A high-mobility two-dimensional electron gas at the spinel/perovskite interface of [gamma]-Al2O3/SrTiO3

Abstract

The discovery of two-dimensional electron gases at the heterointerface between two insulating perovskite-type oxides, such as LaAlO(3) and SrTiO(3), provides opportunities for a new generation of all-oxide electronic devices. Key challenges remain for achieving interfacial electron mobilities much beyond the current value of approximately 1,000 cm(2) V(-1) s(-1) (at low temperatures). Here we create a new type of two-dimensional electron gas at the heterointerface between SrTiO(3) and a spinel γ-Al(2)O(3) epitaxial film with compatible oxygen ions sublattices. Electron mobilities more than one order of magnitude higher than those of hitherto-investigated perovskite-type interfaces are obtained. The spinel/perovskite two-dimensional electron gas, where the two-dimensional conduction character is revealed by quantum magnetoresistance oscillations, is found to result from interface-stabilized oxygen vacancies confined within a layer of 0.9 nm in proximity to the interface. Our findings pave the way for studies of mesoscopic physics with complex oxides and design of high-mobility all-oxide electronic devices.

ARTICLE

Received 4 May 2012 | Accepted 17 Dec 2012 | Published 22 Jan 2013

Y.Z. Chen1, N. Bovet2, F. Trier1, D.V. Christensen1, F.M. Qu3, N.H. Andersen4, T. Kasama5, W. Zhang1, R. Giraud6,7,J. Dufouleur6, T.S. Jespersen8, J.R. Sun3, A. Smith1, J. Nygrd8, L. Lu3, B. Bchner6, B.G. Shen3, S. Linderoth1 & N. Pryds1

The discovery of two-dimensional electron gases at the heterointerface between two insulating perovskite-type oxides, such as LaAlO3 and SrTiO3, provides opportunities for a new generation of all-oxide electronic devices. Key challenges remain for achieving interfacial electron mobilities much beyond the current value of approximately 1,000 cm2 V-1 s-1 (at low

temperatures). Here we create a new type of two-dimensional electron gas at the hetero-interface between SrTiO3 and a spinel g-Al2O3 epitaxial lm with compatible oxygen ions sublattices. Electron mobilities more than one order of magnitude higher than those of hitherto-investigated perovskite-type interfaces are obtained. The spinel/perovskite two-dimensional electron gas, where the two-dimensional conduction character is revealed by quantum magnetoresistance oscillations, is found to result from interface-stabilized oxygen vacancies conned within a layer of 0.9 nm in proximity to the interface. Our ndings pave the way for studies of mesoscopic physics with complex oxides and design of high-mobility all-oxide electronic devices.

1 Department of Energy Conversion and Storage, Technical University of Denmark, Ris Campus, 4000 Roskilde, Denmark. 2 Nano-Science Center, Department of Chemistry, University of Copenhagen, 2100 Copenhagen, Denmark. 3 Institute of Physics, Chinese Academy of Sciences, Beijing 100190, China. 4 Department of Physics, Technical University of Denmark, 2800 Lyngby, Denmark. 5 Center for Electron Nanoscopy, Technical University of Denmark, 2800 Lyngby, Denmark. 6 Leibniz Institute for Solid State and Materials Research, IFW Dresden, D-01171 Dresden, Germany. 7 Laboratoire de Photonique et de Nanostructures-CNRS, Route de Nozay, 91460 Marcoussis, France. 8 Center for Quantum Devices and Nano-Science Center, Niels Bohr Institute, University of Copenhagen, 2100 Copenhagen, Denmark. Correspondence and requests for materials should be addressed to Y.Z.C(email: mailto:[email protected]

Web End [email protected] ).

