Supplementary Information

Strain engineering induced interfacial self-assembly and

intrinsic exchange bias in a manganite perovskite film

B. Cui, C. Song,* G. Y. Wang, H. J. Mao, F. Zeng, and F. Pan*

Key Laboratory of Advanced Materials (MOE), School of Materials Science and

Engineering, Tsinghua University, Beijing 100084, China.

*E-mail: songcheng@mail.tsinghua.edu.cn; panf@mail.tsinghua.edu.cn.

Microstructure of LSMO grown on LaSrAlO4 and SrTiO3 substrates

Electron energy-loss spectroscopy (EELS) La-M4,5 is recorded for the region

shown in the yellow frame of Fig. 1b for the 45 unit cell (u.c.) La2/3Sr1/3MnO3

(LSMO) film grown on LaSrAlO4 (LSAO) substrate, using the same method as

collecting Mn-L2,3 signals discussed in the main text. Obviously, the content of La in

area (1) (d = 0.8 and 2 nm) is much larger than other regions according to its higher

peak intensity. This finding is in accordance with the results of EDX in Fig. 1e. The

shoulder peak marked as “I” in Fig. S1b is attributed to overlapping bands of O-2p

and Mn-3d character (from d = 3.2 nm to d = 6 nm). The absence of this shoulder

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peak also bolsters the shortage of Mn in the self-assembled LaAlO3-like buffer (d =

0.8 and 2 nm). On the other hand, the observation of a protrusion at 541.3 eV at d =

3.2 nm (denoted as “II”) provides clear evidence for the existence of LaSrMnO4,

which is similar to the previous near-edge x-ray absorption fine structure (NEXAFS)

spectrumS1. Therefore, these characteristics in La- and O-EELS strongly support the

formation of three sublattices in LSMO by self-assembly.

FIG. S1. EELS La-M4,5 a and O-K b absorption edges for the region shown in the

yellow area of Fig. 1b. The distance (d) in EELS is the space of the probing place and

substrate surface, and the data are the average results of ±1 u.c. range. The features

“I” and “II” are relevant to overlapping bands of O-2p/Mn-3d and the existence of

LaSrMnO4, respectively.

2

Soft X-ray absorption spectroscopy (XAS) measurements were performed in total

electron yield (TEY) mode at the Beamline BL08U in Shanghai Synchrotron

Radiation Facility. The polarization directions of the linearly polarized x-rays are

tuned by rotating the x-ray incident angle, with 90° and 30° corresponding to

complete in-plane (E//a, Iab) and majority of out-of-plane (E//c, Ic), respectively. X-ray

linear dichroism (XLD), the difference between the two measurements (Iab–Ic), gives

direction insight of the empty Mn 3d statesS2,S3. Given that the probing depth in TEY-

XLD experiments is exponentially decaying within a few nanometers, the 45 u.c.

LSMO film was etched ~7 nm [Etching 1, in area (3)] and ~14 nm [Etching 2, near

area (2)] to obtain the information of corresponding area (Fig. S2a). For the surface of

sample, the preferred orbital ordering stays in 3z2–r2 orbital.S4 As the probe depth

increases, the XLD spectrum at Etching 1 (Fig. S2c) indicates more in-plane empty

states in the eg-band than the surface one, due to the larger compressive strain in the

LSMO bulk, as will be discussed below. For Etching 2, the unique shape of XAS

curves with a prominent peak at 641 eV (Mn L3 edge) in Fig. S2d suggests a distinct

orbital ordering in this scenario. Meanwhile, it is interestingly found that the XLD

signal is markedly distinct from the area above, whose feature is quite characteristic

for XLD of LaSrMnO4 (Refs. S2 and S5). This finding supplies additional evidence

for the existence of LaSrMnO4 in area (2), also suggesting that the film at this

thickness exhibits a 3z2-r2 orbital ordering, which extends perpendicular to the MnO2

sheetS2,S6. The resultant decrease of itineracy of the electron weakens the double-

exchange interaction, which would induce an antiferromagnetic phaseS6. Moreover,

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rather low intensities for both XAS and XLD reflect the ultrathin LaSrMnO4 layer

(only one cell thick) and the lacking of Mn in area (1), as discussed in the main text.

On the contrary, the XLD signal of LSMO grown on SrTiO3 implies that electron

occupies the in-plane x2–y2 orbital mainly due to its tensile strain.

