Dilute Magnetic GaN, SiC and Related Semiconductors

 

Jihyun Kim and F.Ren

Department of Chemical Engineering

University of Florida,Gainesville,FL 32611,USA

 

S.J. Pearton, C.R. Abernathy,M.E.Overberg and G.T.Thaler

Department of Materials Science and Engineering

University of Florida, Gainesville, FL 32611 USA

 

Y.D. Park

Center for Strongly Correlated Materials Research

Seoul National University, Seoul, 151-747 Korea

 

ABSTRACT

  Recent results on achieving ferromagnetism in transition metal-doped GaN,SiC and related materials are discussed.While current generations of semiconductor electronic and photonic devices utilize the charge on electrons and holes in order to perform their specific functionality such as signal processing or light emission,the field of semiconductor spintronics seeks to exploit the spin of charge carriers in new generations of transistors, lasers and integrated magnetic sensors. There is strong potential for new classes of ultra-low power, high speed memory, logic and photonic devices. The utility of such devices depends on the availability of materials with practical magnetic ordering temperatures and most theories predict that the Curie temperature will be a strong function of bandgap.In this chapter we review  the field of wide bandgap dilute magnetic semiconductors,such as GaN,SiC and related materials, exhibiting room temperature ferromagnetism, the origins of the magnetism and its potential applications.

Introduction

            Two of the most successful technologies in existence today have created the Si integrated circuit (IC’s) industry and the data storage industry.  Both continue to advance at a rapid pace.  In the case of IC’s, the number of transistors on a chip doubles about every 18 months according to Moore’s Law.  For magnetic hard disk drive technology, a typical desk-top computer drive today has a 40 Gbyte/disk capacity, whereas in 1995 this capacity was ~1 Gbyte/disk.  Since 1991, the overall bit density on a magnetic head has increased at an annual rate of 60-100% and is currently ~10.7 Gbits/in2 [1].  The integrated circuits operate by controlling the flow of carriers through the semiconductor by applied electric fields.  The key parameter therefore is the charge on the electrons or holes.  For the case of magnetic data storage, the key parameter is the spin of the electron, as spin can be thought of as the fundamental origin of magnetic moment. The characteristics of IC’s include high speed signal processing and excellent reliability, but the memory elements are volatile (the stored information is lost when the power is switched-off, as data is stored as charge in capacitors (i.e. DRAMs).  A key advantage of magnetic memory technologies is that they are non-volatile since they employ ferromagnetic materials which by nature have remanence.

            The emerging field of semiconductor spin transfer electronics (spintronics) seeks to exploit the spin of charge carriers in semiconductors.  It is widely expected that new functionalities for electronics and photonics can be derived if the injection, transfer and detection of carrier spin can be controlled above room temperature.  Among this new class of devices are spin transistors operating at very low powers for mobile applications that rely on batteries, optical emitters with encoded information through their polarized light output, fast non-volatile semiconductor memory and integrated magnetic/electronic/photonic devices (“electromagnetism-on-a-chip”).  Since the magnetic properties of ferromagnetic semiconductors are a function of carrier concentration in the material in many cases, then it will be possible to have electrically or optically-controlled magnetism through field-gating of transistor structures or optical excitation to alter the carrier density.  This novel control of magnetism has already been achieved electronically and optically in an InMnAs metal-insulator semiconductor structure at low temperatures [2,3] and electronically in Mn:Ge [4].  A number of recent reviews have covered the topics of spin injection, coherence length and magnetic properties of materials systems such as (Ga,Mn)As [5-7], (In,Mn)As [5-7] and (Co,Ti)O2 [8] and the general areas of spin injection from metals into semiconductors and applications of the spintronic phenomena [9-12].  The current interest in magnetic semiconductors can be traced to difficulties in injecting spins from a ferromagnetic metal into a semiconductor [13,14], which idea can be traced to fruitful research in epitaxial preparation of ferromagnetic transitional metals on semiconductor substrates [15].  A theory first proposed by Schmidt et al.[16] points out that due to the dissimilar materials properties of a metal and semiconductor, an efficient spin injection in the diffusive transport regime is difficult unless the magnetic material is nearly 100% spin polarized – i.e. half-metallic[17].  Although there have been recent reports of successful and efficient spin injection from a metal to a semiconductor even at room temperature by ballistic transport (i.e. Schottky barriers and tunneling) [18], the realization of functional spintronic devices requires materials with ferromagnetic ordering at operational temperatures compatible with existing semiconductor materials.

Materials Selection

            There are two major criteria for selecting the most promising materials for semiconductor spintronics.  First, the ferromagnetism should be retained to practical temperatures (i.e. >300 K).  Second, it would be a major advantage if there were already an existing technology base for the material in other applications.  Most of the work in the past has focused on (Ga,Mn)As and (In,Mn)As.  There are indeed major markets for their host materials in infra-red light-emitting diodes and lasers and high speed digital electronics (GaAs) and magnetic sensors (InAs). Most of the past attention on ferromagnetic semiconductors focussed on the (Ga,Mn)As[19-42]  and (In,Mn)As[43-50] systems.  In samples carefully grown single-phase by Molecular Beam Epitaxy (MBE), the highest Curie temperatures reported are ~110 K for (Ga,Mn)As and ~ 35 K for (In,Mn)As.  For ternary alloys such as (In0.5Ga0.5)0.93Mn0.07As, the Curie temperature is also low ~110 K[51].  A tremendous amount of research on these materials systems has led to some surprising results, such as the very long spin lifetimes and coherence times in GaAs[4] and the ability to achieve spin transfer through a heterointerface[52-69], either of semiconductor/semiconductor or metal-semiconductor.  One of the most effective methods for investigating spin-polarized transport is by monitoring the polarized electroluminescence output from a quantum-well Light-emitting diode into which the spin current is injected.  Quantum selection rules relating the initial carrier spin polarization and the subsequent polarized optical output can provide a quantitative measure of the injection efficiency [67,69,70].

There are a number of essential requirements for achieving practical spintronic devices in addition to the efficient electrical injection of spin-polarized carriers.  These include the ability to transport the carriers with high transmission efficiency within the host semiconductor or conducting oxide, the ability to detect or collect the spin-polarized carriers and to be able to control the transport through external means such as biasing of a gate contact on a transistor structure.  The observation of spin current-induced switching in magnetic heterostructures is an important step in realizing practical devices[71] Similarly, Nitta et.al [72]demonstrated that a spin-orbit interaction in a semiconductor quantum well could be controlled by applying a gate voltage.  These key aspects of spin injection, spin-dependent transport, manipulation and detection form the basis of current research and future technology.  The use of read sensors based on metallic spin valves in disk drives for magnetic recording is already a $US100 B per year industry.  It should also be pointed out that spintronic effects are inherently tied to nanotechnology, because of the short (~1 nm) characteristic length of some of the magnetic interactions.  Combined with the expected low power capability of spintronic devices, this should lead to extremely high packing densities for memory elements.  A recent review of electronic spin injection, spin transport and spin detection technologies has recently been given by Buhrman[7], as part of a very detailed and comprehensive study of the status and trends of research into spin electronics in Japan, Europe and the US.  The technical issues covered fabrication and characterization of magnetic nanostructures, magnetism and spin control in these structures, magneto-optical properties of semiconductors and magneto-electronics and devices.  The non-technical issues covered included industry and academic cooperation and long-term research challenges.  The panel findings are posted on the web site[7].

