Epitaxial ferromagnetic thin films and superlattices of Mn-based metallic compounds on GaAs

MATERIALS SCIENCE & ENGINEERING ELSEVIER

Materials Science and Engineering B31 (1995) 117-125

B

Epitaxial ferromagnetic thin films and superlattices of Mn-based metallic compounds on GaAs Masaaki Tanaka* Department of Electronic Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113, Japan

Abstract

Epitaxial ferromagnetic thin films grown directly on I I I - V compound semiconductors can lead to the integration of magnetism and high-speed I I I - V electronics/photonics, offering possibilities of hybrid ferromagnetic-semiconductor devices. Despite stringent material requirements for this purpose, we have successfully grown some of Mn-based metallic compounds sharing common atoms with the underlying III-V using molecular beam epitaxy (MBE). First, it is shown that ferromagnetic MnGa thin films with perpendicular magnetization were obtained on (001) GaAs substrates. The c axis, which is the easy axis for magnetization, of the tetragonal unit cell is found to be desirably aligned perpendicular to the film plane. Second, in order to derive more functionality and to increase the freedom in material design, we have also created a new class ot magnetic superlattices (SLs) of metallic compounds, MnGa/NiGa, on GaAs substrates. The MnGa/NiGa SLs show even stronger perpendicular magnetization and evenly-spaced multiple steps in the hysteresis loops in extraordinary Hall effect measurements. Third, ferromagnetic hexagonal MnAs, an As-based compound which is more compatible with existing I I I - V MBE technology, is shown to be successfully grown on GaAs. It was found that the MBE-grown MnAs/GaAs heterostructures have a unique epitaxial relationship and a mono-atomic template plays a critical role in determining the epitaxial orientation and magnetic properties.

Keywords: Magnetic thin films; Magnetic anisotrophy; Mn compounds; Molecular beam epitaxy

I. Introduction

Since Esaki and Tsu proposed a concept of semiconductor superlattices [1], the development of heteroepitaxial growth technology has made it possible to grow a variety of ultrathin heterostructures (HSs) and artificial superlattices (SLs). Most of the HSs and SLs used for electronic and optical devices consist of two semiconductors, such as G a A s / A 1 G a A s , which are very similar in terms of crystal structure, lattice constant, and chemical bonding. Although epitaxial growth of such all-semiconductor heterostructures is now routinely done, they are only a part of the whole potential scope of applications of heterostructures. Recently, some HSs consisting of dissimilar materials have been reported. For example, transition metal (TM) aluminides and gallides (e.g. NiA1 [2], N i G a [3], C o A l [4], C o G a [5]) were epitaxially grown on G a A s by molecular b e a m epitaxy (MBE). These T M - I I I * Also at: Bellcore, 331 Newman Springs Road, Red Bank, NJ 07701, USA. 0921-5107/95/$09.50 © 1995 Elsevier Science S.A. All rights reserved SSDI 0 9 2 1 - 5 1 0 7 ( 9 4 ) 0 8 0 1 3 - 5

type metallic compounds have CsCl-type crystal structures whose lattice constants are close to half the lattice constant of G a A s (0.5a~aAs = 0.283 nm), and therefore they are compatible in heteroepitaxy with the zinc-blende type structure of G a A s and other I I I - V semiconductors. This kind of T M - I I I / I I I - V metal semiconductor HSs can open up new applications. One of the most attractive new directions both for materials science and for device applications will be the integration of magnetic materials and I I I - V semiconductors, offering possibilities of hybrid ferromagn e t i c - s e m i c o n d u c t o r devices. For this purpose, ferromagnets must be (1) grown epitaxially on the I I I - V semiconductor, (2) thermodynamically stable with no interfacial reactions on the I I I - V , and (3) morphologically stable with no agglomerations on the I I I - V . These material requirements are quite stringent, and none of the elemental metals, including Fe and Co on G a A s , can meet all three requirements. We have found that some of the Mn-based metallic compounds sharing c o m m o n atoms with the underly-

M. Tanaka / Materials Science and Engineering B31 (1995) 117-125

118

ing I I I - V are promising. Among these T M - I I I metals, MnxGal x with a Mn composition x of 55-60%, whose crystal structure is tetragonal, is especially interesting, since it is a stable ferromagnetic phase, thermodynamically and morphologically stable on GaAs, and can be epitaxially grown directly on (001) G a A s by M B E [6]. The epitaxial MnGa thin films on G a A s are shown to exhibit perpendicular magnetization due to magneto-crystalline anisotropy. We have also created a new class of SLs consisting of ferromagnetic M n G a and non-magnetic NiGa [7]. Furthermore, we have successfully grown ferromagnetic MnAs, an As-based compound which is more compatible with existing I I I - V M B E technology, on (001) GaAs [8]. In this paper, epitaxial growth, structure, and magnetic and magnetotransport properties of such Mnbased ferromagnetic thin films and magnetic SLs on G a A s are described.

