ZnO nanowire growth and devices

Materials Science and Engineering R 47 (2004) 1–47

ZnO nanowire growth and devices Y.W. Heoa,*, D.P. Nortona, L.C. Tiena, Y. Kwona, B.S. Kangb, F. Renb, S.J. Peartona, J.R. LaRochec a

Department of Materials Science and Engineering, University of Florida, Gainesville, FL 32611, USA Department of Chemical Engineering, University of Florida, Gainesville, FL 32611, USA c Raytheon, Waltham, MA 02451, USA b

Accepted 23 September 2004 Available online 25 December 2004

Abstract The large surface area of ZnO nanorods makes them attractive for gas and chemical sensing, and the ability to control their nucleation sites makes them candidates for micro-lasers or memory arrays. In addition, they might be doped with transition metal (TM) ions to make spin-polarized light sources. To date, most of the work on ZnO nanostructures has focused on the synthesis methods and there have been only a few reports of the electrical characteristics. We review fabrication methods for obtaining device functionality from single ZnO nanorods. A key aspect is the use of sonication to facilitate transfer of the nanorods from the initial substrate on which they are grown to another substrate for device fabrication. Examples of devices fabricated using this method are briefly described, including metal-oxide semiconductor field effect depletion-mode transistors with good saturation behavior, a threshold voltage of
3 V and a maximum transconductance of order 0.3 mS/mm and Pt Schottky diodes with excellent ideality factors of 1.1 at 25 8C and very low (1.5  1010 A, equivalent to 2.35 A cm2, at 10 V) reverse currents. The photoresponse showed only a minor component with long decay times (tens of seconds) thought to originate from surface states. These results show the ability to manipulate the electron transport in nanoscale ZnO devices. # 2004 Elsevier B.V. All rights reserved. Keywords: Nanowires; Nanorods; ZnO; Bandgap

1. Introduction In recent years, significant interest has emerged in the synthesis of nanoscale materials [1]. One of the most attractive classes of materials for functional nanodevices are semiconductors. Various means have been reported for the synthesis of semiconducting nanowires and nanorods [2–4]. Much effort has focused on catalysis-driven bulk synthesis of nanomaterials using approaches that are neither substrate site specific nor compatible with most planar device platforms. Nevertheless, nanodevice functionality has been demonstrated with these materials in the form of electric field-effect switching [5], single electron transistors [6], biological and chemical sensing [7], and luminescence [8] for onedimensional (1-D) semiconducting structures. Included in the semiconductors of interest are semiconducting oxides [9–12]. Of these, zinc oxide is particularly interesting for nanodevice applications. ZnO is an n-type, direct bandgap semiconductor with Eg = 3.35 eV [13,14]. * Corresponding author. Tel.: +1 352 846 1091; fax: +1 352 846 1182. E-mail address: [email protected] (Y.W. Heo). 0927-796X/$ – see front matter # 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.mser.2004.09.001

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ZnO has been effectively used as a gas sensor material based on the near-surface modification of charge distribution with certain surface-absorbed species [15]. ZnO nanorods would provide significant enhancement in sensitivity due to high surface-to-volume ratio. ZnO is also piezoelectric, and is used in surface acoustic wave devices [16]. Huang et al. have reported the site-specific nucleation and growth of ZnO nanorods on deposited Au catalyst using a high-temperature vapor transport process [8]. As with any semiconductor, 1-D ZnO nanostructures provide an attractive candidate system for fundamental quantization and low-dimensional transport studies [17–19]. A large variety of ZnO one-dimensional structures have been demonstrated [20–47]. The large surface area of the nanorods and bio-safe characteristics of ZnO makes them attractive for gas and chemical sensing and biomedical applications, and the ability to control their nucleation sites makes them candidates for micro-lasers or memory arrays. To date, most of the work on ZnO nanostructures has focused on the synthesis methods. There have been only a few reports of the electrical characteristics [20–24]. The initial reports show a pronounced sensitivity of the nanowire conductivity to ultraviolet (UV) illumination and the presence of oxygen in the measurement ambient. There is strong interest in developing solid-state ozone and hydrogen gas sensors for use in both industry and domestic applications. Ideal sensors have the ability to discriminate between different gases and arrays that contain different metal oxides (e.g. SnO2, ZnO, CuO, WO3) on the same chip can be used to obtain this result. The fact that ZnO can be grown at low temperatures on cheap substrates such as glass also makes it attractive for transparent electronics. In this review, we report on the synthesis of both cored and radial heterostructure nanorods and simple, reproducible methods for nanorod device fabrication and give some examples of device functionality. The ability to control the synthesis of high quality ZnO nanowires leads to potential applications in UV photodetection, gas sensing and transparent electronics.

2. ZnO nanorod synthesis Previous effort in the synthesis of ZnO nanowires and nanorods have employed vapor-phase transport via a vapor–liquid–solid (v–l–s) mechanism [46,47], gas reactions [48] and oxidation of metal in the pores of anodic alumina membranes [49,50]. While these materials provide interesting systems for investigating fundamental properties or for exploring device concepts via single prototype device construction, the ability to synthesize nanorods at arbitrary locations at moderate temperatures is needed for nanodevice integration. This requires site-specific nucleation of nanorods, as well as a growth process that remains site specific and is compatible with the device platform of interest. It would be advantageous to achieve nanorod growth from a flux source that could be controlled at the atomic level, thus enabling compositional modulation along the rod length. In this section, we report on the site-selective growth of ZnO nanorods using a catalysis-driven molecular beam epitaxy (MBE) method. Low temperature MBE conditions are identified so that ZnO nucleation and growth occurs only on the deposited metal catalyst. With this approach, site specific, single crystal ZnO nanorod growth is achieved with nanorod diameters as small as 15 nm. The growth experiments were performed using a conventional MBE system. The background base pressure of the growth chamber was
5  108 mbar. An ozone/oxygen mixture was used as the oxidizing source. The nitrogen-free plasma discharge ozone generator yielded an O3/O2 ratio on the order of 1–3%. No effort was made to separate the molecular oxygen from the ozone. The cation flux was provided by a Knudsen effusion cell using high purity (99.9999%) Zn metal as the source. Cation and O2/O3 partial pressure was determined via a nude ionization gauge that was placed at the substrate position prior to growth. The beam pressure of O3/O2 mixture was varied between 5  106 and

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5  104 mbar, controlled by a leak valve between the ozone generator and the chamber. The Zn pressure was varied between 5  107 and 4  106 mbar. The substrates were Si wafers with native SiO2 layer terminating the surface. Site-selective nucleation and growth of ZnO was achieved by coating Si substrates with Ag islands. For thick Ag, a continuous ZnO film could be deposited. For nominal Ag film thicknesses of ˚ , discontinuous Ag islands are realized. On these small metal islands, ZnO nanorods were 20–200 A observed to grow. Efforts to deposit ZnO on Ag-free SiO2/Si surface area under a variety of growth conditions proved unsuccessful. Zn metal deposition could be achieved at substrate temperatures of 25–100 8C, but with no ZnO formation for a wide range of O2/O3 partial pressure. Higher substrate temperatures yield no deposition as the Zn metal vapor pressure rises quickly at moderate temperatures. Typical growth times for ZnO on the Ag-coated silicon was 2 h with growth temperatures ranging from Tg = 300 to 500 8C. After growth, the samples were evaluated by X-ray diffraction, scanning electron microscopy (SEM), transmission electron microscopy (TEM), and photoluminescence (PL). Fig. 1 shows a scanning electron microscopy image of ZnO nanorods grown on a Si wafer that was coated with a nominally 10 nm thick layer of Ag. The Ag was deposited using e-beam evaporation. The images are for ZnO nanorods grown at 400 8C with a Zn pressure of 2  106 Torr Torr and an oxygen/ozone pressure of 5  104 Torr. Under these conditions, ZnO deposition was observed only on the Ag with no growth on regions of the SiO2-terminated Si surface that was devoid of Ag. A dense entangled collection of ZnO nanorods is observed to grow from the surface. Both cylindrical nanorods and faceted whiskers can be observed in the forest of ZnO nanostructures grown at 400 8C. At higher temperatures, only nanorods are observed. In many cases, the length of ZnO nanorods is in excess of 2 mm. Note also that multiple nanorods are observed to nucleate from the relatively large Ag islands. As such, it does not appear that the diameter of the nanorods is determined by the initial radii of the Ag islands. X-ray diffraction of the deposited materials confirms that the

Fig. 1. SEM image of ZnO nanorods nucleated on Ag-coated Si/SiO2 substrate.

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material is ZnO. X-ray diffraction patterns taken along the surface normal, indicating only ZnO peaks. The diffraction pattern for the material grown at 400 8C is consistent with randomly oriented polycrystalline material, although selected area electron diffraction (discussed below) indicates a preferred c-axis orientation of individual nanorods along the long axis. A preferred (0 0 2) orientation seen for nanorods obtained at 500 8C indicates a more vertically aligned growth at this temperature. The most intriguing structures are those that result from isolated Ag nanoparticles. In depositing the Ag catalyst films, certain regions of the SiO2/Si surface were shadowed from deposition, leading to a gradient in Ag thickness, Ag nanoparticle coverage, and average nanoparticle diameter. Within these areas, isolated Ag nanoparticles could be located, thus allowing direct imaging of nanorod formation from individual Ag islands. Clusters of ZnO nanorods were observed to nucleate from these isolated Ag islands. Fig. 2 shows field-emission SEM images of ZnO nanorod clusters, including a highresolution image of a single nanorod. Energy-dispersive spectrometry was used to determine the nanorod composition (ZnO) in addition to confirming the absence of ZnO on regions of the substrate surface that are devoid of Ag. In order to acquire these images, the sample was coated with a thin layer of carbon to avoid charging effects. From the high-resolution image, the nanorod cross-section appears to be cylindrical, although any faceting of the side walls would be obscured by the carbon coating. The thickness of the nanorod shown in Fig. 2 is on the order of 30 nm, although the carbon coating may exaggerate this thickness. In addition to SEM, the nanorods were examined using transmission electron microscopy. Fig. 3 shows a transmission electron microscopy image of an individual ZnO nanorod. The rod was imaged from a cross-sectioned sample. The nanorod shown in Fig. 3 was not carbon coated. An estimate of the rod thickness is 20 nm. Selected area diffraction (SAD) of nanorod specimens indicates that the rods are single crystal ZnO, with the c-axis oriented along the long axis of the rod. Also evident in the image is a small particle embedded at the tip of the rod. This is similar to what is observed for other nanorod synthesis that is driven by a catalytic reaction, where catalyst particles become suspended on the nanorod tip [51,52]. Evidence for termination of the ZnO nanorods tips with catalyst particles is also observed in field-emission SEM images. Local energy-dispersive spectrometry (EDS) measurements indicate that the terminating particle is Ag, although more characterization is needed in order to confirm this. The mechanism for nanorod growth is catalysis driven, and appears to be related to the vapor– liquid–solid model reported for the nanorod synthesis of other materials. ZnO nanoparticle formation

Fig. 2. Deterministic growth of ZnO nanorod clusters formed via catalytically driven MBE; nanorod diameter is 20–30 nm.

