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Epitaxial growth and magnetic properties of MnxPt1−x (x⩽0.38) thin films on (0 0 1)-SrTiO3 substrates
Journal of Magnetism and Magnetic Materials 205 (1999) 151}160
Epitaxial growth and magnetic properties of Mn Pt V \V (x)0.38) thin "lms on (0 0 1)-SrTiO substrates
A. Borgschulte *, D. Menzel , T. Widmer , H. Bremers, U. Barkow , J. Schoenes Institut fu( r Halbleiterphysik und Optik, Technische Universita( t Braunschweig, Mendelssohnstr. 3, D-38106 Braunschweig, Germany Institut fu( r Metallphysik und Nukleare Festko( rperphysik, Technische Universita( t Braunschweig, Mendelssohnstr. 3, D-38106 Braunschweig, Germany Received 28 April 1999
Abstract Thin epitaxial (0 0 1)-oriented Mn Pt "lms with 0)x(0.38 were grown by molecular beam epitaxy on (0 0 1)V \V SrTiO substrates. The structure and the growth mechanism have been investigated. RHEED patterns con"rm the high structural quality of the Mn Pt "lms. A Volmer}Weber model is applied to describe the growth mechanism of MnPt V \V on SrTiO . X-ray di!raction measurements reveal a lattice parameter for MnPt of 3.89(5) As and a strong dependence of the lattice parameter on the Mn content. A possible strain-induced phase transition at higher Mn content is discussed. At room temperature a high saturation magnetization (M ) of 680 emu/cm is found which implies that a highly ordered 1 phase of MnPt occurs. The Curie temperature is enhanced compared to non-strained "lms. 1999 Elsevier Science B.V. All rights reserved. PACS: 75.50.Ss; 81.15.Hi; 75.70.Ak; 78.20.Ls Keywords: Thin "lms; Epitaxy; Magnetic properties; Magneto-optics
1. Introduction The realisation of magneto-optical memories relies on a magnetic medium with, among other important properties, a substantial polar magnetooptical Kerr e!ect. In addition, a high stability of the magnetic domains in the material is required to achieve a long-life readout performance. In order to gain a better understanding of the magnetic and magneto-optical e!ects in transition metal com-
* Corresponding author. Fax: #49-531-391-5155. E-mail address: [email protected] (A. Borgschulte)
pounds containing Pt we studied Mn Pt alloys. V \V They are known to be chemically stable over a wide temperature range. The appeal to study this system is also related to its rather simple crystallographic and electronic structure. This makes it suitable for the application of standard models to interpret the magnetic and magneto-optical e!ects [1,2]. In 1983, Buschow et al. [3] investigated the magneto-optical e!ects of bulk MnPt . The surface of the samples was mechanically polished and the rotation was determined for the laser wavelengths 666 and 830 nm. MnPt showed zero rotation. Since a measurement for two single photon energies may miss rotation peaks, somewhat later
0304-8853/99/$ - see front matter 1999 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 9 9 ) 0 0 4 7 6 - X
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BraK ndle et al. [4] prepared by a similar method bulk samples and measured the full Kerr rotation and ellipticity spectra between 0.6 and 5 eV. They found two extrema with a maximum rotation of 0.133 and noted that the signal was strongly dependent on the treatment of the surface, explaining the results of Buschow et al. Obviously, the mechanical polishing destroys the crystalline structure near the surface of this material in a much more dramatic way than in some other materials. To remove the damaged surface, VergoK hl and Schoenes [5] used a chemical etch after the mechanical polishing. Although the surface of these samples was less shiny than before, the chemical etching led to a much better reproducibility of the results and to an increase of the maximum rotation to 0.553. Recently, Kato et al. [6] have measured the magneto-optical Kerr e!ect of polycrystalline MnPt and found a Kerr rotation h of !1.183 at 1.2 eV. Somewhat ) later, Wierman and co-workers [7] reported a h of ) !2.53 at 95 K and 1.4 eV for magnetron sputtered "lms. This rotation, however, was measured through a protective SiO layer. Thus, the later V data do not represent the intrinsic Kerr rotation of MnPt which is de"ned as the rotation at the interface vacuum/sample. The data have to be scaled down by a factor which, in "rst order, is given by the refractive index of the SiO which V ranges from 1.5 to 1.9. With the extreme surface sensitivity of the magneto-optical Kerr e!ect in MnPt it appeared parti cularly suitable to investigate a sample with as perfect a surface as possible. To this goal we have prepared single-crystalline "lms by molecular beam epitaxy (MBE) and studied the surface and bulk properties of these "lms by re#ection high-energy electron di!raction (RHEED) and X-ray di!raction (XRD) analysis, respectively. While the preparation of (1 1 1)-oriented MnPt "lms has been reported recently by Kokuryu et al. [8] we present here the preparation of (0 0 1)-oriented MnPt "lms on (0 0 1)-SrTiO substrates. The growth mechanism of highly ordered Mn Pt thin "lms on SrTiO V \V (0 0 1) has been systematically studied. Lange et al. [9] have investigated annealed CoPt and MnPt samples and found drastic changes in the Kerr rotation with the quality of the surfaces. As the Kerr rotation is connected with the magnetization
it appears useful to study the magnetic properties of the thin epitaxial "lms. Thus, a magnetic characterization is presented, showing higher ordering temperatures of our "lms. The magnetic data is compared with the Kerr rotation. A detailed optical and magneto-optical analysis will be discussed in a forthcoming paper.
2. Experimental Mn Pt alloy "lms were grown by coevaporaV \V tion from a manganese high-temperature e!usion cell and a platinum e-beam source by molecular beam epitaxy (MBE) onto (0 0 1)-oriented SrTiO substrates. The pressure during deposition was less than 1;10\ Pa. The total deposition rate of 0.6 nm/min was controlled by a beam #ux monitor and a quartz-crystal oscillator. The uncertainty of the growth rate measured by the quartz oscillator was about K10%. We prepared "lms with Mn concentrations ranging from 0 to 38 at%. The "lms were deposited on SrTiO (0 0 1) substrates heated to 500}7003C. Subsequently the "lms were cooled down slowly to room temperature. The structure of the "lms was characterized in situ by re#ecting high-energy di!raction (RHEED) and ex situ by X-ray di!ractometry (XRD). The "lm composition was determined by X-ray #uorescence analysis (EDX). The magnetic properties of the "lms were measured using a SQUID magnetometer over a temperature range of 10}300 K. The Kerr rotation h was determined ) at room temperature and in a magnetic "eld of 1.4 T using a polarized angle-modulation method.
