C. Sudakar
†‡
, Shubra Singh
‡*
, M.S. Ramachandra Rao
‡*
, G. Lawes
†
†
Department of Physics and Astronomy, Wayne State University, Detroit, MI 48201
‡
Department of Physics, Indian Institute of Technology Madras, Chennai, India 600036
*Nano Functional Materials Technology Centre,
Indian Institute of Technology Madras, Chennai, India 600036
The burgeoning interest in nanoscale metal oxides arises from the recognition that the materials properties of these systems depend strongly on morphology, allowing the development of new or enhanced characteristics in geometrically restricted samples[1-3].
Finite size effects can produce significant changes in a number of intrinsic properties in systems having reduced length scales, including the electronic band gap [4], the magnetic coercivity [5], and elastic modulus [6], to name only a few of the characteristics that are strongly sensitive to sample geometry. Simultaneously, the large surface to volume ratio in nanomaterials, realized most dramatically

2 in nanoparticles, can also substantially affect the electronic [7], magnetic [8], optical [9], and elastic properties [10] of these systems. Because of their relatively larger surface to volume ratio, the defect concentration in metal oxide nanostructures is generally higher than that found in bulk systems. These defects can have a profound effect on the physical properties of nanomaterials, so it is crucially important that they be fully considered when characterizing metal oxide nanostructures. There are a number of thorough and accessible reviews on defects in particular metal oxide systems, including ZnO [11-14], TiO
2
[15], and CuO [16] to name a few, along with more comprehensive reports [17-18]. Rather than attempting to provide a general overview of how defects modify the physical properties of oxides, this particular report is more narrowly focused on briefly presenting the entirely new properties and characteristics that can emerge in metal oxide nano-systems due the presence of defects.
The chapter is structured as follows. We begin with a very short review of the basic types of defects in metal oxides. Rather than consider the multitude of possible defect structures, we sharply

3 limit our discussion to point defects. We will specifically focus on oxygen defect vacancies (V
O
), metal ion vacancies (V
M
), and metal interstitials (M
I
) since these are generally the most important and widely studied intrinsic point defects in metal oxides [17-18]. The bulk of the review will center on a discussion of the novel electrical, optical, and magnetic properties that can arise in defect-rich metal oxide systems. We will focus uniquely on the new physical behaviour that emerges due to the presence of defects and do not consider in any depth the rather more widely studied problem of understanding the role of defects in perturbing the existing properties of oxides.
We conclude with a short discussion of how these defect-induced properties can be used to integrate new functionalities into metal oxide nanostructures.
We broadly limit the scope of our discussion to point defects.
As the emphasis of this review is to consider defects in metal oxide nanostructures, we further restrict ourselves to discussing only intrinsic defects, and only very briefly touch on dopant ions as point

4 defects. Within these constraints, a large number point defects can be considered: oxygen vacancies (V
O
), metal vacancies (V
M
), oxygen interstitials (O
I
), metal interstitials (M
I
), and anti-site defects
(M
O
or O
M
). In a large number of cases, the most stable and/or physically important defects are V
O
, V
M
, and M
I
[17-18] so we focus primarily on these specific examples. The interactions among different types of defects can play an important role in determining the physical properties of metal oxides. For example, because Sn is multivalent, Sn
I
point defects can readily form in SnO
2
, which, in turn, supports the formation of V
O
defects leading to n-type conductivity in defect-rich SnO
2
films [19]. Rather remarkably, the presence of these defects alone in metal oxide nanomaterials is sufficient to produce new physical properties that are not simply perturbations of the intrinsic characteristics of the defect-free parent oxide. Heuristically, these emergent properties can, in general, be understood to arise from interactions among the point defects leading to collective behaviour. Certain metal oxide nanostructures are typically able to support a relatively high concentration of V
O
defects, with oxygen non-stoichiometry reaching several percent near the surfaces of na-

5 noscale systems [20]. Because this relatively large defect concentration, defect-defect correlations can affect the response of the system
[20-22].
3.2.1 Defect Structures in Metal Oxide Nanostructures
The concentration and distribution of defects determine a number of properties of crystalline solids. In high quality crystals the concentration of defects can be extremely small, leading to considerable experimental challenges in accurately determining this defect concentration [23]. Crystalline solids contain a number of different types of structural defects. Vacancies defects develop due to the absence of atoms in some lattice sites while interstitials arise from extra atoms occupying the space between the atoms in the lattice [18,
23]. Vacancies and interstitial atoms are point defects as these imperfections are limited to one unit cell and lead to deviations from the crystalline order only in the immediate vicinity of the defect. In addition to point defects, line and plane defects are very often found in real crystal systems [18, 23]. Line defects are dislocations that are characterized by displacements in the crystal structure along specific

6 directions [18, 23]. Examples of plane defects comprise stacking faults, grain boundaries, and internal and external surfaces. A schematic diagram of possible defects in crystalline metal oxide nanostructures is shown in Fig. 1. These defects can profoundly modify the physical properties of materials, including electrical, optical, and magnetic response as we will discuss in the following.
Many metal oxides, including ZnO, TiO
2
, SnO
2
and In
2
O
3
, exhibit marked deviations from stoichiometry under specific annealing conditions including thermal annealing in vacuum [20, 24-25] and under a finite metal vapor pressure [26-27]. A relatively small degree of off-stoichiometry can be supported by the inclusion of point defects, including V
O
and V
M
[23], and annealing can also promote the formation of intersitials and antisite defects. At small concentrations (0.1 to 1 at%) these point defects are typically assumed to be randomly distributed throughout the lattice [23]. At higher defect concentrations, a number of different types of defects can develop, including multiple charge state defects and pairs or complexes of defects [17-18, 23]. However, as we will discuss, many of the properties of defect-rich metal oxides can be understood

7 by considering segregated (though possibly interacting) point defects, so we center our discussion to this class of structures.
3.2.2 Imaging Defects in Metal Oxide Nanostructures
At relatively small concentrations, the distribution of point defects in oxides is determined solely by entropy considerations and consists of randomly distributed defects [23]. At higher defect concentrations, enthalpy begins to affect the distribution, which leads to the formation of new structures including defect clusters, superlattice ordering and extended defects, shear planes, and discrete intermediate phases [23]. Defect-rich TiO
2
(see Fig. 2) is an example of an oxide in which extended defects, including planar defects, are formed by the accumulation and elimination of point defects, such as oxygen vacancies in vacuum annealed samples, along specific crystallographic planes [28]. These vacancies are eliminated by the formation of shear planes in the crystal, which in turn produces a fault in the cation sublattice [28]. Because of this interplay between point defects and extended defect structures, in some metal oxide systems

8 it can be difficult to completely separate the two, as we illustrate in the case of TiO
2
and In
2
O
3
in Fig. 2.
We first consider ideal, isolated point defects in ZnO as a representative metal oxide system. In wurtzite ZnO the possible point defects are: oxygen and zinc vacancies (V
O
, V
Zn
), interstitials
(O i
, Zn i
), and antisite defects (O
Zn
and Zn
O
). V
O
and Zn i
were most generally considered to be the defects responsible for modifying the electric and magnetic properties of the system [29-34]. However recent work [11] suggests that these defects exhibit high formation energies under equilibrium conditions. Zn interstitials (Zn i
) are shallow donors and fast diffusers with a low migration barrier, 0.57 eV, and are therefore not stable at room temperature [11, 35]. V
Zn
, which has a low formation energy, is a deep acceptor, so is able to act as a compensating center in n-type ZnO, and may be relevant for the green luminescence observed in ZnO [11, 35]. O i
has a high energy and acts as a deep acceptor at the octahedral site O i
-1
(oct) in n-type
ZnO [11, 36]. The antisite defects (Zn
O
and O
Zn
) have very high formation energies and are unlikely under equilibrium conditions
[11, 35-36].