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Fig. 3. TEM and selected area diffraction image of a single crystal ZnO nanorod.

via the internal oxidation of Zn in Ag/Zn alloys has previously been reported [53]. In these studies, oxygen is diffused into an Ag/Zn alloy, with nanoscale ZnO precipitates forming in the bulk of the sample. For the present case of nanorod formation, the reaction between ozone/oxygen flux and the Ag islands appears to result in surface and subsurface oxygen diffusion in the metal island, perhaps involving the intermediate formation of Ag2O. Zn atoms impinging on the Ag island surface then diffuse either on the surface or in the bulk of the island, where they react with the Ag2O to form ZnO. The solid solubility of Zn in Ag is on the order of 25 wt.% for the temperatures considered in these experiments. Zn addition to Ag significantly suppresses the melting point to 710 8C at 25 wt.% Zn. Note that the melting point of Zn is rather low (420 8C). Based on these arguments, one might anticipate rather high diffusion rates for Zn in Ag for the temperatures considered. Note that the temperatures at which vapor–liquid–solid growth was previously reported for ZnO are significantly higher than that used with the present work. It should also be noted that the addition of Ag during the growth of complex oxide thin film has been reported to be effective in enhancing the oxidation process for various oxide thin-film compounds [54]. In addition to examining the structure with microscopy and X-ray diffraction, the optical properties of the nanorods were examined using photoluminescence. A He–Cd (325 nm) laser was used as the excitation source. The room temperature luminescence reveals a robust near band edge emission peak located at 375 nm indicating that that rods are highly crystalline. This is consistent with luminescence reported for near-band edge emission in epitaxial films [29,30] and larger diameter ZnO nanorods [55]. A broad, but weak, green emission peak is also observed at
520 nm that is typically associated with trap-state emission attributed to singly ionized oxygen vacancies in ZnO [56].

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The emission is the result of the radiative recombination of photogenerated holes with electrons occupying the oxygen vacancy. Similar results have been observed for ZnO nanorods formed via vapor transport. Enhancement in the green emission in nanorods as compared to bulk may be attributed to a higher density of vacancies in the rods. This may be due to the higher surface area to volume ratio for nanorods as compared to bulk.

3. Structure and optical properties of cored wurtzite (Zn,Mg)O heteroepitaxial nanowires Nanowire growth has been reported using several techniques [57,58], and has included numerous semiconductors [59,60], including the oxides Ga2O3 [61], In2O3 [62], SnO2 [63], and ZnO [64]. Despite significant progress, major challenges in the manipulation of nanowire materials remain. The fabrication of integrated systems using nanowire material requires the site-specific growth or placement of nanowires on relevant device platforms. In addition, the formation of complex, multi-component structures and interfaces are needed for low-dimensional structures and electronic devices. In thin-film semiconductor research, the formation of heteroepitaxial interfaces has proven to be useful in the development of numerous device concepts, as well as in the investigation of lowdimensional phenomena [65]. Unfortunately, such heterostructures have rarely been realized outside of the conventional 2-D planar thin-film geometry [66]. Nevertheless, the synthesis of 1-D linear heterostructures is scientifically interesting and potentially useful, particularly if a technique is employed that allows for spatial selectivity in nanowire placement. Addressing these challenges could prove useful in realizing integrated device functionality involving semiconducting nanowires for a number of applications, including nanoscale electric field-effect transistors [67], single electron transistors [68], biological and chemical sensors [69], electron emitters [70], optical emitters and detectors [69,70]. In this section, the properties of 1-D heteroepitaxial structures are described. In particular, the structural and optical properties of cored (Zn,Mg)O nanowires, formed via self-assembled bimodal growth, are discussed. ZnO is among the more interesting and important semiconducting oxides [71]. ZnO is an n-type, direct bandgap semiconductor with Eg = 3.35 eV. Electron doping via defects originates from Zn interstitials in the ZnO lattice. The intrinsic defect levels that lead to n-type doping lie 0.05 eV below the conduction band. The room temperature Hall mobility in ZnO single crystals is among the highest for the oxide semiconductors, on the order of 200 cm2 V1 s1. The exciton binding energy for ZnO is on the order of 60 meV, yielding efficient luminescence at room temperature. The synthesis of ZnO nanowires and nanorods has been demonstrated using vaporphase transport via a vapor–liquid–solid mechanism [72], gas reactions [73], and oxidation of metal in the pores of anodic alumina membranes [74]. Room temperature ultraviolet lasing via optical pumping has been demonstrated with ZnO nanorods on deposited Au catalyst using a high-temperature vapor transport process [70]. Recently, we reported on catalyst-driven molecular beam epitaxy of ZnO nanorods [75]. The process is site specific, as single crystal ZnO nanorod growth is realized via nucleation on Ag films or islands that are deposited on a SiO2-terminated Si substrate surface. Growth occurs at relatively low substrate temperatures, on the order of 300–500 8C, making it amenable to integration on numerous device platforms. With this approach, nanorod placement can be predefined via location of metal catalyst islands or particles. The heteroepitaxial cored nanostructures described here are based on the (Zn,Mg)O alloy system, and were synthesized using the catalysis-driven molecular beam epitaxy method. Details of the growth experiments are reported elsewhere. An ozone/oxygen mixture was used as the oxidizing source. The cation flux was provided by Knudsen effusion cells using high purity (99.9999%) Zn metal and Mg

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Fig. 4. Field-emission scanning electron microscopy image of cored (Zn,Mg)O nanorods grown on Ag-coated Si. The conditions for growth were Tg: 400 8C; Zn pressure: 3  106 mbar, Mg pressure: 4  107 mbar, O2/O3 pressure: 5  104 mbar. The Mg source was shuttered with a 60 s open/60 s closed cycle.

(99.95%) as the source materials. The substrates were Si wafers with native SiO2 terminating the surface. No effort was made to remove the native oxide or to terminate the surface with hydrogen. Site-selective nucleation and growth of cored nanorods was achieved by coating Si substrates ˚ , discontinuous Ag islands are realized. On with Ag islands. For a nominal Ag film thickness of 20 A these small metal catalyst islands, (Zn,Mg)O nanorods were observed to grow. Fig. 4 shows an FESEM micrograph of these nanorods, indicating a length approaching 1 mm. The growth temperature was 400 8C. Energy-dispersive spectrometry measurement was performed on a single nanowire under the Transmission electron microscopy. EDS confirmed the presence of Mg in the nanorod, as seen in Fig. 5. Typical growth times for (Zn,Mg)O on the Ag-coated silicon was 2 h with growth temperatures ranging from Tg = 300–500 8C. The site specificity for nanowire growth using this technique is evident

Fig. 5. Energy dispersive spectrometry data for (Zn,Mg)O nanorods.

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Fig. 6. Field-emission scanning electron microscopy image of cored (Zn,Mg)O nanorods grown on a patterned Ag-coated Si substrate. The conditions for growth were Tg: 400 8C; Zn pressure: 3  106 mbar; Mg pressure: 4  107 mbar; O2/O3 pressure: 5  104 mbar. Note that (Zn,Mg)O nanorod nucleation occurs only on the catalyst-coated regions.

in Fig. 6, showing FE-SEM images of (Zn,Mg)O nanorods on a Ag-patterned substrate grown in a Zn pressure of 3  106 mbar, a Mg pressure of 4  107 mbar, and an O2/O3 pressure of 5  106 mbar. (Zn,Mg)O nanorods form only on the Ag-coated regions. The potential for growing single nanorods on selected locations is exemplified in Fig. 7, where single ZnO nanorods are nucleated on Ag nanoparticles dispersed on a SiO2-terminated Si surface. In order to acquire these images, the sample was coated with a thin layer of carbon to avoid charging effects. From the high-resolution image, the

Fig. 7. Field-emission scanning electron microscopy image of individual ZnO nanorods grown on Ag nanoparticles dispersed on the Si substrate.

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nanorod cross-section appears to be cylindrical, although any faceting of the side walls may be obscured by the carbon coating. The thickness of the nanorods shown is on the order of 30 nm, although the carbon coating may exaggerate this thickness. The formation of the (Zn,Mg)O nanorods includes the v–1–s mechanism described earlier, although heteroepitaxial growth occurs as well as will be seen. Fig. 8 shows a Z-contrast scanning transmission electron microscopy (Z-STEM) image of an individual (Zn,Mg)O nanorod grown at

Fig. 8. Z-contrast scanning transmission electron microscopy image (a) of a (Zn,Mg)O nanorod with (b) a Ag catalyst particle at the rod tip.

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400 8C, with a Zn pressure of 3  106 mbar, a Mg pressure of 4  107 mbar, and an O2/O3 pressure of 5  104 mbar. The Mg flux was cycled on and off every 60 s, which was inconsequential to the nanorod structure. Evident in the image is a small particle embedded at the tip of the rod. This is similar to what is observed for other nanorod synthesis that is driven by a catalytic reaction, where catalyst particles become suspended on the nanorod tip. The diameter of the catalyst particle is
6 nm. Note that, at the nanorod tip, the contrast in the Z-STEM image is relatively uniform, indicating uniform cation distribution. However, the rod diameter is also tapered along the length, being thicker at the base than on the tip, with an average diameter on the order of 10 nm. As reported elsewhere, a radial segregation of the Zn and Mg occurs during growth [76]. A Znrich core surrounded by a Mg-rich sheath is observed as seen in Fig. 9. As discussed elsewhere, in bulk material, the solubility of Mg in ZnO is relatively low, on the order of 4 at % [77]. In contrast, Mg content as high as Zn0.67Mg0.33O has been reported to be metastable in the wurtzite structure for epitaxial thin films. For this composition, the bandgap of ZnO can be increased to
3.8 eV. For (Zn,Mg)O nanorod growth, it appears that both (hence bimodal) growth modes are relevant, but for different regions in the rod. Under low temperature MBE growth conditions, a solubility-driven segregation occurs during the catalyst-driven core formation, with the core composition determined by bulk solid solubility. Subsequently, an epitaxial sheath grows with Mg content and crystal structure determined by epitaxial stabilization. The net result is the growth of (Zn,Mg)O nanorods that are not uniform in composition across the diameter, but distinctly cored. Fig. 9 shows a highresolution Z-STEM image of a nanorod grown under the conditions described. The lattice image for the nanorod specimen indicates that the rod is crystalline with the wurtzite crystal structure maintained throughout the cross-section. The c-axis is oriented along the long axis of the rod. The higher contrast for the center core region clearly indicates a higher cation atomic mass. The structures consist of a zinc-rich Zn1–xMgxO core (small x) surrounded by a Zn1yMgyO (large y) sheath containing higher Mg content. While the nanorod imaged in Fig. 9 is crystalline across the entire cross-section, other rods exhibit sheath properties that vary along the length. In particular, consider the nanorod shown in Fig. 10. In this case, the core and sheath are both crystalline in one region of the rod. However, as one proceeds along the length, the crystallinity changes. In particular, for the rod considered, the sheath region becomes either polycrystalline or amorphous as one approaches the rod tip. Still further down the nanowire, the image suggests a lack of crystallinity for both the core and sheath, although the lack of crystallinity in the sheath may effectively obscure imaging of the core region. This change in crystallinity change along the length of the rod may reflect the fact the temperature gradient will develop along the nanowire length during growth. This occurs since the substrate is the source of heat during nanorod formation. Sheath material deposited on the tip of longer rods during the latter part of the synthesis process will do so at a lower local temperature than the material closer to the substrate. In order to assess crystalline quality and investigate possible quantization effects, the optical properties of the cored nanorods were examined using photoluminescene. Spectra were taken over the temperature range 6–300 K. A He–Cd (325 nm) laser was used as the excitation source. For the low temperature measurements, the sample was cooled using either a helium flow or closed cycle cryostat. Fig. 11 shows the photoluminescence spectra taken at various temperatures for the cored nanorod specimens. For ZnO rods (no Mg), the photoluminescene results are consistent with luminescence reported for near band edge emission in crystals [78], epitaxial films [79], and larger diameter ZnO nanorods [80]. The free exciton emission dominates luminescence, with a room temperature peak at 3.30 eV. At room temperature, the spectra for the cored (Zn,Mg)O nanorods is also dominated by the free exciton luminescence. However, the peak in luminescence at room temperature is at 3.35 eV, which is blueshifted relative to that seen in pure ZnO (peak at 3.30 eV). As the temperature is

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Fig. 9. Z-contrast scanning transmission electron microscopy images of cored (Zn,Mg)O nanorods. The high contrast of the core (a) indicates a significantly higher Zn content for the core relative to the sheath. The higher resolution image (b) of the core indicates the wurtzite structure for the core and sheath material.

decreased, the free exciton emission shifts to higher energy due to the temperature dependence of the bandgap. At the lower temperatures, the donor-bound exciton (D0,X) dominates luminescence [79]. Also seen at low temperatures is a peak at 3.35 eV, which is assigned as the corresponding longitudinal optical phonon replica (D0,X)-LO. At lower energies, at least four relatively weak peaks are also observed in the visible spectrum, indicated by arrows in the figure. Similar peaks have been reported elsewhere for bulk and thin-film ZnO materials [78,79].

