Pt multilayers and L10 FePt thin films

Thin Solid Films 517 (2008) 531–537

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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / t s f

Dependence of grain sizes and microstrains on annealing temperature in Fe/Pt multilayers and L10 FePt thin films Nikolay Zotov a,⁎, Jürgen Feydt b, Alfred Ludwig a a b

Institute for Materials, Ruhr-University Bochum, 44780 Bochum, Germany Forschungszentrum Caesar, Ludwig-Erhard-Allee 2, 53175 Bonn, Germany

A R T I C L E

I N F O

Article history: Received 23 July 2007 Received in revised form 13 June 2008 Accepted 20 June 2008 Available online 25 June 2008 Keywords: Grain growth FePt Multilayers Thin films Annealing X-ray diffraction

A B S T R A C T Fe/Pt multilayers (MLs) with an overall composition Fe52Pt48, deposited by magnetron sputtering on thermally oxidized Si wafer substrates, were post annealed in vacuum at various temperatures (TA) in the range 573–973 K. The MLs transform directly and completely into polycrystalline hard-magnetic FePt thin films with ordered L10 structure above TA = 573 K. The evolution of the microstructure, the order parameter and the stacking fault density with annealing temperature was investigated by ex-situ X-ray diffraction and line-broadening analysis. The average microstrains beN of the thin films are relatively small (beN ~ 0.2%) and remain practically constant as a function of TA. The thin films show anisotropic size-broadening and grain growth: fast growth rate along the [221] direction and a slow growth rate along the [001] direction. The annealing temperature dependence of the average grain size bDN could be described by a grain growth model with grain growth exponent n = 3 ± 0.5 and activation energy 0.51 ± 0.07 eV. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Due to the very high magnetocrystalline anisotropy of the ordered face-centred (L10) FePt phase [1], hard magnetic FePt thin films and nanoparticles have received considerable attention in recent years. L10 FePt is an attractive candidate for high density magnetic recording applications [2–5] as well as for permanent magnets with high-energy product such as exchange-spring magnets, consisting of exchangecoupled hard and soft magnetic phases [6–7]. For all of these applications, specially tailored microstructures are required with well-defined grain sizes. In order to decrease the media noise in highdensity longitudinal recording media, the average in-plane grain size must be about 6–9 nm [5]. In exchange-spring magnets, the soft regions must be smaller than the Bloch wall thickness of the magnetically hard phase, in order to avoid magnetic reversal [8]. Theoretical calculations suggest further that if the size of the soft phase is twice the domain wall width of the hard phase, both phases would be perfectly exchange coupled [9]. The domain wall width of the L10 FePt phase is estimated to be about 3.8 nm [3]. This means that the soft magnetic regions in exchange-spring magnets involving L10 FePt should be less than or equal to 8 nm. The grain size of polycrystalline thin films is a key factor also for other technologically important properties such as mechanical strength, electrical resistivity, etc.

⁎ Corresponding author. Tel.: +49 234 32 29 109. E-mail address: [email protected] (N. Zotov). 0040-6090/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2008.06.062

Several studies reported data for the grain growth of FePt thin films with as-deposited face-centred cubic (fcc) structure. Ristau et al. [10] determined the grain growth kinetics of 10 nm thick fcc Fe50Pt50 thin films annealed at TA = 973 K, using transmission electron microscopy (TEM). The initial average grain size of the fcc phase was 6 nm. A rapid growth stage was observed to occur within the first 5–10 min, followed by a much slower kinetics during the isothermal annealing [10]. Morgan et al. [11] observed in-situ grain growth using TEM for both nonirradiated and irradiated 25 nm thick fcc Fe50Pt50 thin films. A grain growth exponent n = 3.2 ± 0.04 was determined for the irradiated samples using a grain growth equation modified for irradiation. Insitu TEM observations [12] suggested that ordering of fcc Fe50Pt50 thin films into the L10 structure may have a dependence on the grain size. Caution should be taken with these data, however, because a phase transformation from the as-deposited fcc phase to the L10 phase occurs concurrent with grain growth during annealing in all these studies. Experimental data for the nature of grain growth in pure L10 FePt thin films as well as correlations between grain sizes, microstrains and order parameter of the L10 FePt phase are very scarce. Sun et al. [13] have reported that as-deposited polycrystalline Fe52Pt48 films with fcc structure transform above 623 K into the L10 FePt phase and the grain size of 100 nm thick L10 thin films along the [111] direction increases from 12 nm at TA = 673 K to about 21 nm at TA = 873 K (annealing time 60 min). In-situ X-ray diffraction (XRD) experiments [14] suggested that the growth of L10 FePt at 548 K is reaction controlled. Correspondingly, the aim of this work was to study in more detail the annealing temperature dependence of the grain size and

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microstrains of polycrystalline FePt thin films containing only the L10 FePt phase. The polycrystalline nature of the annealed FePt thin films allowed determination of the grain growth along different crystallographic directions. We have used Fe/Pt multilayers (MLs) as precursors, rather than co-deposited fcc FePt thin films, because the MLs preparation route leads to a decrease of the temperature necessary for the formation of the L10 FePt phase [15–24].