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DOI: 10.1038/ncomms2394

A high-mobility two-dimensional electron gas at the spinel/perovskite interface of c-Al2O3/SrTiO3

ARTICLE NATURE COMMUNICATIONS | DOI: 10.1038/ncomms2394

High-mobility two-dimensional electron gases (2DEGs) conned in epitaxially grown semiconductor heterostructures form the basis of modern electronic and photonic

devices, and have constituted the material basis for the development of quantum transport and mesoscopic physics, for example, the resultant discoveries of the integer and fractional quantum Hall effects1,2. Different from those in semiconductors, strongly correlated electrons in complex oxides with partially occupied d-orbitals give rise to a variety of extraordinary electronic properties, such as high-temperature superconductivity, colossal magnetoresistance, ferromagnetism, ferroelectricity and multiferroicity. Therefore, the high-mobility 2DEGs at atomically engineered complex oxide interfaces not only show promise for multifunctional all-oxide devices with probably even richer behaviour than that in bulk310, but would also provide a wealth of opportunities to study mesoscopic physics with strongly correlated electrons conned in nanostructures. Nevertheless, this requires a large-enough electron mobility, so that the characteristic lengths of the system, such as the mean free path or the phase coherence length, become sizeable with respect to the typical dimension of quantum devices.

The enhancement of electron mobilities for complex oxide 2DEGs, however, meets formidable challenges. To date, these 2DEGs have been fabricated exclusively at oxide interfaces between perovskite bilayers4, such as the (001)-oriented polar LaAlO3 (LAO) lms grown epitaxially on (001)-oriented non-polar SrTiO3 (STO) single crystals with a TiO2 termination3. The two-dimensional (2D) electron mobility in these perovskite-type oxide interfaces is typically B1,000 cm2 V 1 s 1 at 2 K (refs 4, 10), with a sheet carrier density, ns, being 10131014 cm 2.

This Hall mobility is still much lower than those for three-dimensional oxygen-decient STO single crystals11 and La-doped STO epitaxial lms12, amounting to 1.3 104 and 3.2

104 cm2 V 1 s 1, respectively. The 2DEGs at these perovskite-type oxide interfaces are suggested to result from electronic reconstructions due to a polar discontinuity at the interface3; however, mechanisms such as ion transfer across the interface and formation of defects have also been identied to have important roles on the transport properties13,14. Harnessing the impurities and defects at these polar complex oxide interfaces remains elusive15. Despite deliberate efforts, the highest electron Hall mobility observed in the LAO/STO-based oxide interfaces is limited to the order of 5,000 cm2 V 1 s 1 at 2 K (refs 16, 17). Besides interface polarity, we have recently found that chemical redox reactions at the oxide interface between STO single crystals and other complex oxides containing Al, Ti, Zr and Hf elements can provide an alternative approach to creating 2DEGs in complex oxide heterostructures18. Nevertheless, establishing electron connement with increased carrier mobilities in STO-based heterointerfaces remains a challenge18.

Here we present a novel 2DEG with electron Hall mobilities as large as 1.4 105 cm2 V 1 s 1 and ns as high as 3.7

1014 cm 2 at 2 K by creating a spinel/perovskite complex oxide interface between epitaxial alumina (Al2O3) lms and STO single crystals (Fig. 1a). To our knowledge, it is the rst time that complex oxide interfaces based on STO are found to exhibit carrier mobilities larger than any yet reported for either electron-doped STO single crystals11 or optimized epitaxial doped STO lms12. Moreover, such a high mobility opens the door to the design of mesoscopic quantum devices based on complex oxides.

ResultsSubunit cell layer-by-layer growth of c-Al2O3 lms. Al2O3 is a widely used oxide and is also one of the best insulating materials in nature, with a band gap normally above 8.0 eV. The synthesis

of nanoscale Al2O3 usually results in g-Al2O3 with a spinel-type structure, rather than the common a-Al2O3 with a corundum structure, because the g-Al2O3 has a lower surface energy than a-Al2O3 (ref. 19). Remarkably, as illustrated in Fig. 1bd, despite differences in cation sublattices, the oxygen sublattice of the spinel g-Al2O3 matches closely with that of the perovskite STO, as the lattice parameter of g-Al2O3 is twice that of STO (aSTO 3.905 , ag