In this review we focus on a particular and emerging aspect of spintronics, namely recent developments in achieving practical magnetic ordering temperatures in technologically useful semiconductors[73-79].While the progress in synthesizing and controlling the magnetic properties of III-arsenide semiconductors has been astounding, the reported Curie temperatures are too low to have significant practical impact.  A key development that focused attention on wide bandgap semiconductors as being the most promising for achieving high Curie temperatures was the work of Dietl et al.[80].They employed the original Zener model of ferromagnetism[81] to predict TC values exceeding room temperature for materials such as GaN and ZnO containing 5% of Mn and a high hole concentration (3.5x1020 cm-3 ).

 Other materials for which room temperature ferromagnetism has been reported include (Cd,Mn)GeP2 [74], (Zn,Mn)GeP2 [75], ZnSnAs2 [76], (Zn,Co)O [77] and (Co,Ti)O2 [8,78] as well as Eu chalcogenides and others that have been studied in the past [79].  Some of these chalcopyrites and wide bandgap oxides have interesting optical properties, but they lack a technology and experience base as large as that of most semiconductors.

            The key breakthrough that focused attention on wide bandgap semiconductors as being the most promising for achieving practical ordering temperatures was the theoretical work of Dietl et al. [80].  They predicted that cubic GaN doped with ~5at.% of Mn and containing a high concentration of holes (3.5x1020 cm-3) should exhibit a Curie temperature exceeding room temperature.  In the period following the appearance of this work, there has been tremendous progress on both the realization of high-quality (Ga,Mn)N epitaxial layers and on the theory of ferromagnetism in these so-called dilute magnetic semiconductors (DMS).  The term DMS refers to the fact that some fraction of the atoms in a non-magnetic semiconductor like GaN are replaced by magnetic ions.  A key, unanswered question is whether the resulting material is indeed an alloy of (Ga,Mn)N or whether it remains as GaN with clusters, precipitates or second phases that are responsible for the observed magnetic properties[82]

Mechanisms of Ferromagnetism

            Figure 1 shows some of the operative mechanisms for magnetic ordering in DMS materials.  Two basic approaches to understanding the magnetic properties of dilute magnetic semiconductors have emerged.  The first class of approaches is based on mean-field theory which originates in the original model of Zener [81].  The theories that fall into this general model implicitly assume that the dilute magnetic semiconductor is a more-or-less random alloy, e.g. (Ga,Mn)N, in which Mn substitutes for one of the lattice constituents.  The second class of approaches suggests that the magnetic atoms form small (a few atoms) clusters that produce the observed ferromagnetism [82].  A difficulty in experimentally verifying the mechanism responsible for the observed magnetic properties is that depending on the growth conditions employed for growing the DMS material, it is likely that one could readily produce samples that span the entire spectrum of possibilities from single-phase random alloys to nanoclusters of the magnetic atoms to precipitates and second phase formation.  Therefore, it is necessary to decide on a case-by-case basis which mechanism is applicable.  This can only be achieved by a careful correlation of the measured magnetic properties with materials analysis methods that are capable of detecting other phases or precipitates.  If, for example, the magnetic behavior of the DMS is characteristic of that of a known ferromagnetic second phase (such as MnGa or Mn4N in (Ga,Mn)N), then clearly the mean field models are not applicable.  To date, most experimental reports concerning room temperature ferromagnetism in DMS employ x-ray diffraction, selected-area diffraction patterns, transmission electron microscopy, photoemission or x-ray absorption (including extended x-ray absorption fine structure, EXAFS, as discussed later) to determine whether the magnetic atoms are substituting for one of the lattice constituents to form an alloy.  Given the level of dilution of the magnetic atoms, it is often very difficult to categorically determine the origin of the ferromagnetism.  Indirect means such as SQUID magnetometer measurements to exclude any ferromagnetic inter-metallic compounds as the source of magnetic signals and even the presence of what is called the anomalous or extraordinary Hall effect, that have been widely used to verify a single phase system, may be by itself insufficient to characterize a DMS material.  It could also certainly be the case that magnetically-active clusters or second phases could be present in a pseudo-random alloy and therefore that several different mechanisms could contribute to the observed magnetic behavior.  There is a major opportunity for the application of new, element- and lattice position-specific analysis techniques, such as the various scanning tunneling microscopies and Z-contrast scanning transmission electron microscopy (Z-contrast STEM) amongst others for revealing a deeper microscopic understanding of this origin of ferromagnetism in the new DMS materials.

            The mean field approach basically assumes that the ferromagnetism occurs through interactions between the local moments of the Mn atoms, which are mediated by free holes in the material.  The spin-spin coupling is also assumed to be a long-range interaction, allowing use of a mean-field approximation [80,83,84].  In its basic form, this model employs a virtual-crystal approximation to calculate the effective spin-density due to the Mn ion distribution.  The direct Mn-Mn interactions are antiferromagnetic so that the Curie temperature, TC, for a given material with a specific Mn concentration and hole density (derived from Mn acceptors and/or intentional shallow level acceptor doping), is determined by a competition between the ferromagnetic and anti-ferromagnetic interactions.  In the presence of carriers, TC is given by the expression [80,85]

where NOXeff is the effective spin concentration, S is the localized spin state, b is the p-d exchange integral, AF the Fermi liquid parameter, PS the total density of states, kB is Boltzmann’s constant and TAF describes the contribution of antiferromagnetic interactions.  Numerous refinements of this approach have appeared recently, taking into account the effects of positional disorder [86,87], indirect exchange interactions [88], spatial inhomogeneities and free-carrier spin polarization [89,90].  Figure 2 shows a compilation of the predicted TC values, together with the classification of the materials (eg. Group IV semiconductor, etc.).  In the subsequent period after appearance of the Dietl et al [80] paper, remarkable progress has been made on the realization of materials with TC values at or above room temperature.

             The mean-field model and its variants produces reliable estimates of TC for materials such as (Ga,Mn)As and (In,Mn)As and predicts that (Ga,Mn)N will have a value above room temperature [80].  Examples of the predicted ferromagnetic transition temperatures for both (Ga,Mn)As and (Ga,Mn)N are shown in Figure 3 for four different variants of the mean-field approach [91].  These are the standard mean-field theory (TCMF), a version that accounts for the role of Coulomb interactions with holes in the valence band (exchange-enhanced, TCX), another version that accounts for correlations in Mn ion orientations (collective, TCcoll) or an estimate based on where excited spin waves cancel out the total spin of the ground state (TCest) [91].  Note that the dependence of any of the calculated TC values on hole density in the material is much steeper for (Ga,Mn)As than for (Ga,Mn)N.  The range of predicted values for GaAs has a much higher distribution than for GaN.  This data emphasizes the point that the mean field theories produce fairly reliable predictions for (Ga,Mn)As, but at this stage are not very accurate for (Ga,Mn)N.

            A second point largely overlooked in the theoretical work to date is that fact that the assumed hole densities may not be realistic.  While GaAs can be readily doped with shallow acceptors such as C to produce hole densities of around 1021 cm-3 [92] and the Mn acceptors also contribute holes, the p-doping levels in GaN are limited to much lower values under normal conditions.  For example, the ionization level (Ea) of the most common acceptor dopant in GaN, namely Mg, is relatively deep in the gap (EV+0.17eV).  Since the number of holes (P) is determined by the fraction of acceptors that are actually ionized at a given temperature T through a Boltzmann factor

then for Mg at room temperature only a few percent of acceptors are ionized.  While the Mg acceptor concentration in GaN can exceed 1019 cm-3, a typical hole concentration at 25°C is      P~3x1017 cm-3.   Initial reports of the energy level of Mn in GaN show it is very deep in the gap, Ev + 1.4 eV [93], and thus would be an ineffective dopant under most conditions.  Some strategies for enhancing the hole concentration do exist, such as co-doping both acceptors and donors to reduce self-compensation effects [94] or the use of selectively-doped AlGaN/GaN superlattices in which there is transfer of free holes from Mg acceptors in the AlGaN barriers to the GaN quantum wells [95].  These methods appear capable under optimum conditions of increasing the hole density in GaN to >1018 cm-3 at 25°C.