9
The defect structure in metal oxide nanostructures can be imaged using high resolution transmission electron microscopy
(HRTEM). Comparing real-space images of air annealed and vacuum annealed In
2
O
3
thin films clearly demonstrates the effects of point defects on the nanostructure [20]. Fig. 2a shows a HRTEM image of an air annealed In
2
O
3
nanoparticle, showing a well-ordered lattice with no obvious defects. Vacuum annealing this sample introduces oxygen vacancies, as well as other point defects, as shown in Figs 2b and 2c. The agglomeration of point defects leads to a 2–3 nm thick surface disordered layer, as shown in Fig. 2b. Additional point defects, both V
O
and In
I
, can be seen in the bulk of the sample.
These additional point defects are highlighted by arrows in Fig. 2c.
The insets to Fig. 2c show simulated HRTEM images for an oxygen vacancy defect (i) and an oxygen vacancy with and adjacent cluster of two In (ii).
Similar defect-induced structures can be observed in nanoscale TiO
2
(Fig. 2d-g) [25]. Air-annealed TiO
2
thin films consisting particles ranging from 300 to 500 nm show good crystalline order with few defects, as illustrated in Fig. 2d. Conversely, vacuum-

10 annealed films show numerous crystallographic twin boundaries, with individual grains often containing several parallel twins (Fig.
2f). Additionally, these particles exhibit a highly disordered surface phase of few nanometers thick [25]. These are common [28] microstructural features in non-stoichiometric TiO
2
and are intimately related to the formation of shear structures discussed above. The twinning produced in the rutile subcell structure is parallel to (0 1 1), which is the common twinning plane for TiO
2
[28, 37-38]. This twinning is not observed in TiO
2
thin films formed from nanoparticles [25]. However, a substantial non-stoichiometric disordered phase develops at the surface of the nanoparticles (Fig. 2e), which suggests that these planar defects may be more readily diffuse to the surface under thermal annealing in films comprised of smaller particles.
3.2.3 Stability of Intrinsic Point Defects in Metal Oxide
Nanostructures
In order for these point defects to have any meaningful effect on the properties of the metal oxide nanostructures in the context of

11 device applications, they should be stable under ambient conditions.
The determination of whether these defects are stable depends strongly on the details of the specific metal oxide being considered.
In ZnO for example, the V
O
defects are believed to become stable in presence of transition metal dopants such as Co [24]. In this particular study, oxygen defects were introduced in thin film ZnO samples by annealing at high temperatures (600 o C) and low pressures (~10 -6 torr). Raman spectral modes related to –Zn-O-Co- local disordered vibrations in the stoichiometric (or defect-poor) Co:ZnO films disappear after vacuum annealing as the oxygen vacant sites are localized near Co site (–Zn-V
O
-Co-). In a number of systems, however, oxygen vacancy defects may not be stable. Studies on oxygen deficient TiO
2
found that the concentration of V
O
defects decreases rapidly under ambient conditions [25], although these defects can apparently be stabilized by transition metal doping [32, 39-40].
Conversely, in In
2
O
3
nanostructured films, the oxygen vacancy defects are stable for a timescale of years under the same conditions
[41]. Because of this sensitive dependence of the persistence of point defects on the specific compound being considered, it is im-

12 portant to properly characterize the stability of these defects in a particular metal oxide when determining the effects such defects may have on the physical properties of the material.
Metal oxides exhibit a range of electrical transport properties, from metallic to insulating to superconducting [37, 42-43]. The introduction of point defects generically affects all types of electrical transport, through mechanisms ranging from increased scattering in metallic systems to the introduction of additional charge carriers in insulators. We are particularly interested in exploring systems in which the inclusion of point defects qualitatively changes the electrical transport properties. We therefore limit our discussion to considering the onset of metallic or quasi-metallic conductivity induced by point defects in systems where the undoped metal oxide is insulating or semiconducting.
3.3.1 Point Defects and Charge Carriers

13
In general terms, point defects in metal oxide nanostructures act like charge centers [44], which can lead to very high electrical conductivities. Experimentally, a number of metal oxide systems that are insulating when prepared as perfectly stoichiometric samples develop good electrical conductivity with the introduction of point defects [14, 45]. In the simplest models, this increase in conductivity requires shallow donors near the conduction band [46], or acceptors near the valence band [47]. Many, though by no means all, defect-rich oxide materials are found to exhibit n-type conductivity, pointing to an abundance of excess electrons associated with the defects. V
O
sites, which normally act as electron donors, are a possible point defect in all metal oxide systems and it is often believed that the conducting properties in these materials arise from oxygen vacancy defects [48]. While we see that oxygen vacancy defects do play a crucial role in mediating electrical conductivity in many metal oxide materials, other types of point defects can also have a significant effect on transport in defect-rich samples.
ZnO represents one of the most intensely investigated metal oxide system [49-51] in the past decade. While oxygen vacancy de-

14 fects, possibly together with Zn interstitials, had been widely considered to be the origin of the n-type conductivity in this system [12,
48], recent studies suggest that other point defects may be more relevant for determining the electrical transport properties [11]. These investigations find that V
O
sites are deep donors, falling approximately 1 eV below the conduction band, and are thus unlikely to introduce any significant of n-type charge carriers. The concentration of Zn interstitials is found to have a high formation energy in n-type materials, making Zn
I
defects unlikely as the source for increased conductivity in defect-rich ZnO [11]. Upon considering all native point defects in ZnO, the authors conclude that none of these is likely to produce the observed n-type conductivity and instead propose that the charge carriers arise from the accidental inclusion of substitutional hydrogen, H
O
, which can act as a shallow donor [52].
Along a similar line, experimental studies on the conductivity of
ZnO films prepared by pulsed laser deposition provide evidence that nitrogen inclusions may play an important role in the development of n-type conductivity in ZnO [46], with other work pointing to the importance of hydrogen donors [53].