-Al2O3

7.911 (ref. 20), lattice mismatch of1.2%). Such an excellent lattice match between oxygen sublattices, together with the low surface energy of g-Al2O3, makes it compatible to grow epitaxially g-Al2O3/STO spinel/perovskite heterostructures in a persistent 2D layer-by-layer growth mode (see Supplementary Fig. S1). Figure 1e shows typical intensity oscillations of the reection high-energy electron diffraction (RHEED) pattern during the growth of a 3-unit cell (uc) g-Al2O3 lm at a growth temperature of 600 1C. For the epitaxial growth of ionic oxides, when all lm components are supplied simultaneously, the oscillation period corresponds to the minimum unit of the chemical composition needed to ensure charge neutrality2123. For g-Al2O3 grown along the (001) direction, one intensity oscillation corresponds to the growth of one quarter unit cell lm (Fig. 1e), as the g-Al2O3 unit cell consists of four neutral AlOx subunit cells with an interlayer distance of about 0.2 nm. Similar subunit cell layer-by-layer lm growth has been observed in the epitaxial growth of spinel magnetite (Fe3O4)24. The persistent layer-by-layer, 2D lm growth results in a high-quality cubic-on-cubic g-Al2O3/STO epitaxial heterointerface with no obvious dislocations as conrmed by scanning transmission electron microscopy (STEM) (Fig.1f,g).

Electrical transport properties of c-Al2O3/STO interfaces. The investigation of conductivity in our g-Al2O3/STO hetero-structures shows that the interface between the two insulators can become metallic with electrons as the dominant charge carriers (see Supplementary Fig. S2). Of note, under the condition of our lm growth, the bare STO substrate remains highly insulating without lm deposition. More strikingly, 2DEGs with extremely high Hall electron mobilities are obtained when the g-Al2O3 lm is grown at an oxygen background pressure of 10 4 mbar and a growth temperature of 600 1C (Fig. 2ac). As shown in Fig. 2, the interfacial conduction depends critically on the thickness, d, of the g-Al2O3 lm. The heterointerface changes from highly insulating to metallic when d is above a threshold thickness of approximately 2 uc (Fig. 2d,e). At d 2 uc, the interface shows

a sheet resistance, Rs, and a carrier density, ns, in the order of 10 kO/& and 2.3 1013 cm 2 at T 300 K, respectively, similar

to the perovskite-type LAO/STO interface4,69,14,17. Remarkably, we nd a striking Rs decrease of about three orders in magnitude and a Hall mobility as high as mHallB1.1 104 cm2 V 1 s 1 at

T 2K in the spinel/perovskite g-Al2O3/STO interface. By care

fully controlling the lm growth down to a subunit cell level, a great Rs decrease of approximately four orders in magnitude is observed at d 2.5 uc, which is accompanied by the presence

of non-linear Hall resistance with respect to magnetic elds at temperatures below 100 K (see Supplementary Fig. S3). A linear tting to the low-eld Hall resistance gives rise to an impressive mHall of approximately 1.4 105 cm2 V 1 s 1, with

an ns of 3.7 1014 cm 2 at 2 K, which is consistent with those

obtained by tting the entire non-linear Hall effect within a two-band model (see Supplementary Fig. S3). Note that the high-mobility 2DEGs with mHall Z104 cm2 V 1 s 1 at T 2K are

only detected in the thickness range of 2 ucrdo3 uc. Further increasing d deteriorates the electron mobility to less than 1,000 cm2 V 1 s 1, probably due to the signicant outward

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NATURE COMMUNICATIONS | DOI: 10.1038/ncomms2394 ARTICLE

RHEED intensity

(a.u.)

Start Stop

1 uc -Al2O3

a-Al2O3

2DEG

0 200 400 600

3 2 1 0 1

Deposition time (s)

-Al2O3

-Al2O3 (001) face

SrTiO3

TiO2-STO (001) face 2aSTO

Oxygen sublattice

ADF signal

(a.u.)

