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Making Nonmagnetic Semiconductors Ferromagnetic

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Science  14 Aug 1998:
Vol. 281, Issue 5379, pp. 951-956
DOI: 10.1126/science.281.5379.951

Abstract

REVIEW

Semiconductor devices generally take advantage of the charge of electrons, whereas magnetic materials are used for recording information involving electron spin. To make use of both charge and spin of electrons in semiconductors, a high concentration of magnetic elements can be introduced in nonmagnetic III-V semiconductors currently in use for devices. Low solubility of magnetic elements was overcome by low-temperature nonequilibrium molecular beam epitaxial growth, and ferromagnetic (Ga,Mn)As was realized. Magnetotransport measurements revealed that the magnetic transition temperature can be as high as 110 kelvin. The origin of the ferromagnetic interaction is discussed. Multilayer heterostructures including resonant tunneling diodes (RTDs) have also successfully been fabricated. The magnetic coupling between two ferromagnetic (Ga,Mn)As films separated by a nonmagnetic layer indicated the critical role of the holes in the magnetic coupling. The magnetic coupling in all semiconductor ferromagnetic/nonmagnetic layered structures, together with the possibility of spin filtering in RTDs, shows the potential of the present material system for exploring new physics and for developing new functionality toward future electronics.

The mass, charge, and spin of electrons in the solid state lay the foundation of the information technology we use today. Integrated circuits and high-frequency devices made of semiconductors, used for information processing and communications, have had great success using the charge of electrons in semiconductors. Mass storage of information—indispensable for information technology—is carried out by magnetic recording (hard disks, magnetic tapes, magneto-optical disks) using spin of electrons in ferromagnetic materials. It is then quite natural to ask if both the charge and spin of electrons can be used to further enhance the performance of devices. We may then be able to use the capability of mass storage and processing of information at the same time. Alternatively, we may be able to inject spin-polarized current into semiconductors to control the spin state of carriers, which may allow us to carry out qubit (quantum bit) operations required for quantum computing (1). However, there are good reasons why this has not yet been realized. The semiconductors used for devices and integrated circuits, such as silicon (Si) and gallium arsenide (GaAs), do not contain magnetic ions and are nonmagnetic (Fig. 1C), and their magnetic gfactors are generally rather small. In order for there to be a useful difference in energy between the two possible electron spin orientations, the magnetic fields that would have to be applied are too high for everyday use. Moreover, the crystal structures of magnetic materials are usually quite different from that of the semiconductors used in electronics, which makes both materials incompatible with each other.

Figure 1

Three types of semiconductors: (A) a magnetic semiconductor, in which a periodic array of magnetic element is present; (B) a diluted magnetic semiconductor, an alloy between nonmagnetic semiconductor and magnetic element; and (C) a nonmagnetic semiconductor, which contains no magnetic ions.

Ferromagnetism and semiconducting properties coexist in magnetic semiconductors, such as europium chalcogenides and semiconducting spinels that have a periodic array of magnetic elements (Fig. 1A) (2). In these magnetic semiconductors, which were extensively studied in the late 1960s to early 1970s, exchange interactions between the electrons in the semiconducting band and the localized electrons at the magnetic ions lead to a number of peculiar and interesting properties, such as a red shift of band gap when ferromagnetism sets in. Unfortunately, the crystal structure of such magnetic semiconductors is quite different from that of Si and GaAs; in addition, the crystal growth of these compounds is notoriously difficult. To obtain even a small, single crystal requires weeks of preparation and growth.

Making Nonmagnetic Semiconductors Magnetic

The usefulness of semiconductors resides in the ability to dope them with impurities to change their properties, usually top- or n-type. This approach can be followed to introduce magnetic elements into nonmagnetic semiconductors to make them magnetic. This category of semiconductors, called diluted magnetic semiconductors (DMSs; Fig. 1B), are alloys of nonmagnetic semiconductor (Fig. 1C) and magnetic elements (3). Study of DMSs and their heterostructures have centered mostly on II-VI semiconductors, such as CdTe and ZnSe, in which the valence of the cations matches that of the common magnetic ions such as Mn. Although this phenomenon makes these DMSs relatively easy to prepare in bulk form as well as in thin epitaxial layers, II-VI–based DMSs have been difficult to dope to create p- and n-type, which made the material less attractive for applications. The magnetic interaction in II-VI DMSs is dominated by the antiferromagnetic exchange among the Mn spins, which results in the paramagnetic, antiferromagnetic, or spin-glass behavior of the material. It was not possible until very recently to make a II-VI DMS ferromagnetic at low temperature (<2 K) (4).