15
Indium oxide is another widely studied electronic material, but there still remain a number of unanswered questions concerning the fundamental transport properties in this system [54]. In
2
O
3
can exhibit a high degree of non-stoichiometry and shows extremely good n-type dopability [55-58]. It has been suggested that In
2
O
3
is an anion-deficient n-type conductor, but that the small oxygen vacancy defect population, corresponding to approximately 1% of the anions, limits the electron concentration [59]. Experimentally, it is found that oxygen deficient In
2
O
3
is highly compensated, with the ratio of n-type free charge carriers to oxygen vacancy defect sites being approximately 1:5, rather than the 2:1 one would expect is each oxygen vacancy contributes 2 electrons [20, 59]. Recent density functional theory calculations on defect-rich In
2
O
3
find that oxygen vacancies, rather than In interstitials, are the likely source of ntype conductivity [60]. Furthermore, both indium vacancies and oxygen interstitials are identified as possible charge compensation sites.
The dramatic effects of point defects on the electrical transport properties of transition metal oxides are clearly demon-

16 strated by the remarkable change in conductivity of In
2
O
3
thin films upon vacuum annealing, illustrated in Fig. 3. As-prepared In
2
O
3
thin films, which were crystallized by annealing in air and are presumed to be close to stoichiometry, are highly resistive and show insulating behaviour below room temperature. On vacuum annealing, which introduces oxygen vacancies and may also produce other types of point defects, the films develop n-type conductivity with a carrier concentration on the order of n=10
20
cm
-3
. Concomitant with this increase in carrier concentration, the room-temperature resistivity of the films drops by three to four orders of magnitude and the samples exhibit metallic conductivity to low temperatures, with a small upturn in resistivity below ~80 K.
Rather remarkably, defect-rich In
2
O
3
films remain optically transparent, despite the high conductivity. The optical band gap is found to increase from approximately 3.3 eV to 3.6 eV on vacuum annealing, which can be attributed to the Burstein Moss shift, but there is practically no change in the optical transmission, which remains above 80% for visible wavelengths [41]. Similar conducting and optically transparent features are observed in SnO
2
samples,

17 where it is argued that the high oxygen vacancy defect concentration required for producing conductivity is stabilized by the presence of multivalent Sn interstitials [19]. These same density functional studies suggest that the donor electrons are not heavily compensated due to the paucity of acceptor defects (V
Sn
and O
I
). Furthermore, it is found that these donors do not have direct optical transitions in visible wavelengths, so do not directly affect the optical transparency.
Since In has a fixed formal valence of +3, the same mechanism is not likely to apply for In
2
O
3
, but the result on SnO
2
highlights the importance of interstitials in stabilizing oxygen vacancy defects.
3.3.2 Defects and P-Type Conductivity
A number of defect-rich metal oxide systems exhibit p type rather than n type conductivity [37, 42, 47]. First principle calculations on Cu
2
O find that the lowest energy defects are Cu vacancy point defects, V
Cu
, and a point defect complex consisting of a Cu interstitial, Cu
I
, located between two V
Cu
defects [61]. The V
Cu
defects are found to produce de-localized holes near the top of the valence band, leading to p-type conductivity. Measurements on

18 intentionally undoped Cu
2
O thin films find a p-type carrier concentration on the order of 10 15 cm -3 , resulting in a resistivity of approximately 150 [47]. P-type conductivity can also develop in the
Mott insulator NiO. Careful measurements have established that V
Ni sites are the dominant point defect for determining the electrical properties of NiO samples, rather than O
I
sites [62]. It is estimated that the V
Ni
defect concentration can reach 10 16 cm -3 in nanostructured samples, leading to a 6-8 order of magnitude increase in conductivity over undoped NiO single crystal samples [63] .
3.3.3 Defects and Conduction Mechanisms
In addition to understanding the origin of charge carriers in defect-rich metal oxide nanostructures, it is also important to consider the mechanisms for electrical conduction. Depending on the details of the electronic structure, a number of different effects can be relevant for electronic transport. Careful investigations on the low temperature resistivity and Hall resistance of oxygen deficient
TiO
2
, having oxygen vacancy defect concentrations in the range from 4x10 18 cm -3 to 5x10 19 cm -3 , find that the low temperature

19 transport is consistent with hopping conductivity for high and low
V
O
concentrations (n[V
O
]) [64]. However, the transport falls in the intermediate range between hopping and metallic conduction for
8x10
18
cm
-3
<n[V
O
]<2x10
19
cm
-3
, where the donor separation is estimated to be five times the effective Bohr radius of the donor electron [64]. TiO
2
does not develop metallic behaviour because increasing the defect concentration leads to the development of planar defects [64]. Hopping conduction has also been established as the origin of electrical transport in defect rich NiO films [63, 65]. Frequency dependent resistivity studies find that the transport can be well-modeled by the correlated barrier hopping model, which supposes that holes hop from Ni 3+ sites to Ni 2+ sites with a barrier height that depends on the separation between defects. A fit to the data yields a separation of 1.9 eV between the ground state of the defect and the valence band [65]
Band conduction can also be observed in defect-rich metal oxides. ZnO films prepared in an oxygen deficient environment were found to be highly conducting, with a band-like mechanism for conduction having an activation energy of ~1 meV [46]. More stoi-

20 chiometric samples were found to exhibit Arrhenius conductivity, associated with the thermioinic emission of band electrons from grain boundaries at higher temperatures and thermally assisted hopping at lower temperatures [46]. There have also been a number of theoretical studies on the impurity band structure in defect-rich ZnO, as there are proposals that the ferromagnetism in this system (discussed in more detail in Section V) may arise from spin split impurity bands [66]. Ab initio density functional calculations suggest that
V
Zn
point defects should produce metallic behaviour in defect-rich
ZnO, while O
I
defects give rise to a semiconducting electronic structure [67].
Surface effects are also expected to affect the electrical transport properties in metal oxides, which is particularly relevant for nanostructured materials have a high surface area to volume ratio. Theoretical studies on SnO
2
find no evidence for defect-induced states in the gap [68], while subsurface oxygen defect vacancies in
TiO
2
can lead to states falling 0.7 eV below the conduction band edge [69]. A surface conduction layer, presumably arising from defect states, is found in ZnO; the conductivity of this layer is reduced

21 on exposure to oxygen [12]. This surface conduction layer provides an additional channel for electronic transport. In
2
O
3
films and nanostructures can develop a chemical depletion layer near the surface, corresponding to a higher oxygen content at the interface [70].
This system also has a high density of electronic surface states, which produces relatively large band bending.
3.3.4 Plasmon Response in Defect-Rich Oxide Nanostructures
One of the more striking realizations of the collective response of an electron gas is the phenomenon of plasma oscillations.
These plasmons are excited at the plasma frequency p
, given by p
2 2 /m eff
with n the carrier density and m eff
the effective mass. This plasma frequency is typically large for most metals, with p
~11 eV for bulk plasmons in Al [71], and depends strongly on the charge carrier concentration n. Because plasmons reflect collective behavior of the charge carriers, they represent emergent response in insulating metal oxides driven by point defects, which is completely absent in the parent system. To illustrate the clear development of this electronic collective behavior in defect-rich metal

22 oxide nanostructures, we consider the optical response of asprepared (defect poor) and vacuum annealed (V
O
defect rich) In
2
O
3 thin films, as plotted in Fig. 4. The as-prepared sample is insulating and, as expected for transparent materials, has negligible absorption at energies well below the bandgap. Conversely, the vacuum annealed In
2
O
3
sample has a high concentration of V
O
defects leading to a n-type charge carrier concentration of roughly 10 20 cm -3 , as measured by the Hall effect [41]. This high concentration of charge carriers in the defect-rich sample leads to qualitatively different behavior in the low energy optical properties. We observe a clear plasmon peak in the absorption falling at 0.53 eV. While the magnitude of the absorbance associated with this peak falls well below the bandgap absorbance, this plasmon resonance represents the emergence of a distinct electronic response that is absent in the defectpoor parent metal oxide structure.
As discussed in the previous section, charge carriers in oxides arise from a number of different sources including interstitial