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Fig. 10. Z-contrast scanning transmission electron microscopy image of regions of a nanorod showing: (a) a crystalline core and sheath; (b) an amorphous sheath; and (c) an amorphous sheath with no evidence of crystallinity in the core.

Blueshifts in the near band edge luminescence for nanoscale semiconducting structures can reflect quantum confinement effects. With a bandgap range of 3.2–3.8 eV, the ZnO/(Zn,Mg)O heterostructure system affords the opportunity to realize electron confinement due to band offsets. In semiconductors, the onset of reduced dimensionality effects should be observable for structures with dimensions below the relevant length scales. Quantization of the electron states should be observable for low dimensional structures with length scales on the order of the exciton Bohr diameter [81,82]. A significant blueshift in the near-band edge emission was previously observed for InP nanorods with diameters approaching that of the exciton diameter [83]. For ZnO, the exciton Bohr ˚ , which is approximately equal to the core diameters considered. Shifts in the neardiameter is 34 A band edge photoluminescence peak energy have been observed for 0-D ZnO nanoparticles diameters on the order of 2–6 nm [84]. For ZnO/(Zn,Mg)O superlattices, quantum confinement effects are predicted for well widths less than
5 nm, and are reported for ZnO layer thicknesses less than
4 nm [82]. ZnO/(Zn,Mg)O superlattices with a ZnO well width of 3.1 nm show a 100 meV blueshift in nearband edge luminescent peak energy [82]. In the present work, a 50 meV shift is observed for the 1-D nanorod cored structure, the core diameter (4 nm) being roughly equivalent to the bulk exciton diameter (3.4 nm). While a blueshift in near edge emission is consistent with expected quantum confinement in 4 nm diameter ZnO cores, one cannot easily differentiate this from a similar blueshift that would result from a few percent doping of Mg in the core regions. However, from the Mg dependence of the band edge peak, it appears that quantization effects on the PL shift are minimal.

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Fig. 11. Photoluminescence at various temperatures for (a) near band edge and (b) visible emission from cored (Zn,Mg)O nanorods.

Fundamentally, these 1-D ZnO nanostructures provide an attractive system for probing lowdimensional effects such as quantization, interface scattering, and ballistic transport. The relatively low growth temperatures suggest that cored (Zn,Mg)O nanorods could be integrated on device platforms for numerous applications, including chemical sensors, nano-optics, scanning probes, and nanoelectronics. More generally, the spontaneous bimodal growth mode demonstrated for (Zn,Mg)O may prove applicable in the synthesis of heteroepitaxial cored nanowire-materials where thermodynamics and epitaxy impose different cation or anion solubility limits.

4. ZnO/cubic (Mg,Zn)O radial nanowire heterostructures ZnO/(Zn,Mg)O heterostructures have been realized that exhibit quantum confinement in quantum well structures. While the formation of planar semiconductor heterostructures is common for thin films, the synthesis of one-dimensional heterostructures is difficult. Axial heterostructures, in which the chemical modulation is imposed along the length of the nanowire axis, have been reported for a few systems, such as InAs/InP and Si/SiGe nanowires [85,86] ZnO/Zn1xMgxO quantum well nanorods have been grown as axial heterostructures as well [28]. Little work has addressed the synthesis of nanowire heterostructures in which the chemical modulation extends radially from the wire center [87]. In this section, we discuss the formation of ZnO-based one-dimensional radial heterostructure nanowires possessing a radial modulation in composition and structure. In particular, the nanowires consist of wurtzite ZnO cores surrounded by a rock-salt structured (Mg,Zn)O sheath. Despite the

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mismatch in lattice symmetry, the cubic (Mg,Zn)O sheath is found to be epitaxial on the wurtzite ZnO core. The growth experiments were performed using a conventional MBE system. The background base pressure of the growth chamber was
5  108 mbar. An ozone/oxygen mixture was used as the oxidizing source. The nitrogen-free plasma discharge ozone generator yielded an O3/O2 ratio on the order of 1–3%. No effort was made to separate the molecular oxygen from the ozone. The flux of Zn and Mg were provided by Knudsen effusion cells using high purity Zn metal (99.9999%) and Mg metal (99.995%) as the sources. Cation and O2/O3 partial pressures were determined via a nude ionization gauge that was placed at the substrate position prior to growth. The beam pressure of the O3/ O2 mixture was varied between 5  106 and 5  104 mbar, controlled by a leak valve between the ozone generator and the chamber. The Zn pressure was varied between 5  107 and 5  106 mbar and for Mg, the pressure ranged between 1  107 and 1  106 mbar. Si wafers with a native SiO2 layer terminating the surface were used as substrates. Ag served as the catalyst. The ZnO/(Mg,Zn)O nanowires were nucleated and grown on Si substrates coated with Ag for catalytic growth. Typical growth times for nanowires on the Ag-coated silicon were 2 h with growth temperatures ranging from Tg = 300–500 8C. Energy dispersive spectrometry spectra from single nanowires were collected by TEM (Philips 420 EM). Compositional line-scans, profiled across the nanowire, were measured by scanning transmission electron microscopy equipped with EDS. Under continuous Zn, Mg, and O3/O2 flux, the ZnO/(Mg,Zn)O nucleated uniformly on the catalyst-coated substrate. Fig. 12 shows a SEM image of ZnO/(Mg,Zn)O nanowires grown on a Si wafer that was coated with a nominally 2 nm thick layer of Ag. The Ag was deposited using e-beam evaporation. The image is from radial heterostructured ZnO/(Mg,Zn)O nanowires grown with a Zn pressure of 3  106 mbar, a Mg pressure of 4  106 mbar, and an O2/O3 pressure of 5  104 mbar. The growth temperature was 400 8C. Under these conditions, nanowires were observed only on the Ag. No growth occurred on the substrate surface that was devoid of Ag. The length of the ZnO/ (Mg,Zn)O nanowires is in excess of 2 mm.

Fig. 12. SEM image of radial heterostructured ZnO/(Mg,Zn)O nanowires grown on an Ag-coated Si wafer.

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Fig. 13. Radial heterostructured ZnO/(Mg,Zn)O nanowire. (a) TEM image of a nanowire showing a difference in brightness intensity between core and sheath regions. (b) Selected-area diffraction (SAD) from the nanowire that consists of two different crystal structures. (c) Dark field image taken from the diffraction spot of the wurtzite structure showing the ZnO core of nanowire. In (d), a DF image from the ring diffraction pattern of the rock salt structure showing the (Mg,Zn)O sheath.

The morphology and microstructure of ZnO/(Mg,Zn)O nanowires were analyzed by TEM. Fig. 13(a) shows a TEM image of a nanowire, displaying a difference in brightness intensity between the core and sheath regions. This image was obtained without the objective aperture in order to minimize any diffraction contrast. The contrast across the diameter of the nanowire is predominantly mass contrast, reflecting a difference in average atomic number (Z) of the core and sheath region. The darker core region contains more Zn, the lighter sheath region more Mg. Selected area diffraction taken from this single nanowire in Fig. 13(b) shows that the nanowire consists of two different crystal structures. The single crystal diffraction pattern result corresponds to the hexagonal wurtzite structure. Based on the mass contrast discussed earlier, this is a ZnO-rich core. The ring diffraction pattern that is observed corresponds to a cubic rock salt structure with d-spacing similar to that for MgO. The first circle is diffraction from the (Mg,Zn)O {1 1 1}, the second circle by the (Mg,Zn)O {2 0 0}, and the

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third circle by the (Mg,Zn)O {2 2 0} plane of the rock salt structure. A dark field (DF) image shown in Fig. 13(c), formed from the (0 0 0 2) diffraction spot of the wurtzite structure, yields intensity corresponding only to the core part of the nanowire. The discontinuous brightness along the nanowire length results from bending of the nanowire. Clearly, the center core of the nanorod is hexagonal wurtzite ZnO, with the sheath material not exhibiting this structure. Fig. 13(d) shows a DF image formed from the circle diffraction spots of the rock salt structure. The image includes the entire volume of the nanowire, indicating that it comes from the nanowire sheath. Therefore, the core of this wire is ZnO having the wurtzite structure. The surrounding sheath is (Mg,Zn)O having the rock salt structure. While the nanowire heterostructures considered here possess a rock-salt (Mg,Zn)O sheath, we have also observed similar cored nanowires grown with lower Mg flux that display a wurtzite (Zn,Mg)O epitaxial sheath. High-resolution transmission electron microscopy was also performed on the interface region between the ZnO core and (Mg,Zn)O sheath. The image shown in Fig. 14(a) indicates a radial heterostructured nanowire in which the interface between the sheath and core is epitaxial despite the discontinuity in crystal structure and symmetry. The lower-right region of the image is the core, while the upper-left region is the sheath. Growth of the core yields the rod axis along the c-axis direction of the hexagonal ZnO. For the (Mg,Zn)O sheath, the (2 0 0), (1 1 1), and (1 1 1) planes match those of the rock salt structure, according to the lattice spacing and angle between planes. Note that the lattice mismatch between the (Mg,Zn)O (2 0 0) planes and the ZnO (0 0 0 2) planes produces regularly-spaced interfacial edge dislocations. In order to further delineate the structure of these nanorod heterostructures, the compositional distribution of Mg and Zn in a single ZnO/(Mg,Zn)O nanowire was examined by energy-dispersive spectrometry using the electron beam in STEM. Fig. 14(b) plots the intensity of the Mg Ka1 and Zn Ka1 peaks across the nanowire. The compositional line-scan clearly shows a radial segregation of the Zn and Mg cation. Note that the spatially-varying spectra always include irradiation of the sheath region. Hence, the Mg intensity does not drop to zero when irradiating the center. However, in the core region, there is a clear drop in Mg intensity along with a corresponding increase in Zn intensity. This result unambiguously supports the radial segregation of Zn and Mg, forming a ZnO core and a (Mg,Zn)O sheath. The growth of the ZnO/(Mg,Zn)O nanowires is a catalyst-driven reaction similar to ZnO nanowire growth using catalysis-driven molecular beam epitaxy reported earlier [29]. Initially, the ZnO core is nucleated on the Ag catalyst with the reaction between the Zn and the oxygen source. It should be noted that uncored (Mg,Zn)O nanowires with the rock-salt structure throughout the rod were not observed for growth on Ag-coated Si substrates. At the selected growth conditions, Ag does not act as an effective catalyst for growth of cubic-cored (Mg,Zn)O nanowires. According to the phase diagram between ZnO and MgO, cubic (Mg,Zn)O can accommodate a maximum of 56 at.% ZnO at 1600 8C and maintain its rock salt structure with a lattice constant close to that of pure MgO [88]. In the case of ZnO, the solid solubility of Mg is limited to only 2 at.% maximum while keeping its hexagonal structure. In pulsed laser deposited thin film, a non-equilibrium solid solution can accommodate up to 36 at.% Mg solubility in ZnO [89]. However, the ZnO/(Mg,Zn)O heterostructured nanowire synthesis appears to follow equilibrium growth according to the phase diagram. In order to further elucidate the optical properties, photoluminescence of the cored nanowires was measured. This measurement was performed at room temperature with a He–Cd laser as the excitation source. The comparison of band edge photoluminescence between pure ZnO nanorods and cored ZnO/ (Mg,Zn)O nanorods shows only a 50 meV blue shift. As previously reported, a 50 meV blue shift in the near-band luminescent peak energy of cored ZnO compared to that of ZnO nanowires could

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Fig. 14. (a) High-resolution transmission electron microscopy of ZnO and (Mg,Zn)O interface; (Mg,Zn)O sheath in upperleft region, ZnO core in lower-right region. (b) Compositional line-scan across the nanowire probed by STEM EDS spectroscopy.

indicate approximately 2 at.% Mg solubility in the core-ZnO nanowire [90]. However, this radial heterostructured nanowire could also produce quantum confinement, as the bandgap of (Mg,Zn)O is much larger than that of ZnO. The diameter variation of the ZnO core is currently being explored to investigate the quantum confinement effect of the cored ZnO/(Mg,Zn)O nanowire structures. In summary, radial heterostructured nanowires of ZnO/(Mg,Zn)O were selectively grown on an Ag-coated Si wafer. Structural and compositional analyses of the nanowires clearly indicate that, under certain conditions, the core of nanowire is ZnO with the hexagonal wurtzite structure, while the nanowire sheath is (Mg,Zn)O with a cubic rock salt structure.