2. MnGa on GaAs

2.1. Material properties and MBE growth Unlike metastable rMnA1, MnxGa I ., with x 0.55-0.60 (a-phase) is a thermodynamically stable ferromagnetic phase in the bulk M n - G a system [9,10]. In the bulk, the single-phase region was found to extend from 54.5 to 60.0 at.% Mn at 450°C [10]. The crystal structure of the MnGa is tetragonal with an ordered crystal structure of the CuAu type. The M n -

Mn spacing in the basal plane (a~) is 0.274 nm, independent of the Mn content, and along the c axis (co) is 0.365 nm at 56% Mn and 0.369 nm at 59% Mn, respectively. The easy magnetization direction is along the c axis. Since a 0 is close to half of the lattice constant of GaAs (0.283 nm) with a lattice mismatch of only 3.8%, the orientation relationship with the c axis of the M n G a parallel to the surface normal of the (001) GaAs substrates is expected, as shown in Fig. 1. The growth of M n G a on (001) GaAs was performed with a conventional I I I - V M B E machine (Riber 2300) with effusion cells for Mn, Ni, Ga and As, using the multistep technique analogous to that employed in the growth of rMnA1/AIAs/GaAs [11,12]. Since the details of the growth have been reported previously [6], we briefly summarize the growth procedure here. After a 100 nm-thick GaAs buffer layer was grown at 580°C under conventional I I I - V growth conditions, the substrate temperature was cooled to 20-40°C while completely eliminating the As flux. Then a thin (about 0.9 nm thick) amorphous M n G a template was deposited at 20-40°C. The amorphous template was next heated up to 200-250°C to form a monocrystalline template by solid phase epitaxy. After the template with the desired epitaxial relation was thus established, Mn and Ga were subsequently codeposited at 150-200°C at a growth rate of 0.05 p.m h i, to a final MnxGa I • thickness of 3-60 nm. The Mn content x was set to 5 5 - 6 0 % in the present study. Finally, postgrowth annealing was performed at 300400°C for 2 min. Such annealing at a temperature T,

~o

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D

Tetragonal MnGa Ga •

Mn

Zincblende GaAs )O

ao/2 = 0 . 2 8 3 nm

Ga

~ii:iii¸~

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Fig. 1. Left, crystal structures and epitaxial orientation of M n G a on G a A s ; right, cross-sectional T E M image taken from a 10 nm-thick Mn,oGa4, ~ film on G a A s .

M. Tanaka / Materials Science and Engineering B31 (1995) 117-125

of at least 300°C was found to be necessary in order to improve the structural and magnetic properties of the M n G a film.