23 metal ion impurities, substitutional doping ions, and oxygen vacancies. Oxygen vacancies present in the lattice can act as divalent electron donors, with these defects normally acting as shallow donors.
Within the metal oxide systems, point defect ionization occurs in a similar manner to that found in doped semiconductors. Although the scattering of charge carriers in defect-rich oxide systems arises primarily from ionized impurity scattering, the majority of the intrinsic optical phenomena in these materials arise from ionized defects occupying energy states lying in the band gap, at least for wide band gap oxides. When such systems are obtained in the low dimensional form the optical properties of these materials are often modified due to the increase in surface energy and surface defect states, which can lead to a number of interesting possibilities for applications.
The incorporation of metal oxide nanostructures into electrooptical devices relies on their ability to efficiently emit or absorb light; these application prospects are heavily influenced by the energy band structure and lattice dynamics of the system [72-74]. This change in optical response in defect-rich oxides is reflected in the altered band to band transitions and absorption energies [75]. Moreo-

24 ver, oxide nanostructures have lower threshold lasing energies due to quantum effects that increase the density of states near band edges
[76-77]. For a number of varied electro-optical applications it is necessary for all charge carriers, both electrons and holes, to be confined [78]. One-dimensional wide band gap nanostructures are the best candidates for this class of applications due to their remarkable physical and chemical properties. However, it is also crucial to understand how these optical properties may be affected by the almost unavoidable incorporation of point defects in these nanostructured materials.
3.4.1 Photoluminescence from Point Defects in Oxide Nanostructures
Nanocrystalline zinc oxide (nano-ZnO) is a wide band gap semiconductor that is particularly promising for a number of optoelectronic properties, including ultraviolet (UV) light emitting diodes, UV laser diodes, and UV photodetectors because of its very high excitonic binding energy (60 meV) [79] compared to GaN (25 meV) and relative ease of bandgap engineering [80]. Nanostruc-

25 tured ZnO possesses a remarkable photoluminescence (PL) spectrum. The optical response of ZnO changes significantly on the introduction of point defects, rather by doping or the incorporation of intrinsic defects. The PL spectrum of nano-ZnO consists mainly of two emission peaks, one in the ultraviolet, falling near 385 nm, which is ascribed to near-band-edge emission [81], with the other peak located in the visible region, occurring in the green around 500 nm [82-85]. The origin of green luminescence band is still not well understood; this has been attributed to the presence of a variety of different impurities and defects present in the ZnO lattice. ZnO exhibits luminescence defect centers such as oxygen vacancies (located at 50 and 190 meV below the conduction band edge), zinc interstitials (located at 2.5 eV below the conduction band edge) [86] as well as various other native defects [82-85]. However it is interesting to note that the intensity of the green emission can be controlled in a systematic manner by oxidation and reduction [87]. Nanostructures such as nano-islands of ZnO show PL emission whose origin can be explained on the basis of zinc vacancies (V
Zn
) complex defects [88]. The intensity of PL emission from samples containing

26 such islands is much smaller than that from normal thin films due to a smaller area being covered by the islands. It has also been found that upon bio-molecule attachment, nano-ZnO powders exhibit further induced changes in peak intensities and/or peak shifts [89]. Photoluminescence of nano-ZnO particles/SiO
2
aerogels assemblies have shown very strong PL band at 500 nm whose luminescence intensities are 10–50 times higher than that of nanostructured bulk
ZnO. The quantum efficiency is found to lie between 0.2%–1%. This enhancement is attributed to the increase of the singly ionized oxygen vacancies in nano-ZnO particles, which are located in nanopores of the SiO
2
aerogel [90].
Transition metal (TM) ion dopants, such as Ni and V substituting for Zn, suppress the UV emission peak, indicating that the TM doping increases nonradiative recombination processes in this material [91]. These non-radiative transitions arise when free electrons recombine through a process involving a TM ion impurity level instead of populating donor acceptor pairs [92-93]. The suppression of the UV PL peak can also be partially attributed to energy transfer

27 processes from intrinsic donor-acceptor pairs to neighbouring TM ions [92-93].
The conduction band in wurtzite ZnO is constructed mainly from s -type states, while the valence band is formed from p -type states, which is split into three bands due to the influence of crystalfield and spin-orbit interactions [94]. The related free-exciton transitions (FX) from the conduction band to these three valence bands or vice versa are usually denoted by A (also referred to as the heavy hole), B (also referred to as the light hole), and C ( also referred to as crystal-field split band). Our previous studies have suggested that the PL spectrum of Ni doped ZnO nanoneedles at 10 K is dominated by neutral donor bound exciton emissions [95]. We also observe a free A-exciton transition in Ni doped ZnO nanoneedles grown in an
Ar atmosphere at FX
A
= 3.375 eV (Fig. 5) at 10 K. The neutral shallow donor bound exciton dominates because of the presence of donors due to unintentional (or doped) impurities and/or shallow donor-like defects. The free A-exciton bound to a neutral donor is positioned at 3.36 eV (D
0
X
A
). The energy separation between the

28
FX
A
and D
0
X
A
peak gives us the binding energy of the related donor-like defect which is of the order of 15 meV.
We have observed that pure ZnO nanorods calcined at 500 o
C exhibit higher defect emission combined with lower excitonic emission as compared to ZnO nanorods treated at 600 o
C. This is attributed to a sharp increase in the volumetric surface defect concentration with increase in surface area. As the calcination temperature is reduced from 600 o
C to 500 o
C, the surface area to volume ratio for the ZnO nanorods increases by approximately three orders of magnitude due to the small size of these ZnO rods. The dramatic changes in the emission spectra associated with this increase in defect concentration are illustrated in Fig. 6 (a), with structural changes shown in Fig. 6(b); a more complete discussion included in Ref.
[96]. The higher surface defect concentration in the samples calcined at low temperatures (samples with lower dimensions) results a sharp drop in the band-edge intensity near 380 nm and the growth of a very broad peak centered near 475 nm. In this spectrum, the green emission in the range of 450 nm - 500 nm is believed to originate from a transition between the electron in the conduction band and a

29 deep level. This hypothesis is consistent with the luminescence mechanism proposed by Dijken et al. involving an electron in a conduction band and a deeply trapped hole [96].
Besides ZnO a number of other wide band gap semiconductors, including IIIB and IVB group oxides like In
2
O
3
and SnO
2 nanostructures are also actively considered as candidates for fabricating electronic and optoelectronic nanodevices [33, 97]. It is known that bulk In
2
O
3
(E g
= 3.6 eV) does not emit light at room temperature [98-100]. However, In
2
O
3
nanoparticles show PL signals at 430 nm, 480 nm, 520 nm and 637 nm. The origin of most of these peaks from In
2
O
3
films have been attributed to oxygen vacancies [88, 101-102].
3.4.2 Raman Studies on Oxide Nanostructures
Raman spectroscopy provides a powerful tool to probe the structural characteristics of oxide nanostructures. The local symmetry in oxide nanoparticles, specifically ZnO, can be different from that of bulk samples, although the macroscopic crystal structure is identical for both samples [103]. This is illustrated from the Raman

