetocrystalline and magnetoelastic anisotropy (both favor the film normal as the easy axis). Recently, Sander et al. [10] were found that in isotropic strained cubic (1 1 1) Ni films both magnetocrystalline and magnetoelastic anisotropy contribute to the outof-plane anisotropy. These contributions are equal with DKmc ¼ K1 =12 ¼ 0:45 103 J m3 and, respectively, DKme ¼ 5:3 106 e, where e is the magnitude of the in-plane strain. The positive values indicate that the [1 1 1] direction is more favorable for magnetization orientation than another direction. The magnetoelastic coupling anisotropy contribution (DKme ) to PMA is two to three orders of magnitude larger than the magnetocrystalline anisotropy contribution. Therefore, even small isotropic in-plane strains (e), in the subpercent range, are capable of modifying the magnetic anisotropy considerable. For example, a strain of the order of 102 would be sufficient to explain the value of 5.3 104 J m3 of the PMA constant in Ni–W films. Therefore, the magnetoelastic anisotropy contribution to PMA is much more important than the magnetocrystalline anisotropy contribution. Otherwise between the magnetocrystalline anisotropy change and the induced stresses in the films direct correlation exists and, therefore it is difficult to estimate the exact amount of each contribution to the PMA. In short, the magnetoelastic energy associated with in-plane tensile stress (s > 0) and positive magnetostriction (l > 0) is believed to be the primary source of PMA in the Ni–W thin films. The secondary source is the magnetocrystalline anisotropy associated with out-of-plane [1 1 1] easy axes.
4. Conclusions The deviation of the magnetization and magnetic anisotropy in the Ni–W films (normal) from the corresponding direction in pure Ni films
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(in-plane) is attributed to the presence of tungsten and, possibly, its compositional inhomogeneity. The tungsten content dependence of the saturation magnetization, coercivity, anisotropy field and the anisotropy energy in Ni–W films is qualitatively explained by taking in account of the grained structure in the h1 1 1i columnar form surrounded by an amorphous phase, and of the compositional inhomogeneity. Tungsten atoms tend to segregate to the grain boundaries. Also, we consider that the appearance and development of the interface between magnetic core (Ni) and non-magnetic boundary layer (Ni–W) is the driving mechanism of PMA observed. PMA in Ni–W thin films principally arises from the magnetoelastic anisotropy associated with inplane internal stresses and positive magnetostriction. The tungsten content in the films can control the magnitude of the internal stresses. The secondary source of PMA is believed to be the magnetocrystalline anisotropy of h1 1 1i columnar grains and its shape magnetic anisotropy. There is a direct connection between the in-plane internal stress development and the grain columnar growth in the films. It seems that columnar growth is imposed by the necessity of the in-plane internal stress equilibration whatever its nature is. It is considered that the large K? of the Ni–W thin films with relatively large saturation magnetization Ms , may be applicable to the perpendicular magnetic recording media with ultrahigh density.
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Structural origin of perpendicular magnetic anisotropy in Ni–W thin films Nicolae Sulitanu* Department of Solid State Physics, Faculty of Physics, Al.I. Cuza University, Iasi 6600, Romania Received 12 October 2000; received in revised form 28 November 2000
Abstract The magnetic properties and microstructure of electrodeposited Ni–W thin films (0–11.7 at% W in composition) were studied. The film structures were divided into three regions: an FCC nanocrystalline phase (0–2 at% W), a transition region from FCC nanocrystalline to amorphous phase (2–7 at% W), and an amorphous phase (>7 at% W). In the transition region, (4–5 at% W) films with perpendicular magnetic anisotropy (PMA) were found. The saturation magnetization, magnetic anisotropy field, perpendicular magnetic anisotropy and perpendicular coercivity for a typical Ni–W film (4.5 at% W) were 420 kA/m, 451 kA/m, 230 kJ/m and 113 kA/m, respectively. The microstructure of Ni–W films with PMA is composed of isolated columnar crystalline grains (27–36 nm) with the FCC phase surrounded by the Ni–W amorphous phase. The appearance of the interface between the magnetic core of Ni crystalline grains and the Ni–W non-magnetic boundary layer seems to be the driving mechanism for the appearance of PMA. The origin of PMA in Ni–W films is mainly attributed to the magnetoelastic anisotropy associated with in-plane internal stress and positive magnetostriction. The secondary source of PMA is believed to be the magnetocrystalline anisotropy of h1 1 1i columnar grains and its shape magnetic anisotropy. It is concluded that Ni–W electrodeposited films (4–5 at% W) may be applicable for perpendicular magnetic recording media. # 2001 Elsevier Science B.V. All rights reserved. PACS: 61–68; 71–76 Keywords: Microstructure; Magnetic anisotropy; Ferromagnetic films
1. Introduction The demand for high-density recording systems presents a challenge to the physicist, materials scientist and for structural analysis techniques. In the last few decades, the progress of magnetic disc or tape data storage towards higher areal densities has been facilitated by the development of *Tel.: +40-32-20-1173; fax: +40-32-20-1150. E-mail address: [email protected] (N. Sulitanu).
