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Magnetic properties and magnetization reversal of CoSm‖Cr thin films
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ELSEVIER
Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
Magnetic properties and magnetization reversal of CoSm ICr thin films Z.S. Shan a,b, S.S. Malhotra a.b, S.H. Liou a,b, Yi Liu a,c, M. Yu a,b, D.J. Sellmyer
a,b, *
Center for Materials Research and Analysis, Uniz:ersit).'of Nebraska, Lincoln, NE 68588-0113, USA b Behlen Laborator3., of Physics, Unit'ersio~ of Nebraska, Lincoln, NE 68588-0113, USA c Department of Mechanical Engineering, Unit,ersi~, of Nebraska, Lincoln, NE 68588-0113, USA
Received 20 July 1995;revised l0 March 1996
Abstract In this paper the magnetic and structural properties of CoSm thin films with a Cr underlayer (CoSm I[Cr) are presented, with emphasis on the measurements of anisotropy at room and low temperature and magnetization reversal. The grain size of the Cr underlayer is about 250 A and the thin CoSm layer (e.g., 240 A) inherits this grain size. The CoSta layer consists of nanocrystallites, about 50 A in diameter, embedded in an amorphous matrix. The Ar pressure, CoSm layer-thickness, and temperature dependencies of magnetic properties including magnetization, coercivity and especially the anisotropy were investigated systematically. CoSta ILCr with coercivity up to 4.2 kOe at room temperature has been prepared. The intrinsic anisotropy is 4 x 10 6 and 1.4X 10 7 erg/cm 3 at room temperature for CoSm(240 *)llCr and CoSm(960 A)IICr, respectively, and both increase to 3.9 × 107 erg/cm 3 at 10 K. Magnetization reversal studies indicate that the coercivity mechanism changes from wall pinning for samples prepared at lower Ar pressure (5-12 mT) to single-particle coherent rotation for samples prepared at higher pressure (30 mT). The correlations between the microstructure and magnetic properties are discussed.
Keywords: Magnetization reversal; Coercivity mechanism; Magnetic switching volume; Microstructure; Magnetic anisotropy; CoSta film
1. Introduction Thin-film recording media have received continuous attention for high-density hard-disk recording and recent progress has ensured a dominant role of magnetic recording for the foreseeable future. Murdock et al. [1] have predicted that longitudinal recording of 10 G b / i n . 2 requires the media to have
* Corresponding author. Fax: + 1-402-4722879; email: [email protected].
high coercivity (2500-4500 Oe), small grain size ( 8 0 - 1 0 0 ,~), and the individual particles should be isolated from one another in terms of exchange interaction. Sharrock [2] has pointed out that the high anisotropy is required so the thermal stability requirement of KuV/kBT> 100, where K u, V, k B and T are the anisotropy, grain volume, Boltzmann constant and temperature respectively, is satisfied. With Co-based thin-film media, IBM and Hitachi have made hard disks with areal densities of 1 G b / i n . 2 [3] and 2 G b / i n . 2 [4] respectively. Besides the CoCrX (X = Ta, Pt, Ni, etc.) thin-film media, rare-earth
0304-8853/96/$15.00 Copyright © 1996 Elsevier Science B.V. All rights reserved. PII S0304- 8853(96)00296-X
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Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
( R E ) permanent-magnet films or multilayers, such as
2. F a b r i c a t i o n a n d c h a r a c t e r i z a t i o n
CoSm [5-7] and FeSmN [8] films, may be considered as candidates because of their large coercivity and anisotropy. Among RE semi-hard magnetic films, CoSm films prepared at room temperature have gained increasing attention because they may have both small grain size and high uniaxial anisotropy, which is essential for thin-film media of high-density recording. Velu and Lambeth [9-11] reported that CoSta films on Cr underlayers, denoted here as CoSm I[Cr, with coercivity up to 3000 Oe have been prepared by rf sputtering onto 7059 Coming glass substrates at r o o m t e m p e r a t u r e . The high coercivity was attributed to the very large anisotropy of CosSm and the in-plane lattice matching between the CosSm(1120) and Cr(ll0). In their papers the effects of sputtering conditions on the film coercivity have been systematically investigated and the microstructure of CoSta film and Cr underlayer was studied by means of AFM (atomic force microscopy) [11], Their studies show that the CoSmllCr thin films may be a promising medium for high-density longitudinal recording. Recently, Okumura et al. [12,13] reported that a maximum coercivity of 3.5 kOe was obtained for a Co85Sm~5[[Cr film and claimed that coercivity mainly results from both the net-like microstructure and fluctuation in the local anisotropy of the Co85 Sm ~5 layer. Liu et al. [14-16] reported systematic studies of the microstructure of CoSm II Cr including Cr underlayer, CoSm layer, the nanocrystallites in CoSm layer, etc. Singleton et al. [17] reported measurements of the switching volume which is 10- l 8_ 10 ~7 cm 3 and Sellmyer et al. [18] gave a succinct review of the magnetic and structural properties for CoSm IICr films. However, the magnetic properties of CoSm [JCr films have not been reported systematically. For example, a very important property, the anisotropy of CoSm J[Cr films, has not yet been studied intensively. In this paper we report our studies of the magnetic properties of films at room and low temperature, with emphasis on the anisotropy measurement and magnetization reversal analyses. The microstructure and the relationships between the magnetic and microstructure properties will also be discussed briefly.
The CoxSm [[Cr films were sputtered onto microscope cover glass substrates by a dc magnetron sputtering machine. The CoxSm targets were home made by pressing the CoxSm powder and then sintering in vacuum at 1100°C for 30 min. The commercial Cr target had a 99.9% purity. The base pressure of the sputtering system was 1-2 × 10 -7 Torr and the Ar pressure during sputtering varied from 5 to 30 mT to study the Ar-pressure dependence of the magnetic properties, which will be discussed in Section 4.4. The sputtering rate and power for CoSm were ~ 1 .~/s and 15 W and for Cr were ~ 2.5 .~/s and 36 W, respectively. The distance from the substrate to both guns was about 5 cm. Twelve samples were made in one vacuum run to avoid undesired changes of preparation conditions for one series of samples. Using CoxSm (x = 5, 4, 3.5) targets, the compositions of the sputtered films, which were determined using energy dispersive X-ray spectroscopy (EDX) in a scanning electron microscope, were Co5.5Sm, Co4.2Sm and Co3.75m, respectively. The largest coercivities measured were 1.5, 3.7 and 4.2 kOe for the above films. In this paper we mainly focus on studying the Co4.2SmllCr properties, and simply use CoSm to denote Co4.2Sm hereafter. The structural properties were examined by X-ray diffraction, AFM and transition electron microscopy (TEM). The magnetic properties were studied by an AGFM (alternating gradient force magnetometer) from 300 to 10 K and selected samples were measured by a SQUID magnetometer at applied fields up to 55 kOe. In this paper the structural properties and temperature-dependent magnetism were investigated mainly for the samples shown in Fig. 1, i.e. the sputtered Ar pressure which is identical for both Cr and CoSm layers changes from 5 to 30 mT (i.e. samples l, 2 and 3) and the CoSm layer-thickness varies from 60 to 960 A (i.e. samples 4, 2 and 5). Two sets of samples were prepared: one set with no Cr overcoat layer for AFM and TEM studies and the other set with a Cr overcoat layer of 200 ~. for magnetic property studies. Both sets of samples have a Cr underlayer of 1200 ~,. We have also prepared samples with different Ar pressures for the Cr underlayer
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3. Structural properties
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[dcosm:240~, Fig. 1. Samples used for structural and magnetic studies. From samples 1 to 2 to 3, the CoSm layer-thickness is fixed at 240 ,~, while the Ar sputtering pressure increases from 5 to 12 to 30 mTorr. From samples 4 to 2 to 5, the Ar sputtering pressure is fixed at 12 mTorr while the CoSm layer-thickness increases from 60 to 240 to 960 ,~. and the CoSm layer. The Ar pressure was appropriately varied to achieve the highest coercivity values. One example is shown in Fig. 8.