30 spectra comparing bulk and nanostructured samples (Fig. 7). These bulk and nanostructured ZnO samples all have identical crystal structure [104], but markedly different Raman characteristics. Undoped bulk ZnO shows clear Raman peaks at 663 cm
-1
(A
1
(LO)+E
2
(low)), 538 cm
−1
(2LA mode), 437 cm
−1
(attributed to a high frequency nonpolar optical phonon E
2
mode of ZnO), 407 cm
−1
(E
1
(TO) mode) and 381 cm
−1
(A
1
(TO) mode) [105-106]. The E
2
(high) mode at 437 cm
−1
is the strongest mode in the wurtzite crystal structure and any broadening or weakening of this peak indicates the presence of defects in the host lattice. This particular mode, along with the less intense mode at 579 cm
-1
and the A
1
(TO) mode at 381 cm
-1
are strongly suppressed in the nanostructured ZnO brushes and droplets as compared to bulk samples, which is attributed to defectinduced changes in the local symmetry of these samples due to surface defects. In some cases Raman spectra show the presence of
ZnO optical phonon mode, which is red-shifted when compared to bulk ZnO. These are attributed to optical phonon confinement effects [107] or the presence of intrinsic defects on the nanoparticles
[108] . However in the as-grown ZnO nanorods with much bigger

31 size than Bohr exciton radii (~2.34 nm), phonon confinement effect cannot be expected to be the main reason of the shift [109].
Other shifts in the Raman response for ZnO can also be observed in bulk samples with the introduction of substitutional point defects. In the V, Ni, Ti, and Fe doped ZnO bulk samples, the Raman peak frequencies are uniformly red-shifted to lower frequencies
[91]. Such shifts in Raman frequency are believed to depend on residual stress, disorder, and crystal defects present in the samples
[110-111]. The defects induced disorder disrupts long range ordering in the ZnO lattice, which weakens the electric field associated with a mode [111]. Furthermore, the inclusion of point defects in the ZnO lattice can lead to the presence of additional Raman modes, which are referred to as anomalous modes. Two possible mechanisms have been proposed to account for these anomalous modes: disorderactivated scattering or local lattice vibration [106]. Low-frequency
Raman modes have been identified for Fe (19 cm
-1
and 39 cm
-1
) and
Mn (22 cm -1 and 46 cm -1 ) doped ZnO nanoparticles having mean crystallite sizes of ~10 nm and 43 nm respectively [112]. The position of these modes has been connected to the dimension of particles

32 and dopant concentration [112], highlighting the importance of considering the density defects when interpreting or tuning the optical properties of metal oxide nanostructures.
Metal oxide based nanostructures also offer opportunities for promoting new approaches in Raman spectroscopy. Surface enhanced Raman scattering (SERS) is exhibited when a nanoscale dielectric core is surrounded by a metal shell (often called a nanoshell)
[113]. This effect provides a huge increase in the intensity of the
Raman scattering signal, leading to a considerable enhancement of
Raman spectroscopy as a tool for designing biological or chemical sensors [114]. This enhancement in the Raman signal is attributed to a local electromagnetic field enhancement at the metal surface or rough metal structures due to the surface plasmon polaritons [114-
115]. This effect may also be promoted by a chemical enhancement arising from an electronic resonance transfer between surface absorbed molecules and the metal surface [116-118]. Among metal oxide nanostructures SERS has been observed for Au-coated ZnO nanorods having a biomodal size distribution with diameters of 150 and 400 nm prepared on a Si (1 0 0) substrate [119]. These struc-

33 tures show large Raman enhancement factors (EF) values of the order of 10 6 . This enhancement factor is defined as:
EF =
I
SERS
I bulk
/ N ads
/ N bulk where I
SERS
is the intensity of the vibrational mode in the SERS spectrum, I bulk
is the intensity of the same mode in the Raman spectrum and N ads
and N bulk
represent the numbers of the corresponding analytic molecules effectively excited by the laser beam [120].
Highly surface enhanced Raman spectra have also been obtained using indium tin oxide coated gold nanotriangles as well as gold nanoparticles immobilized indium tin oxide [121]. Experimental reports on ZnO crystalline samples covered with Ag-nanoparticles suggest that the resonant Raman scattering process is assisted by metalinduced gap states at the Ag/GaN and Ag/ZnO interfaces. This study provides a view on electron-mediated enhanced Raman scattering
SERS of lattice vibrations in oxide semiconductors [122]. The presence of defect sites such as metal pinholes, can diminish the enhancement [123].
3.4.3 Magneto-Optical Properties of Oxide Nanostructures

34
As will be discussed in more detail in Section V, defects in metal oxide nanostructures can also have strong effects on the magnetic properties of these systems. These induced spin structures can, in turn, affect the optical response of the nanostructures. As an example, the formation of anti-phase boundary defects in metal oxides can give rise to large internal strains [124]. To determine the coercive field of a given sample, longitudinal MOKE magnetometry is measured using a light source. The optical and magneto-optical properties of oxides, such as ZnMnO, in the Faraday configuration give us an estimate of the exchange constant [125]. In this context we emphasize that MOKE effect is very sensitive to the strain, stoichiometry, and film thickness. A large mismatch between the lattice constants of the thin film and substrate can lead to a large residual strain due to the formation of anti-phase boundary defects [126-
127]. In turn, these anti-phase boundary defects may subsequently reduce the net magnetization by changing the exchange interaction across an anti-phase boundary. The magneto-optical properties of metal oxide nanostructures are therefore sensitive to the defect structure in the samples. These structural defect induced changes in the

35 magnetic properties can be investigated by a number of different techniques, including magneto-optic Kerr effect (MOKE) measurements [128].
Magneto-optical probes are provide a powerful tool for identifying the origins of different bound exciton transitions [129]. Exciton bound to ionized impurities can often be identified by the nonlinear splitting of their transitions in an applied magnetic field [130-131].
Such splitting has been observed in ZnO with a magnetic field applied along the c axis. This approach can be valuable in characterizing defects in oxide nanostructures [130].
Metal oxides exhibit a very wide range of magnetic properties, ranging from ferrimagnetism with relatively large saturation magnetizations in Fe
3
O
4
[132], to antiferromagnetic order in NiO
[133] and CoO [134], to simple diamagnetism in ZnO [135] and
TiO
2
[136], to more complex spin structures in Mn
3
O
4
[137]. Expanding this rich set of possible magnetic characteristics, it is wellknown that the magnetic properties of nanostructures materials are

