Pressure-induced tuning of lattice distortion in a high-entropy oxide

As a new class of multi-principal component oxides with high chemical disorder, high-entropy oxides (HEOs) have attracted much attention. The stability and tunability of their structure and properties are of great interest and importance, but remain unclear. By using in situ synchrotron radiation X-ray diffraction, Raman spectroscopy, ultraviolet–visible absorption spectroscopy, and ex situ high-resolution transmission electron microscopy, here we show the existence of lattice distortion in the crystalline (Ce0.2La0.2Pr0.2Sm0.2Y0.2)O2−δ HEO according to the deviation of bond angles from the ideal values, and discover a pressure-induced continuous tuning of lattice distortion (bond angles) and band gap. As continuous bending of bond angles, pressure eventually induces breakdown of the long-range connectivity of lattice and causes amorphization. The amorphous state can be partially recovered upon decompression, forming glass–nanoceramic composite HEO. These results reveal the unexpected flexibility of the structure and properties of HEOs, which could promote the fundamental understanding and applications of HEOs. High entropy oxides can exhibit remarkable properties which may be amenable to pressure tuning. Here a fluorite-type high entropy oxide is shown to undergo pressure-induced lattice distortion with associated changes to optical behaviour.

T o meet the rapidly evolving technology needs, the materials research community has been striving to seek new methodologies for advanced materials design. Increasing the number of components could effectively expand the configurational space for the exploration of novel materials. Recently, multiple principal components (>5) alloys, so-called high-entropy alloys (HEAs) have been successfully developed in metallic systems with solid solution phases. These HEAs are realized based on an idea with deliberately incorporating five or more elemental species into a single solid solution lattice with random site occupancy, which maximizes the configurational entropy to stabilize the system 1- 5 . Surprisingly, this high-entropy strategy also works for oxides with ionic bonds. Rost et al. first synthesized the transition metal (Mg 0.2 Co 0.2 Ni 0.2 Cu 0.2 Zn 0.2 )O high-entropy oxide (HEO) with a rocksalt-type structure 6 . Subsequently, Djenadic et al. synthesized single CaF 2 -type phase rare-earth HEOs (RE-HEOs) (Ce 0.2 La 0.2 Pr 0.2 Sm 0.2 Y 0.2 )O 2−δ 7 . More recently, complex HEOs with multiple cation Wyckoff positions, like perovskites 8,9 or spinels 10 , have also been reported. HEOs provide great flexibility for element-based property design and engineering, therefore, have attracted extensive attention owing to their tremendous potential in various applications. Many remarkable properties of HEOs have been reported, including huge dielectric constant 11 and high room temperature Li ion conductivity 12 , which have advanced several potential applications like defect dipoles 13 , epitaxial interface engineering 14 , and electrode materials in Li-ion battery 15,16 .
High-entropy materials are also expected to possess ultrahigh structural stability due to their maximized configurational entropy in design (high-entropy effect), which could dramatically lower down the Gibbs free energy of the system. On the other hand, the high chemical disorder could induce severe lattice distortion which could raise the energy of the system. Therefore, it is compelling to experimentally clarify the stability and tunability of these novel materials under various conditions, which is crucial for our fundamental understanding of those new materials and also their potential applications 17,18 . High-temperature is usually involved during the formation of single phase HEOs 6,19 . Besides temperature, pressure is another fundamental parameter controlling the structure and properties of materials [20][21][22][23] . Especially, for HEOs, the Gibbs energy (G = PV − ST), the local lattice distortion, the high-level oxygen vacancies, and even electronic states (valency) are all expected to be susceptible to pressure tuning, which, however, have not been explored before. Very recently, the seemingly ultra-stable HEAs have already been found to show rich polymorphism under high-pressure at room temperature. For example, pressure-induced irreversible polymorphic transitions from the fcc to hcp phases were observed in the prototype CoCrFeMnNi HEA 24,25 , and a medium-entropy NiCoCrFe alloy 26 ; In another rare-earth system, an hcp-structured HoDyYGdTb HEA, more complex transitions following the sequence of hcp → Sm-type → dhcp → dfcc were reported 27 ; Even five polymorphs were discovered in the bcc Al 0.6 CoCrFeNi HEA 18,28 . These results demonstrate that the HEAs are far more complex than expected as ultrastable solid solutions. Inspired by these recent results of HEAs, an intriguing question is raised: are the structure and properties tunable in HEOs?