FIG. S2. a, Schematic of experimental configuration for XAS measurements with

different x-ray incident angles. The height of every step produced by etching is

around 7 nm. XAS and XLD at Mn L2,3 edges of surface b, Etching 1 c, Etching 2 d

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states for 45 u.c. LSMO grown on LaSrAlO4 as well as surface case of 45 u.c. LSMO

on SrTiO3 e. For clarity, the intensities of XLD are multiplied 5 times as marked in the

figure.

FIG. S3. a, Representative HAADF-STEM image of the interface area for 45 u.c.

LSMO grown on SrTiO3 (STO). Elemental mappings by EDX measurement for the

area of HAADF-STEM image b are shown in c, La (red), d, Sr (green), e, Mn (blue),

and f, Ti (violet). The white scale bars in b–f represent a length of 4 nm. Note that the

“appearance” of La in SrTiO3 substrates is due to the large width and high intensity of

Ti signals, which interferes the EDX of La (La Kα1,2: 4.6510 eV; Ti Kα1,2: 4.5089 eV).

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Figure S3a displays an atomic resolution high-angle annular dark field scanning

transmission electron microscope (HAADF-STEM) image (also known as Z-contrast)

of the LSMO/SrTiO3 interface for a 45 u.c. sample grown on SrTiO3 (001) substrates.

High quality epitaxial and coherent growth of LSMO films is achieved on SrTiO3, in

sharp contrast to the films deposited on LaSrAlO4. Both HAADF-STEM image and

energy dispersive x-ray spectroscopy (EDX in Fig. S3c–f) confirm that there is no

phase separation within the sensitivity of our measurements. Such a uniform film

should be relevant to the small mismatch between LSMO and SrTiO3.

FIG. S4. a, HRTEM image of LSMO/LaAlO3 cross-section. b, HAADF-STEM image

of the area for EDX measurements, as well as c, La (red), d, Sr (green), e, Mn (blue),

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and f, Al (cyan) elemental applying by EDX. The scale bars in b–f represent a length

of 20 nm.

Figure S4a displays a high resolution transmission electron microscopy (HRTEM)

image of 30 u.c. LSMO grown on LaAlO3(001) substrates. Although the epitaxial

quality of the LSMO film is not as good as that on SrTiO3 and the lattice distortion is

recognizable, we do not observe the layered phase separation as that on LaSrAlO4,

similar to the scenario reported beforeS7. Also, there is no clear composition

separation, as indicated by EDX characterizations in Fig. S4c–f. Hence the EB

behavior likely arises from the coupling between FM LSMO and previously proposed

distorted AFM LSMO interfacial dead layerS7, but the strain between LSMO/LaAlO3

is too weak to induce the delicate interfacial self-assembly.

Magnetization comparison for LSMO on SrTiO3 and LaSrAlO4 substrates

FIG. S5. Magnetic loops measured at 5 K after field cooling from room temperature

in +20 kOe (solid spheres) and –20 kOe (open circles) for the 45 u.c. LSMO film

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grown on SrTiO3 (001) substrates. The growth parameters and measurement details

are identical with those using LaSrAlO4 substrates, but no EB behavior is observed.

FIG. S6. Normalized magnetic loops (M–H) for 45 u.c. LSMO grown on LaSrAlO4

substrates measured at different temperatures (5, 25, 50, 75, and 100 K). The region

near the origin of the loop is blown up to show the extent of exchange bias clearly.

The enhancement of temperature weakens the EB behavior, which is summarized in

Fig. 2b. For clarity, only the data between –4.5 and 4.5 kOe are shown in this figure,

while the actual measurements were carried out between –20 and +20 kOe.

As discussed in the main text, in contrast to symmetric hysteresis loops of LSMO

deposited on SrTiO3 (Fig. S5), Fig. S6 shows the exchange bias behavior in LSMO

grown on LSAO, which can be explained by the exchange coupling between self-

assembled layers. Nanoscale lateral phase separation, the coexistence of

ferromagnetic (FM) metal and charge-ordered antiferromagnetic (CO/AFM) insulator,

is common in manganites, which possibly exists in our samples. However, we note

that the exchange bias effect in our self-assembled tri-layers should not be ascribed to

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unintended lateral phase separation: (i) It is generally accepted that the lateral phase

separation increases dramatically when the temperature is enhanced to the charge-

ordering temperature, accompanied by colossal magnetoresistanceS8. For instance, it is

demonstrated that the glassy mixed-phase is energetically favorable at high

temperaturesS5,S9. Hence the separation tendencies appear stronger with increasing

temperatures, which would strengthen the coupling between FM and AFM phases.