magnetoresistive thin-film recording heads. The present technology uses thin film media which have in-plane magnetization and record information in the longitudinal mode [1–4], but in the future, perpendicular recording may be able to follow the demand for an increased recording density [1,3,5]. Generally, a perpendicular recording medium requires a magnetically uniaxial film in which a strong, perpendicular anisotropy K?; outweighs the demagnetization energy Kd ¼ m0 Ms2 =2 (where Ms is saturation magnetization and m0 is
0304-8853/01/$ - see front matter # 2001 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 0 1 ) 0 0 0 4 1 - 5
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vacuum magnetic permeability), namely K? > Kd and supports magnetization normal to the film. Several types of material in film form have been investigated, including particulate barium ferrite, sputter-deposited crystalline Co-transition metal alloys, amorphous rare earth-transition metal alloy films, Nd–Fe–B alloys, etc [1,3,5]. In this study, we report on the perpendicular magnetic anisotropy (PMA) in electrodeposited Ni–W thin films. These films have in-plane as well as perpendicular anisotropy depending on their composition, thickness and conditions of preparation [6]. The transition between in-plane and outof-plane magnetization can be studied with these films and offers the possibility of understanding the ‘‘real’’ origin of PMA. Assuming a homogeneously film magnetized to saturation with an applied field H in the normal plane, the total energy per unit sample volume is given by the following equation:
the applied field [8]. In accordance with the above, quantitative characterization of the film properties can be obtained from the quality factor Q¼
K? : Kd
ð4Þ
Here, j and y denote the angles subtended by the saturation magnetization and field, respectively, and the film normal. Ku ¼ K? m0 Ms2 =2 signifies the effective (apparent) uniaxial PMA constant including the shape (demagnetizing) energy influence. Therefore, the intrinsic (correct) PMA constant can be written as
This quality factor, Q, is no longer an absolute measure of whether the magnetization in the remanent state lies in-plane or perpendicular to it. Therefore, perpendicular magnetization in a film can occur for Q > 1, that is when Hk > Ms . The PMA constant, K? , contains all first-order (intrinsic) anisotropy contributions. The shape (stray field) anisotropy is a result of the longrange magnetic dipolar interaction that senses the outer boundaries of the sample and does not depend on the film thickness. It is of particular importance in thin films, and favors an in-plane preferential orientation for the magnetization usually observed. The stray field energy density, Kd , must be taken into account to decide when the easy magnetization direction is out-of-plane. For a pure Ni film electrodeposited on polycrystalline copper substrate, Ms ¼ 4:83 105 A m1 and Kd ¼ m0 Ms2 =2 ¼ 1:465 105 A m1. From a phenomenological point of view, the expression of K? can be separated in two major contributions, that of the magnetocrystalline anisotropy, Kmc and that of the magnetoelastic anisotropy, Kme :
K? ¼ Ku 12m0 Ms2 :
K? ¼ Kmc þ Kme :
E ¼ Ku cos2 y m0 Ms H cosðj yÞ:
ð1Þ
ð2Þ
Among the most accurate ways of determining the effective PMA, Ku , is the torque magnetometer method in which the magnetization of the film is always saturated along the direction of the applied field [7]. Referring to Eq. (2) one might define an intrinsic PMA field Hk? ¼
2K? ¼ Hk M s : m 0 Ms
ð3Þ
Here, the first term signifies the effective PMA field and the second term signifies the anisotropy demagnetizing field (Hd ¼ N ? Ms ¼ Ms ). The determination of the anisotropy field, Hk , is very often done by measuring the in-plane hysteresis loops (in-plane is the hard axis of PMA) but it can be more accurately determined by analysis of the rotational hysteresis energy loss as a function of
ð5Þ
The microscopic origin of the magnetocrystalline anisotropy is the spin–orbit interaction which, in turn, is influenced by the type of crystalline lattice. Therefore, the magnetocrystalline energy depends on the orientation of the magnetization vector relative to the crystalline axes. The lowered symmetry of the polycrystalline films strongly modifies the magnetocrystalline anisotropy as compared to the bulk. In principle, the exchange and the dipolar interactions could also contribute to Kme . The exchange interaction, however, cannot give rise to anisotropy since it is proportional to the scalar product of the spin vectors and is therefore independent of the angle between the spins and the crystalline axes. Moreover, for cubic crystalline materials, it can be shown from symmetry arguments that the sum of the dipole–