3. Results and discussion 3.1. Sample preparation Like Cu and Au, Mn and Pt are nearly full miscible. Ordered phases occur for the atomic ratios 1 : 1 and 1 : 3. In the former case the structure is tetragonal (AuCu structure) whereas in the latter case it is cubic (AuCu structure). Stochiometric MnPt has a lattice constant of 3.89 As [10,11] quite similar to the noble metal platinum. Recently, it has
A. Borgschulte et al. / Journal of Magnetism and Magnetic Materials 205 (1999) 151}160
been reported that (0 0 1)-oriented "lms of platinum can be grown epitaxially on SrTiO (0 0 1) by metalorganic chemical vapour deposition [12]. We have attempted to prepare MnPt alloy "lms by the MBE technique and found that highly ordered "lms with AuCu crystal structure can indeed be grown at a temperature of approximately 7003C. In order to describe the correlation between the magnetic anisotropy and the structural properties, a detailed knowledge of the structure and growth mechanism is mandatory. The growth of platinum on insulator crystals has been investigated by several authors [12}14]. Raaijmakers et al. developed a simple model of crystal growth of non-wetting systems [15]. Cillessen et al. [14] used this model to describe the growth of Pt on MgO crystals. In our
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paper, the same model is applied to the analysis of the growth mechanism of MnPt on SrTiO (0 0 1). The RHEED patterns con"rm an island growth in the initial growing stage. The patterns show facetted streaks, which can be explained by the crystalline geometry of the islands. Fig. 1 shows a typical RHEED pattern of the initial growing stage and its geometrical interpretation. The di!erent angles of the streaks in di!erent directions indicate the crystallographic surface planes. The mainly (0 0 1)-oriented Mn Pt islands have a (0 0 1) top V \V facet and (1 1 0) and (1 1 1) facets at the edges. The island morphology is related to the equilibrium shape derived by the Wul! construction. We have estimated the surface energies for the di!erent surface planes using the nearest-neighbour
Fig. 1. Typical RHEED pattern and its geometrical interpretation at the initial growing stage. In the [1 0 0] direction the streak angles are 453, the [1 1 0] direction forms streaks with an angle of about 553. The island morphology is related to the equilibrium shape derived by the Wul! construction. We assume that c 'c .
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broken-bond model [16]. The model predicts that the surface energy of a (1 1 1) plane c is smaller than c , so we expect smaller (1 1 0) surfaces. At early stages of growth, some islands with (1 1 0) orientation are formed (see RHEED pattern in Fig. 1). These (1 1 0)-oriented islands disappear at a thickness of about 5 nm.
The nucleation process of Mn Pt is similar V \V to the behaviour of Pd on MgO as described in a recent paper by Henry et al. [17]. The island formation (Fig. 2a) is followed by a partial coalescence of the islands (Fig. 2b). At a mean critical thickness the "lm becomes coherent and the islands are rapidly transformed into a smooth Mn Pt V \V
Fig. 2. Mechanism of growth of Mn Pt on SrTiO . RHEED patterns reveal an island growth in the initial state of growth (a), V \V followed by a coalescence phase (b). At a critical thickness of K40 nm the islands transform into a smooth "lm (c) resulting in a facetting behaviour at large "lm thicknesses (d).
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"lm, which can be seen from the appearance of Laue circles in the RHEED pattern (Fig. 2c). For a Mn Pt layer deposited at a substrate temperV \V ature of 7003C on SrTiO (0 0 1), the critical thick ness was found to be of the order of 40 nm. The observation of the RHEED patterns con"rms that the Mn Pt (xK0.25) layers are grown on the V \V SrTiO (0 0 1) substrates epitaxially as Mn Pt V \V [1 0 0], (0 0 1)#SrTiO [1 0 0], (0 0 1). At a thick ness much larger than the critical thickness, the "lms begin to form (1 1 1) features (Fig. 2d) followed by a polycrystalline growth for d'60 nm. Using the model of Raaijmakers et al. [15], a wetting angle can be calculated for MnPt deposited on SrTiO (0 0 1) at 7003C. From our RHEED ob servation, no reasonable dependence of the critical thickness on the Mn concentration is indicated. A surface energy for Pt of 1.8 Jm\ and a di!usion constant D"1.5;10\ ms\ was obtained from the empirical relation between the FCC metal-melting temperature (here Pt) and the di!usion constants [18]. The deposition rate used is 10\ ms\. In this way, a wetting angle of about 1003 was calculated which, indeed, is indicative of a non-wetting system. 3.2. X-ray characterization X-ray di!raction measurements of the Mn Pt V \V "lms are consistent with the RHEED analysis. Fig. 3 shows the X-ray di!raction pattern of a 50 nm Mn Pt "lm using CuK radiation a (j"1.541 As ). The measured lattice parameter perpendicular to the "lm plane of (3.90$0.005) As for Mn Pt is in reasonable agreement with the value known from bulk samples [11]. In addition, the lattice parameter depends on the Mn concentration [11]. The formation of the Mn Pt orV \V dered phase can be shown from the existence of the (0 0 l)-superlattice peaks (l"2n#1, n3N). The very small intensities of the di!raction peaks ((1%) of the (1 1 0) orientation con"rm the high perfection of the layers. At this stadium, no (1 1 1) re#ection can be found. Di!raction patterns of "lms of a thickness of 110 nm show (1 1 1) peaks with an intensity of approximately 1}5% of the (0 0 2) peak. This e!ect appears to be more pronounced with increasing "lm thickness. The width
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Fig. 3. X-ray di!raction pattern of a 50 nm MnPt "lm. A long range order parameter of gK1 is derived from the ratio of the intensities I /I . The inset shows a simulation of the (0 0 2) peak using the kinematic di!raction formalism.
of the (0 0 2) di!raction peak for a 50 nm MnPt "lm is in quite good agreement with the broadening of the peaks by the "nite "lm thickness. We have simulated the 2h X-ray di!raction pro"le to the measured (0 0 2) peak using the kinematical di!raction formalism [19]: sin(N(p/j)2a sin h) IJ . sin((p/j)2a sin h)
(1)
The parameters used were the number of atomic planes N and the lattice spacing a. The good agreement with the measurement indicates a small roughness of the MnPt layer (see inset of Fig. 3). However, for the more frequent case of heteroepitaxial "lm growth, the lattice mismatch and the thermal expansion mismatch between the "lm and the substrate materials cause considerable microstructural strain in the "lms. As a result, the properties can be quite di!erent from the intrinsic properties of the corresponding bulk crystalline materials [20]. In order to determine the strain of the epitaxial Mn Pt "lms we measured the outV \V of-plane lattice constant by the above described normal h}2h scan of polycrystalline and epitaxial "lms on quartz and SrTiO substrates, respectively. Fig. 4a shows the out-of-plane lattice parameter of polycrystalline and epitaxial "lms compared to bulk data by Raub et al. [11]. In the range from 0 to
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The thermal expansion coe$cient is a V \VK + . a K1;10\K\ [10], so the thermal expan12 - sion mismatch can be neglected. The lattice mismatch of the system Mn Pt /SrTiO depends V \V strongly on the Mn-content. The cubic SrTiO (perovskite structure) has a lattice constant of 3.905 As which leads to a mismatch for MnPt of #0.4%. Below 20 at% Mn, a negative lattice mismatch occurs, which leads to a tensile strained "lm perpendicular to the surface. Above this content, a positive mismatch occurs, which leads to a compressively strained "lm perpendicular to the surface. E.g. the lattice constant a "3.9231 As of Pt leads to a mismatch of !0.46%. The measured out-of-plane lattice parameter of a 1000 As thick Pt "lm on SrTiO (0 0 1) is a "3.93(5) As . Using , a poisson ratio l"0.37 for Pt the in-plane lattice parameter can be calculated in case of a tetragonal distortion by the formula [21]. 2l e e "! , 1!l ,
Fig. 4. Perpendicular lattice parameters of bulk [11], epitaxial and polycrystalline "lms of Mn Pt (a). At x+38 at% Mn V \V a phase transition from the cubic AuCu structure to the face centered tetragonal AuCu structure occurs. The linear relationship between the Mn content and the lattice parameter of the unstrained cubic phase (Vegard's law) is indicated (dashed line). The perpendicular lattice parameter which is assumed for a fully strained epilayer on SrTiO is also shown (dotted line). (b) shows the face-centered tetragonal MnPt cells. The sites in the planes A and B are occupied by Mn and Pt atoms, respectively. The lattice changes into a cubic phase when the planes A and B are "lled statistically (c).