Ferromagnetic III-V Semiconductors

An approach compatible with the semiconductors used in present-day electronics is to make nonmagnetic III-V semiconductors magnetic, and even ferromagnetic, by introducing a high concentration of magnetic ions. The III-V semiconductors such as GaAs are already in use in a wide variety of electronic equipment in the form of electronic and optoelectronic devices, including cellular phones (microwave transistors), compact disks (semiconductor lasers), and in many other applications. Therefore, the introduction of magnetic III-V semiconductors opens up the possibility of using a variety of magnetic phenomena not present in conventional nonmagnetic III-V semiconductors in the optical and electrical devices already established.

The major obstacle in making III-V semiconductors magnetic has been the low solubility of magnetic elements (such as Mn) in the compounds. Because the magnetic effects are roughly proportional to the concentration of the magnetic ions, one would not expect a major change in properties with limited solubility of magnetic impurities, of the order of 1018 cm–3 or less. A breakthrough was made by using molecular beam epitaxy (MBE), a thin-film growth technique in vacuum that allows one to work far from equilibrium. When a high concentration of magnetic elements is introduced in excess of the solubility limit, formation of the second phase occurs if conditions are near equilibrium. However, when the crystal is grown at low temperature by MBE, there is not enough thermal energy available to form the second phase, and yet there still exists a local potential landscape that allows epitaxial growth of a single-crystal alloy. The effort to grow new III-V–based DMSs by low-temperature MBE was rewarded with successful epitaxial growth of uniform (In,Mn)As films on GaAs substrates in 1989 (5), where partial ferromagnetic order was found (6), and ferromagnetic (Ga,Mn)As in 1996 (7). In the remainder of this review, I describe the preparation and properties of ferromagnetic III-V semiconductors, with particular emphasis on (Ga,Mn)As, and what can be done with the heterostructures based on (Ga,Mn)As.

Molecular Beam Epitaxial Growth

(Ga,Mn)As films have been grown on semi-insulating (001) GaAs substrates in an MBE chamber equipped with solid sources of elemental Ga, Mn, Al, and As. Reflection high-energy electron diffraction (RHEED) patterns were used to monitor the surface reconstruction during growth, which was always carried out under As-stabilized conditions (excess of As). Either a GaAs buffer layer or an (Al,Ga)As buffer layer was then grown before growth of (Ga,Mn)As. For the GaAs buffer, after lowering the substrate temperatureT S to 250°C, a 100-nm GaAs layer was grown before the growth of 150- to 200-nm-thick (Ga,Mn)As, whereas for the (Al,Ga)As buffer, a high growth temperature of 600° to 700°C was maintained, and then T S was lowered for (Ga,Mn)As growth. When the GaAs buffer layer growth was initiated at 250°C, the c(4×4) surface reconstruction pattern of GaAs changed to a (1×1) pattern. No change in the beam fluxes from the high-temperature growth of GaAs was made for this low-temperature GaAs growth. The (Ga,Mn)As growth was started by simply commencing the Mn beam during the low-temperature GaAs growth and keepingT S constant at 250°C. No special precaution was taken at the start of (Ga,Mn)As growth. Typical growth rates were 0.6 μm/hour, with Mn concentration x in (Ga1–xMnx)As films up to 0.07. Although the properties of grown (Ga,Mn)As do depend on growth parameters such as As overpressure and T S, as long as the established growth procedure was followed, the properties of (Ga,Mn)As films were reproducible; for example, for a given Mn concentration x, the ferromagnetic transition temperatureT C was always in the range of 2000x ± 10 K. The surface reconstruction of (Ga,Mn)As was (1×2) during and after growth. When the Mn flux or the substrate temperature, or both, were too high, a complex RHEED pattern appeared that indicated the appearance of the MnAs (NiAs structure) second phase on the surface. A schematic phase diagram of MBE growth is depicted inFig. 2. Details of the growth can be found in (7–9).

Figure 2

Schematic phase diagram showing the relation between growth parameters (substrate temperature and Mn concentration) and the properties of (Ga,Mn)As grown by molecular beam epitaxy. The high concentration of Mn in excess of its solubility limit was introduced by nonequilibrium growth at low temperatures.