O2 Ti4+ Sr2+

Al3+

2

3

Distance (nm)

Figure 1 | High-mobility 2DEGs at epitaxial spinel/perovskite c-Al2O3/STO interfaces. (a) A sketch of the heterostructure. (b) Oxygen sublattices as the backbone to build the spinel/perovskite heterostructure. The compatibility in oxygen sublattices of a g-Al2O3 surface and the TiO2-terminated

STO surface is shown in c and d, respectively. Of note, the tetrahedral cation sites in g-Al2O3 are not shown. (e) Typical RHEED intensity oscillations for the growth of a 3-uc g-Al2O3 on STO in a subunit cell layer-by-layer mode. (f) HAADF STEM image of the epitaxial g-Al2O3/STO interface.

Scale bar, 1 nm. Sr ions are brightest, followed by Ti. The faintly visible Al elements can be determined by the averaged line proles across the interface shown in g. A well-developed TiO2-AlOx heterointerface is dened.

105

2 uc 2.25 uc2.5 uc 2.75 uc

3 uc 4 uc

n s(cm2)

1016

1015

1014

1013

Hall(cm2 V1 s1)

2 uc2.25 uc2.5 uc2.75 uc

3 uc

4 uc

R s (/ )

104

103

102

101

100

-Al2O3/STO 105

104

103

102

101

102

104

106

108

10 100

T (K)

T (K)

T=300 K T=300 K

2 uc2.25 uc2.5 uc2.75 uc

3 uc

4 uc

s (1)

n s(cm2)

1015

1013

1011

109

107

101 Measurement limit

Measurement limit

10 100

0 6

4

2 8

10 100

T (K)

d (uc)

0 6

4

2 8 d (uc)

Figure 2 | Thickness-dependent electronic properties of the c-Al2O3/STO interface. (ac) Temperature dependence of sheet resistance, Rs, carrier density, ns, and low-eld electron Hall mobility, mHall, for the interface conduction at different lm thicknesses. (d,e) Thickness dependence of the sheet conductance, ss, and ns measured at 300 K. High-mobility 2DEGs are obtained at a thickness range of 2 ucrdo3 uc. The lines are guides to the eye.

diffusion of the Ti-cations across the interface as observed by electron energy-loss spectroscopy (EELS; see Supplementary Fig. S4).

2D quantum oscillations of the conduction in c-Al2O3/STO. The 2D nature of the conduction in our spinel/perovskite heterostructures is indicated by angle-dependent Shubnikov-de Haas

(SdH) quantum oscillations, which are superimposed on a huge background of positive magnetoresistance (Fig. 3a). After subtracting the magnetoresistance background, the SdH oscillations become apparent (Fig. 3b) and the extrema positions show a cosine dependence with the angle y between the magnetic eld and the surface normal (Fig. 3c). This reveals the 2D nature of the electron gas formed at our g-Al2O3/STO interfaces. Besides, the

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1.0

20

R xx(T)/R xx(22 mK)

0.6

m* = 1.22 me

0

3

T = 22 mK T = 200 mK

0.8

0.4

0.0

50

335090d =2.5 uc Hall=140,000 cm2 V1 s1

0.6

R xx()

R xx()

2

T = 400 mK T = 600 mK

d =2.25 uc

Hall=43,000 cm2 V1 s1

33

0.4

15

0.2

0.0

0

1

0.2

B=2.04 T

R xx(m)

0.0 0 200 400 600

0.10 0.15 0.20

0

R xx()

T (mK)

10

1/B (T1)

90

0

1

4.4

= 4.96 ps ( SdH=7200 cm2

V1 s1)

ln(R xxsinh( T)/4R 0 T)

0.4

50

4.8

5

2

33

0.2

5.2

3

5.6

T =200 mK

0

0 6

3 9 12 15

6.0 0.4 0.5 0.6 0.7

B (T)

B (T)

0.10 0.15 0.20

1/B (T1)

1/B cos (T1)

Figure 3 | 2D quantum oscillations of the conduction at c-Al2O3/STO interfaces. (a) Longitudinal resistance, Rxx, as a function of magnetic eld with visible SdH oscillations (arrowheads) under different tilt angle, y, at0.3 K for the d 2.5 uc sample. (b,c) Amplitude of the SdH oscillations,

DRxx, under different y versus the reciprocal total magnetic eld and the reciprocal perpendicular magnetic eld component, respectively. The SdH oscillations depend mainly on the reciprocal perpendicular magnetic eld component, particularly in the y angle of 01331, which suggests a 2D conduction nature of the g-Al2O3/STO interface.