Apparently this tendency is opposite to our experiment observation, i.e., the bias field

decreases with increasing temperature, and vanishes at around 110 K. (ii) Although

the lateral phase separation has been detected by several techniquesS10,S11, few

publications provide direct information for the exchange bias arising from this

separation based on magnetization characterization with spatial resolution.

Particularly, it is found that the exchange coupling was observed in AFM-based

manganites with FM nanodomains immersed in the AFM backgroundS12–S14, rather

than extensively studied ferromagnetic systems with AFM clusters imbedded in FM

matrix. Obviously, the major phase in our samples is ferromagnetic according to the

chemical component (La = 2/3 and Sr = 1/3). (iii) Based on the so-called dead layer

theory, manganite films under a critical thickness are AFM insulator and the FM

metal phase dominates gradually with the increase of film thickness associated with

the drop of HEB. The film thickness dependent HEB in Fig. 4a is consistent to this

scenario, revealing that the AFM phase located at the bottom of the film, instead of

lateral distributed one, primarily contributes to the exchange bias.

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Microstructure of LSMO films with different thicknesses

FIG. S7. a, Cross-sectional HAADF-STEM image of 7 u.c. LSMO/LaSrAlO4. b,

HAADF-STEM image of the area for EDX measurements, as well as c, La (red), d, Sr

(green), e, Mn (blue), and f, Al (cyan) elemental applying by EDX. The white scale

bars in b–f represent a length of 1 nm.

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In order to unravel the growth dynamics of LSMO, we investigate a series of

LSMO films with different thicknesses deposited on LaSrAlO4. The cross-sectional

image of 7 u.c. sample is representatively shown in Fig. S7a. The lattice of the films

follows the typical lattice of LSMO, indicating that the films initially behave as real

LSMO without the formation of buffer layers. Meanwhile, the uniform distribution of

all the elements in EDX (Fig. S7c–f) reflects the homogenous LSMO films. This

finding differs dramatically from the scenario of 45 u.c. sample deposited on the same

substrates, as displayed in Fig. 1. This growth mode might be correlated to the space

limitation for element diffusion and the formation of AFM in several monolayers. In

addition, a closer inspection of the image shows that the film is fully strained because

of large film/substrate mismatch.

X-ray diffraction (XRD) patterns concerning (002) diffraction peak of LSMO

grown on LaSrAlO4 substrates with different thicknesses are shown in Fig. S8a,

obtained by the conventional theta-2theta scanning, providing the information along

c-axis. On the other side, the LSMO (111) peak in Fig. S8b was obtained by azimuthal

scanning, which could be used to figure out the in-plane lattice parameterS15. For 30

u.c. sample, in-plane compressive strain elongates the c-axis lattice parameter 3.98 Å,

much larger than its bulk counterpart of 3.87 Å, ascribed to the effect of elastic

energy. Note that the (002) peaks gradually shift to higher angles with the increase of

thickness from 30 u.c. to 150 u.c., reflecting that the c-axis lattice parameter

approaches to its bulk, as illustrated in Fig. S8c, where the lattice parameters (both a

and c) and volume calculated from the results in Fig. S8a and b are shown as a

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function of film thickness. Meanwhile, a-axis lattice parameters are persistently below

the bulk value because of the large compressive strain, which is finally close to that of

bulk for 150 u.c. sample. It is noted that, for the samples thicker than 30 u.c., the

diffraction peaks of LAO-like and LaSrMnO4 overlap in the peaks of LSMO and

substrates due to the continuous transition of structure and parameter. The lattice

parameters and calculated the volume are the average results of the whole films.

FIG. S8. XRD patterns for (002) a and (111) b planes of LSMO with different

thicknesses grown on LaSrAlO4. For clarity, the intensities in b are multiplied by a

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coefficient as shown except for the 150 u.c. sample. c, Thickness dependence of the

lattice parameters and volume calculated from the results from a and b. The

corresponding bulk values are shown by dashed lines for a comparison.