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dipole energies cancels [9, p. 249]. Consequently, the spin–orbit interaction will be primarily responsible for the magnetocrystalline anisotropy in FCC Ni. For pure Ni films (100–500 nm) electrodeposited on a polycrystalline copper substrate the mean value of the Kmc is equal to K1 ¼ 5:4 103 J m3. Strain in a ferromagnetic film changes the magnetocrystalline anisotropy and may thereby alter the direction of magnetization. Recently, Sanders et al. [10] established that the magnetoelastic coupling in strained thin films give rise to out-of-plane magnetic anisotropy. When a thin film is subject to in-plain strain an out-of-plane strain frequently appears which is often assumed to be given by the continuum elasticity as a Poisson-type reaction of the in-plane strain. The magnetoelastic energy contribution to the magnetic anisotropy of an elastically isotropic medium with isotropic magnetostriction, is expressed by the magnetoelastic anisotropy constant 3 Kme ¼ ls: 2
ð6Þ
Here, s is the stress that is related to the strain, e, via the elastic modulus, E, by the following equation: s ¼ Ee. The magnetostriction constant l depends on the orientation and can be positive or negative. For pure Ni films l1 1 1 ¼ 2:4 105 and l1 0 0 ¼ 4:6 105 . The magnetoelastic coupling coefficients in thin films can be two to three order of magnitude larger than the magnetocrystalline anisotropy constants [10,11] so that even small strains in the sub-percent range are capable of considerably modifying the magnetic anisotropy. In epitaxial Ni thin films (tens of nanometers in thickness), a strain of about 102 would be sufficient to explain the K? values of (3– 5.4) 104 J m3. In this article, we shall try to establish the contribution to the Ni–W film PMA of the two types of the magnetic anisotropy namely magnetocrystalline and magnetoelastic. Usually, high-density recording media with low noise characteristics are thin films composed of small crystalline grains and each grain is more or less magnetically isolated [2,5]. This means that not only structural but also magnetic properties are important. Therefore, in developing ideal recording media, it is important to clarify the
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relationships between structural and magnetic properties. Although many studies on the relationship between structural properties and PMA have been reported, the correlation between them does not seem to be sufficiently clear [1,5,12–15]. Up to now, a number of different sources of the structural anisotropy underlying the magnetic anisotropy have been postulated, ranging from stress, to a columnar microstructure, to a chemical or topological short-range order which becomes oriented perpendicular to the film due to some generally unspecified deposition process [5]. However, none of the above mechanisms have provided an entirely satisfactory explanation for the special properties of the Se films. The present study is focused on the correlation between microstructural and magnetic properties of electrodeposited Ni– W thin films and emphasises that the presence of the 5d-transition element, W, is crucial to the microstructure displayed by this kind of material. On the other hand, the W atoms present in the Ni matrix induces a grained structure because W atoms tend to segregate to the grain boundaries, W being known to have a high surface free energy [16]. Tungsten impurity atoms dissolved in a ferromagnetic Ni crystalline matrix change the spin–orbit coupling in the Ni–W magnetic system since the spin–orbit coupling is quite pronounced for these heavy atoms. On the other hand, the Ni ferromagnetic host atoms can induce by hybridisation a sizeable magnetic moment on the W impurity atoms, which are normally non-magnetic [17]. Therefore, the intrinsic origin of PMA in electrodeposited Ni–W thin films is determined by the changes in the electronic and local atomic structures induced by the tailored specific microstructures depending on the W content.
