Voltage-Controlled Oxygen Non-Stoichiometry in SrCoO3âÎ´ Thin

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Voltage-controlled oxygen non-stoichiometry in SrCoO3-# thin films Songbai Hu, Wenqiao Han, Sixia Hu, Jan Seidel, Junling Wang, Rui Wu, Jiaou Wang, Jiali Zhao, Zedong Xu, Mao Ye, and Lang Chen Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.9b01502 • Publication Date (Web): 29 Jul 2019 Downloaded from pubs.acs.org on August 3, 2019

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Chemistry of Materials

Voltage-controlled oxygen non-stoichiometry in SrCoO

thin films

Songbai Hu*,†, Wenqiao Han†, Sixia Hu‡, Jan Seidel§, Junling Wang , Rui Wu , Jiaou Wang , Jiali Zhao , Zedong Xu†, Mao Ye†, Lang Chen*,†

†Department

of physics, Southern University of Science and Technology, Shenzhen 518055,

China ‡Materials

Characterization and Preparation Center, Southern University of Science and Technology, Shenzhen 518055, China §School

of Materials Science and Engineering, UNSW Sydney, Sydney NSW 2052, Australia

School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore Laboratory of Synchrotron Radiation, Institute of High Energy Physics, Chinese Academy of Sciences, Beijing 100039, China.

Abstract: Oxygen non-stoichiometry plays a critical role in determining the physical and chemical functionalities of oxide materials. For wide spread applications involving oxygen transport and exchange with the environment, fast, inexpensive and reversible control of oxygen deficiency is highly desired. This article illustrates voltage-controlled oxygen nonstoichiometry in SrCoO6>A (SCO) thin films, in which the oxygen deficiency ( ) can be tuned between 0 and 0.5 within tens of seconds by a small applied voltage (< 1.7 V). Correspondingly, its magnetism as well as the electrical and optical properties can be tuned accordingly from one end to the other, making it a good candidate for a number of commercial applications, such as oxygen capacitors, catalysts and smart windows etc. This approach can be used as an effective method in imaging the phase diagrams of transition metal oxides, such as ternary ABO6>A (A = Ln, Ca, Sr, Bi; B = Cr, Mn, Co, Fe, Ni), or binary TiOx, WOx, VOx and NiOx etc., paving an avenue to the searching for novel properties in redox materials.

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Oxygen deficiency plays a critical role in determining the physical and chemical functionalities of oxide materials. The static and dynamic properties of oxygen non-stoichiometry switch not only the valence state of the metal ions, but also affect the crystal structures, oxygen defects or ordering. The confluence of these effects can urge metal-insulator transitions and also modify magnetic, electrical and optical properties.1-5 This advantage makes redox oxides good candidates for numerous essential applications, such as gas separation/sensing6-7, electrocatalysis/photocatalysis8-10, solid oxide fuel cells11-12, water splitting13, lithium ion batteries14-17, resistive random-access memory (RRAM)3, 18 or superconductivity19-21, among others. As vacancies are typically the dominant factors leading to enhance working performances, fast and easy control of oxygen deficiency is a critical property for technologies that involve oxygen transport and exchange with the environment. Traditionally, the modulation of oxygen deficiency is realized by high temperature annealing22-23, strain engineering24-25, or chemical oxidation/reduction26-27. Recent works have shown that at room temperature, the vacancy concentration can instead be changed by employing an electrical bias through ionic switching, 2 ACS Paragon Plus Environment

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but they are inherently slow.28-29 Thus, fast, inexpensive and reversible techniques are highly desired for wide-spread applications. In this context, we demonstrated a voltage-controlled oxygen non-stoichiometry in strontium cobaltate thin films in which the oxygen deficiency was tuned between 0 and 0.5 within tens of seconds by a small applied voltage (< 1.7 V). Strontium cobaltate, i.e. SrCoO6>A (SCO, 0< 0.5), undergoes a number of topotactic phase transitions driven by oxygen vacancies.22, 30 Its structural, magnetic, electrical and optical properties present a strong dependence on the oxygen non-stoichiometry.5, 31 At 0 K A K 0.25 SCO adopts a simple cubic perovskite structure (P-SCO, a=3.824 Å for SrCoO330) that is metallic. Bulk stoichiometric SCO is ferromagnetic below 305 K32 and its Curie temperature, Tc, decreases linearly with AA is antiferromagnetic and its Néel temperature amounts to 570 K.33 Bulk BM-SCO is orthorhombic with lattice constants of a = 5.5739 Å, b = 5.4697 Å, and c = 15.7450 Å, which can be viewed as pseudo-tetragonal (at = 3.905 and ct/2 = 3.936 Å). 22, 34 We showed that SCO could be switched continuously in in-plane or out-of-plane direction by either step or constant applied electric field, resembling a variable magneto-ionic switch, in which magnetic properties were controlled by electric fields via ionic movement29. The crystal structures and related properties can be tuned accordingly from one end to the other. Such a continuous change provides possibilities in imaging the full map of phase diagrams for redox metal oxides, such as ternary ABO6>A (A = Ln, Ca, Sr, Bi; B = Cr, Mn, Co, Fe, Ni), or binary TiOx, WOx, VOx and NiOx etc., and optimizing the oxygen content for enhanced working performances.

