University of Groningen Engineering complex oxide ... · of polarization-controlled charge transport [37, 27] and switchable photovoltaic ef-fects. [28, 29] The latter behavior was

University of Groningen

Engineering complex oxide interfaces for oxide electronicsRoy, Saurabh

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S. Roy et al.,Role of terminating layer of substrate in electronic transport across BiFeO3 thin films, inpreparation.

Chapter 6

Influence of ferroelectric polarization on hotelectron transport through BiFeO3

Abstract

Oxide heterointerfaces exhibit a strong correlation between spin, charge and orbital de-grees of freedom among its constituent thin films thus presenting them as strong con-tenders for multifunctional devices. To address such novel functionalities, it is impor-tant to understand the interfacial phenomena and tune them accordingly by externalelectric, magnetic or optical inputs. However, the pristine electronic structure of theseinterfaces till now have been determined using the invasive high resolution scanningtransmission electron microscopy (HR-STEM) or synchrotron-radiation studies. Herewe present a novel technique to ascertain the electronic structure of a buried interface ofmetal-ferroelectric semiconductor, and provide a direct evidence of the ferroelectric po-larization control of the interfacial energy band alignment. We grow heterostructuresof Au/BiFeO3/Nb:SrTiO3 on either terminations of the Nb:SrTiO3 substrates, fabricatedevices and probe hot electron transport across these heterostructures. Exploiting the de-pendence of the ferroelectric polarization of BiFeO3 on the terminating site of a substrate,local Schottky barrier heights (SBH) reveal the influence of the interfacial charges on theenergy band alignment of the device, without application of any external bias across theinterfaces. This novel technique can be extended to further study the detailed interfacephysics of oxides and accordingly tune their fundamental properties.

6.1 Introduction and Motivation

Ferroelectrics form a large sub group of dielectric materials, majorly defined bytwo characteristics. First is the existence of spontaneous electric polarization,

and second is a response of this polarization to an external electric field, i.e. thepolarization must be reversible when a sufficiently large electric field is applied inthe opposite direction. Ferroelectric materials are important contenders for non-volatile electronic devices, including memory devices. Research on these materialsnormally deal with the mechanism of the formation of spontaneous polarization,the relationship of this polarization to the atomic and electronic structure, the con-trol of ferroelectric polarization, and the use of ferroelectric materials in electronic

98 6. Influence of ferroelectric polarization on hot electron transport through BiFeO3

Figure 6.1: Schematic of a typical ferroelectric hysteresis loop. The vertical axes marks theremnant polarization (Pr) and the saturation polarization (PS). Figure adapted from [1].

devices. The ferroelectric polarization directions can be switched by the applicationof an external electric field due to the coupling of the electric field E and the polar-ization P. This coupling reflects the change of relative energy E.P. The response ofthe polarization to the applied electric field produces a hysteresis loop which couldbe measured experimentally, as shown in Figure 6.1.

Perovskites are a large family of materials which exhibit a broad range of physi-cal properties including ferroelectricity, ferromagnetism, magnetoelectric coupling,and superconductivity. [2, 3, 4, 5] An ideal perovskite with a chemical formula ofABO3 has a cubic unit cell. By symmetry this cubic structure does not possess aninherent polarization and thus is not ferroelectric. Within this large family of per-ovskites, a few of them are ferroelectric. Structurally, Goldschmidt presented anempirical theory for a tolerance factor (t) connecting the structure of the perovskiteand the ionic sizes of its constituents, which is defined by

t =RA +RB√2(RB +RO)

(6.1)

where RA, RB and RO are the radii of A, B and O atoms respectively. For a valueof t>1, the structure prefers a ferroelectric state with a polar distortion of B atom (forexample, BaTiO3). For t<1, rotational distortion of the oxygen octahedra suppressesferroelectricity (for example, SrTiO3 and CaTiO3). Thus, ferroelectricity is sensitive

6.1. Introduction and Motivation 99

Metal MetalFerroelectric

Charge distribution

Voltage

Electric �eld

(a)

(b)

(c)

(d)

(e)

Figure 6.2: A schematic of (a) a short circuited structure of an electrode-ferroelectric structurewith the polarization displayed; (b) the charge distribution in presence of perfect electrodes;(c) charge distribution; (d) voltage; and (e) electric field profiles in presence of realistic elec-trodes. Note: the ferroelectric film is completely insulating.

to strain due to its strong effect on the atomic structure of the ABO3 unit. The in-plane strain induced by the substrate influences the phase diagram of a ferroelectricthin film, and this effect is widely used nowadays for the growth of ferroelectric het-erostructures with modified ferroelectric properties. Even in a paraelectric material,such as SrTiO3, epitaxial strain may induce a ferroelectric phase. [6]

With progressive decrease in thicknesses of ferroelectric thin films and their ac-tive use in devices, the measured properties reflect the entire system of films, inter-faces, electrodes, and substrates; and the ferroelectric thin films thus can not be con-sidered in isolation. With the growth of ferroelectric thin films on substrates, theirproperties are often considerably different from those of their bulk parent material.The electronic properties also have a characteristic behavior in thin-film form. Whilebulk ferroelectric materials are usually treated as good insulators, with decreasingfilm thickness, it becomes more appropriate to treat them as semiconductors with afairly large band gap.

