Ferromagnetism and possible application in spintronics ...

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Materials Science and Engineering R 62 (2008) 1–35

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Materials Science and Engineering R journal homepage: www.elsevier.com/locate/mser

Ferromagnetism and possible application in spintronics of transition-metal-doped ZnO ﬁlms F. Pan *, C. Song, X.J. Liu, Y.C. Yang, F. Zeng Laboratory of Advanced Materials, Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China

A R T I C L E I N F O

A B S T R A C T

Keywords: Diluted magnetic oxide ZnO Local structure Defects Spintronics

This review article ﬁrst presents a summary of current understanding of the magnetic properties and intrinsic ferromagnetism of transition-metal (TM)-doped ZnO ﬁlms, which are typical diluted magnetic oxides used in spintronics. The local structure and magnetic behavior of TM-doped ZnO are strongly sensitive to the preparation parameters. In the second part, we discuss in detail the effects of doping elements and concentrations, oxygen partial pressure, substrate and its orientation and temperature, deposition rate, post-annealing temperature and co-doping on the local structure and subsequent ferromagnetic ordering of TM-doped ZnO. It is unambiguously demonstrated that room-temperature ferromagnetism is strongly correlated with structural defects, and the carriers involved in carriermediated exchange are by-products of defects created in ZnO. The third part focuses on recent progress in TM-doped ZnO-based spintronics, such as magnetic tunnel junctions and spin ﬁeld-effect transistors, which provide a route for spin injection from TM-doped ZnO to ZnO. Thus, TM-doped ZnO materials are useful spin sources for spintronics. ß 2008 Elsevier B.V. All rights reserved.

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Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Experimental procedure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Film growth of TM-doped ZnO . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Characterization techniques. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Local structure and ferromagnetism of TM-doped ZnO . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Effect of doping concentration on structure and magnetization . . . . . . . . . . . 3.1.1. Effect of doping concentration on magnetic ordering . . . . . . . . . . . . 3.1.2. Formation of a secondary phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Substrate effects on magnetization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1. Substrate-dependent local structure and magnetic behavior . . . . . . 3.2.2. Substrate orientation-induced distinct magnetization . . . . . . . . . . . . 3.3. Effects of substrate temperature and ﬁlm deposition rate on magnetization . 3.4. Inﬂuence of oxygen partial pressure on magnetic ordering . . . . . . . . . . . . . . . 3.5. Strain-induced ferromagnetism enhancement . . . . . . . . . . . . . . . . . . . . . . . . . 3.6. Inﬂuence of post-annealing on magnetization . . . . . . . . . . . . . . . . . . . . . . . . . 3.7. Effect of co-doping on magnetic properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8. Comparison of ZnO doped with Sc, Ti, V, Cr, Mn, Fe, Co, Ni and Cu . . . . . . . . Origin of ferromagnetism in TM:ZnO . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. Computational work. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. Experimental work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.1. Secondary phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.2. Carrier exchange interactions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

* Corresponding author. Tel.: +86 10 62772907; fax: +86 10 62771160. E-mail address: [email protected] (F. Pan). 0927-796X/$ – see front matter ß 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.mser.2008.04.002

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4.2.3. Defect-based BMP . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3. Effect of bandgap on the Curie temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . TM-doped ZnO-based prototype spintronics and multi-functional materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. Co-doped ZnO-based magnetic tunnel junctions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1.1. Spin-polarized transport in (Zn,Co)O/ZnO/(Zn,Co)O junctions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1.2. Anomalous TMR in (Zn,Co)O-based junctions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. TM-doped ZnO-based spin ﬁeld-effect transistors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3. Ferroelectricity and giant piezoelectric d33 in V and Cr-doped ZnO . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions and outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction A diluted magnetic semiconductor (DMS) is a compound with properties intermediate between a non-magnetic semiconductor and a magnetic element, and is obtained by doping a non-magnetic semiconductor with transition-metal (TM) elements (e.g., Sc, Ti, V, Cr, Mn, Fe, Co, Ni and Cu), including TM-doped III–V (GaAs and InAs) [1–4], II–VI (CdTe) [5] or group IV (Ge and Si) types [6]. However, most DMSs have a low Curie temperature (TC), which limits their use in practical applications [7,8]. In 2000, Dietl et al. [9] theoretically predicted that the TC of DMSs could be increased above room temperature (RT) for p-type DMSs, and ferromagnetism (FM) was stable in DMSs based on wide-bandgap semiconductors, i.e., ZnO and GaN. Using ﬁrst-principle calculations, Sato and Katayama-Yoshida [10] theoretically demonstrated that a ZnO matrix doped with TM atoms such as V, Cr, Fe, Co and Ni exhibited FM ordering, whereas doping with Ti and Cu resulted in a paramagnetic state, which opened a window for experimental attempts to prepare DMSs with room temperature ferromagnetism (RTFM). FM semiconductors with O2 anions, such as TM-doped ZnO, are also termed diluted magnetic oxides (DMOs). Soon after that, consecutive experimental reports revealed that TM-doped ZnO and TiO2 exhibited intrinsic RTFM [11,12]. These observations prompted extensive experimental work and theoretical studies on RTFM DMOs, driven both by an urge to understand the mechanisms involved and by the demand for better materials. In fact, interest in TM-doped ZnO has increased because of promising applications in the ﬁeld of semiconductor spintronics, which seeks to extend the properties and applications of established electronic devices by using the spin of electrons in addition to their charge [7,13–15]. Electron spin can potentially be used to provide efﬁcient injection of spin-polarized carriers for spintronics and to fabricate transparent magnetic materials for spintronics [16–22]. There are two major criteria for selecting the most promising materials for semiconductor spintronics. First, the intrinsic FM should be retained at practical temperatures (i.e. >300 K). Second, it is a major advantage if there is existing technology for the material in other applications [14]. Fortunately, TM-doped ZnO satisﬁes these two criteria. Following the theoretical prediction of RTFM ascribed to a carrier-mediated mechanism in ZnO-based DMOs [9], TM-doped ZnO has been intensively investigated to achieve magnetic ordering above RT and promising magneto-transport properties [11,14,16–20]. As a low-cost, wide-band gap (Eg = 3.37 eV) semiconductor, ZnO itself has been the focus of renewed research for applications such as UV light-emitters, transparent high-power electronics, surface acoustic wave devices, piezoelectric transducers and window materials for display and solar cells [21,22]. Recently, a surge in research into TM:ZnO DMOs has been observed, as shown in Fig. 1, indicating that this area is attracting increasing interest not only because of promising applications in

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spintronics, but also owing to the excellent physical properties. So far, RTFM has commonly been obtained in ZnO systems doped with TM elements such as Sc [23,24], Ti [23,25], V [26–28], Cr [11,23,29,30], Mn [31–40], Fe [41–43], Co [8,23,44–61], Ni [62,63], and Cu [64–68], and co-doping such as CoFe [69,70] and MnCo [71,72], for which Co- and Mn-doped ZnO are the most popular systems. Many groups have reported that TM-doped ZnO ﬁlms exhibit FM ordering at a TC much lower than RT, such as 83 K [73] and 110 K [74] for Mn-doped ZnO. Many previous publications indicate that TM ions at Zn sites is a necessary but not sufﬁcient condition for FM in true TM:ZnO DMOs. Some ZnO ﬁlms doped with Ti [75], V [27], Mn [42], Co [76,77] and Cu [78] only exhibited paramagnetism or superparamagnetism as well as spin glass behavior [79]. Although a considerable amount of experimental data and corresponding mechanisms has been accumulated since the ﬁrst report of RTFM [11,16–20], the origin of FM ordering in DMOs, including TM-doped ZnO, is still a matter of debate, i.e., whether carrier-mediated exchange based on TM2+ replacement of Zn2+ [16,17] or the formation of secondary phases [18,69,80] is involved. Thus, it is critical to distinguish true FM semiconductors from those that merely show magnetic hysteresis. Even more ambiguously, distinct results are frequently obtained in rather similar TM-doped ZnO ﬁlms in the case of Co2+ substitution for Zn2+ [11,81,82], indicating that the magnetization of TM-doped ZnO is strongly dependent on the preparation parameters, and that the preparation process is not easily reproduced from one growth run to another. In this rather contradictory context, several experiments and calculations have been performed to elucidate the nature of RTFM

Fig. 1. Publications per year on TM:ZnO DMOs according to the Web of Science: http://apps.isiknowledge.com/.

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[8,31,44,83]. Mechanism proposed include a bound magnetic polaron (BMP) mechanism involving defects, including oxygen vacancies (VO) and Zn interstitials (Zni) [8,44,57], hole doping involving Zn vacancies [34,49,84], and the formation of FM secondary phases (e.g., Mn2xZnxO3d via Zn diffusion into the Mn oxide [31]) and TM-rich nanocrystals [83]. Moreover, ab initio calculations predicted that competition between FM and anti-FM coupling existed in DMOs [84]. However, there is ongoing debate about these systems. In addition, a trace amount of an external pollutant would also lead to a hysteresis loop during magnetization measurements, further adding to the difﬁculties in determining true DMOs [20,85,86]. So far, publications, including reviews [16–20], have commonly focused on a summary and comparison of DMOs obtained by different groups and the different experimental conditions used (e.g., preparation methods and parameters). The resulting macroscopic magnetic properties, which are critically dependent on growth and processing conditions, are diverse without any common characterizations. It is therefore difﬁcult to identify a consistent mechanism to effectively understand the origin of FM ordering merely by comparison of experimental results achieved by different groups. Half a decade has already passed since the ﬁrst report of RFTM for TM-doped ZnO. Efforts are increasingly shifting not only to the rational design of methods to evaluate this system, but also to the development of TM-doped ZnO-based prototype spintronics [14,85]. When successfully combined with semiconductor functionalities, spintronics will have a considerable impact on future electronic device applications [15]. Spin injection from a TM-doped ZnO layer to pure ZnO and other conventional semiconductors is an important long-term goal of the continuing research into ZnO DMOs [87,88]. This concept has been realized in a representative spintronics application, TM:ZnO-based magnetic tunnel junction (MTJs), which exhibit superior spin-polarized transport, such as a large half voltage (V1/2, at which the tunnel magnetoresistance ratio becomes half of the maximum) [88,89], compared to conventional magnetic metals and magnetic TM-oxide-based MTJs with AlOx and MgO barriers [90]. Another typical application is the spin ﬁeld effect transistor (spin-FET) proposed by Datta and Das [91]. The core idea of this device is to induce spin precession by the Rashba spin–orbit interaction in a two-dimensional electron gas and to use spin-dependent materials. Although a working prototype of the Datta–Das spin-FET has not yet been fabricated, the relevant conditions necessary for the desired spin-FET operation and simpliﬁed devices have been obtained [92,93]. In this review, to study the nature of FM ordering at high TC, our main aim was to extract a single variable from a series of experiments. An example is substrate variation [94], which is of great interest owing to its potential to provide direct information regarding the factors that affect FM ordering. Moreover, since ZnO is a multi-functional material, recent signiﬁcant advances for this system are described, including ferroelectric, piezoelectric, optical, and electric properties, which play important roles in spintronics applications of TM-doped ZnO [21]. We summarize recent progress in TM-doped ZnO-based prototype spintronics, such as MTJs and spin-FETs. It should be pointed out that it is almost impossible to describe all the work performed around the world owing to the rapid pace of studies in this ﬁeld, and therefore some references may have been overlooked. 2. Experimental procedure 2.1. Film growth of TM-doped ZnO TM-doped ZnO ﬁlms are commonly deposited by pulsed laser deposition (PLD) [11,33,41,46–49,95], magnetron co-sputtering,

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including direct current (dc) reactive [44,45,94] and radiofrequency (rf) [32,96–98], combinatorial laser molecular-beam epitaxy [50,51], chemical vapor deposition (metal–organic [34,52,53,99,100], ultrasonic-assisted solution [57] and plasmaenhanced [101]), ion-beam implantation [35,42], ion-beam sputtering [54,55], and sol–gel methods [39,56]. Although the magnetic behavior of TM-doped ZnO ﬁlms is sensitive to the deposition conditions [16–18], there is no conclusion on which deposition method is best for FM ordering of the ﬁlms. For example, Co-doped ZnO ﬁlms deposited by PLD [11,77] or magnetron sputtering [74,94], two widely used deposition methods, could alternatively exhibit RTFM and RT paramagnetism. Therefore, magnetron sputtering is increasingly popular for growing TM-doped ZnO ﬁlms owing to its low cost, high efﬁciency, and easy control, and its production of uniform ﬁlms of large size. The substrate generally used is Al2O3(0 0 1) for low mismatch between the ﬁlm and the substrate (2%) after a 308 in-plane rotation [21], which is suitable for ZnO ﬁlms exhibiting high crystallinity. Some other substrates suitable for the deposition of TM-doped ZnO ﬁlms are Si [37,58,62,67,102], fused quartz [33], glass [54,68,103,104], ZnO [51,105], ScAlMgO4 [106], LiNbO3 [44], LiTaO3 [107], MgO [40], SiO2 [108] and NaCl [94]. Fig. 2 presents typical cross-sectional high-resolution transmission electron microscopy (HRTEM) images of Co-doped ZnO ﬁlms grown by dc reactive magnetron co-sputtering on Al2O3(0 0 1), Si(1 1 1), SiO2(1 0 1) and LiNbO3(1 0 4) substrates. In contrast to the epitaxial growth of Co-doped ZnO ﬁlm on Al2O3 substrate (Fig. 2a), the ﬁlms deposited on Si, SiO2 and LiNbO3 exhibit an (0 0 2) preferred orientation with structural defects (e.g., edge dislocations, marked by ?) in Fig. 2b–d. Large- and small-angle grain boundaries are observed in the ﬁlms on Si (Fig. 2b)/LiNbO3 (Fig. 2d) and SiO2 (Fig. 2c), respectively. Moreover, the (0 0 2) planes are not parallel to the surface of the LiNbO3 substrate; instead, the pillar-like grains tend to grow in two directions, as highlighted by the arrows in Fig. 2d. These observations reveal that the substrate affects the microstructure and the formation of defects in TM-doped ZnO ﬁlm, which are expected to play an important role in FM ordering in this system [94]. For PLD and rf magnetron sputtering techniques, ﬁlms were deposited using a ceramic target prepared using standard ceramic techniques from ZnO and TMO or TM2O3 powders of 4N purity [23]. For example, ZnO and MnO2 were mixed together, ground and calcined for 8 h at 400 8C and sintered at 600–900 8C for 12 h in air to obtain Zn1xMnxO(0 x 0.1) ceramic pellets [33], whereas ZnO and CoO/Co3O4 powders were used to prepare a Zn1xCoxO(0 x 0.1) target. An alternative way to synthesize TM-doped ZnO ﬁlms starts from pure metal targets of Zn and TM [44,59,94] or alternate deposition of TM/ZnO [54,55]. The relative sputtering area of TM chips attached to a Zn target was adjusted to control the TM composition in TM-doped ZnO ﬁlms grown by dc reactive magnetron co-sputtering [44]. Although it is hard to precisely control the TM content in the ﬁlm to the same as that in the target [59], the doping concentration in the ﬁlm can be estimated using extrapolated data, i.e., 5 at.% Co in a ZnO target corresponded to 6 at.% Co-doped ZnO in ﬁlms prepared by PLD [25]. In general, the doping concentration ranges from 1 to 10 at.%, producing typical DMOs exhibiting good crystal quality and strong magnetization; lower doping amounts lead to weak magnetization, whereas higher amounts tend to result in the appearance of secondary phases, depending on the TM elements [50,56,69,80]. The typical deposition conditions for laser ablated ﬁlms are a temperature of 400–600 8C, oxygen partial pressure close to 1 105 to 5 105 torr and laser energy of 1–3 J cm2 [11,25,33,41,46–51,95]. A relatively low (<600 8C) growth temperature leads to homogeneous wurtzite TM-doped ZnO ﬁlms,

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Fig. 2. (a) HRTEM image of Zn0.96Co0.04O/Al2O3 in cross-section. The inset shows the corresponding SAED pattern. (b) Representative columnar crystal grain of 15 nm in ¯ zone for Zn0.96Co0.04O grown on SiO2. (d) Cross-sectional diameter along the ZnO(0 0 2) growth direction in Co:ZnO ﬁlm grown on Si substrate. (c) HRTEM image of the ½1010 HRTEM image of Co:ZnO/128LN. The arrows show the growth direction along ZnO(0 0 2) planes. Typical edge dislocations are marked by white symbols (?).

whereas higher temperatures increase the possibility of an inhomogeneous ﬁlm containing secondary phases, such as hexagonal Co, CoO and Co3O4 phases [17,33,109,110]. For ﬁlms grown by magnetron sputtering, the substrate temperature (200– 400 8C) during deposition is typically lower than that for PLD [44,68,111], and a high substrate temperature (650 8C) results in the appearance of secondary phases [111]. The working gas is a mixture of argon and oxygen at various pressures, which greatly affects the local structure and the electric and magnetic properties of ZnO ﬁlms [112–114]. It was found that high oxygen partial pressure (>3 103 torr) resulted in insulating Co-doped ZnO ﬁlms (>105 V cm), which could provide a reference for elucidation of the nature of RTFM [112,113]. 2.2. Characterization techniques Highly sensitive characterization techniques are critical for understanding the local structure and the corresponding magnetic behavior of TM-doped ZnO and to accurately detect chemical information for doping elements at very low concentrations. X-ray diffraction (XRD) in the Bragg–Brentano, rocking curve, and pole ﬁgure geometries is frequently applied to characterize the ﬁlm structure and crystalline quality of TMdoped ZnO, and to compare the position of diffraction peaks between doped and undoped ZnO ﬁlms to predict the state and site of doping elements [11,41,44,48,59,94]. Moreover, HRTEM, ﬁeld-emission scanning electron microscopy (FE-SEM) and selected-area electron diffraction (SAED) were used in most DMO studies to identify interface bonding and structural characteristics and to observe whether small nanoclusters were present in the ﬁlms [36,44,48,94,110]. During the past 5 years of DMO research, many methods have been introduced to identify the local structure of doping elements, which is considered essential information for elucidating the intrinsic FM and judging whether a system is a genuine DMO. (1) Diffraction peaks due to secondary phases in a TM-doped ZnO

matrix can be observed from the XRD pattern. (2) Small nanoclusters present in the ﬁlm can be observed in HRTEM images. (3) Electron energy loss spectroscopy (EELS) combined with HRTEM [33,108] and X-ray photoelectron spectroscopy (XPS) [44,52] can be used to determine the valence state. (4) X-ray absorption spectroscopy (XAS), including TM K-edge [34,44,55,115], TM L-edge [97,116–118] and O K-edge [37,39,97,109], and its interpretation are used to elucidate the local structure. This technique is currently attracting great interest owing to its potential to provide local chemical information for complex materials and its sensitivity to subtle differences in local arrangement [119,120]. (5) Most importantly, DMOs should not only have a strong magnetic circular dichroism (MCD) signal, but should also exhibit an MCD spectral shape that reﬂects the band structure of the TM-doped ZnO, which is considered the most useful tool to determine whether TM-doped ZnO is a real DMO or not [37,47,52,60,85,121]. (6) Electron spin resonance (ESR) [122,123] and Mo¨ssbauer spectroscopy [42,124] are carried out to determine the local structure. (7) Measurement of optical properties using Raman spectroscopy and optical transmittance (absorption) spectrophotometry are effective in identifying the local structure of the TM dopant [48,108,109,111]. If the state of a trace amount of TM dopant cannot be determined using a single characterization technique, the best way to investigate the local TM structure may be to use several different methods. Magnetization studies are carried out using a superconducting quantum interference device (SQUID) magnetometer in the temperature range 5–350 K. High-temperature magnetization (300 < T < 800 K) was recorded using a vibrating sample magnetometer (VSM), which has also been used to measure TC [11,44,59,125,126]. RTFM can also be examined using other methods, such as alternating gradient magnetometry (AGM). Furthermore, the average atomic magnetic moment is an interesting parameter in evaluating and understanding DMOs. It is therefore important to determine the accurate TM content in samples. X-ray ﬂuorescence (XRF) [103], energy-dispersive X-ray

