The Role of Defects in Functional Oxide Nanostructures

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

3.1 Introduction

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 in-

trinsic 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], mag-

netic [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 ef-

fect on the physical properties of nanomaterials, so it is crucially im-

portant that they be fully considered when characterizing metal ox-

ide nanostructures. There are a number of thorough and accessible

reviews on defects in particular metal oxide systems, including ZnO

[11-14], TiO2 [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

limit our discussion to point defects. We will specifically focus on

3

oxygen defect vacancies (VO), metal ion vacancies (VM), and metal

interstitials (MI) 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 behav-

iour that emerges due to the presence of defects and do not consider

in any depth the rather more widely studied problem of understand-

ing 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 ox-

ide nanostructures.

3.2 Defects in 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 in-

trinsic defects, and only very briefly touch on dopant ions as point

defects. Within these constraints, a large number point defects can

4

be considered: oxygen vacancies (VO), metal vacancies (VM), oxy-

gen interstitials (OI), metal interstitials (MI), and anti-site defects

(MO or OM). In a large number of cases, the most stable and/or phys-

ically important defects are VO, VM, and MI [17-18] so we focus pri-

marily on these specific examples. The interactions among different

types of defects can play an important role in determining the physi-

cal properties of metal oxides. For example, because Sn is multiva-

lent, SnI point defects can readily form in SnO2, which, in turn, sup-

ports the formation of VO defects leading to n-type conductivity in

defect-rich SnO2 films [19]. Rather remarkably, the presence of

these defects alone in metal oxide nanomaterials is sufficient to pro-

duce new physical properties that are not simply perturbations of the

intrinsic characteristics of the defect-free parent oxide. Heuristi-

cally, 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 VO defects, with oxygen

non-stoichiometry reaching several percent near the surfaces of

nanoscale systems [20]. Because this relatively large defect concen-

5

tration, defect-defect correlations can affect the response of the sys-

tem [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 consider-

able 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 ab-

sence of atoms in some lattice sites while interstitials arise from ex-

tra atoms occupying the space between the atoms in the lattice [18,

23]. Vacancies and interstitial atoms are point defects as these im-

perfections 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

directions [18, 23]. Examples of plane defects comprise stacking

6

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, op-

tical, and magnetic response as we will discuss in the following.

Many metal oxides, including ZnO, TiO2, SnO2 and In2O3,

exhibit marked deviations from stoichiometry under specific anneal-

ing 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 VO and VM [23], and annealing can also pro-

mote the formation of intersitials and antisite defects. At small con-

centrations (0.1 to 1 at%) these point defects are typically assumed

to be randomly distributed throughout the lattice [23]. At higher de-

fect concentrations, a number of different types of defects can de-

velop, including multiple charge state defects and pairs or com-

plexes of defects [17-18, 23]. However, as we will discuss, many of

the properties of defect-rich metal oxides can be understood by con-

7

sidering 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 con-

centrations, enthalpy begins to affect the distribution, which leads to

the formation of new structures including defect clusters, superlat-

tice ordering and extended defects, shear planes, and discrete inter-

mediate phases [23]. Defect-rich TiO2 (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 crys-

tallographic planes [28]. These vacancies are eliminated by the for-

mation 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 TiO2 and In2O3 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 (VO, VZn), interstitials

(Oi, Zni), and antisite defects (OZn and ZnO). VO and Zni were most

generally considered to be the defects responsible for modifying the

electric and magnetic properties of the system [29-34]. However re-

cent work [11] suggests that these defects exhibit high formation en-

ergies under equilibrium conditions. Zn interstitials (Zni) are shallow

donors and fast diffusers with a low migration barrier, 0.57 eV, and

are therefore not stable at room temperature [11, 35]. VZn, 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]. Oi has a high energy

and acts as a deep acceptor at the octahedral site Oi-1(oct) in n-type

ZnO [11, 36]. The antisite defects (ZnO and OZn) have very high for-

mation energies and are unlikely under equilibrium conditions [11,

35-36].