Clear RHEED oscillations were observed at the initial stage of growth, which indicated two-dimensional layer-by-layer growth as opposed to island growth. One important finding is that GaAs grown at 250°C could still show very distinct oscillations. Although the oscillations were modified by the presence of Mn, which could act as a surfactant layer, clear oscillations were observed during low-temperature MBE of GaAs even in the growth chamber without a Mn cell. The systematic study on the RHEED oscillations of GaAs in the temperature range from 150° to 700°C was reported elsewhere (10).

Lattice Constant of (Ga,Mn)As

The lattice constants a of the (Ga,Mn)As layers (7) were determined by x-ray diffraction (XRD) as a function of x and are shown in Fig. 3 together with the results on (In,Mn)As (5). Asymmetric XRD on (115) reflection showed that the (Ga,Mn)As layers were fully strained, which indicates high-quality interface. This result was confirmed by the x-ray analysis of GaAs/(Ga,Mn)As superlattice structures (11). As can be seen from Fig. 3, a increases linearly with xfollowing Vegard's law. The extrapolated lattice constants for zincblende MnAs (0.598 nm) are in good agreement with the MnAs lattice constant extrapolated from the (In,Mn)As side (0.601 nm). This lattice constant of hypothetical zincblende MnAs has been reproduced by a recent first-principle calculation (12). The agreement suggests that all of the Mn atoms were incorporated in the zincblende alloy, which was confirmed by Shioda et al. (13), who showed that Mn is indeed substitutionally incorporated into the Ga sublattice by extended x-ray absorption fine-structure measurement.

Figure 3

Lattice constant a versus Mn composition x in (Ga1–xMnx)As films and in (In1–x,Mnx)As films. Extrapolation from the two end materials leads to the same point of 0.6 nm, which is believed to be the lattice constant of hypothetical zincblende MnAs.

Magnetic Properties

Magnetization measurements with a SQUID (superconducting quantum interference device) magnetometer showed the presence of ferromagnetic order in the (Ga,Mn)As films at low temperatures (7). Sharp, square hysteresis loops, indicating a well-ordered ferromagnetic structure, appeared in the magnetization (M) versus magnetic field (B) curves whenB was applied in the plane of the film. This sharp hysteresis was followed by a “paramagnetic” increase that appeared to follow a Brillouin function as B was further increased. This latter response seems to correlate with the transport properties of the films discussed below; the most metallic sample showed a negligibly small paramagnetic contribution, whereas in insulating samples, the paramagnetic contribution reached almost 50% of total saturation magnetization (14).

When the magnetic field was applied perpendicular to the sample surface, an elongated magnetization with little hysteresis was obtained, indicating that the easy-axis for magnetization was not perpendicular to the plane, but in the plane. This result is quite different from that observed in (001) (In,Mn)As, where the perpendicular direction was the only observed easy-axis (15). This difference may be explained by the magneto-elastic effect; the (In,Mn)As layers were under biaxial tensile strain, which makes the lattice spacing perpendicular to the surface smaller than the one in the plane, whereas the present (Ga,Mn)As layers were under compressive strain, which makes the in-plane lattice constant smaller than the perpendicular one. A perpendicular easy-axis was confirmed in the (Ga,Mn)As layer with biaxial tensile stress grown on an (In,Ga)As buffer layer, supporting the present explanation (16).

The low-temperature saturation magnetization,M S, of the (Ga,Mn)As films was consistent with the spin of Mn S = 5/2, although it is difficult to determine S from these experiments alone because of the error involved in determining x and the nonuniformity ofx over the sample.

Magnetotransport Properties

The dependence on temperature T (2 to 300 K) and magnetic field B (up to 7 T) of sheet resistanceR sheet and Hall resistanceR Hall of 150- to 200-nm (Ga,Mn)As layers were measured with a standard dc transport measurement setup. The temperature dependence of R sheet in samples with intermediate Mn composition (x from 0.035 to 0.053) showed that they were on the metal side of the metal-insulator transition, whereas low- and high-x samples were on the insulator side. Although measurements were done on a number of metallic as well as insulating samples, in order to avoid complications arising from the localization effects, I concentrate here on the metallic samples, especially the one with x = 0.053; results for other metallic samples were essentially the same.