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spectroscopy (EDS) and inductively coupled plasma (ICP) atomic emission spectroscopy [44,127] have been used to determine the TM content of samples after measuring the magnetic properties. The average magnetic moment per Co atom was then calculated, for which ICP provided more credible results for the TM content [44,94,127]. To a certain extent, the characterization techniques used are keys in studying DMOs. The rather contradictory situation regarding DMO research means that suitable characterization techniques are important. In contrast, incorrect measurements add further confusion to the ﬁeld, which has already been dogged by conﬂicting papers published over the past several years. It is worth pointing out that checking for basic experimental reproducibility, in addition to correct selection of characterization methods, is crucial in the search for high-TC DMOs and for reliable results and discussion. To elucidate the origin of RTFM and to identify relationships between magnetization and the carriers in TM:ZnO, RT resistivity and Hall effect measurements have been carried out using van der Pauw geometry using indium metal for the contact electrodes and current supplied by a dc-voltage source [109]. For some TM:ZnO ﬁlms with quite high resistance, Hall effect measurements were not suitable and the electric properties were determined using a TF Analyser 2000 ferroelectric testing unit and impedance spectroscopy [44,108,128]. 3. Local structure and ferromagnetism of TM-doped ZnO Doped ZnO ﬁlms are either conducting or semiconducting, as characterized by strong coupling between localized d electrons of the TM ions and the extended s and p carriers of ZnO. Thus, the carriers are spin-polarized and can mediate FM ordering of the magnetic moments of TM ions doped into the oxide lattice, i.e., TM ions replace Zn2+ sites. Furthermore, additional electron doping, such as in (Co,Al)- [49,54,129,130], (Mn,Sn)- [131] and (Cu,Ga)doped ZnO [132], as well as hole doping, such as in (Co,Li)[133,134], (Co,N)- [135,136], (Mn,Cu)- [125], (Fe,Cu)- [126], (Co,Cu)- [137] and (Mn,P)-co-doped ZnO [138], has been investigated to further examine carrier effects on RTFM. It was importantly noted that additional carrier doping had no signiﬁcant impact on RTFM enhancement [125,129,131,132], and some ﬁlms even showed RTFM with an inverse correlation between magnetization and electron density [131,137,138]. Thus, inconsistent phenomena and conclusions have been obtained, indicating that the intrinsic FM of TM-doped ZnO systems remains an open question. It is hence of vital importance to clarify the correlation, if any, between carriers and the mechanism of FM inherent to this class of DMOs. On the other hand, defects are supposed to play a primary role on FM ordering in TM-doped ZnO [8]. In fact, an increasing number of studies have demonstrated that the formation of VO and/or Zni [30,44,57,128,139] could lead to RTFM or FM enhancement. For example, the production of VO by depositing TM:ZnO ﬁlms in vacuum [44,111,117] or annealing ﬁlms in a reducing atmosphere (such as vacuum [137,140,141] and Ar/H2 [55]) could greatly enhance RTFM. Annealing of Co-doped ZnO ﬁlms in a Zn atmosphere also resulted in FM enhancement due to Zni [57]. A similar phenomenon was observed by Schwartz and Gamelin [142] for Co:ZnO nanocrystals. It is worth pointing out that a large number of edge dislocations, which were directly observed by HRTEM, enhanced the concentration of VO and/or Zni point defects, and thus increased the RTFM of TM-doped ZnO ﬁlms and nanowires [58,67,94,143]. Moreover, charge-compensating structural defects can form polarons by interaction between hydrogenic electrons and the defects, leading to coupling of the exchange between polarons [8,139].

5

The mechanism related to BMP is discussed in detail in Section 4.2. Preparing TM-doped ZnO ﬁlms with low carrier concentrations and measuring the RTFM of the samples obtained can provide information on whether RTFM is a result of carrier-mediated exchange and whether the TM-doped ZnO can actually be used in spintronics applications. Thus it is worth studying TM-doped oxides without carrier-mediated exchange (>105 V cm) to further our understanding of the origin of RTFM. Insulating TM-doped oxides with FM ordering have become a new topic of research, namely diluted magnetic insulators (DMIs) [44,144]. On the other hand, spin ordering in tunnel junctions made of barriers that are DMIs such as EuS has interested researchers since Esaki et al. [145] observed a large magnetic ﬁeld effect [146]. The topic has attracted renewed interest because recent experiments revealed that the transverse and longitudinal relaxation times for electron spins in insulating quantum dots are in the nanosecond regime, indicating potential as computing elements in quantum electronics [147,148]. However, EuS failed in the practical sense, because its TC is much lower than RT, with little hope of great improvement [7]. RTFM with a TC of 350 K was ﬁrst observed for highly insulating Co- and Cr-doped TiO2 [144,149,150] and TM-doped SnO2 [41,151]. Interestingly, RTFM explained by the BMP mechanism has been widely observed in TM-doped ZnO DMIs [8,30,44,75,94,108,152,153]. For example, RTFM and a high TC of 790 K were observed in (4 at.%) Co-doped ZnO ﬁlms deposited on 1288 Y–X LiNbO3, which is not carrier-mediated, but co-exists with the dielectric state [44]. This phenomenon can be also reproduced in ZnO ﬁlms doped with different Co concentrations and grown on various substrates. Measurement of RTFM in Zn1xFexO (x 0.04) [43] and a lack of correlation between RTFM and conductivity [154] support this observation. Moreover, Xu et al. [155,156] observed that the saturated anomalous Hall resistivity increased with decreasing electron concentration in Co-doped ZnO. Because the spin splitting was proportional to the macroscopic magnetization induced by the Co ions, the positive magnetoresistance observed and the anomalous Hall effect in the insulating regime suggested that FM could be realized in Co:ZnO with low electron concentrations [157]. As indicated above, substitutional TM ions are a necessary but not sufﬁcient condition for FM ordering [27,42,47,79,158], and large variations in carrier concentration and magnetic properties have been reported for otherwise similar samples, indicating a strong dependence of RTFM on the synthesis and processing parameters [17,159]. Indeed, there is large variability in the magnetic behavior of TM-doped ZnO ﬁlms, and the preparation process is not easily reproduced from one growth run to another. The intimate correlation between carrier concentrations and point defects (i.e., VO and Zni) arising from a large number of structural defects such as edge dislocations and stacking faults leads to difﬁculty in elucidating the relations among carriers, structural defects and FM ordering. It is therefore important to design a series of experiments to distinguish the variation of carriers and structural defects and thus to identify the nature of the RTFM. We discuss the inﬂuence of parameters including the doping elements and concentrations, oxygen partial pressure, substrate and its orientation and temperature, deposition rate, post-annealing temperature and co-doping on the local structure and subsequent FM ordering of TM-doped ZnO in this section. This is important for understanding the origin of RTFM in this system. Such understanding can reveal the potential of a class of promising magnetic materials with strong FM at the atomic level, which is attractive for ﬁelds such as semiconductor spintronics, quantum computation and magneto-optical devices [160].

F. Pan et al. / Materials Science and Engineering R 62 (2008) 1–35

6

3.1. Effect of doping concentration on structure and magnetization 3.1.1. Effect of doping concentration on magnetic ordering Incorporation of substitutional TM ions in II–VI ZnO should be less difﬁcult than incorporating them in other semiconductor hosts, since the valence of Zn (2+) and the ionic radius can be readily adopted by many 3d TM ions [19,50]. Different from TMdoped TiO2 [20,144], TM-doped SnO2 [41,151] and similar systems, this favors substitution of the TM ions at the cation site and helps in achieving higher dopant concentrations. It also means that, in a perfectly stoichiometric sample, substitutional TM introduces a magnetic moment without contributing carriers. Thus, this system can give direct insight into the question of whether the FM observed is carrier-mediated or not [44]. Furthermore, TM atoms have different thermal solubility limits in ZnO: in general, the thermal solubility limits of Co and Mn are much higher than those of Ti, V, Cr, Fe, Ni and Cu [50]. In contrast to the solubility of 15 at.%

for Co in ZnO [50,127], much lower limits (e.g., 5 at.%) have been reported for V, Cr, Fe, Ni and Cu are detected [30,43,50], which is consistent with Raman spectroscopic observations [161]. Table 1 presents a list of TM-doped ZnO DMOs, including the substrate, TC, TM content and the corresponding magnetization features at 300 K. As indicated above, the magnetization of TMdoped ZnO exhibits no correlation with the preparation method, which is therefore not listed in the table. There are many references for Co- and Mn-doped ZnO owing to their strong magnetization signals and the stability of Co and Mn ion conﬁgurations [8,23,44]. It is evident from Table 1 that Co-doped ZnO and Mn-doped ZnO are the most popular systems in the literature [162–166], not only because of their high thermal solubility in ZnO ﬁlms, but also because of their high magnetic moments at RT. The dopant concentration has a deﬁnite inﬂuence on the magnetic properties, with more desirable behavior often obtained at comparatively lower concentrations [23,25,37,62,66], with typical concentrations

Table 1 List of TM:ZnO DMOs System

TM content (at.%)

Substrate

TC (K)

Magnetization features

Reference

Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Co:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Mn:ZnO Sc:ZnO Ti:ZnO Ti:ZnO V:ZnO V:ZnO V:ZnO Cr:ZnO Cr:ZnO Fe:ZnO Ni:ZnO Ni:ZnO Cu:ZnO Cu:ZnO Cu:ZnO Cu:ZnO CoFe:ZnO MnCo:ZnO

5 (5–25) 25 3.5 (3.5–11.5) 5 25 35 7 (7–17) 10 5 25 1.3 (0–5) 5 0.3–0.5 2 3 (1–25) 4 5 2 (2–13) 12 (8–31) 10 5 36 2 (0–4) 7 2 7 1 (1–10) 7 4 2 3 (3–15) 10 2.6 2 (0–6) 5 5 9.9 15 (5–15) 0.1–20 13 5 1.1 4 (0–7) 1 (1–25) 1 (1–7) 1 (1–17) 1 (1–8.3) 2 (2–12) 5 <15 10

Al2O3 (1 1 2 0) Al2O3(0 0 0 1) Glass Al2O3(0 0 0 1) Al2O3(0 0 0 1) Al2O3(0 1 2) Si(0 0 1) Al2O3(0 0 0 1) Al2O3 (1 1 0 2) Al2O3(0 0 0 1) Al2O3(0 0 0 1) Si(1 0 0) Al2O3 Si(1 0 0) Al2O3 (1 1 0 2) LiNbO3(1 0 4) ZnO/Al2O3(0 0 0 1) Al2O3(0 0 0 1) Al2O3(0 0 0 1) ZnO(0 0 0 1) Si(1 0 0) Al2O3(0 0 0 1) Glass Al2O3(1 1 2 0) Al2O3(0 0 0 1) Al2O3(0 0 0 1) Al2O3(0 0 0 1) Al2O3(0 0 0 1) Si(1 0 0) Al2O3(0 0 0 1) Si(1 0 0) Al2O3(0 1 2) Al2O3(0 0 0 1) Si(0 0 1) Al2O3 (1 1 0 2) Al2O3 (1 1 0 2) Al2O3(0 0 0 1) Al2O3 (1 1 2 0) Al2O3(0 0 0 1) Al2O3 (1 1 0 2) Al2O3 (1 1 0 2) – LiNbO3(1 0 4) Al2O3(0 0 0 1) Si(1 0 0) Al2O3(0 0 0 1) Glass Glass Si SiO2/Si Al2O3(0 0 0 1)

>280 – >350 300 >300 >300 >300 >300 >300 300 >300 >300 >300 >300 >300 790 >300 >300 >300 – >300 – 425 – 980 78 >300 83 >300 >300 >300 >300 >350 >300 >300 >300 – >350 – – 400 395 400 – >300 390 390 350 >300 >300 >300

0.01 mB/Co Spin glass 0.21 mB/Co 0.7 mB/mole Co 1.0 mB/Co 0.1 mB/Co 0.06 mB/Co 105 emu 2.6 mB/Co 0.8 mB/Co 0.32 mB/Co 1.04 mB/Co 3 mB/Co 0.4 mB/Co 5.9 mB/Co 6.1 mB/Co 1.5 mB/Co/very weak 0.8 mB/Co 0.04 mB/Co AF 0.5 mB/Co Spin glass 0.16 mB/Mn Para 0.04 mB/Mn 20 emu/cm3 0.6 mB/Mn 1.17 mB/Mn 1.4 mB/Mn 4.36 mB/Mn 105 emu 0.075 mB/Mn 4.8 mB/Mn 0.3 mB/Mn 0.3 mB/Sc 0.5 mB/Ti Para 0.4 mB/V Para Spin glass 0.6 mB/Cr 0.73 mB/Cr 0.15 mB/Fe Para 0.37 mB/Ni 0.4 mB/Cu 1.6 mB/Cu 1.8 mB/Cu 0.037 mB/Cu 15 emu/cm3 0.11 mB

[11] [46] [103] [59] [162] [52] [163] [48] [23] [45] [53] [58] [164] [154] [165] [44] [51] [57] [111] [166] [134] [79] [33] [32] [31] [98] [140] [73] [34] [95] [37] [125] [40] [101] [23] [23] [75] [26] [27] [28] [29] [30] [43] [63] [62] [64] [65] [66] [67] [70] [72]

Magnetization features were recorded at 300 K for the TM content listed in column 2. Para, paramagnetism; AF, anti-ferromagnetism.

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Fig. 3. Reported magnetic moments measured in Co-doped ZnO for different dopant concentrations. The lines are guides for the eyes.

of 5 at.% for Co-doped ZnO [11,23,51,58,59,61,128,133,134,152, 167–170], and 1 at.% for Ni-doped ZnO and Cu-doped ZnO [62,64,65]. Most groups have focused on doping concentrations below the solubility limit to obtain true DMOs. Although distinct magnetic behaviors have been obtained by different groups owing to the effects of the preparation method and parameters, it is worth evaluating the effect of dopant concentrations on magnetization from a series of experiments. Magnetic moments measurements reported for Co-doped ZnO at different dopant concentrations are shown in Fig. 3. It is interesting that Co doping in the concentration range 3–5 at.% shows much stronger RTFM than other 3d TM dopants at the same concentrations, as well as the large solubility limit of Co in ZnO, which explains the many reports on 5 at.% Co:ZnO in the literature [11,23,58–60,95,128,133,134,152] (Table 1). Fig. 3 illustrates the Co composition dependent magnetic moment of Co:ZnO. From Fig. 3 and Table 1, one can see that TM:ZnO materials exhibit giant magnetic moments, such as 5.9 mB/Co and 2.6 mB/Co in semiconducting Zn0.97Co0.03O and Zn0.95Co0.05O ﬁlms, respectively [23,165]. These are much stronger than for the corresponding atom in bulk metal, such as 1.7 mB/Co in Co metal. Similarly, 4.36 mB/Mn [95], 4.8 mB/Mn [40] and 5 mB/Mn [171] were obtained for semiconducting Mn-doped ZnO ﬁlms. Several studies on DMOs such as 5 at.% Co- and Cr-doped SnO2 [172,173] and 5 at.% V-doped TiO2 [174] revealed very large magnetic moment of 7.2 mB/Co, 6.0 mB/Cr and 4.2 mB/V, respectively. On the other hand, 6.1 and 3.2 mB/Co were reported for insulating Zn0.96Co0.04O/LiNbO3 [44], which was attributed to magnetoelectric coupling between the Co:ZnO ﬁlm and the ferroelectric substrate and a subsequent super-coupling mechanism in terms of BMPs [94], indicating that carrier-mediated exchange is not a necessary condition for giant magnetic moments. Using the Lande´ g-factor and the total angular momentum quantum number (J) of Co2+ to evaluate the effective magnetic moment [175] according to pﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃﬃ

meff ¼ g mB JðJ þ 1Þ

7

Note that Co dopants at a concentration of 3–5 at.% can lead to giant magnetic moments [23,44,94,165]. If carrier-mediated exchange were the origin of RTFM, there would be no correlation between the dopant concentration and the average magnetic moment, because carriers are not local; instead, carriers can play a similar role in mediating RTFM in TM-doped ZnO with different dopant concentrations. On the contrary, if RTFM stems from BMP on the basis of defects, the polaron percolaton threshold requires a minimum mismatch between the dopant content and the defect content. In this case, a certain dopant concentration is a precondition for the appearance of a giant magnetic moment. Suitable dopant concentrations and defects effectively couple to form a hydrogenic orbital arising from an electron associated with a particular defect, and the hydrogenic orbital tends to sufﬁciently spread out to overlap with a large number of BMPs, so that few isolated polarons exist. More remote dopants interact too weakly with both polarons to be of any importance. This occurs in samples with lower doping levels, e.g., 1 at.% Co-doped ZnO, which produces a much lower magnetic moment for weak interaction. On the other hand, for ﬁlms doped with higher levels of Co, an increase in dopant concentration (e.g., 15 at.% Co-doped ZnO) leads to a rapid decrease in the moment due to enhanced dopant–dopant association and thus progressive orbital moment quenching [44,172]. In this case, anti-FM or FM appears when there are continuous paths throughout the crystal joining nearest-neighbor magnetic cations [8]. This reveals how RTFM can be explained by a BMP model based on defects. It was reported that Zn0.97Mn0.02Al0.01O ﬁlms showed RTFM with 4.36 mB/Mn, for which the 2 at.% Mn doping to yield a giant magnetic moment is lower than the level of Co dopant required (3–5 at.%). Interestingly, in contrast to the Co doping concentration range (1–25 at.%) reported in many publications [11,45,46,162,165,176], Mn is generally used as a dopant at a much lower concentration of 1–10 at.% [32,140], as illustrated in Table 1. Unlike those of Co- and Mn-doped ZnO, the average magnetic moment of other TM-ZnO system, such as Cr-, Ni- and Cu-doped ZnO, monotonically decreases with increasing dopant concentration from 1 at.% [30,54,62,64–66,71,102,177], as shown in Fig. 4. This phenomenon may be due to the low solubility of these dopants in ZnO; thus, low concentrations (e.g., 1 at.%) correspond to the strongest magnetization and higher concentrations correspond to lower magnetic moments. Therefore, 3–5 at.% Co, 2 at.% Mn and 1 at.% Cr, Ni, and Cu doping levels in ZnO show the most robust magnetic moments, which

(1)

leads to a value of 6.6 mB/Co. In contrast, the net spin moment calculated from the spin-only formula mmax = gmBS is 3 mB. This phenomenon can be explained by spin and orbital angular moment contributions [44,172–174].

Fig. 4. Reported magnetic moments measured in Cu-, Ni- and Cr-doped ZnO for different dopant concentrations. The lines are guides for the eyes.