9

The defect structure in metal oxide nanostructures can be im-

aged using high resolution transmission electron microscopy

(HRTEM). Comparing real-space images of air annealed and vac-

uum annealed In2O3 thin films clearly demonstrates the effects of

point defects on the nanostructure [20]. Fig. 2a shows a HRTEM

image of an air annealed In2O3 nanoparticle, showing a well-ordered

lattice with no obvious defects. Vacuum annealing this sample in-

troduces 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 VO and InI, 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 TiO2 (Fig. 2d-g) [25]. Air-annealed TiO2 thin films con-

sisting particles ranging from 300 to 500 nm show good crystalline

order with few defects, as illustrated in Fig. 2d. Conversely, vac-

10

uum-annealed films show numerous crystallographic twin bound-

aries, 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 TiO2 and are inti-

mately 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 TiO2 [28, 37-38]. This

twinning is not observed in TiO2 thin films formed from nanoparti-

cles [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 parti-

cles.

3.2.3 Stability of Intrinsic Point Defects in Metal Oxide Nanos-

tructures

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 VO defects are believed to become stable in

presence of transition metal dopants such as Co [24]. In this particu-

lar study, oxygen defects were introduced in thin film ZnO samples

by annealing at high temperatures (600 oC) 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 disap-

pear after vacuum annealing as the oxygen vacant sites are localized

near Co site (–Zn-VO-Co-). In a number of systems, however, oxy-

gen vacancy defects may not be stable. Studies on oxygen deficient

TiO2 found that the concentration of VO defects decreases rapidly

under ambient conditions [25], although these defects can apparently

be stabilized by transition metal doping [32, 39-40]. Conversely, in

In2O3 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 spe-

cific compound being considered, it is important to properly charac-

12

terize the stability of these defects in a particular metal oxide when

determining the effects such defects may have on the physical prop-

erties of the material.

3.3 Electrical Response

Metal oxides exhibit a range of electrical transport properties,

from metallic to insulating to superconducting [37, 42-43]. The in-

troduction 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 electri-

cal transport properties. We therefore limit our discussion to consid-

ering 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 sam-

ples develop good electrical conductivity with the introduction of

point defects [14, 45]. In the simplest models, this increase in con-

ductivity 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 conduc-

tivity, pointing to an abundance of excess electrons associated with

the defects. VO sites, which normally act as electron donors, are a

possible point defect in all metal oxide systems and it is often be-

lieved that the conducting properties in these materials arise from

oxygen vacancy defects [48]. While we see that oxygen vacancy de-

fects 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 consid-

ered to be the origin of the n-type conductivity in this system [12,

48], recent studies suggest that other point defects may be more rele-

vant for determining the electrical transport properties [11]. These

investigations find that VO sites are deep donors, falling approxi-

mately 1 eV below the conduction band, and are thus unlikely to in-

troduce 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 ZnI 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 pro-

pose that the charge carriers arise from the accidental inclusion of

substitutional hydrogen, HO, 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]. In2O3 can

exhibit a high degree of non-stoichiometry and shows extremely

good n-type dopability [55-58]. It has been suggested that In2O3 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 In2O3 is highly compensated, with the ratio of

n-type free charge carriers to oxygen vacancy defect sites being ap-

proximately 1:5, rather than the 2:1 one would expect is each oxy-

gen vacancy contributes 2 electrons [20, 59]. Recent density func-

tional theory calculations on defect-rich In2O3 find that oxygen va-

cancies, rather than In interstitials, are the likely source of n-type

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 trans-

port properties of transition metal oxides are clearly demonstrated by

the remarkable change in conductivity of In2O3 thin films upon vac-

16

uum annealing, illustrated in Fig. 3. As-prepared In2O3 thin films,

which were crystallized by annealing in air and are presumed to be

close to stoichiometry, are highly resistive and show insulating be-

haviour below room temperature. On vacuum annealing, which in-

troduces 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=1020 cm-3. Concomitant with this in-

crease 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 up-

turn in resistivity below ~80 K.