The X-ray diffraction patterns of the Cr underlayer on glass substrates show that, as the sputtered Ar pressure varies from 5 to 30 roT, the diffraction patterns only show a single (110) peak, indicating that the Cr underlayer is highly textured. This is a necessary condition for the subsequent CoSm layer to have large coercivity, as has been proved experimentally by Velu and Lambeth [9,10] and Okumura et al. [12,13]. This character was interpreted based on the lattice matching between CosSm(1120) and Cr(ll0). However, recent work has shown that a disordered hexagonal C o - S m alloy is present rather than the C % S m phase [14-16]. X-ray diffraction studies of the CoxSm ( x = 3.5, 4, 5) layer have been performed either for an individual CoxSm layer on a glass substrate or a CoxSm layer on a Cr underlayer which is on the glass
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Fig. 2. Atomic force micrographs for the samples shown in Fig. 1. (a,b,c) The Ar pressure dependence of the surface morphology;(d,b,e) the CoSm layer-thickness dependence of the surface morphology.
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substrate. No crystalline diffraction peaks are observed in both cases. Presumably the crystallite size is too small to show well defined peaks in these cases. To study the surface morphology of CoSta layers, a series of samples as shown in Fig. 1 but without a Cr overcoat layer was specially prepared as has been mentioned in Section 2. The surface morphology revealed by AFM is displayed in Fig. 2. We also prepared a Cr layer of 1200 ,~ on a glass substrate and it displayed nearly the same topographical structure as shown in Fig. 2. This confirms that the surface morphology in Fig. 2a-2e is inherited from the Cr underlayer morphology. One can find from Fig. 2 that the average grain size of the Cr underlayer is about 250 A. As the sputtered Ar pressure increases, i.e. from Fig. 2a to Fig. 2b to Fig. 2c, the surface roughness, which is the roughness superposition of the CoSm layer and Cr underlayer, increases. As the CoSta layer-thickness increases, i.e. from Fig. 2d to Fig. 2b to Fig. 2e, the surface roughness is nearly the same because they are fabricated in the same Ar pressure. However, Fig. 2e shows larger diameters of 'hills' since the 960 ~, thick CoSta layer may fill up some voids and make the 'hills' larger. The AFM micrographs only show the surface morphology of the CoSta layer and the microstructure inside the CoSm layer will be discussed in terms of the TEM results in the next paragraph. Systematic studies of bright-field and high-resolution TEM (HRTEM) for the same series of samples in Figs. 1 and 2 have been performed and the results are presented in earlier papers [14-16]. Here, a brief summary of the microstructure of the Cr underlayer and the CoSta layer is given: the Cr underlayer has a grain size of about 250 A and this grain image can be seen in samples 2, 3 and 4 in Figs. 1 and 2. But this Cr grain image cannot be seen in sample 5 because the CoSm layer is 960 A thick and the Cr grain image cannot be continued through the whole CoSta layer; the Cr grain image also cannot be seen in sample 1, possibly because the Ar sputtering pressure was only 5 mT in this case and the Cr underlayer possesses a smooth surface as shown in Fig. 2a, lacking the columnar structure. The CoSta grains of about 250 A in dimension for samples 2, 3 and 4 consist of CoSta nanocrystallites of typical dimension 50 A embedded in an amorphous CoSta
Grains (250~)
Amorphous matrix •
Crystallites (50/~) /
Fig. 3. Schematic diagram of a CoSm grain of ~ 250 ,~ dimension: the nanocrystallites of dimension ~ 50 ,~ are embedded in an amorphous matrix.
matrix, as shown in Fig. 3. The nanocrystallites have a disordered close-packed structure with varying stacking mode from crystallite to crystallite and sometimes within a single crystallite. The volume fraction of the nanocrystallites in the CoSm layer is shown in Table 1. It is seen that the volume fraction decreases from 91 to 65 to 54% as the Ar sputtering pressure is increased from 5 to 12 to 30 mT (i.e. from sample 1 to 2 to 3) respectively; as the CoSta layer-thickness increases from 60 to 240 to 960 ,~ (i.e. from sample 4 to 2 to 5), the volume fractions of the nanocrystallites remain almost unchanged. These properties are only mentioned here to provide the structural basis for understanding the magnetic properties with the detailed analysis in our other papers published earlier [ 14-16].
4. Magnetic properties at r o o m t e m p e r a t u r e
4.1. Hysteresis loops at room temperature
Both the parallel (i.e. H Jlfilm plane) and perpendicular (i.e. H .1_film plane) magnetization curves for the samples shown in Figs. 1 and 2 are demonstrated in Fig. 4. This figure gives the general features of the Ar pressure dependence (Fig. 4a-4c) and CoSta layer-thickness dependence (Fig. 4d,4e) of the magnetic properties: (i) In this series of samples, sample 2, which was prepared in 12 mT Ar pressure with a 240 ,~ CoSta layer and a 1200 A Cr underlayer, shows the largest coercivity of H c = 2.6 kOe and squareness S = 0.88. Here, Hc was measured from
Z.S. Shan et al./ Journal qf Magnetism and Magnetic Materials 161 (1996) 323-336 Table 1 Ar pressure and CoSm layer-thickness effects on the volume fraction (%) of the nanocrystallites, where Pat is the Ar pressure during sputtering and dsmco the thickness of the SmCo42 layer PAr
Volume ~actionVc(%)
(mT)
dsmco = 6 0 A
dsmco = 240A
dsmco = 960A
5 12 30
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and sample 3 in Fig. 4c); as the CoSm layer-thickness increases (i.e. from Fig. 4d to Fig. 4e), the perpendicular magnetization curves become narrower, which means that the magnetic moments are better aligned in the film plane for the thicker CoSm layer than for the thinner CoSm layer. (iii) The initial curves display a typical wall pinning or particle-rotation feature. The coercivity mechanism and magnetization reversal properties will be discussed in more detail in Section 6. 4.2. Layer-thickness dependence o f magnetic properties
the in-plane magnetization curve. (ii) When looking at the perpendicular magnetization curves one finds that, as the Ar pressure increases (i.e. from Fig. 4a to Fig. 4c), sample 2 in Fig. 4b shows the narrowest magnetization curve, which means that the magnetic moments are well aligned in the film plane compared with the other two samples (i.e. sample 1 in Fig. 4a
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CoSrn layer-thickness (~) Fig. 5. CoSm layer-thickness dependence of (a) coercivity (H~. was measured from the in-plane magnetization curves hereafter), (b) magnetization, and (c) measured anisotropy K', at room temperature. K', is determined from the area surrounded by the parallel and perpendicular magnetization curves.
shows a convex feature and has a maximum at a CoSm layer-thickness of about 200 A. The magnetization shown in Fig. 5b increases about 20% as the CoSm layer-thickness increases from 60 to 960 ~, and this behavior may be explained semiquantitatively in terms of a so-called 'dead layer' which will be discussed below. Fig. 5c shows that the measured anisotropy increases rather rapidly with increasing CoSm layerthickness: K'u varies from ~ 3 × 10 6 to ~ 2 × 107 erg/cm 3 while the CoSm layer-thickness increases from 60 to 960 A. Since samples 2, 4 and 5 show nearly the same size nanocrystallites (50 A) by HRTEM, Fig. 5b shows about the same magnetization and Table 1 shows about the same volume fraction of nanocrystallites; thus they are not the reason for this feature. Furthermore, since all these samples were prepared at the same Ar pressure and with the same thickness of Cr underlayer (1200 A), it
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Z.S. Shah et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
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layer has a microstructure composed of nanocrystallites of ~ 50 ~, which are embedded in an amorphous matrix and the Cr underlayer has a grain size
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of about 250 ,~. Also, our experience indicates that the Co3.7Sm ]lCr samples generally demonstrate better properties, i.e. larger H~ and S, compared with
Co42Sm IlCr. All the hysteresis loops mentioned above were measured by AGFM with maximum applied field Hma~ = 10 kOe. Fig. 9 demonstrates two examples of hysteresis loops of samples 2 and 5 measured by SQUID with H,,,x = 55 kOe. One finds that the loops of sample 2 in Fig. 9a and Fig. 4b or sample 5 in Fig. 9b and Fig. 4e have the same behavior. Also, the K'u values for samples 2 and 5 determined by the areas between parallel and perpendicular curves in Fig. 9a,b agree with those in Fig. 4b,e fairly well.