8

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are critically dependent on the different solubility of these TM dopants in ZnO ﬁlms. 3.1.2. Formation of a secondary phase It is considered necessary to achieve single-phase DMOs for use in devices, in which the carriers can realize spin polarization. If precipitates or clusters below detectable size limits are present and responsible for the FM, and if the carriers do not mediate the FM interaction, this then can limit the usefulness of spintronics fabricated from such materials [18]. Therefore, it is important to detect RTFM arising from TM2+ replacement of Zn2+ sites or secondary phases, which is signiﬁcant for studying practical DMOs. However, the detection of trace amounts of TM atoms requires highly sensitive characterization techniques. The formation of secondary phases mainly depends on the TM solubility limit in ZnO. The thermal solubility limit for Co in ZnO is less than 15 at.% and Co ions systematically replace Zn ions in the ﬁlm without changing the wurtzite structure [50,127]. This is supported by the observations of Prellier et al. [59], who reported that the c-lattice parameter obeys Vegard’s law up to the solubility limit of 10 at.% Co, after which a turning point suggests the breakdown of Vegard’s law. However, different values of the dopant solubility have been obtained; for example, limits in the range 25–50 at.% have been reported for Co-doped ZnO ﬁlms grown by PLD. Kim et al. [178] claimed that the limit is 40 at.%, close to the results reported by Ueda et al. [11]. This difference is most likely attributable to the preparation process, which allows a different distribution of Co within the ZnO wurtzite structure [16,17,19]. The equilibrium solubility limit of Mn in ZnO was determined by Jin et al. [50] to be 20 at.%. The high solubility of dopants seems to result from the growth conditions. It was also found that the formation energy of dopants can have a strong dependence on the local lattice constant, depending on the size mismatch of the dopant and the host atom. For example, several authors claimed that the non-equilibrium nature of the laser material interaction allows doping at high Mn content (up to 36 at.%) [50,79]. Meanwhile, FM ordering has been reported for Mn concentrations between 2.2 and 30 at.% [33,179,180], whereas both the magnetic moment per Mn ion and the FM TC seem to decrease with increasing Mn concentration, i.e., the moment dropped from 0.71 to 0.25 emu cm3 when the Mn content increased from 10 to 30 at.% [76]. In addition, the FM TC was above RT for all concentrations x 5 at.% [33,180], but was usually well below 100 K (45 K) for higher concentrations [179]. Besides RTFM arising from TM ion replacement of Zn2+, secondary phases such as TM metal and TM-based oxides are frequently found to correspond to the magnetization measured. Several different methods have been used to determine the local structure of doping elements and to detect secondary phases, which is important for identifying whether a TM-doped ZnO is a genuine DMO or not. The methods for detecting secondary phases are summarized as follows. (1) In the most simple approach, diffraction peaks of secondary phases present in the TM-doped ZnO matrix can be observed in the XRD pattern. In contrast to Zn1xFexO (x 0.04) samples, which show a single phase with a wurtzite structure (Fig. 5a), as x increases to 0.07, a diffraction peak of a secondary phase appears in the XRD pattern, which is attributed to Fe3O4, as shown for Zn0.93Fe0.07O in Fig. 5b [43]. This observation is in good agreement with the low solubility of Fe in ZnO, e.g., <5 at.% [19,50]. In contrast, no secondary phases are found in Co-doped ZnO, even for 10 at.%, for the same preparation parameters and substrates [44], indicating the different solubility of Co and Fe in ZnO. The presence of Co3O4 and

Fig. 5. (a) XRD pattern of a Zn0.96Fe0.04O ﬁlm. The inset shows XPS spectra of Zn0.96Fe0.04O (dashed line) and Zn0.93Fe0.07O (solid line). (b) XRD pattern of a Zn0.93Fe0.07O ﬁlm. The Fe3O4 secondary phase is marked by solid circles.

CoO peaks has been reported for doping concentrations above 25 and 22 at.% using sol–gel methods and reactive magnetron co-sputtering, respectively [56,163]. The formation of Co metal nanoclusters for x 12 at.% in samples grown by sol–gel and rf sputtering techniques has also been observed [76]. The observation of a CoFe(1 1 0) diffraction peak for Zn1x(CoFe)xO O ﬁlms indicates the formation of CoFe metal particles, which was further conﬁrmed by a clear change in the slope of the d(0 0 2) variation with the CoFe content around x = 0.2 [70]. Likewise, only a Ni(1 1 1) peak was detectable, revealing that fcc Ni nanocrystals are highly oriented with respect to the ZnO host matrix [181]. However, secondary phases in trace amounts and small clusters are hardly detected by XRD, so it should be remembered that the absence of evidence is not evidence of absence [139]. (2) HRTEM and SEM are often performed to investigate the different phases that might have formed in the nanosize range and to determine the state of Co atoms that cannot be detected by XRD. For example, Co3O4 clusters were directly found by both HRTEM [111] and FE-SEM [109] in Co-doped ZnO ﬁlms. ZnCo2O4 and/or CoO clusters with a diameter of 5 nm have been observed by HRTEM [110], which was conﬁrmed by SAD results. Co metal clusters with an average diameter of 10 nm have also been observed by HRTEM in 15 at.% Co-doped ZnO ﬁlms [76]. (3) The valence state of a TM dopant can be determined by EELS and XPS [33,43,44,52,108,182]. The advantage of EELS is that the valence state can be measured in combination with observation of HRTEM images [33,108]. Wei et al. [43] observed that both the Fe 2p1/2 and 2p2/3 peaks of Zn0.93Fe0.07O were between those of Fe3+ and Fe2+, and closer to Fe3+, indicating a

F. Pan et al. / Materials Science and Engineering R 62 (2008) 1–35

Fig. 6. Experimental (solid circles) and MS calculated (solid line) Fe K-edge XANES spectra of (a) Zn0.96Fe0.04O and (b) Zn0.93Fe0.07O ﬁlms.

mix of 2+ and 3+ valence states with predominantly Fe3+, as shown in the inset of Fig. 5a. In contrast, for a doping concentration of x 0.04, such as in Zn0.96Fe0.04O, the positions of the Fe 2p1/2 and 2p2/3 peaks are apparently different from those of Zn0.93Fe0.07O (inset of Fig. 5a), revealing that Fe is in the 2+ valence state [43]. (4) TM K-edge [34,44,55,115], L-edge [78,97,116,117] and O Kedge [37,97,109] XAS and its interpretation can provide a structural ﬁngerprint to determine TM sites in ZnO [8]. As discussed above, for XRD and XPS changes for different doping concentrations, Fe K-edge XANES can be used to determine the local structure of Fe in the ﬁlms and conﬁrm the preliminary conclusion. Experimental (solid squares) and full multiplescattering (MS) ab initio calculated (solid line) XANES for Zn0.96Fe0.04O and Zn0.93Fe0.07O are shown in Fig. 6a and b, respectively [43]. The most interesting result here is that the calculated spectrum has three main features that accurately reﬂect those of the experimental spectrum, denoted by A (A0 ), B (B0 ) and C (C0 ) in order of increasing photon energy in Fig. 6a and b. It should be pointed out that the calculations for Fig. 6a and b were produced for a cluster within a sphere radius of 6 A˚ containing 77 atoms from the central Fe atom that replaces a central Zn in the ZnO atomic arrangement and 89 atoms from the Fe atom that occupies a tetrahedral site in the cubic structure of Fe3O4. There are two main differences between them: (i) the curve for Zn0.96Fe0.04O has an inﬂection at approximately E–E0 = 5 eV (E0 = 7112 eV), whereas the curve for Zn0.93Fe0.07O is ﬂat at that position (A and A0 ); (ii) peak C is

9

slightly more intense than peak B for Zn0.96Fe0.04O, whereas peak B0 is apparently less intense than C0 for Zn0.93Fe0.07O. Moreover, the curve for Zn0.93Fe0.07O has all the features of Fe3O4 [183]. It is hence concluded that when x 0.04, Fe is in the 2+ valence state and substitutes for Zn in the wurtzite lattice of ZnO. When x increases beyond 0.04, Fe mainly exists in the form of Fe3O4. Moreover, Fe2+ substituting for Zn2+ in Zn0.96Fe0.04O and the Fe3O4 secondary phase in Zn0.93Fe0.07O are responsible for 0.15 and 1.0 mB/Fe, respectively. (5) Most importantly, MCD (including X-ray MCD and optical MCD) is considered the best method for evaluating whether a TM-doped ZnO ﬁlm is a real DMO or not [85,139]. X-ray MCD (XMCD) is commonly used, and is the difference in absorption of left- and right-hand-circularly polarized X-rays by a magnetized sample. Thanks to simple sum rules, XMCD can also provide quantitative information about the distribution of spin and orbital angular momenta. Optical MCD measures the relative difference in the circular polarization-dependent optical absorption under an applied magnetic ﬁeld, i.e., 2(K K+)/(K + K+), where K and K+ are optical absorption coefﬁcient for s and s+ polarizations, respectively. The optical absorption depends on the circular polarization ascribed to Zeeman splitting of the band structures. MCD spectroscopy sensitively detects DMOs and reﬂects their electronic structure and identiﬁes secondary phases, taking the strong sp–d exchange interaction of DMOs and subsequent Zeeman splitting of the optical transitions into account [85,158]. The part of the MCD signal that linearly increases with the magnetic ﬁeld (H) represents the paramagnetic component, whereas the MCD signal that persists at H = 0 Oe is the FM component. MCD spectra showing a multiplet structure, unlike spectra of TM metals with a single peak, reveal that TM-doped ZnO shows FM ordering that is not due to metallic TM secondary phases, but arises from TM ion substitution in Zn sites [60,121]. For example, Mn-, Co-, Ni- and Cu-doped ZnO ﬁlms prepared by PLD show pronounced negative optical MCD peaks near 3.4 eV attributed to the bandgap of the host ZnO semiconductor, indicating that ZnO alloyed with Mn, Co, Ni, and Cu is a true DMO with strong exchange interaction between sp-band carriers and localized d electrons [184]. In contrast, the magnetic moment was observed to be 0.1 mB/Co in Co-doped TiO2, but the XMCD spectral line shape is nearly identical to that of Co metal, showing that FM is induced by a small amount of clustered Co, consistent with the Co metal cluster observed by FE-SEM [185]. Likewise, the use of advanced element and orbital selective techniques allows researchers to observe the distinct spectral signature of Co in ionic or metallic states to assign FM phases to the presence of Co metal clusters in ﬁlms doped with Co concentrations >25 at.% [176]. (6) ESR [99,122,123] and Mo¨ssbauer spectra [42,124] also provide information on the local structure of the TM dopant [131]. The ESR spectrum of Zn1xMnxO (x = 0.001) exhibits the hyperﬁne and ﬁne-structure lines of Mn2+, which broaden for higher Mn content (x = 0.035) and change to a very broad single line due to dipolar interactions between Mn ions at higher local concentrations [122]. This phenomenon indicates that ESR can reﬂect local structure variations for a TM dopant. However, a large number of TM neighbors in the ZnO lattice (n 5) have no effect on the line width observed [123]. Mo¨ssbauer spectra cannot only reveal FM (sextet) and paramagnetic (doublet) phases in Fe:ZnO, but are also useful in detecting secondary phases, since the Mo¨ssbauer spectrum of Fe:ZnO is different from the typical sextet of FeOx [124]. (7) Optical properties (e.g., Raman spectroscopy and optical transmittance/absorption measurements) are commonly used

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to characterize TM-doped ZnO because they are intimately related to the defects and electronic structure of TM:ZnO [48,53,108,109,111,161]. For example, spectra show characteristic absorption edges around 567, 613 and 662 nm resulting from the spin-orbit split ligand-ﬁeld transition of tetrahedrally coordinated Co2+ ions due to 4A2(F) ! 2A1(G), 4A2(F) ! 4T1(P) and 4A2(F) ! 2E(G), respectively. Once Co metallic clusters exist in a ﬁlm, the spectrum exhibits an octahedral Co0 4 T1g(F) ! 4T1g(P) at approximately 513 nm [48,109]. Thakur et al. [53] reported that for x 5 at.% in Zn1xCoxO, at which the magnetic moment becomes negligible, broadening of the Elow 2 Raman mode was also negligible, suggesting rejection of Co ions from Zn sites and the appearance of secondary phases. This is consistent with the observation that a decrease in Raman intensity indicates the formation of TM clusters [71,170]. Moreover, as the Cu concentration increased from 1 to 3 at.% in Cu-doped ZnO, Raman spectra showed that besides a broad asymmetric mode at 600 cm1 attributed to the E1 (LO) mode of ZnO, a distinct peak at 634 cm1 arising from the Bg mode of CuO emerged, indicating the formation of a secondary phase of CuO [65]. These methods with different physical bases have been demonstrated to be effective in detecting second phases. However, it is difﬁcult to determine a trace amount of TM dopant using a random selection of characterization techniques. The best way to investigate the state of TM dopants may be to use several methods, and evidence of absence can only become a reality with the advent of a combination of very sensitive tools. 3.2. Substrate effects on magnetization 3.2.1. Substrate-dependent local structure and magnetic behavior As discussed above, several reports and reviews reveal that the magnetic behavior of DMOs appears to be dependent on preparation conditions such as the doping concentration, substrate temperature, and oxygen partial pressure [11,19,44,46]. However, few studies have directly considered the substrate-dependent magnetic behavior of DMOs [94]. The best example is a review in which statistical data imply that the magnetization of Co-doped TiO2 ﬁlms is related to the crystal structure and the substrate [19]. For Co-doped TiO2 ﬁlms grown on Si or Al2O3 substrates, the average magnetic moment per Co atom is <1.0 mB in all cases. For samples using SrTiO3 as a substrate, values were scattered over more or less the whole range of values reported, with a slight accumulation in the lower region. However, for ﬁlms grown on LaAlO3 substrate with an anatase structure, most of the values accumulated in the upper region, i.e. >1.2 mB/Co, but with some scatter between 0.23 and 1.72 mB/Co [19]. The substrates commonly used for deposition of TM:ZnO ﬁlms are Al2O3 [11,23,45,46,52,111,162], Si [58,67,102,134,154,163], fused quartz [33,94,103] and ZnO [51,105,186]. The reported magnetic moment per Co atom for Co:ZnO ﬁlms grown on different substrates is shown in Fig. 7. As the ﬁgure shows, the substrate inﬂuence on magnetic ordering is not clear. The data summarized in Fig. 7 are from results obtained by different groups using different preparation parameters. It is known that the local structure and magnetization of DMOs are intimately related to experimental conditions (e.g., preparation methods and parameters). Therefore, the preparation method, oxygen partial pressure, substrate and temperature are all factors that can explain the scatter from 0 to 5.9 mB/Co and from 0.1 to 1.04 mB/Co on Al2O3 and Si substrates, respectively. However, TM-doped ZnO ﬁlms deposited on different substrates with the same deposition parameters show distinct magnetization. For example, (0.6 at.%)

Fig. 7. Reported magnetic moment per Co atom for Co:ZnO ﬁlms grown on different substrates.

Cu-doped ZnO ﬁlms exhibited 1.6 and 1.2 mB/Cu at 300 K on quartz and sapphire, respectively [65]. RTFM with a magnetic moment of 1.5 mB/Co was observed for Zn0.95Co0.05O ﬁlms deposited on ZnO, whereas for the ﬁlms grown on Al2O3 only weak FM signals were detected that could not be unambiguously separated from those of the substrate. This difference can be explained by different structural properties and defect densities in the ﬁlms [51]. Thus, it seems that the substrate most likely impacts the local TM structure and magnetization of TM-doped ZnO. To systematically study substrate effects on the local structure and magnetic behavior of TM-doped ZnO, a series of experiments with substrate variation was designed [94]. Substrate-dependent magnetization of insulating Zn1xCoxO (x = 1, 4 and 10 at.%) ﬁlms has been highlighted in Fig. 8, for which the substrates were 1288 Y–X LiNbO3 (128LN), 648 Y–X LiNbO3 (64LN) and 368 Y–X LiTaO3 (36LT) ferroelectric crystals, SiO2(1 0 1) piezoelectric crystal, some other single crystals [e.g., Al2O3(0 0 1), Si(1 1 1), and NaCl(1 0 0)], and amorphous glass. It is evident that the substrate has a profound inﬂuence on the magnetic behavior of Co-doped ZnO ﬁlms. Zn0.96Co0.04O ﬁlms are taken as an example. Interestingly, whereas RTFM with giant magnetic moments of 6.1, 2.5, and 3.2 mB/Co is observed for Co:ZnO ﬁlms on 128LN, 64LN and LT ferroelectric crystals, respectively, the ﬁlm on SiO2 piezoelectric

Fig. 8. Room-temperature magnetic moment expressed in Bohr magnetons (mB) per Co atom for representative Zn1xCoxO (x = 1, 4 and 10 at.%) ﬁlms grown on 128LN, 64LN, LT, SiO2, Al2O3, Si, NaCl, and glass substrates.

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crystal shows a much lower magnetic moment of 1.1 mB/Co. Moreover, the magnetic moment decreases to 0.62, 0.66, and 0.24 mB/Co for ﬁlms deposited on Al2O3, Si, and NaCl single crystal, respectively. The ﬁlm on the glass substrate does not exhibit any FM signature [94]. The diamagnetic background of the substrate has been subtracted from all of the magnetization data, indicating that the difference in moments in Co:ZnO ﬁlms does not arise from the substrate. A similar tendency was replicated in Co:ZnO ﬁlms with other Co concentrations (e.g., 1 and 10 at.%). From XPS and qualitative analysis of XANES spectra it is known that the difference in moments cannot be ascribed to different Co forms in Co:ZnO ﬁlms grown on different substrates, since all the Co ions are in the 2+ valence state replacing Zn sites. The rather robust moments on ferroelectric crystals compared to those on the other substrates are attributed to modulation of the magnetic moment of the Co dopant due to magnetoelectric coupling between the Co:ZnO ﬁlm and ferroelectric crystals [94,187]. Capacitance–voltage (C–V) measurements revealed that the typical C–V curve for Co:ZnO samples exhibits hysteresis with a large memory window of 0.6 V, which can be caused by trapping/ detrapping of charges at defect sites. Evidence of defects with trapped electrons in Co-doped ZnO would advance our understanding of the BMP mechanism. The diffusion or motion of defects such as VO is then driven by the presence of stress/strain and by an electric ﬁeld effect arising from spontaneous polarization [149] to realize an appropriate distance between Co2+ and defects and to achieve effective interactions between them. This process of defect diffusion or motion is referred as the reorganization of defects. Only this condition is met, almost all doping Co2+ can effectively interact with the defects producing few isolated Co2+, and the giant magnetic moment of 6.1 mB/Co in Co:ZnO/128LN, which is slightly lower than the ideal magnetic moment of 6.6 mB/Co calculated using Eq. (1), can become a reality. In contrast, some Co2+ and defects with inappropriate distances (e.g., large distances) in Co:ZnO ﬁlms grown on non-ferroelectric substrates (without reorganization of defects) are hard to achieve effective interactions, leading to the existence of more isolated Co2+ and the measured magnetic moment far below the ideal one. It is known that donor spin from well-organized defects strongly correlated with Co2+ within its orbit can mediate effective interactions between them based on a Heisenberg exchange Hamiltonian. The donors then tend to shape the BMP, coupling Co2+ ions within their orbits, which try to sufﬁciently spread out to overlap and interact with adjacent BMPs to realize magnetic ordering, resulting in rather robust moments that can be understood in terms of a supercoupling mechanism [44]. It is worth pointing out that reorganization of the defects could partly occur due to the comparatively weak polarization of the SiO2 piezoelectric substrate, leading to a lower moment per Co atom in Co-doped ZnO ﬁlm, e.g., 1.1 mB/Co. Moreover, although the Co dopant exists in a 2+ state when substituting for Zn2+, randomly distributed defects would lead to very weak moments or no RTFM in Co-doped ZnO ﬁlms on other single substrates (e.g., Al2O3, Si, and NaCl) and glass substrate, respectively [94]. It is thus concluded that the substrate strongly affects the magnetization of TM-doped ZnO ﬁlms. 3.2.2. Substrate orientation-induced distinct magnetization An interesting point in Fig. 8 is that signiﬁcant differences in moments are observed for Co:ZnO ﬁlms grown on the same type of substrate. For example, Zn0.96Co0.04O ﬁlm grown on 64LN exhibits a comparatively lower moment of 2.5 mB/Co compared to 6.1 mB/Co for ﬁlm grown on 128LN. The origin of magnetization differences on the same type of ferroelectric substrate with only a difference in cut

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Fig. 9. (a) Experimental and (b) MS calculated XANES spectra of Co:ZnO ﬁlms grown on 1288 Y–X LiNbO3 (top) and 648 Y–X LiNbO3 (bottom).

direction could be explained by different Co–O bond lengths [94]. Experimental Co K-edge XANES spectra of Co:ZnO ﬁlms are compared in Fig. 9 with the ab initio calculated XANES using full MS theory. The atomic arrangement difference between Co:ZnO on 128LN and on 64LN is the difference in Co–O bond length on the basis of the lattice parameters inferred from XRD results [94]. The calculated spectrum displays three main features accurately replicating those of experimental XANES spectra, denoted by A, B, and C in order of increasing photon energy. For Co-doped ZnO ﬁlm on 128LN, it was found that the Co photoabsorber in tetrahedral coordination is surrounded by a ﬁrst shell comprising one oxygen atom at 1.820 A˚ and three oxygen atoms at 2.069 A˚. This leads to a difference of 0.03 and 0.04 between features B and C in the normalized absorption experimental and MS calculated spectra (upper spectra in Fig. 9a and b). However, for Co-doped ZnO ﬁlm on 64LN, the center Co is surrounded by a ﬁrst shell comprising one oxygen atom at 1.845 A˚ and three oxygen atoms at 2.097 A˚. This results in greater differences of 0.08 and 0.07 for the experimental and calculated spectra (lower spectra in Fig. 9a and b). These results demonstrate that a slight change in Co–O bond length in Co:ZnO ﬁlms has a profound inﬂuence on electron movement and scattering in the vicinity of the Co atom. Subsequently, electron spin and the corresponding moments are likely to be sensitive to the Co–O bond length. In fact, ab initio calculations predicted that FM coupling would be affected by the Mn–O bond length in La0.67Ca0.33MnO3d ﬁlms [94], and that energy differences between FM and anti-FM alignment would be very small in Co-doped ZnO systems, with competition between FM and anti-FM coupling [84,188]. ZnO ﬁlms doped with different Co concentrations were deposited on LT substrates with three different orientations [LT(1 1 0), LT(0 1 2) and LT(0 1 8)] by dc reactive magnetron co-sputtering. The variation in magnetic moment on different orientations as a function of the Co concentration is presented in Fig. 10. Zn0.962Co0.038O ﬁlms deposited on LT(1 1 0), LT(0 1 2) and LT(0 1 8) substrates exhibited different magnetization of 1.21, 2.42, 0.65 mB/Co, respectively, as shown in the inset of Fig. 10. It is likely that the magnetization differences for Co-doped ZnO ﬁlms on three oriented LT substrates can be attributed to the ﬁlm microstructure and the Co–O bond length, which primarily depends on mismatch between the ﬁlm and the different oriented substrates [107]. Therefore, as well as the substrate itself, its orientation has a profound inﬂuence on FM ordering of TM-doped ZnO.