computational simplicity, each XRD line was fitted using 2 pseudoVoigt (pV) functions, in order to describe the Kα1 and Kα2 components of each reflection, plus a linear background. It was shown that the maximum deviation between a Voigtian and its pV representation is only about 0.8% at peak maximum (see Ref. [26] and references therein). The pV function is defined as a linear combination of Cauchy and Gauss functions:

2. Experimental details

pVðxÞ ¼ I0 ½η C ðxÞ þ ð1− ηÞGðxÞ

2.1. Sample preparation

where I0 is the peak intensity, x = (2Θ − 2Θ0) / w, 2Θ0 is the position of the peak maximum, w is the full-width at half maximum (FWHM), C(x) = (1 + x2)− 1 is the Cauchy component, G(x) = exp[−(ln2)x2] is the Gauss component and η is the mixing parameter. The standard deviations of the refined parameters I0, 2Θ0, w and η for each reflection were calculated from least-squares fits. The integral width β (= peak area / peak height, in radians) of a pV profile is equal to [26]:

The [(Pt 1.7 nm/Fe 0.9 nm)5 Fe 0.23 nm]8 MLs were fabricated by magnetron sputtering (von Ardenne CS 730 S) on thermally oxidised 6″ Si (100) wafers using Ar sputtering gas (pAr ~ 0.6 Pa) without additional heating of the substrates. The MLs were deposited by repeatedly moving the substrate past two stationary Fe and Pt elemental targets (99.6% purity). The Fe direct current power density was 5.6 × 10− 2 W/cm2, the Pt radio frequency power density was 23.9 × 10− 2 W/cm2, the Fe deposition rate was about 0.12 nm/s and the Pt deposition rate was about 0.22 nm/s. The thickness (0.23 nm) of the additional Fe layers was selected so that the overall composition is close to Fe50Pt50. The actual composition Fe52Pt48 was determined by energy dispersive X-ray spectroscopy (Leo Supra S55 field emission gun scanning electron microscope operated at 20 keV, Si:Li detector, INCA software, 300 s/measuring point). A detailed characterization of the as-deposited MLs using XRD, X-ray reflectivity and diffuse scattering as well as TEM is given in Ref. [24]. The total thickness 90 ± 6 nm, measured with a surface profilometer (KLA Tencor P-10), is in good agreement with the total thickness 98 nm, determined from XRD simulations of the MLs [24]. Diced samples (5 mm × 5 mm size) were annealed in vacuum (typically in the range 2.0 × 10− 5–2.0 × 10− 7 mbar) at five different annealing temperatures TA = 623, 673, 773, 873 and 973 K using rapid thermal annealing machine (Cryotech) working with halogen lamps. The accuracy of TA was ±10 K. The desired sample temperature, measured separately with a K-type thermocouple, was reached in about 5–6 min. The annealing time was 60 min in order to allow complete transformation of the Fe/Pt MLs into the L10 phase. The initial cooling rates after switching off the heater were about 0.2 K/s at 623 K to 0.4 K/s at 973 K. 2.2. X-ray diffraction measurements All XRD measurements were done on a Bruker D8 Discovery powder diffractometer using symmetric (Bragg–Brentano) scattering geometry with the scattering vector Q = 4πsin(Θ) / λ perpendicular to the film surface, where 2Θ is the scattering angle and λ is the X-ray wavelength. The diffractometer was equipped with a 1/4 Eulerian cradle, two Göbel mirrors (giving a parallel incident X-ray beam both in the equatorial and axial directions with a cross-section of about 1 mm) and flat area detector ‘HI-STAR’ using Cu Kα radiation (λ = 0.15417 nm) from a fine-focused X-ray tube (operated at 40 kV and 40 mA). The requirements for using a parallel beam diffractometer for the analysis of microstructure of materials were recently stressed by Welzel and Mittemeijer [25]. The measuring time was 10 min per frame. The two-dimensional patterns were integrated along the Debye rings and rebinned with a step Δ2Θ = 0.05°. 2.3. Size-strain analysis It is well established by both theory and experiment that in many cases the structurally and instrumentally broadened diffraction profiles from constant-wavelength XRD measurements can be well approximated by Voigt functions. A Voigt function (Voigtian) is the convolution of Cauchy and Gauss functions. In the present study, for

h i 1= β ¼ w 1=2 π η þ 1=2ð1−ηÞðπ=ln2Þ 2

ð1Þ

ð2Þ

The integral widths of the Cauchy and Gauss components of the measured profiles (h), βhC and βhG respectively, can be obtained employing rational approximate expressions: βhC ¼ βh fC ðηÞ