In this work, we use in situ high-pressure synchrotron radiation X-ray diffraction (XRD), pair distribution function (PDF), Raman spectroscopy, and ex situ high-resolution transmission electron microscopy (HRTEM) to investigate the structural evolution of the rare earth (Ce 0.2 La 0.2 Pr 0.2 Sm 0.2 Y 0.2 )O 2−δ RE-HEO ((Ce,La,Pr,Sm,Y)O will be used hereafter for simplicity). Lattice distortion in the initial structure is confirmed by the slight deviation of bond angles from the ideal values. The bond angles are found to be extremely susceptible to pressure, which continuously bend during compression and eventually causes a pressure-induced amorphization (PIA) in HEOs. The optical properties and band gap are also investigated through the transition by in situ high-pressure ultraviolet-vis (UV-vis) absorption spectroscopy, which shows interesting tunable behavior as well accompanying the structural change.

Results
Structure of the initial sample. (Ce,La,Pr,Sm,Y)O HEO powder samples were synthesized using nebulized spray pyrolysis (NSP) method 7 . The initial structure of the sample was characterized by chromatic-aberration corrected HRTEM and synchrotron radiation high-energy XRD. The experimental details are shown in the "Methods" section. HRTEM results ( Supplementary Fig. 1) show that the sample has a particle size of~0.5 µm, a crystalline grain size of~7 nm, and a CaF 2 -type (space group: Fm-3m) structure with its lattice parameter a = 5.6 (1) Å. The structure factor S(q) obtained from the high-energy XRD with the moment transfer q coverage up to 30 Å −1 is shown in Fig. 1a, the structure can be well indexed into a CaF 2 -type structure with a refined lattice parameter a = 5.507(1) Å. These results of the initial structure of the (Ce,La,Pr,Sm,Y)O HEO are consistent with the previous reports 7 . Fig. 1b shows the PDF data, G(r), which was obtained by a Fourier transformation of the S(q) data. Bond lengths can be accurately determined by fitting the peaks in G(r), then, bond angles can be calculated. The initial average bond angles between oxygen and rare-earth elements, θ RE-O-RE (~108.6°) and θ O-RE-O (~68.3°), both are slightly smaller than the ideal values of a  Supplementary  Fig. 1c), which may be attributed to the local lattice distortion in the (Ce,La,Pr,Sm,Y)O HEO. The lattice distortion could be caused by atomic size mismatch between neighboring RE elements, owing to the fact that cationic radii of the constituent rareearth elements vary considerably, for example, from 0.97 Å (for Ce 4+ ) to 1.116 Å (for La 3+ ). Additionally, the high level of oxygen vacancies could also play an unneglectable role resulting in the observed lattice distortion 29 .
In situ high-pressure XRD. To address the structural stability of the (Ce,La,Pr,Sm,Y)O HEO, its structural evolution up to~31 GPa was monitored by in situ high-pressure synchrotron radiation XRD in a symmetric diamond anvil cell (DAC). A schematic illustration of a DAC is shown in Fig. 2a. Figure 2b presents the XRD patterns during compression and decompression. During compression, the initial CaF 2 -type structure loses its crystal coherence above~16 GPa indicated by the disappearance of the initial sharp Bragg peaks and the emerging of amorphous-like broad peaks. The amorphous-like phase exists up to the highest pressure of~31 GPa and seems recoverable to ambient pressure during decompression without the reappearance of the sharp Bragg peaks. The two-dimensional (2D) XRD images are shown in Supplementary Fig. 2. Below 16 GPa, with increasing pressure, all the Bragg peaks of the initial CaF 2 -type structure shift to the larger two-theta values below~9 GPa simply due to the pressure-induced lattice shrinkage (Fig. 2c). At~4 GPa, there is a slight slope change, and the peak position stops shifting above~9 GPa. Below~9 GPa, the intensity of all XRD peaks gradually but significantly weakens (Fig. 2d), and all the peak widths continuously increase as well (Fig. 2e). Both peak intensity and peak width show a slope change at~4 GPa and become almost constant above 9 GPa. The peak intensity weakening and peak width increasing start at the very beginning of compression. This result indicates that the initial CaF 2 -type structure of the (Ce,La,Pr,Sm,Y)O HEO is susceptible to pressure and continuously becomes more and more disordered (topologically) during compression. Below~16 GPa, the structure evolution has three stages (one till~4 GPa, another between~4 and~9 GPa, and lastly above~9 GPa) according to the peak position, intensity, and peak width changes as a function of pressure.