38 at% Mn a linear relationship between the Mncontent of the bulk or polycrystalline alloys and the lattice constant is observed (Vegards law). At larger Mn fractions ('38 at%) the alloy crystallizes in a face-centered tetragonal AuCu structure [11].
(2)
with the biaxial stress e "e "e and e "e . VV WW , XX , The calculated in-plane constant is a "3.91(3) As , indicating a strained epilayer. For the stochiometric compound MnPt we found a lattice constant a "3.89(5) As as expected for a slightly strained , layer. At higher Mn contents the out-of-plane lattice parameter of epitaxial "lms decreases more than expected for an elastically strained epilayer. To explain this discrepancy we developed a model which describes a change of the atomic distribution in the lattice induced by the strain. We consider that the di!erent lattice planes A and B are not "lled statistically with Mn atoms (Fig. 4b and c) which would be the case for the bulk alloy. The prefered "lling of Mn atoms of the planes A (ordered AuCu structure) is found for more than 38 at% Mn. Below this concentration, this behaviour could occur under the in#uence of the strain. Thus, the perpendicular lattice constant is smaller than the in-plane constant because Mn is smaller than Pt. Therefore, the strain induces a continuous phase transition from the cubic to the tetragonal phase. This e!ect is an additional indication that strain may change thermodynamical quantities [22].
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157
Fig. 6. The Curie temperature (¹ ) increases with the Mn con! tent x. The ¹ of epitaxial "lms is enhanced in comparison to ! polycrystalline "lms [25] and bulk material [10].
Fig. 5. Hysteresis loops of Mn Pt "lms with di!erent Mn V \V content x at room temperature (a). The saturation magnetization M is strongly dependent on the Mn content x. The 1 maximum of the magnetization occurs at 25 at% Mn. The curve of the Kerr rotation (*) is very similar to the dependence of M 1 (䢇) on the Mn content (b).
3.3. Magnetic properties Mn Pt is known to be a strong ferromagnet V \V [10]. Pickart et al. found a saturation magnetization M of MnPt bulk samples of (4.12$0.08)l 1 per unit cell at 77 K by neutron di!raction measurements [23]. At the same temperature Antonini et al. measured a magnetic moment of (4.04$0.08)l using ballistic methods and (4.42 $0.17)l performing polarized neutron measurements [24]. By using a SQUID magnetometer, we "nd a saturation magnetization of epitaxial MnPt "lms of 4.76 and 4.31l at 77 and 300 K, respectively. Fig. 5a shows SQUID hysteresis measurements of the in plane magnetization of our epitaxial thin Mn Pt "lms with di!erent Mn concentrations. V \V As one can see in Fig. 5b, the saturation magneti-
zation (z) depends strongly on the Mn content. Below the percolation limit of about 15 at% Mn, no magnetic ordering is found in Mn Pt , beV \V cause Mn has too few nearest and next-nearest Mn neighbours to achieve magnetic ordering. Magnetic ordering occurs at a Mn content of more than 15 at%, and M increases up to the stochiometric 1 MnPt alloy. In the ordered phase, the Mn atoms occupy the corner sites of the FCC cubic cell, whereas the Pt atoms are found on the facecentered sites. In this con"guration, the largest ferromagnetic coupling of the Mn atoms occurs. If the Mn content grows any further, Mn atoms start to occupy Pt sites and the magnetization decreases. The reason is obviously due to the fact, that the magnetic interaction between Mn atoms sitting on a Mn site and the nearest Pt site is antiferromagnetic. The Curie temperature (¹ ) of 375 K for bulk ! MnPt is well established [25]. Our measurements on epitaxial thin "lms reveal an enhanced Curie temperature of 500 K. Fig. 6 shows the Curie temperatures for bulk samples, magnetron sputtered thin "lms and epitaxial thin "lms as function of the Mn content [25,10]. While the bulk samples and sputtered "lms show a rather constant positive slope of ¹ with increasing Mn concentration, the ! epitaxial "lms behave quite di!erently. At a Mn content of +20 at% the Curie temperature grows
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Fig. 7. In plane (straight line) and out of plane (dashed and dotted lines) magnetization of MnPt "lms with various thick nesses (room temperature). Even in the thinnest "lm (17.5 nm) the easy axis of magnetization lies in the "lm plane.
Fig. 8. Experimental Kerr spectra of MnPt "lms at room temperature [8] and theoretical Kerr spectrum at ¹"0 K [26] in comparison to our epitaxial thin "lm data at room temperature.
rapidly with decreasing slope, when the concentration increases. This can be understood, if the strain of the "lms is considered. At a Mn content of 20 at%, the lattice parameters of Mn Pt and the V \V SrTiO substrate are identical. Therefore, the Curie temperatures of bulk samples and epitaxial "lms are similar. Beyond 20 at% Mn, a lattice mismatch occurs, which leads to a tensile strained "lm parallel to the surface and a compressively strained "lm perpendicular to the surface. Due to the changing Mn distances, the magnetic ordering is in#uenced, so that the exchange interaction increases or decreases depending on the direction. This leads to an enhanced Curie temperature because the larger exchange determines the ordering temperature. The ¹ curve of the epitaxial "lms comes closer to the ! curve for the bulk samples and the magnetron sputtered "lms at a Mn content of about 33 at%. MnPt crystallizes in the cubic AuCu structure. Therefore, no favored direction of magnetization along any of the cubic axes is expected. In a thin "lm, however, the easy axis of magnetization lies in the "lm plane due to the shape anisotropy. This prefered direction does not change, at least not down to a "lm thickness of 17.5 nm (Fig. 7). An easy axis of magnetization perpendicular to the "lm plane can be achieved, if thin Co layers are used as spacer layers between 1.9 nm thick MnPt layers [8].
3.4. Magneto-optical properties In Fig. 8 a theoretical Kerr spectrum (LDA calculations) [2,26] and two experimental Kerr spectra measured by di!erent authors [8,25] are compared with our results for an epitaxial (0 0 1)-oriented MnPt "lm. At "rst, we compare the theoretical data with our experimental data. The shape of the two spectra are quite similar to each other. The di!erent maximum values of the two spectra can be attributed in part to the reduction of the sample magnetization at room temperature (see Section 3.3). The theoretical Kerr spectrum is calculated for zero temperature. If the Kerr spectra are measured at room temperature, where the magnetization is reduced, then the overall size of the measured Kerr rotation will be reduced, too (see below). The reason for the shift of the theoretical spectrum to lower energies remains debatable. It is known that LDA calculations fail to describe properly exchange and correlation e!ects, underestimating these e!ects for occupied states and overestimating these e!ects for unoccupied states. The best agreement of the experimental spectra can be found between the (1 1 1)-grown MnPt "lms [8] and the present results from (0 0 1)-grown "lms (Fig. 9). The magnitude of the peak at about 1.1 eV increases nearly linearly with the magnetization of
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159
Fig. 9. Kerr spectra at room temperature of epitaxially grown Mn Pt "lms with di!erent Mn concentrations. The maxV \V imum of the Kerr rotation occurs at a Mn content of 25 at% (see Fig. 5b).