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5. Ferromagnetism in ZnO The manipulation of spin in semiconductors (so-called spintronics) presents a new paradigm for functionality in electronic materials. Ideas currently under development offer intriguing opportunities to pursue novel device concepts based on the discrimination and manipulation of spin distributions. Of the semiconducting materials, ZnO offers significant potential in providing charge, photonic, and spinbased functionality. ZnO is a direct, wide bandgap semiconductor with potential utility in UV photonics and transparent electronics. For spintronics, theoretical predictions suggest that room temperature carrier-mediated ferromagnetism should be possible in ZnO, albeit for p-type material. Unfortunately, the realization of p-type ZnO has proven difficult. Ab initio calculations do predict ferromagnetism in n-type ZnO doped with most transition metal (TM) ions, including Co and Cr, but predict no ferromagnetism for Mn-doped ZnO. This is consistent with experimental results where ferromagnetism is not observed in Mn-doped ZnO that is n-type due to group III impurities. However, we recently observed ferromagnetism in n-type Mn-implanted, Sn-doped ZnO crystals, where Sn (a group IV element) serves as a doubly ionized donor impurity. The Curie temperature is quite high, approaching 250 K. It is unclear what special role the Sn dopant plays, as compared to group III elements, in enabling ferromagnetism in the Mn-doped ZnO system. The Sn may simply provide carriers, albeit electrons, that effectively mediate the spin interactions. The Sn dopant might alternatively form complexes with Mn, resulting in both Mn2+ and Mn3+ sites that could yield a ferrimagnetic ordering. Nevertheless, ferromagnetism is observed in Mn-doped ZnO (co-doped with Sn). In recent years, the injection and manipulation of spin-polarized electrons in semiconductor materials has become the focus of intense research activity [91–93]. A functional spin offers an opportunity to develop new device concepts based on the discrimination and manipulation of spin distributions. Fundamental challenges related to the lifetime, manipulation, and detection of spinpolarized electrons in a semiconductor host must be addressed in order to realize spintronic technology based on semiconductors. Of particular importance is the realization of spin-polarized currents in a semiconductor matrix. Efforts directed at investigating the injection of electrons through a ferromagnetic metal/semiconductor junction have shown this structure to be ineffective. However, recent experiments indicate that injection of spin-polarized electrons from a ferromagnetic semiconductor into a non-ferromagnetic semiconductor is possible without detrimental interface scattering. Unfortunately, ferromagnetism in semiconductors is rare and poorly understood, with ferromagnetic transition temperatures well below room temperature for most known materials. Clearly, the discovery of ferromagnetism above 300 K in a semiconductor would prove enabling in realizing practical spinbased electronic devices. More generally, the study of any semiconducting material that supports spinpolarized electron distributions would be useful in understanding and developing spintronic concepts. 5.1. Ferromagnetism in semiconductors Magnetism in semiconductor materials has been studied for many years, and includes spin glass and anti-ferromagnetic behavior in Mn-doped II–VI compounds, as well as ferromagnetism in europium chalcogenides and Cr-based spinels [93–99]. In recent years, ferromagnetism in semiconductors has received renewed attention, partly due to interest in spintronic device concepts [91]. Due to transition metal solubility and technological interest, contemporary research has primarily focused on II–VI and III–V materials. Strong ferromagnetic interaction between localized spins has been observed in Mn-doped II–VI compounds with high carrier densities. In Pb1xySnyMnxTe (y > 0.6) possessing a hole concentration on the order of 1020–1021 cm3, ferromagnetism has been

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achieved [88]. In addition, it has been shown that free holes in low-dimensional structures of II–VI dilute magnetic semiconductors can induce ferromagnetic order [95]. For many semiconductor materials, the bulk solid solubility for magnetic and/or electronic dopants is not conducive with the coexistence of carriers and spins in high densities. However, the low solubility for transition metals in semiconductors can often be overcome via low temperature epitaxial growth. This approach has been used with Mn-doped GaAs [96–99] in achieving ferromagnetism with a transition temperature of 110 K, which is remarkable high compare to traditional dilute magnetic semiconductor material. Behavior indicative of ferromagnetism at temperatures above 300 K has recently been reported for GaN and chalcopyrite semiconductors doped with transition metals, illustrating the potential of achieving room temperature spintronics technologies [100,101]. Despite recent experimental success, a fundamental description of ferromagnetism in semiconductors remains incomplete. Recent theoretical treatments have yielded useful insight into fundamental mechanisms [102]. Dietl et al. [101,102] have applied Zener’s model for ferromagnetism, driven by exchange interaction between carriers and localized spins, to explain the ferromagnetic transition temperature in III–V and II–VI compound semiconductors. The theory assumes ferromagnetic correlations mediated by holes from shallow acceptors in a matrix of localized spins in a magnetically doped semiconductor. Specifically, Mn ions substituted on the group II or III site provide the local spin. In the case of III–V semiconductors, Mn also provides the acceptor dopant. High concentrations of holes are believed to mediate the ferromagnetic interactions among Mn ions. Direct exchange among Mn is anti-ferromagnetic as observed in fully compensated (Ga,Mn)As that is donordoped. In the case of electron doped or heavily Mn doped materials, no ferromagnetism is detected. Theoretical results suggest that carrier-mediated ferromagnetism in n-type material is relegated to low temperatures, if it occurs at all, while it is predicted at higher temperatures for p-type materials [102]. In p-type GaAs grown by low temperature MBE, Mn doping in the concentration range 0.04    0.06 results in ferromagnetism in GaAs. The model described has been reasonably successful in explaining the relatively high transition temperature observed for (Ga,Mn)As. Carrier-mediated ferromagnetism in semiconductors is dependent on the magnetic dopant concentration as well as on the carrier type and carrier density. As these systems can be envisioned as approaching a metal-insulator transition when carrier density is increased and ferromagnetism is observed, it is useful to consider the effect of localization on the onset of ferromagnetism. As carrier density is increased, the progression from localized states to itinerant electrons is gradual. On the metallic side of the transition, some electrons populate extended states while others reside at singly occupied impurity states. On crossing the metal-insulator boundary, the extended states become localized, although the localization radius gradually decreases from infinity. For interactions on a length scale smaller that the localization length, the electron wavefunction remains extended. In theory, holes in extended or weakly localized states could mediate the long-range interactions between localized spins. This suggests that for materials that are marginally semiconducting, such as in heavily doped semiconducting oxides, carrier-mediated ferromagnetic interactions may be possible. This theoretical treatment presents several interesting trends and predictions. For the materials considered in detail (semiconductors with zinc-blende structure), magnetic interactions are favored in hole-doped materials due to the interaction of Mn2+ ions with the valence band. This is consistent with previous calculations for the exchange interaction between Mn2+ ions in II–VI compounds [94,103] showing that the dominant contribution is from two-hole processes. This superexchange mechanism can be viewed as an indirect exchange interaction mediated by the anions, thus involving the valence band [104,105]. Note that valence band properties are primarily determined by anions in II–VI compounds. The model by Dietl et al. predicts that the transition temperature will scale with a reduction in the atomic mass of the constituent elements due to an increase in p–d hybridization and a

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reduction in spin-orbit coupling. Most importantly, the theory predicts a Tc greater than 300 K for ptype GaN and ZnO, with Tc dependent on the concentration of magnetic ions and holes. Recent experimental evidence for ferromagnetism in GaN appears to substantiate the theoretical arguments [104,105]. 5.2. Spin polarization in ZnO As pointed out earlier, the theory by Dietl predicts room temperature ferromagnetism for Mndoped p-type ZnO. In addition to Dietl’s prediction, ferromagnetism in magnetically doped ZnO has been theoretically investigated by ab initio calculations based on local density approximation [106]. Again, the results suggest that ferromagnetic ordering of Mn is favored when mediated by hole doping. However, for V, Cr, Fe, Co, and Ni dopants, ferromagnetic ordering in ZnO is predicted to occur without the need of additional charge carriers. Several groups have investigated the magnetic properties of TM-doped ZnO. In all of these studies, the ZnO material was n-type. The magnetic properties of Ni-doped ZnO thin films were reported [107]. For films doped with 3–25 at.% Ni, ferromagnetism was observed at 2 K. Above 30 K, superparamagnetic behavior was observed. Fukumura et al. have shown that epitaxial thin films of Mn-doped ZnO can be obtained by pulsed-laser deposition, with Mn substitution as high as 35% while maintaining the wurtzite structure [108]. This is well above the equilibrium solubility limit of
13%, and illustrates the utility of lowtemperature epitaxial growth in achieving metastable solubility in thin films. Codoping with Al resulted in n-type material with carrier concentration in excess of 1019 cm3. Large magnetoresistance was observed in the films, but no evidence for ferromagnetism was reported. However, Jung et al. recently reported ferromagnetism in Mn-doped ZnO epitaxial films, with a Curie temperature of 45 K [109]. The discrepancy appears to lie in differing film-growth conditions. Recently, researchers have utilized ion implantation to survey the magnetic properties of a number of transition metal dopants in various semiconducting oxide materials. Among the system investigated, the researchers have observed high temperature ferromagnetism in ZnO crystals that are implanted with transition metal dopants, including Co and Mn, co-doped with Sn. In the case of Mn co-doped with Sn, the Sn ions are provided as doubly ionized donors. This result differs from that reported for Mn-doped ZnO doped n-type with Al or Ga, and suggests that group IV dopants may behave differently than shallow group III donors in terms of interaction with magnetic dopants. If carrier-mediated mechanisms are responsible, one must explain why the behavior depends on the specific cation dopant species chosen (Sn versus Al,Ga). Insight into this issue may reside in the fact that doping via a multi-ionized impurity likely introduces relatively deep donor levels in the energy gap. Conduction from deep donors is often due to impurity band and/or hopping conduction, as opposed to conventional free electrons excited to the conduction band. Any carrier-mediated processes would be dependent on the relevant conduction mechanisms. This result appears to contradict the expectation that, in the absence of a shallow acceptor level, the dominant exchange mechanism is short-range superexchange which, for Mn2+ in ZnO, should favor anti-ferromagnetic ordering. The results will be discussed in detail in a later section. Despite the uncertainty in the mechanism, these results (high temperature ferromagnetism in Co- and Mn,Sn-doped ZnO) indicate a pathway for exploring spintronics in ZnO materials. Table 1 shows the valence and ionic radii for a number of dopant candidates. The ionic radius of ˚ ) is relative close to that for Zn (0.60 A ˚ ), suggesting moderate solid solubility without Mn2+ (0.66 A phase segregation. As such, the primary transition metal dopant of interest will be Mn. Chromium and cobalt presents the possibility of achieving ferromagnetism in ZnO via doping with magnetic ions for which the net superexchange coupling is ferromagnetic. Low temperature ferromagnetic behavior has

