Journal of Materials Science & Technology, 2020, 47(0): 122-130 DOI: 10.1016/j.jmst.2019.12.024

Research Article

Carbide precipitation behavior and mechanical properties of micro-alloyed medium Mn steel

Luhan Haoa, Xiang Jia,b, Guangqian Zhanga, Wei Zhaoa, Mingyue Sun,b,*, Yan Peng,a,*

aNational Engineering Research Center for Equipment and Technology of Cold Strip Rolling, Yanshan University, Qinhuangdao 066004, China

bShenyang National Laboratory for Material Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China

Corresponding authors: * E-mail addresses:mysun@imr.ac.cn(M. Sun),pengyan@ysu.edu.cn(Y. Peng).

Received: 2019-10-28   Accepted: 2019-12-26   Online: 2020-06-15

Abstract

The carbide precipitation behavior and mechanical properties of advanced high strength steel deformed at different temperatures are investigated by X-ray diffractometer (XRD), scanning electron microscope (SEM), transmission electron microscope (TEM) equipped with an energy dispersing spectroscopy (EDS), and tensile tests. The medium Mn steel was subjected to controlled deformation up to 70% at 750 °C, 850 °C, 950 °C, and 1050 °C, and then quenched with water to room temperature, followed by intercritical annealing at 630 °C for 10 min. In comparison with the undeformed and quenched specimen, it can be concluded that acicular cementite precipitates during the quenching and cooling process, while granular NbC is the deformation induced precipitate and grows during the following annealing process. As the deformation temperature increases from 750 °C to 1050 °C, the product of strength and elongation increases at first and then decreases. The smallest average size of second phase particles (20 nm) and the best mechanical properties (32.5 GPa.%) can be obtained at the deformation temperature of 950 °C.

Keywords: Advanced high strength steel ; Medium Mn steel ; Thermal deformation ; Intercritical annealing ; The product of strength and elongation

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Cite this article

Luhan Hao, Xiang Ji, Guangqian Zhang, Wei Zhao, Mingyue Sun, Yan Peng. Carbide precipitation behavior and mechanical properties of micro-alloyed medium Mn steel. Journal of Materials Science & Technology[J], 2020, 47(0): 122-130 DOI:10.1016/j.jmst.2019.12.024

1. Introduction

Medium manganese (Mn) steels have been extensively investigated as the candidate alloys for the third generation of advanced high strength steels (AHSS) [[1], [2], [3]]. The microstructure of medium Mn steels is typically composed of a cubic α’ martensite (matrix) and some retained/reversed austenite obtained by intercritical annealing (reversion heat-treatment) [[4], [5], [6]]. A wide range of mechanical properties can be obtained for medium Mn steels by optimizing alloy composition and adjusting processing technology and heat treatment process. In order to meet the requirements of ultra-high strength and good toughness, the Nb, V and Mo micro-alloyed elements are introduced to medium Mn steels which can be combined with C or N to form precipitates of different type, size and distribution. The precipitates play a significant role in strengthening steels by grain boundary strengthening and precipitation hardening [[7], [8], [9]]. Zhou et al. [10] made a review concerning the microstructural control involving matrix and second phase particles. It was found that both the precipitation behavior and coarsening rate of MX is relevant to the composition adjustment of martensitic/ferritic steels. Moreover, by adopting proper heat treatment, the initial size and amounts of MX can be further tailored. Chen et al. [11] found that the molybdenum addition of 0.38 wt% could greatly increase strengthening. The Vickers hardness can be mainly increased through enhanced hardenability and associated modification of the resultant transformation microstructure. Wu et al. [12] used nano-indentation test to evaluate the hardness of various VC in Ni-Cr cast iron, which showed that cubic VC possesses a relatively high Vickers hardness. And the data of Vickers hardness show an increase in hardness, which is associated with the formation of a large amount of nanometer-sized carbides [13].