            A further issue that needs additional exploration in the theories is the role of electrons, rather than holes, in stabilizing the ferromagnetism in DMS materials.  All of the reports of ferromagnetism in (Ga,Mn)N, for example, occur for material that is actually n-type.  Since the material has to be grown at relatively low temperatures to avoid Mn precipitation and therefore only Molecular Beam Epitaxy (MBE) can be used, there is always the possibility of unintentional n-type doping from nitrogen vacancies, residual lattice defects or impurities such as oxygen.  Therefore stoichiometry effects, crystal defects or unintentional impurities may control the final conductivity, rather than Mn or the intentionally-introduced acceptor dopants.  Once again, this is much less of an issue in materials such as GaAs, whose low temperature growth is relatively well-understood and controlled.

            While most of the theoretical work for DMS materials has focused on the use of Mn as the magnetic dopant, there has been some progress on identifying other transition metal atoms that may be effective.  Figure 4 shows the predicted stability of ferromagnetic states in GaN doped with different 3d transition metal atoms [96].  The results are based on a local spin-density approximation which assumed that Ga atoms were randomly substituted with the magnetic atoms and did not take into account any additional carrier doping effects.  In this study it was found that (Ga,V)N and (Ga,Cr)N showed stable ferromagnetism for all transition metal concentrations whereas Fe, Co or Ni doping produced spin-glass ground states [96].  For the case of Mn, the ferromagnetic state was the lowest energy state for concentrations up to ~20%, whereas the spin-glass state became the most stable at higher Mn concentrations.

A.       (Ga,Mn)P

Ferromagnetism above room temperature in (Ga,Mn)P has been reported for two different methods of Mn incorporation, namely ion implantation[97] and doping during MBE growth[97,98] The implantation process is an efficient one for rapidly screening whether particular combinations of magnetic dopants and host semiconductors are promising in terms of ferromagnetic properties.  We have used implantation to introduce ions such as Mn, Fe and Ni into a variety of substrates, including GaN, SiC and GaP

The temperature-dependent magnetization of a strongly p-type (p~1020), carbon-doped GaP sample implanted with ~6 at.% of Mn and then annealed at 700C, is shown in Figure 5.  The diamagnetic contribution was subtracted from the background.  A Curie temperature (TC) of ~270 K is indicated by the dashed vertical line, while the inset shows a ferromagnetic Curie temperature, of 236 K.

            Examples of hysteresis loops from MBE grown samples are shown in Figure 6.  The hysteresis could be detected to 330 K.   No secondary phases (such as MnGa or MnP) or clusters were determined by transmission electron microscopy, x-ray diffraction or selected-area diffraction pattern analysis.

           While mean-field theories predict relatively low Curie temperatures (<110K) for (Ga,Mn)P [97,98], recent experiments show ferromagnetism above 300K [97,98].  In other respects, the magnetic behavior of the (Ga,Mn)P was consistent with mean-field predictions.  For example, the magnetization versus temperature plots showed a more classical concave shape than observed with many DMS materials.  In addition, the Curie temperature was strongly influenced by the carrier density and type in the material, with highly p-type samples showing much higher values than n-type or undoped samples.  Finally, the Curie temperature increased with Mn concentration up to ~6 at.% and decreased at higher concentrations [98].  No secondary phases or clusters could be detected by transmission electron microscopy, x-ray diffraction or selected area diffraction patterns.  Similar results were achieved in samples in which the Mn was incorporated during MBE growth or directly implanted with Mn.

            GaP is a particularly attractive host material for spintronic applications because it is almost lattice-matched to Si.  One can therefore envision integration of (Ga,Mn)P spintronic magnetic sensors or data storage elements to form fast non-volatile Magnetic Random Access Memories (MRAM).  Although it has an indirect bandgap, it can be made to luminescence through addition of isoelectric dopants such as nitrogen or else one could employ the direct bandgap ternary InGaP, which is lattice matched to GaAs.  The quaternary InGaAlP materials system is used for visible light-emitting diodes, laser diodes, heterojunction bipolar transistors and high electron mobility transistors.  An immediate application of the DMS counterparts to the component binary and ternary materials in this system would be to add spin functionality to all of these devices.  A further advantage to the wide bandgap phosphides is that they exhibit room temperature ferromagnetism even for relatively high growth temperatures during MBE. 

      Obviously, the Mn can also be incorporated during MBE growth of the (Ga,Mn)P.  The p-type doping level can be separately controlled by incorporating carbon from a CBr4 source while P is obtained from thermal cracking of PH3. A phase diagram for the epi growth of this materials system has been developed and this can be used to tailor the magnetic properties of the (Ga,Mn)P.  For samples grown at 600C with 9.4 at.% Mn, hysteresis is still detectable at 300 K, with a coercive field of ~39 Oe .

            B.        (Ga,Mn)N

            The initial work on this material involved either microcrystals synthesized by nitridization of pure metallic Ga in supercritical ammonia or bulk crystals grown in reactions of Ga/Mn alloys on GaN/Mn mixtures with ammonia at ~1200°C[99].These samples exhibit paramagnetic properties over a broad range of Mn concentrations, as did some of its early MBE-grown films.  By contrast, Figure 7 shows room temperature ferromagnetism from more recent n-type (Ga,Mn)N samples.

          The first reports of the magnetic properties of (Ga,Mn)N involved bulk microcrystallites grown at high temperatures (~1200°C), but while percent levels of Mn were incorporated, the samples exhibited paramagnetic behavior [99].  By sharp contrast, in epitaxial GaN layers grown on sapphire substrates and then subjected to solid state diffusion of Mn at temperatures from 250-800°C for various periods, clear signatures of room temperature ferromagnetism were observed [100,101].  Figure 8 shows anomalous Hall effect data (left) at 323K and the temperature dependence of sheet resistance at zero applied field for 2 different Mn-diffused samples  and an undoped GaN control sample (right)[101].  The Curie temperature was found to be in the range 220-370K, depending on the diffusion conditions.  The use of ion implantation to introduce the Mn produced lower magnetic ordering temperatures [102].