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Table 1 Valence and ionic radii for candidate dopant atoms Atom

Valence

˚) Ionic radius (A

Zn Sn Li Ag Mn Cr Fe Co V Ni Mg

+2 +4 +1 +1 +2 +3 +2 +2 +3 +2 +2

0.60 0.55 0.59 1.00 0.66 0.62 0.63 0.38 0.64 0.55 0.57

been observed in Cr-based spinel semiconductors. Theoretical results by Blinowski, Kacman, and Majewski predict that Cr doping in II–VI semiconductors should result in ferromagnetism [110]. Ab initio calculations based on the local density of states approximation specifically predicts ferromagnetism in Co- and Cr-doped ZnO without the need for additional doping [106]. Fig. 15 shows the magnetization versus field behavior at 10 K for Sn-doped ZnO samples implanted with 3 and 5 at.% Mn. Hysteretic behavior is clearly observed, consistent with ferromagnetism. At 10 K, the coercive field in the 3 at.% Mn-doped sample is 250 Oe. It must be noted that other possible explanations for hysteretic M versus H behavior that are remotely possible include superparamagnetism and spin-glass effects [111–114]. Magnetization measurements were also performed on Sn:ZnO crystals that were not subjected to the Mn implant. This was done to eliminate the possibility that spurious transition metal impurities might be responsible for the magnetic response. The Sn-doped ZnO crystals exhibit no magnetic hysteresis, showing that the Mn doping is responsible for the behavior. To track the hysteretic behavior in the implanted samples as a function of temperature, both field-cooled and zero field-cooled magnetization measurements were performed from 4.2 to 300 K. By taking the difference between these two quantities, the para- and diamagnetic

Fig. 15. M vs. H curve for Mn implanted ZnO:Sn single crystals showing ferromagnetism in (a) 3 at % Mn and (b) 5 at.% Mn implantation doses.

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contributions to the magnetization can be subtracted, leaving only a measure of the hysteretic ferromagnetic regime. When assigning the origin of ferromagnetism in transition metal doped semiconductors, one must carefully consider the possibility that secondary phase formation is responsible. High-resolution transmission electron microscopy is the most direct means of exploring this issue. These activities are currently being pursued for the Mn-implanted ZnO samples. Nevertheless, one can also consider what known ferromagnetic impurity phases are possible [111–114]. First, metallic Mn is anti-ferromagnetic, with a Ne´ el temperature of 100 K. In addition, nearly all of the possible Mn-based binary and ternary oxide candidates are anti-ferromagnetic. The exception to this is Mn3O4, which is ferromagnetic with a Curie temperature of 46 K in thin films [115]. X-ray diffraction measurements on the implanted samples show no evidence for Mn-O phases, although it is recognized that diffraction is limited in detecting secondary phases that may represent a fraction of a percent of total volume in the implanted region. However, even if this phase were present, it could not account for the remarkably high ferromagnetic transition temperature of
250 K observed for the Mn-implanted ZnO:Sn crystals. Note also that the Mn concentrations used in this study were well below the solid solubility limit of Mn in ZnO. It should also be noted that increasing the Mn content from 3 to 5 at.% resulted in a significant decrease in the relative magnetization response as is seen in Fig. 15. This provides strong evidence, albeit indirect, that the magnetization is not due to any precipitating secondary phase. If the formation of a secondary Mn-related phase was responsible for the ferromagnetic behavior, an increase in Mn concentration would presumably increase the secondary phase volume fraction and related magnetization signature. Instead, the opposite behavior is observed. One can postulate as to the mechanism by which Sn doping yields ferromagnetic interactions among the Mn ions. First, the Sn ions may simply provide carriers or bound donor states with extended wavefunctions that mediate interactions among the Mn ions. The Sn dopants may alternatively form complexes with the Mn ions, yielding a distribution of Mn3+ sites. In this case, the mixture of Mn2+/ Mn3+ sites could yield ferrimagnetic ordering through superexchange interaction. With this, the magnetic moment/Mn ion should increase with Sn doping even if the measured conductivity changes very little. It should be noted that very little is known about the behavior of Sn in ZnO. Identifying the location of the energy state associated with Sn in ZnO would be most useful in the interpretation of magnetic properties when co-doped with Mn. One could also look for shifts in the Sn or Mn dopant state location as the other dopant is added. An additional means of elucidating the mechanism for ferromagnetism in ZnO co-doped with a deep donor is to consider an alternative transition metal. For this, Cr is particularly attractive. First, Cr provides an opportunity to realize ferromagnetism in ZnO via ions for which the net superexchange coupling is ferromagnetic. As with Mn, Cr itself is anti-ferromagnetic, thus eliminating any role of Cr precipitates in yielding spurious ferromagnetism. Third, theory predicts that Cr-doped n-type ZnO should be ferromagnetic. As such, experiments similar to those described above for ZnO co-doped with Mn and Sn should be performed with Cr replacing Mn. Ferromagnetism in Co-doped semiconducting oxides has been reported for several materials, including ZnO and TiO2. The observation that a doubly ionized donor (Sn) leads to ferromagnetic interactions among Mn atoms in ZnO raises the question as to whether deep states are, in general, effective in mediating ferromagnetism. Of particular interest are the theoretical predictions that ferromagnetism above room temperature should be possible in Mn-doped ZnO that is p-type. Given the difficulty in realizing a shallow acceptor in ZnO, an obvious experiment, given the results for the Mn co-doped with a deep donor (Sn), is to co-doped Mn with a deep acceptor. Attractive deep acceptors for addressing this question are Cu and As. Copper doping introduces an acceptor level with an energy
0.17 eV below ˚ ) on the O (1.38 A ˚ ) site, the conduction band. There is a large ionic radii mismatch for As (2.22 A

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suggesting limited solid solubility for these anions [116]. Nevertheless, p–n junction-like behavior has been reported between an n-type ZnO films on GaAs subjected to annealing [117]. In this case, a ptype layer was reportedly produced at the GaAs/ZnO interface. Further understanding of group V substitution in ZnO requires the study of doped materials that are free from complications of reactive substrates or interfacial layers. The primary interest is to investigate whether the deep acceptor states from Cu or As mediate ferromagnetic interactions in TM-doped ZnO. For practical application in spintronic devices, the Curie temperature should be well above room temperature [118,119]. As discussed earlier, theory suggests that the Curie temperature will tend to increase with decreasing cation mass. In addition, there is phenomenological evidence that Tc increases with increasing gap. The observed trend is for Tc to increase with increasing bandgap. While the equilibrium solid solubility of Mg in wurtzite ZnO is only
2%, epitaxial Zn1xMgxO thin films have been realized with x as large as 0.35 [120]. A key requirement in understanding ferromagnetism in transition metal doped semiconductors, including ZnO, is to delineate whether the magnetism originates from substitutional dopants on cation sites, or from the formation of a secondary phase that is ferromagnetic. The importance of this issue cannot be understated. The concept of spintronics based on ferromagnetic semiconductors assumes that the spin polarization exists in the distribution of semiconductor carriers. Localized magnetic precipitates might be of interest in nanomagnetics, but is of little utility for semiconductor-based spintronics. The question of precipitates versus carrier-mediated ferromagnetism is complex, and is a central topic of discussion for other semiconducting oxides that exhibit ferromagnetism, in particular the Co-doped TiO2 system [121,122]. Several issues must be addressed in order to gain insight into the possible role of secondary phase precipitates in the magnetic properties of transition metal doped semiconductors, specifically for ZnO films. First, one should identify all candidate magnetic phases possible from the assemblage of elements. The coincidence of Tc with a known candidate secondary ferromagnetic phase indicates a likely source of at least part of the magnetic signature. The UF group has been investigating a number of TM-doped semiconducting oxide materials in order to understand ferromagnetism in semiconductors, and to identify promising candidates for spintronics. As an example, we recently investigated the magnetic properties of Mn-doped Cu2O. Cu2O is a p-type semiconductor with a bandgap of 2.0 eV and a hole mobility of 100 cm2 V1 s1. The growth of Mndoped Cu2O films was achieved via pulsed laser deposition from a Mn-doped Cu-O target. These epitaxial Cu2O films doped with Mn are clearly ferromagnetic with a Tc of
48 K. However, some question as to the origin of the ferromagnetism arises when it is recognized that the measured Curie temperature close to that of Mn3O4, which has a Tc of
46 K. Obviously, the simplest explanation for ferromagnetic behavior in this material is Mn3O4 precipitates. However, there is no evidence for this phase in X-ray diffraction. Additional work is needed in order to delineate the possible formation of Mn3O4 in the Cu2O films. Nevertheless, knowledge of the candidate ferromagnetic secondary phases is invaluable in sorting out the ferromagnetic response. Fortuitously, for Mn-doped ZnO:Sn, the only ferromagnetic secondary phase candidate is the Mn3O4 spinel. None of the other possible secondary phases involving combinations of Zn, Mn, O, and Sn yield a known ferromagnetic material. The high temperature ferromagnetism in Mn-implanted ZnO:Sn crystals cannot be attributed to Mn3O4 as the Tc is much higher in the Mn-doped ZnO than for the Mn3O4 phase. In addition to TEM, one can also use X-ray diffraction to search for secondary phases within the films. Despite the obvious sensitivity limitations involved in detecting impurity phases that represent only 1–5% vol.% of a thin-film sample, we have been able to detect nanoscale precipitates in transition metal implanted samples. This is demonstrated specifically for another transition metal doped semiconducting oxide that we have investigated, namely Co-implanted ZnO. It has been reported that epitaxial (Zn1xCox)O (x = 0.05–0.25) exhibit high temperature ferromagnetism with a Tc greater

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than 300 K [123]. The ferromagnetic behavior is assigned to substitutional Co on the Zn site. Monte Carlo simulations on the indirect exchange interaction of Co-doped ZnO also predicts ferromagnetism in these materials [124]. SQUID magnetometry measurements of Co-doped ZnO clearly indicate ferromagnetism. However, a careful examination of the X-ray diffraction u–2u scan along the surface normal, clearly indicates the presence of Co precipitates. A Co (1 1 0) peak is clearly evident. From the width of the peak, the size of the cobalt precipitates can be estimated to be approximately 3.6 nm. Clearly, the possibility of precipitate formation must be carefully considered for each dopant investigated. It should be noted that the presence of magnetic precipitates may mask an underlying carriermediated ferromagnetism due to substitutional doping. In this case, a more direct means of measuring the spin properties of the semiconductor carrier population is needed. Much of the research activity being presently pursued in spintronics is directed at demonstrating device structures (spin-LED, spinFET) that differentiates spin polarization distribution among the electron/hole population in the semiconductor. Future work would employ these materials in such device structures. Some techniques, such as magnetic circular dichroism [125] and magneto-optical Kerr effect [126], also yield information regarding global spin distributions. The 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. This can be accomplished under ambient conditions via optical pumping with appropriately polarized laser light [91]. However, ultimate device integration will require electrical spin injection. Electrical spin injection can be accomplished either by injecting from a spin-polarized source or by spin-filtering unpolarized carriers at the interface. Despite persistent efforts by many groups, spin injection from a conventional ferromagnetic metal into a semiconductor has proven highly inefficient. By sharp contrast, efficient spin injection has recently been successfully demonstrated in all-semiconductor tunnel diode structures either by using a spin-polarized dilute magnetic semiconductor (DMS) as the injector or by using a paramagnetic semiconductor under high magnetic field as a spin filter. These experiments, coupled with the earlier discoveries of ferromagnetic ordering in (Ga,Mn)As at an unprecedented temperature of 120 K and the long spin coherence length and time in a variety of semiconductors offer the possibility of a breakthrough in spintronics. In order to implement these applications, however, DMS with an ordering temperature over 300 K must be synthesized.