RESULTS AND DISCUSSION High quality SCO thin films were deposited on (001) SrTiO3 (STO) substrate by pulsed laser deposition (PLD). The thickness of the film was estimated to be 10 nm after 1200 pulses of deposition according to transmission electron microscopy (TEM) cross-sectional image (see figure S1). The structures, strain state and surface topography of the thin films were characterized by XRD and AFM, respectively. Figure 1 (a) shows the XRD 2 scans for the as-deposited thin films. Except those peaks near the STO substrate, half order peaks labelled as (002) & (006) were detected as well, indicating a BM structure SCO.22 The crystallinity of the thin film was investigated by XRD rocking curves (see figure S2). The full width at the half maximum (FWHM) for BM-SCO thin film on STO substrates was approximately 0.05o, indicating a good crystallinity for this 10 nm thin film. The strain state of the thin films was investigated by reciprocal space maps (RSMs), which was shown in figure 1 (b). No horizontal __ __ shift was observed between STO (103) and SCO ( 1 03), demonstrating that our thin films were fully strained by the substrates and no relaxation between them. It is worthy to mention that the strain plays an important role in controlling the oxygen diffusion kinetics and thereby affects significantly the redox reaction. Our previous work investigated thoroughly the oxygen diffusion kinetics in SCO under broad epitaxial strain conditions (-1.2% ~ +3.9%).25 It is found 3 ACS Paragon Plus Environment


that moderate tensile strains of ~+2% provide the optimal conditions for reversible oxygen uptaken and liberation; while under moderate compressive (-1.2%) or large tensile (3.9)% strains the oxygen liberation is highly favourable but rather than the reverse up-taken. Here STO was selected as the substrate because it provides a tensile strain of 1.8% where exhibits the best reversibility. From the AFM topography in figure 1 (c) the sample surfaces show a step-andterrace for SCO thin films. The profile of each single step gives a 0.75 nm height difference which is quite close to the d-spacing of the BM-SCO (002) (see figure S3).

(a)

(c)

0.79 0.78 STO (-103)

qz (Å-1)

BM-SCO (006)

BM-SCO (008) STO (002)

(b) BM-SCO (004) STO (001)

BM-SCO (002)

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0.77 0.76

BM-SCO (-103)

0.75

10

20

30

40

-0.26 -0.25 qx (Å-1)

50

2 (o) Figure 1. (a) XRD substrates.

0.5 m

scan, (b) RSM and (c) AFM of BM-SCO thin films grown on STO

In as-grown samples the oxygen stoichiometry was determined to be SCO2.5 as the half order peaks, i.e., BM-SCO (002) & (006) which represent the alternating stack of oxygen octahedra and tetrahedra, can be seen from the XRD.22 As long as the oxygen content exceeds 2.75 the BM to P phase transition occurs35. In perovskite SCO the linear decrease of lattice constants on oxygen stoichiometry has been established by Taguchi H. et al.31 and Takeda Y.30 et al. decades ago, and confirmed by Jeen H. et al.1, Petrie J. R. et al.24 and Hu S. et al.25 recently. For nearly stoichiometric SrCoO2.95 the lattice constant was found to be 3.8289 Å.32 That means the lattice constants of bulk SCO decrease from 3.845 Å at SrCoO2.75 to 3.8289 Å at SrCoO2.95. With these numbers, the oxygen content in SCO can be estimated through the d-spacing. The evolution of oxygen deficiency in SCO was studied by monitoring the redox reactions in-situ using X-ray diffraction (XRD) as a function of voltage. The results are given in figure 2. Figure 2 (a) shows the XRD peak evolution of SCO thin film by increasing the voltage from 0.5 V to 1.7 V. The black dashed line denotes the d-spacing of SCO2.75 (002) on STO. It was estimated to be 3.837 Å from the bulk 3.845 Å30 by excluding the strain effect using a typical Poisson’s ratio of 0.3 for perovskite oxides36. As indicated by XRD, a small shoulder on the right of the STO peak starts to grow immediately after 0.5 V. In figure 2 (b) a plateau is captured by the current-time (I-t) mapping between 0.5 V and 0.7 V, which implies a drastic oxygen up4 ACS Paragon Plus Environment