ð3Þ

βhG ¼ βh fG ðηÞ

ð4Þ

where fC (η) and fG (η) are 3rd-order polynomials of η [27]. The integral widths βgC and βgG of the instrumental broadening (g) can be calculated in the same way. In this study, these parameters were determined from measurements on National Institute of Standards and Technology (USA) corundum standard reference material 1976. The integral widths of the true structurally broadened profiles f can be calculated then by:  1=2 βfG ¼ β2hG −β2gG

ð5Þ

and βfC ¼ βhC −βgC

ð6Þ

Assuming further that the Cauchy component of the structurally broadened profiles f arises from small grain size and that the Gauss component is due to microstrains, the grain size Dhkl and the microstrain ehkl for a single diffraction line with indices (hkl) can be obtained accordingly [28–29]: Dhkl ¼ λ=βfC cosðΘhkl Þ

ð7Þ

and ehkl ¼ 1=4βfG ctgðΘhkl Þ

ð8Þ

The grain sizes Dhkl determined from Eq. (7) refer to a volume average of columns lengths perpendicular to the film surface. The errors in Dhkl and ehkl were calculated using full error-propagation analysis and Eqs. (2)–(8). More accurate and sophisticated methods (e.g. Fourier methods) for analysis of diffraction line-broadening exist (a critical review of such methods is given, for example, in Ref. [30– 31]), but they exceed the goals of this study to get an initial estimate of the line broadening and to follow changes in the size, microstrain and stacking fault parameters with annealing temperature. 3. Results The diffraction patterns from strongly-textured crystalline MLs (also called superlattices) with scattering vector Q perpendicular to

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Fig. 1. X-ray diffraction pattern of an as-deposited Fe/Pt multilayers (full circles), profile fit (full line) and difference plot (+).

the multilayer's surface are characterized by Bragg peaks of the average structure, centred at Q00l = (2π / do)l, l = 0, 1, 2,...(where do is the average interplanar distance along the growth direction) and satellites around them, arising from the compositional modulation of the MLs. For MLs without variations of the interplanar spacings and/or the superlattice period Λ, the positions of the satellites are at Qm = Q00l ± (2π / Λ)m, m = 1, 2,.... (for more details on scattering from superlattices, see for example [32]). The as-deposited MLs show a strong (001) Bragg peak at about 2.85 Å− 1. This peak is denoted as the 0-satellite in Fig. 1. In addition, a relatively weak (001)− satellite is observed at Q− 1 ~ 2.59 Å− 1. The (001)+ satellite at Q+ 1 ~ 3.11 Å− 1 as well as higher-order satellites Qm (m = ±2, ±3, ...) are not observed. The presence of interdiffusion in the asdeposited MLs smears out the compositional profile of the superlattice and greatly reduces the intensity of these satellites (for more details see Ref. [24]). Numerical simulations of the XRD patterns of metallic MLs with a Gaussian distribution of grain sizes along the growth direction [33] revealed that both the FWHM and the background signal decrease with increasing average grain size. The Monte Carlo simulation algorithm used in Ref. [33] would give qualitatively similar results for other (non-Gaussian) grain-size distribution functions as well. Experimentally, there are only few TEM cross-sectional studies of grain-size distributions in as-deposited metallic superlattices. Junhua

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Fig. 3. Profile fits of the (311), (113) and (222) reflections of samples annealed at TA = 673, 773 and 973 K: experimental points (full circles), fit (full lines). The patterns at TA = 773 and 973 K are shifted along the y-axis for clarity.

[34] reported that the distribution of nano-sized grains in [Pd 1.0 nm/ CoCr18B4 0.3 nm] MLs can be described well by a Gaussian function. Although the actual grain-size distribution in the investigated Fe/Pt MLs may not be exactly Gaussian, the above studies [33,34] justify the use of integral width line-broadening methods for a semi-quantitative estimation of grain sizes in the MLs. A fit of the (001)− and (001)0 satellites using pV functions is also given in Fig. 1. The goodness-of-fit factor (GOF) is 4.0%. Annealing of the MLs at 573 K does not lead to significant changes (see below). In-situ XRD of similar Fe/Pt MLs [14] also showed that up to 534 K there is almost no line broadening of the MLs. The coercivity of the as-deposited MLs, measured with a Lake Shore vibrating sample magnetometer (VSM) is HC = 12 ± 4 Oe. Annealing of the MLs for 60 min at TA = 623 K already leads to the complete transformation of the MLs into the tetragonal (L10) FePt phase — see Fig. 2. The coercivity of this sample, measured with VSM, is about 10 kOe [23]. A significant increase of Hc (Hc ≥ 5 kOe) of Fe/Pt MLs annealed at TA N 573–623 K was observed also in Ref. [19,20,21,35]. For easier comparison with previous studies, the Bragg peaks of the tetragonal L10 phase are indexed in Fig. 2 using the equivalent fcc unit cell (atet = acub√2 / 2, ctet = acub). At all annealing temperatures TA = 623, 673, 773, 873 and 973 K, the samples are polycrystalline with [111] texture and 15 reflections were measured in the range 20–90° 2Θ, which were then fitted with pV functions. An example of the profile fitting of the (311), (113) and (222) reflections is given in Fig. 3 for three different annealing temperatures. The average GOF and the mixing parameters bηN, averaged over all fitted reflections, are given in Table 1. Good profile fits were achieved for all Bragg peaks (bGOFN between 3.0 and 4.5%). At all TA a decrease of the out-of-plane microstrains ehkl was observed with increasing grain size Dhkl (see Fig. 4), except for TA = 673 K, where ehkl is approximately constant as a function of Dhkl. In the case of Gaussian strains, the microstrain parameter ehkl is proportional to the root mean square (r.m.s.) strain bε 2oN1/2 (see Ref. [31] and references therein). The r.m.s. strain bε 2oN1/2 itself is proportional to b(ΔD)2N1/2 /D,