Rather remarkably, defect-rich In2O3 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 re-

mains above 80% for visible wavelengths [41]. Similar conducting

and optically transparent features are observed in SnO2 samples,

where it is argued that the high oxygen vacancy defect concentration

17

required for producing conductivity is stabilized by the presence of

multivalent Sn interstitials [19]. These same density functional stud-

ies suggest that the donor electrons are not heavily compensated due

to the paucity of acceptor defects (VSn and OI). Furthermore, it is

found that these donors do not have direct optical transitions in visi-

ble 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 In2O3, but the result on SnO2 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 calcula-

tions on Cu2O find that the lowest energy defects are Cu vacancy

point defects, VCu, and a point defect complex consisting of a Cu in-

terstitial, CuI, located between two VCu defects [61]. The VCu defects

are found to produce de-localized holes near the top of the valence

band, leading to p-type conductivity. Measurements on intentionally

undoped Cu2O thin films find a p-type carrier concentration on the

18

order of 1015 cm-3, resulting in a resistivity of approximately 150

cm [47]. P-type conductivity can also develop in the Mott insulator

NiO. Careful measurements have established that VNi sites are the

dominant point defect for determining the electrical properties of

NiO samples, rather than OI sites [62]. It is estimated that the VNi

defect concentration can reach 1016 cm-3 in nanostructured samples,

leading to a 6-8 order of magnitude increase in conductivity over un-

doped 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 con-

sider 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

TiO2, having oxygen vacancy defect concentrations in the range

from 4x1018 cm-3 to 5x1019 cm-3, find that the low temperature trans-

port is consistent with hopping conductivity for high and low VO

19

concentrations (n[VO]) [64]. However, the transport falls in the in-

termediate range between hopping and metallic conduction for

8x1018 cm-3<n[VO]<2x1019 cm-3, where the donor separation is esti-

mated to be five times the effective Bohr radius of the donor elec-

tron [64]. TiO2 does not develop metallic behaviour because in-

creasing 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]. Fre-

quency dependent resistivity studies find that the transport can be

well-modeled by the correlated barrier hopping model, which sup-

poses that holes hop from Ni3+ sites to Ni2+ 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-

chiometric samples were found to exhibit Arrhenius conductivity,

20

associated with the thermioinic emission of band electrons from

grain boundaries at higher temperatures and thermally assisted hop-

ping 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 (dis-

cussed in more detail in Section V) may arise from spin split impu-

rity bands [66]. Ab initio density functional calculations suggest

that VZn point defects should produce metallic behaviour in defect-

rich ZnO, while OI defects give rise to a semiconducting electronic

structure [67].

Surface effects are also expected to affect the electrical trans-

port properties in metal oxides, which is particularly relevant for

nanostructured materials have a high surface area to volume ratio.

Theoretical studies on SnO2 find no evidence for defect-induced

states in the gap [68], while subsurface oxygen defect vacancies in

TiO2 can lead to states falling 0.7 eV below the conduction band

edge [69]. A surface conduction layer, presumably arising from de-

fect states, is found in ZnO; the conductivity of this layer is reduced

on exposure to oxygen [12]. This surface conduction layer provides

21

an additional channel for electronic transport. In2O3 films and

nanostructures can develop a chemical depletion layer near the sur-

face, 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 re-

sponse of an electron gas is the phenomenon of plasma oscillations.

These plasmons are excited at the plasma frequency p, given by

p2=4ne2/meff with n the carrier density and meff the effective mass.

This plasma frequency is typically large for most metals, with

ћp~11 eV for bulk plasmons in Al [71], and depends strongly on

the charge carrier concentration n. Because plasmons reflect collec-

tive behavior of the charge carriers, they represent emergent re-

sponse in insulating metal oxides driven by point defects, which is

completely absent in the parent system. To illustrate the clear devel-

opment of this electronic collective behavior in defect-rich metal ox-

ide nanostructures, we consider the optical response of as-prepared

22

(defect poor) and vacuum annealed (VO defect rich) In2O3 thin films,

as plotted in Fig. 4. The as-prepared sample is insulating and, as ex-

pected for transparent materials, has negligible absorption at ener-

gies well below the bandgap. Conversely, the vacuum annealed

In2O3 sample has a high concentration of VO defects leading to a n-

type charge carrier concentration of roughly 1020 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 ab-

sorbance associated with this peak falls well below the bandgap ab-

sorbance, this plasmon resonance represents the emergence of a dis-

tinct electronic response that is absent in the defect-poor parent

metal oxide structure.