6. Ferromagnetism in ZnO nanorods An SEM is shown of a single rod is shown in Fig. 16(top). The growth time was
2 h at 400 8C. The typical length of the resultant nanorods was
2 mm, with typical diameters in the range of 15– 30 nm. It is also possible to grow cored nanorods with ZnMgO composition varying across the rod diameter (Fig. 16 bottom). The samples were subsequently implanted with Mn+ and Co+ ions at a fixed energy of 250 keV and doses of 1–5  1016 cm2, while the samples were held at
350 8C to avoid ˚ , with peak transition metal amorphization. The projected range for both types of ion is
1500 A concentrations corresponding to roughly 1–5 at.%. The samples were then annealed at 700 8C for 5 min under flowing N2 to promote migration of the transition metal ions into substitutional positions. Even at this highest dose condition, the nanorods are stable to the implant/anneal cycle. A comparison of micrographs before and after this cycle did not show any observable change in the nanorods. Fig. 17 shows the magnetization versus field behavior at 300 K for a Mn-implanted (5  1016 cm2 dose) nanorod sample after the 700 8C, 5 min anneal. Hysteretic behavior is clearly present, with a coercive field at 100 K of 100 Oe. The possible explanations for this data include

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Fig. 16. TEM image of single-crystal ZnO nanorod(top) and TEM micrographs showing cored (Zn1xMgx)O nanorods having Zn-rich phase surrounded by another (Zn1xMgx)O phase (bottom).

ferromagnetism, superparamagnetism or spin-glass effects [91,127]. The magnetization of unimplanted, annealed nanorods was as much as three orders of magnitude lower than in the implanted, annealed samples, demonstrating that the transition metals are responsible for the observed magnetic properties. The calculated moment from the saturation magnetization was
2.2 Bohr magnetons per Mn ion, showing that a significant fraction of the implanted Mn is contributing to the magnetization. None of the potential second phases can account for the observed magnetic behavior. Thus, for Mn-implanted nanorods, it does not appear that secondary phases plays a significant role in the magnetic properties. Indeed, the magnetization results for the Mn-implanted nanorods are very similar to those reported previously for Mn-implanted, bulk n-type ZnO single crystals, in which highresolution X-ray diffraction did not observe any secondary phases. We note also that solubility limits

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Fig. 17. Magnetization vs. field at 300 K for Mn-implanted ZnO nanorod.

for Mn in ZnO are between 13% for equilibrium conditions to 35% for incorporation by pulsed laser deposition. These values are well beyond the concentrations employed here. By sharp contrast, in the case of Co implanted ZnO, macroscopic Co precipitates are ferromagnetic with Curie temperatures up to 1382 K when large enough to exhibit bulk properties. Bulk, single-crystals of ZnO implanted with Co under the same conditions as employed here showed the ˚ , well below the presence of (1 1 0)-oriented hexagonal Co nanocrystals with average diameter
35 A ˚ superparamagnetic critical diameter of Co near room temperature of
50 A. This is consistent with the magnetization data for the Co-implanted nanorods, which do not show ferromagnetism at room temperature. Thus, in the case of nanorods examined here, it appears the magnetic properties most likely result from the presence of Co nanocrystals. The results are similar to those obtained by implantation of these same species into bulk, single-crystal ZnO. One can envision implanting Mn into p–n junction nanorods to investigate the possibilities of spin-polarized UV light emission or using Mndoped nanorods as magnetic storage elements. The most effective measurement of the quality of the oxide-based ferromagnetic materials will be in the operation of device structures, such as spin-FETs or photo-induced ferromagnets. While promising advances in spin-injection efficiency from metal contacts into semiconductors have been made recently (using injection from ohmic contacts or ballistic point contacts, tunnel injection or hot electron injection), in general there are difficulties in reproducibility. The use of ferromagnetic semiconductors as the injection source in device structures should allow a more direct measurement of the efficiency and length scale of spin transport. There are two test structures based on ZnO, that will demonstrate whether novel aspects of ferromagnetism that can be exploited in devices. The most common structure is shown in Fig. 18 and was suggested by Sato and KatayamaYoshida (Mater. Res. Soc. Symp. Proc., vol. 666, F4.6.1, 2001) as an analog of the spin-FET. The structure takes advantage of the fact that (Zn,Mn)O can be grown as an anti-ferromagnetic spin-glass

Fig. 18. Schematic of ZnO-based spin-fet.

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insulator, while hole- and Mn-co-doped ZnO can be a half-metallic ferromagnet. In the structure of Fig. 18, the application of negative gate bias brings holes into the (Zn,Mn)O and converts it to the halfmetallic ferromagnetic state, using ferromagnetic (Zn,Co)O as the source and drain contact material, it should be possible to have a 100% spin-polarized electron flow in the (Zn,Mn)O channel. The device can be fabricated by growing the source/drain materials on top of the (Zn,Mn)O and then etching away in the gate region for selective growth of the gate oxide and deposition of the gate metal. It may be possible to realize such a device using ferromagnetic ZnO nanorods.

7. Properties of Mn- and Co-implanted bulk ZnO The properties of transition metal impurities, together with studies of the nature of shallow donors and shallow acceptors, are of particular interest due to the predicted room temperature ferromagnetism of heavily TM doped ZnO. TM dopants can be incorporated into ZnO both during growth (from vapor phase, by solid state recrystallization, by hydrothermal method or e.g. from PbF2 solutions [128–132] and by ion implantation. The work done in early 80 s was mostly driven by the general interest to the behavior of TMs in various semiconductors stimulated by the success of semi-insulating GaAs and InP. At that period, no group succeeded in observation of room temperature magnetism. The results of electron spin resonance (ESR) and photoluminescence spectra studies of Fe, Ni, Mn, Co doped ZnO were rather confusing. The most thoroughly studied were the properties of Fe doped samples. ESR of these samples showed clear indications of the presence of the Fe3+ ionized donors in semi-insulating material with the Fermi level deep in the bandgap. The signal was photosensitive and could be quenched by light with photon energies peaked near 2.75 eV which was explained by optically stimulated transition from Fe3+ to Fe2+ state by capturing an electron from the valence band. Detailed studies placed the optical ionization transition threshold from the Fe3+ state to the conduction band at about 1.4 eV suggesting a strong lattice coupling of the deep donor state of Fe in ZnO [131]. A somewhat similar photosensitive ESR signal was also observed for Ni doped ZnO [128]. At the same time, ESR signals in the Mn doped and Co doped ZnO were found not to be photosensitive. However, for the Co doped material a strong optical band at 0.55–0.66 mm was detected in absorption and photocurrent spectra and attributed to an optical transition from the ground state to the excited states of the Co ion with subsequent thermal excitation of the electrons into the conduction band. Photoluminescence spectra studies in Fe, Ni and Co doped ZnO crystals revealed the presence of red luminescence bands at respectively 0.63, 0.68 and 0.6850 mm that were attributed to transitions involving excited states of the TM ions [128]. In more recent papers, the electrical properties of Mn and Co were studied by means of current– voltage (I–V), capacitance–voltage (C–V) and isothermal capacitance transients (ICTS) in ZnO crystals solution grown from PbF2 [132]. No indication of the presence of deep states that could be related to Mn and Co were observed in ICTS spectra which could point to respective centers being deeper than the Schottky barrier height (about 0.7 eV). Some indirect manifestations of the influence of deep states due to Mn and Co could be observed in C–V and especially I–V measurements [131]. ZnFeO crystals co-doped with Cu (but not simple ZnFeO crystals) grown by solid state recrystallization showed a Curie temperature of 500 K and the authors suspected that some role in that was played by Cu acceptors [131]. For Co and Mn implanted and 700 8C annealed samples, the room temperature van der Pauw characteristics were virtually the same as for the control sample. The temperature dependence of the conductivity showed the same activation energy of 13 meV. The optical transmission spectrum of the virgin ZnO crystal showed no strong absorption bands anywhere but at the bandedge region. The first

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band observed with the threshold near 0.75 eV and the second with the threshold near 1.4 eV are very similar for all samples. In addition to those the Mn and Co implanted ZnO crystals show strong absorption bands with optical threshold near 2 eV and all samples show defect absorption bands with the threshold near 2.3–2.5 eV. The TM-related bands near 2 eV show the apparent threshold energies of about 2 eV for Mn and of about 1.9 eV for Co. The energy of the yellow-green defect absorption band is close to 2.5 eV for both the Mn and the Co implanted ZnO crystals and is somewhat lower, close to 2.35 eV for the heavily proton implanted sample. It is clear that no p-type layer is formed at the surface of Mn, Co implanted and annealed samples and thus the near-room temperature ferromagnetism observed in such implanted layers is not promoted by free holes as assumed in the model proposed by Dietl [101,102]. This correlates very well with the results obtained on high-Curie-temperature bulk ZnOFe(Cu) crystals that were shown to be n-type [131]. Therefore, the mechanism of room temperature ferromagnetism in ZnO doped with TM is unconventional as, e.g. for GaN doped with TM. The next question is whether or not any features in the absorption or MCL spectra can be associated with Mn and Co. The most likely candidates in that sense are the absorption bands at 1.9 and 2 eV and the red MCL bands at 1.84 and 1.89 eV. The absorption and the MCL bands seem to be closely related in the sense that their energies are reasonably close and the relative shift of the bands’ energies when changing Mn to Co is similar although not quite equal in both bands. The positions of these MCL bands for different TM ions are very close to each other which brings to mind the possibility that the bands are due to intracenter transitions between the states of TM ions. However, we do not see any fine structure confirming such an interpretation either in absorption or in MCL and this question requires further study. For example, as mentioned earlier, a similar band near 2 eV observed in photocurrent and absorption spectra of Co doped samples was attributed to intra-center transitions from the ground to excited states with subsequent thermal excitation of electrons into the conduction band [127]. The intensity of absorption in the red band is considerably higher and the MCL peak is considerably broader for the Mn implanted sample compared to the Co implanted sample. Since the concentrations of implanted Mn and Co are virtually the same this is most likely related to the higher density of substitutional Mn achieved after implantation and annealing compared to Co due to higher solubility of Mn at 700 8C. In summary, we have shown that the Mn and Co implanted and 700 8C annealed regions of ZnO crystals remain n-type and even show a higher electron concentration than the virgin samples. Therefore, the high Curie temperature of about 250 K measured on similarly implanted and annealed crystals is not due to hole-mediated spin alignment but rather due to some other mechanism. The implanted crystals show optical absorption and MCL bands near 2 eV that can be attributed to TM ions on the grounds that such bands are not observed in the virgin or proton implanted samples. The position of the bands slightly shifts for various TM ions (e.g. corresponding MCL bands are at 1.84 eV for Co, 1.89 eV for Mn (respective absorption bands are at 1.9 and 2 eV). The origin of the bands is not quite clear at the moment. We somewhat tentatively attribute them to internal transitions between the crystal-field-split states of substitutional TM ions. If so one has to conclude that the concentration of such ions is considerably higher after implantation and annealing for the Mn implanted samples compared to the Co implanted samples. This could be due to a higher solubility of the Mn in ZnO at 700 8C. This higher concentration of deep donor defects in the Mn implanted samples stimulates formation of acceptor-type Zn vacancies giving rise to the violet MCL band near 3 eV not observed in the Co implanted or proton implanted samples. In addition to these directly or indirectly TM-related bands, we also observed in Mn and Co implanted samples absorption bands near 0.75, 1.4, 2.5 eV that are due to radiation damage defects as follows from comparison with the results obtained on heavily proton implanted samples.