2.2. Structure and magnetic properties of epitaxial MnGa on GaAs

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The epitaxial M n G a with its c axis aligned normal to the substrate was confirmed first by in-situ reflection high energy electron diffraction ( R H E E D ) and also by cross-sectional transmission electron microscopy (TEM). Fig. 1 also shows a (110) T E M lattice image of the 10 nm-thick Mn60Ga40 film grown on GaAs, indicating that a monocrystalline M n G a layer was indeed grown with the c axis properly oriented perpendicular to the GaAs substrate. The interface between the M n G a and the GaAs is found to be very smooth and abrupt, and no interfacial transition layer was observed. The horizontal lattice fringes with a spacing of 0.31 nm correspond to the (001) planes of the tetragonal structure of the MnGa. This value of the c parameter in the present epitaxial M n G a film is significantly smaller than that of bulk M n G a (0.360.37 nm) [9,10]. T o investigate the magnetic properties of the epitaxial M n G a films, we performed vibrating sample magn e t o m e t e r (VSM) measurements, as a function of magnetic field, perpendicular to the substrate at room temperature. Fairly square hysteresis loops were observed at room temperature [6], offering direct evidence of the desired perpendicular magnetization. Similar results were obtained in all the samples of M n G a with the Mn content ranging from 55% to 60%. The remnant magnetization M r along the c axis was estimated to be 113-225 emu cm -3, and the coercive field H c was 0.85-3.15 kOe, depending on the thickness of M n G a and the annealing temperature [13]. The Hall effect measured in a ferromagnetic film is known to have an "extraordinary" component resulting from the asymmetric scattering of carriers with magnetic atoms [14[. The extraordinary Hall effect ( E H E ) is very useful in this case because the Hall resistivity measured in a patterned Hall bar in the M n G a films is proportional to the magnetic moment perpendicular to the film plane. Fig. 2 shows an example of Hall resistance Rxy vs. perpendicular magnetic field measured at room temperature on a 200 /xm-wide Hall bar fabricated in a 10 nm-thick Mn58Ga42 film prepared with T a = 300°C. One can see a fairly square hysteresis loop with nearly 100% remanence, a Hall resistance Rxy of 3.5 f~ (corresponding a Hall resistivity Pxy of 3.5/xf~cm) and H c of 1.46 kOe. Similar square-like hysteresis loops were observed in the E H E measurements in all the present M n G a samples, indicating a large component of perpendicular magnetization, consistent with the VSM

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measurement results. The E H E resistivity Pxy was found to vary from 0.1 ~f~cm to 4 / z I l c m , depending on the growth parameters [13].

3. M n G a / N i G a Superlattices on GaAs

3.1. Formation of MnGa/NiGa magnetic superlattices on GaAs In order to increase the freedom in material design, we have created a new class of ultrathin epitaxial superlattices consisting of ferromagnetic tetragonal M n G a and non-magnetic cubic (CsCl-type) NiGa on (001) GaAs substrates, as shown in Fig. 3. NiGa with a stoichiometric composition is a nonmagnetic metal, with cubic CsCl-type crystal structures having the lattice constant a o of 0.289 nm. Since the ao value of this T M - I I I intermetallic is close to half (0.283 nm) the lattice constant of GaAs, one can expect that M n G a / N i G a SLs should be compatible in heteroepitaxy with GaAs and related I I I - V semiconductors, with the desired "cube-on-cube" crystal orientations of [110](001) T M - I I I / / [ l l 0 ] ( 0 0 1 ) GaAs. The growth of M n G a / N i G a SLs on GaAs (001) substrates was performed with a conventional I I I - V M B E machine [7]. Here we briefly summarize the growth procedure. After growing a 100 nm-thick undoped GaAs buffer layer at 580°C on the semiinsulating GaAs substrate, the substrate temperature T s was cooled to 200-220°C while completely eliminating the As flux in the growth chamber. The Ni shutter was opened to deposit one monolayer (ML) of Ni on the c(4 x 4) GaAs surface, and then Ni and Ga

M. Tanaka / Materials Science and Engineering B31 (1995) 117-125

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were codeposited to grow n MLs of NiGa, followed by the codeposition of Mn and G a to grow m MLs of M n G a . The N i / G a and M n / G a ratio were set to 50/50 and 60/40, respectively. After growing the M n G a / N i G a SLs with p periods at 200-220°C, a postgrowth annealing was p e r f o r m e d at 300°C for 1-2 rain to i m p r o v e the structural and magnetic properties of the MnGa/NiGa SLs. T h r o u g h o u t the growth, the R H E E D patterns were very streaky and sharp, indicating that epitaxial monocrystalline metallic layers with good quality are successfully grown [7]. Fig. 3 shows cross-sectional T E M images of the M n G a / N i G a SL with rn = n = 6 and p = 5, giving the direct evidence of the successful formation of the compositional SLs with atomically abrupt and smooth interfaces. We have confirmed the " c u b e - o n - c u b e " epitaxial relationship as expected, where the c axis of the tetragonal M n G a is aligned perpendicular to the film plane, by these R H E E D and T E M analyses. From the T E M diffractions, the M n G a in the SLs is found to have reduced tetragonality, with a c,, value of 0.35 nm, the value being smaller than those of bulk. This reduced tetragonality influences the magnetic properties, particularly the coercive field of the ferromagnetic films.