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taken at the initial. Meanwhile, the resistivity of the thin film in figure 1 (c) begins to drop. That means with a small voltage even as low as 0.5 V the thin film could be oxidized to P-SCO. However, the expansion of the shoulder slows down after 0.7 V, implying a moderate oxygen up-taken. The corresponding current flow in figure 2 (b) drops to the 6U level again and starts to accelerate thereafter. Interestingly, the resistivity of the thin film shows little change in the range of 0.7 V ~ 1.5 V. This is because the overall conductivity of the thin film is limited by BM-SCO which manifests a high resistivity. Actually, the oxygen up-taken occurs firstly in the vicinity of the electrodes, and then pushes on radically to the middle in in-plane direction. Such a mechanism can be identified clearly by the colour change in large samples, e.g. 10×10 mm2. As the BM-SCO narrows down in the middle, the P-SCO at the two sides links up together at 1.55 V and the full conversion of BM to P phase is accomplished. Simultaneously, a sharp increase of current and drop of resistivity were observed in figure 1 (b) and (c), respectively. As the voltage increases, the d-spacing decreases step by step, indicating that the oxygen deficiency can be reduced linearly for A K 0.25. We noticed that the BM to P phase transition occurs as the oxygen content exceeds 2.75.26 At 1.7 V, the d-spacing of SCO reduces down to 3.77 Å in out-of-plane direction. It corresponds to 3.8 Å in bulk by excluding the strain effect using a typical Poisson’s ratio of 0.3 for perovskite oxides36. Such a small c parameter has not been reported yet in other preparation methods, indicating an extremely low oxygen deficiency in SCO. Further increase of voltage could not decrease the d-spacing any more, which means the oxygen content has been saturated in the thin film. As the lattice constants decreases a lot more than from SCO2.75 to SCO2.95 it is postulated that the thin film is almost fully oxidized. The corresponding RSM and AFM of saturated P-SCO is given in figure S4. The film remains fully strained as before, although the surface gets roughened. The reversible control of oxygen deficiency was achieved by a negative voltage. The results are given in figure 1 (d) ~ (f). The oxygen liberation occurs immediately at the loading of the voltage by looking at the shrinking P-SCO (002) peak in figure 2 (d) and the large current in figure 1 (e). The oxygen saturated SCO depleted quickly at voltages of even less than -0.1 V by noticing the d-spacing increases from 3.77 Å to 3.812 Å. Meanwhile, a leap of resistivity was detected in figure 1 (f). Such a trend persists till the end of the phase transition. Contrary to figure 1 (c), it shows no plateau here. This is probably due to the P-SCO depletes firstly from the top and the bottom layer works as an electrode. In other words, the oxygen liberation occurs in out-of-plane direction. Such a direction-selective reaction mechanism offers an effective method for the 3D interfacial control at the nanometre scale for novel oxide devices.



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BM-SCO (008)

Voltage (V)

(a)

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10-2

10-1

100

101

0.0 -0.2 -0.4 -0.6

Figure 2. Control of oxygen vacancies in SCO by linearly increasing voltage. The XRD peak evolution of SCO as a function of positive and negative voltage are given in (a) and (d), respectively. T The black dashed line denotes the d-spacing of SCO2.75 (002) on STO. The insitu I-t mapping during the redox reaction are recorded in (b) and (e). The resistivity of the thin film after each step are given in (c) and (d) as well. The oxygen deficiency in the thin film can be tuned between 0 and 0.5 easily and reversibly by linearly increasing voltage.