2. Experiment The Ni–W films (320 nm in thickness) were prepared by electrodeposition from aqueous plating baths containing sodium tungstate [18]. The films were deposited on copper substrates in discoid form. Deposits of different composition (0–11.7 at% W) were obtained by varying the sodium tungstate content of the fresh plating
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solution. The composition of Ni–W films was carried out by electron probe microanalysis (EPMA). The structural state of samples was defined by X-ray diffraction (XRD) in MoKa radiation. The grain morphology and grain size were directly determined from scanning electron microscopy (SEM) micrographs and from the broadening of XRD peaks using the Scherrer peaks broadening formula [19]. The magnetization and effective perpendicular anisotropy constant of the samples were determined with an automated torque-magnetometer (ATQM) in a maximum external field of 480 kA/m. Magnetization curves (hysteresis loops) were measured using a vibrating sample magnetometer (VSM). A magnetic field up to 500 kA/m was applied parallel or perpendicular to the film surface to measure both in-plane and perpendicular magnetization curves.
Fig. 1. Tungsten content dependence of the saturation magnetization Ms and of the anisotropy field Hk .
3. Results and discussion Fig. 1 presents the values of Ms and Hk for Ni– W films 320 nm thick depending on the W content. The saturation magnetization Ms slightly and monotonically decreases until 460 kA/m when the W content increases in the range below 2 at%. Ms begins to decrease more rapidly and linearly at W contents higher than 2 at% and is 175 kA/m for 11.7 at% W (films with higher W content are under investigation). Such a rapid decrease of Ms cannot be explained by simple dilution of the Ni matrix. The fast decrease suggests that W co-deposition enhances the formation of the less crystallised phase mainly composed of Ni and W. The anisotropy field, Hk; increases with W content increasing and reaches a maximum value of 455 kA/m for 4.5 at% W. In the whole region of W contents between 4–5 at%, one has Hk > Ms and, therefore, these films clearly exhibit PMA. Fig. 2 shows XRD diagrams of the films with various W contents (0, 2, 4.5 and 8 at%). The XRD diagrams clarified that the Ni–W electrodeposited films with W content below 2 at% are composed of FCC Ni crystalline grains whose [1 1 1] axes are preferentially oriented in the
Fig. 2. X-ray diffraction diagrams of films with thickness D of 320 nm and various W contents.
direction perpendicular to the film plane. Although the films with higher W content of 2 at% still present diffraction peaks corresponding to the (1 1 1) planes of FCC Ni grains, they have an amorphous phase volume fraction because the (1 1 1) plane intensity rapidly decreases with an increase of W content above 2 at%. This finding is in accordance with SEM investigations. Fig. 3 shows SEM micrographs for films with two different W contents: 4.5 at% W (Fig. 3a) and 6.5 at% W (Fig. 3b). The film structures were divided into three regions: FCC nanocrystalline phase (0–2 at% W), transition region from FCC nanocrystalline to amorphous phase (2–7 at% W) and amorphous phase (>7 at% W). The crystalline Ni–W films (52 at% W) consists of FCC
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Fig. 3. Scanning electron micrographs of Ni–W films: (a) 2 at% W and (b) 4.5 at% W.