In situ high-pressure PDF. Since structural disordering occurs in the (Ce,La,Pr,Sm,Y)O HEO during compression, to derive more detailed information of the structural evolution, another experiment with larger Q coverage (up to 12 Å) was performed (more details are shown in the "Methods" section). The in situ highpressure structure factor S(q) was obtained after careful background subtracting and normalization (Fig. 3a). The peak position, intensity, and width of the S(q) exhibit consistent behavior with the raw XRD data shown in Fig Fig. 2 In situ high-pressure XRD. a A schematic illustration of the DAC used to generate high-pressure. b Representative XRD patterns of (Ce,La,Pr,Sm,Y)O during compression and decompression at room temperature. Inset is an image of the sample loaded inside a DAC, the scale bar is 100 μm. The peaks of the amorphous phase are marked by asterisk symbols. The X-ray wavelength is 0.3218 Å. c Pressure dependence of d/d 0 of the CaF 2 -type cubic phase upon compression. Pressure dependence of the peak height d and peak width e of the (111) peaks (at~5.8°) upon compression obtained. Meanwhile, the peak intensity and width do not change with pressure anymore, suggesting the completion of the amorphization process.
Structural information in real space is presented by PDF (Fig. 4a). Below~20 GPa, the PDF patterns show strong oscillations without noticeable damping up to 25 Å, which is a typical feature of crystalline materials with long-range translational symmetry. In contrast, above~20 GPa, all the oscillations are significantly weakened and gradually smeared out above~10 Å, which is in line with the amorphization process indicated by S (q) data in reciprocal space. All the peaks of PDF which correspond to the distances between different atom pairs shift to smaller r values with increasing pressure below~16 GPa. For convenience, the first nine peaks (illustrated in Fig. 4b) of the initial CaF 2 structure are denoted as P 1 -P 9 in Fig. 4a. Some peaks gradually weaken and merge into one new broad peak during compression, e.g., P 1 and P 2 , P 4 and P 5 . Above~20 GPa, only four broad and weak peaks remain due to amorphization. The changes of the nine peak positions, the calculated atomic pair distances, and bond angles are shown in Fig. 5. The atomic pair distances shrink quite consistently below~4 GPa, and then diverge at higher pressures, which could be the reason for the slight slope change at~4 GPa observed in Fig. 2c-e. It should be noted that although the q coverage in the in situ high-pressure PDF measurement is limited (q max ≈ 12 Å −1 ) due to the restricted opening of DACs, the S(q) and G(r) are reliable according to their consistent feature with the high-energy PDF of the ambient samples (q max ≈ 30 Å −1 , see Fig. 1).  (Fig. 5a, b). With accurately determined bond lengths by fitting the peaks in G(r), the pressure dependence of bond angles inside one cube θ O-RE-O and between the nearest cubes θ RE-O-RE can be calculated as shown in Fig. 5c. During compression, both of the average θ O-RE-O and θ RE-O-RE decrease continuously. The sample volume compression is achieved mainly by bending the bond angles. Eventually, the long-range coherence between REO 8 cubes breaks down when the pressure is above the mechanical instability limit, hence, leading to amorphization.
Upon decompression, both S(q) and G(r) seem to retain their amorphous feature. However, regardless of the peak shifting due to elastic recovery, the S(q) pattern is slightly different from the high-pressure pattern at 34.6 GPa especially in the range from 2.5 to 4.5 Å −1 . This result suggests that the low-pressure recovered state is different from the high-pressure amorphous state. G(r) also looks different after decompression. To clarify the structure change after the pressure release, HRTEM was performed on the recovered samples ( Supplementary Fig. 3). The recovered sample mostly remains in a highly disordered amorphous state, but some fine nano-grains of the initial CaF 2 crystalline structure show up in the amorphous matrix. Therefore, a composite HEO with amorphous matrix plus some super fine nanocrystals has been synthesized via compression and the following decompression processes of the initial HEO.