3.2. Magnetic and magnetotransport properties To investigate the magnetic properties of the M n G a / N i G a SLs, we p e r f o r m e d both VSM and E H E m e a s u r e m e n t s at room t e m p e r a t u r e [15]. Both VSM and E H E hysteresis loops are square for the magnetic field applied normal to the substrates, indicating the

desired perpendicular magnetization of the superlatrice samples. VSM m e a s u r e m e n t s on these M n G a / N i G a ultrathin superlattices showed higher values of remnant magnetization of 267-302 emu cm 3 at r o o m temperature. More interestingly, E H E hysteresis loops of the M n G a / N i G a SLs are found to have multiple steps at room temperature. Fig. 4 shows the E H E hysteresis loops of the M n G a / N i G a SLs with (Figs. 4(a)-4(c)) m = n = 10 and p = 3, and (Figs. 4(d) and 4(e)) with r n - n = 6 and p = 5, measured at r o o m t e m p e r a t u r e with differing m a x i m u m values H ..... of the applied magnetic field perpendicular to the substrate plane. The H ..... value was set to (Fig. 4(a)) 2.7 k O e , (Fig. 4(b)) 4.0 kOe and (Fig. 4(c)) 7.0 k O e for the (MnGa)~0/(NiGa)L0 SL with 3 periods, and (Fig. 4(d)) to 7.0 k O e and (Fig. 4(e)) 2.6-7.0 k O e for the (MnGa)6/(NiGa)6 SL with 5 periods. W h e n the applied field is as high as 7.0 k O e , the entire magnetic structure is saturated for both SLs. Note that the E H E hysteresis loops have evenly-spaced multiple steps when H ..... is 7.0 k O e , as shown in Figs. 4(c) and 4(d), where three steps and five steps are clearly seen in the E H E loops of M n G a / N i G a SLs with 3 periods and 5 periods, respectively. Since the n u m b e r of steps in the loops are the same as the n u m b e r of ferromagnetic M n G a layers, each step is formed by the reversal of magnetization in one of the M n G a layers. This is because the M n G a layers have different coercive forces H~ from one to another. In principle, these p h e n o m e n a can be applied to multi-level signal recording, which can lead to much higher storage

M. Tanaka

Materials Science and Engineering B31 (1995) 117-125

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Fig. 4. E H E hysteresis loops of the M n G a / N i G a SLs measured at room temperature with differing maximum values Hm, x of the applied magnetic field perpendicular to the substrate plane, and with: ( a ) - ( c ) m = n = 10 and p = 3; (d), (e) m = n = 6 and p = 5. The Hm, x value was set to (a) 2.7 kOe, (b) 4.0 kOe and (c) 7.0 kOe for the (MnGa)10/(NiGa)H, SL with 3 periods, and (d) 7.0 KOe and (e) 2.6-7.0 kOe for the (MnGa)6(NiGa)6 SL with 5 periods.

density. T h e capability of reading and writing multilevel signals is d e m o n s t r a t e d for the M n G a / N i G a SL with 3 periods in Figs. 4(a)-4(c), where four different static values of Hall resistance are realized on varying the value H ..... . Likewise, in Fig. 4(e) six different values of Hall resistance are shown for the M n G a / N i G a SL with 5 periods. This artificial material appears to be a promising candidate for non-volatile magnetic m e m o r y coupled with underlying I I I - V circuitry. The possible origin of the different H c values in the M n G a layers of the SLs is the difference in lattice strain. Since the bulk a o values of tetragonal M n G a and CsCl-type N i G a have the lattice mismatch o f 3.8% and + 2.1%, respectively, the c/a ratio in the lattice of the epitaxial M n G a layers of the SLs can be slightly different due to the difference in strain, leading to the difference in H c. A n o t h e r possible reason for the different H c may be the difference in the interface roughness or defect density. The full understanding of the mechanisms and the artificial

"design" of the E H E hysteresis will be interesting future issues.