A large number of studies have shown that size, distribution and precipitation of precipitated phase in high-strength steel have an important influence on the material properties [14,15]. For example, Pouraliakbar et al. [16], Khorrami et al. [17] and Ranibarnodeh et al. [18] have studied the precipitation of carbides in different kinds of steels and found that various types of carbides such as M2C, M23C6 and MC can be precipitated. Chen et al. [19] discussed the effects of adding Nb, Ti and Mo to steel on the precipitation intensification effect of carbide. Misra et al. [20] found that high dislocation density and small precipitated particles are the main reasons for high strength (yield strength) of 770 MPa by studying the Nb-Ti composites reinforced hot rolled steel. For Nb-V-Mo micro-alloyed steel, the increase in the deformation temperature would restrict the precipitation of Nb in austenite and enhance its precipitation in ferrite, which would make the Nb-rich precipitates much finer and the precipitation strengthening much more effective. In the meantime, the raise of deformation temperature would decrease the effect of grain refinement strengthening [21,22]. An optimized design of thermo-mechanical process is required in order to take advantages of both precipitation hardening and grain refinement strengthening. However, the carbide precipitation mechanism and interactions with the thermal deformation parameters in micro-alloyed medium Mn steel have not been reported yet.

In the present study, the effects of deformation temperature on the microstructure, the behavior of carbides precipitation and mechanical properties of the micro-alloyed medium Mn steel were investigated with the aim of optimizing the thermo-mechanical process and improving the mechanical properties of the AHSS.

2. Experimental details

A 25 kg ingot of micro-alloyed low-carbon medium Mn steel was cast in vacuum induction furnace. The composition of the ingot was measured by wet chemical analysis, which is shown in Table 1. The material was homogenized for 4 h at 1200 °C in Ar-rich atmosphere, and subsequently hot forged to billet of section size 60 mm × 60 mm at different temperatures ranging from 1200 to 850 °C, followed by air cooling to ambient temperature. Then, the forged steel was normalized at 1100 °C for 80 min. The as-quenched microstructure of the hot forged specimen was lath martensitic without retained austenite, as shown in Fig. 1.

Table 1   Chemical compositions of the experimental steel (mass %).

CMnSiMoNbVPSFe
0.154.81.20.110.050.050.0040.002Bal.

New window| CSV


Fig. 1.

Fig. 1.   Microstructure of the studied steel at hot forging state under SEM.


The phase transition temperature Ac1 of the material was tested to be 620 °C and Ac3 was 818 °C, and non-recrystallization temperature of the austenite was below 980 °C which is obtained by the experiment of austenitic dynamic recrystallization. The billet was machined into a number of cylindrical specimens with the diameter of 8 mm and the height of 12 mm for thermal-mechanical simulation experiments on Gleeble 3800 machine (USA). The schematic diagram of experiment setup, position for microstructure observation and dimensions of the micro-tensile specimens are schematically shown in Fig. 2(a), and thermal mechanical schedule is shown in Fig. 2(b). After austenitizing at 1200 °C for 5 min, the specimens were cooled at the cooling rate of 5 °C /s to the deformation temperature of 750 °C, 850 °C, 950 °C, and 1050 °C and held for 1 min. Then uniaxial compression was applied to the specimens at the true strain rate (1/s) with the height reduction of 70% (true strain 1.2). Finally, the deformed specimens were quenched by water immediately after compression. All the quenched deformation specimens were annealed at 630 °C for 10 min, and then air cooled to room temperature. The specimens were then sectioned through the center parallel to the axial direction for microstructure observations. After the center sections were ground, polished, and etched with alcohol solution containing 4% of nitric acid (volume fraction) for 15 s, the microstructures were observed by optical microscopy (OM, ZEISS AXIOVERT 200MAT, Germany) and scanning electron microscope (SEM, JMS-6301-F, Japan). Furthermore, the morphology and distribution of nanometer-sized precipitates in the specimens were analyzed by transmission electron microscopy (TEM, FEI Tecnai G2 F20, USA). Samples for TEM examination were sliced to 500 μm thicknesses and mechanically polished to 50 μm, and then electrically thinned by a twin-jet polisher in a 7 vol% perchloric acid and 93 vol% ethanol solutions. In addition, the chemical compositions of the particles were analyzed by the Energy Dispersive X-Ray spectrum (EDS) of the TEM. The identification of precipitated phase can be determined by X-Ray diffraction (XRD, Rigaku D/Max-2400PC, Japan) patterns of the powdered extract specimens. The test samples were extracted by the electrochemical extraction with 10% hydrochloric acid methanol solution and the extract current remaining at about 0.25 mA/mm2 and voltage of 2.8 V. During the XRD measurements, Cu Kα radiation was used and the specimens were step-scanned from 10° to 90° at a scanning speed of 1°/min. Tensile test at room temperature was carried out on electronic universal testing machine (MTS E45.105, China) with a stretching speed of 0.5 mm/min. The specimen size was shown in Fig. 2(a) and it was recommended that the thickness should be about 1.0 mm with the size tolerance of ± 0.03 mm. Three samples for the same process were tested and the average values were calculated.