3.4 Optical response

As discussed in the previous section, charge carriers in ox-

ides arise from a number of different sources including interstitial

metal ion impurities, substitutional doping ions, and oxygen vacan-

23

cies. Oxygen vacancies present in the lattice can act as divalent elec-

tron 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 pri-

marily from ionized impurity scattering, the majority of the intrinsic

optical phenomena in these materials arise from ionized defects oc-

cupying 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 electro-

optical devices relies on their ability to efficiently emit or absorb

light; these application prospects are heavily influenced by the en-

ergy band structure and lattice dynamics of the system [72-74]. This

change in optical response in defect-rich oxides is reflected in the al-

tered band to band transitions and absorption energies [75]. More-

over, oxide nanostructures have lower threshold lasing energies due

24

to quantum effects that increase the density of states near band edges

[76-77]. For a number of varied electro-optical applications it is nec-

essary for all charge carriers, both electrons and holes, to be con-

fined [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 un-

derstand 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 Nanostruc-

tures

Nanocrystalline zinc oxide (nano-ZnO) is a wide band gap

semiconductor that is particularly promising for a number of opto-

electronic 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-

tured ZnO possesses a remarkable photoluminescence (PL) spec-

25

trum. The optical response of ZnO changes significantly on the in-

troduction 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 ex-

hibits luminescence defect centers such as oxygen vacancies (lo-

cated at 50 and 190 meV below the conduction band edge), zinc in-

terstitials (located at 2.5 eV below the conduction band edge) [86] as

well as various other native defects [82-85]. However it is interest-

ing to note that the intensity of the green emission can be controlled

in a systematic manner by oxidation and reduction [87]. Nanostruc-

tures such as nano-islands of ZnO show PL emission whose origin

can be explained on the basis of zinc vacancies (VZn) complex de-

fects [88]. The intensity of PL emission from samples containing

such islands is much smaller than that from normal thin films due to

26

a smaller area being covered by the islands. It has also been found

that upon bio-molecule attachment, nano-ZnO powders exhibit fur-

ther induced changes in peak intensities and/or peak shifts [89]. Pho-

toluminescence of nano-ZnO particles/SiO2 aerogels assemblies

have shown very strong PL band at 500 nm whose luminescence in-

tensities 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 oxy-

gen vacancies in nano-ZnO particles, which are located in nanopores

of the SiO2 aerogel [90].

Transition metal (TM) ion dopants, such as Ni and V substi-

tuting for Zn, suppress the UV emission peak, indicating that the TM

doping increases nonradiative recombination processes in this mate-

rial [91]. These non-radiative transitions arise when free electrons

recombine through a process involving a TM ion impurity level in-

stead of populating donor acceptor pairs [92-93]. The suppression

of the UV PL peak can also be partially attributed to energy transfer

processes from intrinsic donor-acceptor pairs to neighbouring TM

ions [92-93].

27

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 crystal-

field and spin-orbit interactions [94]. The related free-exciton transi-

tions (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 FXA = 3.375 eV (Fig. 5) at 10 K. The neutral shal-

low 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 (D0XA). The energy separation between the

FXA and D0XA peak gives us the binding energy of the related donor-

like defect which is of the order of 15 meV.

28

We have observed that pure ZnO nanorods calcined at 500 oC

exhibit higher defect emission combined with lower excitonic emis-

sion as compared to ZnO nanorods treated at 600 oC. This is attrib-

uted to a sharp increase in the volumetric surface defect concentra-

tion with increase in surface area. As the calcination temperature is

reduced from 600 oC to 500 oC, the surface area to volume ratio for

the ZnO nanorods increases by approximately three orders of magni-

tude due to the small size of these ZnO rods. The dramatic changes

in the emission spectra associated with this increase in defect con-

centration 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 emis-

sion in the range of 450 nm - 500 nm is believed to originate from a

transition between the electron in the conduction band and a deep

level. This hypothesis is consistent with the luminescence mecha-

29

nism 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 semiconduc-

tors, including IIIB and IVB group oxides like In2O3 and SnO2

nanostructures are also actively considered as candidates for fabri-

cating electronic and optoelectronic nanodevices [33, 97]. It is

known that bulk In2O3 (Eg = 3.6 eV) does not emit light at room tem-

perature [98-100]. However, In2O3 nanoparticles show PL signals at

430 nm, 480 nm, 520 nm and 637 nm. The origin of most of these

peaks from In2O3 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 symme-

try 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

spectra comparing bulk and nanostructured samples (Fig. 7). These

30

bulk and nanostructured ZnO samples all have identical crystal

structure [104], but markedly different Raman characteristics. Un-

doped bulk ZnO shows clear Raman peaks at 663 cm-1 (A1(LO)