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Fig. 10. Saturated magnetization of Co:ZnO as a function of Co concentration for ﬁlms grown on LiTaO3 substrate in different orientations: LT(1 1 0), LT(0 1 2) and LT(0 1 8). The inset shows hysteresis loops of Zn0.962Co0.038O ﬁlms deposited on LT(1 1 0), LT(0 1 2) and LT(0 1 8).

3.3. Effects of substrate temperature and ﬁlm deposition rate on magnetization Sharma et al. [33] reported that substrate temperature affects the magnetization of Mn-doped ZnO ﬁlms. Films prepared at a low processing temperature exhibited RTFM, whereas samples at standard high temperatures (T > 700 8C) showed clustering and no RTFM. This is in agreement with work carried out by Kim et al. [46], who found that a high substrate temperature of 700 8C led to the appearance of secondary phases. In addition, Fe-doped ZnO prepared by ion implantation at 253 K was a true DMO, whereas preparation at 623 K led to the formation of particles inside the host matrix responsible for the FM properties [42]. Very recently, it was found that Co-doped ZnO ﬁlms grown at <200 8C exhibited stronger FM than those deposited at higher temperatures [111]. These observations indicate that a high preparation temperature is detrimental to RTFM of DMOs, and even results in the formation of secondary phases within the host matrix. On the other hand, different from other DMOs such as TM-doped TiO2 [189,190], the local structure and magnetic properties of TM-doped ZnO ﬁlms are more sensitive to the substrate temperature and ﬁlm deposition rate, primarily due to temperature-dependent Zn loss during the deposition process [58]. Song et al. [109] investigated how the electrical properties and structural defects of TM:ZnO samples are related to FM ordering to identify factors that affect RTFM. A suitable combination of deposition rate and substrate temperature is key in determining the crystallinity and corresponding magnetization of Co-doped ZnO ﬁlms deposited by rf magnetron co-sputtering at 0.15 A˚/s (slow growth) and 0.6 A˚/s (fast growth) at 573 K [high substrate temperature (Tsub)] and RT (low Tsub), as presented in Table 2 [109]. For this series of ﬁlms, the sample prepared with slow growth and

high Tsub exhibits the lowest saturation magnetization of approximately 0.25 mB/Co. For the ﬁlm grown at RT, the magnetic moment slightly increases to 0.39 mB/Co. For faster deposition rates, much larger magnetic moments of 0.73 and 0.91 mB/Co are observed for the samples grown with high and low Tsub, respectively. XRD results show that the intensity of the ZnO(0 0 2) diffraction peak decreases with increasing deposition rate and/or decreasing substrate temperature, reﬂecting a decrease in ZnO crystallinity. This was demonstrated by F scans, for which comparison of the full width at half-maximum (FWHM) intensity for peaks with sixfold symmetry revealed a decrease in crystal quality and an increase in structural defects with increasing deposition rate and decreasing Tsub [109]. Experimental and calculated Co K-edge XANES spectra are presented in Fig. 11a and b, respectively. The height difference between the characteristic B1 and C1 peaks increases as deposition rate increases and Tsub decreases. This is reproduced by the MS calculation (Fig. 11b). The introduction of oxygen vacancies in the ﬁrst coordination sphere results in an obvious height difference between B1 and C1 for the sample deposited with slow growth and low Tsub, as well as the sample grown with fast growth and high Tsub. For the sample prepared with fast growth and low Tsub, the existence of oxygen vacancies in the ﬁrst coordination sphere and a Zn interstitial between the ﬁrst two coordination spheres would result in line shapes consistent with the experimental data. This is supported by experimental and calculated O K-edge XAS spectra (Fig. 11c and d, respectively). The calculations reﬂect less of a Zn loss during deposition for the sample prepared with fast growth and low Tsub [58], whose spectrum is similar to that of an ideal stoichiometric ZnO, and shoulder A2 at 531.2 eV appears and increases in intensity from the sample prepared with slow growth and low Tsub to the sample grown with fast growth and low Tsub, corresponding to an increase in defect concentration in the ﬁlms, in good agreement with microstructural characterizations [191]. The magnetic moment of Co-doped ZnO ﬁlm increases with increasing structural defects and electron concentration (ne). It is worth pointing out that this was the basis for the previous concept of carrier-mediated RTFM in DMOs, because the correlation of FM ordering and carrier concentration was emphasized. The true origin of FM ordering, i.e., structural defects, was neglected in this carrier-mediated hypothesis, which seriously misguided extensive research into DMOs. To determine the intrinsic FM in a Co-doped ZnO system, the sample prepared with slow growth and high Tsub was annealed for 1 h at 400 8C, resulting in an increase in ne to 2.9 1019 cm3. The crystallinity of the Co-doped ZnO ﬁlm also improved on annealing, revealing that the number of structural defects such as edge dislocations and stacking faults decreased [109]. The saturation magnetization slightly decreased from 0.25 to 0.19 mB/Co, revealing that an increase in ne cannot enhance the magnetization of Co-doped ZnO after annealing at 400 8C. Instead, the magnetic moment per Co atom decreased with increasing crystallinity, which allows identiﬁcation of a direct relationship between the FM of Co-doped ZnO and the presence of structural defects. This implies that the established theory of carriermediated exchange, which works well for p-type GaMnAs, is not

Table 2 Resistivity (R), electron concentration (ne) and magnetic moments (M) for Co:ZnO samples prepared by rf magnetron co-sputtering at low/high deposition rate and low/high substrate temperature (Tsub) Sample

Rate (A˚/s)

Tsub (K)

R (V cm)

ne (cm3)

M (mB/Co)

Slow growth high Tsub Slow growth low Tsub Fast growth high Tsub Fast growth low Tsub

0.15 0.15 0.6 0.6

573 RT 573 RT

>103 5.6 1.9 101 1.7 102

No Hall signal 2.3 1018 1.5 1019 5.2 1019

0.25 0.39 0.73 0.91

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Fig. 11. (a) Experimental and (b) calculated Co K-edge XANES spectra for Co:ZnO ﬁlms (E0 = 7708.8 eV). (c) Experimental and (d) calculated O K-edge XAS for Co:ZnO ﬁlms. Spectra correspond to samples grown with slow growth and high Tsub (LH), slow growth and low Tsub (LL), fast growth and high Tsub (HH) as well as fast growth and low Tsub (HL) from bottom to top.

generally applicable to TM-doped ZnO. Instead, structural defects play a primary role in RTFM of TM-doped ZnO, and carriers involved in carrier-mediated exchange are natural by-products of the creation of defects, which is demonstrated in publications with increasing frequency [60,76]. 3.4. Inﬂuence of oxygen partial pressure on magnetic ordering In general, both the oxygen partial pressure (PO2 ) and the substrate temperature (Tsub) play important roles in the microstructure of DMO ﬁlms, which subsequently affects the magnetic properties. PO2 and Tsub affect the formation of secondary phases [46,33]. Lee et al. [98] reported that ne dramatically decreased with increasing PO2 up to an Ar/O2 ratio of 1:1 for the preparation of Mn:ZnO ﬁlms by rf magnetron sputtering, which showed intrinsic FM with TC of 78 K. However, ﬁlms grown under O2 overpressure showed various secondary phases, such as ZnMn3O4, MnO2 and ZnMnO3. Films deposited at lower temperature (600 8C) and higher oxygen pressures (105) showed a homogeneous distribution of Zn0.75Co0.25O, leading to spin-glass behavior, whereas mixed phases of wurtzite ZnO with rock-salt CoO and hexagonal Co phases coexisted at higher temperatures and lower pressures, accompanied by RTFM [46]. On the other hand, variation of PO2 can lead to a change in magnetic behavior and electric properties. On increasing PO2 from 1.0 104 to 5 Pa, the resistivity of Zn0.95Co0.05O ﬁlms increased from 2.15 102 to 2.06 104 V cm, corresponding to a change in magnetization from RTFM to paramagnetism, revealing that oxygen vacancies are absolutely necessary to induce FM coupling in TM-doped ZnO [168]. Philip et al. [243] observed that the electric and magnetic behavior of Cr-doped In2O3 ﬁlms – from FM metal-like to FM semiconducting to paramagnetic insulators – can be controllably tuned by P O2 . It seems that PO2 plays an important role in the magnetization of DMOs. A transition from DMI to DMO was realized in Zn0.96Co0.04O ﬁlms grown by dc reactive magnetron co-sputtering by decreasing P O2 from 4.2 103 to 1.0 103 torr for samples 1 and 2,

corresponding to a deposition rate of 3.0 and 2.0 nm/min, respectively [112]. Fig. 12 shows the PO2 -dependent RTFM and resistivity (R) of the ﬁlms. Both the magnetization and resistivity decrease with decreasing PO2 . Insulating sample 1 (0.62 mB/Co) exhibits more robust RTFM than semiconducting sample 2 (0.45 mB/Co), revealing that an increase in carrier concentration does not enhance RTFM. A similar trend was observed by Chakraborti et al. [137] and Ivill et al. [131,138]: an increase in conductivity did not increase the magnetic moment; instead, an increase in carrier concentration led to weak magnetization, suggesting that a mechanism mediated by free carriers is not a feasible explanation for FM in TM:ZnO ﬁlms. To interpret the intrinsic RTFM in this system, microstructural and local structural characterizations were conducted [112]. XRD spectra in Bragg–Bretano, rocking curve and F scan geometries suggested that the ﬁlms were highly crystalline, as shown in Fig. 13. The XRD pattern of sample 1 (Fig. 13a) has a weaker Co:ZnO(0 0 2) peak than that for sample 2 (Fig. 13b). Comparison of

Fig. 12. Magnetic moment and resistivity of Zn0.96Co0.04O ﬁlms as a function of P O2 . The inset shows magnetic hysteresis loops (M–H) for Co:ZnO ﬁlms 1 and 2 prepared at high and low P O2 , respectively.

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Fig. 13. XRD spectra of Co:ZnO ﬁlm for (a) sample 1 and (b) sample 2. X-ray rocking curves from the wurtzite (0 0 2) plane for (c) sample 1 and (d) sample 2. F scan of {1 0 1} family peaks for (e) sample 1 and (f) sample 2.

FWHM values for X-ray rocking curves for the wurtzite(0 0 2) plane for samples 1 and 2 reveals an improvement in crystallinity of sample 2 (0.188 decrease in the FWHM value; Fig. 13c and d). The F scan for the (1 0 1) plane of sample 2 grown at lower PO2 is similar to that of sample 1 to a certain extent, but with more symmetrical peaks. These results indicate that sample 2 is more crystalline and has fewer structural defects such as edge dislocations and stacking faults due to the comparatively slow rate of ﬁlm growth [189,190]. Fig. 14 presents Co K-edge XANES for samples 1 and 2 and for reference materials (Co metal, CoO, Co2O3 and Co3O4). It is evident that Co in Co:ZnO is not in the form of Co metal or Co-based oxides such as CoO, Co2O3 and Co3O4. The calculated spectra displays three main features accurately replicating those of the experimental spectra, denoted by A, B, and C in order of increasing photon energy (insets a and b of Fig. 14). These spectra reveal that for sample 1 the Co photoabsorber in tetrahedral coordination is surrounded by a ﬁrst shell comprising four oxygen atoms, with no difference in normalized absorption between features B and C (upper spectra in insets a and b). For sample 2, the Co center is surrounded by a ﬁrst shell of three oxygen atoms and a VO, leading to an obvious difference between features B and C for the experimental (0.13) and calculated (0.11) spectra (lower spectra in insets a and b). It is therefore concluded that there are more VO and thus a higher carrier concentration in sample 2 due to lower P O2 . However, sample 1 contains more structural defects than sample 2, resulting in stronger RTFM for sample 1, demonstrating that structural defects are responsible for RTFM in Co-doped ZnO.

RTFM in Co:ZnO ﬁlms by controlling the ﬁlm thickness from 15 to 900 nm (15, 30, 60, 100, 150, 300, 600 and 900 nm). The ﬁlms were deposited on Si(1 0 0) and LiNbO3(1 0 4) substrates by rf magnetron sputtering. The thinnest ﬁlms (15 nm) with greatest lattice strain had the highest saturated magnetic moments of 2.96 and 5.52 mB/Co for Co:ZnO/Si and Co:ZnO/LiNbO3, respectively (Fig. 15) The saturated FM rapidly decreased with increasing ﬁlm thickness;

3.5. Strain-induced ferromagnetism enhancement Strain engineering, which offers the opportunity to enhance the physical properties of a chosen material in thin ﬁlm form by judicious use of growth parameters, represents a new direction for nanoscale material research [192–194]. It has been found that a slight strain-induced break in symmetry has a great inﬂuence on the single spin and the corresponding magnetic behavior of typical DMSs such as Mn:GaAs [195]. Given that growing ﬁlms of different thickness is a simple and effective method to modulate the lattice strain within ﬁlms, Liu et al. [196] studied the effects of strain on

Fig. 14. Co K-edge XANES spectra for Co:ZnO and references materials Co metal, CoO, Co2O3 and Co3O4. The insets show (a) experimental and (b) MS calculated XANES spectra for samples 1 (square) and 2 (circle).

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accumulated state of strain energy due to the large lattice mismatch (40.10%) between the ﬁlm and the Si substrate, and the 15-nm thick Co-doped ZnO ﬁlm exhibits the greatest tensile strain of 6.65% (Fig. 16c). As the ﬁlm thickness increases, the accumulated strain energy progressively relaxes, i.e., the c-axis tensile strain continuously decreases, and is only 0.69% for the 900nm ﬁlm, corresponding to weak FM. The FM enhancement in Codoped ZnO ﬁlms originates from a combination of an increase in Co–O bond length and an increase in the number of defects induced by lattice strain [196]. These results are consistent with the observation that lattice deformation plays a signiﬁcant role in the high-TC FM of DMOs [197]. 3.6. Inﬂuence of post-annealing on magnetization

Fig. 15. Saturated magnetization as a function of thickness for Co:ZnO ﬁlms deposited on Si and LiNbO3 substrates measured at room temperature. The inset shows the hysteresis loop for the 15-nm thick Co:ZnO/LiNbO3 ﬁlm.

the results for the 900-nm ﬁlm was 0.02 mB/Co, which is approximately two orders of magnitude lower than for the 15nm thick ﬁlm. Fig. 16a shows XRD patterns for Co:ZnO/Si ﬁlms of different thickness. Comparison of the (0 0 2) peaks reveals a shift from 33.848 for the 15-nm ﬁlm to 34.458 for unstrained ZnO, corresponding to a change in c-axis lattice constant from 5.207 to 5.553 A˚, indicating that the unit cell in the Co:ZnO ﬁlm is elongated along the c axis. With increasing ﬁlm thickness, the ZnO(0 0 2) peak position shifts towards the value for unstrained ZnO. The c-axis lattice constant and calculated lattice strain are plotted as a function of ﬁlm thickness in Fig. 16b and c. The c-axis lattice constant decreases slightly with increasing thickness. Thus, the initial stage of Co-doped ZnO growth represents a strongly

Annealing has a great effect on the magnetization of TM-doped ZnO, largely because it can change the microstructure and local structure of ﬁlms [38,52,55,57,73,109,140,141,154]. It is commonly observed that annealing in vacuum or a reducing atmosphere results in enhanced FM of the samples due to the introduction of VO [52,55,140,141], with Zni critical for RTFM in TM-doped ZnO [30,57,141,154,198]. For example, a study of annealing effects on the local structure and magnetism of Codoped ZnO ﬁlms under air, Ar, and Ar/H2 atmospheres revealed that ﬁlm annealed in Ar/H2 had a pronounced pre-edge Zn K-edge XANES peak and contained more VO (1 1021 cm3) and stronger RTFM, indicating that RTFM is strongly dependent on VO [55]. Furthermore, a change in Raman spectra (appearance/disappearance of an additional broad vibrational mode) for Co-doped ZnO ﬁlms on alternate annealing in air/high vacuum suggested that VO is crucial in the development of RTFM [100]. Using alternate annealing of Co-doped ZnO ﬁlms in air and Zn vapor, Kittilstved et al. [198] demonstrated quantitative reversible cycling of saturated magnetic moments, and direct correlation between oxidative FM quenching and the diffusion and oxidation of Zni. Similarly, annealing of ﬁlms in different oxidation atmospheres in

Fig. 16. (a) XRD patterns of Co:ZnO/Si ﬁlms of 15–900 nm in thickness. (b) Lattice constant and (c) c-axis lattice strain as a function of ﬁlm thickness.