Fig. 2.

Fig. 2.   Schematic diagram of experiment setup and positions of the studied specimen (a) and thermal deformation process (b). DT: deformation temperature; Min: minutes; S: seconds; WQ: water quenching; AC: air cooling.


3. Results and discussion

3.1. Microstructure and substructure characterization

The room-temperature microstructure of the quenched specimens deformed under austenite recrystallized temperature (1050 °C), unrecrystallized temperature (950 °C and 850 °C) and intercritical region temperature (750 °C) was analyzed. As shown in Fig. 3, the microstructure of the quenched specimens after deformation is typical martensite laths with little amount of second phase. It can be seen that the length of the martensite lath deformed at 750 °C is shorter than of one deformed at 850 °C, and the width of the lath is almost the same. This indicates that the microstructure gets more refined with the deformation temperature getting lower. Since at the low deformation temperature, the prior austenite grains have a higher dislocation density and larger deformation energy. The martensite nucleation is favored and the growth of the martensite structure can be hindered. As the compression temperature increased to 950 °C, the martensite lath size got coarsened. The martensite lath size is larger when deformed at 1050 °C because the austenite structure of the large deformation undergoes dynamic recrystallization to consumeformation energy and it is easy for grains to grow up at high temperature. However, the compression of the non-recrystallized zone is beneficial for maintaining the dislocation density in the austenite, thereby increasing the nucleation point of the grain and ree dfining the grain.

Fig. 3.

Fig. 3.   SEM microstructure of the experimental steel at quenched state after deformation at (a) 750 °C; (b) 850 °C; (c) 950 °C; (d) 1050 °C.


Fig. 4 shows the SEM microstructure of the experimental steel after 70% compression at 750 °C, 850 °C, 950 °C and 1050 °C, followed by intercritical annealing. Along with the tempered martensite, a large amount of fine precipitate was observed in the annealed specimens. Fig. 5 shows the microstructure and substructure of the annealed specimens under TEM. It can be seen that the microstructure of the sample after thermo-mechanical treatment is martensite laths with different orientations and is arranged substantially parallel in the martensite block, while the morphology of martensite is different at different deformation temperatures. When the low temperature deformation occurs, the martensite lath is deformed and twisted, and the significant distortion of substructure occurs, thereby dislocation density gets increased. As the deformation temperature increases, the size of the martensite lath increases slightly. A large number of dislocations accumulate to form a stable dislocation cell structure and fine precipitates are scattered between the laths. When the deformation temperature is increased to 1050 °C, the martensite lath is coarser and the number and density of dislocations in the microstructure are reduced. It is probably because that as the deformation temperature increases, the deformation resistance of the experimental steel gets decreased, and the movements and proliferation of dislocations are also reduced. In addition, a small amount of retained austenite (yellow arrow area in Fig. 5) was found in the sample, which has a great effect on the toughness and plasticity of the material.