            In (Ga,Mn)N films grown by MBE at temperatures between 580-720°C with Mn contents of 6-9 at.%, magnetization (M) versus magnetic field (H) curves showed clear hysteresis at 300K, with coercivities of 52-85 Oe and residual magnetizations of 0.08-0.77 emu/g at this temperature [103].  Figure 5 shows the temperature dependence of the magnetization for a sample with 9 at.% Mn, yielding an estimated TC of 940K using a mean field approximation.  Note that while the electrical properties of the samples were not measured, they were almost certainly n-type.  As we discussed above, it is difficult to obtain high Curie temperatures in n-type DMS materials according to the mean-field theories and this is something that needs to be addressed in future refinements of these theories.  Room temperature ferromagnetism in n-type (Ga,Mn)N grown by MBE has also been reported by Thaler et al. [104].  In that case, strenuous efforts were made to exclude any possible contribution from the sample holder in the superconducting quantum interference device (SQUID) magnetometer or other spurious effects.  It is also worthwhile to point out that for the studies of (Ga,Mn)N showing ferromagnetic ordering by magnetization measurements, a number of materials characterization techniques did not show the presence of any second ferromagnetic phases within detectable limits.  In addition, the values of the measured coercivities are relatively small.  If indeed there were undetectable amounts of nano-sized clusters, due to geometrical effects, the expected fields at which these clusters would switch magnetically would be expected to be much larger than what has been observed.  Extended x-ray absorption fine structure (EXAFS) measurements performed on (Ga,Mn)N samples grown by MBE on sapphire at temperatures of 400-650°C with Mn concentrations of ~7x1020 cm-3 (i.e. slightly over 2 at.%) are shown in Figure 10 [105].  The similarity of the experimental data with simulated curves for a sample containing this concentration of Mn substituted for Ga on substitutional lattice positions indicates that Mn is in fact soluble at these densities.  In the samples grown at 650°C, £ 1 at.% of the total amount of Mn was found to be present as Mn clusters.  However at lower growth temperatures (400°C), the amount of Mn that could be present as clusters increased up to ~36 at.% of the total Mn incorporated.  The ionic state of the substitutional Mn was found to be primarily Mn(2), so that these impurities act as acceptors when substituting for the Ga with valence three.  However, when the electrical properties of these samples were measured, they were found to be resistive [105].  This result emphasizes how much more needs to be understood concerning the effects of compensation and unintentional doping of (Ga,Mn)N, since the EXAFS data indicated the samples should have shown very high p-type conductivity due to incorporation of Mn acceptors.

            Other reports have also recently appeared on the magnetic properties of GaN doped with other transition metal impurities.  For initially p-type samples directly implanted with either Fe or Ni, ferromagnetism was observed at temperatures of ~200K [106] and 50K [107], respectively.  (Ga,Fe)N films grown by MBE showed Curie temperatures of £100K, with EXAFS data showing that the majority of the Fe was substitutional on Ga sites [108].  (Ga,Cr)N layers grown in a similar fashion at 700°C on sapphire substrates showed single-phase behavior, clear hysteresis and saturation of magnetization at 300K and a Curie temperature exceeding 400K [109].

            Epi growth of (Ga,Mn)N has produced a range of growth conditions producing single-phase material and the resulting magnetic properties[100-107].  In general, no second phases are found for Mn levels below ~10% for growth temperatures of ~750°C.  The (Ga,Mn)N retains n-type conductivity under these conditions. 

            In accordance with most of the theoretical predictions, magnetotransport data showed the anomalous Hall effect, negative magnetoresistance and magnetic resistance at temperatures that were dependent on the Mn concentration.  For example, in films with very low (<1%) or very high (~9%) Mn concentrations, the Curie temperatures were between 10-25 K.  An example is shown in Figure 7 for an n-type (Ga,Mn)N sample with Mn ~7%.  The sheet resistance shows negative magnetoresistance below 150 K, with the anomalous Hall coefficient disappearing below 25 K.  When the Mn concentration was decreased to 3 at.%, the (Ga,Mn)N showed the highest degree of ordering per Mn atom[104]  Figure 7(a) shows hysteresis present at 300 K, while the magnetization as a function of temperature is shown in Figure 7(b).  Data from samples with different Mn concentrations is shown in Figure 7(c) and indicates ferromagnetic coupling, leading to a lower moment per Mn.  Data from field-cooled and zero field –cooled conditions was further suggestive of room temperature magnetization[104].  The significance of these results is that there are many advantages from a device viewpoint to having n-type ferromagnetic semiconductors.

The local structure and effective chemical valency of Mn in MBE-grown (Ga,Mn)N samples has been investigated by Extended X-Ray Absorption Fine Structure[105]. It was concluded that most of the Mn was incorporated substitutionally on the Ga sub-lattice with effective valency close to +2 for samples with ~2 at.% Mn[105]. There was also evidence that a fraction (from 1-36%,depending on growth condition) of the total Mn concentration could be present as small Mn clusters[105].

C.Chalcopyrite Materials

            The chalcopyrite semiconductors are of interest for a number of applications.  For example, ZnGeP2 exhibits unusual non-linear optical properties and can be used in optical oscillators and frequency converters.  ZnSnAs2 shows promise for far-IR generation and frequency converters.  The wide bandgap chalcopyrites ZnGeN2 and ZnSiN2 have lattice parameters close to GaN and SiC, respectively and the achievement of ferromagnetism in these materials would make it possible for direct integration of magnetic sensors and switches with blue/green/UV lasers and light-emitting diodes, UV solar-blind detectors and microwave power electronic devices fabricated in the GaN and SiC.  The bandgap of ZnGexSi1-xN2 varies linearly with composition from 3.2 eV (x=1) to 4.46 eV (x=0).

            Numerous reports of room temperature ferromagnetism in Mn-doped chalcopyrites have appeared.  A compilation of these results and those from transition metal doped GaN and GaP are shown in Table I.  The ZnSnAs2 is somewhat of an anomaly due to its small bandgap, but little theory is available at this point on the chalcopyrites and their expected magnetic properties as a function of bandgap, doping or Mn concentration.  In the only case in which electrical properties were reported, the ZnGeSiN2:Mn was n-type [110].  There is also no information available on the energy level of Mn in the bandgap.

D.SiC

     Very little attention has been paid to potential dilute magnetic semiconductor behavior in SiC, which is at a relatively mature state of development for high power, high temperature electronics.  With its wide bandgap (3.0 eV for the 6H polytype), excellent transport properties and dopability, it would be a good candidate for spintronic applications.  In this section, we report on the structural and magnetic properties of 6H-SiC implanted with Ni, Fe or Mn at doses designed to produce peak concentrations of these elements up to ~5 at.%.  At these concentrations, ferromagnetic ordering temperatures between 50-270 K were observed.

       Bulk, Al-doped (p = 1017 cm-3 at 25°C) 6H-SiC substrates were used for all these experiments.  The samples were directly implanted with Fe, Ni or Mn ions into the Si face at doses from 3-5x1016 cm-2 at 250 keV energy in all cases.  To avoid amorphization the sample temperature was held at ~ 350°C during the implant step.  Calculated ion profiles showed a projected range of ~1300 Å (an example is shown in Figure 11 for Ni at a dose of 3x1016 cm-2).  The peak concentration corresponded to roughly 3 or 5 at.% for the respective doses of 3 and 5x1016 cm-2.  Following implantation, the samples were annealed for 5 mins at 700-1000°C under flowing N2 in a Heatpulse 410T system.  The structural quality of the material was examined by x-ray diffraction (XRD), selected area diffraction pattern (SADP) analysis and cross-sectional transmission electron microscopy (TEM).  The magnetic properties were measured on a Quantum Design SQUID magnetometer.

(a)    Ni implantation

Figure 12 (top) shows a cross-sectional TEM micrograph from a SiC sample implanted with 3x1016 cm-2 Ni+ and annealed at 700°C.  There is a buried band of defects formed at ~2500Å deep, commonly referred to as end-of-range damage.  This consists of a variety of defect types, including dislocation loops and other extended clusters of point defects.  The same basic features were present for the higher dose (5x1016 cm-2) samples and for those annealed at 1000°C, indicating that the damage is stable to at least that temperature.  The bottom of Figure 12 shows the temperature dependence of the difference in magnetization signal between field-cooled and zero field-cooled conditions at a field of 500 Oe.  The large paramagnetic background signal has been subtracted in this data.  The transition temperature is ~50 K based on an extrapolation of where the magnetization is equal to zero within experimental error and similar results were obtained for the higher Ni dose and higher annealing temperature.