RHall can be expressed as Embedded Image(1) where R0 is the ordinary (normal) Hall coefficient, RS is the anomalous Hall coefficient, d is the sample thickness, and M is the magnetization of the sample. RS is proportional to Rsheet in the present samples [skew scattering (17)] and thusRS/d =cRsheet, where c is a constant. Because the anomalous Hall term is the dominant term up to room temperature, M of the sample can be determined fromRHall [M ∼ (1/c)RHall/Rsheet]. In order to determine the conduction type and the carrier concentration, the ordinary Hall coefficient was measured as the slope of theRHall-B curve at low temperature under high magnetic field, where M saturates.

The results of magnetotransport measurements on a sample withx = 0.053 are shown in Fig. 4, A and B. The T andB dependence of R Hall reflects that of M, confirming the dominating contribution of the anomalous Hall effect. The sheet resistivityR sheet first increases as T decreases with an increase of negative magnetoresistance (a decrease in resistance with increasing B). The zero-field resistivity peaks at around T C and then decreases. The negative magnetoresistance also peaks at T C. Using Arrott plots, in which (R Hall/R sheet)2is plotted against [B/(R Hall/R sheet)] at each temperature to obtain a quantity that is proportional to the saturation magnetization M S from the extrapolated intercept (Fig. 4C; M S is zero when the intercept is at the origin), the T dependence ofM S and T C can be determined; for the present sample T C = 110 K, which was the highest T C for the studied samples. Note thatR Hall/R sheet is proportional to M. The T dependence ofM S can be fitted with a standard Brillouin function, which seems to show that the ferromagnetism in (Ga,Mn)As can be understood in the framework of a mean-field theory. The paramagnetic Curie temperature θ was obtained from the T dependence of the inverse of the zero field slope ofR Hall/R sheet(proportional to susceptibility χ). A straight line characteristic of the Curie-Weiss law was obtained. For all of the samples, θ was very close to T C. The slope ofR Hall-B measured at 10 K revealed that the conduction type was p-type with a hole concentration of 1.0 × 1020 cm–3.

Figure 4

Magnetic field dependence of (A) the Hall resistivity R Hall and (B) the sheet resistivity R sheet of (Ga,Mn)As (x = 0.053) with temperature as a parameter. Because the anomalous Hall term proportional to magnetization is dominant,R Hall reflects the field and temperature dependence of magnetization perpendicular to the plane. From (A) and (B), and knowing thatR Hall/R sheet is proportional to magnetization, the saturation magnetization and magnetic transition temperature can be determined with the Arrott plot shown in (C), which shows that the transition temperature is 110 K.

The R sheet-T curve showed a maximum at around T C, which moved to higher Twith increasing B. I have attributed this behavior ofR sheet to the scattering of carriers by magnetic fluctuation through exchange interactions (18), which has been observed in magnetic semiconductors (19), as opposed to metal-insulator transition proposed by Van Esch et al.(20). This interpretation is further supported by the high-temperature part of R sheet being proportional to χT, as expected from the critical scattering. The observed negative magnetoresistance can be understood as the reduction of scattering by aligning the spins byB. Well above T C, the Bdependence of R sheet can be fit to the following critical scattering resistivity formula Embedded Image(2) where k F is the Fermi wavevector, n is the hole concentration, mis the effective mass, Γ is the p-d exchange,n s is the Mn concentration, S = 5/2, <S> is the average spin on Mn, e is the charge of an electron, and h is Planck's constant (21). By using the measured hole concentration and the effective mass of 0.5 m 0 (the free electron mass), the fit of Eq. 2 to experimental results for all of the metallic samples yields Γ = 150 ± 40 eV Å3 orN 0β ≈ 3.3 eV in terms ofN 0β commonly used to describep-d interaction in DMSs. Typicalp-d exchange (N 0β) in II-VI DMSs is about 1 eV. The origin of this large exchange is not clear at the moment; the effect of weak localization, which enhances the ferromagnetic interaction through electron-electron interaction, might have to be considered, and in the immediate vicinity of the metal-insulator transition, the description of spin-disorder scattering may have to be modified. Large p-d exchange in GaAs doped with Mn was reported by Szczytko et al.(22), who investigated magneto-optical properties. On the other hand, smaller exchange was inferred from photoemission experiments on (Ga,Mn)As (23).

