Mn alloys

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Physica E 16 (2003) 90 – 98 www.elsevier.com/locate/physe

Novel ferromagnetism in digital GaAs/Mn and GaSb/Mn alloys B.D. McCombea;∗ , M. Naa , X. Chena , M. Cheona , S. Wanga , H. Luoa , X. Liub , Y. Sasakib , T. Wojtowiczb , J.K. Furdynab , S.J. Potashnikc , P. Schi9erc a Center

for Advanced Photonic and Electronic Materials, Department of Physics, University at Bualo, The State University of New York, Bualo, NY 14260, USA b Department of Physics, University of Notre Dame, Notre Dame, IN 46556, USA c Department of Physics and Materials Research Institute, Pennsylvania State University, University Park, PA 16802, USA

Abstract Ferromagnetic III-Mn-V semiconductors are under intensive investigation for spintronic applications. In the present work digital alloys of GaAs/Mn and GaSb/Mn were fabricated by a combination of molecular beam epitaxy and atomic layer epitaxy. The Mn fraction in the layers was varied at constant III–V spacer thickness, and the III–V spacer-layer thickness was varied at constant Mn fraction (0.5 monolayer). Transmission electron micrographs showed good crystal quality with no evidence of three-dimensional (3D) Mn-V precipitates. The GaAs/Mn samples have Curie temperatures in the vicinity of 40 K and exhibit an anomalous Hall e9ect (AHE) similar to that seen in GaMnAs random alloys. These samples all show thermally activated resistance at zero Aeld, characteristic of hopping conduction with evidence of Coulomb-gap-like behavior. The GaSb/Mn samples exhibit ferromagnetism (with temperature dependent hysteresis loops) and a strong AHE up to 400 K. At low temperatures the remanent magnetization initially drops rapidly with increasing temperature, indicative of Curie temperatures between 10 and 50 K depending on the Mn concentration. However, a substantial remanent magnetization persists to high temperature, suggesting a second phase with a Curie temperature above 400 K. These samples all showed essentially metallic behavior and weak negative magnetoresistance at low temperatures. These results are discussed in the context of a model of quasi-2D MnSb islands (for which there is direct evidence) embedded in a matrix of GaMnSb in the Mn-containing layers. ? 2002 Elsevier Science B.V. All rights reserved. PACS: 75.50.Pp; 72:20: − I; 85:75: − d Keywords: Ferromagnetic semiconductors; Digital alloys; GaMnSb; GaMnAs; Magnetic and magnetotransport properties

1. Introduction The promise of spintronic applications of III-Mn-V semiconductors and the observation of Curie temperatures as high as 110 K in GaMnAs and above

∗

Corresponding author. Fax: +1-716-645-5994. E-mail address: [email protected] (B.D. McCombe).

room temperature in GaMnN random alloys has generated great interest in ferromagnetism in semiconductors [1–4]. The hole-mediated exchange interaction, which leads to ferromagnetism in random III-Mn-V alloys [2,4], also leads to some interesting and unique device possibilities. In addition, with these materials spin-related e9ects can, in principle, be integrated into the well-established III–V-based electronic/electro-optic technologies. However, the

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optical quality of the ferromagnetic materials is generally very poor due to the low-temperature growth, and the spin injection eJciency from GaMnAs is very low [5]. Recent experiments with II-Mn-VI spin polarizers/aligners on III–V non-magnetic structures (spin LEDs) have demonstrated eJcient spin injection [6,7]. However, the II-Mn-VI spin injectors are paramagnetic, necessitating low-temperature operation combined with modest-to-high magnetic Aelds to achieve large spin polarization. Spin injection from Fe into GaAs through tunneling barriers has shown promise at room temperature [8,9]. It is clear that approaches to improving the quality, enhancing the Curie temperature, TC , and increasing the spin polarization of ferromagnetic III-Mn-V materials are of great interest. It is now well accepted [2,4] that the mechanism for ferromagnetism in these materials is carrier(hole)-mediated exchange, which o9ers the opportunity to modify the magnetic properties either optically or electrically [10,11]. However, basic materials problems limit the Mn concentration that can be incorporated to produce usable ferromagnetism. Two important factors in carrier mediated exchange interactions are the distance between Mn ions and the itinerant carrier concentration. We have employed digital alloys as an approach to providing more control of Mn and hole concentrations, as well as a way of improving basic material quality. Our recent e9orts have focused on growth of GaAs/Mn and GaSb/Mn digital alloys by a combination of atomic layer epitaxy (ALE) and molecular beam epitaxy (MBE). In this approach quasi-two-dimensional (2D) layers of a dopant (in this case Mn) are periodically inserted in the GaAs or GaSb host lattices. We [12] and others [13] have demonstrated that GaAs/Mn digital alloys can be grown with high crystal quality, with an upper limit on Mn concentration of 50% in each Mn-containing monolayer (ML) [14]. In the following sections we describe the growth and structural characterization of digital alloys and present results on the dependence of magnetotransport and magnetic properties of GaAs/Mn digital alloys as a function of Mn fraction in the MLs and as a function of both Mn fraction and the spacing between the Mn-containing layers in GaSb/Mn digital alloys. We summarize the results and present a qualitative