A variety of secondary phases could potentially be present in Ni-implanted SiC, including Ni1NixSiy and NixCy components with cubic, hexagonal, orthorhombic, tetragonal, monoclinic or rhombohedral symmetries.  We did not observe any additional peaks in the XRD spectrum of the SiC after Ni implantation, no precipitates visible to the 20 Å resolution of the TEM and no extra spots in the SADP.  As an example, Figure 13 shows a SADP from the implanted region of a sample implanted with 5x1016 cm-2 Ni+ and annealed at 700°C.The rings originate from some polycrystalline regions in the implanted region.

From these results, Ni does not appear to be a promising candidate for producing high transition temperatures in SiC.  There is as yet no theory to guide the choice of magnetic dopants for this material.

(b)   Fe implantation

Figure 14 (top) shows a magnetization curve at 10 K for SiC implanted with 5x1016 cm-2 Fe+ and annealed at 700°C.  Once again the diamagnetic background has been subtracted.  The coercive field is ~ 50 Oe at 10 K, and the magnetization went to zero at ~ 270 K for this Fe+ dose.  The samples implanted with the lower dose of 3x1016 cm-2 were paramagnetic.  The existing theories suggest that the presence of ferromagnetism and its associated ordering temperature are strongly dependent on the concentration of the magnetic ions.  In past work on Mn-implanted GaN we have observed that doses up to 1 at.% lead to paramagnetic behavior, then the ferromagnetic ordering temperature increases with dose up to 3 at.% and then decreases at higher concentrations.  The  bottom part of Figure 14 shows a close-up view of the buried damage layer noted at the end-of-range.  The width of this layer is ~ 260Å and once again no other phases were detected in the implanted region.

(c)    Mn implantation

Figure 15 shows the temperature dependence of the difference between the field cooled (FC) and zero field-cooled (ZFC) magnetization for a SiC sample implanted with 5x1016 cm-2 Mn+ and annealed at 700°C.  The apparent ordering temperature is ~ 250 K at this dose(again based on where the magnetization goes to zero within experimental error), while the samples implanted at the lower dose remained paramagnetic.  The hysteresis loop for the 5x1016 cm-2 sample showed a coercive field of ~ 150 Oe at 10 K.

Table II shows a comparison of the characteristics of the samples implanted with the three different elements.  It is clear that the results for Ni are much less promising than those for Fe and Mn.In summary,the direct implantation of Fe,Ni and Mn into SiC produced significantly different magnetic characteristics.While fairly high apparent ordering temperatures were observed for Mn and Fe(between 250 and 270K),the Ni led to low values of the ordering temperature(about 50K).The origin of the ferromagnetic contributions in implanted SiC is still to be determined, as it is in other materials systems such as (Ga,Mn)N and (Ga,Fe)N that show similar behavior.  Future work should focus also on the effects of carrier density and type on the magnetic properties.  While the Dietl theory requires a high hole concentration for achievement of  high TC values, recent papers indicate that ferromagnetism can be observed in n-type or insulating semiconductors.  The use of ion implantation to introduce the magnetic dopants is attractive because of its versatility in controlling the element implanted and its concentration.

Potential Device Applications

            Previous articles have discussed some spintronic device concepts such as spin junction diodes and solar cells [9], optical isolators and electrically-controlled ferromagnets [10].  The realization of light-emitting diodes with a degree of polarized output has been used to measure spin injection efficiency in heterostructures [111-113].  Such structures can reveal much about spin transport through heterointerfaces after realistic device processing schemes involving etching, annealing and metallization.  The spin transfer in such situations has proven surprisingly robust [114].  It is obviously desirable that spintronic devices are operable at or above room temperature.  As an initial demonstration that (Ga,Mn)N layers can be used as the n-type injection layer in GaN/InGaN blue light-emitting diodes, Figure 16 shows the LED during operation (top) and the spectral output (left).  It is necessary to next establish the extent of any degree of polarization of the light emission, which might be difficult to observe in GaN/InGaN LEDs, since it has been shown that the free exciton components in the EL spectrum contribute mostly to the observed circular polarization of the emitted light [115].  While the expected advantages of spin-based devices include non-volatility, higher integration densities, lower power operation and higher switching speeds, there are many factors still to consider in whether any of these can be realized.  These factors include whether the signal sizes due to spin effects are large enough at room temperature to justify the extra development work needed to make spintronic devices and whether the expected added functionality possible will materialize.

In addition to active and/or optical devices, wide bandgap DMS materials may also be used as passive devices. LeClair et al.[116] have recently shown an artificial half-metallic structure by using a polycrystalline sputtered ferromagnetic semiconductor (EuS) as a tunneling barrier.  This barrier can function as an effective spin filter, since a tunneling electron encounters a differing barrier height depending on its spin below Tc of the barrier material (for EuS, TC ~16.8K).  At low temperatures, the spin-filtering efficiencies were found to be ~90%.  Room temperature DMS materials for these spin-filtering effects could be used to increase magnetoresistance changes in current magnetic tunnel junctions and metallic spin-valve structures. 

Issues to be Resolved

            As described earlier, there are a number of existing models for the observed ferromagnetism in semiconductors.  The near-field models consider the ferromagnetism to be mediated by delocalized or weakly localized holes in the p-type materials.  The magnetic Mn ion provides a localized spin and acts as an acceptor in most III-V semiconductors so that it can also provide holes.  In these models, the TC is proportional to the density of Mn ions and the hole density.  Many aspects of the experimental data can be explained by the basic mean-field model.  However, ferromagnetism has been observed in samples that have very low hole concentrations, in insulating material and more recently in n-type material.  Models in these regimes are starting to appear [117].   

            An alternative approach using local density functional calculations suggests that the magnetic impurities may form small nano-size clusters that produce the observed ferromagnetism [82].  These clusters would be difficult to detect by most characterization techniques.  Clearly there is a need to more fully characterize the materials showing room temperature ferromagnetism and correlate these results to establish on a case-by-case basis which is the operative mechanism and also to refine the theories based on experimental input.  More work is also needed to establish the energy levels of the Mn, whether there are more effective magnetic dopant atoms and how the magnetic properties are influenced by carrier density and type.  Even basic measurements such as how the bandgap changes with Mn concentration in GaN and GaP have not been performed.  The control of spin injection and manipulation of spin transport by external means such as voltage from a gate contact or magnetic fields from adjacent current lines or ferromagnetic contacts is at the heart of whether spintronics can be exploited in device structures and these areas are still in their infancy.  A concerted effort on the physics and materials science of the new dilute magnetic semiconductors is underway in many groups around the world, but fresh insights, theories and characterization methods would greatly accelerate the process.

 

Acknowledgments     

            The work at UF was partially supported by NSF-DMR 0101438 and by the US Army Research Office under grants nos.ARO DAAD 19-01-1-0710 and DAAD 19-02-1-0420,while the work at SNU was partially supported by KOSEF and Samsung Electronics Endowment through CSCMR and by the Seoul National University Research Foundation.  The authors are very grateful to their collaborators A.F. Hebard, N.A. Theodoropoulou, R.G. Wilson, J.M. Zavada, D.P. Norton, S.N.G. Chu, J.S. Lee and Z.G. Khim.