Fig. 4.

Fig. 4.   SEM microstructure of the intercritical annealed specimens after deformation at (a) 750 °C; (b) 850 °C; (c) 950 °C; (d) 1050 °C.


Fig. 5.

Fig. 5.   TEM microstructure and substructure of the annealed specimens after deformation at (a) 750 °C; (b) 850 °C; (c) 950 °C; (d) 1050 °C.


3.2. Analysis of the second phase precipitation behavior

Firstly, the precipitation characteristics of the undeformed and quenched specimen are studied by TEM - EDS (Fig. 6) and the corresponding Selected Area Diffraction (SAD) pattern is shown in the inset of Fig. 6(a). It can be concluded that the acicular carbides with orthorhombic structure are M3C and rich in Fe and Mn.

Fig. 6.

Fig. 6.   TEM images with SAD pattern of M3C (a) and EDS analysis of the precipitates (b) of the quenched specimen without deformation.


Fig. 7 demonstrates the precipitates distribution of the quenched specimens, which have been deformed beforehand separately at 750 °C, 850 °C, 950 °C and 1050 °C. It can be found that there are two classical morphologies of the precipitates - granular and acicular. The precipitates are mainly precipitated inside the crystal grains with a few along the grain boundary. With the deformation temperature increasing from 750 °C to 950 °C, the granular precipitate gets increased and becomes finer circular particles, with very small amount of acicular carbides around. However, the precipitated particles get rapidly increased and large needle-like precipitates are occured when the deformation temperature is increased to 1050 °C.

Fig. 7.

Fig. 7.   TEM images of precipitates in quenched specimens after 70% deformation at (a) 750 °C; (b) 850 °C; (c) 950 °C; (d) 1050 °C.


Compared with the undeformed quenched specimen, it can be concluded that the needle-like M3C precipitates are precipitated during the quenching process and the granular precipitates are occured during deformation process. According to the following XRD results, the granular second phase particles are identified to be micro-alloyed carbonitrides, which have a small size ranging from a few to a few dozen nanometers. At higher deformation temperature the dislocations are annihilated by dislocation climbing. Therefore, the micro-alloy precipitates are mainly precipitated in the martensite matrix instead of the mobile dislocations.

Fig. 8 shows the TEM analysis of the annealed specimens after thermal deformation at different temperatures with the subsequent EDS analysis of the precipitates and martensite matrix. With the increase of the deformation temperature, the number of precipitates gets increased and becomes finer, and the smallest and most dispersed distribution of the precipitates occurred at 950 °C relative to 750 °C and 850 °C, as shown in Fig. 8(a)-(c). The precipitates grow up rapidly at the deformation temperature of 1050 °C. According to transmission electron microscopy, it can be observed that the distribution of the precipitates is not homogeneous for the specimen deformed at 750 °C and 1050 °C. Some precipitates gather together forming a cluster and the orientation of the precipitation is almost the same inside one cluster. And there is always a certain distance between the clusters. When the distance between the clusters gets smaller, the precipitates will coalesce into bigger size as shown in Fig. 8(a) and (d).

Fig. 8.

Fig. 8.   TEM images of precipitates of the annealed specimens after 70% deformation at (a) 750 °C; (b) 850 °C; (c) 950 °C; (d) 1050 °C and EDS analysis of precipitates and matrix of area 1 (e), area 2 (f) and area 3 (g).