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model based on these and other measurements and some recent unpublished theoretical calculations. 2. Growth and structural characterization All samples used in this study were grown on (0 0 1) GaAs substrates. The GaAs/Mn samples were grown under the same growth conditions, other than the Mn deposition time. The initial growth follows the typical procedure for GaMnAs alloys, starting N of GaAs, at a subwith an initial layer of 200 A strate temperature of 580◦ C. The substrate temperature is then brought down to 275◦ C and a layer of low-temperature GaAs is grown to a thickness of a few tens of nanometers. This is followed by growth of the GaAs/Mn digital alloy, the GaAs layers by MBE and the Mn layers by ALE. This combination of Mn and GaAs layers is repeated for a desired number of periods. For this study, the GaAs spacer-layer thickness in the digital alloys was kept at 9 ML (≈2:5 nm), large enough to minimize layer–layer magnetic interactions. The Mn fraction was varied from about 0.02 to 0:5 ML. Each of the resulting samples has 50 repetitions of the basic structure and the following fractional coverages in the Mn-containing layers: 0.14 (sample 1), 0.23 (sample 2), 0.4 (sample 3) and 0.5 (sample 4). The e9ective 3D alloy compositions are: 0.014 (sample 1), 0.023 (sample 2), 0.04 (sample 3), and 0.05 (sample 4). The large lattice mismatch between GaSb and GaAs (7.5%) necessitates growth of a thick GaSb bu9er layer (nominally 500 nm) upon which the GaSb/Mn digital alloys were grown. The growth was monitored by reOection high-energy electron di9raction (RHEED). The GaSb/Mn digital alloys consist of 50 periods of 0:5 ML Mn layers separated by GaSb layers of various thicknesses. None of the samples discussed showed any indication of 3D precipitate formation in the RHEED patterns. Both the GaAs/Mn and the GaSb/Mn samples were characterized by cross-sectional transmission electron microscopy (TEM) and by magnetic force microscopy (MFM). Examples of the TEM results are shown in Fig. 1. The samples showed good structural quality and clear evidence of well-deAned superlattice formation with superlattice periods in

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Fig. 1. Cross-sectional Transmission Electron Micrographs of (a) a GaAs/Mn digital alloy with an 8 GaAs ML spacer-layer thickness and 0:2 ML Mn nominal coverage in each Mn-containing layer, and (b) a GaSb/Mn sample with 12 GaSb ML spacer and 0:4 ML Mn layer coverage.

good agreement with nominal values and the RHEED measurements. No evidence for 3D MnAs or MnSb precipitates was found in the as-grown samples. High temperature (¿ 500◦ C) annealing leads to the formation of 3D precipitates that are clearly observed in cross-sectional TEM and, in the case of the GaSb/Mn samples, MFM measurements. Because a maximum of 12 ML of Mn can be deposited with this technique, [14] lateral, quasi-2D islands of MnAs or MnSb (depending on the host material) can be expected in addition to randomly distributed (laterally isolated) Mn ions within the Mn-containing layers. Some redistribution of Mn into adjacent layers of the host lattice is also expected. Magnetization measurements with a superconducting quantum interference device (SQUID) magnetometer and electrical magnetotransport measurements in the van der Pauw and Hall conAgurations were used to study the magnetic properties and the interaction between the Mn spins and the itinerant carriers, respectively. For the electrical measurements contacts were made by Au/Zn/Au or Au metallization and subsequent di9usion at 250◦ C.