References

1.           see for example http://www.almaden.ibm.com/sst/

2.           H. Ohno, D. Chiba, F. Matsukura, T. Omiya, E. Abe, T. Dietl, Y. Ohno and K. Ohtani, Nature 408, 944 (2000).

3.           A. Oiwa, Y. Mitsumori, R. Moriya, T. Slupinski and H. Munekata, Phys. Rev. Lett. 88, 137202 (2002).

4.           Y.D. Park, A.T. Hanbicki, S.C. Erwin, C.S. Hellberg, J.M. Sullivan, J.E. Mattson, A. Wilson, G. Spanos and B.T. Jonker, Science 295, 651 (2002).

5.           H. Ohno, J. Vac. Sci. Technol. B18, 2039 (2000).

6.           S.A. Wolf, D.D. Awschalom, R.A. Buhrman, J.M. Daughton, S. von Molnar, M.L. Roukes, A.Y. Chtchelkanova and D.M. Treger, Science 294, 1488 (2001).

7.           Von Molnar, S., et al., “World Technology (WTEC) Study on Spin Electronics: Highlights of Recent US Research and Development Activities,” http://www.wtec.org/spin_US_summary.pdf  (2001).

8.           S.A. Chambers, Materials Today, pp. 34-39, April 2002.

9.           S. Das Sarma, American Scientist 89, 516 (2001).

10.       H. Ohno, F. Matsukura and Y. Ohno, JSAP International 5, 4 (2002).

11.       D.D. Awschalom and J.M. Kikkawa, Science 287, 473 (2000).

12.       C. Gould, G. Schmidt, G. Richler, R. Fiederling, P. Grabs and L.W. Molenkamp, L.W., Appl. Surf. Sci. (in press, 2002).

13.       PR. Hammar, B.R. Bennett, M.J. Yang and Mark Johnson,  Phys. Rev. Lett. 83, 203 (1999).

14.       F.G. Monzon, H.X. Tang and M.L. Roukes,  Phys. Rev. Lett. 84, 5022 (2000).

15.       G.A. Prinz, “Magnetic Metal Films on Semiconductor Substrates” in Ultrathin Magnetic Structures II, eds. Heinrich, B. and Bland, J.A.C. (Springer-Verlag, New York, 1994).

16.       G. Schmidt, D. Ferrand, L.W. Molenkamp, A.T. Filip and B.J. van Wees, Phys. Rev. B 62, R4793 (2000).

17.       E.I. Rashba, E. I., Phys. Rev. B 62, R16267 (2000).

18.       H.J. Zhu, M. Ramsteiner, H. Kostial, M. Wassermeier, H.-P. Schönherr and K.H. Ploog, Phys. Rev. Lett. 87, 016601 (2001).

19.       F. Matsukura, H. Ohno, A. Shen and Y. Sugawara, Phys. Rev. B57, R2037 (1998)

20.       H. Ohno, A. Shen, F. Matsukura, A. Oiwa, A. Endo, S. Katsumoto and Y. Iye, Appl. Phys. Lett. 69, 363 (1996).

21.       R. Shioda, K. Ando, T. Hayashi and M. Tanaka, Phys. Rev. B58, 1100 (1998).

22.       Y. Satoh, N. Inoue, Y. Nishikawa and J. Yoshino, Proc. 3rd Symp. Physics and Applications of  Spin-Related Phenomena in Semiconductors, ed. H. Ohno, J. Yoshino and Y. Oka, Nov. 1997, Sendai, Japan, p. 23.

23.       T. Hayashi, M. Tanaka, T. Nishinaga, H. Shimoda, H. Tsuchiya and Y. Otsuka, J. Cryst. Growth 175, 1063 (1997).

24.       A. Van Esch, L. Van Bockstal, J. de Boeck, G. Verbanck, A.S. vas Steenbergen, R.J. Wellman, G. Grietens, R. Bogaerts, F. Herlach and G. Borghs, Phys. Rev. B56, 13103 (1997).

25.       B. Beschoten, P.A. Crowell, I. Malajovich, D.D. Awschalom, F. Matsukura, A. Shen and H. Ohno, Phys. Rev. Lett. 83, 3073 (1999).

26.       M. Tanaka, J. Vac. Sci. Technol. B16, 2267 (1998).

27.       Y. Nagai, T. Kurimoto, K. Nagasaka, H. Nojiri, M. Motokawa, F. Matsukura, T. Dietl and H. Ohno, Jap. J. Appl. Phys. 40, 6231 (2001).

28.       J. Sadowski, R. Mathieu, P. Svedlindh, J.Z. Domagala, J. Bak-Misiuk, J. Swiatek, M. Karlsteen, J. Kanski, L. Ilver, H. Asklund and V. Sodervall, Appl. Phys. Lett. 78, 3271 (2001).

29.       A. Shen, F. Matsukura, S.P. Guo, Y. Sugawara, H. Ohno, M. Tani, A. Abe and H.C. Liu, J. Cryst. Growth 201/202, 379 (1999).

30.       H. Shimizu, T. Hayashi, T. Nishinaga and M. Tanaka, Appl. Phys. Lett. 74, 398 (1999).

31.       B. Grandidier, J.P. Hys, C. Delerue, D. Stievenard, Y. Higo and M. Tanaka, Appl. Phys. Lett. 77, 4001 (2000).

32.       R.K. Kawakami, E. Johnson-Halperin, L.F. Chen, M. Hanson, N Guebels, J.S. Speck, A.C. Gossard and D.D. Awschalom, Appl. Phys. Lett. 77, 2379 (2000).

33.       K. Ando, T. Hayashi, M. Tanaka and A. Twardowski, J. Appl. Phys. 83, 65481 (1998).

34.       D. Chiba, N. Akiba, F. Matsukura, Y. Ohno and H. Ohno, Appl. Phys. Lett. 77, 1873 (2000).

35.       H. Ohno, F. Matsukura, T. Owiya and N. Akiba, J. Appl. Phys. 85, 4277 (1999).

36.       T. Hayashi, M. Tanaka, T. Nishinaga and H. Shimada, J. Appl. Phys. 81, 4865 (1997).

37.       T. Hayashi, M. Tanaka, K. Seto, T. Nishinaga and K. Ando, Appl. Phys. Lett. 71, 1825 (1997).

38.       A. Twardowski, Mat. Sci. Eng. B63, 96 (1999).

39.       T. Hayashi, M. Tanaka and A. Asamitsu, J. Appl. Phys. 87, 4673 (2000).

40.       N. Akiba, D. Chiba, K. Natata, F. Matsukura, Y. Ohno and H. Ohno, J. App. Phys. 87, 6436 (2000).

41.       S.J. Potashnik, K.C. Ku, S.H. Chun, J.J. Berry, N. Samarth and P. Schiffer, Appl. Phys. Lett. 79, 1495 (2001).

42.       G.M. Schott, W. Faschinger and L.W. Molenkamp, Appl. Phys. Lett. 79, 1807 (2001).

43.       H. Munekata, H. Ohno, S. von Molnar, A. Segmuller, L.L. Chang and L. Esaki, Phys. Rev. Lett. 63, 1849 (1989).

44.       Akai, Phys. Rev. Lett. 81, 3002 (1998).

45.       Ohno, H. Munekata, T. Penney, S. von Molnar and L.L Chang, Phys. Rev. Lett. 68, 2864 (1992).

46.       H. Munekata, A. Zaslevsky, P. Fumagalli and R.J. Gambino, Appl. Phys. Lett. 63, 2929 (1993).

47.       S. Koshihara, A. Oiwa, M. Hirasawa, S. Katsumoto, Y. Iye, C. Urano, H. Takagi and H. Munekata, Phys. Rev. Lett. 78, 4617 (1997).