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the presence or absence of Zn vapor revealed that RTFM is not linked to conductivity and that a small level of Zni defects is important in mediating RTFM in TM:ZnO [154,102]. Similar phenomena have been observed for TM-doped ZnO nanocrystals [20,199,142]. In contrast, a decrease in VO concentration due to high-temperature annealing in oxygen had a large detrimental effect on FM, which tends to disprove a mechanism involving free carriers [200,177]. These experimental annealing results demonstrate that the magnetic behavior observed is directly related to the presence of intrinsic defects, notably VO and Zni. Electrons locally trapped by these defects occupy an orbital overlapping with the d shells of TM neighbors to form a BMP, which accounts for the nature of intrinsic RTFM. The microstructure and magnetic properties of Zn0.92Co0.08O ﬁlms annealed at different temperatures (200, 400 and 600 8C) in vacuum and air were systematically investigated by Liu et al. [141]. The magnetic moment increased after annealing in vacuum and decreased after annealing in air at the same temperature. A typical plot of hysteresis loops for as-grown ﬁlm and ﬁlms annealed in air and vacuum at 400 8C is shown in Fig. 17. The vacuum-annealed ﬁlm shows saturation magnetization (Ms) of approximately 0.45 mB/Co, which is slightly greater than for the as-grown ﬁlm (0.34 mB/Co), whereas the value for the air-annealed ﬁlm (0.21 mB/ Co) is much lower than Ms of the as-grown ﬁlm. The value of Ms increases with annealing temperature up to 400 8C due to the creation of VO during annealing in vacuum [149]. However, annealing at 600 8C greatly decreased the FM of Co-doped ZnO ﬁlms, likely due to an improvement in crystal quality and a corresponding decrease in structural defects in the ﬁlm [141]. This is supported by the result indicating that FM ordering increased with annealing of Zn0.88Co0.12O ﬁlms in vacuum up to 500 8C, whereas a higher annealing temperature resulted in a decrease in magnetization [111]. On annealing in air at high temperature (>700 8C), Mn2+ and 2+ Co ions may segregate from the crystal and form impurity phases or precipitates, which can lead to a decrease in Raman peak intensity [71]. The XANES spectrum of Co:ZnO ﬁlm annealed at 750 8C differs from that of an as-deposited ﬁlm, but closely matches that of Co3O4 (Fig. 18a) in terms of a similar absorption edge, characteristic peaks, and normalized absorption intensity. This can be ascribed to a transition from Co2+ replacement of Zn2+ to the formation of Co3O4-based compounds. This is consistent with a FE-SEM image showing small islands of precipitates (Fig. 18a inset) that are almost randomly distributed and range

Fig. 18. (a) Co K-edge XANES spectra for as-deposited and annealed (750 8C) sample LH, Co metal, CoO and Co3O4 (E0 = 7708.8 eV). The inset shows an FE-SEM image of LH after annealing at 750 8C. (b) Magnetization of as-deposited and annealed (750 8C) LH.

in size from 60 to 120 nm. This transition in local structure (from Co2+ replacement of Zn2+ to Co3O4-based compounds) during 750 8C annealing means that the behavior of Co-doped ZnO ﬁlms changes from RTFM to paramagnetism, as illustrated in Fig. 18b. These phenomena observed after annealing are consistent with the principles described for substrate temperature-dependent magnetization. A combination of a comparatively low temperature (400 8C) and a reducing atmosphere increases defects (VO or Zni) and enhances FM ordering [33,94], whereas high-temperature processing results in a decrease in RTFM. 3.7. Effect of co-doping on magnetic properties

Fig. 17. Hysteresis loops for as-grown and vacuum- and air-annealed (400 8C) Zn0.92Co0.08O thin ﬁlms measured at room temperature.

Controlling the experimental parameters, e.g., decrease in P O2 , can increase the carrier concentration and thus enhances FM [11,48,62,95,103,152,201]. However, some researchers have reported that the creation of free carriers is not sufﬁcient for FM interactions [77,202]. Moreover, several groups observed RTFM in the absence of free carriers [8,30,41,44,45,75,94,108,109,152, 153,186]. To elucidate whether the carrier concentration can enhance RTFM or if there is a correlation between them, co-doping has been applied to generate additional electron/hole doping, as presented in Table 3. The possible effects of co-doping, i.e., simultaneous doping of ZnO with two or more TM ions, are interesting. Typical

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Table 3 List of co-doped ZnO systems, including ﬁlms, bulk and nanocrystals System

Content (at.%)

Methods

Magnetism

Reference

CoAl CoAl CoAl CoAl CoAl MnAl MnSn MnSn CuGa CoLi CoLi CoLi CoN CuN MnN MnN MnN MnP MnCu FeCu CoCu

9/1 10/(0,1.6) 5/(0–0.5) 5/(0–3) 5/(0–2) 2/(0–4) (3,5)/– 3/(0–0.1) 5/1 (5–10)/(0–10) 5/5 5/10 1.5/– 10/– 4/– 5/7 3/– 3/2 10/(0,5) 5/(0–1) 5/(0–10)

PLD PLD rf sputtering Sol–gel PLD PLD Ion implantation PLD PLD Sol–gel Spin-coating Sol–gel rf sputtering dc sputtering ICP-CVD rf sputtering rf sputtering PLD PLD Solid state reaction PLD

0 ! 0.1 mB 0.33 ! 0.11 mB 1.7 ! 3.36 mB 0.30 ! 0.48 mB 0.88 ! 1.6 mB 0.1 ! 4.36 mB 5.5 105 emu 0.1 ! 0.5 mB 1.45 ! 0 mB 50 ! 375 emu/mole 0 ! 0.5 mB 0.009 ! 0.048 mB 0.085 ! 0.129 mB 0.36 ! 0.15 mB 0.3 ! 1.4 mB 0 ! 0.3 mB 0 ! 0.92 mB 0.2 ! 0.07 mB 0.075 ! 0.1 mB 0 ! 0.75 mB 5 ! 15 emu/cm3

[52] [186] [129] [206] [95] [95] [35] [131] [132] [133] [134] [203] [204] [66] [205] [135] [136] [138] [125] [126] [137]

Content is presented as x/y for the main TM element/additional dopant. Magnetism results indicate the change in parameter before ! after co-doping.

electron doping elements in TM-doped ZnO are Al, Sn and Ga [52,95,129,131,132,186,35], whereas N, Li, and Cu are used for hole doping [125,126,129,203–205]. It has been reported that an increase in carrier concentration due to trace amounts of additional dopants can greatly enhance RTFM [95,126,206]. Additional hole doping, using Li or N promotes strong hybridization of Mn 3d/N 2p states and long-range FM coupling, as it brings the d electron/TM closer to the optical values for double exchange [133,134], although some authors reported that co-doping plays a negligible role in enhancing the magnetic moment [125]. Ivill et al. [131] designed a series of experiments on (Mn,Sn)-co-doped ZnO with varying Sn concentration (0–0.1 at.%). They found that the ﬁlms exhibited FM, with an inverse correlation between magnetization and electron density, as a function of the Sn doping. This phenomenon is most consistent with the BMP model in which bound acceptors mediate FM doping. In this case, an increase in electron density decreases the acceptor concentration and thus quenches the FM exchange. Similar conclusions have been reported for (Mn,P)- and (Cu,Ga)-co-doped ZnO [132,138].

FM enhancement arising from an increase in donor defects or electron carrier concentration is ambiguous because these two factors coexist in most doped ZnO systems and are hard to distinguish. It is hence of vital importance to design an experiment that can separate the two factors [129,204]. Zn0.95xCo0.05AlxO (x = 0, 0.001, 0.002, 0.004, 0.005) thin ﬁlms were deposited on Si(1 0 0) substrates by rf magnetron co-sputtering [129]. Fig. 19 shows XRD spectra of Zn0.95xCo0.05AlxO (x = 0, 0.002) ﬁlms. The peak corresponding to ZnO(0 0 2) changes little except for a decrease in intensity after additional Al doping in the Zn0.95Co0.05O thin ﬁlm, implying a decrease in crystalline quality and an increase in structural defects on inclusion of a trace amount of Al. This was conﬁrmed by optical transmittance spectra of pure ZnO and Zn0.95xCo0.05AlxO (0 x 0.005) [129]. Fig. 20 shows curves of the magnetization versus magnetic ﬁeld (M–H) for Zn0.95xCo0.05AlxO (0 x 0.005) thin ﬁlms. The Al concentration dependence of Ms and ne is shown in the inset. It is evident that a trace amount of Al can enhance the magnetic moment per Co ion of Co:ZnO ﬁlms. The increase in Ms with

Fig. 19. Comparison of XRD patterns for pure ZnO, Zn0.95Co0.05O and Zn0.95xCo0.05AlxO (x = 0, 0.002) thin ﬁlms on Si(1 0 0) substrates.

Fig. 20. Electron carrier concentration and saturation magnetization as a function of Al concentration. The inset shows magnetization hysteresis curves of Zn0.95xCo0.05AlxO thin ﬁlms measured at RT.

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additional Al doping, combined with the increase in ne, reveals that the free carrier concentration is important, as expected for DMOs. However, it should be noted that ne (1.597 1020 cm3) is greater and Ms (2.34 mB/Co) is smaller for the Zn0.95xCo0.05AlxO (x = 0.004) ﬁlm than for the Zn0.95xCo0.05AlxO (x = 0.001) ﬁlm (ne = 4.597 1019 cm3, Ms = 2.71 mB/Co), indicating that electrons are not the real mediators of FM exchange. It is clear that additional Al doping leads to a decrease in both the crystalline quality and the number of crystal grains in Co:ZnO ﬁlms, indicating that a large number of structural defects are introduced, mainly located in grain boundary areas in the samples. In addition, the concentration of donor defects (notably VO and Zni) increases to a certain extent [58,109], accounting for the FM enhancement. It is noted that the ne dramatically increases up to x = 0.002 and then remains relatively constant with further increases in Al concentration, implying that Al is not the only candidate responsible for introducing electron carriers, as shown in Fig. 21a (red circle). It is well known that ZnO is a natural ntype semiconductor owing to the existence of native defects such as VO and Zni. An increase in structural defects could enhance the concentrations of point defects such as VO and Zni in the ﬁlms. Such an increase in point defects, as well as the native defects, is regarded as another primary donor source of electron carriers. Since free electrons can be introduced into Co:ZnO ﬁlms not only by additional Al doping, but also by an increase in donor defects, the ratio of the carrier concentration to the Al density (ne/nAl) is an indirect measure of the number of defects in a ﬁlm, with a larger ne/ nAl value indication a greater number of donor defects. For Zn0.95xCo0.05AlxO (x = 0.001), ne/nAl = 1.09 is greater than the

value for x = 0.004 (ne/nAl = 0.95), implying a greater number of donor defects in the former ﬁlm and corresponding to the stronger Ms of 2.71 mB/Co, indicating that donor defects most likely play a critical role in governing FM ordering. A systematic study by Mokhtari et al. [207] revealed that additional doping could ionize other defects, which played an important role in FM enhancement. Fig. 21 presents the change in carrier density as a function of the Al concentration in (a) Co- and (b) Mn-doped ﬁlms. Data obtained for ﬁlms prepared by PLD at 0.05 mtorr and by magnetron sputtering are denoted by black and red circles, respectively. The circles around the data points indicate the magnitude of the magnetic moment measured at RT. The solid lines correspond to addition of one carrier per Al atom. It is evident that all data lie signiﬁcantly above this line, indicating that the presence of Al has ionized other defects, perhaps some Co2+ and Mn2+, which provides carriers in the ﬁlms. Acceptor co-doping using Li and N has also been extensively studied to determine the origin of RTFM [134,204]. Fig. 22 shows unpolarized Raman scattering spectra measured at RT for (Co,N):ZnO ﬁlms. Vibrational modes are evident at 273, 470, 639, 691 and 854 cm1, in addition to the host phonons of ZnO, indicating the substitutional behavior of N3 for O2 and an increase in donor defects for (Co,N)-co-doped ZnO ﬁlms [204]. The higher intensity of the vibrational mode at 577 cm1 indicates a higher concentration of donor defects in (Co, N):ZnO ﬁlms than in Co:ZnO ﬁlms. Acceptor doping with nitrogen compensates the electron carriers and hence decreases ne [135]. Magnetization measurements reveal that co-doping increases Ms for (Co,N)-codoped ZnO (0.129 mB/Co) compared to Co:ZnO (0.085 mB/Co)

Fig. 21. Carrier density of (a) Zn0.95xCo0.05AlxO and (b) Zn0.98xMn0.02AlxO ﬁlms. The solid line corresponds to addition of one carrier per Al atoms. Films were grown by PLD at 0.05 mtorr (black) [207] and magnetron sputtering (red), respectively [129]. The thickness of the circles indicates the magnetic moment in Bohr magnetons. The circles lie signiﬁcantly above the line, indicating that the presence of Al has ionized other defects and also possibly some of the Co2+. (For interpretation of the references to color in this ﬁgure legend, the reader is referred to the web version of the article.)

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relate this difference directly to the magnetization of ZnO doped with different TM dopants. The Lande´ g-factor can be expressed as [175]: g¼

Fig. 22. Raman spectra of ZnO, Co:ZnO and (Co,N):ZnO ﬁlms. ZnO and silicon substrate modes are marked by (*) and (#), respectively.

[204]. Combined with FM enhancement, donor defects other than electron carriers are responsible for RTFM in TM-doped ZnO ﬁlms. 3.8. Comparison of ZnO doped with Sc, Ti, V, Cr, Mn, Fe, Co, Ni and Cu More than 800 reports on TM-doped ZnO DMOs have been published from 2000 to date. A large number of these focus on Co[8,11,23,44–55,58–60,97,103,111,126,141,152,154,164,165,198] and Mn-doped ZnO ﬁlms [31–37,39,40,73,98,117,140,186]. There have also been numerous experiments on ZnO ﬁlms doped with other TM elements, including Sc [23,24], Ti [23,25], V [26–28], Cr [11,23,29,30,208], Fe [41–43,124], Ni [62,63] and Cu [64–67,132]. Some of the magnetic moments measured at room temperature reported for 5 at.% TM-doped ZnO ﬁlms are summarized in Fig. 23. ZnO ﬁlms doped with 5 at.% Co commonly exhibit much stronger RTFM than ZnO doped with other TM elements. This is consistent with results obtained by Ueda et al. [11] and Venkatesan et al. [23] for ZnO doped with a series of TM elements; such studies can exclude the inﬂuence of the preparation parameters on magnetic ordering. It is evident that ZnO ﬁlms doped with different TMs exhibit RTFM with different Ms values, as well as very small magnetic moments indicating negligible FM. It is tempting to

Fig. 23. Magnetic moments reported for (5 at.%) TM:ZnO at room temperature.

3 1 LðL þ 1Þ SðS þ 1Þ ; 2 2 JðJ þ 1Þ

(2)

where J is the total angular momentum, including the spin (S) and orbit (L) angular momentum. The Lande´ g-factor and J can be used to evaluate the effective magnetic moment of TM ions in TM:ZnO ﬁlms. For Co2+ in Co:ZnO, the Lande´ g-factor is 4/3, the effective magnetic moment calculated using Eq. (1) is 6.6 mB/Co and the net spin moment calculated from the spin-only formula mmax ¼ g mB S is 3 mB. Similarly, Mn2+ has both an effective and net spin magnetic moment of 5 mB. These different ideal magnetic values for TM ions are most likely one of the reason for different moments of TM:ZnO. On the other hand, it is generally considered that Co is more soluble than other TM atoms in ZnO ﬁlms [11,44,105]. In terms of concentration-dependent magnetization, the largest moments have been observed for ZnO doped with different TM concentrations, i.e., 3–5 at.% for Co [44,58,103], 2 at.% for Mn [33,95], and 1 at.% for Cr, Ni and Cu [62,64,66], and ZnO ﬁlms with other dopant contents tend to show much lower magnetic moments. Therefore, besides ideal moment discrepancies for these TM ions in ZnO ﬁlms, different solubility and different maximum dopant concentrationdependent magnetization would also account for different magnetic moments in 5 at.% TM-doped ZnO ﬁlms.

4. Origin of ferromagnetism in TM:ZnO As discussed in Section 3, almost all of the publications that reported magnetization of TM:ZnO referred to the origin of magnetic ordering. This is very important not only for understanding the nature of TM:ZnO, but also for applications in useful spintronics. Many points and mechanisms have been proposed for intrinsic magnetic ordering. However, it is hard to experimentally characterize and quantitatively analyze the nature of RTFM because so many factors affect the magnetic behavior of TMdoped ZnO, including the microstructure, local structure, electronic structure, coupling between localized magnetization, and trace amounts of TM dopant atoms in complicated oxide ﬁlms. On the other hand, FM values reported seem to be rather weak, often many orders of magnitude smaller than the expected spin-only saturation moments. The samples that exhibit FM often do so at levels low enough so that no technique other than magnetization itself can reveal whether the phenomenon is intrinsic or arises from impurities [209]. Magnetization is measured using a SQUID magnetometer with a high sensitivity of 108 emu; trace FM contamination of the substrate can lead to ﬁlms with magnetization signals, which means that measurements may be ambiguous. Detection of the local structure and electronic structure of TM elements in TM:ZnO ﬁlms is generally limited by the sensitivity and efﬁciency of the equipment used [85,139], which can lead to false conclusions. Reports that cast doubt on the very existence of any type of long-range magnetic order in clean samples of TMdoped ZnO at low doping concentrations are also appearing with increasing frequency [209]. There is ongoing debate about these systems. Therefore, from the broad spread of experimental data it is difﬁcult to draw general conclusions about the nature of magnetic interactions in TM-doped ZnO. The confusing situation clearly calls for a critical and even subjective position to be taken on this topic. So far, the most popular mechanisms suggested for magnetic ordering in TM-doped ZnO are carrier-mediated exchange (electron and hole carriers), and the BMP model with

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respect to defects, including VO and Zni, and Zn vacancies, which are described in detail in this section. 4.1. Computational work Computational work based on ab initio calculations can provide some initial predictions and explanations of the electronic structure and the nature of FM ordering in TM-doped ZnO. To simulate realistic dopant concentrations, calculation are required for supercells containing a large number of atoms, in which the doping concentrations are usually higher than 5 at.% [84,210]. Hence, computational results for TM:ZnO have been produced using a range of electronic structure implementations, generally using a density functional theory (DFT) approach based on pseudopotentials with localized atomic-orbital basis sets, in which the total energy of a many-electron system is described as a function of the electron density [19,84,210,211]. This method combines accuracy with low computational cost, particularly compared to other approaches such as plane-wave pseudopotential or all-electron methods [84]. The most widely used approximation is the local density approximation (LDA) exchange correlation function [10]. Relaxation of the constraint of an equal occupation of spin-up and spin-down states and minimization of the total energy with respect to spin occupancy results in the local spin-density approximation (LSDA) [19,133]. It is known that LDA or LDSA calculations fail in predicting the insulating behavior of many TM oxides, but produce the metallic ground state. Incorporation of the effective on-site Coulomb interaction, characterized by Hubbard U (LDA + U), between TM 3d electrons into LDA or LDSA calculations greatly improves the interaction for TM oxides [212,213]. On the other hand, the highly precise projected augmented wave (PAW) method within the generalized gradient approximation (GGA) or GGA + U scheme, which includes an explicit dependence of exchange-correlation energy on the electron density gradient, has also been used in calculations [33,10,214,215]. In 2000, Dietl et al. [9] theoretically predicted that p-type ZnO doped with 5 at.% Mn and similar TM-doped wide-bandgap semiconductors would show RTFM, which focused much attention on TM:ZnO research. The prediction used a mean ﬁeld approach, assuming hole-mediated FM interactions between Mn local moments. TC was then determined from competition between this long-range carrier-mediated FM coupling and short-range Mn–Mn anti-FM exchange interactions. Following this prediction, Sato and Katayama-Yoshida [10] used the Korringa–Kohn– Rostoker (KKR) Green function method based on LDA of density functional theory to calculate the properties of ZnO doped with TMs such as V, Cr, Mn, Fe, Co, and Ni. Additional holes were introduced by replacing O with N, and electrons were added by replacing some of the Zn atoms with Ga. The energy difference between FM and anti-FM ordering was taken as the upper TC limit. This energy difference (EFM EAFM) was in the region of 0.2 eV, which corresponds to a high TC. An FM ground state would be favored at a doping level of 25 at.% V, Cr, Fe, Co, or Ni in ZnO without additional dopants to provide carriers, whereas Mn-doped ZnO was anti-FM. Spaldin [84] identiﬁed potential FM ground states within (Zn,Co)O or (Zn,Mn)O using the DFT approach with an LSDA exchange-correlation function. This is achieved for 32-atom supercells containing two dopant ions in different positional arrangements: a ‘‘close’’ conﬁguration in which TM atoms were separated by a single O ion, and a ‘‘separated’’ conﬁguration in which they were connected by an –O–Zn–O– bond. The calculations show that the energy difference between the anti-FM and FM spin orders was only a few meV, producing paramagnetic behavior