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8. Fabrication approaches to ZnO nanowire devices The preparation of the nanowires has been described in detail earlier in this review. In brief, discontinuous Au droplets were used as the catalyst for ZnO nanowires growth and formed by ˚ ) on Al2O3 wafers at 1100 8C. ZnO nanowires were annealing e-beam evaporated Au thin films (
20 A 8 deposited by MBE with a base pressure of 5  10 mbar using high purity (99.9999%) Zn metal and an O3/O2 plasma discharge as the source chemicals. The Zn pressure was varied between 4  106 and 2  107 mbar, while the beam pressure of the O3/O2 mixture was varied between 5  106 and 5  104 mbarr The growth time was
2 h at 600 8C. The typical length of the resultant nanowires was
5 mm, with typical diameters in the range of 30–100 nm. Selected area diffraction patterns showed the nanowires to be single-crystal. They were released from the substrate by sonication in ethanol and then transferred to SiO2-coated Si substrates. E-beam lithography was used to pattern sputtered Al/Pt/Au electrodes contacting both ends of a single nanowire. The contacts were sintered at 300 8C for 1 min to obtain a low contact resistance. The separation of the electrodes was
3 mm. Ebeam evaporated Pt/Au was used as the gate metallization by patterning a 1 mm wide strip orthogonal to the nanowire. We found that a simple and effective method for removing the nanorods from the initial growth substrate was by sonification in a solvent such as ethanol. Fig. 19 shows scanning electron microscopy pictures of the ZnO nanorods after growth on the Si substrate (left) and after 5 min sonification in ethanol (right). The acoustic energy supplied to the solvent is enough to dislodge a large fraction of the nanorods and disperse them into the solution. We found that >90% of the nanorods could be harvested in this manner. Fig. 20 shows a higher magnification picture of the Si substrate after removal of most of the nanorods. Transfer of the nanorods from the ethanol solution to a new substrate is achieved by dispersing the solution onto the new substrate, followed by evaporation of the ethanol. This is shown in the SEM micrographs of Fig. 21. The obvious advantage of this approach is simplicity but the main drawback is the random nature of where the nanorods are placed. Our approach is to initially prepare an SiO2-

Fig. 19. SEM micrographs of ZnO nanorods before (left) and after (right) sonification.

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Fig. 20. Highly magnified view of ZnO substrate surface after sonication.

coated Si wafer and etch alignment marks into the SiO2. Once the nanorods are on this new substrate, a mask design for the particular device being fabricated is written using software on an e-beam writer and then transferred lithographically so that the ends of the nanorod are covered by Ohmic contact pads. A schematic of this process is shown in Fig. 22. We have used this method to fabricate a range of different nanorod devices. An example of a pattern for measuring the transport properties of a single nanorod is shown in the SEM pictures of

Fig. 21. Nanowires transferred by sonication and suspension in solvent to new Si substrate.

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Fig. 22. Prototype device fabrication sequence.

Fig. 23. In the temperature range from 25 to 150 8C, the resistivity of nanorods treated in H2 at 400 8C prior to measurement showed an activation energy of 0.089  0.02 eV and was insensitive to the ambient used (C2H4, N2O, O2 or 10% H2 in N2). By sharp contrast, the conductivity of nanorods not treated in H2 was sensitive to trace concentrations of gases in the measurement ambient even at room temperature, demonstrating their potential as gas sensors. In summary, a simple method for growth of ZnO nanorods and their removal and transfer to another substrate for subsequent device fabrication has been described. The approach provides a high yield method for studying the electrical transport properties of nanorods and their functionalization into diodes, photdetectors and transistors. Much more sophisticated approaches will be needed for making nanorod device arrays.

9. Gas sensing ZnO has a long history for use as a gas sensing material [132–134]. There is a strong interest in the development of wide bandgap semiconductor gas sensors for applications including detection of combustion gases for fuel leak detection in spacecraft, automobiles and aircraft, fire detectors, exhaust diagnosis and emissions from industrial processes. Of particular interest are methods for detecting ethylene (C2H4), which offers problems because of its strong double bonds and hence the difficulty in dissociating it at modest temperatures. Wide bandgap semiconductors such as ZnO are capable of operating at much higher temperatures than more conventional semiconductors such as Si. Diode or field-effect transistor structures are sensitive to gases such as hydrogen and hydrocarbons. Ideal sensors have the ability to discriminate between different gases and arrays that contain different metal oxides (e.g. SnO2, ZnO, CuO, WO3) on the same chip can be used to obtain this result [134]. The gas sensing mechanism suggested include the desorption of adsorbed surface oxygen and grain boundaries in poly-ZnO [11–26], exchange of charges between adsorbed gas species and the ZnO surface leading to changes in depletion depth [133] and changes in surface or grain boundary conduction by gas adsorption/desorption [134].

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Fig. 23. SEM micrographs of structure for transport measurements of ZnO nanowires (top) and close-up of central region (bottom).

Pt/ZnO Schottky diodes fabricated on bulk ZnO show changes in forward current of 0.3 mA at a forward bias of 0.5 Vor alternatively a change of 50 mV bias at a fixed forward current of 8 mA when 5 ppm of H2 is introduced into a N2 ambient at 25 8C, as shown in Fig. 24. The rectifying current– voltage characteristic shows a non-reversible collapse to Ohmic behavior when as little as 50 ppm of H2 is present in the N2 ambient. At higher temperatures, the recovery is thermally activated with an activation energy of
0.25 eV. This suggests that introduction of hydrogen shallow donors into the ZnO is a contributor to the change in current of the diodes.

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Fig. 24. I–V characteristics at 25 8C of Pt/ZnO diode measured in N2 or H2-containing ambients.

Fig. 25 shows the I–V characteristics at 50 and 150 8C of the Pt/ZnO diode both in pure N2 and in ambients containing various concentrations of C2H4. At a given forward or reverse bias, the current increases upon introduction of the C2H4, through a lowering of the effective barrier height. One of the main mechanisms is once again the catalytic decomposition of the C2H4 on the Pt metallization, followed by diffusion to the underlying interface with the ZnO. In conventional semiconductor gas sensors, the hydrogen forms an interfacial dipole layer that can collapse the Schottky barrier and produce more Ohmic-like behavior for the Pt contact. The recovery of the rectifying nature of the Pt contact was many orders of magnitude longer than for Pt/GaN or Pt/SiC diodes measured under the same conditions in the same chamber. For measurements over the temperature range 50–150 8C, the activation energy for recovery of the rectification of the contact was estimated from the change in forward current at a fixed bias of 1.5 V. This was thermally activated through a relation of the type IF = Io exp(Ea/kT) with a value for Ea of
0.22 eV, comparable for the value of 0.17 eV obtained for the diffusivity of atomic deuterium in plasma exposed bulk ZnO. This indicates suggests that at least some part of the change in current upon hydrogen gas exposure is due to in-diffusion of hydrogen shallow donors that increase the effective doping density in the near-surface region and reduce the effective barrier height.

Fig. 25. I–V characteristics at 150 8C of Pt/ZnO diodes measured in different gas ambients.

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10. Transport in ZnO nanowires ZnO is attractive for forming various types of nanorods, nanowires and nanotubes for observation of quantum effects [135–147]. The initial reports show a pronounced sensitivity of the nanowire conductivity to ultraviolet illumination and the presence of oxygen in the measurement ambient. The photoresponse of the nanowires shows that the potential barrier between the ZnO and the contact is lowered by illumination with above bandgap light. The preparation of the nanorods proceeded as follows; discontinuous Au droplets were used as the catalyst for ZnO nanorods growth and formed by annealing e-beam evaporated Au thin films ˚ ) on p-Si (100) wafers at 1100 8C. ZnO nanorods were deposited by MBE with a base pressure (
20 A of 5  108 mbar using high purity (99.9999%) Zn metal and an O3/O2 plasma discharge as the source chemicals. The Zn pressure was varied between 4  106 and 2  107 mbar, while the beam pressure of the O3/O2 mixture was varied between 5  106 and 5  104 mbar The growth time was
2 h at 400
600 8C. The typical length of the resultant nanorods was 2
10 mm, with typical diameters in the range of 30–150 nm. Selected area diffraction patterns showed the nanorods to be single-crystal. They were released from the substrate by sonication in ethanol and then transferred to SiO2-coated Si substrates. E-beam lithography was used to pattern sputtered Al/Pt/Au electrodes contacting both ends of a single nanorod. The separation of the electrodes was
3.7 mm. A scanning electron micrograph of the completed device is shown in Fig. 26. Au wires were bonded to the contact pad for current–voltage measurements performed over the range 25–150 8C in a range of different ambients (C2H4, N2O, O2 or 10% H2 in N2). The as-grown nanorods were resistive and passed very small currents (<1010 A). To make them more conducting, we annealed them in hydrogen gas at 400C. Fig. 27 shows the I–V characteristics as a function of measurement temperature for these conducting nanorods. It is expected that the hydrogen increases the n-type doping in the ZnO. The hydrogen-annealed nanorods are indeed more conducting, with currents of
3  108 A at 0.5 V bias voltage of either polarity. The current is thermally activated of the form I = Io exp(–Ea/kT),where Ea is the activation energy, k is Boltzmann’s constant

Fig. 26. SEM micrograph of single ZnO nanorod bridging two Al/Pt/Au Ohmic contact pads.

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Fig. 27. I–V characteristics of single nanorod measured at different substrate temperatures.

and T is the absolute measurement temperature. The nanorods showed a strong, reversible response to above bandgap (366 nm) ultraviolet light, with the UV-induced current being approximately a factor of 7 larger than the dark current at a given voltage. The intensity of the photocurrent was independent of the illumination time for a fixed modulation frequency. Previous reports have suggested that photocurrents excited in ZnO nanowires by above bandgap light are surface-related and do not represent true bulk conduction. Fig. 28 shows the resistivity the nanorods obtained from the I–V characteristics, in Arrhenius form. The apparent activation energy derived from this plot is 0.089  0.02 eV. This does not correspond to any of the known donor dopant or native defect ionization energies in ZnO and it is possible that the conduction is again surface-related, as in the UV photo-response measurements. In the hydrogen-annealed nanorods, the currents were insensitive to the measurement ambient, as shown in Fig. 29. One might expect the current to be dependent on this ambient if the conduction were truly dominated by the surface. Various gas sensing mechanism in ZnO-based gas sensors have been suggested, including the desorption of adsorbed surface oxygen and grain boundaries in poly-ZnO, exchange of charges between adsorbed gas species and the ZnO surface leading to changes in depletion depth and changes in surface or grain boundary conduction by gas adsorption/desorption. In all of these cases, the mechanism is surface-related. By sharp contrast to the data for the hydrogenannealed nanorods, the unannealed samples showed a strong sensitivity to measurement ambient, with increased currents in the presence of hydrogen. Clearly the control of both bulk and surface quality is a key area for practical applications of the nanorods in sensing of gases, chemicals or UV light.

Fig. 28. Arrhenius plot of single nanorod resistivity obtained from I–V measurements in air.

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Fig. 29. I–V characteristics of nanorod at 25 8C as a function of the measurement ambient.