Conventionally, ferroelectricity has been believed to exist just in bulk materi-als and disappear as the dimensions of a ferroelectric material goes down to the


nanoscale. This is because the depolarizing field in ferroelectric increases as thedimension of the materials is reduced. The depolarization effect on nanoscale fer-roelectric can be understood from Figure 6.2. [7] The depolarizing field is oppositeto the direction of ferroelectric polarization. With ferroelectric polarization beingnormal to the interface, the screening charges from the electrode compensate theferroelectric polarization charge at each interface. For ideal metal, the polarizationcharges are perfectly compensated at both interfaces and therefore no depolarizingfield emerges in the ferroelectric film. However, in reality the effect of interface po-larization charge is not fully compensated due to the incomplete screening by theelectrode. The screening charges are distributed in a vicinity of the interface. Thenet charge distribution results in a voltage jump at each interface from the electrode.With both electrodes short-circuited, the voltage drop in the ferroelectric film fromone interface to the other leads to the depolarizing field opposite to the polarization.This field depends inversely on the thickness of the ferroelectric film. Therefore,when the dimension of the ferroelectric film is down to a critical size, the depolariz-ing field will be strong enough to suppress the ferroelectricity.

Switching the polarization of such films requires a metallic contact, raising fun-damental issues on the behavior of the interface between the ferroelectric layer andthe electrode. In addition to the depolarizing field, the partially covalent nature ofthe bonds in the ferroelectric changes the band structure with respect to that of a per-fectly ionic compound. Intense theoretical activity has been deployed to understandthe electronic structure and band alignment at the interface. [8, 9, 10, 11, 12, 13, 14]First principles calculations can even predict an Ohmic barrier. Pintilie et.al. haveextended semiconductor theory of the metal-insulator interface to the case of ferro-electric capacitors, including the effect of the polarization on the band line-up andconsequential transport properties. [15]

Realizing the formation of the electrode interface and the relation between thepolarization charges and the electronic properties of the interface is crucial in con-trolling the polarization switching and the magnitude of leakage current. These twoare important considerations in the read-write processes in nonvolatile memories.[16] The influence of the polarization and electrode interface is better studied usinggood quality epitaxial ferroelectrics, avoiding in this way the deleterious effects ofthe grain boundaries (large extrinsic contribution to the dielectric constant, back-switching effect, low carrier mobility, etc. [17, 32, 19] In the ferroelectric capacitorsbased on high-quality epitaxial films the hysteresis loop is rectangular, which canbe attributed to the absence of pinning centers for the ferroelectric domains, associ-ated with grain boundaries. This also represents a very efficient mechanism for the

6.1. Introduction and Motivation 101

(a) (b)

Figure 6.3: PFM images (topography and phase) of electrically polarized locations. (a) PFMimages after polarizing a 3x3 µm2 area. The top image shows the topography image andthe bottom one, the phase contrast. (b) PFM images after five successive scans of readingout the ferroelectric polarization of the written location. The phase contrast image shows noimpression of polarization. Applied voltage 8 V to the sample.

screening of the depolarization field during the switching process. The last one maybe related to the height of the potential barriers at the electrode interfaces, whichseems to be influenced by the magnitude and orientation of the ferroelectric polar-ization. [37, 28]

Despite these important theoretical advances, there is little direct experimentaldata, due to the intrinsic difficulty of measuring the electronic structure of a buriedinterface. Determination of the electronic structure of such a device is even morechallenging in traditional electrical characterizing methods. Several groups haveconducted electrical measurement on these systems [22, 23, 24] but they do not di-rectly probe the microscopic interfacial electronic structure. A recent study exhib-ited the resistive switching behavior of semiconducting ferroelectric BiFeO3. [25]The semiconducting behavior of BiFeO3 has been exploited in the demonstrationof polarization-controlled charge transport [37, 27] and switchable photovoltaic ef-fects. [28, 29] The latter behavior was qualitatively explained by a transition from aSchottky to an Ohmic contact at the interface, driven ferroelectrically. [30] Here,we demonstrate a novel non-destructive technique to probe the ferroelectric po-larization dependence of the transport of hot electrons through a metal-BiFeO3-semiconductor stack and reveal the interfacial electronic structure in a non-invasive