5. Temperature-dependent magnetic properties Two examples of the temperature-dependent hysteresis are shown in Fig. 10a for sample 2 (CoSm(240 ~,)[JCr) and Fig. 10b for sample 5 (CoSm(960 o A)JICr). These figures display the outline of the temperature-dependent properties of the magnetism:
as temperature decreases, the loops become broader (i.e. H~ increases) and perpendicular magnetization curves become more horizontal (i.e. K', increases as the temperature decreases). The initial magnetization curves in Fig, 10a show that the threshold field Hth, which corresponds to the applied field required to raise the magnetization rapidly, increases as the temperature decreases. Fig. l la shows the temperature dependence of magnetization for samples with a 240 A CoSta layer 1000
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Z.S. Shan et al. / Journal qf Magnetism and Magnetic Materials 161 (1996) 323-336
prepared at 5 mT (sample 1), 12 naT (sample 2) and 30 mT (sample 3) of Ar pressure. Compared with samples 1 and 2, sample 3 shows a stronger temperature dependence. This is possibly due to the fact that the CoSta layer of sample 3 has a larger volume fraction of amorphous material, as revealed by TEM (see Table 1), which causes a stronger temperaturedependent magnetization. Fig. l lb shows the temperature dependence of magnetization for samples with 60 A (sample 4), 240 A (sample 1) and 960 ,~ (sample 5) CoSta layers prepared at the same Ar pressure of 12 mT. Compared with samples 4 and 2, sample 5 shows a stronger temperature dependence. The reason for this behavior is still not clear. Fig. 12a,b shows the temperature dependence of the measured anisotropy for the same samples as in Fig. l la,b. It is found that: (i) Sample 2, which shows the largest H~ and S in Fig. 4, displays the strongest K'u temperature dependence. As T varies from 300 to 10 K, K'u increases from ~ 7.2 × 106 to ~ 4.2 × 107 erg/cm ~. (ii) Sample 5 has the largest K'u ~ 2 × 107 e r g / c m 3 at 300 K; K', increases with decreasing ir and approaches about the same K', value as sample 2 at low T. (iii) Sample 4 has the thinnest CoSta layer of 60 A; however, it shows the weakest K', temperature dependence compared with samples 2 and 5 with thicker CoSta layers. This indicates that the interfaces of CoSta and Cr show only the weaker temperature-dependent feature. This feature is different from conventional multilayers, such as C o / P t [22] and C o / A u [23], which have a stronger temperature-dependent feature for the thinner Co layer samples.
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The temperature-dependent coercivity for the same samples as shown in Figs. 11 and 12 are given in Fig. 13a,b. The coercivity increases with decreasing temperature for all samples: H c increases by a factor of ~ 2 . 5 as T varies from 300 to 50 K for all samples except sample 5 which increases by a factor of 3.3. Comparing Fig. 13a,b with Fig. 12, one finds the temperature-dependent feature of K', is different from that of He. This implies that the anisotropy is not the only source controlling the coercivity and this will be discussed in more detail in the next section. Plots of 'logarithmic H c vs. T' for the same samples as in Fig. 13a,b are given in Fig. 13c,d: this shows good straight lines for all samples. This behavior was considered as evidence for the wall-pinning coercivity mechanism [13,24] which will be discussed in more detail in Section 6.
6. Magnetization reversal
(K)
Fig. 12. Temperature dependence of the measured anisotropy for the samples shown in Fig. 1 la,b.
It is generally accepted that the magnetization reversal or coercivity mechanism can be investigated
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Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
by means of: (i) the initial magnetization curve; (ii) minor loops; (iii) curves of 'logarithmic H c vs. T'; (iv) the 'H~. vs. Hmax' curve, where H~ and Hma x are the coercivity and maximum applied field of minor loops; (v) temperature-dependent magnetic properties studied using Kronmiiller's model [25,26]; and (vi) measurements of the switching volume in terms of the time decay of magnetization or the field-sweep rate dependence of coercivity. In this paper we use all these approaches to study the magnetization reversal behavior of our samples. As is well known, the characteristics of initial magnetization curves, minor loops, 'H~ vs. Hma×' curves and Kronmiiller fitting curves vary with different types of magnetization reversal mechanisms. For example, curves a, b and c in Fig. 14 show the initial curves of nucleation, wall pinning and singleparticle coherent rotation types of reversal mechanisms [26,27]. Examples of minor loops corresponding to the above mechanism can be found in the reports of Wang [27] and Becker [28]. For the ' H c vs. HmaX' curves, a common feature is that there is a region within which Hc is higher than HmaX [27,28] for the nucleation type of reversal and H~ is always less than HmaX for the wall pinning and particle-rotation reversal. Fig. 15 demonstrates the initial magnetization curves and minor loops for samples 1, 2 and 3, which were prepared in an Ar pressure of 5, 12 and 30 mT, respectively. One finds that both samples 1 and 2 show the wall-pinning type of reversal feature, and sample 3 shows the particle-rotation type of reversal feature. The 'He vs. Hma~' curves show that H~ is less than HmaX for all three samples. There-
(a)
(~)
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Fig. 14. Schematic diagram of the initial magnetization curves for (a) nucleation, (b) wall pinning and (c) particle-rotation type mechanisms.
i
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.
i
,
i
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i
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.
.
.
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,
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fore, one may conclude that nucleation is not the major source of magnetization reversal, and the magnetization reversal is dominated by wall pinning at lower Ar pressure (i.e. 5-12 mT) and single-particle rotation at higher Ar pressure (i.e. 30 mT). This transition may be related to the Ar pressure dependence of the microstructure, as shown in Fig. 2a-c: since the film roughness increases with increasing Ar pressure, the particle-like feature of reversal behavior becomes stronger. The straight line feature of the 'logarithmic H~ vs. T' curve as shown in Fig. 13c,d was regarded as a criterion of the pinning mechanism [13,24]. However, Fig. 15c shows particle-rotation behavior and it shows a straight line feature in Fig. 13c, which implies that this criterion is not sufficient and one should analyze the magnetization reversal behavior using different approaches before coming to a conclusion.
Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336 12
Therefore, the curve of ' H e l M ~ vs. 2 K I u 5 / M 2' should be a straight line if this model is applicable. Assuming the slope of the above curve is tan a, which can be determined from the experimental data, then the size of the pinning site may be expressed as
10
0"1
0
/
lOOK
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ro
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4.0[3
333
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and the condition 'r o < 6 B' may be expressed as
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where 6 B is the wall thickness, r 0 the size of the pinning site, Neff the local effective demagnetization factor, and A the exchange constant for the host region. Substituting the 6 n expression of Eq. (5) into Eq. (4) and dividing both sides of Eq. (4) by M~, we have
We notice that there is no parameter 'A' in Eq. (8), i.e. the criterion of ' r 0 < 6 B' may be checked irrespective of the value of the exchange constant A. The ' H c ( T ) / M ~ vs. 2 K I 5 / M [ ' curve for sample 2 is depicted in Fig. 16a. It shows that: (i) The experimental data at different temperatures fits the straight line fairly well, which implies that domain wall pinning is the dominant mechanism. (ii) The calculated size of the pinning site based on Eq. (7) is ~ 3.7 A assuming the exchange constant A --- 3.5 X 10 - 6 erg/cm. (iii) The small pinning-site assumption of ' r 0 < 6 B' is satisfied if the experimental data of tan a and K u are substituted into Eq. (8). (iv) Neff has a small negative value of - 3 . 6 and such behavior has been found for several CoSm Ilfr films in this paper and C o / P t films [29] as well. The macroscopic demagnetization factor N is always positive and varies from 0 to 47r for a continuous medium. The reason for a negative Neff here is attributed to the fact that Neff is an averaged local effective demagnetization factor which is sensitive to the local microstructure and is controlled by the stray fields at the edges and comers of the grains [26]. (v) We have performed a Kronmtiller fitting with the assumption of r 0 > 6 B for the same sample as shown in Fig. 16b, but a straight line fit could not be obtained and this implies that the small pinning site assumption, i.e. r o < 6 B as shown in Fig. 16a, is reasonable. The magnetic switching volume V * is an important parameter for understanding the behavior of magnetization reversal since it influences the media noise and thermal stability in high density magnetic recording. V * may be determined by the 'fieldsweep rate dependence of coercivity', i.e. [29,30]
He(T ) 2"rr r o 2K~ 5 M--"~ = 3v/3 ~ M2
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.
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.
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.
.
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Fig. 16. Experimental curve for (a) ' H c / M ~ vs. 2KIS/M~" and (b) ' H c / M ~ vs. 2K°5/Mf" for sample 2.