Fig. 10. Room temperature Kerr rotation of a MnPt "lm annealed at di!erent temperatures. The Kerr rotation was measured through a protective SiO layer at a magnetic "eld of 1 T V and a wavelengh of 1200 nm.
the samples (Fig. 5b). This e!ect was predicted in the work of Oppeneer et al. [26] who studied the in#uence of the exchange splitting on the Kerr rotation h of XPt compounds (X"V, Cr, Mn, ) Fe, and Co) and found an increase of h with ) increasing exchange splitting. This dependence can also explain why the Kerr rotation of our MnPt "lms is larger than e.g. that of Kato et al. [6]. We "nd a large maximum value of !0.823 around 1.1 eV and at least one additional peak of h K0.23 ) at approximately 2.6 eV. The sign of h is reversed ) at 2.2 eV. To our knowledge, this is the largest intrinsic Kerr rotation found at room temperature in MnPt , although Wierman et al. [7] quote larger values. The obvious reason is that we present the intrinsic Kerr signal at the interface MnPt /air, whereas the data from Ref. [7] are for the interface MnPt /SiO . Thus, the later data do not represent V the Kerr rotation of MnPt which is de"ned as the rotation at the interface vacuum/sample, but the data have to be scaled down by a factor which, in "rst order, is given by the refractive index of the SiO . In addition, interference e!ects may have to V be taken into account. In thin "lms even without a protective layer interference e!ects may occur if the "lm thickness is in the order of the optical penetration depth. To exclude this possibility we measured the Kerr rotation of "lms with di!erent thicknesses between 200 and 1000 As . No signi"cant di!erences of the rotation were found.
To show the dependence of the Kerr e!ect on the long-range order parameter, we grew a MnPt thin "lm at a temperature of about 853C on SrTiO substrates and coated it with a SiO layer. At this temperature, the "lm does not grow epitaxially, but with a long-range order parameter of nearly zero as determined by X-ray di!raction analysis. The "lm was heated for half an hour at di!erent temperatures and then cooled down to room temperature. Fig. 10 shows the increase of the absolute value of the Kerr rotation measured at room temperature at a wavelength of 1200 nm with increasing annealing temperature. The intrinsic Kerr rotation has to be scaled down by a factor of 1.5}1.9 due to the refractive index of the protective SiO layer on this V sample. After annealing at 8003C the maximum Kerr rotation is obtained. As already pointed out by Wierman and Kirby [25] this corroborates the large long-range order parameter of the epitaxially grown "lms.
4. Conclusions Thin epitaxial Mn Pt "lms have been preV \V pared by molecular beam epitaxy. The growth mechanism can be described by a Volmer}Weber model. The high structural quality of the "lms has been con"rmed by RHEED and X-ray di!raction. Magnetic characterization of the "lms with
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a SQUID magnetometer reveals a strong dependence of the saturation magnetization on the Mn content x. The maximum magnetization is obtained for the stochiometric highly ordered compound MnPt . In contrast to bulk material, the Curie temperature ¹ is enhanced and this is at! tributed to strain. The high structural and magnetic ordering is re#ected in a large room temperature Kerr rotation of !0.823 at 1.2 eV. We shows that the magnetization and the Kerr rotation is dependent on the Mn content and on the long-range order parameter. Acknowledgements We acknowledge the help of the Institut fuK r Halbleitertechnik (TU Braunschweig) for the EDX measurements. References [1] E. Kren, G. Kadar, L. Pal, J. Solyom, P. Szabo, T. Tarnoczi, Phys. Rev. 171 (1968) 574. [2] P.M. Oppeneer, V.N. Antonov, T. Kraft, H. Eschrig, A.N. Yaresko, A.Y. Perlov, Solid State Comm. 94 (1995) 255. [3] K.H.J. Buschow, P.G. van Engen, R. Jongebreur, J. Magn. Magn. Mater. 38 (1983) 1. [4] H. BraK ndle, J. Schoenes, F. Hulliger, W. Reim, unpublished. Data given by J. Schoenes in: R.W. Cahn, P. Haasen, E.J. Kramer (Eds.), Materials Science and Technology, Vol. 3, Verlag Chemie, Weinheim, 1992, p. 147. [5] M. VergoK hl, J. Schoenes, J. Magn. Soc. Jpn. 20 (1996) 141. [6] T. Kato, H. Kikuzawa, S. Iwata, S. Tsunashima, S. Uchiyama, J. Magn. Magn. Mater. 140}144 (1995) 713.
[7] K.W. Wierman, J.N. Hil"ker, R.F. Sabiryanov, S.S. Jaswal, R.D. Kirby, J.A. Woollam, Phys. Rev. B 55 (1997) 3093. [8] M. Kokuryu, T. Kato, S. Iwata, S. Tsunashima, J. Appl. Phys. 81 (1997) 4779. [9] R.J. Lange, S.J. Lee, D.W. Lynch, P.C. Can"eld, B.N. Harmon, S. Zollner, Phys. Rev. B 58 (1998) 351. [10] M. AuwaK rter, A. Ku{mann, Ann. Phys. 6 (1950) 169. [11] E. Raub, W. Mahler, Z. Metallk. 4 (1955) 282. [12] B.S. Kwak, P.N. First, A. Erbil, B.J. Wilkens, J.D. Budai, M.F. Chisholm, L.A. Boatner, J. Appl. Phys. 72 (1992) 3735. [13] B.M. Lairson, M.R. Visokay, R. Sinclair, S. Hagstrom, B.M. Clemens, Appl. Phys. Lett. 61 (1992) 1390. [14] J.F.M. Cillessen, R.M. Wolf, D.M. de Leeuw, Thin Solid Films 226 (1993) 53. [15] I.J.M.M. Raaijmakers, R.A.A. Hack, A.G. Dirks, Suppl. Vide, Couches Mines, Proceedings of the International Symposium on Trends and New Applications in Thin Films, 1987. [16] J.M. Howe, Interfaces in Materials, Wiley, New York, 1997. [17] C.R. Henry, C. Chapton, C. Duriez, S. Giorgio, Surf. Sci. 253 (1991) 177. [18] N.A. Gjostein, in: J.J. Burke, M.L. Reed, V. Weiss (Eds.), Surfaces and Interfaces I, Syracuse University Press, New York, 1967, p. 271. [19] B.E. Warren, X-ray di!raction, Addison-Wesley, Reading, MA, 1969. [20] S. Jin, T.H. Tiefel, M. McCormack, H.M. O'Bryan, L.H. Chen, R. Ramesh, D. Schurig, Appl. Phys. Lett. 67 (1995) 557. [21] J.Y. Tsao, Materials fundamentals of molecular beam epitaxy, Academic Press, San Diego, 1993. [22] R.A. Rao, Q. Gan, C.B. Eom, R.J. Cava, Y. Suzuki, J.J. Krajewski, S.C. Gausepohl, M. Lee, Appl. Phys. Lett. 70 (1997) 3035. [23] S.J. Pickart, R. Nathans, J. Appl. Phys. 33 (1962) 1336. [24] B. Antonini, F. Lucari, F. Menzinger, A. Paoletti, Phys. Rev. 187 (1969) 611. [25] K.W. Wierman, R.D. Kirby, J. Magn. Magn. Mater. 154 (1996) 12. [26] P.M. Oppeneer, V.N. Antonov, H. Eschrig, A.N. Yaresko, A.Y. Perlov, J. Phys.: Condens. Matter 8 (1996) 5769.