crystalline grains (45–54 nm) and, therefore, we should properly call them nanocrystalline films. The Ni–W films with 2–7 at% W consisted of isolated crystalline grains (12–45 nm) with FCC phase surrounded by an amorphous phase and, therefore, these films will be called nanocrystalline–amorphous films. XRD data for films with more than 7 at% W show only a single broad peak (1 1 1) and no sharp (crystalline) peaks; i.e. consistent with a ‘‘X-ray amorphous’’ structure. This means that the films have a higher degree of disorder, but are still crystalline at least on a nanoscale. Otherwise, SEM investigations show that films with a W content somewhat above this critical value still contain small volume fractions of precipitate with diameters less than 10 nm that cannot be detected by large angle XRD. As a general remark, it was observed from XRD diagrams that diffraction peak corresponding to Ni (1 1 1) planes shifted to the side of the lower diffraction angle 2y with an increase of W content
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above 2 at% (Fig. 2). This means that lattice spacing between the adjacent Ni (1 1 1) planes is expanded due to the large internal stress resulting from the W atom penetration into the Ni matrix. The W is not soluble in the Ni lattice because of the high difference in atom diameters, >18% [20]. The nanocrystalline – amorphous Ni–W films with a (1 1 1) highly favored crystallographic orientation indicate the presence of a strong texture in the films, together with a strong magnetocrystalline anisotropy which is associated with a strong [1 1 1] uniaxial magnetic anisotropy (Fig. 1). Since the films have a grainy structure, there is a close correlation between the crystalline texture and the columnar grain structure [21]. Otherwise, a columnar grain structure was observed in the course of SEM investigation on the Ni–W films crosssection. Actually, this finding is in accordance with Fig. 3a, which shows a plan view W map of the Ni–4.5 at% W film. Tungsten enriched regions can be observed with brighter contrast in this figure. Almost no contrast was observed for the sample with lower W content than 2 at%. Fig. 3a shows that W clear segregated to the Ni grain boundaries and this effect is important for W which has high surface free energy [16,22] (bright contrast around of small region, 36 nm, with weak contrast). The width of the W enriched region was about 7 nm. It seems certain that Surely the crystal lattice in the boundary regions is fairly disordered and consists of very fine grains and/or partially amorphous phase. If so, the W enriched region can effectively isolate grains, and consequently can reduce exchange coupling between neighboring grains. Consequently, the film coercivity increases. Fig. 4 shows the in-plane Hck and, respectively, perpendicular Hc? coercivity as function of W content. In the composition region below 4.5 at% W, Hck increases up to about 48 kA/m. However, Hck is very small (maximum 48 kA/m) comparatively with Hc? , which also increases with an increase of W content and has a sharp peak (113 kA/m) at about 4.5 at% W. The small values for Hck and the similar variations of Hck and Hc? indicate that the magnetization is preponderantly oriented out-of-the-film plane. Beside the isolation of the Ni grains, the coercivity is influenced by the in-plane strain through a magnetostriction
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Fig. 5. Typical M2H curves (in-plane, k and perpendicular, ?) for a typical Ni–W film (4.5 at% W) with perpendicular magnetic anisotropy .
Fig. 4. In-plane Hck and, respectively, perpendicular Hc? coercivity versus W content for Ni–W films.
mechanism [23]. The qualitative confirmation that the stress really acts on Ni–W films was obtained from the large differences in the character of the in-plane hysteresis loops exhibited by samples with different W content [24]. Samples of pure Ni have tensile stress and exhibit almost square hysteresis loops with low coercivity of approximately 6 kA/m whereas the Ni–W films having high tensile stress show severely sheared rhombic hysteresis loops with high coercivity (15–48 kA/m), and these films have PMA. From Fig. 4 it is evident that the coercivity of the fine grained (12–45 nm) Ni–W films is much larger than the coercivity observed in coarse-grained (50–54 nm) pure Ni films. This would be one reason why Ni–W fine grained films retain a high magnetic anisotropy energy. Fig. 5 shows typical M2H hysteresis loops of 320 nm thick Ni–W film with 4.5 at% W measured along two orthogonal directions of the magnetization of the magnetic field relative to the sample plane. The area enclosed between the parallel and perpendicular loops gives the magnetic anisotropy energy. The in-plane hysteresis loop is a characteristic loop for films with stripe domain structure [25], and the perpendicular hysteresis loop is similar to that of films which exhibit PMA and magnetization reversal by magnetic domain wall displacement [26]. The film has a clear PMA, characterised by K? ¼ 230 kJ/m and Hc? ¼ 113 kA/m. These values of K? and Ms ¼
Fig. 6. Magnetic torque curve (clockwise curve) measured while rotating a 480 kA/m magnetic field in the normal plane to a Ni–W (4.5 at% W) film.