36 often very different than what is observed in bulk systems [138].
Given the diverse nature of metal oxides and specific nanostructures, there is a bewildering array of magnetic properties that are manifested in metal oxide nanostructures, before even considering modifications arising from defects. Rather than attempting to completely summarize the effects of point defects on the magnetic properties of all categories of metal oxide nanostructures, we instead consider only specific examples of systems in which these defects can induce weak ferromagnetic behaviour. We will first briefly discuss some results concerning the development of superparamagnetism and weak ferromagnetic moments in antiferromagnetic metal oxide nanoparticles before visiting the emergence of ferromagnetic order in diamagnetic semiconducting metal oxides induced by point defects.
3.5.1 Magnetism in Metal Oxide Nanoparticles
Measurements on the weak ferromagnetism in nanoscale antiferromagnetic metal oxide systems can be challenging, because of the possibility of accidentally incorporating ferromagnetic secondary phases, such metallic Co inclusions in CoO nanoparticles or thin

37 films [139]. Despite these difficulties, there is a growing realization that antiferromagnetic metal oxide nanostructures often exhibit ferromagnetic properties that cannot necessarily be ascribed to impurity phases [138]. The presence of surfaces (or interfaces) in antiferromagnetic materials can typically lead to the presence of uncompensated spins arising from the incomplete cancellation of the sublattice magnetizations in the antiferromagnetic spin structure [140]. As the surface-to-volume ratio is exceeding large in nanostructures, the fraction of such uncompensated spins can be a considerable fraction of the total. In addition to such “native” uncompensated moments, produced solely by geometrical restrictions, the inclusion of point defects can also yield uncompensated spins, which can also exhibit paramagnetic or weak ferromagnetic behaviour [22]. More complicated magnetic effects, including the onset of multi-sublattice antiferromagnetic order in nanoparticles [141] or modifications of the electronic orbitals at the metal oxide surface [142] have also been proposed, although we omit any discussion of these properties in the following.

38
The magnetic properties of CoO nanoparticles are widely studied [143-146], in part because Co/CoO core/shell nanoparticle represent a model system for the investigation of exchange bias coupling [145]. It has been found that the uncompensated moments present in CoO can be roughly divided into two classes: those moments that are strongly coupled to the antiferromagnetic lattice and those that are not [146]. As least a portion of spins falling in the latter category have been attributed to point defects, and these are believed to produce a paramagnetic or superparamagnetic response at low temperatures. Similar effects have been observed in NiO nanoparticles, which have been shown to exhibit a superparamagnetic response that increases with decreasing particle size [147]. Careful measurements on small NiO nanoparticles find that the effective moment arising from these uncompensated spins exceeds 2000

B
[148], which is considerably larger than what would be expected simply from uncompensated surface spins.
More generally, it has been suggested that low temperature superparamagnetic behaviour is a general characteristic of antiferromagnetic metal oxide nanoparticles [138], including MnO and

39
NiO [149]. Measurements on both NiO and MnO nanoparticles find evidence for superparamagnetic behavior, with saturation magnetizations for the ferromagnetic component on the order of a few emu/g depending on particle size [149]. These samples show hysteretic behaviour at low temperatures, which vanishes at higher temperatures, characteristic of superparamagnetism. These investigations also find that the superparamagnetic blocking temperature increases with increasing particle size for the NiO nanoparticles but, surprisingly, decreases with increasing size for the MnO nanoparticles [149]. This difference in the size dependence of the magnetic properties at least hints at the possibility that the origins for superparamagnetism in the two samples may be different. Because weak ferromagnetism in antiferromagnetic systems can arise both from discontinuities in the magnetic structure at the surface and from point defects, it is challenging to unambiguously assign the observed superparamagnetic moments to one mechanism or the other. Nevertheless, it is clear that structural defects can significantly modify the magnetic response in antiferromagnetic metal oxide nanostructures.

40
3.5.2 Ferromagnetism in Defect-Rich Semiconducting Metal Oxides
A more dramatic example of how point defects can affect the magnetic properties of metal oxides can be found in the observation of ferromagnetism in defect-rich, intrinsically diamagnetic semiconducting oxides [21, 39, 138, 150-152]. The original studies on this class of materials highlighted the development of ferromagnetism in metal oxide films doped with magnetic transition metal ions, in particular, Co substituted into TiO
2
[153] and Mn substituted into ZnO
[154]. Measurements on this class of materials found considerable sample-to-sample variation in the magnetic properties [155-156], leading to suggestions that the magnetic properties were produced by precipitates of a secondary ferromagnetic phase [157]. Furthermore, measurements on nearly stoichiometric Co and Mn doped
ZnO samples found no evidence for ferromagnetism and identified only weak antiferromagnetic nearest-neighbor coupling between the dopant ions [158-159]. Subsequent experiments on Co doped ZnO found evidence for the crucial role played by oxygen vacancy defects in the development of ferromagnetic order in this class of mate-

41 rials, with air annealed films (low oxygen vacancy defect concentration) having negligible magnetizations while vacuum annealed films
(high oxygen vacancy defect concentration) exhibiting distinct ferromagnetism [24].
Despite the recognition that point defects play an important role in the development of ferromagnetic order in transition metal doped semiconducting oxides, it is difficult to disentangle the effects of defects from the contributions arising from the magnetic dopant ions[160]. However, over the past several years it has become apparent that ferromagnetism can develop in diamagnetic metal oxides, which is believed to be driven solely by the presence of point defects
[161]. Signatures of ferromagnetic order were observed in undoped
HfO
2
[151] and subsequently in a range of other metal oxides, including TiO [39-40], In
2
O
3
[21], ZnO [162], and CeO
2
[163-164], among many others [138, 152]. Because most of these systems do not have thermodynamically stable magnetic compositions, it is unlikely that the ferromagnetism arises from the precipitation of ferromagnetic impurity phases. It is found that the magnetic properties of these systems depend strongly on the nature of the point defects

42 present [165], leading to suggestions of defect mediated ferromagnetism in metal oxide nanostructures [166].
It is known that point defects in diamagnetic metal oxides can introduce local moments [22]. The details of this local formation depend sensitively on the compound. For example, oxygen vacancies [167] and zinc interstitials have been predicted to be nonmagnetic in wurzite ZnO, although there are reports of Zn interstitials enhancing the magnetic properties in doped ZnO [168], while oxygen interstitials and zinc vacancies are expected to show sizeable moments, ranging from roughly 0.2
B
[167] to 2
B
[67]. However, both oxygen and cerium vacancies are expected to contribute to the magnetic moment in defect-rich CeO
2
[164]. However, in the complete absence of interactions, defect induced moments would be expected to result in paramagnetic, rather than ferromagnetic behaviour, so the emergence of ferromagnetism in these metal oxide materials is rather surprising.
The strong connection between defects and ferromagnetism in metal oxide nanostructures is demonstrated by studies on TiO
2 thin films [25]. The as-prepared TiO
2
thin films are relatively defect