1.0 1.5 2.0 2.5 3.0

Figure 4 | Small-eld and low-temperature behaviour of the SdH oscillations. (a) Temperature dependence of the SdH oscillations at y 01

for the d 2.25 uc sample. (b) Temperature dependence of the scaled

oscillation amplitude at B 2.04 T, giving a carrier effective mass of 1.22 me.

(c) Dingle plot of the SdH oscillations at 200 mK, giving a total scattering time t 4.96 10 12 s, a related Dingle temperature TD 0.24 K and

a consequent quantum mobility mSdH 7.2 103 cm2 V 1 s 1.

absence of oscillations at y 901 further conrms that the spatial

width of the 2DEG is smaller than at least the cyclotron radius at15 T, the typical value of which is below 10 nm for our hetero-structures. Moreover, the angular dependence of the SdH oscillations measured at high magnetic elds suggests a multiple-subband contribution to charge transport. For instance, an extra feature is observed at y 501 with Bcosy 7.2 T, which may

result from a p shift of the oscillations due to a spin-split band. Such a phase shift has been observed in the high-mobility 2DEG of GaN/AlGaN interfaces when the Zeeman energy (depending on the total B) and the cyclotron energy (depending on the perpendicular component of B) are equal25.

To conrm the high mobility achieved in our g-Al2O3/STO 2DEGs, we increased the visibility of the SdH oscillations by cooling one sample (d 2.25 uc) down to 22 mK in a dilution

refrigerator. Ultra-low noise measurements allow us to evidence the oscillations down to about 1 T (Fig. 4a), which directly shows that the quantum mobility extracted from the SdH oscillations, mSdH, is in the range of 104 cm2 V 1 s 1, as inferred from the onset of oscillations. Importantly, the low-eld dependence of the

SdH oscillations reveals the typical behaviour due to a single band. According to theory26, the oscillations amplitude DRxx can be described as:

DRxx 4R0e aTDaT/ sinhaT

where, a 2p2kB/ hoc, oc eB/m is the cyclotron frequency,

m is the carrier effective mass, kB is Boltzmanns constant and h is Plancks constant divided by 2p. R0 is the classical resistance in zero eld. TD h/2pkBt is the Dingle temperature, t is the

total scattering time. At a xed magnetic eld, m can be deduced by tting the temperature-dependent oscillation amplitude with

DRxxT/DRxxT0 T sinhaT0/T0 sinhaT(T0 22 mK). As

shown in Fig. 4b, for B 2.04 T the t leads to an effective

mass of m (1.220.03) me (me is the bare electron mass),

consistent with those reported for other STO-based hetero-structures16,2730. At a xed temperature, TD or t can be deduced from the slope of the Dingle plot, that is, lnDRxx sinhaT/4R0aT

versus 1/B (Fig. 4c for T 200 mK), which gives a

t 4.96 10 12 s or TD 0.24 K, corresponding to a quantum

mobility mSdH et/m of 7.2 103 cm2 V 1 s 1. Such an unprece

dented high mSdH in our g-Al2O3/STO 2DEGs is more than one order of magnitude higher than those observed in the perovskite/ perovskite LAO/STO16,17,27 and GaTiO3/STO30 heterostructures, which are typically below 300 cm2 V 1 s 1. Note that the difference between mHall and mSdH in our g-Al2O3/STO heterostructures could come from a different scattering time (that is, the transport scattering time and the total scattering time, respectively), which has also been reported in the LAO/STO16,17,27 and d-doped STO heterostructures28,29, as well as the GaAs/AlGaAs heterostructures31. In short, the SdH measurements support the formation of high-mobility 2DEGs at our spinel/perovskite heterointerfaces (see also Supplementary Fig. S5).