420 kA/m are almost equal to those of Co–Cr thin films [27]. Moreover, the Ni–W films exhibit a low anisotropy dispersion switching field distribution, S * ¼ 4 kA/m that gives rise to the right shoulder in the perpendicular M2H hysteresis loop. The squareness ratio of the in-plane M2H loop was low, S ¼ 0:6, and indicated a weak interaction field between grains [28]. These results indicate that there are regions with strong uniaxial anisotropy in the direction perpendicular to the film plane. Further, the high-field (480 kA/m) torque curves obtained by rotating the magnetic field round of the film plane normal axis were measured to determine the magnitude and direction of the magnetic anisotropy. Fig. 6 shows the torque curve measured in the plane normal to the Ni–4.5 at% W film plane. Since the magnetic easy and hard axes correspond to the direction where the magnetic torque
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becomes zero with negative and, respectively, positive slope, the film has PMA. Furthermore, a detailed examination of the torque curve reveals that the magnetic easy and hard axes are not exactly in respective normal and in-plane directions, and consequently the angle between these axes is not exactly 908. This shows that direction of the hard axis anisotropy is inside the film plane and that of the easy axis anisotropy is tilted from the film normal direction. Since the normal plane, in which the external magnetic field rotates in this experiment is arbitrary, no more details about the dependence of tilted angle on this film can be provided, although this tilt may depend on the deposition conditions. However, the tilting angle for PMA is less than 38. Fig. 7 shows the PMA constant, K? and demagnetizing anisotropy constant Kd , in relation to the W content. In the composition range between 4 and 5 at% W, K? > m0 Ms =2, and has the same variation with Hk (Fig. 1) and Hc? (Fig. 4). This finding is a new confirmation that these films exhibit PMA. We see that K? in Fig. 7 increases steeply when the W content is less than 4.5 at%, while K? decreases steadily as the W content increases above 4.5 at%. This suggests that the origin of magnetic aniso-
Fig. 7. Tungsten content dependence of the effective perpendicular anisotropy K? and the demagnetizing anisotropy Kd .
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tropy for each composition region is different and this behavior is reflecting different structural states of the Ni–W films. The nanocrystalline phase predominates in films with less than 4.5 at% W while the amorphous phase begin to predominate the nanocrystalline phase in films with more than 4.5 at% W. Moreover, films with PMA (4–5 at% W) consisted of isolated crystalline grains (27– 36 nm) with the FCC phase surrounded by the amorphous phase (Fig. 3a). If we suppose that Ni–W films with PMA, having thickness D, consist of cylindrical grains with total diameter d, the contribution of the shape anisotropy of columnar crystalline grains, each surrounded by a nonmagnetic W-rich segregated layer with width a at the grain boundary, can be expressed as [29] Kdg ¼ a
D d m0 Ms2 : dD 2
ð7Þ
The Ms reduction can be attributed to the ferromagnetism reduction in the grain boundary layer in consequence of not only local structural variation but also from an enhanced W solubility. We can assume that Ms0 for pure Ni film and Ms for Ni–W film are related through the following expression: aS Ms ¼ Ms0 1 : ð8Þ V Here, S and V represent the mean grain surface and, the mean grain volume respectively. For Ni–W films with PMA the width of nonmagnetic surface layer should be approximately 1–1.5 nm, and the Kdg contribution to PMA is negligible. This finding is in agreement with the theoretical estimate of Snyder and Kryder [30] for sputtered Co–Cr columnar films. Therefore, if shape columnar grain anisotropy is not responsible for PMA, it would have to be associated with a more complicated mechanism like stress along the grain boundary and magnetostriction, or perhaps ferromagnetic–antiferromagnetic coupling. It is more likely that PMA would be related to magnetocrystalline anisotropy, to over-all film stresses and magnetostriction or possibly related to pair ordering. Generally, it is known that a Ni single crystal has a cubic magnetocrystalline anisotropy
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(Kmc ¼ 1:41 103 J m3) which is averaged out in polycrystalline Ni with completely random grain orientation. However, nanocrystalline-amorphous Ni–W films with high crystallographic (1 1 1) texture have magnetocrystalline anisotropy and this is confirmed by the specific in-plane M–H hysteresis loop (rhombic character, Fig. 5). Ni–W films with PMA are perfectly (1 1 1) textured, the film normal is the easy axis and the maximum value of magnetocrystalline anisotropy field is 455 kA/m. Therefore, the magnetocrystalline anisotropy of the small crystalline grains has a welldefined alignment axis i.e. the magnetocrystalline anisotropy in Ni–W films (4–5 at% W) can be considered to be uniaxial in nature. Our results on the correlation between the magnetocrystalline