+E2(low)), 538 cm−1 (2LA mode), 437 cm−1 (attributed to a high fre-

quency nonpolar optical phonon E2 mode of ZnO), 407 cm−1 (E1(TO)

mode) and 381 cm−1 (A1(TO) mode) [105-106]. The E2 (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 A1(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 defect-induced

changes in the local symmetry of these samples due to surface de-

fects. In some cases Raman spectra show the presence of ZnO opti-

cal 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 size than Bohr ex-

31

citon radii (~2.34 nm), phonon confinement effect cannot be ex-

pected to be the main reason of the shift [109].

Other shifts in the Raman response for ZnO can also be ob-

served in bulk samples with the introduction of substitutional point

defects. In the V, Ni, Ti, and Fe doped ZnO bulk samples, the Ra-

man 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: disorder-acti-

vated scattering or local lattice vibration [106]. Low-frequency Ra-

man 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 posi-

tion of these modes has been connected to the dimension of particles

32

and dopant concentration [112], highlighting the importance of con-

sidering 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 en-

hanced Raman scattering (SERS) is exhibited when a nanoscale di-

electric core is surrounded by a metal shell (often called a nanoshell)

[113]. This effect provides a huge increase in the intensity of the Ra-

man scattering signal, leading to a considerable enhancement of Ra-

man spectroscopy as a tool for designing biological or chemical sen-

sors [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 ab-

sorbed molecules and the metal surface [116-118]. Among metal ox-

ide 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 or-

der of 106. This enhancement factor is defined as:

EF =

where ISERS is the intensity of the vibrational mode in the SERS spec-

trum, Ibulk is the intensity of the same mode in the Raman spectrum

and Nads and Nbulk represent the numbers of the corresponding ana-

lytic molecules effectively excited by the laser beam [120]. Highly

surface enhanced Raman spectra have also been obtained using in-

dium tin oxide coated gold nanotriangles as well as gold nanoparti-

cles 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 metal-induced

gap states at the Ag/GaN and Ag/ZnO interfaces. This study pro-

vides 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].

34

3.4.3 Magneto-Optical Properties of Oxide Nanostructures

As will be discussed in more detail in Section V, defects in

metal oxide nanostructures can also have strong effects on the mag-

netic properties of these systems. These induced spin structures can,

in turn, affect the optical response of the nanostructures. As an ex-

ample, the formation of anti-phase boundary defects in metal oxides

can give rise to large internal strains [124]. To determine the coer-

cive 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, stoi-

chiometry, 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 struc-

35

ture in the samples. These structural defect induced changes in the

magnetic properties can be investigated by a number of different

techniques, including magneto-optic Kerr effect (MOKE) measure-

ments [128].

Magneto-optical probes are provide a powerful tool for identify-

ing 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 ap-

plied along the c axis. This approach can be valuable in characteriz-

ing defects in oxide nanostructures [130].

3.5 Magnetic Response

Metal oxides exhibit a very wide range of magnetic proper-

ties, ranging from ferrimagnetism with relatively large saturation

magnetizations in Fe3O4 [132], to antiferromagnetic order in NiO

[133] and CoO [134], to simple diamagnetism in ZnO [135] and

TiO2 [136], to more complex spin structures in Mn3O4 [137]. Ex-

panding this rich set of possible magnetic characteristics, it is well-

36

known that the magnetic properties of nanostructures materials are

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 mani-

fested in metal oxide nanostructures, before even considering modi-

fications 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 de-

fects.