48.       Y.L. Soo, S.W. Huang, Z.H. Ming, Y.H. Kao and H. Munekata, Phys. Rev. B53, 4905 (1996).

49.       A. Oiwa, T. Slupinski and H. Munekata, Appl. Phys. Lett. 78, 518 (2001).

50.       Y. Nishikawa, A. Tackeuchi, M. Yamaguchi, S. Muto and O. Wada, IEEE J. Sel. Topics Quantum Electron. 2, 661 (1996).

51.       H. Munekata, presented at ICCG-13, August 2001.

52.       I. Malajovich, J.M. Kikkawa, D.D. Awschalom, J.J. Berry and N. Samarth, Phys. Rev. Lett. 84, 1015 (2000).

53.       P.R. Hammar, B.R. Bennet, M.Y. Yang and M. Johnson, J. Appl. Phys. 87, 4665 (2000).

54.       A. Hirohata, Y.B. Xu, C.M. Guetler and J.A.C. Bland, J. Appl. Phys. 87, 4670 (2000).

55.       M. Johnson, J. Vac. Sci. Technol. A16, 1806 (1998).

56.       S. Cardelis, C.G. Smith, C.H.W. Barnes, E.H. Linfield and J. Ritchie, Phys. Rev. B 60 7764(1999).

57.       R. Fiederling, M. Kein, G. Rerescher, W. Ossan, G. Schmidt, A. Wang and L.W. Molenkamp, Nature 402, 787 (1999).

58.       G. Borghs and J. De Boeck, Mat. Sci. Eng. B84, 75 (2001).

59.       Y. Ohno, D.K. Young, B. Bescholen, F. Matsukura, H. Ohno and D.D. Awschalom, Nature 402, 790 (1999).

60.       B.T. Jonker, Y.D. Park, B.R. Bennett, H.D. Cheong, G. Kioseoglou and A. Petrou, Phys. Rev. B62, 8180 (2000).

61.       Y.D. Park, B.T. Jonker, B.R. Bennett, G. Itskos, M. Furis, G. Kioseoglou and A. Petrou, App. Phys. Lett. 77, 3989 (2000).

62.       G. Schmidt, D. Ferrand, L.W. Molenkamp, A.T. Filip and B.J. van Wees, Phys. Rev. B62, R4790 (2000).

63.       Y.Q. Jin, R.C. Shi and S.J. Chou, IEEE Trans. Magn. 32, 4707 (1996).

64.       C.M. Hu, J. Nitta, A. Jensen, J.B. Hansen and H. Takayanagai, Phys. Rev. B63, 125333 (2001).

65.       F.G. Monzon, H.X. Tang and M.L. Roukes, Phys. Rev. Lett. 84, 5022 (2000).

66.       S. Gardelis, C.G. Smith, C.H. W. , F. Matsukura and H. Ohno, Jap. J. Appl. Phys. 40, L1274 (2001).

67.       H. Breve, S. Nemeth, Z. Liu, J. De Boeck and G. Borghs, J. Magn. Magn. Materials. 226-0230, 933 (2001).

68.       H. Munekata, H. Ohno, S. von Molnar, A. Segmuller, L.L. Chang and L. Esaki, Phys. Rev. Lett. 63, 1849 (1989).

69.       H.J. Zhu, M. Ramsteiner, H. Kostial, M. Wassermeier, H.P. Schononherr and K.H. Ploog, Phys. Rev. Lett. 87, 016601 (2001).

70.       M. Kohda, Y. Ohno, K. Takamura, F. Matsukura and H. Ohno, Jap. J. Appl. Phys.40 L1274 (2001).

71.       J.A. Katine, F. J. Albert, R.A. Buhrman, E.D. Myers and D.C. Ralph, Phys. Rev. Lett. 84, 319 (2000).

72.       J. Nitta, T.  Ahazaki, H. Takayanngi and T. Enoki, Phys. Rev. Lett. 78, 1335 (1997).

73.       S. Cho., S. Choi, G.B. Cha, S.C. Hong, Y. Kim, Y.-J. Zhao ,A.J. Freeman, J.B. Ketterson, B.J. Kim,     Y.C. Kim and B.C.Choi,Phys.Rev.Lett.88 257203-1 (2002).

74.       G.A. Medvedkin, T. Ishibashi, T. Nishi and K. Hiyata, Jap. J. Appl.  Phys. 39, L949 (2000).

75.       G.A. Medvedkin, K. Hirose, T. Ishibashi, T. Nishi, V.G. Voevodin and K. Sato, K., J. Cryst. Growth 236, 609 (2002).

76.       S. Choi, G.B. Cha, S.C. Hong, S. Cho, Y. Kim, J.B. Ketterson, S.-Y. Jeong  and G.C. Yi,  Solid-State Commun. 122, 165 (2002).

77.       K. Ueda, H. Tahata and T. Kawai, Appl. Phys. Lett. 79, 988 (2001).

78.       Y. Matsumoto, M. Murakami, T. Shono, H. Hasegawa, T. Fukumura,  M. Kawasaki, P. Ahmet, T. Chikyow, S. Koshikara, and H. Koinuma, Science 291, 854 (2001).

79.       F. Holtzberg, S. von Molnar, and J.M.D. Coey, “Handbook on Semiconductors,” ed. T. Moss (North-Holland, Amsterdam, 1980).

80.       T. Dietl, H. Ohno, F. Matsukura, J. Cibert, and D. Ferrand, Science 287, 1019 (2000).

81.       C. Zener, Phys. Rev. B81, 440 (1951).

82.       M. Van Schilfgaarde and O.N. Myrasov, Phys. Rev. B63, 233205 (2001).

83.       T. Dietl, H. Ohno and F. Matsukura, Phys. Rev. B63, 195205 (2001).

84.       T. Dietl, J. Appl. Phys. 89, 7437 (2001).

85.       T. Jungwirth, W.A. Atkinson, B. Lee, and A.H. MacDonald, Phys. Rev. B59, 9818 (1999).

86.       M. Berciu, and R.N. Bhatt, Phys. Rev. Lett. 87, 108203 (2001).

87.       R.N. Bhatt, M. Berciu, M.D. Kennett, and X. Wan, J. Superconductivity: Incorporating Novel Magnetism 15, 71 (2002).

88.       V.I. Litvinov and V.A. Dugaev, Phys. Rev. Lett. 86, 5593 (2001).

89.       J. Konig, H.H. Lin, and A.H. MacDonald, Phys. Rev. Lett. 84, 5628 (2001).

90.       J. Schliemann, J. Konig, and A.H. MacDonald, Phys. Rev. B64, 165201 (2001).

91.       T. Jungwirth, J. Konig, J. Sinova, J. Kucera, and A.H. MacDonald, (in press).

92.       C.R. Abernathy, Mat. Sci. Rep. R16, 203 (1995).

93.       R.Y. Korortiev J.M. Gregie and B.W. Wessels, Appl. Phys. Lett. 80, 1731 (2002).

94.       H. Katayama-Yoshida, R. Kato and T. Yamamoto, J. Cryst. Growth 231, 438 (2001).

95.       I.D. Goepfert, E.F. Schubert, A. Osinsky, P.E. Norris, and N.N. Faleev, J. Appl. Phys. 88, 2030 (2000).

96.       K. Sato and H. Katayama-Yoshida, Jap. J. Appl. Phys. 40, L485 (2001).

97.       N. Theodoropoulou, A.F. Hebard, M.E. Overberg, C.R. Abernathy, S.J. Pearton, S.N.G. Chu, and R.G. Wilson, Phys. Rev. Lett. 89 107203(2002)

98.       M.E. Overberg, B.P. Gila, G.T. Thaler, C.R. Abernathy, S.J. Pearton, N. Theodoropoulou,  K.T. McCarthy, S.B. Arnason, A.F. Hebard, S.N.G. Chu, R.G. Wilson,  J.M. Zavada, and Y.D. Park, J. Vac. Sci. Technol. B 20 969 (2002).