Based on Kronmtiller's model [25,26], one example of the temperature dependence of the magnetic properties for sample 2 is given in Fig. 16. When wall pinning is the dominant mechanism, the Kronmiiller formula gives Hc(T)=~
21r ( r o ) 2 K u ~ M~
-- N e f f M s
for r o < 6 B,
(4) 6.= 7r
(5)
A
kBT Neff;
r° < 6B.
(6)
dH~
(9)
334
ZS. Shah et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
where dHa/dt is the sweep rate of the applied field, and k~, T and M~ are the Boltzmann constant, temperature and saturation magnetization, respectively, k is a constant independent of the field sweep rate. Therefore, the switching volume V * may be determined from the slope of the ' H c vs. ln(d H i d t)' curve. The measured ' H c vs. ln(dHa/dt)' curves for the samples in Fig. 1 are shown in Fig. 17, and the V * values are also listed. It can be seen that: (i) As the CoSm layer thickness increases from 60 to 960 (i.e. from sample 4 to 2 to 5) the V * values increase from ~ 2 . 6 × 1 0 ~8 to ~ 4 . 5 × 10 - I s to ~ 1 2 × 10-Is cm 3. It is understandable that the magnetization reversal is in a larger volume while the magnetic layer becomes thicker. (ii) As the Ar pressure during sputtering increases from 5 to 30 mT (i.e. from sample 1 to 2 to 3) V* shows a minimum of ~ 4 . 5 × 1 0 -J8 cm 3 for sample 2 and increases rapidly for both samples 1 and 3. This may be related to the fact that sample 2 has the weaker interaction among grains or nanocrystallites, which causes the larger H~ compared with both samples 1 and 3, and then leads to a smaller switching volume. (iii) V * has an order of magnitude of 1 0 - ] 8 - 1 0 i7 cm 3. This value is much larger than the volume of
the nanocrystallites of 50 A dimension and is close to the grain volume of 250 A dimension. This means that the magnetization reversal roughly corresponds to the switching unit of a CoSm grain. The switching volume V* for samples 1-5 is also determined by the time decay of the magnetization and remanence curve measurements [31,17]. The measured V * for each sample is listed within parentheses in Fig. 17, e.g. V* = 5 . 3 × 10 -18 cm 3 for sample 2. More information about the determination of V * and a discussion of the behavior of V * can be found in an ealier paper [17]. We note that the V * values determined by these two methods are consistent with each other within reasonable error, which perhaps indicates that the V* values in Fig. 17 are acceptable for analyzing the process of magnetization reversal. o
7. C o n c l u s i o n s The following results were obtained based on systematic studies of the structural and magnetic properties of CoSm IICr films. (i) The Cr underlayer shows important effects on the microstructure of the subsequent CoSta layer and
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ZS. Sh.n et al. / J o u o u d ~l Magnetism . m l Magnetic Materials 161 (1996) 3 2 3 - 3 3 6
consequently its magnetic propert!es. The grain size of the Cr underlayer is about 250 A. The CoSta layer on the Cr underlayer continues a similar grain-like morphology, at least up to 240 A thickness. The microstructure of the CoSta layer is composed of nanocrystallites of close-packed structure with about 50 A diameter distributed in an amorphous matrix. The volume fraction of nanocrystallites decreases with increasing Ar sputtering pressure. (ii) By adjusting the preparation conditions, Co37Sm [ICr films with coercivity up to 4.2 kOe'and squareness S = 0.95 have been prepared. (iii) Systematic studies of the CoSta layer-thickness and Ar pressure dependence of magnetic properties have been performed at room temperature. It was found that the maximum intrinsic anisotropy K u at room temperature (K~ = K', - NdM2/2, where NO is the demagnetization factor) is about 4 X 106 and 1.4× 107 erg/cm 3 for CoSm(240 ~,)ILCr and CoSm(960 ~,)[[Cr films, respectively. While these values are large, they are considerably smaller than the 1.1 × 10 s erg/cm 3 for the anisotropy of bulk C%Sm. In the design of future high density recording media, it is essential to actually measure the K u values of the films as-prepared because the thin film structures produced will not in general have Ko values equal to the bulk values of RE-TM intermetallic compounds. (iv) The temperature dependencies of magnetization, coercivity, and measured anisotropy have been studied. The intrinsic anisotropy increases from 4 X 10 6 to 3.9 × 107 erg/cm 3 as the temperature varies from 300 to 10 K for sample 2 (CoSta 240 A I!Cr 1200 ,~) and studies of temperature-dependent magnetic properties using Kronmfiller's formula indicate that wall pinning of the small pinning site is the major coercivity mechanism for this sample. (v) Magnetization reversal studies show that the coercivity mechanism changes from a wall-pinning dominant mechanism at lower Ar pressures (e.g., 5-12 roT) to a particle-rotation dominant mechanism at high pressure (30 roT), and there is no evidence for nucleation-controlled coercivity. (vi) The switching volume is of the order of 1 0 - t T _ 1 0 - 1 8 c m ~ and is influenced by the preparation conditions, and the magnetization reversal takes place roughly at the level of CoSm grains(~ 25(I A) rather than small nanocrystallites ( ~ 50 A).
335
Acknowledgements The authors are grateful to Profs. S.Y. Jeong, B.W. Robertson, and Dr. E. Singleton for their helpful assistance and discussions, Mr. Y.B. Zhang for help in preparing the AFM pictures, F. Foong for measuring the CoSm composition by EDX, B. Jones for X-ray diffraction, and Alex Nolte and Sam Rankin for their help in building the sputtering system. We also acknowledge Prof. H. Kronmtiller for his informative discussions on the coercivity mechanism. This work was supported by the National Science Foundation under grants DMR-9222276 and OSR-9255225, by ARPA/NSIC under grant MDA972-93-1-0009, and by the Center for Materials Research and Analysis at the University of Nebraska.
References [I] E. Murdock, R.F. Simmons and R. Davidson, IEEE Trans. Magn. 28 (1992) 3078. [2] M.P. Sharrock, IEEE Trans. Magn. 26 (1990) 193. [3] T. Yogi, C. Tsang, T. Nguyen, K. Ju, G. Gorman and G. Castilo, IEEE Trans. Magn. 26 (1990) 2271. [4] M. Futamoto, F. Kugiya, M. Suzuki, H. Takano, Y. Matsuda, N. Inaba, Y. Migamura, K. Akagi, T. Nakao, H. Sawagucbi, H. Fukuoka, T. Mnemoto and T. Takagaki, IEEE Trans. Magn. 27 (1991) 5280. [5] S.H. Aly, T.D. Cheung, L. Wickramasekara and F.J. Cadieu, J. Appl. Phys. 57 (1985) 2147. [6] F.J. Cadieu, H. Hedge and K. Chen, J. Appl. Phys. 67 (1990) 4969. [7] F.J. Cadieu, J. Vac. Sci. Technol. A 6 (1988) 1668. [8] D. Wang and W.D, Doyle, IEEE Trans. Magn. 30 (1994) 4032. [9] E.M.T. Velu and D.N. Lambeth, J. Appl. Phys. 69 (1991) 5175. [10] E.M.T. Velu and D.N. Lambeth, IEEE Trans. Magn. 28 (1992) 3249. [I 1] E.M.T. Velu, D.N. Lambeth, J.T. Thornton and P.E, Russel, J. Appl. Phys. 75 (1994) 6132. [12] Y. Okumura, H. Fujimori, O. Suzuki, N. Hosoya, X.B. Yang and H. Morita, IEEE Trans. Magn. 30 (1994) 4038. [13] Y. Okumura, O. Suzuki, H. Morita, X.B. Yang and H. Fujimori, J. Magn. Magn. Mater. 146 (1995) 5. [14] Y. Liu, B. Robertson, Z.S. Shah, S. Malhotra, M.J. Yu, S.K. Renukunta, S.H. Liou and D.J. Sellmyer, IEEE Trans. Magn. 30 (I 994) 4035. [15] Y. Liu, B. Robertson, Z.S. Shan, S.H. Liou and D.J. Sellmyer, J. Appl. Phys. 77 (1995) 3831. [16] Y. Liu, D.J. Sellmyer, B.W. Robertson, Z.S. Shan and S.H. Liou, IEEE Trans. Magn. 31 (1995) 2740.