in the absence of free carriers. Since the total energy calculated is very close for FM and anti-FM orders [84,210], actual TM-doped ZnO systems observed in experiments may remain in different metastable states that exhibit FM, anti-FM, and spin-glass behavior, depending on subtle differences in experimental conditions, even for similar ﬁlms. More recently, Ye et al. [211] found that the FM state was 43 meV lower than the anti-FM state, and thus 12.5 at.% Cu-doped ZnO was predicted to be in the ground state with a net magnetic moment of 1 mB/Cu and an estimated TC of 380 K. In contrast, it was found that the close case for Mn:ZnO favors an anti-FM state, whereas the separated conﬁguration tends to be an FM state, stable by only 0.01 eV [216]. Furthermore, different geometries with FM, ferrimagnetic and anti-FM conﬁgurations were found to be nearly energetically degenerate for both bulk V:ZnO and subsurface layers of the ﬁlms, where V atoms couple ferromagnetically when they occupy surface sites of the ﬁlms [210]. On the other hand, Lee and Chang [188] calculated the energy difference between FM and anti-FM alignments for two Co atoms substituted in ZnO using the GGA approximation in supercells of 2 2 1, 2 1 2 and 4 1 1 to allow different Co–Co distances. The total energy difference between the FM and anti-FM states was <3 meV per Co atom, indicating that competition between FM and anti-FM coupling exists in Co-doped ZnO and that spontaneous magnetization is not possible in intrinsic Co-doped ZnO, which theoretically supports FM at 300 K in some Co-doped ZnO ﬁlms, but the reproducibility was less than 10% [11]. Meanwhile, high concentrations of both Co ions and electron carriers are needed to achieve FM [188,217]. These diverse magnetic behaviors predicted by calculations could be one of the main reasons for the conﬂicting experimental results reported. However, for both Co- and Mn-doped ZnO, the FM state was stabilized by the introduction of holes, such as the creation of Zn vacancies in (Co,Cu)-doped ZnO. Following this prediction, the effects of Zn vacancies on FM were experimentally and theoretically investigated [34,218,219]. Some groups demonstrated that Zn vacancies promote the appearance of RTFM by ﬁrst-principle calculations [34,218,219]. Comparison of experimental and calculated O K-edge XANES revealed that Zn vacancies indeed exist in Mn-doped ZnO ﬁlms and account for the RTFM. Likewise, the presence of Fe3+ is a signature of hole doping induced by Zn vacancies. A deep acceptor trap due to Zn vacancies with thermal activation energy of 0.815 eV has also been detected by deep-level transient spectroscopy [49]. Karmakar et al. [218] revealed experimentally and theoretically that hole doping is crucial in promoting FM in Fe-doped ZnO systems, and the presence of Fe3+ is a signature of hole doping induced by Zn vacancies. Furthermore, Yan et al. [34] recently demonstrated that Zn vacancies favor RTFM in Mn-doped ZnO ﬁlms using simulations of O K-edge XANES and ﬁrst-principle calculations; similar calculations and results were reported by Ius¸an et al. [219]. To elucidate the effects of co-doping and defects on magnetic ordering in TM-doped ZnO, Gopal and Spaldin [213] systematically investigated four cases: (a) single-TM (Cr, Mn, Fe, Co, Ni and Cu) substitution at Zn sites; (b) substitutional magnetic TM ions combined with additional Cu and Li dopants; (c) substitutional magnetic TM ions combined with VO; and (d) pairs of magnetic ions (Co and Fe, Co and Mn). The results indicate that FM ordering of TM ions is not induced by addition of substitutional TM impurities or VO. Incorporation of interstitial or substitutional Li is favorable for FM, as are Zn vacancies, which is consistent with experimental results for (Co,Li)-co-doped ZnO [133,134]. We cannot draw deﬁnite conclusions regarding intrinsic RTFM from the computational work carried out. Electron carriers [188,220] and hole doping [49,84,133,134,213,221] seem to be

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necessary for FM ordering. Gopal and Spaldin [213] proposed that the calculated magnetic behavior is strongly dependent both on the computational details (with FM disfavored by improved convergence) and on the choice of exchange correlation function (with FM disfavored by the more appropriate LSDA + U method). This observation explains the large spread of computational results in the literature. Fortunately, there is some common ground: most publications agree that intrinsic FM does not occur in TM-doped ZnO at reasonable doping concentrations, but in many cases the addition of carriers or point defects stabilizes the FM state [222,223]. These additional carriers or point defects are in fact byproducts of structural defects in TM-doped ZnO, which have frequently been demonstrated by recent experiments and appear with increasing frequency. 4.2. Experimental work Fueled by theoretical predictions, claims of FM at and above room temperature in TM-doped ZnO have abounded. Unfortunately, neither the true nature of these materials nor the physical causes of the magnetism have been adequately determined. It is known that it is difﬁcult to introduce factors that may inﬂuence RTFM, such as defects and carriers, into calculations. Moreover, it is not easy to realize low doping concentrations (<2 at.%) in TMdoped ZnO supercells because of the limitation of the calculation efﬁciency of ﬁrst-principle calculations. So far, hundreds of studies have experimentally investigated intrinsic RTFM by focusing on correlations among the magnetization, microstructure, local structure and electric properties. Besides the extrinsic origin of RTFM arising from FM secondary phases, several mechanisms have been proposed for intrinsic magnetic ordering, including carriermediated exchange (electron and hole carriers) and a BMP model in terms of structural defects (VO and Zni), which are described in details in the following subsections. 4.2.1. Secondary phases Before elucidating the mechanisms of intrinsic FM, it must ﬁrst be conﬁrmed whether the magnetic hysteresis loops measured are intrinsic or extrinsic. Magnetic contamination [86], measurement errors [139] and secondary phases [31,110] are all possibilities for FM signals. Low levels of FM contamination in the substrate and ﬁlm can yield magnetization signals detectable by magnetometers, which have a high sensitivity of 108 emu. The mass of a 100-nm thick ﬁlm on a substrate of 5 mm 5 mm 0.5 mm is approximately 10–15 mg, whereas the substrate mass is 50 mg. Approximately 200 ng of FM contaminant, equivalent to a speck of radius of 20 mm, would be sufﬁcient to produce the magnetic moment observed for a typical FM thin-ﬁlm specimen [139]. A common feature is that all samples tend to exhibit phase segregation at high dopant concentrations and rather high annealing temperatures (>700 8C in general). These secondary phases are often likely to be responsible for the FM observed, which are key points for the study of DMOs. A key requirement in elucidating RFTM in TM-doped ZnO is to judge whether the magnetism originates from substitutional dopants on Zn sites or from the formation of a secondary FM phase. A number of FM or ferrimagnetic Co-based phases have been detected. For example, FM Mn2xZnxO3d has been proposed as the origin of high-temperature FM in (Zn,Mn)O [31]. Both metals and their oxides have been identiﬁed as being responsible for FM behavior in TM-doped ZnO ﬁlms, such as Co [18,224,225], Mn [122], CoFe [70], CoO [56,163], Co3O4 [56,109,111,163], Fe3O4 [43], and CuO [65]. These types of TM-doped ZnO ﬁlms are not genuine DMOs. However, these non-magnetic semiconductors with embedded secondary FM phases may exhibit an anomalous Hall

21

effect (AHE) [226], which was previously considered an indicator of whether a doped sample was a DMO or not. AHE was observed in FM ZnO and in non-FM Cu-doped ZnO ﬁlms, indicating that AHE does not uniquely demonstrate FM behavior [227]. Therefore, checking for basic experimental reproducibility, in addition to carrying out careful high-sensitivity characterizations, is highly warranted. 4.2.2. Carrier exchange interactions Since RTFM arising from carrier exchange interactions has been predicted for several TM-doped ZnO systems [9,10,228], a great deal of attention has been focused on doped ZnO systems as potential DMOs to provide efﬁcient injection of spin-polarized carriers for spintronics devices. Meta-analysis reveals that early publications (2001–2004) generally favored carrier exchange interactions [11,81,48], which are characterized by strong coupling between localized d electrons of TM ions and the extended sp carriers of ZnO. The term carrier-mediated exchange refers to interactions between localized magnetic moments that are mediated by free carriers in the system, which can be divided into three cases: the Rudermann–Kittel–Kasuya–Yosida (RKKY) interaction; Zener carrier-mediated exchange; and Zener double exchange. RKKY interaction formally describes the magnetic exchange between a single localized magnetic moment and a free electron gas. Zener carrier-mediated exchange proposes that carriers can mediate FM interaction between local moments in the ﬁlms and both local magnetic moments and itinerant carriers [9]. FM ordering is driven by a decrease in carrier energy due to redistribution between spin sub-bands split by the exchange interaction. Finally, the Zener double-exchange model refers to indirect coupling mediated by oxygen atoms between neighboring FM ions in different states. The kinetic energy of the system decreases if the magnetic moments align in parallel, since parallel alignment allows electron transfer from ions in the low state to those in the higher state [19]. In TM-doped ZnO, the addition of Al as a dopant, or a decrease in PO2 during ﬁlm growth, seems to increase the carrier concentration and enhance FM [46,206]. This suggests the presence of carriermediated exchange. It was found that two distinct mechanisms could give rise to FM in TM-doped ZnO prepared by PLD: magnetic polarons and carrier-mediated exchange, giving rise to DMI and DMO behavior, respectively [152]. Recently, some groups reported that additional doping cannot signiﬁcantly enhance FM ordering [109,112,125,129,131,138]. For ZnO, it is known that annealing in a reducing atmosphere results in the creation of VO, with the formation of one charge-compensating electron for every VO, thus leading to an increase in conductivity. That is, an increase in carriers generally accompanies the production of defects, leading to misguided and shortsighted conclusion of a carrier-mediated mechanism. By separating these two factors, several reports have determined that FM enhancement is most likely due a decrease in crystalline quality and an increase in structural defects with a trace amount of additional doping (e.g., Al and Sn) [129,130,138]. 4.2.3. Defect-based BMP Note that conducting or semiconducting TM-doped ZnO ﬁlms with a perfect single-crystal structure show negligible FM, which is considered to represent the absence of structural defects [27,77,79,153,158,184]. There is a dramatic increase in resistivity for Co:ZnO ﬁlms on Cu doping, but this change in resistivity does not affect the magnetic moment, suggesting that a carriermediated mechanism is not a feasible explanation for FM in these ﬁlms [137,139]. Cu-doped ZnO ﬁlms showed RTFM with 1.45 mB/ Cu, whereas FM completely vanished when additional carriers were introduced into the ﬁlms [132]. Coey et al. [8] observed RTFM

22

F. Pan et al. / Materials Science and Engineering R 62 (2008) 1–35

in insulating TM:ZnO and elucidated a BMP mechanism. Furthermore, an increasing number of experimental results demonstrate that carriers are not necessary for RTFM, and FM interaction can be unambiguously observed in ﬁlms without free carriers but in a dielectric state [8,30,41,44,45,75,94,108,109,152,153,186]. This insulating nature is signiﬁcant, in that a magnetic coupling interaction other than carrier-mediated exchange is apparently operative. In fact, many speciﬁcally designed experiments appearing with increasing frequency reveal that defects such as VO and Zni account for RTFM in TM:ZnO [44,58,94,167]. For example, Co:ZnO ﬁlms annealed in Ar/H2 with a pronounced Zn K-edge XANES pre-edge peak showed an increase in VO (1 1021 cm3) that was responsible for much stronger RTFM [55]. Annealing of Zn0.92Co0.08O ﬁlms in vacuum and air also demonstrated that VO and Zni are responsible for the RTFM of the ﬁlms [141]. More interestingly, quantitative reversible cycling of magnetic moments was achieved by annealing TM:ZnO ﬁlms in alternating air and Zn vapor (vacuum), demonstrating a direct correlation between FM and Zni (VO) [102,154,198]. By annealing ﬁlms in different oxidation atmospheres and in the presence or absence of Zn vapor, it has been demonstrated that RTFM is not linked to conductivity and that a small level of Zni defects and consequent BMP plays a crucial role in mediating RTFM in TM:ZnO. Similarly, post-annealing leads to a decrease in the c lattice parameter by affecting the Zni concentration, corresponding to a decrease in RTFM [57]. Using XAS and XMCD, Matsumura et al. [197] obtained direct evidence that the absence of anomalous dispersion of the atomic scattering factor for Co is probably caused by signiﬁcant lattice deformation of the local structure around Co due to VO, accounting for RTFM in Co:ZnO, whereas Co ions indeed substitute in Zn sites in paramagnetic Co:ZnO thin ﬁlms [197]. This observation also demonstrates that VO is necessary for the appearance of RTFM. Elimination of defects, as well as ﬁlling up VO, would degrade the FM ordering [29,141,200]. For example, Cr:ZnO ﬁlms fabricated under various conditions revealed that elimination of structural defects and ﬁlling up of VO might lead to a certain level of FM degradation [29]. Likewise, annealing of Cu-doped ZnO ﬁlms in oxygen led to a decrease in VO and had a large detrimental effect on FM [177]. The large amount of edge dislocations observed would enhance the concentration of VO and Zni point defects, which contribute to FM ordering [58,94]. Therefore, these results imply

that structural defects have a signiﬁcant inﬂuence on magnetism in these systems. Under conditions in which acceptor dopants are activated, leading to a decrease in free-electron density, magnetization is enhanced. The result is consistent with hole-mediated FM in Mndoped ZnO, in which bound acceptors mediate the FM ordering. Increasing the electron density decreases the acceptor concentration, thus quenching the FM exchange [131,138]. In addition, ZnO and similar TM oxides show paramagnetism at RT, whereas it has been claimed that RTFM can be obtained in certain d0 TM oxides without a dopant [229–231]. It should be pointed out that pure HfO2 powder develops a weak magnetic moment on heating in vacuum, which is eliminated on annealing in oxygen [229]. Lattice defects are proposed to be responsible for the FM properties in these semiconducting and insulating ﬁlms without doping. It is worth pointing out that the absence of free spin-polarized carriers would restrict the application of these materials in RT spintronics [139,154,232]. Fig. 24a shows the lattice structure of TM:ZnO, in which TM ions are incorporated into the wurtzite lattice at Zn2+ ion sites. The highly non-equilibrium process of ﬁlm deposition makes it possible that structural defects in the doped ﬁlm are located throughout the lattice at arbitrary distances with respect to Co sites, which can play an important role in the system, as illustrated in Fig. 24b. The formation of one charge-compensating electron for every VO produces a polaron. An electron associated with a particular defect is conﬁned in a hydrogenic orbital of radius rH = er(m/m*)a0 and g = er(m/m*), where er is the dielectric constant, m the electron mass, m* the effective mass of the donor electrons, and a0 is the Bohr radius (0.53 A˚). For Co:ZnO, m/m* = 3.57 and er is in the range 4–37 [8,44,107,141], which is dependent on the fabrication method and parameters. For example, er = 10 corresponds to large rH and g, i.e., rH = 18.9 A˚ and g = 35.7. Likewise, the polaron size is estimated to be approximately 7.6 and 7.8 A˚ in Codoped ZnO [8,100]. In contrast, the disappearance of RTFM after hydrogen implantation can be ascribed to the enhanced effective mass and subsequent decrease in g, which also supports the above equation [38]. It should be noted that if the donor orbital extends over a sufﬁcient number of cations, there is nothing to be gained by any further increase in er [8]. Interestingly, an electron trapped in a defect is demonstrated by C–V measurement, which exhibits hysteresis with a large memory

Fig. 24. (a) Lattice structure of TM:ZnO. (b) Three-polaron subsystem representing supercoupling between polarons.

F. Pan et al. / Materials Science and Engineering R 62 (2008) 1–35

window of 0.6 V [94], effectively providing evidence of a BMP mechanism: the charge-compensating electron in each defect forms a polaron. The formula for Co-doped ZnO can be written as (Zn0.96Co0.04)(O&d), where & represents a donor (defect) and dp is the defect percolation threshold [8]. For (4 at.%) Co:ZnO ﬁlms with er = 22 and me/m 0.3, dp can be calculated from g3dp = 4.3 [8]. Using the value of dp calculated, the critical concentration of oxygen vacancies (n&; defects) can be obtained from the relation d = n&/no, where no is the oxygen density of no = 1.4 1022 for a close-packed oxygen lattice in ZnO. In the present case, the defect percolation threshold n& is thus approximately 1.6 1017 [108]. FM ordering is due to superexchange between complexes in DMI, i.e., VO and magnetic impurities, which are stabilized by charge transfer from vacancies to impurities. The percolation threshold for magnetic ordering is determined by the radius of vacancy levels, but the exchange mechanism does not require free carriers [228]. If there is a sufﬁciently large orbital radius, overlap between a hydrogenic electron and the cations within its orbit leads to FM supercoupling between them, as shown in Fig. 24b. Interaction between the hydrogenic electron and the cations is represented by a Heisenberg exchange Hamiltonian [233]: Hˆ i j ¼

X

J i j Sˆi sˆ j ;

(3)

ij

where S is the spin of Co2+ and s is the donor electron spin. The donors tend to form BMPs, coupling Co2+ within their orbits. The Hamiltonian of a two-polaron subsystem is given by Eq. (3), where the donor electron spin index j takes only two values, j1 and j2, corresponding to the two polarons under consideration. The hydrogenic orbital tends to spread out sufﬁciently to overlap with a large number of BMPs, but some isolated polarons cannot be covered and thus cannot achieve macro-FM ordering. It is therefore concluded that the origin of FM in DMOs does not necessarily involve free charge carriers and that RKKY-type indirect exchange interactions are not necessary for magnetic ordering. In contrast, TM ions form deep levels in the bandgap and a certain concentration of defects in the ﬁlms may play a more active role in mediating FM exchange than in traditional RKKY systems, as increasingly supported by related publications. 4.3. Effect of bandgap on the Curie temperature It is generally reported that TM-doped ZnO ﬁlms show FM ordering at or above RT (Table 1); these materials have been considered typical DMOs with high TC, and should be promising for use in spintronics. In fact, using a mean ﬁeld approach, calculations predicted TC for oxides and semiconductors, assuming holemediated FM interactions between Mn local moments. TC was then determined from competition between this long-range carriermediated FM coupling and short-range TM–TM anti-FM exchange interactions. The major breakthrough in such calculations is magnetic ordering temperatures in excess of RT for Mn-doped GaN and ZnO (Fig. 25, plotted according to Ref. [9]). Given that TC is one of the most important parameters for DMOs, many researchers have tried to enhance it and to elucidate the factors that inﬂuence it. This has led to TC enhancement from much lower than RT [158] to 790 K [44] for TM:ZnO and to 700 K for TM-doped TiO2 [190,234], and to 110 K [3] to 173 K [235] for Mn:GaAs [236]. The highest TC values for typical DMOs with intrinsic magnetization are listed in Table 4, including Mn:InAs [237], Mn:Ge [238,6], Mn:GaAs [1,2], Cr:ZnTe [239], (Mn,C):GaP [240], Fe/Co/Ni:TiO2 [234], Cr:TiO2 [190], Mn:ZnO [33], Co:ZnO [44], Cu:ZnO [65], Co: SnO2 [172], Fe:SnO2 [41], Co:CeO2 [241], Mn:GaN [242], Cr:In2O3 [243], Co:LiNbO3 [120], Co:HfO2 [244], Cr:AlN [245], in order of

23

Fig. 25. Theoretical Curie temperature of DMOs [9].

increasing bandgap (Eg). It should be noted that high TC values attributed to secondary phases are excluded from the table, such as TC = 980 K for Mn-doped ZnO and 1180 K for Co:TiO2 due to the formation of Co metal [226] and Mn2xZnxO3d [31], respectively. In contrast to the confusion in the literature regarding magnetic moments of DMOs, according to theoretical studies based on the BMP mechanism [8,236], the TC of DMOs is commonly quantitatively given by [8]: T C ¼ ½ðS þ 1Þs2 xd=3

1=2

3

Ji j f O ðrceff =r O Þ =kB ;