In figure 3 we tried to control the oxygen deficiency by constant voltage. It was set to 0.7 V/-0.05 V for oxygen up-taken/liberation, respectively. Comparing with the linearly increasing voltage in figure 2 (a), the P-SCO (002) emerges after only 10 seconds in figure 3 (a). At this stage only a small partial of the thin film was transformed to P-SCO as a strong BM peak was still observed. After 40 seconds, the sample was converted to P-SCO completely, and saturated in approximately 20 seconds. The oxidization plateau in figure 2 (b) was not observed in figure 3 (b) probably due to the drastic oxygen up-taken at higher voltage. Similarly, in figure 3 (c) a plateau was observed for resistivity, and then decreased at the phase transition. The resistivity of the thin film can be lowered further with time duration. The saturated sample was then subjected to negative constant voltage to increase the oxygen deficiency. As shown in figure 3 (d), the sample was fully transformed to BM-SCO in 60 s by a small voltage of -0.05 V, during which the P-SCO (002) peak moved almost linearly with the duration. The in-situ I-t mapping in figure 3 (e) exhibits a similar linear down trend over time as figure 2 (e) before the phase transition. Interestingly, the resistivity in figure 3 (f) shows little change at the initial 20 s, although the (002) peak shifts to lower angle. This is probably due to the out-of-plane oxygen liberation. The current is bypassed by the bottom conductive layer. Similar behavior is seen in figure 2 (f). Figure S5 shows the reproducibility of oxygen up-taken/liberation by linearly increasing voltage and constant voltage, according to which the quality of the film shows little 6 ACS Paragon Plus Environment

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change after being treated back and forth. Combining figure 2 & 3, it is found that the oxygen up-taken occurs in in-plane direction due to the initial high resistivity state of the BM-SCO, while it liberates in out-of-plane direction in low resistive P-SCO, which paves the way for interfacial control in asymmetrical heterojunctions37.

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20

-0.05 V

40 60 80

80 -140nA

Figure 3. Control of oxygen deficiency in SCO by constant voltage. The XRD peak evolution of SCO as a function of voltage duration at 1.5 V and -0.05 V are given in (a) and (c), respectively. The in-situ I-t mapping during the redox reaction are shown in (b) and (d); The resistivity of the thin film after each redox reaction are presented in (c) and (f) as well. The oxygen deficiency in the thin film can be tuned between 0 and 0.5 back and forth.

Figure 4 gives the soft X-ray absorption spectroscopy (XAS) at the Co L-edge and O Kedge for SCO2.5 (as-grown BM-SCO), SCO2.75 (SCO right after the BM to P phase transition), SCO3 (fully-oxidized SCO). The Co L3 peak in figure 4 (a) shifts to higher photon energy end as the oxygen content increases, which exhibits an analogous trend to those previous reports5, 24. The branch ratio I(L )/(I(L )+I (L )) decreases from 0.75±0.01 at SCO 3 2 3 2.5 to 0.69±0.01 at SCO2.75, and further 0.68±0.01 at SCO3, indicating a transfer from high-spin Co3+ (t2g4eg2, S=2) to low-spin Co4+ (t2g5eg0, S=1/2).38 Figure 4 (b) gives the corresponding spectroscopy at O Kedge, in which feature A originates from the transition from the O 1s core-level state to the hybridization of unoccupied O 2p with Co 3d t2g, and B from that with Co 3d eg in octahedron coordination.39 The O K-edge of SCO2.5 looks the same as other PLD prepared samples5, 22, but rather different at higher oxygen content, i.e. SCO2.75, since the strong pre-peak from the highspin Co4+ (t2g3eg2, S=5/2) does not show.40 That the intensity ratio I(A)/I(B) keeps decreasing reflects changes in band occupancy, i.e. t2g4eg2 for SCO2.5 to low-spin t2g5eg0 for SCO3, as the



oxygen stoichiometry improves. In particular, feature A2 almost disappears in SCO3, which means nearly all of the high-spin Co3+ have been transferred into low-spin Co4+.

L3

SCO3 SCO2.75

L2

SCO2.5

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Normalized intensity (a. u.)

(b)

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B2 A2 B1 A1 A0 B0

520

805

530

540

550

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570

Photon energy (eV)

Photon energy (eV)

Figure 4. XAS at (a) the Co L-edge and (b) O K-edge for SCO2.5, SCO2.75 and SCO3. The feature A originates from the transition from the O 1s core-level state to the hybridization of unoccupied O 2p with Co 3d t2g, and B from that with Co 3d eg in octahedron coordination.