Epitaxial growth and magnetic properties of Mn Pt V \V (x)0.38) thin "lms on (0 0 1)-SrTiO substrates
A. Borgschulte *, D. Menzel , T. Widmer , H. Bremers, U. Barkow , J. Schoenes Institut fu( r Halbleiterphysik und Optik, Technische Universita( t Braunschweig, Mendelssohnstr. 3, D-38106 Braunschweig, Germany Institut fu( r Metallphysik und Nukleare Festko( rperphysik, Technische Universita( t Braunschweig, Mendelssohnstr. 3, D-38106 Braunschweig, Germany Received 28 April 1999
Abstract Thin epitaxial (0 0 1)-oriented Mn Pt "lms with 0)x(0.38 were grown by molecular beam epitaxy on (0 0 1)V \V SrTiO substrates. The structure and the growth mechanism have been investigated. RHEED patterns con"rm the high structural quality of the Mn Pt "lms. A Volmer}Weber model is applied to describe the growth mechanism of MnPt V \V on SrTiO . X-ray di!raction measurements reveal a lattice parameter for MnPt of 3.89(5) As and a strong dependence of the lattice parameter on the Mn content. A possible strain-induced phase transition at higher Mn content is discussed. At room temperature a high saturation magnetization (M ) of 680 emu/cm is found which implies that a highly ordered 1 phase of MnPt occurs. The Curie temperature is enhanced compared to non-strained "lms. 1999 Elsevier Science B.V. All rights reserved. PACS: 75.50.Ss; 81.15.Hi; 75.70.Ak; 78.20.Ls Keywords: Thin "lms; Epitaxy; Magnetic properties; Magneto-optics
1. Introduction The realisation of magneto-optical memories relies on a magnetic medium with, among other important properties, a substantial polar magnetooptical Kerr e!ect. In addition, a high stability of the magnetic domains in the material is required to achieve a long-life readout performance. In order to gain a better understanding of the magnetic and magneto-optical e!ects in transition metal com-
* Corresponding author. Fax: #49-531-391-5155. E-mail address: [email protected] (A. Borgschulte)
pounds containing Pt we studied Mn Pt alloys. V \V They are known to be chemically stable over a wide temperature range. The appeal to study this system is also related to its rather simple crystallographic and electronic structure. This makes it suitable for the application of standard models to interpret the magnetic and magneto-optical e!ects [1,2]. In 1983, Buschow et al. [3] investigated the magneto-optical e!ects of bulk MnPt . The surface of the samples was mechanically polished and the rotation was determined for the laser wavelengths 666 and 830 nm. MnPt showed zero rotation. Since a measurement for two single photon energies may miss rotation peaks, somewhat later
0304-8853/99/$ - see front matter 1999 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 9 9 ) 0 0 4 7 6 - X
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BraK ndle et al. [4] prepared by a similar method bulk samples and measured the full Kerr rotation and ellipticity spectra between 0.6 and 5 eV. They found two extrema with a maximum rotation of 0.133 and noted that the signal was strongly dependent on the treatment of the surface, explaining the results of Buschow et al. Obviously, the mechanical polishing destroys the crystalline structure near the surface of this material in a much more dramatic way than in some other materials. To remove the damaged surface, VergoK hl and Schoenes [5] used a chemical etch after the mechanical polishing. Although the surface of these samples was less shiny than before, the chemical etching led to a much better reproducibility of the results and to an increase of the maximum rotation to 0.553. Recently, Kato et al. [6] have measured the magneto-optical Kerr e!ect of polycrystalline MnPt and found a Kerr rotation h of !1.183 at 1.2 eV. Somewhat ) later, Wierman and co-workers [7] reported a h of ) !2.53 at 95 K and 1.4 eV for magnetron sputtered "lms. This rotation, however, was measured through a protective SiO layer. Thus, the later V data do not represent the intrinsic Kerr rotation of MnPt which is de"ned as the rotation at the interface vacuum/sample. The data have to be scaled down by a factor which, in "rst order, is given by the refractive index of the SiO which V ranges from 1.5 to 1.9. With the extreme surface sensitivity of the magneto-optical Kerr e!ect in MnPt it appeared parti cularly suitable to investigate a sample with as perfect a surface as possible. To this goal we have prepared single-crystalline "lms by molecular beam epitaxy (MBE) and studied the surface and bulk properties of these "lms by re#ection high-energy electron di!raction (RHEED) and X-ray di!raction (XRD) analysis, respectively. While the preparation of (1 1 1)-oriented MnPt "lms has been reported recently by Kokuryu et al. [8] we present here the preparation of (0 0 1)-oriented MnPt "lms on (0 0 1)-SrTiO substrates. The growth mechanism of highly ordered Mn Pt thin "lms on SrTiO V \V (0 0 1) has been systematically studied. Lange et al. [9] have investigated annealed CoPt and MnPt samples and found drastic changes in the Kerr rotation with the quality of the surfaces. As the Kerr rotation is connected with the magnetization
it appears useful to study the magnetic properties of the thin epitaxial "lms. Thus, a magnetic characterization is presented, showing higher ordering temperatures of our "lms. The magnetic data is compared with the Kerr rotation. A detailed optical and magneto-optical analysis will be discussed in a forthcoming paper.
2. Experimental Mn Pt alloy "lms were grown by coevaporaV \V tion from a manganese high-temperature e!usion cell and a platinum e-beam source by molecular beam epitaxy (MBE) onto (0 0 1)-oriented SrTiO substrates. The pressure during deposition was less than 1;10\ Pa. The total deposition rate of 0.6 nm/min was controlled by a beam #ux monitor and a quartz-crystal oscillator. The uncertainty of the growth rate measured by the quartz oscillator was about K10%. We prepared "lms with Mn concentrations ranging from 0 to 38 at%. The "lms were deposited on SrTiO (0 0 1) substrates heated to 500}7003C. Subsequently the "lms were cooled down slowly to room temperature. The structure of the "lms was characterized in situ by re#ecting high-energy di!raction (RHEED) and ex situ by X-ray di!ractometry (XRD). The "lm composition was determined by X-ray #uorescence analysis (EDX). The magnetic properties of the "lms were measured using a SQUID magnetometer over a temperature range of 10}300 K. The Kerr rotation h was determined ) at room temperature and in a magnetic "eld of 1.4 T using a polarized angle-modulation method.