43 free, as shown in the high resolution transmission electron microscopy images in Fig. 2d, and have a very small magnetization (Fig.
8a). This small moment can be attributed to residual oxygen defect vacancies that remain after air annealing [25]. Conversely, vacuum annealed TiO
2
films, presumably having a much higher concentration of oxygen vacancies, exhibit a much higher concentration of defects, leading to an amorphous structure at the surface (Fig. 2e). Introducing oxygen vacancy defects leads to a considerable enhancement in the magnetization. This increase depends on film thickness, pointing to an intimate connection among microstructure, point defects, and the emergence of ferromagnetism, and reaches 40 emu/cm 3 for films having a thickness of 25 nm (Fig. 8b). Comparing the size of the magnetic signal with an estimate of the total volume occupied by the surface disordered layer in the TiO
2
films leads to the suggestion that the magnetism in these samples may develop solely in the defect-rich regions of the sample [25]. Most significantly, the magnetization decreases systematically when the films are exposed to air under ambient conditions [25], again highlighting both the role of oxygen vacancy defects in developing magnetic or-

44 der and the importance of properly characterizing the stability of point defects when considering their effects on metal oxide nanostructures.
Investigations on CeO
2
, Al
2
O
3
, ZnO, In
2
O
3
, and SnO
2
[138,
169] nanoparticles, among others, find evidence for weak ferromagnetism, having small saturation magnetizations but clear hysteresis loops, albeit often with almost negligible coercivities. The moments in these nanostructured samples is very small, on the order of only
10
-4
to 10
-3
emu/g [138], but significantly larger than the completely negligible magnetizations observed in diamagnetic bulk samples.
Sintering the samples at high temperatures in the presence of oxygen completely suppresses the magnetization [138], leading to the conclusion that the ferromagnetism may be intimately connected with the defect structure, specifically including oxygen vacancy defects, at the surface of the nanoparticles. It is suggested that unpaired electrons trapped on oxygen vacancies may be relevant for the development of ferromagnetism and, furthermore, that such ferromagnetic order may be a general characteristic of all metal oxide nanoparticles
[161, 166].

45
3.5.3 Spin Polarization in Defect-Rich Metal Oxide Nanostructures
There is considerable debate concerning the observations of ferromagnetism in undoped metal oxide nanostructures, including the concern that these magnetic features may arise from the accidental incorporation of ferromagnetic impurities during sample preparation or handling [160, 170]. This uncertainty arises mainly because the very small magnetizations observed in these measurements could, in many cases, be produced by almost negligibly small amounts of contaminants, which could easily be missed by even the most thorough sample characterization. It is therefore desirable to probe the development of magnetic order in these systems using some technique that is not sensitive to trace amounts of ferromagnetic impurity phases. A number of different approaches to this problem have been considered, including magnetotransport measurements [171] and magnetic dichroism spectroscopy [172]. In the following, we discuss another approach based on measurements of

46 the charge carrier spin polarization at normal/superconducting interface [173].
Thin films of undoped In
2
O
3
exhibit a small but distinct ferromagnetic signature with the inclusion of oxygen vacancy defects, introduced by vacuum annealing [21]. Room temperature magnetic hysteresis loops, showing a saturation moment of 0.3±1 emu/cm
3 and a coercive field of 50-200 Oe, are shown in Fig. 9a. Because these vacuum annealed In
2
O
3
films remain conducting to low temperatures, as discussed in Section III, it is possible to probe the spin polarization of the charge carriers using Point Contact Andreev Reflection (PCAR) [173]. The results of PCAR measurements made at
T=2 K with a Nb tip are shown in Fig. 9b. The zero voltage dip in conductance is characteristic of a finite spin polarization of the In
2
O
3 charge carriers. A more careful analysis of the conductance curve yields an estimated spin polarization of approximately 50%, indicative of ferromagnetism in these defect-rich metal oxide nanostructures [21]. While these investigations do not unambiguously prove the existence of intrinsic, carrier mediated ferromagnetic order, they do firmly establish that the measured magnetization is at least

47 strongly coupled to the conduction electrons. Evidence for a finite spin polarization in Co doped ZnO films has also been inferred from low temperature tunneling magnetoresistance measurements on
Co/Al
2
O
3
/Co:ZnO heterostructures [174].
3.5.4 Mechanisms for Magnetism in Metal Oxide Nanostructures
There are a number of proposals for the origin of weak ferromagnetism in defect-rich semiconducting metal oxide nanostructures [161, 166]. As the stoichiometric parent compounds are diamagnetic, the point defects must both provide the magnetic moments and introduce interactions among these defect moments. While oxygen vacancy defects are predicted to yield magnetic moments of approximately 2
B
[67], this relatively small moment is insufficient to produce the high Curie temperatures, typically well above room temperature [175], observed in many defect-rich metal oxide nanostructures. It has been suggested that cation defects may offer much larger magnetic moments, leading to correspondingly larger
Curie temperatures [21]. For example, density functional calculations on CeO
2
find a moment of 2
B
per oxygen vacancy, associat-

48
B
attached to Ce vacancies arising from O 2p orbitals [176]. Similar computations on SnO
2
find that Sn vacancies are magnetic, carrying
B
, while oxygen vacancy defects are non-magnetic [177].
The presence of defect-induced magnetic moments alone is insufficient to ferromagnetism in these semiconducting metal oxide nanostructures; room temperature ferromagnetism requires relatively large interactions among these moments. Local density approximation calculations on SnO
2
find evidence for an oscillating exchange interaction between moments associated with Sn vacancies (V
Sn
), with strong ferromagnetic coupling arising for an average separation of 0.55 nm [177]. This exchange coupling can be modeled approximately by a Ruderman-Kittel-Kasuya-Yosida (RKKY) type interaction, with a k
F
of 0.12 nm
-1
, suggesting the importance of charge carriers in mediating the ferromagnetism. The possible role of defectinduced conduction electrons in promoting ferromagnetic order in defect-rich metal oxides has also been discussed for In
2
O
3
thin films
[178-179]. It has been proposed that the n-type carriers from oxy-

49 gen vacancy defects are highly, though not completely, compensated by indium vacancy defects. The residual n-type charge carriers may have a relatively high density of states at the Fermi level, due to the quasi-localized nature of the donors, promoting high temperature ferromagnetic order [21].
It has recently been suggested that the ferromagnetism in metal oxide nanostructures can be produced by Stoner-type band splitting rather than exchange coupling between local moments
[180]. This is motivated, in part, by the argument that, for reasonable materials parameters, RKKY mediated ferromagnetic order in these systems should develop only below 20 K. In this model, the defects produce a density of states N
S
(E), with a peak lying close to the Fermi level. The introduction of a local charge reservoir can shift the Fermi level to align with a peak in N
S
(E), which can satisfy the Stoner criterion leading to a spin-split impurity band. In this proposal, the charge reservoir is introduced by mixed valence transition metal dopants [66], although other mechanisms are also suggested [22, 138, 180]. In principle, metal ion vacancies, which can act as electron acceptors, could provide the charge reservoir. This

50 model is particularly relevant for nanostructured metal oxide systems, as it is suggested that the ferromagnetism would arise only in defect-rich areas or surfaces, where the defect density of states would be large [138].
While there remains considerable uncertainty surrounding the mechanisms giving rise to ferromagnetism in metal oxide nanostructures, experimentally it is clear that weak ferromagnetism is very generally observed in both antiferromagnetic and diamagnetic systems. Furthermore, a number of studies confirm that this magnetization is correlated with the density of point defects, whether oxygen vacancies, metal ion vacancies, or other. Although the saturation magnetizations associated with this ferromagnetism is generally small, and the coercive fields are also small or vanishing at room temperature for superparamagnetic systems, such ferromagnetic order represents an emergent property driven solely by point defects, which is relevant for understanding how these defects modify the behaviour of metal oxide nanostructures.