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[17] E.W. Singleton, Z.S. Shan, S.Y. Jeong and DJ. Sellmyer, IEEE Trans. Magn. 31 (1995) 2743. [18] D.J. Sellmyer, Z.S. Shan, Y. Liu, S.H. Liou, S.S. Malhotra and B.W. Robertson, Scripta Metal. Mater. 33 (1995) 1545. [19] Z.S. Shan, S. Nafis, J. Woollam, S.H, Liou and D.J. Sellmyer, J. Appl. Phys. 73 (1993) 6347. [20] C.D. England, W.R. Bennett and C.M. Falco, J. Appl. Phys. 64 (1988) 5757. [21] G. Xiao and C.L. Chien, J. Appl. Phys. 61 (1987) 4061. [22] Z.S. Shan, J.X. Shen, R.D. Kirby and D.J. Sellmyer, J. Appl. Phys. 75 (1994) 6418, [23] Z.S. Shan, J.X. Shen and D.J. Sellmyer (to be published). [24] Y.X. Bo and T. Miyazaki, J. Magn. Magn. Mater. 86 (1990) 37.
[25] H. Kronmiiller, K.D. Durst and M. Sagawa, J. Magn. Magn. Mater. 74 (1988) 291. [26] H. Kronmiiller, in: Proc. Magneto-Optic Recording Int. Symp., J. Magn. Soc. Jpn. 17 (S1) (1992) 260. [27] D. Wang, Ph.D. Dissertation, University of Nebraska at Lincoln, 1993. [28] J. Becker, IEEE Trans. Magn. 12 (1976) 965. [29] T, Suzuki, IEEE Trans. Magn. 31 (1995) 4085. [30] P. Bruno, G. Bayreuther, P. Beauvillain, C. Chappert, G. Lugert, D. Renard, J.P. Renard and J. Seiden, J. Appl. Phys. 68 (1990) 5759. [31] K. O'Grady, T. Thompson, S.J. Greaves and G. Bayreuther, J. Appl. Phys. 75 (1994) 6849.
~
Journal of Jan nlognoUc mMerlals
ELSEVIER
Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
Magnetic properties and magnetization reversal of CoSm ICr thin films Z.S. Shan a,b, S.S. Malhotra a.b, S.H. Liou a,b, Yi Liu a,c, M. Yu a,b, D.J. Sellmyer
a,b, *
Center for Materials Research and Analysis, Uniz:ersit).'of Nebraska, Lincoln, NE 68588-0113, USA b Behlen Laborator3., of Physics, Unit'ersio~ of Nebraska, Lincoln, NE 68588-0113, USA c Department of Mechanical Engineering, Unit,ersi~, of Nebraska, Lincoln, NE 68588-0113, USA
Received 20 July 1995;revised l0 March 1996
Abstract In this paper the magnetic and structural properties of CoSm thin films with a Cr underlayer (CoSm I[Cr) are presented, with emphasis on the measurements of anisotropy at room and low temperature and magnetization reversal. The grain size of the Cr underlayer is about 250 A and the thin CoSm layer (e.g., 240 A) inherits this grain size. The CoSta layer consists of nanocrystallites, about 50 A in diameter, embedded in an amorphous matrix. The Ar pressure, CoSm layer-thickness, and temperature dependencies of magnetic properties including magnetization, coercivity and especially the anisotropy were investigated systematically. CoSta ILCr with coercivity up to 4.2 kOe at room temperature has been prepared. The intrinsic anisotropy is 4 x 10 6 and 1.4X 10 7 erg/cm 3 at room temperature for CoSm(240 *)llCr and CoSm(960 A)IICr, respectively, and both increase to 3.9 × 107 erg/cm 3 at 10 K. Magnetization reversal studies indicate that the coercivity mechanism changes from wall pinning for samples prepared at lower Ar pressure (5-12 mT) to single-particle coherent rotation for samples prepared at higher pressure (30 mT). The correlations between the microstructure and magnetic properties are discussed.
Keywords: Magnetization reversal; Coercivity mechanism; Magnetic switching volume; Microstructure; Magnetic anisotropy; CoSta film
1. Introduction Thin-film recording media have received continuous attention for high-density hard-disk recording and recent progress has ensured a dominant role of magnetic recording for the foreseeable future. Murdock et al. [1] have predicted that longitudinal recording of 10 G b / i n . 2 requires the media to have
* Corresponding author. Fax: + 1-402-4722879; email: [email protected].
high coercivity (2500-4500 Oe), small grain size ( 8 0 - 1 0 0 ,~), and the individual particles should be isolated from one another in terms of exchange interaction. Sharrock [2] has pointed out that the high anisotropy is required so the thermal stability requirement of KuV/kBT> 100, where K u, V, k B and T are the anisotropy, grain volume, Boltzmann constant and temperature respectively, is satisfied. With Co-based thin-film media, IBM and Hitachi have made hard disks with areal densities of 1 G b / i n . 2 [3] and 2 G b / i n . 2 [4] respectively. Besides the CoCrX (X = Ta, Pt, Ni, etc.) thin-film media, rare-earth
0304-8853/96/$15.00 Copyright © 1996 Elsevier Science B.V. All rights reserved. PII S0304- 8853(96)00296-X
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( R E ) permanent-magnet films or multilayers, such as
2. F a b r i c a t i o n a n d c h a r a c t e r i z a t i o n
CoSm [5-7] and FeSmN [8] films, may be considered as candidates because of their large coercivity and anisotropy. Among RE semi-hard magnetic films, CoSm films prepared at room temperature have gained increasing attention because they may have both small grain size and high uniaxial anisotropy, which is essential for thin-film media of high-density recording. Velu and Lambeth [9-11] reported that CoSta films on Cr underlayers, denoted here as CoSm I[Cr, with coercivity up to 3000 Oe have been prepared by rf sputtering onto 7059 Coming glass substrates at r o o m t e m p e r a t u r e . The high coercivity was attributed to the very large anisotropy of CosSm and the in-plane lattice matching between the CosSm(1120) and Cr(ll0). In their papers the effects of sputtering conditions on the film coercivity have been systematically investigated and the microstructure of CoSta film and Cr underlayer was studied by means of AFM (atomic force microscopy) [11], Their studies show that the CoSmllCr thin films may be a promising medium for high-density longitudinal recording. Recently, Okumura et al. [12,13] reported that a maximum coercivity of 3.5 kOe was obtained for a Co85Sm~5[[Cr film and claimed that coercivity mainly results from both the net-like microstructure and fluctuation in the local anisotropy of the Co85 Sm ~5 layer. Liu et al. [14-16] reported systematic studies of the microstructure of CoSm II Cr including Cr underlayer, CoSm layer, the nanocrystallites in CoSm layer, etc. Singleton et al. [17] reported measurements of the switching volume which is 10- l 8_ 10 ~7 cm 3 and Sellmyer et al. [18] gave a succinct review of the magnetic and structural properties for CoSm IICr films. However, the magnetic properties of CoSm [JCr films have not been reported systematically. For example, a very important property, the anisotropy of CoSm J[Cr films, has not yet been studied intensively. In this paper we report our studies of the magnetic properties of films at room and low temperature, with emphasis on the anisotropy measurement and magnetization reversal analyses. The microstructure and the relationships between the magnetic and microstructure properties will also be discussed briefly.