Spatial connement of the c-Al2O3/STO interface 2DEG. To determine the origin and depth-prole for the conduction in the g-Al2O3/STO heterostructures, angle-resolved X-ray photoelectron spectroscopy (XPS) measurements are performed. We nd that the electrons are exclusively accumulated on the otherwise empty 3d shell of Ti4 on the STO side. The most remarkable

XPS result is that the Ti3 signal in g-Al2O3/STO heterointer-faces shows strong dependence on the photoelectrons detection angle, j, with respect to the surface normal. An increase of the

Ti3 signal with increasing j, as shown in Fig. 5a, is clearly detected for d 2.5 uc with the highest Hall mobility. This further

conrms that the conduction in our g-Al2O3/STO heterointerface is highly conned at the interface region. To make more quantitative analyses, we assume a simple case that the 2DEG extends from the interface to a depth, t, into the STO substrate32. The interface region is further assumed to be stoichiometric, sharp and characterized by a constant fraction, p, of Ti3 per STO unit cell. Taking into account the attenuation length of photoelectrons,

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0.24

d =2.5 uc

60 50 30 0

Experimental Best fitting

Electron confinement:0.9 0.2 nm

Intensity (a.u.)

I(Ti3+)/I(Ti4+)

0.18

Ti3+

0.12

462 460 458 456 0 20 40 60

Binding energy (eV)

Emission angle (degree)

Figure 5 | Spatial connement of the 2DEG at the c-Al2O3/STO heterointerface determined by angle-resolved XPS. (a) The Ti 2p3/2 XPS spectra at various emission angles j for the d 2.5 uc sample. (b) The

angle dependence of the ratio of Ti3 to Ti4 signal, I(Ti3 )/I(Ti4 ), indicates a strong connement of the conduction layer within 0.9 nm. Error bars indicate deviations, 10%, of experimental values.

the ratio of Ti3 to Ti4 signal, I(Ti3 )/I(Ti4 ), as a function of j can be calculated as32:

ITi3 ITi4

p 1 exp t/l cos j

1 p 1 exp t/l cos j

where, l is the electron escape depth in STO. According to the NIST database (NIST Standard Reference Database 71, version1.2), l is approximately 2.2 nm for our setup. As shown in Fig. 5b, the best tting of the experimental I(Ti3 )/I(Ti4 ) ratios gives a pB0.31, which equals to an ns B2.1 1014 cm 2 and a t of

0.9 nm. Therefore, the electrons at our g-Al2O3/STO hetero-interface are strongly conned within approximately the rst 2 uc of STO surface in proximity to the interface. Note that the ns deduced here is slightly lower than that obtained from Hall data (Fig. 2c). This could be due to the presence of outward diffusion of the Ti-cations into alumina lms, where Ti4 is the dominant component (see Supplementary Figure S4). Such concern is also consistent with the fact that the out-diffused Ti is found to have a negligible contribution to the measured interface conduction. For example, the interface conduction remains unaffected when the capping alumina lm is etched away by a 4-M aqueous NaOH solution. This strongly suggests that the effective charge carriers are mainly located on the STO side.

DiscussionAs each layer of the (001)-oriented g-Al2O3/STO heterointerface is nominally charge neutral, the polar discontinuity-induced electronic reconstruction as expected in the LAO/STO interface3 may not contribute here. The presence of Ti3 is probably a signature of the formation of oxygen vacancies on the STO side.