There is a lath-like structure which contains high-density dislocation lines and dislocation cells in the experimental steel. The precipitated phase of the annealed sample has the irregular, nearly spherical or rod-like morphology. The large-sized precipitated phase is precipitated from the deformed austenite during the hot deformation process and grows and coarsens during the subsequent annealing process; while the small precipitated phase is precipitated during the subsequent annealing process. Whether it is deformation-induced precipitation phase or phase transformation-induced precipitation phase, the precipitation phase needs to minimize the specific surface area and interfacial energy per unit area in order to reduce the interfacial energy at the beginning of precipitation. Therefore, fine precipitation is generally spherical or granular, but the small size of the precipitated particles means that they have a large specific surface area and a high interfacial energy. From the viewpoint of thermodynamics, these precipitated particles are mutually annexed and coarsened to reduce the interface energy. Nano-carbides exist in the original azimuth boundary and the matrix. Some of the particles in the matrix are close to the dislocations which cause pinning effect, and some are on the sub-grain boundary. The carbides on the grain boundaries are distributed in a chain shape which can prevent the sliding and migration of grain boundaries at high temperatures. Carbides in the crystal are obstacles to dislocation motion, which can pin the dislocations and is beneficial to strengthen the steel.

Fig. 8(e-g) is the TEM-EDS result of white rectangular area. Combined with the TEM-EDS result of area 3 (Fig. 8(g)) and the corresponding diffraction pattern of granular carbides (the inset of Fig. 8(d)), the precipitate of area 3 is confirmed to be NbC, which is also enriched with Mo and V. The EDS results also show that area 2 (Fig. 8f) is concentrated with Mn, which differs from the matrix (area 1, Fig. 8e), and it was confirmed by the corresponding diffraction pattern that the elongated particles is (Fe, Mn)3C (the inset of Fig. 8(a)). XRD analysis results of the precipitation extracted from the annealed specimens are shown in Fig. 9. It can be found that most of the precipitates were face-centered cubic NbC.

Fig. 9.

Fig. 9.   X-ray diffraction pattern of precipitates extracted from the annealed specimens after 70% deformation at 750 °C, 850 °C, 950 °C, and 1050 °C.


More than 5 TEM images of the specimens at each deformation temperature were used to measure the average particle size of the carbides. The distributions of the carbide particle sizes for the annealed specimens are presented in Fig. 10. As the heat compression temperature increased, the second phase particles got refined, while the coarsening of the carbide particles was severe at 1050 °C. When the deformation temperature was 750 °C, the particle size distribution of the precipitate particles was mainly in the range of 10-36 nm and the average carbide particle size was 30 nm, while it was 26 nm at 850 °C. When the compression temperature was at 950 °C, the particle size of the precipitated phase was concentrated below 20 nm. This phenomenon indicates that thermal deformation at the austenite region near the nose temperature (which has been measured to be 925 °C) can promote the precipitation of the small-sized second phase. However, as the temperature increases to 1050 °C, the proportion of precipitate particles with a size of 1-36 nm decreases dramatically, and coarsed particles of 36-60 nm in diameter get increased and the average carbide particle size of 36 nm appears. This is mainly because at high temperatures, the diffusion rate is faster, and the second phase particles are easily coarsened by ostwald ripening.

Fig. 10.

Fig. 10.   Particle size distribution of the precipitates in the annealed specimens.


3.3. Mechanical properties

Micro-tensile tests were performed to evaluate the mechanical properties of the experimental steel. The strength and plasticity for the quenched specimens and the annealed specimens after deformation at different temperatures are shown in Fig. 11. For the annealed specimens, the lowest yield strength 1072 MPa occurs at the deformation temperature of 750℃ and the lowest elongation 23.5% occurs at the deformation temperature of 850 ℃. The deformed specimens without annealing exhibit a relatively higher yield strength and tensile strength but a lower elongation in comparison with the annealed specimens. The mechanical properties of the specimens deformed at 950 ℃and annealed subsequently are the optimal parameters at which yield strength of 1204 MPa, elongation of 27% and the highest product of strength and elongation (32.5 GPa.%) could be obtained. The mechanical properties of advanced high strength steel are determined by the microstructural characteristics and the precipitation behavior, which will be discussed in detail.

Fig. 11.

Fig. 11.   The strength (a) and plasticity (b) of the experimental steel at the quenched and annealed state after deformation.