Table 1 Average goodness-of-fit factors bGOF)N and averaged mixing parameters bηN

Fig. 2. X-ray powder diffraction pattern of a sample annealed at TA = 623 K.

T (K)

bGOFN, %

bηN

623 673 773 873 973

3.0 2.9 4.4 4.3 2.9

0.6(3) 0.7(3) 0.6(3) 0.4(1) 0.4(3)

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Fig. 4. Microstrains ehkl of FePt thin film (TA = 623 K) as a function of the grain size Dhkl.

α-Fe2O3 [37], ball-milled MnCO3, CdCO3 and PbCO3 powders [38] as well as SrBi2Nb2O9 epitaxial thin films [39]. On the contrary, Depero et al. [40] observed an 1 / bDN2 dependence of the microstrains in TiO2 nanopowders. The temperature dependencies of the grain sizes averaged over the different crystallographic directions bDN and the average microstrains beN are given in Fig. 5(a) and (b), respectively. The D and e parameters for the MLs are also included in the plots. Annealing of the MLs up to about 573 K does not lead to significant changes in the average grain size of the MLs, but does induce significant mircostrains. As soon as the L10 FePt phase is formed, the out-of-plane microstrains in the thin films drop substantially and remain practically constant, independent of TA. On the contrary, bDN of the L10 FePt phase increases systematically with TA. The bDN values are smaller than the total thickness of the films (90 ± 6 nm), suggesting that the grain growth is not limited by the thin film thickness. Similar TA behaviour was observed for grain growth in 100 nm thick Fe52Pt48 thin films along the [111] direction (Ref. [13]), as well as in 10 nm thick CoPt thin films [10] and for Cu grain growth in immiscible Cu/Cr MLs [41]. 4. Discussion

where b(ΔD)2N1/2 is the r.m.s. deviation of the average grain size (hkl subscripts are skipped for clarity). Therefore, the inverse relation between ehkl and Dhkl (see Fig. 4) indicates that the r.m.s deviations b(ΔD)2N1/2 for most of the investigated Fe52Pt48 thin films are, in first approximation, independent of the grain size. Similar inverse correlation between r.m.s. strains and grain sizes was observed for ZrO2 [36],

4.1. Anisotropic grain growth The grain sizes Dhkl of grains with different crystallographic orientations, perpendicular to the film surface, exhibit broad distributions at all TA. This is the reason for the relatively large errors of bDN in Fig. 5(a) and suggests the presence of strong anisotropic sizebroadening. Indeed, the grain sizes along the [110] and/or [310] directions are larger than along the other directions. Minimum grain sizes are observed either along [221] or along the [200] directions, except for TA = 623 K, where D has a minimum along the [001] direction. Anisotropic grain growth was also observed in Mo–Cu thin films between 573 K and 773 K [42]. On the contrary, the microstrains are almost isotropically distributed. Only the microstrains along the [001] direction (e001 ~ 0.4– 0.6%) are slightly larger than the average microstrains beN ~ 0.2%. Evidently, there is an increased density of lattice imperfections (defects) along the [001] direction. The L10 FePt structure consists of alternating Fe and Pt atomic layers along the [001] direction (for more details on the structure of L10 phases see Ref. [43]). That is why every structural imperfection, which leads to a local disruption of the Fe–Pt ordering, will affect the microstrains along the [001] direction more than along other directions. To semi-quantitatively compare the growth rates, linear fits of Dhkl as a function of TA were made, although the temperature dependence of bDN is actually more complex (see below). The fit results are given in Table 2. The grains along the [221] direction show the fastest growth rate, the grains along the [001] direction show the slowest growth rate. This leads also to changes in the texture of the samples with increasing TA. Table 2 Growth rates for grains with different crystallographic orientation

Fig. 5. (a) Average grain sizes bDN and (b) average microstrains beN as a function of annealing temperature. Empty circles – Fe/Pt MLs, full triangles – L10 FePt thin films. The lines are only a guide for the eye.

(hkl)

dDhkl/dT (nm/K)

001 110 111 200 002 201 112 202 221 311 113

0.06(2) 0.08(4) 0.08(3) 0.07(2) 0.13(2) 0.07(3) 0.12(1) 0.12(3) 0.17(6) 0.14(5) 0.13(3)

Standard deviations are given in parentheses.