(4)

where S is the localized core spin, s the electron donor spin, x the doping concentration, d the donor concentration, Jij the s–d exchange parameter (bandgap Eg), fO the oxygen packing fraction for the oxide, rceff the effective cation radius, and kB is Boltzmann’s constant. Theoretical predictions according to Eq. (4) reveal that Eg and the doping and donor concentrations are important factors for TC of a given system. For example, TC can be estimated from the following parameters for Co:ZnO: S = 3/2, s = 1/2, x = 0.04, Jij = 3.6 eV [44], d = 0.01, rceff ¼ 0:20 nm ¼ 0:20 nm, and r0 = 0.14 nm [8]. The TC given Table 4 Curie temperatures (TC) reported for DMOs Band gap (eV)

System

TC (K)

InAs

0.36

Mn:InAs

30

[237]

Ge

0.75

Mn:Ge Mn:Ge

116 150

[238] [6]

GaAs

1.55

Mn:GaAs Mn:GaAs

150 173

[1] [2]

ZnTe GaP

2.38 2.87

Cr:ZnTe (Mn,C):GaP

300 300

[239] [240]

TiO2

3.0 (rutile) 3.2 (anatase)

Fe/Co/Ni:TiO2 Cr:TiO2

700 700

[234] [190]

ZnO

3.37

Mn:ZnO Co:ZnO Cu:ZnO

425 790 390

[33] [44] [65]

SnO2

3.4

Co: SnO2 Fe:SnO2

650 610

[172] [41]

CeO2 GaN In2O3 LiNbO3 HfO2 AlN

3.4 3.5 3.75 3.8 4.5 6.2

Co: CeO2 Mn:GaN Cr–In2O3 Co–LiNbO3 Co:HfO2 Cr:AlN

740–875 940 850–930 710 800 900

[241] [242] [243] [120] [244] [245]

Reference

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F. Pan et al. / Materials Science and Engineering R 62 (2008) 1–35

Fig. 26. Curie temperature reported for DMOs as a function of bandgap. The correlation between them is described by a simple linear ﬁt: TC = 226.3Eg 107.4.

by Eq. (4) can then be estimated as 822 K, which is close to the experimental result of 790 K for (4 at.%) Co:ZnO [44]. TC values reported for DMOs as a function of semiconductor Eg are summarized in Fig. 26. Interestingly, the reported data reveal that TC is proportional to the semiconductor Eg. This correlation can be described by a simple linear ﬁt: TC = 226.3Eg 107.4, demonstrating that Eg has a signiﬁcant inﬂuence on the TC of TM-doped ZnO. This relation provides an insight into the TC values measured for DMOs ﬁlms and opens the possibility of increasing TC by modifying the Eg of DMOs, especially for Mn-doped GaAs, which has potential applications in spintronics owing to its strong spin–orbit coupling and high spin polarization [2]. 5. TM-doped ZnO-based prototype spintronics and multifunctional materials Spintronics is a topic that has attracted much interest in recent years. Within the context of spintronics, the spin of electrons, and not just their electrical charge, is controlled during the operation of information circuits. The transport and manipulation of spinpolarized electrons or holes in semiconductors offer huge potential for novel devices that combine non-volatile information storage with high processing speed at low power, which may even be useful for quantum computation [7,246]. It was the discovery of giant magnetoresistance (GMR) and tunnel magnetoresistance (TMR) that initiated spintronics research and guided a ﬁrstgeneration device in the form of spin valves based on FM multilayers, such as Co/Ru [247], which are now used in the read heads of most hard-disk storage devices. Second-generation spintronics integrating semiconductor materials and magnetic elements involve new ﬂexible devices such as spin-FET and spin logic devices, which not only improve the capabilities of electronic transistors, but also have new functionalities [248]. Successful realization of most spintronics applications depends critically on the ability to create spin-polarized charge carriers in a conventional semiconductor in a device structure [22]. To successfully incorporate spin into existing semiconductor technology, technical issues such as efﬁcient injection, transport, control and manipulation, and detection of spin polarization and spin-polarized currents have to be solved [7]. A major issue that needs to be resolved is the transfer of spinpolarized carriers from a magnetic contact into a non-magnetic semiconductor. Highly efﬁcient electrical spin injection was realized for the ﬁrst time in 1999 [249]. Spin injection can be

accomplished under ambient conditions via optical pumping with appropriately polarized laser light [249,250]. This method is well known and has been used to achieve the drift transport of optically induced spin-polarized carriers in semiconductors [251]. However, ultimate device integration will require electrical spin injection, which can be accomplished either by injection from a spinpolarized source or by spin-ﬁltering unpolarized carriers at the interface [22,246]. Several experiments have successfully demonstrated spin injection into semiconductors and additional concepts for spin ﬁlters and spin aligners have been proposed [252]. It is envisioned that the merging of electronics, photonics and magnetics will ultimately result in new spin-based multifunctional devices, e.g., MTJs, spin-FETs, spin light-emitting diodes (spinLEDs), spin resonant tunneling devices, optical switches operating at terahertz frequency, and modulators and quantum bits for quantum computation and communication [7]. For example, MTJs and spin-FETs fabricated by DMO-based transparent heterostructures are not only effective models for studying spin transport and spin injection, but are also promising in information processing [237]. Spin-polarized electrons are injected via an n-contact to a pcontact in a magnetic semiconductor-based LED with a TM-doped ZnO n-contact and a conventional non-magnetic p-contact. ZnO is expected to be one of the most promising materials for RT polarized light emission, but optical spin polarization in ZnO has not been detected to date owing to the short spin relaxation time, which likely results from the Rashba effect. Possible solutions involve either cubic-phase ZnO or the use of additional stressor layers to create greater spin splitting to obtain polarized light emission from these structures, as well as the consideration of alternative semiconductors and fresh device approaches [159,253,254]. Despite persistent efforts by many groups, spin injection from a conventional FM metal into a semiconductor has proved highly inefﬁcient, and is limited to <0.1% [246,255]. Although promising advances in spin-injection efﬁciency from metal contacts into semiconductors have been recently been achieved (using injection from ohmic contacts or ballistic point contacts, tunnel injection or hot electron injection), there are generally difﬁculties in reproducibility [256]. Some experiments also yielded high spin-injection efﬁciencies; however, other experiments revealed no unambiguous results, and for many of the concepts proposed even a proof of principle is still lacking [91]. In contrast, efﬁcient spin injection has recently been successfully demonstrated in all semiconductor tunnel diode structures using either a spin-polarized DMO as the

F. Pan et al. / Materials Science and Engineering R 62 (2008) 1–35

injector or a paramagnetic semiconductor under high magnetic ﬁeld as a spin ﬁlter [257–259]. According to this concept, making carriers ﬂow from DMOs to a conventional semiconductor, such as from TM:ZnO to ZnO, would be an effective way to realize high spin injection efﬁciency because of their minimum crystal structure mismatch, good chemical compatibility and similar electric properties. 5.1. Co-doped ZnO-based magnetic tunnel junctions The main trigger for the search for spin injection was the discovery of GMR in 1988 by Gru¨nberg [260] and Fert [261], who were awarded the 2007 Nobel Prize for Physics. In the simplest case, a very thin, non-magnetic metal layer is sandwiched between two FM metal layers. Later, CoFe/Al2O3/Co and NiFe junctions were found to show TMR of 11.8% at RT [262]. These publications prompted a focus on the study of GMR and TMR [247,263,264]. TM:ZnO and similar DMOs with RTFM have been prepared to provide materials in which the spin freedom of electrons can be manipulated for spin transport electronics [16]. Steady progress is being made on this front, but few reports have been published on spin-polarized transport in heterostructures made from doped ZnO [88,89,265]. In fact, spin injection from a doping ZnO layer to pure ZnO and other conventional semiconductors is a key longterm goal of the extensive study of ZnO DMOs [87,88]. 5.1.1. Spin-polarized transport in (Zn,Co)O/ZnO/(Zn,Co)O junctions Large TMR at low temperature (e.g., 4 K) has been obtained in (Ga,Mn)As-based MTJs [266–269]. However, the operation of these junctions is reduced to a rather low temperature (<30 K) owing to the low TC (170 K) of (Ga,Mn)As [1,2,235]. It is important to develop MTJs of high-TC DMOs (such as TM-doped ZnO) to enhance the operation temperature, and eventually to achieve RT device

25

paradigms. In fact, besides TM:ZnO, Ti0.95Co0.05O and (Zn,Cr)Te with TC above RT have been used as one of the magnetic electrodes in MTJs, possessing TMR up to 180 and 250 K, respectively, but are limited by inelastic tunneling conduction due to the low quality of the amorphous AlOx barrier and/or the junction interface [270– 272]. The TMR temperature and spin polarization could be enhanced by optimization of the interface and barrier quality in MTJs [264,270–272], in analogy to single-crystal Fe/MgO/Fe MTJs with large TMR of 180% at 300 K [273]. Based on the concepts described above, TM:ZnO-based heterostructures have been developed to study spin injection and spin-polarized transport [88,89,274]. In particular, (Zn,Co)O was selected for fabrication of two FM electrodes, with a pure insulating ZnO layer as the barrier, which is compatible with (Zn,Co)O and allows epitaxial growth [88,89]. (Zn,Co)O/ZnO/ (Zn,Co)O epitaxial structures satisfy two primary requirements for spin injection: a minimum conductance mismatch and high spin polarized contacts [265]. In this scenario, TMR is expected to persist up to RT and exhibit high V1/2 to overcome one of the main shortcomings of MTJs for technical applications: a dramatic decrease in TMR with bias. Moreover, this MTJ structure indeed provides an approach for spin injection from a (Zn,Co)O electrode to ZnO, detected by measuring SP tunneling into the other (Zn,Co)O electrode. Fig. 27a shows a cross-sectional view of (Zn,Co)O (50 nm)/ZnO (4 nm)/(Zn,Co)O (25 nm) junctions. The two (Zn,Co)O layers serve as the bottom (50 nm) and top (25 nm) magnetic electrodes (r = 1.7 102 V cm), which are separated by an insulating ZnO barrier with a ﬁlm thickness of 4 nm (107 V cm). For vertical transport measurement, the ZnO barrier and the Zn0.94Co0.06O top electrode were patterned using shadow mask into a circular junction with a diameter of 300 mm (Fig. 27b). Fig. 27c shows a typical atomic force microscopy image of the junction surface. The

Fig. 27. (a) Schematic representation of a (Zn,Co)O/ZnO/(Zn,Co)O junction. (b) SEM image of an MTJ array. (c) AFM image of junctions for a scan area of 3 mm 3 mm. (d) XRD spectrum of the MTJ. The inset displays a pole ﬁgure for the wurtzite(1 0 1) plane.

26

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Fig. 28. (a) Low-magniﬁcation image of MTJ cross-section. The dashed line denotes the (Zn,Co)O/Al2O3 interface and areas ‘‘b’’ and ‘‘c’’ correspond to the images in (b) and (c), respectively. (b) HRTEM image of (Zn,Co)O/ZnO/(Zn,Co)O heterostructure. (c) Cross-sectional HRTEM image of (Zn,Co)O/Al2O3 interface. (d) SAED pattern of the junction (F, ﬁlm; S, substrate).

root-mean-square of the surface roughness is approximately 0.9 nm for a scan area of 3 mm 3 mm, indicating that the junction has a smooth surface. The sharp and intense wurtzite ZnO(0 0 2) peak observed in Fig. 27d indicates that the trilayered ﬁlms are highly crystalline. This is further conﬁrmed by a pole ﬁgure of the ﬁlms acquired from the (1 0 1) planes, which was obtained by azimuthally rotating the sample from 08 to 3608 at tilt angles from 208 to 908. Six poles are evident, in agreement with the sixfold symmetry of hexagonal ZnO. The low-magniﬁcation image in Fig. 28a reveals the uniform thickness of the junction of 80 nm. The HRTEM image in Fig. 28b shows the good crystallinity of the heterostructure acquired for area ‘‘b’’ in Fig. 28a. It was found that both the (Zn,Co)O electrodes and the ZnO barrier behave as singlecrystal structures with well-deﬁned interfaces. A typical crosssectional image of the (Zn,Co)O/Al2O3 interface is shown in Fig. 28c for area ‘‘c’’ in Fig. 28a. The SAD pattern in Fig. 28d indicates that the orientation relationship is wurtzite(0 0 2)jjsapphire(0 0 3). It is thus concluded that epitaxial growth of (Zn,Co)O-based MTJs occurs from the bottom (Zn,Co)O electrode via the ZnO barrier to the top (Zn,Co)O electrode, although some defects exist. The spin injection and spin transport characteristics are manifested directly in the magnetotransport properties. Fig. 29 shows the TMR dependence on the magnetic ﬁeld at four typical temperatures. The TMR ratio is deﬁned as TMR% = [R(H) R(0)]/ R(0), where R(H) and R(0) are the ﬁeld-dependent resistance and zero-ﬁeld resistance, respectively. A common feature is that all curves show positive TMR with symmetrical behavior and a high saturation ﬁeld >2 T. On the other hand, the TMR curves do not exhibit a butterﬂy shape; instead, the TMR monotonously increases with increasing magnetic ﬁeld. It is noted that Co:ZnO junctions have a coercivity ﬁeld of 11 mT and saturation of 0.2 T. Although the magnetization is accompanied by an increase in parabolic resistance, it is clear that the high TMR saturation

(>1 T) does not reﬂect any speciﬁc correlation with magnetization data. This result is supported by work carried out by Ramachandran et al. [265], who reported that the positive TMR apparently saturates and starts to decrease at intermediate magnetic ﬁelds, and that the value of this magnetic ﬁeld decreases with decreasing temperature. The positive TMR can be attributed to spin splitting of the conduction band and a subsequent redistribution of conduction electrons when an s–d exchange interaction is present [265]. The diffusion of Co ‘‘dopes’’ the ZnO barrier near the interface and likely leads to alignment of paramagnetic Co states and a decrease in spin-ﬂip scattering; subsequent suppression of spin-ﬂip

Fig. 29. TMR curves of junctions measured at a current of 1 mA at 4, 50, 100, 200, and 300 K.

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scattering at or near the interface has been identiﬁed as a possible source of high-ﬁeld TMR [88]. A possible explanation for the intrinsic positive TMR can be found by invoking spin orientations in (Zn,Co)O. In the low ﬁeld, the entire density of states near the Fermi energy (EF) arise from the minority spin band for (Zn,Co)O [275]; the spin (S) is parallel in the bottom and top electrodes, because both electrodes are fabricated from the same (Zn,Co)O, as marked in Fig. 28. In this parallel case, the electron can easily tunnel from one electrode to the other, giving rise to low resistance in the junction. As the ﬁeld is enhanced, domain rotation becomes prevalent, accompanied by rotation of S. It is noted that the domains, as well as S, in magnetic materials are rotated sequentially rather than simultaneously, particularly for the bottom and top (Zn,Co)O electrodes of different ﬁlm thickness in our scenario. Consequently, S in the two electrodes would not be totally parallel when applying H, leading to difﬁcult transport and producing high resistivity. Moreover, these descriptions concerning sequential rotation of S most likely correspond to a monotonic increase rather than butterﬂy TMR curves.

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The magnitude of a TMR ratio is commonly deﬁned using Jullie`re’s formula, TMR = 2P1P2/(1 P1P2), where P1 and P2 are the spin polarization of FM1 and FM2, respectively [276]. Thus, SP was calculated as 31% at 4 K. This is apparently lower than the value for (Ga,Mn)As-based MTJs [266], presumably because Co:ZnO is a controllable n-type DMO with very low spin–orbit coupling. This is different to GaMnAs, which is a p-type DMO with very strong spin– orbit coupling [41]. Interestingly, the SP is estimated to be 4% at RT, revealing an effective but very small spin polarization in fully epitaxial (Zn,Co)O/ZnO/(Zn,Co)O junctions at RT. Clearly, higher SP and a better understanding of the physical phenomena require further efforts. 5.1.2. Anomalous TMR in (Zn,Co)O-based junctions Fig. 30a shows non-linear current–voltage (I–V) characteristics for MTJs up to 2 V under H = 0 and 1 T at 4 K, which is expected for a tunneling device. This is indicative of carrier injection by tunneling, taken the linear I–V curve for a single (Zn,Co)O ﬁlm into account. On the basis of I–V curves obtained at H = 1 and 0 T, TMR versus bias (TMR%–V) at different temperatures is plotted in Fig. 30b–e, where

Fig. 30. (a) MTJ I–V characteristics under applied ﬁelds of H = 0 and 1 T at 4 K. Typical bias dependence TMR curves at (b) 4, (c) 100, (d) 200, and (e) 300 K.

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TMR% = [Rjun(1 T) Rjun(0 T)]/Rjun(0 T). Although both the TMR ratio and the signal-to-noise ratio drastically decrease with increasing temperature, the parabolic TMR%–V curve remains at 300 K, revealing that the effects of bias on transport in the junction can persist up to RT. The most outstanding property of all the TMR%–V curves from 4 to 300 K is their weak bias dependence. V1/2, the voltage at which the TMR ratio is half of the maximum, exceeds 2 V over a wide temperature range (4–300 K). These values are much larger than for conventional magnetic metal and for DMO-based MTJs with AlOx, MgO and similar barriers [270–273]. The decrease in TMR with bias can be explained by inelastic scattering by magnon excitations at the FM/insulator interface [264]. The much lower bias dependence is likely due to the coherence of the (Zn,Co)O/ZnO interface owing to both the minimum conductance mismatch and the high spinpolarized contacts. The typical V1/2 value is approximately 0.2 V [262,277]. V1/2 enhancement has generally been achieved by fabrication of double barriers [278] and by improving control of the single-crystal barrier [273,279], yielding up to 1.44 V in fully epitaxial MgO double-barrier MTJs [280]. As discussed above, all-epitaxial (Zn,Co)O/ZnO/(Zn,Co)O junctions show a very large V1/2 far in excess of 2 V [88]. The double-barrier junctions grown by coherent sputtering on Al2O3(0 0 1) substrate are composed of two FM (Zn,Co)O electrodes (50 and 25 nm for bottom and top, respectively) separated by a ZnO (4 nm)/(Zn,Co)O (6 nm)/ZnO (4 nm) trilayer [89]. On the basis of I–V curves obtained at H = 2 and 0 T (Fig. 31a), TMR%–V at different temperature is plotted in Fig. 31.

Four interesting phenomena are observed at different temperatures: a very large V1/2 in excess of 4 V at T > 6 K; nearly constant TMR up to 2 V at T = 5 K; a combination of higher (at low voltage) and lower (at high voltage) TMR with increasing voltage at T = 3– 4 K; and an anomalous increase in TMR with increasing voltage at T = 2 K. The anomalous voltage-dependent TMR can be explained by the large energy separation between the Fermi level and the mobility edge [89]. A non-metal ! metal transition caused by the magnetic ﬁeld can bring the Fermi level above the mobility edge, so that the energy separation between the Fermi level and the mobility edge is large enough to cause the V1/2 behavior observed [281]. 5.2. TM-doped ZnO-based spin ﬁeld-effect transistors The most effective measurement of the quality of the DMObased FM materials is the operation of device structures such as the MTJs described above, spin-FETs and photo-induced ferromagnets [16,22,88,282]. The use of DMOs as the spin source in spintronics could allow high spin injection efﬁciency and long spin coherence. The spin-FET proposed by Datta and Das [91] and subsequent numerous clones have motivated extensive study of spin precision-controlled electronic semiconductor devices [237,282–284]. Huang et al. [285] claimed that a silicon spin transport device showed output current modulation through voltage control of spin precision, with 37% electron current spin polarization after transport through 10-mm undoped silicon, but no spin-FET based

Fig. 31. (a) MTJ current–bias voltage (I–V) characteristics up to 5 V under applied ﬁelds of H = 0 and 2 T at 3 K. Typical bias voltage dependence of TMR curves at (b) 20, (c) 10, (d) 6, (e) 5, (f) 4, (g) 3 and (h) 2 K.