The CoxSm [[Cr films were sputtered onto microscope cover glass substrates by a dc magnetron sputtering machine. The CoxSm targets were home made by pressing the CoxSm powder and then sintering in vacuum at 1100°C for 30 min. The commercial Cr target had a 99.9% purity. The base pressure of the sputtering system was 1-2 × 10 -7 Torr and the Ar pressure during sputtering varied from 5 to 30 mT to study the Ar-pressure dependence of the magnetic properties, which will be discussed in Section 4.4. The sputtering rate and power for CoSm were ~ 1 .~/s and 15 W and for Cr were ~ 2.5 .~/s and 36 W, respectively. The distance from the substrate to both guns was about 5 cm. Twelve samples were made in one vacuum run to avoid undesired changes of preparation conditions for one series of samples. Using CoxSm (x = 5, 4, 3.5) targets, the compositions of the sputtered films, which were determined using energy dispersive X-ray spectroscopy (EDX) in a scanning electron microscope, were Co5.5Sm, Co4.2Sm and Co3.75m, respectively. The largest coercivities measured were 1.5, 3.7 and 4.2 kOe for the above films. In this paper we mainly focus on studying the Co4.2SmllCr properties, and simply use CoSm to denote Co4.2Sm hereafter. The structural properties were examined by X-ray diffraction, AFM and transition electron microscopy (TEM). The magnetic properties were studied by an AGFM (alternating gradient force magnetometer) from 300 to 10 K and selected samples were measured by a SQUID magnetometer at applied fields up to 55 kOe. In this paper the structural properties and temperature-dependent magnetism were investigated mainly for the samples shown in Fig. 1, i.e. the sputtered Ar pressure which is identical for both Cr and CoSm layers changes from 5 to 30 mT (i.e. samples l, 2 and 3) and the CoSm layer-thickness varies from 60 to 960 A (i.e. samples 4, 2 and 5). Two sets of samples were prepared: one set with no Cr overcoat layer for AFM and TEM studies and the other set with a Cr overcoat layer of 200 ~. for magnetic property studies. Both sets of samples have a Cr underlayer of 1200 ~,. We have also prepared samples with different Ar pressures for the Cr underlayer
325
ZS. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
3. Structural properties
#1
PAr: 5 m r dcos~ :240A #5 # 4 rr~TPAr:12 # 2 m ] PAr:12 m [ PAr:12 Jdcosm: 60A [dcosm:240Adcos~:960A #3 IPAr:30 m T
[dcosm:240~, Fig. 1. Samples used for structural and magnetic studies. From samples 1 to 2 to 3, the CoSm layer-thickness is fixed at 240 ,~, while the Ar sputtering pressure increases from 5 to 12 to 30 mTorr. From samples 4 to 2 to 5, the Ar sputtering pressure is fixed at 12 mTorr while the CoSm layer-thickness increases from 60 to 240 to 960 ,~. and the CoSm layer. The Ar pressure was appropriately varied to achieve the highest coercivity values. One example is shown in Fig. 8.
The X-ray diffraction patterns of the Cr underlayer on glass substrates show that, as the sputtered Ar pressure varies from 5 to 30 roT, the diffraction patterns only show a single (110) peak, indicating that the Cr underlayer is highly textured. This is a necessary condition for the subsequent CoSm layer to have large coercivity, as has been proved experimentally by Velu and Lambeth [9,10] and Okumura et al. [12,13]. This character was interpreted based on the lattice matching between CosSm(1120) and Cr(ll0). However, recent work has shown that a disordered hexagonal C o - S m alloy is present rather than the C % S m phase [14-16]. X-ray diffraction studies of the CoxSm ( x = 3.5, 4, 5) layer have been performed either for an individual CoxSm layer on a glass substrate or a CoxSm layer on a Cr underlayer which is on the glass
#1 PAr: 5 mT
15,0
j
d costa:240/~
#5 Ip~:# 4 #2 2 m,TIP,~.:I2 m~ P~.:12 m'[ [ dcosm: 60A [dcosm:240A dcosm:960A
7.5
PAr:30 mT [dcosm:2401
500
~
508
Fig. 2. Atomic force micrographs for the samples shown in Fig. 1. (a,b,c) The Ar pressure dependence of the surface morphology;(d,b,e) the CoSm layer-thickness dependence of the surface morphology.
326
Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
substrate. No crystalline diffraction peaks are observed in both cases. Presumably the crystallite size is too small to show well defined peaks in these cases. To study the surface morphology of CoSta layers, a series of samples as shown in Fig. 1 but without a Cr overcoat layer was specially prepared as has been mentioned in Section 2. The surface morphology revealed by AFM is displayed in Fig. 2. We also prepared a Cr layer of 1200 ,~ on a glass substrate and it displayed nearly the same topographical structure as shown in Fig. 2. This confirms that the surface morphology in Fig. 2a-2e is inherited from the Cr underlayer morphology. One can find from Fig. 2 that the average grain size of the Cr underlayer is about 250 A. As the sputtered Ar pressure increases, i.e. from Fig. 2a to Fig. 2b to Fig. 2c, the surface roughness, which is the roughness superposition of the CoSm layer and Cr underlayer, increases. As the CoSta layer-thickness increases, i.e. from Fig. 2d to Fig. 2b to Fig. 2e, the surface roughness is nearly the same because they are fabricated in the same Ar pressure. However, Fig. 2e shows larger diameters of 'hills' since the 960 ~, thick CoSta layer may fill up some voids and make the 'hills' larger. The AFM micrographs only show the surface morphology of the CoSta layer and the microstructure inside the CoSm layer will be discussed in terms of the TEM results in the next paragraph. Systematic studies of bright-field and high-resolution TEM (HRTEM) for the same series of samples in Figs. 1 and 2 have been performed and the results are presented in earlier papers [14-16]. Here, a brief summary of the microstructure of the Cr underlayer and the CoSta layer is given: the Cr underlayer has a grain size of about 250 A and this grain image can be seen in samples 2, 3 and 4 in Figs. 1 and 2. But this Cr grain image cannot be seen in sample 5 because the CoSm layer is 960 A thick and the Cr grain image cannot be continued through the whole CoSta layer; the Cr grain image also cannot be seen in sample 1, possibly because the Ar sputtering pressure was only 5 mT in this case and the Cr underlayer possesses a smooth surface as shown in Fig. 2a, lacking the columnar structure. The CoSta grains of about 250 A in dimension for samples 2, 3 and 4 consist of CoSta nanocrystallites of typical dimension 50 A embedded in an amorphous CoSta
Grains (250~)
Amorphous matrix •
Crystallites (50/~) /
Fig. 3. Schematic diagram of a CoSm grain of ~ 250 ,~ dimension: the nanocrystallites of dimension ~ 50 ,~ are embedded in an amorphous matrix.
matrix, as shown in Fig. 3. The nanocrystallites have a disordered close-packed structure with varying stacking mode from crystallite to crystallite and sometimes within a single crystallite. The volume fraction of the nanocrystallites in the CoSm layer is shown in Table 1. It is seen that the volume fraction decreases from 91 to 65 to 54% as the Ar sputtering pressure is increased from 5 to 12 to 30 mT (i.e. from sample 1 to 2 to 3) respectively; as the CoSta layer-thickness increases from 60 to 240 to 960 ,~ (i.e. from sample 4 to 2 to 5), the volume fractions of the nanocrystallites remain almost unchanged. These properties are only mentioned here to provide the structural basis for understanding the magnetic properties with the detailed analysis in our other papers published earlier [ 14-16].
4. Magnetic properties at r o o m t e m p e r a t u r e
4.1. Hysteresis loops at room temperature
Both the parallel (i.e. H Jlfilm plane) and perpendicular (i.e. H .1_film plane) magnetization curves for the samples shown in Figs. 1 and 2 are demonstrated in Fig. 4. This figure gives the general features of the Ar pressure dependence (Fig. 4a-4c) and CoSta layer-thickness dependence (Fig. 4d,4e) of the magnetic properties: (i) In this series of samples, sample 2, which was prepared in 12 mT Ar pressure with a 240 ,~ CoSta layer and a 1200 A Cr underlayer, shows the largest coercivity of H c = 2.6 kOe and squareness S = 0.88. Here, Hc was measured from
Z.S. Shan et al./ Journal qf Magnetism and Magnetic Materials 161 (1996) 323-336 Table 1 Ar pressure and CoSm layer-thickness effects on the volume fraction (%) of the nanocrystallites, where Pat is the Ar pressure during sputtering and dsmco the thickness of the SmCo42 layer PAr
Volume ~actionVc(%)
(mT)
dsmco = 6 0 A
dsmco = 240A
dsmco = 960A
5 12 30
65
91 65 54
63
and sample 3 in Fig. 4c); as the CoSm layer-thickness increases (i.e. from Fig. 4d to Fig. 4e), the perpendicular magnetization curves become narrower, which means that the magnetic moments are better aligned in the film plane for the thicker CoSm layer than for the thinner CoSm layer. (iii) The initial curves display a typical wall pinning or particle-rotation feature. The coercivity mechanism and magnetization reversal properties will be discussed in more detail in Section 6. 4.2. Layer-thickness dependence o f magnetic properties
the in-plane magnetization curve. (ii) When looking at the perpendicular magnetization curves one finds that, as the Ar pressure increases (i.e. from Fig. 4a to Fig. 4c), sample 2 in Fig. 4b shows the narrowest magnetization curve, which means that the magnetic moments are well aligned in the film plane compared with the other two samples (i.e. sample 1 in Fig. 4a
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328
Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336 3000
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4.3. Determination of the 'interface thickness'
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The CoSm layer magnetization in the presence of the interface between the CoSm and Cr layers (both under and coating layers) can be expressed as [19]
CoSrn layer-thickness (~) Fig. 5. CoSm layer-thickness dependence of (a) coercivity (H~. was measured from the in-plane magnetization curves hereafter), (b) magnetization, and (c) measured anisotropy K', at room temperature. K', is determined from the area surrounded by the parallel and perpendicular magnetization curves.