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grain growth in thin films (see for a review Ref. [46]) and in bulk materials (see for a review Ref. [51]). In bulk materials the most common grain growth law (normal growth) is diffusion-controlled and can be written in the form: bDN2 −D20 ¼ tc0 expð−HGG =kB T Þ

ð9Þ

where D0 is the average grain size at time t = 0, HGG is the activation energy of grain growth, kB is the Boltzmann constant and c0 is a constant. In thin films two empirical grain-growth models are commonly used: bDN ¼ D0 þ ct m

ð10Þ

or more generally  1=n bDN ¼ Dn0 þ ct

ð11Þ

where again Fig. 6. Comparison of experimental (full squares) and calculated (thick line) average grain sizes bDN as a function of annealing temperature.

Several effects can lead to anisotropic grain growth: anisotropic (hkl-dependent) surface free energies, anisotropic strain energy, presence of anisotropic interdiffusion in L10 FePt and/or anisotropic elastic grain interactions. Indeed, Kushida et al. [44] have observed that the activation energies for interdiffusion in bulk L10 FePt along the [001] and the [100] directions are different (3.80 and 3.17 eV, respectively). Kim et al. [45] reported that in 25 nm thick Fe55Pt45 thin films, large transformation microstrains as well as in-plane tensile strains induced strong (001) texture and selective growth of (001) grains at the early stages (annealing times b 100 s) of annealing at 823 K. If all the grains are uniformly strained with equibiaxial in-plane strain ε||, then the strainenergy component ΔW of the free energy for two neighbouring grains with orientations (001) and (111) will be [46,47] : ΔWε =W001 −W111 = (M001 −M111)ε2||, where Mhkl is the biaxial modulus for a grain with texture (hkl). On the other hand, assuming that the grains are uniformly stressed with equibiaxial in-plane stress σ, then ΔWσ = W001 − W111 = (1 / M001 − 1 / M111)σ2. If the Poissons's ratio ν of the material is 1/3, then the in-plane strain ε|| = −ε⊥ [48]. Recent ab-initio calculations [49] have shown that ν of the L10 FePt phase (ν = 0.31) is very close to ν = 1/3. Therefore, the calculated out-of-plane microstrains can be used to estimate the inplane strains. In order for the (001) grains to grow faster than the (111) grains, the free energy ΔG = ΔWε and/or ΔG = ΔWσ must be negative. The biaxial moduli M001 and M111 for the L10 FePt phase were calculated using the approach given in Ref. [47] and the recently determined elastic constants for L10 FePt (Ref. [49]); M001 is equal to 494 GPa and M111 is equal to 425 GPa. These values show qualitatively that: (1) for small in-plane strains (as in the present case, where | ε|| | = | ε⊥ | ~ beN ~ 0.2%) ΔG = ΔWε N 0 and the (111) texture is favoured; (2) in the case of large in-plane stresses and strains (as most probably is the case reported in Ref. [45]) ΔG = ΔWσ b 0 and (001) texture could be formed, although minimization of the surface and interface free energy in fcc metals and fcc-related alloys favour the formation of (111) texture [46]. Computer simulations of strain energy versus surface energy effects on grain growth in thin films also show that growth of grains with (001) texture can develop only in highly strained thin films [50]. 4.2. Grain growth model Fig. 5(a) shows that annealing of the MLs at different temperatures TA ≥ 623 K leads to nucleation and subsequent growth of the L10 FePt phase. Different models have been proposed in the literature for the