3. Results: GaAs/Mn Magnetization measurements were performed on all samples at temperatures between 5 K and room temperature. As the fraction of Mn is increased to x = 0:1 ML (nominal), ferromagnetism becomes observable above 5 K as evidenced by (typically square) magnetization hysteresis loops (see Fig. 2). However, the saturation magnetization does not increase monotonically with Mn coverage up to its maximum of x = 0:5 M; it peaks at about x = 0:23 ML. This behavior is indicative of either strong antiferromagnetic coupling among an increasing fraction of the Mn ions, or a larger fraction of compensating donor-like defects (As antisites) and concomitantly reduced carrier (hole) density at higher Mn coverage. High Aeld magnetotransport measurements (up to 33 T) indicate that the latter is unlikely—the hole density increases with increasing Mn coverage for the two samples that have been studied to date. With the assumption that substitutional Mn has spin 5=2 and with g = 2 one can deAne an eective spin density (ESD) from the saturation magnetization at a high Aeld (¿ 2 T). This provides a measure of the total number of spins that can

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Fig. 2. Magnetization hysteresis loop at 5 K for a GaAs/Mn sample with a 9 ML spacer and Mn layer coverage of 0:4 ML (sample 3). The applied Aeld is in the sample plane (the easy axis lies in the plane).

be aligned within this Aeld range. The saturation magnetization and the Curie temperature are highly correlated, both reaching a maximum at x = 0:23 for this particular group of samples. Our studies of samples grown under di9erent conditions showed that the ESD also depends on growth conditions; two samples with nominally identical Mn fraction (and the same GaAs spacer-layer thickness) can have signiAcantly di9erent ESD and TC [12]. Other studies [14] have shown that the dependence of TC on GaAs spacer-layer thickness is relatively weak. All GaAs/Mn digital alloys studied so far are non-metallic in the sense that their resistance is thermally activated at low temperatures. These results show that in the digital alloys the ferromagnetic interaction is mediated by holes that are localized to some extent, rather than fully itinerant (metallic). Unannealed GaMnAs random alloys exhibit metallic conductivity (R increases with T ) for 0:035 ¡ x ¡ 0:07 and ferromagnetism at temperatures up to 110 K (x = 0:053). For x = 0:01– 0.025 they are non-metallic and ferromagnetic with lower Curie temperatures; thus, localized holes mediate the ferromagnetic interaction in this case also. Recent studies [15] have

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shown that over the range x = 0:01– 0.85 properly annealed Ga1−x Mnx As samples are both metallic and ferromagnetic, with Curie temperatures that increase monotonically up to x = 0:05 and saturate at 110 K for samples in the range 0:05 ¡ x ¡ 0:085. The Hall resistance RH , shown in Fig. 3(a) for sample 3, exhibits a strong anomalous Hall e9ect (AHE) that is correlated with the observed ferromagnetism. The decrease of the magnitude of the Hall resistance with B for B ¿ 1 T is due to the fact that RH ˙(Rsheet )n , where n ¡ 1 in the present case, and Rsheet exhibits a strong negative magnetoresistance in this region as shown in Fig. 3(b). The general behavior of the magnetoresistance is qualitatively similar to that observed in ferromagnetic GaMnAs random alloys; there is a small initial positive magnetoresistance region followed by a large negative magnetoresistance at high Aelds. The positive magnetoresistance is attributed to rotation of the magnetization from easy axis (in-plane) to hard axis (perpendicular to the place). The resistance changes by a factor of 3– 4 between B = 0 and 9 T. For T ¿ TC , the low-Aeld positive magnetoresistance disappears, leaving negative magnetoresistance at all Aelds over a range of temperatures (up to 200 K). All samples studied show this qualitative behavior. The magnitude of the negative magnetoresistance, [Rsheet (0) − Rsheet (B = 7 T)]=Rsheet (0)] ≈ 2– 4, di9ers from that seen in ferromagnetic GaMnAs random alloys [16]. For metallic random alloy samples the magnitude of the negative magnetoresistance is much smaller, increases with temperature up to TC , and then decreases with further increase of T above TC . The negative magnetoresistance in the non-metallic, random alloy samples at high Mn composition (x ¿ 0:07) is much larger than that seen in the digital alloys. At low concentrations (x ¡ 0:03) the negative magnetoresistance in the random alloys is closer to that seen in the digital alloys. The sheet resistance at B = 0 also shows interesting behavior. The temperature dependence of Rsheet (0) is shown in Figs. 4(a) and (b), for two samples (samples 1 and 3). In these Agures ln R is plotted vs. T −n , where n is 1 or a fraction, to show more clearly the form of the activated behavior at low temperatures. The straight lines are Ats of ln R = ln R0 − [T0 =T ]n for comparison. It is apparent that the behavior of sample 1 (x = 0:14) is best At with n = 14 , while that

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Fig. 3. Magnetotransport measurements for the sample of Fig. 2 at various temperatures: (a) Hall resistance; (b) magnetoresistance. The carrier density from the slope at 300 K is 1:07×1013 cm−2 per layer. The behavior is qualitatively characteristic of all samples in this series.