way. A combined study of STM topography and BEEM current provides a quan-titative picture of the spatial uniformity of the potential landscape of the buriedmetal-semiconductor interface at the nanometer scale. However, a fundamental is-sue lies in switching the ferroelectric polarization of an ultra-thin film of BiFeO3

without the use of a metallic bottom electrode. This arises as a result of incompletescreening of the charges at the interface with the semiconductor causing the genera-tion of a depolarizing field that strongly impacts the switchability of its ferroelectricstate. An attempt of switching the polarization state of a thin BiFeO3 film grownon Nb:SrTiO3 is as depicted in Figure 6.3. After polarizing a selected area, succes-sive PFM phase images show progressive decrease in the phase contrast; therebyconfirming a local charging phenomena to appear in the contrast initially, and not achange in ferroelectric polarization. Several such locations were probed for switch-ing the local polarization state, all yielding similar results. A solution to studyingthe differently polarized state of BiFeO3 films was arrived at by growing the filmon different terminations of the substrate, which yields the ferroelectric polarizationeither towards the interface or away from it.

6.2 Interface control of polarization state

Atomically sharp interfaces between complex oxides reveal newer functionalities inheterostructures. Recently, in a study by Yu et.al. [31], it was demonstrated thatthe ferroelectric polarization of a thin film is inherently controlled by the interfacialchemical termination. They explored this phenomena by growing BiFeO3 thin filmson conducting La0.7Sr0.3MnO3 (LSMO) bottom electrode and tuned the valence mis-match at the interface. With the aid of RHEED, thin films of BiFeO3 were grownwith atomic sharpness on either terminations of LSMO. When the two structures(FeO2-BiO-MnO2-La,SrO and BiO-FeO2-LaSrO-MnO2) were probed for their ferro-electric properties using a PFM, the as grown polarization of the sample directlycorrelated with the interface termination of the heterostructure. They observed thatsamples possessing the MnO2 termination had polarization upwards (pointing upoutwards the interface) whereas the ferroelectric polarization was pointing down-wards (pointing down into the interface) for the films possessing LaSrO termina-tions (Figure 6.4).

Such a dependence was studied by density functional theory (DFT) calculationsbased on the atomic scale electrostatic potentials across two non-equivalent inter-faces, as shown in Figure 6.5. They observed an increase in the potential gradientwhen the direction of polarization was against the preferred direction of electrostatic

6.3. Sample growth, fabrication and measurement technique 103

P P

Figure 6.4: A schematic of two heterointerfaces between LaSrMnO3 and BiFeO3, and theirrespective atomic plane stacking. The first interface (left) is FeO2-BiO-MnO2-LaSrO and thesecond is BiO-FeO2-LaSrO-MnO2. The charge associated with each atomic plane is also in-dicated. The arrows indicate the as-grown ferroelectric polarizations of the BiFeO3 films.Adapted from [31].

potential and it resulted in an enhancement of the depolarizing field for a specificpolarization direction. This potential gradient were composed of two major con-tributions, ferroelectric depolarizing field and the electrostatic potential differencebetween two non-equivalent interface, however the major contribution to the pref-erence of ferroelectric polarization direction originated from the valence mismatchat the interface.

They extended this study for different interfaces and tabulated the presence ofpreferred polarization direction for thin films of BiFeO3 grown on different substrateterminations. Thus, this result provides us a route to obtain different bulk polariza-tion directions of BiFeO3 when grown on differently terminated substrates therebyproviding us with an opportunity to circumvent the issue of polarizing different re-gions of the same ferroelectric film in different directions. A growth of BiFeO3 onthe TiO2 and SrO terminations of Nb:SrTiO3 would eventually lead to ferroelectricpolarizations towards and away from the interface.

6.3 Sample growth, fabrication and measurement tech-nique

BiFeO3 thin films were grown using pulsed laser deposition on both SrO and TiO2

terminations of Nb:SrTiO3 substrates (Nb = 0.01 wt. %). The initial steps of chem-ically terminating the substrate prior to deposition has been mentioned before indetails. To attain an uniform SrO termination of the substrate, 1 unit cell (u.c.) ofSrO is deposited on the previously TiO2 terminated substrate. Subsequently, 4 nm


1

LSMO LSMOBiFeO3

V (FE)

2

V

(Interface)

+

V

(FE+Interface)

-

(a)

(b)

(c)