Origin of Ferromagnetism

In the absence of holes, the magnetic interaction among Mn has been shown to be antiferromagnetic in n-type (In,Mn)As (24) and in fully carrier compensated (Ga,Mn)As (25). These results show that the ferromagnetic interaction is hole induced. I have examined whether the ferromagnetism fits into the framework of the Ruderman-Kittel-Kasuya-Yosida (RKKY) interaction, which was shown to be responsible for the carrier-induced ferromagnetism in a IV-VI compound (Pb,Sn,Mn)Te (26). T C can be calculated from the exchange constant, and the hole concentration can be determined from the magnetotransport measurements. Although the result depends slightly on the cut-off length of the RKKY interaction, the calculated T C was in good agreement with the experimentally determined T C (18). Because of this quantitative agreement, I believe the RKKY interaction is most likely responsible for the appearance of ferromagnetism in (Ga,Mn)As. The oscillating nature of the RKKY interaction does not show up in the magnetism, presumably because of the low hole concentration. The sign of RKKY interaction is, in effect, only ferromagnetic because the first zero of the oscillation, beyond which the interaction changes its sign and becomes antiferromagnetic, occurs at a much greater distance (because of the low hole concentration) than the cut-off length of the interaction.

Ferromagnetism observed in insulating samples can probably be understood in the same framework, the carrier-mediated and RKKY-like interaction. In insulating samples close to the metal-insulator transition, the localization length is not extended over the sample size (millimeters), but is still quite long in comparison to the length scale of magnetic interactions (nanometers). Thus, the RKKY-like interaction can still be in effect in the insulating samples.

The understanding of the ferromagnetism of (Ga,Mn)As is not adequate, however. There are issues remaining to be studied such as to what extent the “pure” RKKY interaction is applicable to the present material system: It was pointed out (4), for example, that the behavior of the critical scattering may be qualitatively different when the spin-spin interaction is of long range (present case) as opposed to the short range interaction (magnetic semiconductors).

Spin-Dependent Resonant Tunneling

The fabrication of a ferromagnetic semiconductor compatible with the well-established GaAs-AlAs lattice–matched heterostructure system makes it possible to probe what can be achieved by combining semiconductor heterostructures and ferromagnetism. When semiconductors become ferromagnetic, spin splitting of the conduction as well as of the valence bands occurs because ofs-d and p-d exchange interactions. I have chosen AlAs/GaAs/AlAs double-barrier resonant tunneling diode (RTD) structures with ferromagneticp-type (Ga,Mn)As on one side and p-type GaAs on the other to see if holes of one spin type can be filtered using the energy difference of spin splitting in ferromagnetic (Ga,Mn)As (27).

The structure I studied consists of (from the surface side) 150-nm-thick (Ga0.965Mn0.035)As; 15-nm undoped GaAs spacer; 5-nm undoped AlAs barrier; 5-nm undoped GaAs quantum well; 5-nm undoped AlAs barrier; 5-nm undoped GaAs spacer; 150-nm Be-doped GaAs (p = 5 × 1017cm–3); 150-nm Be-doped GaAs (p = 5 × 1018 cm–3); and p+ GaAs substrates. All of the layers were grown at 650°C except for the last (Ga,Mn)As layer, which was grown at 250°C. A schematic zero-bias valence band diagram of the structures is depicted in the inset of Fig. 5. T C of the (Ga,Mn)As layer is expected to be ∼70 K.

Figure 5

Derivative of current (dI/dV) versus voltage (V) of a resonant tunneling diode with ferromagnetic (Ga,Mn)As emitter. The labeling indicates the relevant resonant state in the GaAs well. When holes are injected from the (Ga,Mn)As side (positive bias), a spontaneous splitting of resonant peak HH2 is observed below 80 K. The transition temperature of (Ga,Mn)As is expected to be 70 K. The splitting is attributed to spin splitting of (Ga,Mn)As valence band states.

As seen in Fig. 5, a total of six peaks have been observed in the dI/dV versus V curve of the present RTD. Each label in Fig. 5 indicates the resonance level in the GaAs well. When holes were injected from the (Ga,Mn)As side (positive bias), a spontaneous resonant peak splitting of the peak labeled HH2 was observed below T C of (Ga,Mn)As without applying a magnetic field, as indicated in Fig. 5. The magnitude of the splitting is shown to be proportional to M Scalculated from the Brillouin function showing the origin of peak splitting as the spin splitting in the valence band of ferromagnetic (Ga,Mn)As. I therefore believe that the splitting observed in theI-V curves is due to the spin splitting of the valence band associated with the development of spontaneous magnetization in (Ga,Mn)As. The observation of splitting suggests that Fermi energies of the spin-split holes are greater than the energy separation of the spin-split bands and that the holes are not fully spin-polarized in (Ga,Mn)As. The reason why only HH2 shows a pronounced splitting is not understood at present. This RTD result shows the possibility of filtering one type of spin and injecting it into a nonmagnetic semiconductor by using an RTD energy filter combined with the energy difference of the spontaneous spin-split bands.