The magnetization of the SCO thin film depending on the oxygen deficiency is given in figure 5. The evolution of M-T curves with the duration of constant voltage is shown in figure 5 (a). The corresponding magnetic hysteresis loops are given in Fig 5 (b). To make it simple, the oxygen content in the thin film is denoted as O3- .The film remains antiferromagnetic before 20 s at 0.7 V since the M-T curves shows no spontaneous magnetization and the remnant magnetization (Mr) equals zero according to the hysteresis loops. Note that BM-SCO is antiferromagnetic and the Néel temperature amounts to 570 K33. It then changes to ferromagnetic after 20 s of oxygen up-taken. However, the Curie temperature (Tc = 154 K) and moment is relatively low compared to the fully oxidized sample, indicating a high oxygen deficiency in the thin film. As discussed in figure 2 & 3, the partially oxidized thin film is actually a physical mixture of SCO2.5 and SCO2.75. Thus, the week moment stems from the asformed SCO2.75. As more and more SCO2.5 is transformed to SCO2.75, the magnetization becomes larger and larger. Therefore, the Tc does not increase much before the full conversion of the thin film. After 60 s the Tc increases up to 205 K for fully-oxidized SCO3. It looks there is a plateau for Tc as well, corresponding to that of resistivity. A reverse trend is seen for a voltage of -0.05 V. The magnetization decreases step by step with the durations, and to nearly zero after 90 s of oxygen liberation. In figure 5 (b) the saturation magnetization (Ms) and coercive field (Mc) of the thin film changes almost linearly with the durations. However, the ratio of Mr/Ms exhibits a jump at 45 s/0.7 V and 50 s/-0.05 V, where the phase transition occurs, indicating a strong spontaneous magnetization for phase pure P-SCO.


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The magnetization measurement in figure 5 shows that the Tc of fully oxidized SCO3 amounts to approximately 205 K. It is comparable to those PLD prepared samples which are annealed by high pressure O2 (Tc=200 K)22, but still lower than the 250 K of ozone annealed SCO2.88 thin films24 and 305 K of bulk SCO2.9532. From the XRD in figure S6 the quality of the sample shows little change before and after the oxidization. Such a difference probably stems from the different electronic structures of SCO3 (low-spin Co4+, t2g5eg0, S=1/2) compared with those PLD prepared samples (high-spin Co4+, t2g3eg2, S=5/2). As the Curie temperature depends directly on the spin number, the decrease from high-spin state to low-spin state results in a lowered Tc. (a)

T (K)

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100 50 0 -50 -100 -150 -10k 0 10k

H (Oe)

Figure 5. Dependence of magnetization on the oxygen deficiency in SCO thin films. The temperature (M-T) and field (M-H) dependent magnetization of SrCoO3- thin film after being treated by 0.7/-0.05 V with different durations is given in (a) and (b), respectively. The oxygen content in the thin film is denoted as O3- . The thin film can be tuned continuously from antiferromagnetic to ferromagnetic by a constant voltage.

Besides the magnetization, the optical properties of SCO thin films can be tuned by voltage as well. The optical transmittance changes of SCO thin films upon the durations of constant voltage are given in figure 6. In figure 6 (a), the transmittance of as-deposited SCO2.5 decreases gradually from approximately 98% in the vicinity of 2000 nm to 32% near 500 nm. The film exhibits a absorption edge in the visible light range, according to which the band gap can be estimated to be ~2 eV, which is similar to liquid-gating treated samples5. The transmittance in the infrared range decreases more than the visible range as the more part of thin film was oxidized to SCO2.75. At 40 s where the thin film has been converted to SCO2.75 completely, the 9 ACS Paragon Plus Environment


transmittance decreases to below 60%. Further up-taken of oxygen to O2.84 depresses it down to 40%. However, the transmittance below 500 nm improves with the oxygen stoichiometry. Figure 6 (b) shows the evolution of transmittance under -0.05 V, which can be recovered with the duration. Surprisingly, it improves significantly (~20%) in the visible range after 90 s of reduction, indicating that the thin film becomes more transparent compared with the asdeposited SCO2.5.

Transmittance (%)

100

(b)

Visible Infrared

+0.7 V

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O3 (60/0s)

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2000

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Wavelength (nm)

Figure 6. Optical transmittance changes of SrCoO3- thin film during the voltage-controlled (a) oxidization and (b) reduction.