Figure 6.5: Schematic representation of electrostatic potentials through the supercell dueto contributions from (a) the ferroelectric depolarization field; (b) the asymmetric metal-ferroelectric interfaces; and (c) a resultant contribution from both. The green and the bluearrow represent the direction of ferroelectric polarization and the red arrows indicate the in-terface dipole due to valence mismatch at the interfaces. This resultant asymmetry dictatesthe ferroelectric polatization of the BiFeO3 film grown. Adapted from [31].

thick BiFeO3 films were deposited on either samples. Epitaxially, these films weregrown at a substrate temperature of 670◦C from a target with 11 % excess Bi to main-

6.4. Polarization control of energy band alignments 105

tain the stoichiometry of volatile Bi. The background oxygen pressure was kept at0.3 mbar, the fluence at 2 J/cm2 and a pulse repetition rate of 0.5 Hz. The targetto substrate distance was kept at 58 mm. Simultaneously during the growth of thefilm, reflection high energy electron diffraction (RHEED) specular spots were usedto monitor the quality of film grown. Figure 6.6 (a, b) shows the AFM topographyscan and their corresponding RHEED spots after the growth of BiFeO3 on SrO andTiO2 terminations, respectively. This was followed by patterning device structuresof 250 µm x 650 µm using standard UV lithography and ion beam etching technique.Thereafter, the sample was patterned again to selectively deposit 10 nm of Au onthe device area followed by the deposition of top and bottom contacts of Au (100nm) and Ti (75 nm)/Au (75 nm), respectively. For BEEM studies, a modified com-mercial STM system from RHK Technology was used. The samples were studiedboth for their macroscopic I-V characteristics as well as with BEEM at a local scale.The primary difference in studying the electrical transport across the interface liesin the application of voltage across the metal-semiconductor (M-S) interface. Themeasurements were all done at room temperature (300 K), and thermionic emissiontheory was used to ascertain the Schottky barrier height and the quality of the M-Sinterface. Additionally, since the injection current in BEEM experiments is in the or-der of nanoamperes (nA), the zero bias resistance of the device assumes importancein collecting the BEEM current. In the devices shown above, the zero bias current issufficiently low to obtain spectroscopy data. In BEEM, no bias is applied to the M-Sinterface and thus, the injected hot electrons enter the conduction band by virtue oftheir kinetic energy after traveling through the film and across the interface. TheBEEM current, IB is recorded over a local area of a few nanometers (nm), at a con-stant tunnel current, IT , while varying the VT . IB clearly depends on the thickness ofthe metallic over layer and decreases exponentially with increasing thickness. Eachspecific location on the device was probed with 150-200 spectra and many locationson different devices were averaged to obtain the representative BEEM spectrum.Figure 6.6 (e) shows the schematic of the measurement set-up. At each of these loca-tions, the local SBH was extracted from the onset of their respective BEEM currentusing the Bell-Kaiser (B-K) model.

6.4 Polarization control of energy band alignments

In order to understand the formation of electrode interface and the relation be-tween polarization charges and the electronic properties of the device, a compos-ite energy band diagram of the system is required. Here we consider an asym-metric heterostructure consisting of thin metallic films, an n-doped semiconductor


(a) (b)

Ti (75 nm)/Au(75 nm)Ohmic contact

BiFeO3 (4 nm)

Au (10 nm)

STM contact

IB

IT

VT

(c)

Figure 6.6: (a, b) AFM image of 4 nm BiFeO3 deposited by PLD on SrO and TiO2 terminatedsubstrates (left and right). Their corresponding root mean square (rms) roughness is 0.26 nmand 0.42 nm, respectively. The corresponding images of the specular RHEED spots are shownbelow. (c) Schematic representation of the device and the BEEM measurement set-up.

(Nb:SrTiO3) and a ferroelectric thin film. Although BiFeO3 is electrically insulatingin bulk, oxygen vacancies play a very important role in controlling the conduction inthin films. These oxygen vacancies present themselves as n-type defect states in theband gap that enables thermally activated electrons to be promoted to the conduc-tion band. This conduction through BiFeO3 thin films can be controlled by differentoxygen content, either by altering the cooling rate or oxygen pressure after growth,but otherwise growing under similar conditions. This renders the device to havetwo interfaces, i.e. metallic thin film-BiFeO3 and BiFeO3-Nb:SrTiO3.