3. Results and discussion 3.1. Sample preparation Like Cu and Au, Mn and Pt are nearly full miscible. Ordered phases occur for the atomic ratios 1 : 1 and 1 : 3. In the former case the structure is tetragonal (AuCu structure) whereas in the latter case it is cubic (AuCu structure). Stochiometric MnPt has a lattice constant of 3.89 As [10,11] quite similar to the noble metal platinum. Recently, it has
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been reported that (0 0 1)-oriented "lms of platinum can be grown epitaxially on SrTiO (0 0 1) by metalorganic chemical vapour deposition [12]. We have attempted to prepare MnPt alloy "lms by the MBE technique and found that highly ordered "lms with AuCu crystal structure can indeed be grown at a temperature of approximately 7003C. In order to describe the correlation between the magnetic anisotropy and the structural properties, a detailed knowledge of the structure and growth mechanism is mandatory. The growth of platinum on insulator crystals has been investigated by several authors [12}14]. Raaijmakers et al. developed a simple model of crystal growth of non-wetting systems [15]. Cillessen et al. [14] used this model to describe the growth of Pt on MgO crystals. In our
153
paper, the same model is applied to the analysis of the growth mechanism of MnPt on SrTiO (0 0 1). The RHEED patterns con"rm an island growth in the initial growing stage. The patterns show facetted streaks, which can be explained by the crystalline geometry of the islands. Fig. 1 shows a typical RHEED pattern of the initial growing stage and its geometrical interpretation. The di!erent angles of the streaks in di!erent directions indicate the crystallographic surface planes. The mainly (0 0 1)-oriented Mn Pt islands have a (0 0 1) top V \V facet and (1 1 0) and (1 1 1) facets at the edges. The island morphology is related to the equilibrium shape derived by the Wul! construction. We have estimated the surface energies for the di!erent surface planes using the nearest-neighbour
Fig. 1. Typical RHEED pattern and its geometrical interpretation at the initial growing stage. In the [1 0 0] direction the streak angles are 453, the [1 1 0] direction forms streaks with an angle of about 553. The island morphology is related to the equilibrium shape derived by the Wul! construction. We assume that c 'c .
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broken-bond model [16]. The model predicts that the surface energy of a (1 1 1) plane c is smaller than c , so we expect smaller (1 1 0) surfaces. At early stages of growth, some islands with (1 1 0) orientation are formed (see RHEED pattern in Fig. 1). These (1 1 0)-oriented islands disappear at a thickness of about 5 nm.
The nucleation process of Mn Pt is similar V \V to the behaviour of Pd on MgO as described in a recent paper by Henry et al. [17]. The island formation (Fig. 2a) is followed by a partial coalescence of the islands (Fig. 2b). At a mean critical thickness the "lm becomes coherent and the islands are rapidly transformed into a smooth Mn Pt V \V
Fig. 2. Mechanism of growth of Mn Pt on SrTiO . RHEED patterns reveal an island growth in the initial state of growth (a), V \V followed by a coalescence phase (b). At a critical thickness of K40 nm the islands transform into a smooth "lm (c) resulting in a facetting behaviour at large "lm thicknesses (d).
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"lm, which can be seen from the appearance of Laue circles in the RHEED pattern (Fig. 2c). For a Mn Pt layer deposited at a substrate temperV \V ature of 7003C on SrTiO (0 0 1), the critical thick ness was found to be of the order of 40 nm. The observation of the RHEED patterns con"rms that the Mn Pt (xK0.25) layers are grown on the V \V SrTiO (0 0 1) substrates epitaxially as Mn Pt V \V [1 0 0], (0 0 1)#SrTiO [1 0 0], (0 0 1). At a thick ness much larger than the critical thickness, the "lms begin to form (1 1 1) features (Fig. 2d) followed by a polycrystalline growth for d'60 nm. Using the model of Raaijmakers et al. [15], a wetting angle can be calculated for MnPt deposited on SrTiO (0 0 1) at 7003C. From our RHEED ob servation, no reasonable dependence of the critical thickness on the Mn concentration is indicated. A surface energy for Pt of 1.8 Jm\ and a di!usion constant D"1.5;10\ ms\ was obtained from the empirical relation between the FCC metal-melting temperature (here Pt) and the di!usion constants [18]. The deposition rate used is 10\ ms\. In this way, a wetting angle of about 1003 was calculated which, indeed, is indicative of a non-wetting system. 3.2. X-ray characterization X-ray di!raction measurements of the Mn Pt V \V "lms are consistent with the RHEED analysis. Fig. 3 shows the X-ray di!raction pattern of a 50 nm Mn Pt "lm using CuK radiation a (j"1.541 As ). The measured lattice parameter perpendicular to the "lm plane of (3.90$0.005) As for Mn Pt is in reasonable agreement with the value known from bulk samples [11]. In addition, the lattice parameter depends on the Mn concentration [11]. The formation of the Mn Pt orV \V dered phase can be shown from the existence of the (0 0 l)-superlattice peaks (l"2n#1, n3N). The very small intensities of the di!raction peaks ((1%) of the (1 1 0) orientation con"rm the high perfection of the layers. At this stadium, no (1 1 1) re#ection can be found. Di!raction patterns of "lms of a thickness of 110 nm show (1 1 1) peaks with an intensity of approximately 1}5% of the (0 0 2) peak. This e!ect appears to be more pronounced with increasing "lm thickness. The width
155
Fig. 3. X-ray di!raction pattern of a 50 nm MnPt "lm. A long range order parameter of gK1 is derived from the ratio of the intensities I /I . The inset shows a simulation of the (0 0 2) peak using the kinematic di!raction formalism.
of the (0 0 2) di!raction peak for a 50 nm MnPt "lm is in quite good agreement with the broadening of the peaks by the "nite "lm thickness. We have simulated the 2h X-ray di!raction pro"le to the measured (0 0 2) peak using the kinematical di!raction formalism [19]: sin(N(p/j)2a sin h) IJ . sin((p/j)2a sin h)
(1)
The parameters used were the number of atomic planes N and the lattice spacing a. The good agreement with the measurement indicates a small roughness of the MnPt layer (see inset of Fig. 3). However, for the more frequent case of heteroepitaxial "lm growth, the lattice mismatch and the thermal expansion mismatch between the "lm and the substrate materials cause considerable microstructural strain in the "lms. As a result, the properties can be quite di!erent from the intrinsic properties of the corresponding bulk crystalline materials [20]. In order to determine the strain of the epitaxial Mn Pt "lms we measured the outV \V of-plane lattice constant by the above described normal h}2h scan of polycrystalline and epitaxial "lms on quartz and SrTiO substrates, respectively. Fig. 4a shows the out-of-plane lattice parameter of polycrystalline and epitaxial "lms compared to bulk data by Raub et al. [11]. In the range from 0 to
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The thermal expansion coe$cient is a V \VK + . a K1;10\K\ [10], so the thermal expan12 - sion mismatch can be neglected. The lattice mismatch of the system Mn Pt /SrTiO depends V \V strongly on the Mn-content. The cubic SrTiO (perovskite structure) has a lattice constant of 3.905 As which leads to a mismatch for MnPt of #0.4%. Below 20 at% Mn, a negative lattice mismatch occurs, which leads to a tensile strained "lm perpendicular to the surface. Above this content, a positive mismatch occurs, which leads to a compressively strained "lm perpendicular to the surface. E.g. the lattice constant a "3.9231 As of Pt leads to a mismatch of !0.46%. The measured out-of-plane lattice parameter of a 1000 As thick Pt "lm on SrTiO (0 0 1) is a "3.93(5) As . Using , a poisson ratio l"0.37 for Pt the in-plane lattice parameter can be calculated in case of a tetragonal distortion by the formula [21]. 2l e e "! , 1!l ,
Fig. 4. Perpendicular lattice parameters of bulk [11], epitaxial and polycrystalline "lms of Mn Pt (a). At x+38 at% Mn V \V a phase transition from the cubic AuCu structure to the face centered tetragonal AuCu structure occurs. The linear relationship between the Mn content and the lattice parameter of the unstrained cubic phase (Vegard's law) is indicated (dashed line). The perpendicular lattice parameter which is assumed for a fully strained epilayer on SrTiO is also shown (dotted line). (b) shows the face-centered tetragonal MnPt cells. The sites in the planes A and B are occupied by Mn and Pt atoms, respectively. The lattice changes into a cubic phase when the planes A and B are "lled statistically (c).