4. M n A s on G a A s

4.1. Material properties and M B E growth In the bulk M n - A s system, M n A s at the stoichiometric composition (MnsoA%0) has a t h e r m o d y namically stable ferromagnetic phase ( a - p h a s e ) at r o o m t e m p e r a t u r e [16]. The Curie t e m p e r a t u r e of the ferromagnetic a M n A s is 40°C, slightly above r o o m temperature. The crystal structure of the a M n A s is hexagonal of NiAs-type, with lattice constants of a = 0.3725 nm and c = 0.5713 nm. At the stoichiometric composition, this ferromagnetic M n A s phase extends up to 40°C, and p a r a m a g n e t i c / 3 M n A s of a MnP-type orthorhombic structure (with lattice constants of a = 0.366 nm, b = 0.636 nm and c = 0.572 nm) and para-

122

M. Tanaka / Materials Science and Engineering B31 (1995) 117-125

magnetic y M n A s of a NiAs-type hexagonal structure exist at 40-125°C and above 125°C, respectively. H e r e we will show that MBE-grown ferromagnetic M n A s thin films take two different epitaxial orientations depending on the growth procedure. Very recently, we have found that epitaxial MnAs can be grown on (001) G a A s with the growth plane of (1100), which we refer to as type A, as shown in Fig. 5(a). H e r e we clarify the process of template formation to obtain this type A epitaxial relationship. We also show how changing the template results in a different epitaxial growth plane, the (3101) of MnAs, as shown in Fig.5(b), which we refer to as type B. M B E growth of the M n A s thin films was p e r f o r m e d with a conventional I I I - V M B E machine (Riber 2300) with a Mn effusion cell and an As-valved cracker. After growing a 100 nm-thick undoped G a A s buffer layer at 580°C on a (001) G a A s substrate, the sub-

strate t e m p e r a t u r e was cooled to 200-250°C and the As flux was shut off. At this t e m p e r a t u r e , the G a A s surface shows a c(4 x 4) reconstruction in R H E E D . We then grew two M n A s samples with different templates as follows. Type A: A s 2 flUX was first supplied for 1-2 min on this c(4 x 4) G a A s surface at 200°C, after which the surface showed disordered c(4 × 4) (d(4 x 4)) reconstruction [17], where the twofold streaks in the R H E E D are vague and surface reconstruction is disordered due to an extra As layer on top of the c(4 x 4) surface. On this d(4 x 4) surface, which is more As-rich than c(4 x 4), Mn and As 2 fluxes were supplied to grow epitaxial M n A s with a growth rate of 50 nm h t Type B: One monolayer of Mn was first deposited on the clear c(4 x 4) G a A s surface at 200-220°C, causing the twofold streaks to disappear and the

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M. Tanaka / Materials Science and Engineering B31 (1995) 117-125

pattern to turn to (1 x 1) in the R H E E D . On this surface, Mn and A S 2 fluxes were supplied to grow MnAs with a growth rate of 50 nm h-1. For both type A and type B samples, the substrate temperature was set to 250°C 10 rain after forming the template and starting the MnAs growth. The total thickness of the MBE-grown MnAs was set to 20-100 nm. Unlike the MBE growth of M n G a and MnA1, the As2/Mn flux ratio is not very critical in the growth of MnAs and was set at 2.0-5.0, an As-rich condition, since excess As atoms are re-evaporated from the surface and stoichiometric MnAs is easily obtained under the present growth conditions in a manner analogous to the MBE growth of GaAs. After completing the growth, postgrowth annealing was done under the As 2 flux at 400°C for 1-2 min to improve the structural and magnetic properties.

4.2. Epitaxial orientations Fig. 6 shows schematic sample layer structures, and in-situ R H E E D patterns of the 50 nm-thick MnAs of type A (Figs. 6(a) and 6(b)) and of type B (Figs. 6(c) and 6(d)), with the incident electron beam azimuth along the [110] and ~10] of the GaAs substrates. Though epitaxial MnAs thin films with a smooth surface were obtained for both type A and type B growth, the different patterns show the different epitaxial orientations. The analysis of the R H E E D patterns along various incident angles revealed that

123

the growth plane of type A and type B is ~100) and (1101), respectively, as shown in Fig. 5. It was also found from the R H E E D that the epitaxial in-plane relationship is [--H20]MnAs//[llO]GaAs and [O001]MnAs//[llO~aAs for type A, and l[]-]-20]MnAs/ and [ll02]MnAs//[ll0] GaAs for type B. Note here that the direction of the [H20] axis of MnAs of type A and type B is 90 ° different with respect to the GaAs substrates. We speculate that this is related to the difference in the atomic structure of the c(4 x 4) and d(4 x 4) GaAs surfaces. The direction of the As dimers on the top layer of the c(4 x 4) is known to be along the [110] [18], whereas the atomic structure of the d(4 x 4) surface is unknown. However, it is likely that the d(4 x 4) surface has an extra As layer on top of the c ( 4 x 4) surface, with the direction of the top As bonds rotated 90 ° [17], leading to the 90 ° rotation of the l[-H20] axis of the epitaxially grown MnAs films. X-ray measurements (0-20 scan) were made on both types of MnAs thin films, and we confirmed the R H E E D results described above. For type A samples, a (]-100) MnAs peak was found in the 20 spectra, indicating that the growth direction is normal to the (1100). In contrast, type-B MnAs thin films showed a (]-101) peak and a much smaller (1102) peak. Since the intensity of the (1101) peak is ten times as high as the (1102) peak, the growth direction of the type-B MnAs is dominantly normal to the (H01), and additionally normal to the ([102).