6.4. Polarization control of energy band alignments 107

These two interfaces are discussed separately. We start with the BiFeO3-Nb:SrTiO3

interface. From previous studies, the energetic positions of the conduction band(CB) and valence band (VB) edges of Nb:SrTiO3 without the applied sample bias (atzero sample bias) are extracted to be -0.06 V and -3.26 V, respectively. [34] The en-ergetic positions of the CB and VB edges of BiFeO3 without the applied sample biasare 1.33 V and -1.37 V. [34] With presence of no additional dopant concentrations,the Fermi level of BiFeO3 is situated close to the middle of the gap. A high levelof doping in Nb:SrTiO3 makes it a degenerate semiconductor. The built-in poten-tial at the BiFeO3/Nb:SrTiO3 interface can be determined based on the difference ofthe work function between them. The work function of Nb:SrTiO3 is 4.08 eV andBiFeO3 is 4.70 eV. [35] Consequently, when Nb-SrTiO3 and the as-grown P-downstate in BiFeO3 are joined together, the ideal value of built-in potential (φbi) is 0.62eV. Some of the electrons from Nb:SrTiO3 spntaneously move into BiFeO3 becausethe Fermi level of BiFeO3 is lower than that of Nb-SrTiO3 and such a movement ofelectrons causes negative charges to accumulate in BiFeO3 depletion layer. Withoutconsidering the contribution of the as-grown polarization, the band structure of vir-gin BiFeO3/Nb:SrTiO3 can be schematically plotted in Figure 6.7

In accordance to earlier theoretical estimations, the surface charge neutralitylevel (φ0) is located 2.20 eV above the valence band edge. [36] The φbi value of0.62 eV, suggests that the Fermi level at the interface is below φ0 resulting in a pos-itively charged BiFeO3-Nb:SrTiO3 interface. However, in order for it to be a stablesystem, the as-grown interface trap charge prefers to be zero, suggesting the Fermilevel of BiFeO3 approaches the surface charge neutrality level at the interface. Thisconsideration motivates the as-grown polarization state to be downward. The posi-tive bound charges of the as-grown P-down state in BiFeO3 will assist the reductionof the unwanted positive trapped charges at the hetero-interface. Consequently, thedepolarization field resulting from the as grown polarization contributes to a widerelectron region of the band at the interface, as schematically indicated in Figure6.7(c).

According to the model proposed Pintilie et. al. [32], the ferroelectric polariza-tion can influence the built-in potential as ∆φ = φ

′

bi - φbi = ± Pδ/ε0εS , where φ′

bi isthe built-in potential with contribution from polarization, φbi is the built-in poten-tial without contribution from polarization, ε0 is the permittivity of free space, εS isthe static dielectric constant, P is ferroelectric polarization, and δ is the thickness ofan interface layer between the polarization surface charge and the physical interfacewith the metallic system. For BiFeO3 films, the ferroelectric polarization value P is


EC

EV

EF

Nb:SrTiO3

EC

EV

BiFeO3

3.20 eV

4.08 eV

2.70 eV

3.30 eV

EV

Nb:SrTiO3BiFeO3

EC

EC

EV

EF

4.08 eV 3.30 eV

0.62 eV

0.62 eV

2.70 eV

bi =

BiFeO3

EV

Nb:SrTiO3

EC

EC

EV

EF

4.08 eV 3.30 eV

1.00 eV

2.70 eV

’bi =

Pdown

BiFeO3

EV

Nb:SrTiO3

EC

EC

EV

EF

4.08 eV 3.30 eV

0.20 eV

2.70 eV

’bi =

Pup

(a) (b)

(c) (d)

Figure 6.7: Schematic diagram of the energy bands at the interface of Nb:SrTiO3 and BiFeO3.(a) Energy bands when BiFeO3 and Nb:SrTiO3 kept separate. (b) BiFeO3 and Nb:SrTiO3 incontact with eachother, but without any applied bias and the ferroelectric polarization of theBiFeO3 is not considered. (c) Energy band alignments when the ferroelectric polarization isdownwards. (d) Energy band alignment with ferroelectric polarization upwards, causing achange in the built-in voltage of 0.24 V.

approximately 60 µ C/cm2, and the reported low-frequency dielectric constant isapproximately 100. The variation in the built-in potential ∆φ can then be estimatedas 0.38 eV for a δ value of the order of 0.4 nm. With these values of φbi (0.62 V)and ∆φ (0.38 V), the predicted φ

′

bi for a downward polarized BiFeO3 is 1.00 V. Simi-larly for the polarization up scenario, the ideal built-in voltage φ

′

bi is deduced to be0.24 V from φbi (0.62 V) and ∆φ (0.38 V). Such a behavior, previously observed, isattributed to the screening of the ferroelectric polarization at the interface by oxy-gen vacancies. [37] [38] Figuratively, the band alignment at the heterointerface forpolarization up situation is shown in Figure 6.7(d). However, in this scenario the

6.5. Results and discussions 109

EV

EC

EF

5.1 eV

2.7 eV

BiFeO3Au

Vacuum

3.3 eV

Pup

EV

EC

EF

5.1 eV

2.7 eV

BiFeO3Au

3.3 eV

Pdown

Figure 6.8: Schematic diagram of the energy bands at the interface of Au and BiFeO3 for thedifferent polarization states of BiFeO3.

depolarization field from the negatively bound charges result in an upward bend-ing of the band and produces a narrow depletion region near the interface.