This scenario is consistent with the fact that the interfacial conductivity can be completely removed when the Ti3 content is signicantly suppressed by suitable annealing in 1 bar pure O2 at a temperature higher than 200 1C (see Supplementary Fig. S6). Such an oxygen-vacancy-dominated 2DEG is expected to be formed as a consequence of chemical redox reactions occurring on the STO surface during the lm growth of g-Al2O3, analogous to what has been observed in metallic amorphous STO-based heterostructures grown at room temperature18. Note that the

2DEG at the crystalline g-Al2O3/STO heterointerface is formed at a high temperature of 600 1C, where the oxygen ions in STO are already highly mobile. This is normally expected to level out any difference in the depth-prole of oxygen distribution in STO18,33. However, this is not the case in the crystalline g-Al2O3/STO heterostructures as inferred from both Figs 3 and 5. Moreover, the conduction at the interface of thick lms, for example, at d 8 uc, can survive the annealing at 300 1C for 24 h in 1 bar

pure O2 with only negligible changes in the conductivity (see Supplementary Fig. S6). These features strongly suggest that the oxygen vacancies and the 2DEGs are stabilized by an interface effect, such as by the formation of a space charge region near the heterointerface. It is worth noting that an inherent oxygen ion deciency has been observed at the grain boundary of STO bicrystals34, where a considerable electron accumulation has also been predicted if the barrier height of the grain boundary is deliberately controlled35. The high electron mobility of STO-based oxide materials at low temperatures is generally related to the polarization shielding of the ionized defect scattering centres driven by the large dielectric constant of STO36. The higher mobility of our spinel/perovskite oxide interface compared with the perovskite-type oxide heterointerface may be due to the better lattice match and, thereby, a more perfect structure and well-dened interface. Though further investigations are needed to reveal how the interface properties increase the mobility and the associated strong suppression of the defect and impurity scattering, our results strongly suggest that defect engineering of oxygen vacancies is crucial for the high mobility of 2DEGs conned at the interface between complex oxides.

In conclusion, we have demonstrated that high-mobility 2DEGs with clear quantum magnetoresistance oscillations and strong spatial connement can be created at well-dened spinel/ perovskite g-Al2O3/STO oxide interfaces. The strongly spatial connement of charge carriers achieved directly in the as-deposited spinel/perovskite oxide heterostructures without any post annealing provides the possibility to fabricate multilayers of complex oxides with several 2DEGs. Furthermore, by combining two of the largest groups of oxides, plenty of new physical properties, for instance, interfacial magnetism6 and super-conductivity7 as observed in the perovskite-type LAO/STO interface, may be found at the g-Al2O3/STO heterointerface.

Finally, with a large enhancement of the electron mobility, the g-Al2O3/STO heterointerface probably enables the design of mesoscopic quantum devices based on complex oxide 2DEGs and opens new avenues for oxide nanoelectronics and mesoscopic physics.

Methods

Sample growth. The g-Al2O3 thin lms were grown by pulsed laser deposition37 using a KrF laser (l 248 nm) with a repetition rate of 1 Hz and laser uence of

1.5 J cm 2. The target-substrate distance was xed at 5.6 cm. Commercial a-Al2O3 single crystals were used as targets. Singly TiO2-terminated (001) STO crystals with a size of 5 5 0.5 mm3 were used as substrates. Note that the TiO2 termination

of our substrates is obtained by chemical etching using HCL-HNO3 as acidic solution37, which is found to produce less defects on the STO surface compared with the conventional buffered hydrouoric acid etch method38,39. The lm growth process was monitored by in-situ high pressure RHEED. During deposition, the oxygen pressure was xed at 10 4 mbar with the deposition temperature changing from room temperature (20 1C) to 700 1C. After lm deposition, the samples were cooled down to room temperature at the deposition pressure. The lm thickness was determined by both RHEED oscillations and X-ray reectivity measurements.