c ¼ c0 expð−HGG =kB T Þ

ð12Þ

If m = n = 1, then Eqs. (10) and (11) correspond to reactioncontrolled (linear) growth. In contrast with bulk materials, significant variations in m (n) have been reported in nanocrystalline and thin film materials. Values of n between 2 and 4 were reported for nanocrystalline Cu [52]. Both parabolic (m = n = 2) and linear grain growth have been observed in Au–In thin films [53]. Huang et al. [54] have obtained a grain growth exponent n = 2.3 ± 0.2 by fitting the measured in-plane grain sizes of PdIn thin films as a function of time to Eq. (11) at three temperatures between 773 K and 823 K. Admon et al. [41] used m = 0.5 in Eq. (10) for the analysis of out-of-plane Cu grain growth in Cu–Cr MLs. According to Eq. (11) the growth exponent n can be determined from a least-squares (LS) fit of bDN data as a function of time at fixed temperature (in this case c is a constant) or from LS fit of bDN as a function of temperature at fixed annealing time t. We have used the second approach in the present study. Best fit is obtained for n = 3 ± 0.5, D0 ~ 16.7 nm, c0 = 1.47 × 106 nm3/min and HGG = 0.51 ± 0.07 eV (see Fig. 6). The value D0 ~ 16.7 nm is close to the average grain size of the MLs along the normal to the film surface (16.4 ± 0.3 nm) — see Fig. 5(a). We have also fitted the grain size versus time data of Ristau et al. [10] for Fe50Pt50 thin films at TA = 973 K using Eq. (11); n = 3.5 ± 0.3 and c = 1694 ± 160 nm3.5/min gave the best fit. This growth exponent is, within error limits, equal to the result obtained in the present study. The experimentally established growth exponent n = 3 ± 0.5 can be indicative of several grain growth processes: pore-controlled lattice diffusion, boundary-controlled coalescence of the L10 FePt phase in the Fe/Pt MLs matrix by lattice diffusion, boundary-controlled impurity drag, etc (for more details on these mechanisms, see Ref. [51] and references therein). The value of the activation energy for grain growth HGG = 0.51 ± 0.07 eV is similar to the values 0.56 eV [54] and 0.86 eV [41] for PdIn and CuCr thin films, respectively. The value HGG = 0.51 ± 0.07 eV is much lower than the activation energy for interdiffusion in Fe/Pt MLs (1.7 ± 0.6 eV) [14] as well as in bulk L10 FePt at high temperatures (3.17–3.80 eV) [44] and fairly consistent with available grainboundary diffusion data. For example, the activation energies for grain-boundary diffusion (HGB) in Cu/Ni MLs is HGB ~ 0.31 eV [55] and in Ta/NiFe MLs HGB = 0.55 ± 0.11 eV [56]. 4.3. Stacking faults and order parameter According to the theory of Warren [57], planar defects (stacking faults) affect the diffraction profiles of fcc metals (as well as fcc-related alloys) only if h + k + l = 3 m. This means that the (200) and (002) reflections should be broader than the (111) peak, if there are stacking faults along the [111] direction. We do observe that the integral

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Fig. 7. Dependence of the total stacking fault density on the annealing temperature TA. The insert shows the order parameter S as a function of TA. The dashed line is only a guide for the eye.

breadths β200 N β022 N β111 for TA N 623 K, which suggests the presence of stacking faults in the annealed Fe52Pt48 samples. {111} stacking and twin faults were observed by TEM in both FePt thin films [58–59] and nanoparticles [60]. Stacking faults are often responsible for Cauchy [η = 1 in Eq. (1)] diffraction line shapes [61]. That is why the decrease of the average mixing parameter bηN with TA (see Table 1) may indicate a lowering of the stacking fault density with increasing TA. The total stacking fault density (TSFD) 1.5(α' + α'') + β was calculated as follows [62–63]: 1:5ðα Vþ αWÞ þ β ¼ 1:7677a0 ðD111 −D200 Þ=D111 D200

ð13Þ

where α' is the intrinsic, α'' is the extrinsic, β is the twin fault stacking density and a0 is the corresponding fcc lattice parameter. Fig. 7 shows that indeed the TSFD decreases with increasing TA. At the same time the increase of TA leads to a rapid increase of the order parameter S of the L10 FePt phase, as can be seen from the inset in Fig. 7. S was calculated from the lattice parameters using the expression [64]: h i S ¼ ð1−c=aÞ= 1−ðc=aÞf 2

ð14Þ

where (c/a)f is the ratio of the lattice parameters for fully ordered L10 L10 FePt. The lattice parameters were determined from Rietveld refinements using the FULLPROF® program. The order parameter is already quite large for annealing at 623 K (S = 0.95 ± 0.01). This indicates that the transformation of the MLs is very rapid and the observed broadening of the XRD lines is mainly due to grain size effects. A similar increase of S with TA was also observed by other authors [19,21,65,66]. It should be noted, however, that the order parameter of the annealed samples in the present study is higher than the S values reported in Ref. [19,21] for annealed [Fe 2.5 nm/Pt 2.5 nm]10 MLs, although the deposition rates were similar and the annealing time was the same in all three cases (60 min). Most probably the as-deposited Fe/Pt MLs in Ref. [19,21] do not have 50:50 composition, although the nominal thicknesses are equal. This could explain the reported lower S values in Ref. [19,21], because the order parameter of off-stoichiometric Fe50 − xPt50 + x alloys decreases for x ≠ 0.0. 5. Conclusions The microstructural evolution of magnetron sputtered [(Pt 1.7 nm/ Fe 0.9 nm)5 Fe 0.23 nm]8 multilayers annealed at different temperatures TA for 60 min was investigated using XRD and integral width linebroadening analysis. The MLs transform above 573 K directly into polycrystalline thin films with L10 structure and (111) texture. The L10