On the other hand, the interface of the metallic thin film on semiconductingBiFeO3 aligns it’s energy bands in accordance to the conventional Schottky barriertheory. This situation is far simplified; a change in the ferroelectric polarizationstate of BiFeO3 causes a change in the interface potential of BiFeO3 thereby caus-ing a change in the band alignment. From the above discussions, we can herebyconstruct the composite band diagram of the entire system as shown in Figure 6.8.

6.5 Results and discussions

The electronic transport studies on the samples mainly comprised of both macro-scopic current-voltage (I-V) measurements and further local scale investigations us-ing BEEM. Figure 6.9 shows the I-V characteristics of the samples at 300 K. Eachplot is a representative of multiple measurements on several devices on the samesample and all of them show good diodes. The I-V characteristics show clear rec-tifying behavior, where the forward biased current shows a linear increase in thesemilog scale, whereas no current flows in the reverse bias. This linear increase anda strong asymmetry in I is consistent with the thermionic emission over a Schottkybarrier, as expected from the asymmetry of the energy profiles. However, a measureof dependence on this charge transport phenomena is given by the ideality factor


-1.0 -0.5 0.0 0.5 1.0 1.5

10-11

10-9

10-7

10-5

10-3

|Cur

rent

(A)|

Voltage (V)

I

V

SrO-Nb:SrTiO3

BFOAu

T = 300 KB = 1.25±0.02 eVn = 1.2

-1.0 -0.5 0.0 0.5 1.0 1.510-12

10-10

10-8

10-6

10-4

|Cur

rent

(A)|

Voltage (V)

I

V

TiO2-Nb:SrTiO3

BFOAu

T = 300 KB = 1.06±0.02 eVn = 1.3

(a) (b)

Figure 6.9: Electrical (I-V) characteristics of Au/BiFeO3/SrO-Nb:SrTiO3 andAu/BiFeO3/TiO2-Nb:SrTiO3 devices at RT (300 K). Inset shows a schematic of the electricalconnections. The straight line fit is the thermionic emission fit allowing an extraction of theSchottky barrier height and the ideality factor.

(n). Classical thermionic emission current does not depend on the Schottky barrierwidth but only on the barrier height; which yields an ideality factor of 1. The for-ward bias characteristics were fitted with thermionic emission theory to extract theSchottky barrier height (SBH) at zero bias and the ideality factor.

The measurement schematic for the I-V characterization is as shown in Figure6.9 (inset). As the Au-BiFeO3 forms a M-S interface, a peak in electrical potentialprofile of height eΦB is formed, and this barrier height is modified by the polar-ization state of BiFeO3 (up or down). The macroscopic SBH for the two polariza-tion states of BiFeO3 are extracted to be 1.25±0.02 eV (Au/BiFeO3/SrO-Nb:SrTiO3)and 1.06±0.02 eV (Au/BiFeO3/TiO2-Nb:SrTiO3). A simple Schottky-Mott modelfor a M-S junction would expect a SBH of 1.8 eV considering the work functionof Au to be 5.1 eV and the electron affinity of BiFeO3 as 3.3 eV. The experimen-tally extracted SBH deviates from the expected value and clearly shows an addi-tional interface effects determining the final energy band line-ups. The idealityfactors for Au/BiFeO3/SrO-Nb:SrTiO3 (1.2) and Au/BiFeO3/TiO2-Nb:SrTiO3 (1.3)can attribute to charge transport mechanisms like tunneling or thermally assistedfield emission being prevalent across the interface. However, the barrier heightsextracted for BiFeO3 grown on SrO and TiO2 terminations were different therebyindicating an influence of polarization in the SBH obtained from macroscopic I-Vmeasurements.