99.       M. Zajac, J. Gosk, M. Kaminska, A. Twardowski, T. Szyszko, and S. Podliasko, Appl. Phys. Lett. 79, 2432 (2001).

100.   M.L. Reed, M.K. Ritums, H.H. Stadelmaier, M.J. Reed, C.A. Parker, S.M. Bedair and N.A. El-Masry, Mater. Lett. 51, 500 (2001).

101.   M.L. Reed, N.A. El-Masry, H. Stadelmaier, M.E. Ritums, N.J. Reed, C.A. Parker, J.C. Roberts, and S.M. Bedair, Appl. Phys. Lett. 79, 3473 (2001).

102.   N. Theodoropoulou, A.F. Hebard, M.E. Overberg, C.R. Abernathy, S.J. Pearton, S.N.G. Chu,  and R.G. Wilson,  Appl. Phys. Lett. 78, 3475 (2001).

103.   S. Sonoda, S. Shimizu, T. Sasaki, Y. Yamamoto and H. Hori, J. Cryst. Growth 237-239 1358 (2002).

104.   G.T. Thaler, M.E. Overberg, B. Gila, R. Frazier, C.R. Abernathy, S.J. Pearton, J.S. Lee,  S.Y. Lee,  Y.D. Park, Z.G. Khim,  J. Kim and F. Ren,  Appl. Phys. Lett. 80 3964 (2002).

105.   Y.L. Soo, G. Kioseouglou., S. Kim, S. Huang, Y.H. Kaa, S. Kubarawa, S. Owa, T. Kondo and H. Munekata, Appl. Phys. Lett. 79, 3926 (2001).

106.   N.A. Theodoropoulou, A.F. Hebard, S.N.G. Chu, M.E. Overberg, C.R. Abernathy, S.J. Pearton, R.G. Wilson and J.M. Zavada, Appl. Phys. Lett. 79, 3452 (2001).

107.   S.J. Pearton, M.E. Overberg, G. Thaler, C.R. Abernathy, N. Theodoropoulou, A.F. Hebard, S.N.G. Chu, R.G. Wilson, J.M. Zavada, A.Y. Polyakov, A. Osinsky and Y.D. Park, J. Vac. Sci. Technol. A20, 583 (2002).

108.   H. Akinaga, S. Nemeth, J. De Boeck, L. Nistor, H. Bender, G. Borghs, H. Ofuchi and M. Oshima, Appl. Phys. Lett. 77, 4377 (2000).

109.   M. Hashimoto, Y.Z. Zhou, M. Kanamura and H. Asahi,  Solid-State Commun.122 37 (2002)

110.   Pearton, S.J., Overberg, Abernathy, C.R. Theodoropoulou, N.A., Hebard, A.F., Chu, S.N.G., Osinsky, A., Zuflyigin, V., Zhu, L.D., Polyakov, A.Y. and Wilson, R.G., J. Appl. Phys. 92 2047(2002).

111.   R. Fiederling, M. Kein, G. Resescher, W. Ossau, G. Schmidt, W. Wang, and L.W. Molenkamp, Nature 402, 787 (1999).

112.   Y. Ohno, D.K. Young, B. Beschoten, F. Matsukura, H. Ohno, and D.D. Awschalom, Nature 402, 790 (1999).

113.   B.T. Jonker, Y.D. Park, B.R. Bennet, H.D. Cheong, G. Kioseoglou and A. Petrou, Phys. Rev. B62, 8180 (2000).

114.   Y.D. Park, B.T. Jonker, B.R. Bennet, G. Itzkos, M. Furis, G. Kioseoglou and A. Petrou, Appl. Phys. Lett. 77, 3989 (2000).

115.   B.T. Jonker, A.T. Hanbicki, Y.D. Park, G. Itskos, M. Furis G. Kioseoglou and A. Petrou, Appl. Phys. Lett. 79, 3098 (2001).

116.   P. LeClair, J.K. Ha H.J.M. Swagten, J.T. Kohlhepp, C.H. van de Vin and W.J.M. de Jonge,  Appl. Phys. Lett. 80, 625 (2002).

117.   A. Kaminski and S. Das Sarma, Phys. Rev. Lett. 88 247202-1 (2002).

Figure Captions

Figure 1.         Semiconductor matrix with high concentrations of magnetic impurities (i.e., Mn),

randomly distributed (defects), can be insulators (a) for group II-VI materials where divalent Mn ions occupy group II sites.  At high concentrations, Mn ions are antiferromagnetically coupled, but at dilute limits, atomic distances between magnetic ions are large, and antiferromagnetic coupling is weak.  For the cases where there is high concentrations of carriers (b) (i.e., (Ga,Mn)As where Mn ions behave as acceptors and provide magnetic moment as they occupy trivalent Ga sites), the carriers are thought to mediate ferromagnetic coupling between magnetic ions.  Between near insulating and metallic cases, at low carrier regimes, hole carrier concentrations are localized near the magnetic impurity.  Below certain temperatures, a percolation network (c) is formed in which clusters the holes are delocalized and hop from site to site, which energetically favors maintaining the carriers’ spin orientation during the process, an effective mechanism for aligning Mn moments within the cluster network.  Alternatively, at percolation limits, localized hole near the magnetic impurity is polarized, and the energy of the system is lowered when the polarization of the localized holes are parallel (d).

Figure 2.         Predicted Curie temperatures as a function of bandgap(after ref.80)

 

Figure 3.         Predicted ferromagnetic transition temperatures in (Ga,Mn)As (left) or (Ga,Mn)N containing 5 at.% Mn, as a function of hole density.  The four different curves in each graph represent results obtained from different variants of mean-field theory (after ref. 91).

 

Figure 4.         Predicted stability of the ferromagnetic states of different transition metal (TM) atoms in GaN as a function of transition metal concentration.  The vertical axis represents the energy difference between the ferromagnetic and spin glass states for each metal atom (after ref. 96).

 

Figure 5.      Field-cooled magnetization of (Ga,Mn)P as a function of temperature.  The solid line shows a Bloch law dependence, while the dashed lines are 95% confidence bands.  The vertical dashed line at TC = 270 K is the field-independent inflection point and the vertical arrows in the main panel and inset mask to ferromagnetic Curie temperature Θf.  The inset shows the temperature dependence of difference in magnetization between field-cooled and zero field-cooled conditions.

 

Figure 6.        Magnetization versus field for MBE grown (Ga,Mn)P with 9.4 at.% Mn.

 

Figure 7.      (a)B-H from MBE grown (Ga,Mn)N with 9.4at.%Mn(closed circles) and from sapphire substrate(upper circles),(b) M-T of (Ga,Mn)N,(c)B-H from (Ga,Mn)N as a function of Mn concentration.

Figure 8.         Temperature dependence of sheet resistance at zero magnetic field for Mn-diffused GaN and as-grown GaN (right) and room temperature anomalous Hall effect hysteresis curves for Mn-diffused GaN (left) (after ref.101).

 

Figure 9.        Magnetization versus temperature for (Ga,Mn)N sample grown by MBE with ~9 at.% Mn.  The extrapolation of the curve is based on a mean-field approximation (after ref. 103).