thin films are highly ordered already at TA = 623 K (S = 0.95± 0.01). The increase of S correlates with a decrease of the total stacking fault density. The investigated Fe52Pt48 thin films show almost isotropic (hklindependent) microstrain broadening. Only the microstrains along the [001] direction are larger than the average microstrains beN, most probably due to increased density of anti-site defects along that direction. The values of beN remain small (~ 0.2%) and practically constant as a function of TA, which suggests that strain energy factors do not play significant role during grain growth in the present study. On the contrary, the investigated Fe52Pt48 thin films show both anisotropic size line-broadening and anisotropic grain growth. At most of the annealing temperatures the grains along the [110] and/or the [310] directions have the largest grain sizes. Evidently, orientationdependent grain-boundary and surface energy factors lead to specific textures during grain growth of the L10 FePt thin films. The grains along the [221] direction show the fastest, while the grains along the [001] direction, show the slowest growth rates. The temperature dependence of the grain size bDN, averaged over different crystallographic directions, can be tentatively modelled with grain-growth law bDN3 − bD0N3 = tc0exp(− HGG / kBTA) and apparent activation energy HGG = 0.51 ± 0.07 eV. Additional cross-sectional TEM studies would be necessary to distinguish between different mechanisms, which might lead to the observed grain-growth exponent n = 3 ± 0.5 and will be the subject of a future study. The calculated activation energy is about 1/6 of that for bulk interdiffusion, which might indicate that grain growth in the investigated Fe52Pt48 thin films is controlled mainly by grain-boundary diffusion. Acknowledgement The authors thank A. Savan for the careful reading of the manuscript. References [1] O.A. Ivanov, L.V. Solina, V.A. Demshina, L.M. Magat, Phys. Met. Metallogr. (Phiz. Metal. Metalloved.) 35 (1973) 92. [2] K.R. Coffey, M.A. Parker, J.K. Howard, IEEE Trans. Magn. 31 (1995) 2737. [3] D. Weller, A. Moser, L. Folks, M.E. Best, W. Lee, M.F. Toney, M. Schwickert, J.-U. Thiele, M.F. Doerner, IEEE Trans. Magn. 36 (2000) 10. [4] J. Yu, U. Ruediger, A.D. Kent, R.F.C. Farrow, R.F. Marks, D. Weller, L. Folks, S.S.P. Parkin, J. Appl. Phys. 87 (2000) 6854. [5] K. Shibata, Mater. Trans. 44 (2003) 1542. [6] J.P. Liu, C.P. Luo, Y. Liu, D.J. Sellmyer, Appl. Phys. Lett. 72 (1998) 483. [7] X. Rui, J.E. Shield, Z. Sun, L. Yue, Y. Xu, D.J. Sellmyer, Z. Liu, D.J. Miller, J. Magn. Magn. Mater. 305 (2006) 76. [8] R. Skomski, J.M.D. Coey, IEEE Tran. Magn. 29 (1993) 2860. [9] E.F. Kneller, R. Hawig, IEEE Trans. Magn. 27 (1991) 3588. [10] R.A. Ristau, K. Barmak, K.R. Coffey, J.K. Howard, J. Mater. Res. 14 (1999) 3263. [11] N.W. Morgan, R.C. Birtcher, G.B. Thompson, J. Appl. Phys. 100 (2006) 124904. [12] Y.K. Takahashi, M. Ohnuma, K. Hono, J. Appl. Phys. 93 (2003) 7580. [13] A.C. Sun, P.C. Kuo, S.C. Chen, C.Y. Chuo, H.L. Huang, J.H. Hsu, J. Appl. Phys. 95 (2004) 7264. [14] N. Zotov, J. Feydt, A. Savan, A. Ludwig, J. Appl. Phys. 100 (2006) 073517. [15] B.J. Daniels, W.D. Nix, B.M. Clemens, Mater. Res. Soc. Symp. Proc. 343 (1994) 549. [16] J.P. Liu, C.P. Luo, Y. Liu, D.J. Sellmyer, Appl. Phys. Lett. 72 (1998) 483. [17] T.C. Hufnagel, M.C. Kautzky, B.J. Daniels, B.M. Clemens, J. Appl. Phys. 85 (1999) 2609. [18] C.P. Luo, D.J. Sellmyer, Appl. Phys. Lett. 75 (1999) 3162. [19] Y. Endo, N. Kikuchi, O. Kitakami, Y. Shimada, J. Appl. Phys. 89 (2001) 7065. [20] M.L. Yan, N. Powers, D.J. Sellmyer, J. Appl. Phys. 93 (2003) 8292. [21] Y. Endo, K. Oikawa, T. Miyazaki, O. Kitakami, Y. Shimada, J. Appl. Phys. 94 (2003) 7222. [22] Y.F. Ding, J.S. Chen, E. Liu, Thin Solid Films 474 (2005) 141. [23] A. Ludwig, N. Zotov, A. Savan, S. Groudeva-Zotova, Appl. Surf. Sci. 252 (2006) 2518. [24] N. Zotov, J. Feydt, T. Walther, A. Ludwig, Appl. Surf. Sci. 253 (2006) 128. [25] U. Welzel, E.J. Mittemeijer, Powder Diffr. 20 (2005) 376. [26] S. Enzo, G. Fagherazzi, A. Benedetti, S. Polizzi, J. Appl. Crystallogr. 21 (1988) 536. [27] R. Delhez, Th.H. de Keijser, E.J. Mittemeijer, Fresenius Z. Anal. Chem. 312 (1982) 1. [28] P. Scherrer, Nachr. Ges. Wiss. Gött. 2 (1918) 98. [29] A.R. Stokes, A.J.C. Wilson, Proc. Phys. Soc. 56 (1944) 174. [30] P. Scardi, M. Leoni, R. Delhez, J. Appl. Crystallogr. 37 (2004) 381. [31] I. Lucks, P. Lamparter, E.J. Mittemeijer, J. Appl. Crystallogr. 37 (2004) 300.