3.5.1 Magnetism in Metal Oxide Nanoparticles

Measurements on the weak ferromagnetism in nanoscale an-

tiferromagnetic metal oxide systems can be challenging, because of

37

the possibility of accidentally incorporating ferromagnetic secondary

phases, such metallic Co inclusions in CoO nanoparticles or thin

films [139]. Despite these difficulties, there is a growing realization

that antiferromagnetic metal oxide nanostructures often exhibit fer-

romagnetic properties that cannot necessarily be ascribed to impurity

phases [138]. The presence of surfaces (or interfaces) in antiferro-

magnetic materials can typically lead to the presence of uncompen-

sated 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 compli-

cated magnetic effects, including the onset of multi-sublattice anti-

ferromagnetic order in nanoparticles [141] or modifications of the

electronic orbitals at the metal oxide surface [142] have also been

38

proposed, although we omit any discussion of these properties in the

following.

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 cou-

pling [145]. It has been found that the uncompensated moments

present in CoO can be roughly divided into two classes: those mo-

ments that are strongly coupled to the antiferromagnetic lattice and

those that are not [146]. As least a portion of spins falling in the lat-

ter category have been attributed to point defects, and these are be-

lieved 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 superparamag-

netic response that increases with decreasing particle size [147].

Careful measurements on small NiO nanoparticles find that the ef-

fective moment arising from these uncompensated spins exceeds

2000 B [148], which is considerably larger than what would be ex-

pected simply from uncompensated surface spins.

39

More generally, it has been suggested that low temperature

superparamagnetic behaviour is a general characteristic of antiferro-

magnetic metal oxide nanoparticles [138], including MnO and NiO

[149]. Measurements on both NiO and MnO nanoparticles find evi-

dence for superparamagnetic behavior, with saturation magnetiza-

tions for the ferromagnetic component on the order of a few emu/g

depending on particle size [149]. These samples show hysteretic be-

haviour at low temperatures, which vanishes at higher temperatures,

characteristic of superparamagnetism. These investigations also find

that the superparamagnetic blocking temperature increases with in-

creasing particle size for the NiO nanoparticles but, surprisingly, de-

creases 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 an-

tiferromagnetic systems can arise both from discontinuities in the

magnetic structure at the surface and from point defects, it is chal-

lenging to unambiguously assign the observed superparamagnetic

moments to one mechanism or the other. Nevertheless, it is clear

40

that structural defects can significantly modify the magnetic re-

sponse in antiferromagnetic metal oxide nanostructures.

3.5.2 Ferromagnetism in Defect-Rich Semiconducting Metal Ox-

ides

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 semicon-

ducting 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 par-

ticular, Co substituted into TiO2 [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]. Further-

more, 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

41

dopant ions [158-159]. Subsequent experiments on Co doped ZnO

found evidence for the crucial role played by oxygen vacancy de-

fects in the development of ferromagnetic order in this class of mate-

rials, with air annealed films (low oxygen vacancy defect concentra-

tion) having negligible magnetizations while vacuum annealed films

(high oxygen vacancy defect concentration) exhibiting distinct ferro-

magnetism [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 ap-

parent 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

HfO2 [151] and subsequently in a range of other metal oxides, in-

cluding TiO [39-40], In2O3 [21], ZnO [162], and CeO2 [163-164],

among many others [138, 152]. Because most of these systems do

not have thermodynamically stable magnetic compositions, it is un-

42

likely that the ferromagnetism arises from the precipitation of ferro-

magnetic impurity phases. It is found that the magnetic properties

of these systems depend strongly on the nature of the point defects

present [165], leading to suggestions of defect mediated ferromag-

netism 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 forma-

tion depend sensitively on the compound. For example, oxygen va-

cancies [167] and zinc interstitials have been predicted to be non-

magnetic in wurzite ZnO, although there are reports of Zn intersti-

tials 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]. How-

ever, both oxygen and cerium vacancies are expected to contribute

to the magnetic moment in defect-rich CeO2 [164]. However, in the

complete absence of interactions, defect induced moments would be

expected to result in paramagnetic, rather than ferromagnetic behav-

iour, so the emergence of ferromagnetism in these metal oxide mate-

rials is rather surprising.