38 at% Mn a linear relationship between the Mncontent of the bulk or polycrystalline alloys and the lattice constant is observed (Vegards law). At larger Mn fractions ('38 at%) the alloy crystallizes in a face-centered tetragonal AuCu structure [11].
(2)
with the biaxial stress e "e "e and e "e . VV WW , XX , The calculated in-plane constant is a "3.91(3) As , indicating a strained epilayer. For the stochiometric compound MnPt we found a lattice constant a "3.89(5) As as expected for a slightly strained , layer. At higher Mn contents the out-of-plane lattice parameter of epitaxial "lms decreases more than expected for an elastically strained epilayer. To explain this discrepancy we developed a model which describes a change of the atomic distribution in the lattice induced by the strain. We consider that the di!erent lattice planes A and B are not "lled statistically with Mn atoms (Fig. 4b and c) which would be the case for the bulk alloy. The prefered "lling of Mn atoms of the planes A (ordered AuCu structure) is found for more than 38 at% Mn. Below this concentration, this behaviour could occur under the in#uence of the strain. Thus, the perpendicular lattice constant is smaller than the in-plane constant because Mn is smaller than Pt. Therefore, the strain induces a continuous phase transition from the cubic to the tetragonal phase. This e!ect is an additional indication that strain may change thermodynamical quantities [22].
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Fig. 6. The Curie temperature (¹ ) increases with the Mn con! tent x. The ¹ of epitaxial "lms is enhanced in comparison to ! polycrystalline "lms [25] and bulk material [10].
Fig. 5. Hysteresis loops of Mn Pt "lms with di!erent Mn V \V content x at room temperature (a). The saturation magnetization M is strongly dependent on the Mn content x. The 1 maximum of the magnetization occurs at 25 at% Mn. The curve of the Kerr rotation (*) is very similar to the dependence of M 1 (䢇) on the Mn content (b).
3.3. Magnetic properties Mn Pt is known to be a strong ferromagnet V \V [10]. Pickart et al. found a saturation magnetization M of MnPt bulk samples of (4.12$0.08)l 1 per unit cell at 77 K by neutron di!raction measurements [23]. At the same temperature Antonini et al. measured a magnetic moment of (4.04$0.08)l using ballistic methods and (4.42 $0.17)l performing polarized neutron measurements [24]. By using a SQUID magnetometer, we "nd a saturation magnetization of epitaxial MnPt "lms of 4.76 and 4.31l at 77 and 300 K, respectively. Fig. 5a shows SQUID hysteresis measurements of the in plane magnetization of our epitaxial thin Mn Pt "lms with di!erent Mn concentrations. V \V As one can see in Fig. 5b, the saturation magneti-
zation (z) depends strongly on the Mn content. Below the percolation limit of about 15 at% Mn, no magnetic ordering is found in Mn Pt , beV \V cause Mn has too few nearest and next-nearest Mn neighbours to achieve magnetic ordering. Magnetic ordering occurs at a Mn content of more than 15 at%, and M increases up to the stochiometric 1 MnPt alloy. In the ordered phase, the Mn atoms occupy the corner sites of the FCC cubic cell, whereas the Pt atoms are found on the facecentered sites. In this con"guration, the largest ferromagnetic coupling of the Mn atoms occurs. If the Mn content grows any further, Mn atoms start to occupy Pt sites and the magnetization decreases. The reason is obviously due to the fact, that the magnetic interaction between Mn atoms sitting on a Mn site and the nearest Pt site is antiferromagnetic. The Curie temperature (¹ ) of 375 K for bulk ! MnPt is well established [25]. Our measurements on epitaxial thin "lms reveal an enhanced Curie temperature of 500 K. Fig. 6 shows the Curie temperatures for bulk samples, magnetron sputtered thin "lms and epitaxial thin "lms as function of the Mn content [25,10]. While the bulk samples and sputtered "lms show a rather constant positive slope of ¹ with increasing Mn concentration, the ! epitaxial "lms behave quite di!erently. At a Mn content of +20 at% the Curie temperature grows
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Fig. 7. In plane (straight line) and out of plane (dashed and dotted lines) magnetization of MnPt "lms with various thick nesses (room temperature). Even in the thinnest "lm (17.5 nm) the easy axis of magnetization lies in the "lm plane.
Fig. 8. Experimental Kerr spectra of MnPt "lms at room temperature [8] and theoretical Kerr spectrum at ¹"0 K [26] in comparison to our epitaxial thin "lm data at room temperature.
rapidly with decreasing slope, when the concentration increases. This can be understood, if the strain of the "lms is considered. At a Mn content of 20 at%, the lattice parameters of Mn Pt and the V \V SrTiO substrate are identical. Therefore, the Curie temperatures of bulk samples and epitaxial "lms are similar. Beyond 20 at% Mn, a lattice mismatch occurs, which leads to a tensile strained "lm parallel to the surface and a compressively strained "lm perpendicular to the surface. Due to the changing Mn distances, the magnetic ordering is in#uenced, so that the exchange interaction increases or decreases depending on the direction. This leads to an enhanced Curie temperature because the larger exchange determines the ordering temperature. The ¹ curve of the epitaxial "lms comes closer to the ! curve for the bulk samples and the magnetron sputtered "lms at a Mn content of about 33 at%. MnPt crystallizes in the cubic AuCu structure. Therefore, no favored direction of magnetization along any of the cubic axes is expected. In a thin "lm, however, the easy axis of magnetization lies in the "lm plane due to the shape anisotropy. This prefered direction does not change, at least not down to a "lm thickness of 17.5 nm (Fig. 7). An easy axis of magnetization perpendicular to the "lm plane can be achieved, if thin Co layers are used as spacer layers between 1.9 nm thick MnPt layers [8].
3.4. Magneto-optical properties In Fig. 8 a theoretical Kerr spectrum (LDA calculations) [2,26] and two experimental Kerr spectra measured by di!erent authors [8,25] are compared with our results for an epitaxial (0 0 1)-oriented MnPt "lm. At "rst, we compare the theoretical data with our experimental data. The shape of the two spectra are quite similar to each other. The di!erent maximum values of the two spectra can be attributed in part to the reduction of the sample magnetization at room temperature (see Section 3.3). The theoretical Kerr spectrum is calculated for zero temperature. If the Kerr spectra are measured at room temperature, where the magnetization is reduced, then the overall size of the measured Kerr rotation will be reduced, too (see below). The reason for the shift of the theoretical spectrum to lower energies remains debatable. It is known that LDA calculations fail to describe properly exchange and correlation e!ects, underestimating these e!ects for occupied states and overestimating these e!ects for unoccupied states. The best agreement of the experimental spectra can be found between the (1 1 1)-grown MnPt "lms [8] and the present results from (0 0 1)-grown "lms (Fig. 9). The magnitude of the peak at about 1.1 eV increases nearly linearly with the magnetization of
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159
Fig. 9. Kerr spectra at room temperature of epitaxially grown Mn Pt "lms with di!erent Mn concentrations. The maxV \V imum of the Kerr rotation occurs at a Mn content of 25 at% (see Fig. 5b).