51
When considering defects in metal oxide nanostructures, it is crucial to recognize that all of the negative connotations of the word
“defect” are strictly associated only with imperfections in the crystal lattice of the system and should not prejudice the interpretation of the induced changes in the physical properties. While lattice defects can have a large effect on the properties of oxide materials, these changes can be beneficial, detrimental, or neutral depending on the specific application being considered. An enhanced conductivity produced by charge hopping between defects sites is certainly a drawback for a metal oxide insulating layer in a MOSFET [181-
182], but can be crucial for developing optically transparent thin film electrodes [41]. The weak ferromagnetism arising from surface defects in nanoparticles may offer a route to developing spintronic devices [183], but could also produce spurious magnetic signals in nanoscale sensors. It is, however, essential to consider the potential effects of defect-induced properties when investigating metal oxide nanostructures and to recognize the opportunities presented by defects in introducing new functionalities into the system.

52
Controlling the defect chemistry in metal oxide nanostructures offers a route to tuning existing materials properties or incorporating new characteristics. This can be an attractive approach for modifying materials, since this does not involve the addition of new elements, which reduces the potential for the formation of spurious secondary phases. The defect structure in nanomaterials can normally be tuned by varying the conditions during sample preparation, or by post preparation techniques. The possibility of reversibly controlling the defect structure, such as vacuum annealing to introduce oxygen vacancy defects then annealing in an oxygen-rich environment to remove these defects, leads to a tunability that is completely absent when controlling the materials properties by doping. However, it can be exquisitely difficult to experimentally parameterize defect-induced properties because of problems associated with quantifying the concentration of native defects. It is much more straightforward to measure a small percentage of Co doped into ZnO
[32] than it is to determine the degree of oxygen non-stoichiometry for slightly oxygen deficient samples [59]. While simulations provide a great deal of insight into the relationship between the concen-

53 tration of native defects and their effects on the physical properties of oxides, the ability to tailor the defect structure in metal oxides for specific applications will require the development of improved tools to more precisely parameterize the point defects in real systems.
One of the central themes of this short review is that native point defects in oxides can lead to the emergence of completely new materials properties that are absent in the stoichiometric parent compound. Since nanostructured materials can typically develop much higher defect concentrations than bulk systems, researchers working with oxide nanomaterials should remain cognizant of this effect.
One further complication associated with the defect chemistry of metal oxides arises from the fact that in some, but by no means all, cases, the native point defects can be unstable. In some materials, oxygen vacancy defects are removed as materials are held under ambient conditions, while in other systems, metal ion interstitials can readily diffuse even at room temperature [12]. This can lead to a very complicated time dependence for the induced physical properties. For example, weak ferromagnetism in oxygen deficient TiO
2 vanishes after only a few hours for samples stored under ambient

54 conditions [25], while the magnetic signal in In
2
O
3
films persists for years [21]. This dynamical evolution in physical properties is likely to be detrimental for many applications, so it is important to fully investigate the stability of defect-rich metal oxide nanostructures, along with their response, before they can be considered for incorporation into devices.
We have presented a brief and somewhat idiosyncratic overview of defects in metal oxide nanostructures and how these defects modify the physical properties of these systems. We have focused primarily on native point defects in binary oxides, mainly considering only defects and interstitials. The specific nature of the most relevant point defects, whether V
O
, M
I
, or other, varies considerably among different materials, and, in many cases, remains a topic of lively debate. Beyond perturbing the existing materials properties, such as introducing a impurity paramagnetic response or increasing the conductivity, the modifications offered by these point defects can lead to the emergence of entirely new physical behaviour. The

55 presence of point defects can produce metallic conductivity, at least over some range of temperatures, optical responses characteristic of collective behaviour, and weak ferromagnetism. While the specific mechanisms producing these features vary considerably from system to system, the broad nature of the response is somewhat universal, particularly considering the magnetic characteristics. The ability to not only modify but build new physical properties into metal oxide nanostructures through defect chemistry greatly expands the funcationality of these materials and is expected to play a crucial role in the next generation of oxide devices.
We have greatly benefitted from many conversations with A. Dixit,
P. Kharel, B. Nadgorny, R. Naik, V.M. Naik, R. Panguluri, R.
Seshadri, R. Suryanarayanan, and J. Thakur. We acknowledge support from the National Science Foundation through DMR-0644823, from the Jane and Frank Warchol Foundation, and from the Institute for Manufacturing Research at Wayne State University.

56
Fig. 1 Schematic illustration of possible structural defect in metal oxides (adapted from Ref. [23])
Fig. 2 Defects in In
2
O
3
[(a) to (c)] and TiO
2
nanoparticles [(d)-(e)] and micron sized particles [(f) – (g)]. High resolution transmission electron micrographs of typical surface regions for (a) as-deposited and (b) vacuum annealed In
2
O
3
samples. (c) A magnified view of a section of figure (b) with arrows showing typical of several distortions in the crystal lattice. The square region of HRTEM in (c) corresponds to a unit cell of In
2
O
3
shown in the ball and stick model projected along (100) plane. The insets (i and ii) in (c) are the simulated HRTEM images with (i) a oxygen vacancy and (ii) a oxygen vacancy with two adjacent In atoms clustering models. Bright field
TEM images show surface regions of TiO
2
nanoparticles for air (d) and vacuum (e) annealed samples. The TEM (f) and HRTEM (g) images of vacuum-annealed sputter-deposited films show the interior of the crystallites with large number of parallel twin running along the (0 1 1)

57
Fig. 3 Temperature dependent resistivity for an air annealed defect poor In
2
O
3
thin film (open symbols), and the same film after vacuum annealing (defect rich, closed symbols).
Fig. 4 Optical absorption spectra for defect rich (upper curve) and defect poor (lower curve) In
2
O
3
thin films
Fig. 5 Logarithmic plot of low temperature PL spectrum of Ni: ZnO grown in Ar atmosphere at 10 K.
Fig. 6 (a) PL spectra of ZnO nanorods calcined at different temperatures. (b) SEM images of ZnO nanorods synthesized at different temperatures [a more complete discussion is included in Ref 96].
Fig. 7 A Comparison of E
2
phonon shift in Raman spectra of the asdeposited nanostructured and the bulk ZnO samples. The peak at
579 cm
−1
and 381 cm
−1
occurring in bulk ZnO powder are sup-

58 pressed for the brushes as well as droplets. Inset shows shift in the peak occurring at 437 cm -1 .
Fig. 8 Room temperature magnetization curves for an air annealed defect poor TiO
2
film (a) and for a vacuum annealed defect rich
TiO
2
film (b).
Fig. 9 (a) Room temperature magnetization curve for an oxygen deficient In
2
O
3
thin film. (b) Point contact Andreev reflection measurement on the same oxygen deficient In
2
O
3
thin film, measured at
T=2 K using a Nb tip. A fit to this curve yields an estimate spin polarization of P=45%.

59
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