The structural stability of the prototype (Ce,La,Pr,Sm,Y)O HEO has been studied from room temperature to~1023 K at ambient pressure, but no amorphization or any other phase transition has ever been observed 7 . A symmetry lowering phase transition from Fm3m to Ia-3 space groups was observed only after calcination at 1273 K for 1 h 7 . Its phase stability under highpressure was not explored before. In contrast, the phase transitions under high-pressure have been extensively studied in all of its constituent rare-earth oxides in nano or bulk forms, for example, Ce 2 O 3 [37][38][39] . Among them, however, the PIA process was only observed in the Y 2 O 3 nanoparticles with sizes <16 nm 39 . The Y 2 O 3 nanoparticles have a cubic structure (space group: Ia-3) at ambient condition, which consists of YO 6 octahedra as the building block. The YO 6 octahedra are less compressible than the global Y 2 O 3 lattice. Therefore, during compression, the densification is mainly carried by bending the polyhedral-linking angles to squeeze the space between the octahedra. When the bending is beyond a critical mechanical limit, the structure collapses and transforms into an amorphous state at~25 GPa. Smaller grain size usually results in higher transition pressures for the PIA 39 . This mechanism was also proposed for the PIA in many other polyhedral-linked materials such as the prototype oxide, SiO 2 40 . PIA refers to an unusual phenomenon with direct transformation of a crystal into an amorphous solid, which was firstly reported in ice I h , afterwards, has attracted intense interest and observed in a wide range of substances 41 . Our results in this work demonstrate that a similar mechanism proposed for PIA in conventional oxides may also hold for the case of the chemically much more complex HEO according to the bonding length and angle changes during compression.
The (Ce,La,Pr,Sm,Y)O HEO consists of five cations with different ionic radii, and more importantly different oxidation states (valence), e.g., 3+ oxidation for La, Sm, Y; 4+ for Ce, and mixed multiple oxidation states (3+ and 4+) for Pr 29,42 . CeO 2 is believed to be the critical component which stabilizes the singlephase of the HEO and provides a CaF 2 -type host lattice for the other cations. However, except Ce 4+ and Pr 4+ , the other rareearth cations including partial Pr are present in the 3+ oxidation states. To keep the charge balance of the system, high-level oxygen vacancies are naturally required and created in the (Ce,La, Pr,Sm,Y)O HEO 29 . Therefore, the oxygen vacancies and cations valence are closely associated in the HEO. The different ionic radii and a high level of oxygen vacancies could both contribute to the lattice distortion in the initial (Ce,La,Pr,Sm,Y)O HEO sample which may be the reason for the slight deviation of the bond angels from ideal values. However, according to the G(r) data, the first atomic shell consists of RE-O bonding is rigid without obvious bond shrinkage during compression. Therefore, the size of different cations may not change much as well during the entire compression. On the other hand, the atoms (both the RE cations and other oxygen inions) surrounding a given oxygen vacancy site are expected to displace from their original positions in a perfect lattice, which could also result in bond angle deviations. Therefore, the behavior of oxygen vacancies during compression could dramatically affect the structure evolution under high-pressure.
In situ high-pressure Raman spectroscopy. Oxygen vacancies of the (Ce,La,Pr,Sm,Y)O HEO can be readily detected by Raman spectroscopy 29 . Thus, we performed in situ high-pressure Raman spectroscopy (Fig. 6) on the (Ce,La,Pr,Sm,Y)O HEO in a DAC (details can be found in the "Methods" section), which could provide another indicator of the development of the lattice distortion of (Ce,La,Pr,Sm,Y)O HEO during compression. The F 2g symmetric vibration mode of the eight-fold RE-O bond in the (Ce,La,Pr,Sm,Y)O which is typically located at~450 cm −1 is barely visible. The Raman peak weakening for F 2g can be explained by the symmetry breaking (the severe local lattice distortion) in the (Ce,La,Pr,Sm,Y)O caused by chemical disorder 29 . The Raman peak at~600 cm −1 (V O band) at ambient conditions was confirmed to be the characteristic peak of oxygen vacancies of the (Ce,La,Pr,Sm,Y)O HEO which is relatively strong at ambient conditions. During compression, the V O band shifts to higher wavenumbers, and its intensity decreases significantly until ~20 GPa. This observation indicates that the oxygen vacancies may change their distribution and concentration during compression