/[Tl0]GaAs

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As

GaAs (001)

Type-B I

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M. Tanaka / Materials Science and Engineering B31 (1995) 117-125

124

4.3. Magnetic properties To investigate the magnetic properties of the type-A and type-B MnAs grown on GaAs, we have performed magnetization measurements using a vibrating sample magnetometer (VSM), on the 50 nm-thick films of both types at room temperature, with various directions of applied magnetic field. The results of the M - H measurements are summarized in Fig. 7. The type-A MnAs films show almost perfectly square hysteresis characteristics when the applied magnetic field H was along the [110] of GaAs which is parallel to the l[H-20] of MnAs, as shown in Fig. 7(a). The remnant magnetization M~ is 416 emu cm 3, and the coercive field H~ is 196 Oe. However, when H was applied along the []-10] of GaAs, which is parallel to the [0001] of MnAs, no hysteresis loop was observed, as shown in Fig. 7(b). When H was along the [001] GaAs (perpendicular to the substrate plane), only a little hysteresis was seen. This indicates that the easy magnetization axis of the type-A MnAs was in plane, along the l[-[-fZ0]MnAs//[ll0] of GaAs. Similar measurements were performed on the typeB MnAs films. In contrast to type A, no hysteresis was

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type-A

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seen when H was along the [110] of GaAs, as shown in Fig. 7(c), or when H was along the [001] of GaAs. However, a square-like hysteresis loop was recorded when H was applied along the []-10] of GaAs, which is parallel to the l[H-20] of MnAs, as shown in Fig. 7(d). The values of M r and H~ of the type-B MnAs are 331 emu cm 3 and 319 Oe, respectively. Thus the easy magnetization direction of the type B 'is along the l[H-201 of MnAs (the same as type A), which is parallel to the [110] of GaAs (90 ° different from type A). These results indicate the importance of the very first monolayer and the interface stoichiometry in determining the epitaxial relationship and magnetic properties in M n A s / G a A s systems.

5. Conclusions

We have successfully grown epitaxial ferromagnetic thin films and superlattices of Mn-based metallic compounds on (001) GaAs substrates by MBE. Careful control of the growth processes, including appropriate template approaches, leads to the desired epitaxial orientations and good magnetic properties. These novel heterostructures consisting of Mn-based magnetic compounds and III-V semiconductors are promising materials for potential applications to a new class of devices using both their magnetic and semiconducting properties.

Acknowledgments

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This work has been performed mainly at Bellcore in collaboration with J.P. Harbison, T. Sands, T.L. Cheeks, J. De Boeck, L.T. Florez, V.G. Keramidas, and G.M. Rothberg. The author wishes to acknowledge all these coworkers. Useful discussions with C.J. Palmstrom are appreciated. Thanks are due to constant encouragement from Prof. T. Nishinaga of the University of Tokyo. This work is partly supported by the Japan Society for the Promotion of Science.

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,

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, i , i

, i , i , i , i , E , i

-1.0 - 0 . 8 - 0 . 6 - 0 . 4 - 0 2 Magnetic

0

0.2

, i ,

0.4 0.6 0.8

1.0

Field (kOe)

Fig. 7. Room temperature M - H characteristics of the MnAs thin films of ((a) and (b)) type A and ((c) and (d)) type B with thickness 5(1 nm grown on GaAs. The magnetic field was applied ((a) and (c)) along the [110], and ((h) and (d)) along the [110] of GaAs substrates. Note that the range of the magnetic field H is from - 1.0 to 1.0 kOe in (a) and (d), where H is parallel to the easy magnetization axis and square hysteresis loops are seen, while it is from - 1 0 . 0 to 10.0 kOe in (b) and (c) where no hysteresis is seen.

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