CONCLUSIONS In summary, high quality SCO thin films were grown epitaxially on STO substrate by PLD. The oxygen non-stoichiometry of the sample was switched continuously by either step or constant voltage. At positive voltage, the up-taken of oxygen leads to a topotactic BM-to-P phase transition; while at small negative voltage the liberation of oxygen results in a reverse Pto-BM transition. For high resistivity BM-SCO, the oxygen fills in the in-plane direction; however, it liberates in the out-of-plane direction for low resistivity P-SCO. With such a technique, the film can be switched continuously between insulator and metal, antiferromagnetic and ferromagnetic, or transparent and opaque. The continuous tuning of oxygen deficiency provides possibilities in imaging the full map of phase diagrams for redox metal oxides, such as such as ternary ABO6>A (A = Ln, Ca, Sr, Bi; B = Cr, Mn, Co, Fe, Ni), or binary TiOx, WOx, VOx and NiOx etc., paving an avenue to the searching for novel properties in redox materials. The direction-selective reaction mechanism offers an effective method for 3D interfacial control at the nanometre scale in novel oxide devices, for example, asymmetrical oxide heterojunctions and memristors. The fast voltage-controlled redox reaction makes SCO 10 ACS Paragon Plus Environment

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a good candidate for a number of cutting-edge commercial applications, such as oxygen capacitors, catalysts and smart windows etc..

METHODS The SCO thin films were deposited on (001) SrTiO3 substrates (cubic, a = 3.905 Å) by PLD with a 248 nm wavelength KrF excimer laser. The growth conditions were a temperature of 750 oC, oxygen pressure of 100 mTorr, laser energy density of 1.5 J/cm2, and laser frequency of 1 Hz. The crystal structures and strain state of the thin films were characterized by XRD (Rigaku, SmartLab 9 KW, Cu K , =1.5413 Å). Three silver electrodes were extended from the edges of the thin film, and were then sealed by a Ted Pella crystal bond at 125 oC. Two of them were used for the voltage loading in redox reaction, and the other for conductivity measurement. The measurement geometry is shown in figure S7 in the support information. It was then immersed into a home-made Teflon bath in which 1 mol/L of KOH solution was filled. The sample surface was placed horizontally for X-ray characterization. Pulses of small voltage was applied on the sample and the current flow through the solution were monitored by a Chenhua CHI66E electrochemical workstation to create an in-situ current-time (I-t) mapping. The resistance of the sample after each treatment was measured by a Keithley 6517B electrometer. Prior to the that, the solution level was lowered down to expose the dry surface with a syringe sucking at the bottom of the bath. The structure evolution of the sample was monitored by X-ray at the meanwhile. With such a step up, the structural and electrical changes of the sample can be monitored in-situ during the oxygen up-taken/liberation. The XAS were taken at Beijing Synchrotron Radiation Facility (BSRF) in total fluorescence yield mode (TFY). The magnetization of the thin film was measured by a Quantum Designed MPMS3. The optical transmittance was obtained by a Perklin Elmer Lambda 950 spectrometer.

ASSOCIATED CONTENT Supporting Information. TEM cross-sectional image of thick SCO film on STO substrate; Rocking curves of as-grown BM-SCO thin film along (008) direction; AFM topography of BM-SCO thin film on STO substrate and the height profile; The RSM and AFM of oxygen saturated P-SCO after being oxidized from BM-SCO by applying a linearly increasing voltage; The reproducibility of voltage-controlled oxygen up-taken/liberation by linearly increasing voltage and constant voltage; XRD scans and rocking curves of SCO thin films before and after the oxidization; Measurement geometry of SCO thin film during the oxygen up-taken/liberation

AUTHOR INFORMATION 11 ACS Paragon Plus Environment


Corresponding Authors *E-mail:

[email protected] (S. B. H).

*E-mail:

[email protected] (L. C).

ORCID Songbai Hu: 0000-0002-0183-9559 Lang Chen: 0000-0002-8762-892X Author contributions S. B. H and W. Q. H contributed equally to this work. Notes The authors declare no conflict of interests.

ACKNOWLEDGEMENTS We acknowledge support by the National Natural Science Foundation of China (Grant Nos. 11804145, 11604140, 61601217), Natural Science Foundation of Guangdong Province of China (No. 2018A030310221), the Science and Technology Research Items of Shenzhen (Grant Nos. JCYJ20170412153325679, JCYJ20170817110302672). J. S. acknowledges support by the Australian Research Council through Discovery Grants.

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