43

The strong connection between defects and ferromagnetism

in metal oxide nanostructures is demonstrated by studies on TiO2

thin films [25]. The as-prepared TiO2 thin films are relatively defect

free, as shown in the high resolution transmission electron mi-

croscopy 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 TiO2 films, presumably having a much higher con-

centration 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 en-

hancement 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/cm3 for films having a thickness of 25 nm (Fig. 8b). Compar-

ing the size of the magnetic signal with an estimate of the total vol-

ume occupied by the surface disordered layer in the TiO2 films leads

to the suggestion that the magnetism in these samples may develop

solely in the defect-rich regions of the sample [25]. Most signifi-

44

cantly, 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-

der and the importance of properly characterizing the stability of

point defects when considering their effects on metal oxide nanos-

tructures.

Investigations on CeO2, Al2O3, ZnO, In2O3, and SnO2 [138,

169] nanoparticles, among others, find evidence for weak ferromag-

netism, 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 con-

clusion 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 elec-

trons trapped on oxygen vacancies may be relevant for the develop-

45

ment of ferromagnetism and, furthermore, that such ferromagnetic

order may be a general characteristic of all metal oxide nanoparticles

[161, 166].

3.5.3 Spin Polarization in Defect-Rich Metal Oxide Nanostruc-

tures

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 acciden-

tal incorporation of ferromagnetic impurities during sample prepara-

tion 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 ferromag-

netic impurity phases. A number of different approaches to this

problem have been considered, including magnetotransport measure-

46

ments [171] and magnetic dichroism spectroscopy [172]. In the fol-

lowing, we discuss another approach based on measurements of the

charge carrier spin polarization at normal/superconducting interface

[173].

Thin films of undoped In2O3 exhibit a small but distinct fer-

romagnetic 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/cm3

and a coercive field of 50-200 Oe, are shown in Fig. 9a. Because

these vacuum annealed In2O3 films remain conducting to low tem-

peratures, as discussed in Section III, it is possible to probe the spin

polarization of the charge carriers using Point Contact Andreev Re-

flection (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 In2O3

charge carriers. A more careful analysis of the conductance curve

yields an estimated spin polarization of approximately 50%, indica-

tive of ferromagnetism in these defect-rich metal oxide nanostruc-

tures [21]. While these investigations do not unambiguously prove

47

the existence of intrinsic, carrier mediated ferromagnetic order, they

do firmly establish that the measured magnetization is at least

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/Al2O3/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 ferro-

magnetism in defect-rich semiconducting metal oxide nanostructures

[161, 166]. As the stoichiometric parent compounds are diamag-

netic, 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 approxi-

mately 2 B [67], this relatively small moment is insufficient to pro-

duce the high Curie temperatures, typically well above room temper-

ature [175], observed in many defect-rich metal oxide nanostruc-

tures. It has been suggested that cation defects may offer much

larger magnetic moments, leading to correspondingly larger Curie

48

temperatures [21]. For example, density functional calculations on

CeO2 find a moment of 2 B per oxygen vacancy, associated with

the Ce 4f electrons, with a much larger moment of 4 B attached to

Ce vacancies arising from O 2p orbitals [176]. Similar computations

on SnO2 find that Sn vacancies are magnetic, carrying a moment of

approximately 4 B, while oxygen vacancy defects are non-mag-

netic [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 approxima-

tion calculations on SnO2 find evidence for an oscillating exchange

interaction between moments associated with Sn vacancies (VSn),

with strong ferromagnetic coupling arising for an average separation

of 0.55 nm [177]. This exchange coupling can be modeled approxi-

mately by a Ruderman-Kittel-Kasuya-Yosida (RKKY) type interac-

tion, with a kF of 0.12 nm-1, suggesting the importance of charge car-

riers in mediating the ferromagnetism. The possible role of defect-

induced conduction electrons in promoting ferromagnetic order in

49

defect-rich metal oxides has also been discussed for In2O3 thin films

[178-179]. It has been proposed that the n-type carriers from oxy-

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 reason-

able 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 NS(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 NS(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 transi-

tion metal dopants [66], although other mechanisms are also sug-

50

gested [22, 138, 180]. In principle, metal ion vacancies, which can

act as electron acceptors, could provide the charge reservoir. This

model is particularly relevant for nanostructured metal oxide sys-

tems, 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 nanos-

tructures, experimentally it is clear that weak ferromagnetism is very

generally observed in both antiferromagnetic and diamagnetic sys-

tems. Furthermore, a number of studies confirm that this magnetiza-

tion 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 or-

der 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

3.6 Defect Engineering in Metal Oxide Nanostructures

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

52

nanostructures and to recognize the opportunities presented by de-

fects in introducing new functionalities into the system.