Fig. 10. Room temperature Kerr rotation of a MnPt "lm annealed at di!erent temperatures. The Kerr rotation was measured through a protective SiO layer at a magnetic "eld of 1 T V and a wavelengh of 1200 nm.
the samples (Fig. 5b). This e!ect was predicted in the work of Oppeneer et al. [26] who studied the in#uence of the exchange splitting on the Kerr rotation h of XPt compounds (X"V, Cr, Mn, ) Fe, and Co) and found an increase of h with ) increasing exchange splitting. This dependence can also explain why the Kerr rotation of our MnPt "lms is larger than e.g. that of Kato et al. [6]. We "nd a large maximum value of !0.823 around 1.1 eV and at least one additional peak of h K0.23 ) at approximately 2.6 eV. The sign of h is reversed ) at 2.2 eV. To our knowledge, this is the largest intrinsic Kerr rotation found at room temperature in MnPt , although Wierman et al. [7] quote larger values. The obvious reason is that we present the intrinsic Kerr signal at the interface MnPt /air, whereas the data from Ref. [7] are for the interface MnPt /SiO . Thus, the later data do not represent V the Kerr rotation of MnPt which is de"ned as the rotation at the interface vacuum/sample, but the data have to be scaled down by a factor which, in "rst order, is given by the refractive index of the SiO . In addition, interference e!ects may have to V be taken into account. In thin "lms even without a protective layer interference e!ects may occur if the "lm thickness is in the order of the optical penetration depth. To exclude this possibility we measured the Kerr rotation of "lms with di!erent thicknesses between 200 and 1000 As . No signi"cant di!erences of the rotation were found.
To show the dependence of the Kerr e!ect on the long-range order parameter, we grew a MnPt thin "lm at a temperature of about 853C on SrTiO substrates and coated it with a SiO layer. At this temperature, the "lm does not grow epitaxially, but with a long-range order parameter of nearly zero as determined by X-ray di!raction analysis. The "lm was heated for half an hour at di!erent temperatures and then cooled down to room temperature. Fig. 10 shows the increase of the absolute value of the Kerr rotation measured at room temperature at a wavelength of 1200 nm with increasing annealing temperature. The intrinsic Kerr rotation has to be scaled down by a factor of 1.5}1.9 due to the refractive index of the protective SiO layer on this V sample. After annealing at 8003C the maximum Kerr rotation is obtained. As already pointed out by Wierman and Kirby [25] this corroborates the large long-range order parameter of the epitaxially grown "lms.
4. Conclusions Thin epitaxial Mn Pt "lms have been preV \V pared by molecular beam epitaxy. The growth mechanism can be described by a Volmer}Weber model. The high structural quality of the "lms has been con"rmed by RHEED and X-ray di!raction. Magnetic characterization of the "lms with
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a SQUID magnetometer reveals a strong dependence of the saturation magnetization on the Mn content x. The maximum magnetization is obtained for the stochiometric highly ordered compound MnPt . In contrast to bulk material, the Curie temperature ¹ is enhanced and this is at! tributed to strain. The high structural and magnetic ordering is re#ected in a large room temperature Kerr rotation of !0.823 at 1.2 eV. We shows that the magnetization and the Kerr rotation is dependent on the Mn content and on the long-range order parameter. Acknowledgements We acknowledge the help of the Institut fuK r Halbleitertechnik (TU Braunschweig) for the EDX measurements. References [1] E. Kren, G. Kadar, L. Pal, J. Solyom, P. Szabo, T. Tarnoczi, Phys. Rev. 171 (1968) 574. [2] P.M. Oppeneer, V.N. Antonov, T. Kraft, H. Eschrig, A.N. Yaresko, A.Y. Perlov, Solid State Comm. 94 (1995) 255. [3] K.H.J. Buschow, P.G. van Engen, R. Jongebreur, J. Magn. Magn. Mater. 38 (1983) 1. [4] H. BraK ndle, J. Schoenes, F. Hulliger, W. Reim, unpublished. Data given by J. Schoenes in: R.W. Cahn, P. Haasen, E.J. Kramer (Eds.), Materials Science and Technology, Vol. 3, Verlag Chemie, Weinheim, 1992, p. 147. [5] M. VergoK hl, J. Schoenes, J. Magn. Soc. Jpn. 20 (1996) 141. [6] T. Kato, H. Kikuzawa, S. Iwata, S. Tsunashima, S. Uchiyama, J. Magn. Magn. Mater. 140}144 (1995) 713.
[7] K.W. Wierman, J.N. Hil"ker, R.F. Sabiryanov, S.S. Jaswal, R.D. Kirby, J.A. Woollam, Phys. Rev. B 55 (1997) 3093. [8] M. Kokuryu, T. Kato, S. Iwata, S. Tsunashima, J. Appl. Phys. 81 (1997) 4779. [9] R.J. Lange, S.J. Lee, D.W. Lynch, P.C. Can"eld, B.N. Harmon, S. Zollner, Phys. Rev. B 58 (1998) 351. [10] M. AuwaK rter, A. Ku{mann, Ann. Phys. 6 (1950) 169. [11] E. Raub, W. Mahler, Z. Metallk. 4 (1955) 282. [12] B.S. Kwak, P.N. First, A. Erbil, B.J. Wilkens, J.D. Budai, M.F. Chisholm, L.A. Boatner, J. Appl. Phys. 72 (1992) 3735. [13] B.M. Lairson, M.R. Visokay, R. Sinclair, S. Hagstrom, B.M. Clemens, Appl. Phys. Lett. 61 (1992) 1390. [14] J.F.M. Cillessen, R.M. Wolf, D.M. de Leeuw, Thin Solid Films 226 (1993) 53. [15] I.J.M.M. Raaijmakers, R.A.A. Hack, A.G. Dirks, Suppl. Vide, Couches Mines, Proceedings of the International Symposium on Trends and New Applications in Thin Films, 1987. [16] J.M. Howe, Interfaces in Materials, Wiley, New York, 1997. [17] C.R. Henry, C. Chapton, C. Duriez, S. Giorgio, Surf. Sci. 253 (1991) 177. [18] N.A. Gjostein, in: J.J. Burke, M.L. Reed, V. Weiss (Eds.), Surfaces and Interfaces I, Syracuse University Press, New York, 1967, p. 271. [19] B.E. Warren, X-ray di!raction, Addison-Wesley, Reading, MA, 1969. [20] S. Jin, T.H. Tiefel, M. McCormack, H.M. O'Bryan, L.H. Chen, R. Ramesh, D. Schurig, Appl. Phys. Lett. 67 (1995) 557. [21] J.Y. Tsao, Materials fundamentals of molecular beam epitaxy, Academic Press, San Diego, 1993. [22] R.A. Rao, Q. Gan, C.B. Eom, R.J. Cava, Y. Suzuki, J.J. Krajewski, S.C. Gausepohl, M. Lee, Appl. Phys. Lett. 70 (1997) 3035. [23] S.J. Pickart, R. Nathans, J. Appl. Phys. 33 (1962) 1336. [24] B. Antonini, F. Lucari, F. Menzinger, A. Paoletti, Phys. Rev. 187 (1969) 611. [25] K.W. Wierman, R.D. Kirby, J. Magn. Magn. Mater. 154 (1996) 12. [26] P.M. Oppeneer, V.N. Antonov, H. Eschrig, A.N. Yaresko, A.Y. Perlov, J. Phys.: Condens. Matter 8 (1996) 5769.