below 20 GPa, which could be associated with the relatively large compressibility of the vacancy spot and also the tunable valences of the Pr 3+ to Pr 4+ during compression 43,44 . Above~20 GPa, the amorphization is completed, therefore, it is reasonable that the intensity of V O band remained almost constant. Naturally, the bond angles and local lattice distortion will be altered by the collapse of oxygen vacancies and oxidation stateinduced change in the cationic radii of Pr cation, which is consistent with the results obtained by in situ high-pressure XRD. During decompression, the intensity and peak position of the V o band gradually recovered to the initial values. This result suggests that the collapse of oxygen vacancies may be caused by elastic deformation and the change of cations valence during compression, which is reversible upon decompression and likely independent of the lattice disorder 45 . To clarify the role of each constitutent element of this HEO in its structural evolution under high-pressure, element-specific technique, such as x-ray absorption spectroscopy (XAS) is needed. However, in situ highpressure XAS expecially extended x-ray absorption fine structure (EXAFS) is challenging to the (Ce,La,Pr,Sm,Y)O HEO in a DAC due to the either too high X-ray energy (K-edge) for the thin samples (<20 μm) or too low X-ray energy (L-III edge) for the thick diamond anvils (~2.4 mm), and too close absorption edges of some constitutent elements. During decompression, most of the amorphous state could be kinetically retained by the possible large barrier required for the recovery of the initial cubic crystalline structure. During heating at ambient pressure, the (Ce,La,Pr,Sm,Y)O HEO will have a phase transition from the fluorite-type to the bixbyite-type phase transition with lower symmetry 7 . However, we did not observe this polymorphic crystal-to-crystal transition under high-pressure and room temperature. Instead, an amorphization takes place. We speculate that the crystal-to-crystal transitions may be hindered by a large energy barrier in this system, therefore, we also performed in situ laser heating XRD measurements at highpressure (more experimental details can be found in ref. 46 ) on the (Ce,La,Pr,Sm,Y)O HEO (Supplementary Fig. 4). Thermodynamically more stable crystal phases (do not match the bixbyite-type phase) do exist at pressures before the amorphization and after the amorphization if the sample is heated abovẽ 1100 K. A possible phase separation also cannot be neglected under high-pressure and high-temperature. More efforts are needed to solve the structure of the high-pressure and hightemperature crystalline phase in the future.
Pressure tuning of band gap. Since the structure of (Ce,La,Pr, Sm,Y)O HEO is extremely susceptible to pressure, we expected tunable properties during compression as well. In this work, we studied its band gap as a function of pressure using in situ highpressure UV-vis spectroscopy. The indirect and direct gap values were obtained by fitting the absorption spectra using the Tauc relation ( Fig. 7) 29 is the Kubelka-Mubelka-Munk function, h is the Planck's constant, v is the frequency, A is a constant, and E g is the band gap energy. Details of the fitting of E g can be found in Supplementary  Fig. 5. Before amorphization (<~16 GPa), the sample exhibits enhanced band gap with increasing pressure, this is unusual compared to the typical band gap decreasing in traditional oxides during compression 47,48 . It may be caused by the accumulated lattice distortion energy. Above 16 GPa, once the sample starts amorphization, the band gap turns to decrease sharply until the amorphization completes at~22 GPa. Above 22 GPa, the band gap increases with pressure again. The complex pressure dependence of the band gap observed in the (Ce,La,Pr,Sm,Y)O HEO is consistent with its structure evolution during compression revealed by XRD. The enhanced band gap below 16 GPa and above 22 GPa during compression seems to be an intrinsic feature in both the cubic crystalline and amorphous states during compression. The drop of band gap between 16 and 22 GPa is likely associated with the amorphization process, which may be attributed to the distortion (strain) release due to the amorphization (bond breaking and atom rearrangement). Although there is no reliable data collected during decompression in this work due to our experimental issue, we expected a partially reversible band gap change during decompression according to its atomic structure evolution. help to understand its underlying mechanism and even achieve band gap engineering via high-pressure treatment 49,50 .