Controlling the defect chemistry in metal oxide nanostruc-

tures offers a route to tuning existing materials properties or incor-

porating 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 nor-

mally be tuned by varying the conditions during sample preparation,

or by post preparation techniques. The possibility of reversibly con-

trolling the defect structure, such as vacuum annealing to introduce

oxygen vacancy defects then annealing in an oxygen-rich environ-

ment to remove these defects, leads to a tunability that is completely

absent when controlling the materials properties by doping. How-

ever, it can be exquisitely difficult to experimentally parameterize

defect-induced properties because of problems associated with quan-

tifying the concentration of native defects. It is much more straight-

forward to measure a small percentage of Co doped into ZnO [32]

than it is to determine the degree of oxygen non-stoichiometry for

53

slightly oxygen deficient samples [59]. While simulations provide a

great deal of insight into the relationship between the concentration

of native defects and their effects on the physical properties of ox-

ides, the ability to tailor the defect structure in metal oxides for spe-

cific 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 com-

pound. 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 am-

bient 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 proper-

54

ties. For example, weak ferromagnetism in oxygen deficient TiO2

vanishes after only a few hours for samples stored under ambient

conditions [25], while the magnetic signal in In2O3 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 in-

vestigate the stability of defect-rich metal oxide nanostructures,

along with their response, before they can be considered for incorpo-

ration into devices.

3.7 Conclusions

We have presented a brief and somewhat idiosyncratic over-

view 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 consider-

ing only defects and interstitials. The specific nature of the most rel-

evant point defects, whether VO, MI, 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

55

the conductivity, the modifications offered by these point defects

can lead to the emergence of entirely new physical behaviour. The

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 funca-

tionality of these materials and is expected to play a crucial role in

the next generation of oxide devices.

Acknowledgements

We have greatly benefitted from many conversations with A. Dixit,

P. Kharel, B. Nadgorny, R. Naik, V.M. Naik, R. Panguluri, R. Se-

shadri, R. Suryanarayanan, and J. Thakur. We acknowledge sup-

port from the National Science Foundation through DMR-0644823,

56

from the Jane and Frank Warchol Foundation, and from the Insti-

tute for Manufacturing Research at Wayne State University.

57

Fig. 1 Schematic illustration of possible structural defect in metal

oxides (adapted from Ref. [23])

Fig. 2 Defects in In2O3 [(a) to (c)] and TiO2 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 In2O3 samples. (c) A magnified view of a

section of figure (b) with arrows showing typical of several distor-

tions in the crystal lattice. The square region of HRTEM in (c) corre-

sponds to a unit cell of In2O3 shown in the ball and stick model pro-

jected along (100) plane. The insets (i and ii) in (c) are the simulated

HRTEM images with (i) a oxygen vacancy and (ii) a oxygen va-

cancy with two adjacent In atoms clustering models. Bright field

TEM images show surface regions of TiO2 nanoparticles for air (d)

and vacuum (e) annealed samples. The TEM (f) and HRTEM (g)

images of vacuum-annealed sputter-deposited films show the inte-

rior of the crystallites with large number of parallel twin running

along the (0 1 1)

58

Fig. 3 Temperature dependent resistivity for an air annealed de-

fect poor In2O3 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) In2O3 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 tempera-

tures. (b) SEM images of ZnO nanorods synthesized at different

temperatures [a more complete discussion is included

in Ref 96].

Fig. 7 A Comparison of E2 phonon shift in Raman spectra of the as-

deposited nanostructured and the bulk ZnO samples. The peak at

579 cm−1 and 381 cm−1 occurring in bulk ZnO powder are sup-

59

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 TiO2 film (a) and for a vacuum annealed defect rich TiO2

film (b).

Fig. 9 (a) Room temperature magnetization curve for an oxygen de-

ficient In2O3 thin film. (b) Point contact Andreev reflection mea-

surement on the same oxygen deficient In2O3 thin film, measured at

T=2 K using a Nb tip. A fit to this curve yields an estimate spin po-

larization of P=45%.

60

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