shows a convex feature and has a maximum at a CoSm layer-thickness of about 200 A. The magnetization shown in Fig. 5b increases about 20% as the CoSm layer-thickness increases from 60 to 960 ~, and this behavior may be explained semiquantitatively in terms of a so-called 'dead layer' which will be discussed below. Fig. 5c shows that the measured anisotropy increases rather rapidly with increasing CoSm layerthickness: K'u varies from ~ 3 × 10 6 to ~ 2 × 107 erg/cm 3 while the CoSm layer-thickness increases from 60 to 960 A. Since samples 2, 4 and 5 show nearly the same size nanocrystallites (50 A) by HRTEM, Fig. 5b shows about the same magnetization and Table 1 shows about the same volume fraction of nanocrystallites; thus they are not the reason for this feature. Furthermore, since all these samples were prepared at the same Ar pressure and with the same thickness of Cr underlayer (1200 A), it
M'=M 0 1-
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(3)
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Z.S. Shah et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
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layer has a microstructure composed of nanocrystallites of ~ 50 ~, which are embedded in an amorphous matrix and the Cr underlayer has a grain size
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of about 250 ,~. Also, our experience indicates that the Co3.7Sm ]lCr samples generally demonstrate better properties, i.e. larger H~ and S, compared with
Co42Sm IlCr. All the hysteresis loops mentioned above were measured by AGFM with maximum applied field Hma~ = 10 kOe. Fig. 9 demonstrates two examples of hysteresis loops of samples 2 and 5 measured by SQUID with H,,,x = 55 kOe. One finds that the loops of sample 2 in Fig. 9a and Fig. 4b or sample 5 in Fig. 9b and Fig. 4e have the same behavior. Also, the K'u values for samples 2 and 5 determined by the areas between parallel and perpendicular curves in Fig. 9a,b agree with those in Fig. 4b,e fairly well.
5. Temperature-dependent magnetic properties Two examples of the temperature-dependent hysteresis are shown in Fig. 10a for sample 2 (CoSm(240 ~,)[JCr) and Fig. 10b for sample 5 (CoSm(960 o A)JICr). These figures display the outline of the temperature-dependent properties of the magnetism:
as temperature decreases, the loops become broader (i.e. H~ increases) and perpendicular magnetization curves become more horizontal (i.e. K', increases as the temperature decreases). The initial magnetization curves in Fig, 10a show that the threshold field Hth, which corresponds to the applied field required to raise the magnetization rapidly, increases as the temperature decreases. Fig. l la shows the temperature dependence of magnetization for samples with a 240 A CoSta layer 1000
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331
Z.S. Shan et al. / Journal qf Magnetism and Magnetic Materials 161 (1996) 323-336
prepared at 5 mT (sample 1), 12 naT (sample 2) and 30 mT (sample 3) of Ar pressure. Compared with samples 1 and 2, sample 3 shows a stronger temperature dependence. This is possibly due to the fact that the CoSta layer of sample 3 has a larger volume fraction of amorphous material, as revealed by TEM (see Table 1), which causes a stronger temperaturedependent magnetization. Fig. l lb shows the temperature dependence of magnetization for samples with 60 A (sample 4), 240 A (sample 1) and 960 ,~ (sample 5) CoSta layers prepared at the same Ar pressure of 12 mT. Compared with samples 4 and 2, sample 5 shows a stronger temperature dependence. The reason for this behavior is still not clear. Fig. 12a,b shows the temperature dependence of the measured anisotropy for the same samples as in Fig. l la,b. It is found that: (i) Sample 2, which shows the largest H~ and S in Fig. 4, displays the strongest K'u temperature dependence. As T varies from 300 to 10 K, K'u increases from ~ 7.2 × 106 to ~ 4.2 × 107 erg/cm ~. (ii) Sample 5 has the largest K'u ~ 2 × 107 e r g / c m 3 at 300 K; K', increases with decreasing ir and approaches about the same K', value as sample 2 at low T. (iii) Sample 4 has the thinnest CoSta layer of 60 A; however, it shows the weakest K', temperature dependence compared with samples 2 and 5 with thicker CoSta layers. This indicates that the interfaces of CoSta and Cr show only the weaker temperature-dependent feature. This feature is different from conventional multilayers, such as C o / P t [22] and C o / A u [23], which have a stronger temperature-dependent feature for the thinner Co layer samples.
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The temperature-dependent coercivity for the same samples as shown in Figs. 11 and 12 are given in Fig. 13a,b. The coercivity increases with decreasing temperature for all samples: H c increases by a factor of ~ 2 . 5 as T varies from 300 to 50 K for all samples except sample 5 which increases by a factor of 3.3. Comparing Fig. 13a,b with Fig. 12, one finds the temperature-dependent feature of K', is different from that of He. This implies that the anisotropy is not the only source controlling the coercivity and this will be discussed in more detail in the next section. Plots of 'logarithmic H c vs. T' for the same samples as in Fig. 13a,b are given in Fig. 13c,d: this shows good straight lines for all samples. This behavior was considered as evidence for the wall-pinning coercivity mechanism [13,24] which will be discussed in more detail in Section 6.
6. Magnetization reversal
(K)
Fig. 12. Temperature dependence of the measured anisotropy for the samples shown in Fig. 1 la,b.
It is generally accepted that the magnetization reversal or coercivity mechanism can be investigated
332
Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
by means of: (i) the initial magnetization curve; (ii) minor loops; (iii) curves of 'logarithmic H c vs. T'; (iv) the 'H~. vs. Hmax' curve, where H~ and Hma x are the coercivity and maximum applied field of minor loops; (v) temperature-dependent magnetic properties studied using Kronmiiller's model [25,26]; and (vi) measurements of the switching volume in terms of the time decay of magnetization or the field-sweep rate dependence of coercivity. In this paper we use all these approaches to study the magnetization reversal behavior of our samples. As is well known, the characteristics of initial magnetization curves, minor loops, 'H~ vs. Hma×' curves and Kronmiiller fitting curves vary with different types of magnetization reversal mechanisms. For example, curves a, b and c in Fig. 14 show the initial curves of nucleation, wall pinning and singleparticle coherent rotation types of reversal mechanisms [26,27]. Examples of minor loops corresponding to the above mechanism can be found in the reports of Wang [27] and Becker [28]. For the ' H c vs. HmaX' curves, a common feature is that there is a region within which Hc is higher than HmaX [27,28] for the nucleation type of reversal and H~ is always less than HmaX for the wall pinning and particle-rotation reversal. Fig. 15 demonstrates the initial magnetization curves and minor loops for samples 1, 2 and 3, which were prepared in an Ar pressure of 5, 12 and 30 mT, respectively. One finds that both samples 1 and 2 show the wall-pinning type of reversal feature, and sample 3 shows the particle-rotation type of reversal feature. The 'He vs. Hma~' curves show that H~ is less than HmaX for all three samples. There-
(a)
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Fig. 14. Schematic diagram of the initial magnetization curves for (a) nucleation, (b) wall pinning and (c) particle-rotation type mechanisms.
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fore, one may conclude that nucleation is not the major source of magnetization reversal, and the magnetization reversal is dominated by wall pinning at lower Ar pressure (i.e. 5-12 mT) and single-particle rotation at higher Ar pressure (i.e. 30 mT). This transition may be related to the Ar pressure dependence of the microstructure, as shown in Fig. 2a-c: since the film roughness increases with increasing Ar pressure, the particle-like feature of reversal behavior becomes stronger. The straight line feature of the 'logarithmic H~ vs. T' curve as shown in Fig. 13c,d was regarded as a criterion of the pinning mechanism [13,24]. However, Fig. 15c shows particle-rotation behavior and it shows a straight line feature in Fig. 13c, which implies that this criterion is not sufficient and one should analyze the magnetization reversal behavior using different approaches before coming to a conclusion.