The hot electron transmission at several locations through the metallic overlay-


-2.4 -2.0 -1.6 -1.2 -0.8

0.0

1.0x10-4

2.0x10-4

3.0x10-4

4.0x10-4

BEEM

tran

smis

sion

(pA/

nA)

Bias Voltage (V)

SrO-Nb:SrTiO3

BFOAu

STM tip

VT

IT

IB

-2.4 -2.0 -1.6 -1.2 -0.8

0.0

2.0x10-3

4.0x10-3

6.0x10-3

8.0x10-3

BEEM

Tra

nsm

issi

on (p

A/nA

)

Bias Voltage (V)

TiO2-Nb:SrTiO3

BFOAu

STM tip

VT

IT

IB

Figure 6.10: BEEM spectra for the Au/BiFeO3/SrO-Nb:SrTiO3 and Au/BiFeO3/TiO2-Nb:SrTiO3 devices at RT (300 K). Each spectrum is a representative of several spectra takenat the same location. For each device, three such BEEM spectra are represented which pro-vides an idea of the uniformity in BiFeO3 thin films growth on SrO and TiO2 terminations ofSrTiO3.

ers and across the Au/BiFeO3/TiO2-Nb:SrTiO3 and Au/BiFeO3/SrO-Nb:SrTiO3 areshown in Figure 6.10. In both cases, the BEEM transmission is found to increasesharply with increasing VT beyond a certain tip bias, while IT is kept constant.Many such spectra were taken to ascertain the uniformity of the hot electron trans-port property of the devices. These hot electrons are injected by the STM tip acrossthe vacuum tunnel barriers into the Au overlayer. These hot electrons propagatethrough the film and are scattered both elastically and inelastically in the process,and reach the unbiased Au-BiFeO3 interface. A few of these electrons that possessenough energy and satisfy the stringent moment conservation criteria, overcome theSchottky barrier and enter the conduction band of the 4 nm thick BiFeO3 thin film.As the carriers enter the BiFeO3 thin film and travel across it by virtue of their ki-netic energy, they will, due to the band bending at the BiFeO3-Nb:SrTiO3 interface,be accelerated away from the interface into the conduction band of the Nb:SrTiO3

and constitute the BEEM current. As is evident from the energy band diagrams, theelectric field at the BiFeO3-Nb:SrTiO3 interface always aids in transmitting chargecarriers away from the interface. A change in the sign of polarization charge onlychanges the magnitude of the electric field at this interface. Consequently, the high-est energy barrier that hot electrons encounter in their transmission is the interfacebetween the metallic over-layer and the BiFeO3. Since the electric potential at thisinterface is directly related to the ferroelectric polarization state of BiFeO3, measur-ing a change in the SBH at this interface reflects its polarization state. This can be


obtained from the onset of the BEEM transmission beyond a certain threshold volt-age (corresponding to the SBH). Using the B-K model [33], we obtain the local SBHfor Au/BiFeO3 interface by extrapolating the straight line so as to intersect the volt-age axis in the plot of the square root of the BEEM transmission versus the appliedvoltage, as shown in Figure 6.11. The measurements indicated an extracted SBH of1.56±0.03 eV for up Au/BiFeO3/SrO-Nb:SrTiO3 devices and a value of 1.07±0.03eV for Au/BiFeO3/TiO2-Nb:SrTiO3. The SBH directly corresponds to the collectedBEEM current for the two samples. We observe a higher IB for Au/BiFeO3/TiO2-Nb:SrTiO3, indicating a larger number of transmitted electrons that can overcomethe SBH and contribute to the BEEM current. In either devices, the injected hotelectrons travel the same thickness of Au overlayer, and thus undergo similar at-tenuation. A difference of an order of magnitude in the BEEM transmission for thetwo devices point to the dissimilar energy barriers. Furthermore, a wide spread inthe BEEM current in Au/BiFeO3/TiO2-Nb:SrTiO3 hints the non-uniform growth ofBiFeO3 on TiO2 termination. The SBH extracted from different locations and devicesfrom either samples are represented in the histogram shown in Figure 6.11 (inset).

Although location-wise, a slight deviation in the extracted SBH was observedfor the two kinds of samples, a clear difference in the energy band line-ups at theAu and BiFeO3 interface was observed, indicating a tuning of the barrier potentialby virtue of its ferroelectric polarization. In order to the confirm the observed phe-nomena, we grew a heterostructure of metallic LaSrMnO3 (LSMO)/BiFeO3/SrO-Nb:SrTiO3. Here in place of Au, a thin film of LSMO (3.5 nm) is grown on a similar4 nm thick BiFeO3 film on SrO terminated Nb:SrTiO3 substrate. Devices were fab-ricated on the sample in a similar fashion as mentioned earlier. Preliminary I-Vcharacterization of the sample shows a good rectifying diode with an extracted SBHof 1.33±0.02 eV and an ideality factor of 1.3, Figure 6.12(a). The leakage current atzero bias is sufficiently low (10−10 Amps) and hot electron transport studies wereperformed on this sample. The representative BEEM spectra are obtained by sweep-ing the tip bias while simultaneously recording the IB at a constant tunnel currentinjection. The local SBH extracted from these measurements, using the B-K model[33] is 1.74±0.03 eV as shown in Figure 6.12(b), which are comparable to the pre-vious sample with similar direction of polarization. We thus clearly establish thedependence of the grown polarization state of BiFeO3 on the extracted SBH fromthe devices.