anisotropy and structural anisotropy in Ni–W films (4–5 at% W) support the theoretical results of Hua et al. [31] and Garcia-Otero et al. [32] regarding equilibrium micromagnetic configurations in a cubic Ni particle system. However, it is believed that the observed PMA in Ni–W films is not caused mainly by the magnetocrystalline anisotropy of Ni small grain even if its value sensibly changes in Ni–W films. The W grain boundary segregation could cause strong elastic distortions of the Ni crystalline lattice near the grain boundary accompanied by an exchange coupling between neighboring grains. Therefore, the small Ms of Ni–W films (compared with pure Ni films) results in domination of the straininduced anisotropy over the magnetocrystalline and microshape anisotropies. For nanocrystallineamorphous Ni–W films (2–7 at% W), both Hc? = Hck and S? =Sk ratios increase with increasing W content (Figs. 4 and 5). On the other hand, the diffraction peak corresponding to Ni (1 1 1) planes is shifted to the side of the lower diffraction angle 2y with an increase of W content (Fig. 2). This finding can be understood as follows: during Ni and W alloying, the W atoms induce an in-plane tensile stress (s > 0) acting parallel to the film surface. This stress results in a stress-induced magnetic anisotropy. The magnitude of the Ni magnetostriction constant, l, decreases with dilution, reflecting in a decrease of Ms (Fig. 1). It is known that l of polycrystalline Ni film can change its sign from negative to
positive when the Ni lattice is doped in low concentration with a metallic non-magnetic impurity like, for example, 2.7 at% Cr [9, p. 284]. A similar behavior also should occur in Ni–W alloy films. Therefore, in Ni–W thin films the strains can both change the magnetocrystalline and magnetoelastic anisotropy (both favor the film normal as the easy axis). Recently, Sander et al. [10] were found that in isotropic strained cubic (1 1 1) Ni films both magnetocrystalline and magnetoelastic anisotropy contribute to the outof-plane anisotropy. These contributions are equal with DKmc ¼ K1 =12 ¼ 0:45 103 J m3 and, respectively, DKme ¼ 5:3 106 e, where e is the magnitude of the in-plane strain. The positive values indicate that the [1 1 1] direction is more favorable for magnetization orientation than another direction. The magnetoelastic coupling anisotropy contribution (DKme ) to PMA is two to three orders of magnitude larger than the magnetocrystalline anisotropy contribution. Therefore, even small isotropic in-plane strains (e), in the subpercent range, are capable of modifying the magnetic anisotropy considerable. For example, a strain of the order of 102 would be sufficient to explain the value of 5.3 104 J m3 of the PMA constant in Ni–W films. Therefore, the magnetoelastic anisotropy contribution to PMA is much more important than the magnetocrystalline anisotropy contribution. Otherwise between the magnetocrystalline anisotropy change and the induced stresses in the films direct correlation exists and, therefore it is difficult to estimate the exact amount of each contribution to the PMA. In short, the magnetoelastic energy associated with in-plane tensile stress (s > 0) and positive magnetostriction (l > 0) is believed to be the primary source of PMA in the Ni–W thin films. The secondary source is the magnetocrystalline anisotropy associated with out-of-plane [1 1 1] easy axes.
4. Conclusions The deviation of the magnetization and magnetic anisotropy in the Ni–W films (normal) from the corresponding direction in pure Ni films
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(in-plane) is attributed to the presence of tungsten and, possibly, its compositional inhomogeneity. The tungsten content dependence of the saturation magnetization, coercivity, anisotropy field and the anisotropy energy in Ni–W films is qualitatively explained by taking in account of the grained structure in the h1 1 1i columnar form surrounded by an amorphous phase, and of the compositional inhomogeneity. Tungsten atoms tend to segregate to the grain boundaries. Also, we consider that the appearance and development of the interface between magnetic core (Ni) and non-magnetic boundary layer (Ni–W) is the driving mechanism of PMA observed. PMA in Ni–W thin films principally arises from the magnetoelastic anisotropy associated with inplane internal stresses and positive magnetostriction. The tungsten content in the films can control the magnitude of the internal stresses. The secondary source of PMA is believed to be the magnetocrystalline anisotropy of h1 1 1i columnar grains and its shape magnetic anisotropy. There is a direct connection between the in-plane internal stress development and the grain columnar growth in the films. It seems that columnar growth is imposed by the necessity of the in-plane internal stress equilibration whatever its nature is. It is considered that the large K? of the Ni–W thin films with relatively large saturation magnetization Ms , may be applicable to the perpendicular magnetic recording media with ultrahigh density.
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