Electrical transport measurement. The transport properties of the buried interface were measured using a four-probe Van der Pauw method, with ultrasonically wire-bonded aluminium wires as electrodes, placed at the corners of the square sample. The temperature-dependent electrical transport and Hall-effect measurements were performed in a CRYOGENIC cryogen-free measurement system, with the temperature ranging from 300 K down to 2 K and magnetic elds up to 16 T. To conrm the carrier density and mobility, some Hall-bar patterned samples were

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also measured, which were prepared directly through a mechanical mask18. Note that the use of a mechanical mask at deposition temperatures higher than 500 1C may have a deleterious effect on the carrier mobility, as the high oxygen ion diffusion can unintentionally disturb the oxygen equilibrium for realizing high mobility. The angle-dependent SdH measurements were performed in a sorption-pumped 3He cryostat with standard lock-in technique at 0.3 K, with magnetic elds up to 15 T by changing the angles manually. The temperature-dependent SdH measurements were performed in a dilution refrigerator with a base temperature of 22 mK and an improved temperature stability, using ultra-low noise electronics. During all the transport measurements, the applied currents were within 110 mA (for AC current, the frequency was 327 Hz). Special care was taken to avoid heating effect.

XPS measurement. The XPS measurements were performed in a Kratos Axis UltraDLD instrument, using a monochromatic Al Ka X-ray source with photon energy of 1,486.6 eV. This leads to a kinetic energy of Ti 2p electrons of roughly 1,025 eV. According to the NIST database(NIST Standard Reference Database 71, version 1.2), the electron escape depth is approximately 22 in STO at this kinetic energy. The pass energy used for the high resolution scan was 20 eV. The detection angle of the electrons varied between 01 and 601 with respect to the sample normal. For analysing the Ti 2p3/2 peaks (Ti4 is at a binding energy of 459.5 eV, whereas the Ti3 is 1.6 eV0.1 eV lower), a Shirley background was subtracted and the spectra were normalized to the total area below the Ti peaks ([Ti] [Ti4] [Ti3 ] 100%).

STEM and EELS measurements. Aberration-corrected STEM measurements were performed by an FEI Titan 80300ST TEM equipped with a high brightness Shottky emitter (XFEG) and a Gatan Image Filter (Tridiem). High-angle annular dark eld (HAADF) images were acquired at 300 kV, where the probe size, convergence angle and HAADF collection angle were 0.81 , 20 mrad and 46291 mrad, respectively. For EELS in the STEM, an accelerating voltage of120 kV (probe size of 1.52.0 ) was used to reduce knock-on damage to the specimen. The energy resolution of EELS was B0.9 eV. Spectrum imaging was used to collect spectra across the interface. We typically recorded the spectrum images consisting of 40 ten-analysis point lines (that is, 10 40 pixel) parallel to

the interface and acquired each line by an increment of 0.28 nm. Each spectrum was obtained at a dispersion of 0.1 eV for 0.20.4 s. Then the spectra along the lines were summed after removing the spectra from beam-damaged regions according to the HAADF contrast to increase signal/background ratio.

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Acknowledgements

We thank J. Fleig, F. W. Poulsen, N. Bonanos, S. Stemmer and Y.Q. Li for helpful discussions. We also thank K. Thydn, Z.I. Balogh, J.W. Andreasen, E. Johnson,Y. Zhao, X. Tang, W.W. Gao, N.Y. Wu, J. Geyti, K.V. Hansen, K. Engelbrecht and L. Theil Kuhn for their help.

Author contributions

Y.Z.C. contributed to the concept design, lm growth, transport measurements, data analysis, interpretation and writing of the manuscript. N.P. and S.L. contributed to the concept design. N.B. contributed to the XPS measurements and analysis. F.T., D.V.C., N.H.A. and T.S.J. contributed to the transport measurements and analysis. F.M.Q, R.G. and J.D. contributed to the SdH measurements and analysis. T.K. contributed to the STEM and EELS measurements and analysis. W.Z. contributed to the HRTEM measurements and analysis. J.R.S., A.S., J.N., L.L., B. B. and B.G.S discussed the data. All authors extensively discussed the results and the manuscript.

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How to cite this article: Chen. Y. Z. et al. A high-mobility two-dimensional electron gas at the spinel/perovskite interface of g-Al2O3/SrTiO3. Nat. Commun. 4:1371 doi: 10.1038/

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