11. ZnO nanowire Schottky diodes E-beam lithography was used to pattern sputtered Al/Pt/Au electrodes contacting both ends of a single nanowire. The separation of the electrodes was
3 mm. E-beam evaporated Pt/Au was used as the gate metallization by patterning a 1 mm wide strip orthogonal to the nanowire. A scanning electron micrograph of the completed device is shown in Fig. 30. Au wires were bonded to the contact pads for current–voltage measurements at 25 8C. In some cases, the diodes were illuminated with above band edge ultraviolet light during the measurement. Fig. 31 shows I–V characteristics from the Pt/ZnO nanowire diode, measured both in the dark and with UV illumination. While the measurement in the dark shows rectifying behavior, the UV illumination produces an Ohmic characteristic. This is similar to the result reported by Keem et al. who showed that the above bandgap light lowered the potential barrier between the Schottky contact and the ZnO nanowire. In their case however, the photoresponse recovery times were very long (>104 s), which was ascribed to the dominance of surface states, whereas our recovery times were

Fig. 30. SEM micrograph of ZnO nanowire Schottky diode.

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Fig. 31. I–V characteristics from the ZnO Schottky nanowire diodes both in the dark and with UV illumination.

limited by turn-off of the Hg lamp and suggest the conduction in our nanowires is dominated by bulk transport. Fig. 32 shows the reverse current characteristic in more detail. The reverse current was still only 1.5  1010 A at 10 V bias, corresponding to a current density of 2.35 A cm2. If one defines the reverse breakdown voltage as that bias at which the reverse current density is 1 mA cm2, this value is
2 V for the nanowire diodes. The forward I–V characteristics in the dark (shown in Fig. 33) were fit to the relation for thermionic emission over a barrier     ef eV JF ¼ A  T 2 exp  b exp nkT kT where J is the current density, A* Richardson’s constant for n-GaN, T the absolute temperature, e the electronic charge, fb the barrier height, k Boltzmann’s constant, n the ideality factor and V the applied voltage. We could not drive our nanowires to the saturation current condition because of fears the high current density would create self-heating problems and therefore were unable to extract the barrier height. However, previous measurements of the barrier height of Pt on ZnO have yielded values of 0.75 eV reported by Mead using internal photoemission yield spectroscopy to obtain the barrier height

Fig. 32. Reverse current of ZnO nanowire diode measured in the dark.

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Fig. 33. Forward current of ZnO nanowire diode measured in the dark.

of the laterally homogeneous portions of the contacts and more recently 0.61 eV on bulk ZnO. From the measured data, we derived an ideality factor of n = 1.1. This shows that there is little recombination in our nanowires. The on-state resistance, RON,was 1.65 V cm2 (the resistivity of the nanowire was 22.6 V cm) for an individual nanowire device,with an on-off current ratio of
6 at 0.15/5 V bias. These values can be further improved by optimizing the growth process of the nanowires and the subsequent contact metallization processes.

12. ZnO nanowire MOSFET There is great current interest in fabrication of ZnO channel thin-film transistors for applications in transparent flat panel displays [138–147]. Other transparent conducting oxides such as Sn-doped In2O3, Al-doped ZnO, and Sb-doped SnO2 have been widely applied as transparent electrodes for liquid crystal displays, organic light-emitting diodes, and solar cells [147]. E-beam lithography was used to pattern sputtered Al/Pt/Au electrodes contacting both ends of a single nanowire. The separation of the electrodes was
7 mm. Au wires were bonded to the contact pad for current–voltage measurements. (Ce,Tb)MgAl11O19 with thickness 50 nm was selected as the gate dielectric as it exhibits a large bandgap sufficient to yield a positive band offset with respect to ZnO. The top gate electrode was e-beam deposited Al/Pt/Au. Fig. 34 shows a schematic of the ZnO nanowire MOSFET. A scanning electron micrograph of the completed device is shown in Figs. 35 and 36. Note that less that 50% of the nanowire channel is covered by the gate metal. Current–voltage

Fig. 34. Schematic of ZnO nanowire depletion-mode FET.

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Fig. 35. SEM micrograph of fabricated FET.

characteristics were performed at room temperature using an Agilent 4155A Semiconductor Parameter Analyzer. Fig. 37 shows the drain current–drain voltage (ID–VDS) characteristics (top) and the transfer characteristics (bottom) measured at room temperature in the dark. The modulation of the channel conductance indicates that the operation of the device is an n-channel depletion mode. The gate leakage current is low and the nanowire MOSFETs exhibit excellent saturation and pinch-off characteristics, indicating that the entire channel region under the gate metal can be depleted of electrons. The threshold voltage is
3 V with a maximum transconductance of
3 mS/mm. The on/ off current ratio at VG of 0–2.5 V and VD of 10 V was of order 25. The field-effect mobility mFE can be determined from the transconductance using the relation IDS = (W/L) mFECOX(VGS–VT)VDS where W is the channel width, L is the channel, COX the gate oxide capacitance and VT is the threshold voltage. The extracted mobility was
0.3 cm2/V s, which is comparable to that reported previously for thinfilm ZnO enhancement mode MOSFETs [147]. The carrier concentration in the channel is estimated to be
1016 cm3. The nanowire showed strong photoresponse to UV illumination. Fig. 38 shows the shows the drain current–drain voltage (ID–VDS) characteristics (top) and the transfer characteristics (bottom) measured at room temperature under 366 nm illumination. The drain–source currents increase by approximately a factor of 5 and the maximum transconductance increases to 5 mS/mm. Part of the increase in drain–source current reflects photoconductivity in the portion of the wire uncovered by the gate metal. The on-off current ratio at VG of 0–3 V and VD of 10 V increases to
125. In this case, the extracted mobility was still
3 cm2/V s, but the carrier density in the channel was
5  1016 cm3. The photoresponse was rapid, being limited by the switching characteristics of the UV lamp, which may suggest the photocurrent is bulk-related and there is not a strong contribution from surface states that produce long recovery times. In conclusion, ZnO nanowire depletion mode MOSFETs show excellent pinch-off and saturation characteristics and a strong UV photoresponse. These devices look promising for transparent transistor applications requiring low leakage current and indicate that the nanowires can be grown and transferred to another substrate without major degradation of their electrical transport properties.

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Fig. 36. SEM micrographs of ZnO MOSFET structure (top) and close-up of the actual nanorod and its contacts (bottom).

13. UV nanowire photodetectors Recent reports have shown the sensitivity of ZnO nanowires to the presence of oxygen in the measurement ambient and to ultraviolet illumination [20–23]. In the latter case, above bandgap illumination was found to change the current–voltage characteristics of ZnO nanowires grown by thermal evaporation of ball-milled powders between two Au electrodes from rectifying to Ohmic. By contrast, there was no change in effective built-in potential barrier between the ZnO nanowires and the contacts for below bandgap illumination. The slow photoresponse of the nanowires was suggested to

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Fig. 37. IDS–VDS (top) and transfer characteristics (bottom) of ZnO nanowire FET at room temperature in the dark.

originate in the presence of surface states which trapped electrons with release time constants from milliseconds to hours [22,23]. Fig. 39 shows the I–V characteristics of the nanowires in the dark and under illumination from 366 nm light. The conductivity is greatly increased as a result of the illumination, as evidenced by the higher current. No effect was observed for illumination with below bandgap light. Transport measurements to be reported elsewhere show that the ideality factor of Pt Schottky diodes formed on the nanowires exhibit an ideality factor of 1.1,which suggests there is little recombination occurring in the nanowire. Note also the excellent Ohmicity of the contacts to the nanowire, even at low bias. On blanket films of n-type ZnO with carrier concentration in the 1016 cm3 range we obtained contact resistance of 3–5  105 Ohm cm2 for these contacts. In the case of ZnO nanowires made by thermal evaporation, the I–V characteristics were rectifying in the dark and only became Ohmic during above bandgap illumination. The conductivity of the nanowire during illumination with 366 nm light was 0.2 Ohm cm. Fig. 40 shows the photoresponse of the single ZnO nanowire at a bias of 0.25 V under pulsed illumination from a 366 nm wavelength Hg lamp. The photoresponse is much faster than that reported for ZnO nanowires grown by thermal evaporation from ball-milled ZnO powders and likely is due to the reduced influence of the surface states seen in that material. The generally quoted mechanism for the photoconduction is creation of holes by the illumination that discharge the negatively charged oxygen ions on the nanowire surface, with detrapping of electrons and transit to the electrodes. The recombination times in high quality ZnO measured from time-resolved photoluminescence are short,

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Fig. 38. IDS–VDS (top) and transfer characteristics (bottom) of ZnO nanowire FET at room temperature measured with illumination from UV (366 nm) light.

on the order of tens of ps, while the photoresponse measures the electron trapping time. There is also a direct correlation reported between the photoluminescence lifetime and the defect density in both bulk and epitaxial ZnO. In our nanowires, the electron trapping times are on the order of tens of seconds and these trapping effects are only a small fraction of the total photoresponse recovery characteristic. Note

Fig. 39. I–V characteristics from single ZnO nanowire measured at 25 8C in the dark or under illumination from a 366 nmHg lamp.

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Fig. 40. Time dependence of photocurrent at as the 366 nm light source is modulated.

Fig. 41. Time dependence of photocurrent in the nanowire as either a 254 nm or 366 nm light source is modulated.

also the fairly constant peak photocurrent as the lamp is switched on, showing that that any traps present have discharged in the time frame of the measurement. Fig. 41 shows the photoresponse from the nanowires during pulsed illumination from either 254 or 366 nm light. The lower peak photocurrent in the former case may be related to the more efficient absorption near the surface of the nanowires. Once again, we see an absence of the very long time constants for recovery seen in nanowires prepared by thermal evaporation. In conclusion, single ZnO nanowires with excellent photocurrent response and current–voltage characteristics have been prepared by site selective MBE. The I–V characteristics are linear even when measured in the dark and the photoconduction appears predominantly to originate in bulk conduction processes with only a minor surface trapping component. These devices look very promising for UV detection.

14. Conclusions and summary Major advances have been made by many groups around the world recently in ZnO nanowire synthesis and also in conventional epitaxial thing film growth where there have been demonstrations of p-type ZnO, improved Ohmic contacts, high resolution dry etch pattern transfer processes and implant isolation processes and in realizing room temperature ferromagnetism in ZnO. There are still many areas that need additional work, including:

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1. Higher p-type doping levels in nanowires. This will require even better control of the background ntype conductivity in the material that arises from native defects and impurities such as hydrogen. 2. Realization of high quality p–n nanowire junctions, with good breakdown characteristics. These are the building blocks for devices such as nanoscale UV LEDs or lasers. 3. Improved Schottky contacts with higher barrier height than reported to date. This will need a better understanding of surface cleaning processes and metal schemes that do not disrupt the ZnO surface during their deposition. In particular, it is necessary to understand that nature of the defects present on ZnO surfaces that may be pinning the Fermi level. 4. P-type Ohmic contacts with lower specific contact resistance. A desirable goal is to achieve a contact resistance of 105 Ohm cm2 or less for ZnO-based light emitters so that self-heating in the contact region does not degrade their reliability. 5. Clear demonstrations of spin-polarized carrier distributions in Mn- or Co-doped ZnO nanorods such as the presence of the anomalous Hall effect or spin polarized light emission. It is likely that many of these areas will see significant progress in the near future as the promise of ZnO nanorods for low cost transparent electronics and UV light emission becomes more widely known.

Acknowledgments The work at UF is partially supported by AFOSR grant under grant number F49620-03-1-0370 (T. Steiner), NSF (CTS-0301178, monitored by Dr. M. Burka and Dr. D. Senich), by NASA Kennedy Space Center Grant NAG 10-316 monitored by Mr. Daniel E. Fitch, ONR (N00014-98-1-02-04, H.B. Dietrich), and NSF DMR 0400416. The authors would like to thank C.R. Abernathy, B.P. Gila, K. Pruessner, K.N. Seibin, M. Kaufman, W. Sigmund, M. Overberg and M.F. Chisholm.

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