N. Zotov et al. / Thin Solid Films 517 (2008) 531–537 [32] Y. Fujii, in: T. Shinjo, T. Takada (Eds.), Metallic Superlattices, Elsevier, Amsterdam, 1987. [33] Z. Mitura, P. Mikolajczak, J. Phys. F. Met. Phys. 18 (1988) 183. [34] W. Junhua, Nanotechnology 13 (2002) 720. [35] H. Ito, K. Kusunoki, H. Saito, S. Ishio, J. Magn. Magn. Mater. 272 –276 (2004) 2180. [36] N.H. Brett, M. Gonzalez, J. Bouillot, J.C. Niepce, J. Mater. Sci. 19 (1984) 1349. [37] J. Morales, J.L. Tirado, M. Macias, Solid State Chem. 53 (1984) 303. [38] J.M. Criado, M. Gonzalez, C. Real, J. Mater. Sci. Lett. 5 (1986) 467. [39] A. Boulle, C. Legrand, R. Guinebretiere, J.P. Mercurio, A. Dauger, Thin Solid Films 391 (2001) 42. [40] L.E. Depero, L. Sangaletti, A. Allieri, E. Bontempi, A. Marino, M. Zocchi, J. Cryst. Growth 198-199 (1999) 516. [41] U. Admon, I. Dahan, M.P. Dariel, G. Kimmel, J. Sariel, A. Shtechman, B. Yahav, L. Zevin, D.S. Lashmore, Thin Solid Films 251 (1994) 105. [42] G. Ramanath, H.Z. Xiao, L.C. Yang, A. Rockett, L.H. Allen, J. Appl. Phys. 78 (1995) 2435. [43] D.E. Laughin, K. Srinivasan, M. Tanase, L. Wang, Scr. Mater. 53 (2005) 383. [44] A. Kushida, K. Tanaka, H. Numakura, Mater. Trans. 44 (2003) 59. [45] J.-S. Kim, Y.-M. Koo, N. Shin, J. Appl. Phys. 100 (2006) 093909. [46] C.V. Thompson, Annu. Rev. Mater. Sci. 20 (1990) 245. [47] F. Huang, M.L. Weaver, J. Appl. Phys. 98 (2005) 073505. [48] C. Noyan, J.B. Cohen, Residual Sterss Measurement by Diffraction and Interpretation, Spinger-Verlag, Berlin, 1986.

[49] [50] [51] [52] [53] [54] [55] [56] [57] [58] [59] [60] [61] [62] [63] [64] [65] [66]

537

N. Zotov, A. Ludwig, Intermetallics 16 (2008) 113. R. Carel, C.V. Thompson, H.J. Frost, Acta Mater. 44 (1996) 2475. H.V. Atkinson, Acta Metall. 36 (1988) 469. S.K. Ganapthi, D.M. Owen, A.H. Chokshi, Scr. Metall. Mater. 25 (1991) 2699. M. Millares, B. Pieraggi, E. Lelievre, Solid State Ionics 63–65 (1993) 575. M. Huang, Y. Wang, Y.A. Chang, Thin Solid Films 449 (2004) 113. R. Venos, W. Palmer, H. Hoffman, Thin Solid Films, 162 (1988) 155. I. Hashim, H.A. Atwater, K.T.Y. Kung, R.M. Valletta, J. Appl. Phys. 74 (1993) 458. B.E. Warren, Prog. Met. Phys. 8 (1959) 147. M.-G. Kim, S.-C. Shin, K. Kang, Appl. Phys. Lett. 80 (2002) 3802. G. Safran, T. Suzuki, K. Ouchi, P.B. Barna, G. Radnoszi, Thin Solid Films 496 (2006) 580. B. Bian, D.E. Laughin, K. Sato, Y. Hirotsu, J. Appl. Phys. 87 (2000) 6962. A.J.C. Wilson, X-ray Optics, 2nd ed.Methuen, London, 1962. S.K. Chatterjee, S.K. Halder, S.P. Sen Gupta, J. Appl. Phys. 48 (1977) 1442. F.-E. Teng, B. Yang, Y. Wang, Metall. Trans. A, 23 (1992) 2859. B.W. Robarts, Acta Metall. 2 (1954) 597. A.C. Sun, S.C. Chen, P.C. Kuo, C.Y. Chou, Y.H. Fang, J.-H. Hsu, H.L. Huang, H.W. Chang, IEEE Trans. Magn. 41 (2005) 3772. Y.K. Takahashi, K. Hono, Scr. Mater. 53 (2005) 403.