Interfaces between two oxides are known generate novel electronic phases thatare not accessible in bulk. In the case of BiFeO3 grown on TiO2/SrO-Nb:SrTiO3, ourstudy provides an opportunity understand the role played by the heterointerfaceis stabilizing the polar phase of BiFeO3; whose origin lies in the coupling between


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Figure 6.11: The extracted SBH at the Au/BiFeO3 interface Schottky interface, obtained byplotting the square root of the normalized BEEM current IB vs. VT and fitting it with theB-K model at 300 K. The inset shows the corresponding occurrences of the distribution ofSBH obtained at different locations for both the devices. The error bars reflect the standarddeviation of the measured onset of the BEEM current at a location.

polarization, oxygen octahedral rotation and their tilt.

The concept of interfacial coupling rests on the idea that the substrate tends totranslate its octahedral pattern to the film grown on top of it. Magnitude and sense(in-phase or out-of-phase) of these rotations determines the electronic and magneticproperties. The stability of the polar phase of BiFeO3 lies in the interface charge,misfit strain and octahedral tilt at the interface of substrate. The interface selectionof the ferroelectric polarization is due to the preference of SrO or TiO2 interfaceto stabilize either an in-phase or out-of-phase rotation pattern in BiFeO3 about thefilm-normal direction. The TiO2 terminated interface enables the BiFeO3 film tohave a complete FeO6 octahedron, and as in the bulk the film prefers to adopt anout-of-phase rotation of the octahedron. This results in a more energetically favoredstate which causes a polar atomic displacement. In contrast, the SrO terminationleads to FeO6 octahedra with nearest neighbor interactions determined by the Sror Bi cations. Since SrFeO3 has a structural tendency to adopt an in-phase rota-tion about the film normal which presents a lowest energy configuration, whichcauses a stronger interaction between the Bi cation and the oxygen octahedra, al-tering the cationic displacement thereby manifesting as a different polarized state.Such a change in ferroelectic polarization in our samples is captured by a respec-tive variation in the energy band alignment at its interface with a metal. Under


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Figure 6.12: (a) Electrical (I-V) characteristics of LSMO/BiFeO3/SrO-Nb:SrTiO3 devices atRT (300 K). Inset shows the representative schematic of the set-up. (b) Representative BEEMspectra of the device, an average of several spectra taken at various locations and on differentdevices. The inset shows the extracted SBH of 1.74±0.04 eV. The error bars reflect the standarddeviation of the measured onset of the BEEM current at a location.

normal circumstances, detection of out-of-plane polarization in ultra-thin films be-comes quite challenging and even more without a bottom electrode. Recently, suchstudies have been conducted using annular dark field (ADF) scanning transmissionelectron microscopy (STEM) on thin films which determine the atomic positions atthe heterointerface to determine the polarization state. Our technique, on the otherhand gives an indirect knowledge of the ferroelectric polarizations in ultra-thin films[39].

6.6 Conclusion

In conclusion, we have demonstrated a novel way of studying the different po-larization states of BiFeO3 thin films obtained by atomically precise control of theinterface using BEEM. The as-grown polarization states of BiFeO3 thin films wereengineered by its growth on SrO and TiO2 terminated Nb:SrTiO3 substrates. Macro-scopic I-V characteristics of the both the samples showed good rectifying diode, butthe extracted SBH of samples where BiFeO3 was grown on SrO terminations (down-polarized) consistently showed a higher value than the one grown on TiO2/Nb:SrTiO3

(up-polarized). The samples were characterized for hot electron transport across theheterointerface. From the collected BEEM current, the extracted SBH from the B-K

6.6. Bibliography 115

fit clearly showed a higher potential barrier for Au/BiFeO3/SrO-Nb:SrTiO3 than forthe sample with BiFeO3 grown on TiO2 terminations. To confirm this observation,we fabricated another sample (LSMO/BiFeO3/SrO-Nb:SrTiO3), which is also polar-ized upwards, and studied hot electron transport through it. The extracted SBH wassimilar to the sample with BiFeO3 grown on SrO termination. Our experiments haveextended the conventional device design in BEEM to detect the ferroelectric polar-ization effects, allowing a manifestation of the polarization states as different SBH inour measurements. Utilizing the unique capabilities of BEEM, the electronic poten-tial landscape of the buried interface can be mapped out at a nanoscale. Such studiescould be instrumental in understanding the role of adsorbates acting as screeningcharges.

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University of Groningen Engineering complex oxide ... · of polarization-controlled charge transport [37, 27] and switchable photovoltaic ef-fects. [28, 29] The latter behavior was