We then discuss the anomalous XRD results and corresponding lattice volume in

ultrathin 7 u.c. sample. At the initial growth stage (7 u.c.), both a- and c-axis lattice

parameters of LSMO are obviously smaller than their bulk counterparts, producing

the small lattice volume (53.4 Å3) compared to its bulk (57.9 Å3). We have not yet

been able to determine the origin of such anomaly, which might be caused by Al

penetration from the substrate as detected by the EDX in Fig. S7f. The self-assembled

layers might start to appear in our 30 u.c. sample, indicated by the EB effect, leading

to the release of the interfacial strainS16 and the sharp transition of c-axis lattice

parameter (Fig. S8c).

Magnetic properties of LSMO grown on LaAlO3 and SrTiO 3 substrates

To obtain more information about the effect of strain on the properties of LSMO, we

also deposited LSMO films on LaAlO3 (001) and SrTiO3 (001) substrates with an in-

plane lattice of 3.789 and 3.904 Å, respectively. The EB effect was also observed in

the LSMO films grown on LaAlO3 substrates (Fig. S9a), but with bias fields much

smaller (e.g., ~140 Oe for the 45 u.c. sample) than the ones with LaSrAlO4 substrates.

This comparison suggests that the EB behavior could be tuned by interface

engineering by an intentional strain design. The ferromagnetism of 7 u.c. sample is

greatly suppressed (its magnetization is even multiplied by 5 times in Fig. S9b),

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indicating that there is a nonmagnetic interface in this system, in analogy with

previous publicationsS17,S18. We then deduce that the EB effect most likely originates

from the coupling of ferromagnetic LSMO and interfacial dead layer just above

LaAlO3, as discussed above. When the SrTiO3 substrate is employed, the EB effect

cannot be observed any more irrespectively how thick the LSMO film is (Fig. S9c), as

mentioned in Fig. S5. The temperature dependent magnetization measured in 7 u.c.

sample shows the same qualitative behavior as the thicker ones, but with a reduced

overall magnitude (Fig. S9d), suggesting that the influence of the so-called dead layer

on the magnetic property of LSMO on SrTiO3 is limited, because of a small

film/substrate mismatch. These results imply that the dead layer and magnetic

properties of LSMO films could be effectively tailored by interface engineering.

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FIG. S9. Normalized M–H curves measured at 5 K after field cooling from room

temperature in 20 kOe and M–T curves for 7 u.c., 30 u.c., and 45 u.c. thick LSMO

grown on LaAlO3 (a and b) and SrTiO3 (c and d) substrates, respectively. For clarity,

the magnetization of 7 u.c. samples in b and d are multiplied by 5 times.

Details for experiment method

Let us discuss the possible influence of LaSrAlO4 substrate treatment on the

formation of LaAlO3-like layer just above the film/substrate interface. LaSrAlO4

substrates were etched in NH4F-HF solution and annealed in an ultrahigh vacuum

with a base pressure of 10–9 Torr at 700 oC for 1 hour before film deposition. The

chemical treatment and the high temperature annealing, which were confirmed to

induce the loss of surface La in LaSrAlO4 (Ref. S19), possibly produce AlO2-

dominated surface, in analogy to our TiO2 terminated SrTiO3. This pre-condition

might benefit the formation of LaAlO3-like layer through the diffusion of Al from the

substrates at the high growth temperature.

The growth dynamics of LSMO films was investigated by monitoring the

intensity variations of RHEED (reflection high-energy electron diffraction) patterns.

Figure S10a shows clear RHEED oscillations recorded during the growth of LSMO

on TiO2 terminated SrTiO3, suggesting a growth rate of 1.16 nm/min. This also

provides the thickness information for the simultaneous grown LSMO films on

LaSrAlO4 and LaAlO3. The inset is the RHEED patterns at two typical stages of

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growth, which is well streaky. The high-quality of LSMO film is also confirmed by

atomic force microscopy morphology with atomically flat terraces as defined by the

underlying vicinal SrTiO3 substrates and separated by steps with nearly one unit cell

step height (Fig. S10b).

FIG. S10. a, RHEED intensity oscillations and patterns (the inset) of LSMO

grown on TiO2-terminated SrTiO3 substrates. b, The morphology of LSMO on

STO with a thickness of 30 u.c.

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