Interlayer Magnetic Interactions

Motivated by intense ongoing research on metallic multilayers (28), carrier-mediated magnetic interactions between two ferromagnetic (Ga,Mn)As layers separated by a nonmagnetic semiconducting layer was studied using (Ga,Mn)As-based heterostructures (29). The all-semiconductor ferromagnetic/nonmagnetic/ferromagnetic trilayer structures studied here consist of a 30-nm (Ga,Mn)As (x= 0.04) layer and a 30-nm (Ga,Mn)As (x = 0.02) layer separated by a nonmagnetic(Al,Ga)As layer. The thickness of the intermediary layer was fixed to 10 monolayers, and the Al composition was varied (x Al = 0.16 and 0.29), which varied the barrier in the valence band. Magnetic measurements (M-B curves, with B applied in the plane) revealed that the two layers were magnetically decoupled and that the M-B curve was a simple addition of the two individual M-B curves measured on separately grown samples for x Al = 0.29. However, forx Al = 0.16, the magnetization curve showed only one step, indicating that the two magnetic layers were now ferromagnetically coupled. The magnetic coupling between the two ferromagnetic (Ga,Mn)As films separated by a nonmagnetic GaAs layer was also shown to be a function of thickness of the intermediary GaAs layer. Both sets of results indicate the critical role of the holes in the intermediary layer on the magnetic coupling. This result is consistent with the RKKY interaction as the origin of the magnetic coupling in the present material system.

Prospects of Spin-Related Phenomena in Semiconductors

Driven by the thrust for faster and denser integrated circuits, semiconductor technology has experienced a continuous reduction in its working dimension, which now has reached a few tens of atomic spacing, if not less, at the most advanced structures. Spin of carriers become increasingly important in these small structures because the exchange interaction can become appreciable, even if the structure is made of nonmagnetic semiconductors (30, 31). In order to take advantage of this trend and use the spin degree of freedom in semiconductors, one has to be able to create, sustain, control, and detect the spin polarization of carriers. The most straightforward way to create spin polarization electrically is by “spin-injection,” that is, by injection of spin-polarized carriers. To do this with ferromagnetic metal/semiconductor junctions has not been easy, presumably because of the presence of scattering at the Schottky barrier interface, although tunneling from a ferromagnetic metal through vacuum into a semiconductor was shown to provide a high degree of polarization (32) and, more recently, room-temperature operation of Si-based spin-valve transistors was demonstrated using spin-dependent transport over Schottky barriers (33). A very good interface between ferromagnet and semiconductor is critical for this application, and (Ga,Mn)As appears to be a promising candidate. How long the injected spin can exist depends on the spin relaxation time, which can be quite long in lightly doped nonmagnetic semiconductors (34). For control of spin, carrier-induced ferromagnetism might be used; by using field-effect to control the carrier density, ferromagnetism may be turned on and off. In fact, photo-generated carrier was recently used to induce ferromagnetism in (In,Mn)As (35). Detection requires spin-selective junction, which can again be provided by ferromagnetic materials with good interface to semiconductors.

Conclusion

The magnetic element Mn has been introduced into the nonmagnetic host-lattice of GaAs, widely used in semiconductor electronics, in excess of its solubility limit by low-temperature MBE. In this homogeneous alloy of GaAs and Mn—(Ga,Mn)As—Mn occupies Ga sites and provides magnetic moments as well as holes, which makes (Ga,Mn)As conducting. The hole-mediated ferromagnetic interaction results in ferromagnetism with a transition temperature as high as 110 K. (Ga,Mn)As can be grown on GaAs-related heterostructures coherently, which makes it possible to bring ferromagnetism and semiconductor heterostructure together. (Ga,Mn)As-based RTDs revealed the possibility of spin-filtering, and all-semiconductor ferromagnetic/nonmagnetic/ferromagnetic trilayer structures were used to investigate the interlayer magnetic interaction in semiconductors. The new III-V–based DMSs can thus be used to explore a new field in semiconductor physics and technology, where both semiconducting and magnetic properties play critical roles.

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