Discussion
The structural stability of the (Ce,La,Pr,Sm,Y)O HEO is studied during compression up to~35 GPa using in situ high-pressure synchrotron radiation XRD. A crystal to amorphous phase transition is observed which starts at~16 GPa and completes at 22 GPa. PDF analysis reveals the microscopic mechanism of the amorphization, which is attributed to the breakdown of longrange connectivity between REO 8 cubes caused by continuous bending/tilting of bond angles inside and between the REO 8 cubes. The atomic distortion caused by chemical disorder and high-level oxygen vacancies is suggested to exist even in the initial crystalline HEO (the bond angles deviate from ideal values) and is extremely sensitive to pressure. The development of lattice distortion plays a vital role in the amorphization and band gap tuning. These results suggest that although the HEOs are very stable at high temperatures and ambient pressure, they are susceptible to pressure loading. Their properties could be highly tunable under high-pressure as well. Pressure-induced transformation directly from crystalline states to the amorphous solid states at temperatures below the melting point or glass transition temperature is well-known as the "PIA" phenomenon since the 1980s in various materials 51 . However, it has never been observed in the multicomponent HEO systems before. For HEO, the PIA also means the direct transition from ceramics to glasses, which could lead to the development of useful nano-ceramic-glass composite materials with tunable properties. Since the HEOs may commonly involve severe lattice distortion due to cation size mismatch and/or a high level of oxygen vacancies, we expect interesting pressure-induced transitions including PIA in other HEOs.

Methods
In situ high-pressure synchrotron XRD measurements. In situ high-pressure angle-dispersive XRD experiments were mainly performed at the GSECARS beamline 13-ID-D, Advanced Photon Source (APS), Argonne National Laboratory (ANL). A monochromatic X-ray beam with a wavelength of 0.3218 Å was focused by a Kirkpatrick-Baez (KB) mirror system down to~2.5 × 3.5 μm 2 at the rotation center of the sample stage. The detector was a Mar165 charge-coupled device (CCD). The position and orientation of the detector to the X-ray beam were calibrated using the diffraction pattern of a LaB 6 standard. High-pressure was generated using a symmetric DAC with a culet size of~400 μm in diameter. The gasket was T301 stainless steel that was pre-indented to~20 GPa. A hole with a diameter of~150 μm was drilled to serve as the sample chamber by a laser drilling system at the center of the gasket indent. The powder was pre-pressed into a thin pellet and cut into flakes with proper sizes for DACs. Then, the tiny flake was loaded into a DAC along with a tiny ruby ball and Au foil beside the samples to calibrate pressure. Methonal:ethonal = 4:1(volume ratio) mixture was used as the pressure-transmitting medium. The background scattering was collected at each pressure point by shining the X-ray beam on the empty area inside the sample chamber, which only passed through the pressure medium and two diamond anvils. The exposure time was set at 10 s for normal XRD. For PDF measurements, the exposure time for each pattern was set to be 30 s. One-dimensional (1D) XRD patterns were obtained by integrating the 2D patterns along the azimuth angle from 0°to 360°with the Dioptas software 52 . The resulting diffraction patterns were refined using the GSAS package.
In situ high-pressure Raman spectroscopy measurements. Raman spectroscopy measurements were conducted using a Renishaw inVia spectrometer with a 532 nm laser excitation. High-pressure was generated using a symmetric DAC with a culet size of~400 μm in diameter. The powder sample was pre-pressed into a pellet and then loaded into a DAC, along with two tiny ruby balls beside the samples to calibrate pressure. Silicone oil was used as the pressure-transmitting medium. The exposure time for each pattern was~50 s. The signal from background was obtained by shining a laser beam on the spot inside the sample chamber but away from the sample.
In situ high-pressure UV-vis spectroscopy measurements. UV-vis spectroscopy measurements were conducted using an Ocean Optics DH-2000-BAL spectrometer. High-pressure was generated using a symmetric DAC with a culet of 300 μm in diameter. The powder sample was pre-pressed into thin flakes with right thickness and then loaded into a DAC, along with two tiny ruby balls beside the sample to calibrate pressure. Silicone oil was used as the pressure-transmitting medium. Each spectrum was collected for 1 s with wavelength ranged from 180 to 860 nm. The background signal was collected at the empty area (with pressure medium only) of the sample chamber.
HRTEM. The powder samples were dispersed in ethyl alcohol and picked up using a holey carbon grid. HRTEM experiments were carried out using the chromaticaberration-corrected TEM (ACAT) at ANL with an image corrector to correct both spherical and chromatic aberrations. An accelerating voltage of 80 kV was used to minimize the possible electron irradiation effect on the sample structure.

Data availability
The data that support the findings of this study are available from the corresponding author upon reasonable request.