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Fig. 32. Schematic diagram of an all ZnO-based spin-FET [14].

on FM semiconductors has been experimentally demonstrated yet. The primary obstacle to experimental achievement is the inability to achieve high spin injection efﬁciency at the interface between the source and channel and high spin detection efﬁciency at the interface between the channel and drain, which are key for spinFETs. It is known that the spin injection and detection efﬁciencies have to be of the order of 99.9995% to achieve an on–off conductance ratio of 105, which is a typical value for modern transistors [283]. However, the maximum spin injection efﬁciency demonstrated is only 90% at low temperature [249], which is far below the requirement for spin-FETs. Because any scattering at the interface between the ferromagnet and the semiconductor can cause spin relaxation, a minimum conductance mismatch and high chemical compatibility between these layers are necessary. Fortunately, ZnO-based DMO heterostructures can be used to fabricate spin-FETs with potential to show high spin injection efﬁciency and spin transport coherence [22,93]. Following the idea of the fully epitaxial growth of (Zn,Co)O-based MTJs and the concept of ZnO-based spin-FETs by Pearton et al. [22], a schematic diagram of a spin-FET on the basis of ZnO DMOs is presented in Fig. 32. In this scheme, (Zn,Mn)O can be a half-metallic ferromagnet, which serves as the channel, and (Zn,Co)O is used to fabricate the source [30 nm (Zn,Co)O injector] and drain [60 nm (Zn,Co)O detector], as well as a thin insulating ZnO layer acting as a gate oxide. Using FM (Zn,Co)O as the source

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and drain contact material, it should be possible to have a 100% spin-polarized electron ﬂow in the (Zn,Mn)O channel [22]. These heterostructures could be grown by magnetron sputtering to realize fully epitaxial growth and shadow mask patterning. By manipulating the relative orientation of the magnetization of the injector and detector with a different in-plane external magnetic ﬁeld and gate voltage, the current in the detector can be changed correspondingly, in analogy to the in-plane spin valve [93,285]. P AP AP The magnetocurrent ratio is MC% ¼ ðISD ISD Þ=ISD , where ISD is the magnetocurrent and the superscripts P and AP refer to parallel and antiparallel injector/detector magnetization conﬁguration, respectively. Subsequently, this magnetocurrent ratio corresponds to a P conduction electron current spin polarization of P ¼ ðISD AP P AP ISD Þ=ðISD þ ISD Þ with respect to Jullie`re’s formula [276]. The fabrication of TM-doped ZnO-based spin-FETs and the realization of reliable spin detection are highly warranted and should be the focus of intense efforts. 5.3. Ferroelectricity and giant piezoelectric d33 in V and Cr-doped ZnO Although FM in TM-doped ZnO systems has been discussed in detail, ferroelectricity, which usually exists in perovskites and ilmenites, has also been found in these systems. Recently, Yang et al. [286,287] reported the observation of ferroelectric behavior in V- and Cr-doped ZnO ﬁlms. In their experiments, ferroelectric measurements were performed on Ag/Zn1xTMxO/n-Si (TM = V,Cr) sandwich structures in which Ag was deposited on ﬁlms as the top electrodes and a conductive Si substrate served as the bottom electrode. Fig. 33a and b show polarization–electric ﬁeld (P–E) hysteresis loops for V- and Cr-doped ZnO ﬁlms, respectively. Although the characteristics are not fully saturated, the hysteretic behavior still indicates the presence of ferroelectricity. The shift of the loop center observed for these Ag/Zn1xTMxO/n-Si structures is considered to result from the asymmetric electrode structure. The ferroelectric behavior in V- and Cr-doped ZnO ﬁlms was further conﬁrmed by well-shaped displacement-applied ﬁeld (D–E) ‘‘butterﬂy’’ loops, as shown in Fig. 33c and d, respectively

Fig. 33. (a) Frequency dependence of the P–E characteristics of Zn0.975V0.025O ﬁlms measured at room temperature. (b) P–E characteristics of Ag/Zn0.94Cr0.06O/n-Si FET measured at 2 kHz and room temperature. Piezoelectric hysteresis loop (black) and D–V curve (red) for (c) Zn0.975V0.025O and (d) Zn0.94Cr0.06O ﬁlms. (For interpretation of the references to color in this ﬁgure legend, the reader is referred to the web version of the article.)

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Fig. 35. Effective d33 as a function of V concentration x in Zn1xVxO ﬁlms, bulk ZnO and ZnO nanobelts.

Fig. 34. (a) Room-temperature magnetic hysteresis of Zn0.985V0.015O ﬁlms. The inset shows the same parameters on an expanded scale. (b) Temperature dependence of (&) FC and (~) ZFC magnetizations for Zn0.985V0.015O ﬁlms. The inset shows the temperature dependence of the difference between the FC and ZFC data.

[287,288]. It should be noted that ferroelectricity in ZnO-based systems has also been reported for Li- and Mg-doped ZnO bulk and/or ﬁlms [289–291]. Piezoelectric hysteresis loops (d33-applied ﬁeld) could be calculated from D–E curves and clearly showed that the Zn1xTMxO ﬁlms are switchable and the ferroelectricity is retained. More importantly, high piezoelectric d33 coefﬁcients were observed for ferroelectric V- and Cr-doped ZnO ﬁlms. The observation of ferroelectricity in V- and Cr-doped ZnO is particularly intriguing in light of recent reports of FM in these systems, suggesting the possibility of engineering a multiferroic material with simultaneous FM and ferroelectricity. Fig. 34a displays a typical magnetic hysteresis loop (M versus H) for a Zn0.985V0.015O ﬁlm measured at 300 K. The sample exhibits weak FM and saturation magnetization of 0.13 mB/V. A coercive ﬁeld of 8 mT can be observed on the expanded scale of the hysteresis loop. Fig. 34b shows the temperature dependence of the ﬁeldcooled (FC) and zero-ﬁeld-cooled (ZFC) magnetization for the Zn0.985V0.015O ﬁlm. The data were taken in a magnetic ﬁeld of 1500 Oe applied parallel to the ﬁlm surface. The inset shows the magnetization difference between FC and ZFC data as a function of temperature, i.e., DM(T) = MFC(T) MZFC(T). The DM–T curve exhibits a decreasing trend overall, but the DM value is still reasonably positive near RT. Other mechanisms such as spin glass effects and superparamagnetic clusters could also present hysteresis, but they would exhibit a value of DM = 0 below 100 K [292]. Therefore, the hysteresis arises from intrinsic FM of insulating Vdoped ZnO, which is a special multiferroic. Multiferroics that show

simultaneous FM and ferroelectric ordering have attracted extensive interest since they have potential applications in important devices such as transducers and memory [292,293]. The piezoelectric properties of ZnO ﬁlms doped with different V concentrations have been systematically investigated by SPM characterization by Yang et al. [288] and Zhao et al. [294]. Fig. 35 shows the piezoelectric d33 coefﬁcients reported for bulk ZnO and quasi-one-dimensional nanobelts and the average d33 coefﬁcient as a function of vanadium concentration in Zn1xVxO ﬁlms. The d33 coefﬁcient for bulk ZnO and a nanobelt is 9.9 and 26.7 pC/N, respectively; the latter was the highest d33 value for ZnO-based systems before ferroelectric TM-doped ZnO systems were found. The d33 value measured for an undoped ZnO ﬁlm was 11.8 pC/N. However, when 2.5 at.% V was introduced, a peak d33 value of 110 pC/N was obtained, approximately one order of magnitude greater than the value for pure ZnO ﬁlm and fourfold greater than for the ZnO nanobelt. A comparably high d33 value of 120 pC/N was obtained in ferroelectric Zn0.94Cr0.06O ﬁlms. The intriguing piezoelectric properties mean that V- and Cr-doped ZnO are promising candidates for sensors, actuators, and transducers. The origin of ferroelectric behavior in ZnO-based systems has been the subject of much discussion, with two main interpretations: the so-called size-mismatch model and the electronic origin model. It is commonly believed that the probable reason for the ferroelectric behavior is that smaller substitutes can occupy offcenter positions and form permanent electric dipoles, which results in spontaneous polarization. Four ZnO-based systems have been found to exhibit ferroelectricity so far: Li:ZnO, Mg:ZnO, V:ZnO, and Cr:ZnO. The ionic radius of Li+, Mg2+, V5+, and Cr3+ is 0.6, 0.65, 0.54, and 0.63 A˚, respectively, which are all smaller than that of Zn2+ (0.74 A˚) [286,287,289,295]. Some researchers have suggested that a change in electronic conﬁguration could also be a possible reason for this novel ferroelectricity [286,295]. For instance, the replacement of Zn2+ (1s2 2s2 2p6 3s2 3p6 3d10) by V5+ (1s2 2s2 2p6 3s2 3p6) with no 3d electrons likely affects the 3d–2p hybridization between Zn and O in V-doped ZnO systems, which could change the Zn(V)–O bond length and induce local dipole moments. However, both models are still assumptions and convincing evidence is required for conﬁrmation. XANES provides advantageous local and chemical information and is sensitive to subtle differences in local arrangement, which can yield a better understanding of ferroelectric behavior. Therefore, XANES was recently used to characterize ferroelectric Vdoped ZnO systems [286]. Qualitative analysis of XANES V K-edge spectra revealed that V substitutes replacing Zn are in the 5+ state.

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Fig. 36. Experimental and MS calculated V K-edge XANES spectra for V:ZnO ﬁlms.

To verify the assumption that V5+ ions are off-center, V K-edge XANES spectra were calculated via the full MS ab initio method. It was found that three main features (peaks A, B and C) match the experimental curve to a certain extent by simply replacing the central Zn by V in a designed V-doped ZnO cluster (Fig. 36). However, the relative height of peaks B and C differ, and there is a much lower pre-edge peak compared to the experimental spectrum. Constriction of local VO4 units (C) and displacement of V5+ ions from the cell center (D) were then introduced as parameters in the calculation. Introduction of VO4 constriction improved the ﬁt to the experimental curve. When V5+ displacement was also considered, the closest ﬁt was obtained. This accurate replication of the characteristics of the experimental spectrum implies that V5+ displacement likely occurs and is responsible for ferroelectricity in the V:ZnO system. Thus the XANES calculations provided a direct evidence for the sizemismatch model and clariﬁed the ferroelectric origin in V:ZnO system. For ferroelectrics with a centrosymmetric paraelectric phase, d33 can be expressed as follows: d33 ¼ 2Q eff e0 er Ps ;

(5)

where Qeff is the effective electrostriction coefﬁcient, and e0 and er are the permittivity of free space and relative permittivity, respectively. Eq. (5) indicates that the piezoresponse in ferroelectrics is directly dependent on Ps via a factor linking the electrostrictive characteristics and the dielectric constants. On the other hand, this type of dependence also exists in ferroelectrics with a non-centrosymmetric paraelectric phase [288]. Therefore, the giant d33 in V- and Cr-doped ZnO was attributed to the emergence of switchable Ps. Moreover, the relatively high er in Vand Cr-doped ZnO ﬁlms induced by switchable Ps may also contribute to the high d33. A microscopic explanation was also

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proposed for a V-doped ZnO system [288]. Since the dominant effect of the electric ﬁeld in wurtzite semiconductors is to rotate bonds that are non-collinear with the polar c-axis, i.e., Zn2–O1 bonds [287], towards the direction of the applied ﬁeld, thus producing piezoelectric strain, the electromechanical response in wurtzite materials should be mainly governed by the ease of bond bending [296]. In V-doped ZnO ﬁlms, since V5+ ions have a higher positive charge than Zn2+ ions, the non-collinear V–O1 bonds ought to have a stronger polarity than Zn2–O1 bonds and hence can rotate more easily in an applied ﬁeld. On the other hand, (2 at.%) Cu-, Ni-, Co- and Fe-doped ZnO ﬁlms have much smaller d33 values of 13.5, 9.8, 10.7 and 5.8 pC/N, respectively, which may be primarily because the radii of Cu2+ (0.72 A˚) and Ni2+ (0.72 A˚) are almost equal to that of Zn2+ (0.74 A˚), whereas those of Co2+ (0.79 A˚) and Fe2+ (0.76 A˚) are greater than that of Zn2+ [297,298], different from the scenario for Li+, Mg2+, V5+, and Cr3+ ions. V- and Cr-doped ZnO show simultaneous FM and ferroelectricity, and are thus promising multiferroic materials. TM-doped ZnO is also considered as a promising optical material that could be extensively used for transparent electrodes/devices and the window of solar cells. Some optoelectronic applications of ZnO overlap with those of GaN, which is now widely used for LEDs. However, ZnO has some advantages over GaN: the large exciton binding energy (60 meV) of ZnO could lead to lasing action based on exciton recombination, even above RT, which is widely used for production of green, blue-UV, and white LEDs, as discussed in numerous reviews and reports [21,22,299,300,301]. It is worth pointing out that ZnO has much simpler crystal-growth technology, resulting in potential lower costs for ZnO-based devices. The transmittance is usually greater than 80% for TM-doped ZnO in the visible region [48,109,129], indicating that these materials are suitable for transparent spintronics applications with integrated TM-doped ZnO layers [88,89,302,303]. Thus, ZnO could be used as materials with multi-functions, including FM, ferroelectric, piezoelectric, optical, and electric features, which could play important roles in spintronics applications of TM-doped ZnO. 6. Conclusions and outlook In conclusion, we have presented a comprehensive review of the preparation parameters that affect the local TM structure and magnetization of TM-doped ZnO-based spintronics prototypes. ZnO ﬁlms have been the subject of varying degrees of research effort over recent decades. This effort has recently been intensiﬁed for the preparation and elucidation of TM-doped ZnO DMOs with RTFM and the development of spin-polarized injection, transport and detection experiments on TM-doped ZnO-based heterostructures for spintronics applications. Although it is clear that TMdoped ZnO systems exhibit RTFM, there are still many open questions regarding its origin and use in spintronics. By extracting a single variable from a series of experiments, carriers and defects can be separated and direct relationships between FM and the presence of structural defects can be identiﬁed. It has been unambiguously demonstrated that structural defects are responsible for RTFM in TM-doped ZnO, whereas carriers involved in carrier-mediated exchange are natural by-products of the creation of defects in ZnO. In contrast, an increase in electron density decreases the acceptor concentration, thus quenching FM exchange. We have described recent signiﬁcant advances in TMdoped ZnO-based prototype spintronics, in which TM-doped ZnObased MTJ structures provide a compass for the spin injection from (Zn,Co)O to ZnO at RT and shows a large ‘‘half voltage’’ (>4 V). In addition, second generation spintronics integrating TM-doped ZnO ﬁlms with promising ferromagnetic, ferroelectric, piezoelectric and optical properties tend to fabricate new ﬂexible devices with

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the ability to create and manipulate spin-polarized carriers in a conventional semiconductor. Although FM ordering arising from defects in TM-doped ZnO have been determined, defects are difﬁcult to characterize experimentally and measure quantitatively, which is important to profoundly understand the origin of magnetic behaviors in this system. Considerably designing experiments to accurately control the formation of defects and quantitatively build the correlations between defects and magnetic moments are highly warranted. The best way to investigate the innumerable possible defects in TMdoped ZnO may be with the help of a computer. More importantly, spin-polarized transport in TM-doped ZnO and the efﬁciency of spin injection from TM-doped ZnO to conventional semiconductors would draw extensive attention, and only in this way, DMOs can move to the real applications in useful spintronics. Spin polarized injection, transport and detection experiments in TMdoped ZnO-based heterostructures are lacking at present, the spin state at the interfaces of heterostructures would require a more precise control and careful quantum analysis. In this case, efforts are increasingly shifting not only to the design of TM-doped ZnObased prototype spintronics, but also to the task of thinking calmly methods to evaluate these devices. Acknowledgements The authors are grateful to Prof. B.X. Liu of Tsinghua University for helpful suggestions on this paper, and to Prof. G.A. Gehring of the University of Shefﬁeld for fruitful discussions. The authors also acknowledge the cooperation of X.X. Wei in providing some data, and of Dr. W.S. Yan and B. He of the National Synchrotron Radiation Laboratory in local structure measurements. This work was supported by the National Natural Science Foundation of China (Grant nos. 50772055 and 50325105), the Ministry of Science and Technology of China (Grant no. 2006CB605201-1) and a National Hi-tech (R&D) Project of China (Grant no. 2006AA03Z435). References [1] A.H. Macdonald, P. Schiffer, N. Samarth, Nat. Mater. 4 (2005) 195. [2] T. Jungwirth, J. Sinova, J. Masˇek, J. Kucˇera, A.H. MacDonld, Rev. Mod. Phys. 78 (2006) 809. [3] H. Ohno, A. Shen, F. Matsukura, A. Oiwa, A. Endo, S. Katsumoto, Y. Iye, Appl. Phys. Lett. 69 (1996) 363. [4] H. Munekata, H. Ohno, S. von Molnar, A. Segmu¨ller, L.L. Chang, L. Esaki, Phys. Rev. Lett. 63 (1989) 1849. [5] S. von Molna´r, D. Read, J. Magn. Magn. Mater. 242–245 (2002) 13. [6] Y.D. Park, A.T. Hanbicki, S.C. Erwin, C.S. Hellberg, J.M. Sullivan, J.E. Mattson, T.F. Ambrose, A. Wilson, G. Spanos, B.T. Jonker, Science 295 (2002) 651. [7] S.A. Wolf, D.D. Awschalom, R.A. Buhrman, J.M. Daughton, S. von Molna´r, M.L. Roukes, A.Y. Chtchelkanova, D.M. Treger, Science 294 (2001) 1488. [8] J.M.D. Coey, M. Venkatesan, C.B. Fitzgerald, Nat. Mater. 4 (2005) 173. [9] T. Dietl, H. Ohno, F. Matsukura, J. Cibert, D. Ferrand, Science 287 (2000) 1019. [10] K. Sato, H. Katayama-Yoshida, Jpn. J. Appl. Phys. 39 (2000) L555. [11] K. Ueda, H. Tabata, T. Kawai, Appl. Phys. Lett. 79 (2001) 988. [12] Y. Matsumoto, M. Murakami, T. Shono, T. Hasegawa, T. Fukumura, M. Kawasaki, P. Ahmet, T. Chikyow, S. Koshihara, H. Koinuma, Science 291 (2001) 854. [13] I. Zˇutic´, J. Fabian, S. Das Sarma, Rev. Mod. Phys. 76 (2004) 323. [14] S.J. Pearton, C.R. Abernathy, D.P. Norton, A.F. Hebard, Y.D. Park, L.A. Boatner, J.D. Budai, Mater. Sci. Eng. R. 40 (2003) 137. [15] D.D. Awschalom, M.E. Flatte´, Nat. Phys. 3 (2007) 153. [16] S.J. Pearton, C.R. Abernathy, M.E. Overberg, G.T. Thaler, D.P. Norton, N. Theodoropoulou, A.F. Hebard, Y.D. Park, F. Ren, J. Kim, L.A. Boatner, J. Appl. Phys. 93 (2003) 1. [17] W. Prellier, A. Fouchet, B. Mercey, J. Phys.: Condens. Matter. 15 (2003) R1583. [18] S.J. Pearton, W.H. Heo, M. Ivill, D.P. Norton, T. Steiner, Semicond. Sci. Technol. 19 (2004) R59. [19] R. Janisch, P. Gopal, N.A. Spaldin, J. Phys.: Condens. Matter 17 (2005) R657. [20] S.A. Chambers, Surf. Sci. Rep. 61 (2006) 345. ¨ zgu¨r, Ya.I. Alivov, C. Liu, A. Teke, M.A. Reshchikov, S. Dog˘an, V. Avrutin, S.-J. ¨. O [21] U Cho, H. Morkoc¸, J. Appl. Phys. 98 (2005) 041301. [22] S.J. Pearton, D.P. Norton, K. Ip, Y.W. Heo, T. Steiner, Prog. Mater. Sci. 50 (2005) 293.

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