Z.S. Shan et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336 12
Therefore, the curve of ' H e l M ~ vs. 2 K I u 5 / M 2' should be a straight line if this model is applicable. Assuming the slope of the above curve is tan a, which can be determined from the experimental data, then the size of the pinning site may be expressed as
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We notice that there is no parameter 'A' in Eq. (8), i.e. the criterion of ' r 0 < 6 B' may be checked irrespective of the value of the exchange constant A. The ' H c ( T ) / M ~ vs. 2 K I 5 / M [ ' curve for sample 2 is depicted in Fig. 16a. It shows that: (i) The experimental data at different temperatures fits the straight line fairly well, which implies that domain wall pinning is the dominant mechanism. (ii) The calculated size of the pinning site based on Eq. (7) is ~ 3.7 A assuming the exchange constant A --- 3.5 X 10 - 6 erg/cm. (iii) The small pinning-site assumption of ' r 0 < 6 B' is satisfied if the experimental data of tan a and K u are substituted into Eq. (8). (iv) Neff has a small negative value of - 3 . 6 and such behavior has been found for several CoSm Ilfr films in this paper and C o / P t films [29] as well. The macroscopic demagnetization factor N is always positive and varies from 0 to 47r for a continuous medium. The reason for a negative Neff here is attributed to the fact that Neff is an averaged local effective demagnetization factor which is sensitive to the local microstructure and is controlled by the stray fields at the edges and comers of the grains [26]. (v) We have performed a Kronmtiller fitting with the assumption of r 0 > 6 B for the same sample as shown in Fig. 16b, but a straight line fit could not be obtained and this implies that the small pinning site assumption, i.e. r o < 6 B as shown in Fig. 16a, is reasonable. The magnetic switching volume V * is an important parameter for understanding the behavior of magnetization reversal since it influences the media noise and thermal stability in high density magnetic recording. V * may be determined by the 'fieldsweep rate dependence of coercivity', i.e. [29,30]
He(T ) 2"rr r o 2K~ 5 M--"~ = 3v/3 ~ M2
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Based on Kronmtiller's model [25,26], one example of the temperature dependence of the magnetic properties for sample 2 is given in Fig. 16. When wall pinning is the dominant mechanism, the Kronmiiller formula gives Hc(T)=~
21r ( r o ) 2 K u ~ M~
-- N e f f M s
for r o < 6 B,
(4) 6.= 7r
(5)
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r° < 6B.
(6)
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(9)
334
ZS. Shah et al. / Journal of Magnetism and Magnetic Materials 161 (1996) 323-336
where dHa/dt is the sweep rate of the applied field, and k~, T and M~ are the Boltzmann constant, temperature and saturation magnetization, respectively, k is a constant independent of the field sweep rate. Therefore, the switching volume V * may be determined from the slope of the ' H c vs. ln(d H i d t)' curve. The measured ' H c vs. ln(dHa/dt)' curves for the samples in Fig. 1 are shown in Fig. 17, and the V * values are also listed. It can be seen that: (i) As the CoSm layer thickness increases from 60 to 960 (i.e. from sample 4 to 2 to 5) the V * values increase from ~ 2 . 6 × 1 0 ~8 to ~ 4 . 5 × 10 - I s to ~ 1 2 × 10-Is cm 3. It is understandable that the magnetization reversal is in a larger volume while the magnetic layer becomes thicker. (ii) As the Ar pressure during sputtering increases from 5 to 30 mT (i.e. from sample 1 to 2 to 3) V* shows a minimum of ~ 4 . 5 × 1 0 -J8 cm 3 for sample 2 and increases rapidly for both samples 1 and 3. This may be related to the fact that sample 2 has the weaker interaction among grains or nanocrystallites, which causes the larger H~ compared with both samples 1 and 3, and then leads to a smaller switching volume. (iii) V * has an order of magnitude of 1 0 - ] 8 - 1 0 i7 cm 3. This value is much larger than the volume of
the nanocrystallites of 50 A dimension and is close to the grain volume of 250 A dimension. This means that the magnetization reversal roughly corresponds to the switching unit of a CoSm grain. The switching volume V* for samples 1-5 is also determined by the time decay of the magnetization and remanence curve measurements [31,17]. The measured V * for each sample is listed within parentheses in Fig. 17, e.g. V* = 5 . 3 × 10 -18 cm 3 for sample 2. More information about the determination of V * and a discussion of the behavior of V * can be found in an ealier paper [17]. We note that the V * values determined by these two methods are consistent with each other within reasonable error, which perhaps indicates that the V* values in Fig. 17 are acceptable for analyzing the process of magnetization reversal. o
7. C o n c l u s i o n s The following results were obtained based on systematic studies of the structural and magnetic properties of CoSm IICr films. (i) The Cr underlayer shows important effects on the microstructure of the subsequent CoSta layer and
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Fig. 17. Plots of ' H c vs. l n ( d H j d t ) ' for samples shown in Fig. 1 (note: the values on the abscissa are ( d H J d t ) ) . V * values determined by 'H~ vs. l n ( d H j d t ) " and the time decay of the magnetization are listed at the top and bottom (within parentheses) of the figure for each sample, respectively.
ZS. Sh.n et al. / J o u o u d ~l Magnetism . m l Magnetic Materials 161 (1996) 3 2 3 - 3 3 6
consequently its magnetic propert!es. The grain size of the Cr underlayer is about 250 A. The CoSta layer on the Cr underlayer continues a similar grain-like morphology, at least up to 240 A thickness. The microstructure of the CoSta layer is composed of nanocrystallites of close-packed structure with about 50 A diameter distributed in an amorphous matrix. The volume fraction of nanocrystallites decreases with increasing Ar sputtering pressure. (ii) By adjusting the preparation conditions, Co37Sm [ICr films with coercivity up to 4.2 kOe'and squareness S = 0.95 have been prepared. (iii) Systematic studies of the CoSta layer-thickness and Ar pressure dependence of magnetic properties have been performed at room temperature. It was found that the maximum intrinsic anisotropy K u at room temperature (K~ = K', - NdM2/2, where NO is the demagnetization factor) is about 4 X 106 and 1.4× 107 erg/cm 3 for CoSm(240 ~,)ILCr and CoSm(960 ~,)[[Cr films, respectively. While these values are large, they are considerably smaller than the 1.1 × 10 s erg/cm 3 for the anisotropy of bulk C%Sm. In the design of future high density recording media, it is essential to actually measure the K u values of the films as-prepared because the thin film structures produced will not in general have Ko values equal to the bulk values of RE-TM intermetallic compounds. (iv) The temperature dependencies of magnetization, coercivity, and measured anisotropy have been studied. The intrinsic anisotropy increases from 4 X 10 6 to 3.9 × 107 erg/cm 3 as the temperature varies from 300 to 10 K for sample 2 (CoSta 240 A I!Cr 1200 ,~) and studies of temperature-dependent magnetic properties using Kronmfiller's formula indicate that wall pinning of the small pinning site is the major coercivity mechanism for this sample. (v) Magnetization reversal studies show that the coercivity mechanism changes from a wall-pinning dominant mechanism at lower Ar pressures (e.g., 5-12 roT) to a particle-rotation dominant mechanism at high pressure (30 roT), and there is no evidence for nucleation-controlled coercivity. (vi) The switching volume is of the order of 1 0 - t T _ 1 0 - 1 8 c m ~ and is influenced by the preparation conditions, and the magnetization reversal takes place roughly at the level of CoSm grains(~ 25(I A) rather than small nanocrystallites ( ~ 50 A).
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Acknowledgements The authors are grateful to Profs. S.Y. Jeong, B.W. Robertson, and Dr. E. Singleton for their helpful assistance and discussions, Mr. Y.B. Zhang for help in preparing the AFM pictures, F. Foong for measuring the CoSm composition by EDX, B. Jones for X-ray diffraction, and Alex Nolte and Sam Rankin for their help in building the sputtering system. We also acknowledge Prof. H. Kronmtiller for his informative discussions on the coercivity mechanism. This work was supported by the National Science Foundation under grants DMR-9222276 and OSR-9255225, by ARPA/NSIC under grant MDA972-93-1-0009, and by the Center for Materials Research and Analysis at the University of Nebraska.
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