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of Sample 3 (x = 0:4) is best At with n = 12 . The n = 12 dependence is seen in all the samples of this sequence having x ¿ 0:2. This behavior is consistent with conduction via variable range hopping in the presence of a coulomb gap [17]. Thus, the samples with Mn coverage ¿ 0:2 ML/layer appear to behave as strongly correlated and disordered metals. The behavior of sample 1 is consistent with Mott variable range hopping in 3D [17]. Other samples with high Mn content (¿ 0:2) grown under di9erent conditions that lead to a smaller ESD (more defects) have n = 14 . These results thus appear to be related to proximity to the metallic state.

1 4

for two GaAs/Mn samples: (a) sample 1; (b) sample 3.

4. Results: GaSb/Mn All samples studied showed ferromagnetism at temperatures up to 400 K, the highest temperature achievable with our magnetometer, as indicated by hysteresis loops in the magnetization. Thus there is a maximum Curie temperature, which we denote as TC2 ¿ 400 K. Data for one sample is shown in Fig. 5(a). Fig. 5(b) shows the remanent magnetization vs. temperature. There is an initial rapid decrease of remanent magnetization with temperature and then a leveling o9 to a nearly constant value above about 50 K, which is suggestive of another phase with a much lower Curie

B.D. McCombe et al. / Physica E 16 (2003) 90 – 98 8

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Fig. 5. Magnetization results for one of the GaSb/Mn digital alloy samples: (a) hysteresis loops at three temperatures; (b) remanent magnetization at B = 0 vs. T after the samples was magnetized at low temperature.

temperature, TC1 . The hysteresis loops of Fig. 5(a) show strong temperature dependence over the range of temperatures studied. For example, the coercive Aelds at 5 and 285 K di9er by a factor of 2, 0.01 and 0:005 T, respectively, while the magnetization at high Aelds decreases signiAcantly over this temperature range. This behavior is in sharp contrast to the ferromagnetism that has been seen previously at room temperature in GaMnSb samples containing 3D precipitates (in the NiAs structure), which is characterized by hysteresis loops with almost no variation in coercive Aeld between 5 and 300 K. Furthermore, the shapes of the hysteresis loops for the present digital alloys di9er signiAcantly (at all temperatures) from those of 3D MnSb precipitates [18]. Such precipitates can be produced in our samples by annealing at high (500◦ C) temperatures. After annealing, the greatly changed hysteresis loops seen in our samples are very similar those of MnSb precipitates. We conclude from our SQUID measurements, cross-sectional TEM measurements and MFM measurements (discussed below) that the observed ferromagnetism in the digital alloys is not due 3D MnSb precipitates. The properties of the charge carriers in III–V dilute magnetic semiconductors are particularly important, since the interaction between itinerant carriers and localized Mn moments is the determining factor in the ferromagnetism of these materials. Magnetotransport measurements can be used to estimate the coupling between charge carriers and magnetic ions, to deter-

mine scattering mechanisms, and under certain conditions to measure carrier densities in ferromagnetic magnetic semiconductors [1,16]. All GaSb/Mn samples in the present study exhibit metallic behavior, in contrast to the thermally activated resistance for the GaAs/Mn digital alloys discussed above [14]. The zero-Aeld sheet resistance depends weakly on temperature, but shows interesting non-monotonic behavior (not discussed here). The carrier density estimated from the low-temperature, high-Aeld (linear) region of the AHE (data up to 33 T) is 3–6 × 1013 cm−2 per Mn layer at 4:2 K, corresponding to roughly 20% of the nominal 0:5 ML Mn per Mn-containing layer. A large hole density is important for the hole-mediated exchange interaction between Mn ions (and possibly between lateral quasi-2D MnSb islands). The AHE provides information about the magnetization and spin–orbit interaction, as well as scattering mechanisms for the itinerant carriers. As shown in Fig. 6(b) the magnitude of the anomalous contribution to the Hall e9ect initially decreases with increasing temperature but persists and regains strength at temperatures above 200 K—the slope of the Hall resistance near zero-Aeld is a measure of the magnetization, as well as the strength of its coupling with the carriers. The sign of the AHE and the slope at low Aelds is related to the band structure of the itinerant carriers and the spin–orbit interaction [20]. We note that two-carrier (electrons and holes) conduction, can in principle lead to similar low-Aeld

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Fig. 6. Magnetotransport measurements from 5 to 320 K for a GaSb/Mn digital alloy sample having 0:5 ML Mn coverage per layer and a spacer thickness of 12 GaSb MLs: (a) magnetoresistance; (b) Hall resistance.

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behavior. However, this mechanism requires that electrons dominate the conductivity at low Aelds at both low and high temperatures (negative slope of RH ), and have a negligible e9ect at high Aelds (where the slope is positive). We have explored this possibility and indeed And electrons in some samples (apparently in the GaSb near the AlSb bu9er layer), but the density is about 0.4% of the total hole sheet density, while the mobility is about 10 times that of the holes. In addition, all samples show a very similar low-temperature AHE independent of whether there are electrons present. Most importantly, the hysteretic behavior (demonstrated by the open loops in the inset to Fig. 7 at 4 and 400 K) is a signature of both ferromagnetism and its interaction with the holes. The fact that the hysteresis loop in the AHE at 400 K is stretched out further than the corresponding hysteresis loop in the magnetization (Fig. 5(a)) is not understood. Previous work on GaMnSb samples grown at high temperatures and containing 3D MnSb precipitates showed no clear AHE even at 4 K [18]. The AHE observed in low-temperature-grown GaMnSb random alloys with no MnSb precipitates decreased rapidly above TC (25 K), vanishing completely around 50 K [19]. The AHE coeJcient for the GaSb/Mn digital alloys is negative at all temperatures, as found also for GaMnSb random alloys at low temperature [19]. The sign is presently not understood. The AHE, even

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in “conventional” ferromagnetic metals, is not fully understood (see, e.g., Ref. [20]). To further support our assertion that 3D MnSb precipitates are not the source of the observed phenomena, we have carried out high temperature annealing studies. Samples were annealed at 400◦ C and 500◦ C for 5 min and examined by SQUID magnetometry and atomic force microscopy (AFM)/MFM. Both the magnetization and MFM measurements

B.D. McCombe et al. / Physica E 16 (2003) 90 – 98

showed the appearance of 3D precipitates in samples annealed at 500◦ C. The coercive Aeld at 5 K increased from 0:01 T for the as-grown sample to 0:04 T after annealing. The shape of the hysteresis loop also became identical to that reported for MnSb precipitates [18]. Large numbers of precipitates were observed for the sample annealed at 500◦ C by MFM, while there was no visible change of surface structure from AFM images. Temperature dependent magnetization and magnetotransport data reveal additional complexity associated with ferromagnetism in the as-grown samples. The remanent magnetization shows an initial rapid drop with temperature at low temperatures, which changes to a much slower decrease at higher temperatures. This change occurs between 30 and 50 K, as shown in Fig. 5(b) for one sample. This behavior suggests the coexistence of more than one phase. The magnetoresistance, shown in Fig. 6(a) changes from negative to positive with increasing temperature, with a crossover between 40 and 50 K for this sample. Below 40 K the behavior is very similar to that seen previously in GaMnSb random alloys [19]. In the case of LT-grown GaMnSb random alloys, the negative magnetoresistance decreases rapidly above TC (25 K), and the sheet resistance becomes independent of magnetic Aeld above 50 K. The present results and unpublished cross-sectional scanning tunneling microscopy (STM) measurements at the Naval Research Laboratory [21] that show MnSb islands, suggest the following qualitative picture. The Mn-containing layers consist of quasi-2D GaMnSb random alloys and small quasi-2D islands (lateral dimensions about 4 –5 nm) of zinc-blende MnSb. The matrix (quasi-2D GaMnSb random alloy) is ferromagnetic below 30 –50 K, depending on Mn coverage, and is responsible for similarities between the GaSb/Mn digital alloys and GaMnSb random alloys at low temperatures. The observed ferromagnetism at higher temperatures is associated with the 2D MnSb islands, and the two contributing regions interact via the holes [22]. MFM studies of these samples at room temperature show a lateral stripe-domain pattern, indicating a collective state at room temperature. A detailed theoretical

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study is needed to fully understand the mechanisms involved. 5. Conclusions Digital III–V/Mn alloys appear to be promising for producing ferromagnetic III–V semiconducting materials and structures. The structural quality is very good, and appropriately processed samples show cooperative ferromagnetic behavior at room temperature and above, with a strong anomalous Hall e9ect. Irrespective of the details of the mechanism, the facts that carrier transport is strongly modiAed by the ferromagnetic properties and hysteresis is observed in the AHE at 400 K, suggest that such structures may be useful as spin polarizers and injectors at room temperature. Acknowledgements This work was supported in part by the DARPA/ ONR SpinS program under N000140010951 and ONR under N000140010819.

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