The microstructure of the quenched specimens exhibits almost martensite structure that helps to increase the ultimate strength as a strengthening phase, but the corresponding plasticity is poor. With the increase of the hot compression temperature, the size of the martensite lath gradually grows, and the tensile strength decreases slightly. When the deformation temperature reaches 1050 °C, the strength decreases significantly due to the coarsening of martensite lath. The finer the grain is, the more grains the deformation energy can be dispersed into during the plastic deformation process, resulting in a more uniform deformation without stress concentration and correspondingly the increase of the steel strength.

The tensile strength and yield strength of the annealed specimens got decreased, while the elongation and the product of strength were improved at the same trend along with the raise of the deformation temperature. After annealing, a certain amount of carbide is precipitated which would consume the carbon in the martensite and weaken the solid solution strengthening effect of carbon in the martensite matrix, and correspondingly lower the strength of the steel. The finely dispersed carbide on the martensite matrix can increase the yield strength and tensile strength to some extent by interacting with mobile dislocations and preventing them from moving, but it is not as significant as the weakening effect of solid solution strengthening. However, the fine precipitates get elongation and the product of strength and elongation increased effectively in comparison to the quenched samples.

Hong et al. [23] reported that the strength and toughness of 22 Mn (0.6C-22 Mn) steels were slightly enhanced after annealing at 650 °C for 30 min, although the intergranular cementite did grow. The toughness was considerably reduced, while the strength did not change much after annealing at 650 °C for 24 h due to the coarsening of the intergranular cementite. In general, the size and distribution of the cementite at the grain boundaries can affect the mechanical properties of the steels. As reported by Gutierrez-Urrutia et al. [24], the impingement of twins on the coarse grain boundary of carbides during the deformation produced local stress concentrations at the carbide interfaces. These grain boundary carbides provided the nucleation sites and propagation paths for cracking, causing intergranular cracking. Thus, the smaller the size of the grain boundary carbides was, the smaller the stress concentrations were, and thus, a smaller effect of the grain boundary of the carbides on the mechanical properties was observed. The presence of the cementite precipitated in the austenite grain is significantly helpful to increase the mechanical properties of the materials [25]. The superior mechanical properties of the advanced high strength steel can be explained in terms of the dispersion strengthening of intragranular nano-sized carbides and their interactions with the dislocation substructures during deformation.

In addition to the cementite, a high density of NbC precipitated in martensite matrix in the temperature range of 750 °C-1050 °C. The diffusion rate of the Nb atoms increased with the temperature. Nb atoms associated with vacancies and then formed localized clusters. As a result of the local migration of the niobium-vacancy groupings, the growth process from clusters through stable nuclei to precipitates took place and then, the NbC began to precipitate [[26], [27], [28]]. At temperature values of 750 °C to 950 °C, the increasing density of dislocations accelerated the diffusion of C and alloy atoms which was beneficial to the growth of the NbC precipitates in the martensite matrix. These dispersed precipitates blocked the dislocation lines, effectively pinning the grain boundaries, which refined the microstructure as well as improved the tensile strength in the hot compression. The formation of carbide precipitates consumed a large amount of C and other alloying elements (such as Nb, V, Fe, Mo, Mn) in the matrix. When heat compression temperature increased to 1050 °C, the fine NbC precipitates dissolved in order to decrease the interfacial energy. Although the dissolution rate of the second particle was high, these atoms did not have enough time to diffuse because of the rapid cooling and short duration of the deformation process. The solubility of the Nb and C atoms near coarse particles increased rapidly, and then the coarse precipitates continued to grow. As reported by Kamikawa et al. [29], the precipitate strengthening was lower for larger carbide radii. Hence, the decrease in strength at a temperature of 1050 °C may be attributed to the presence of a lower volume fraction of coarse carbides. However, the product of strength and elongation of the steel was improved because of the formation of the NbC nanoparticles. The dispersion of fine carbides improved the combined mechanical properties of steels, while the coarsening of these particles could decrease the toughness and strength. Hence, it is imperative to study the coarsening kinetics of the precipitates. The chosen microstructures and mechanical properties can be obtained by controlling the size of the precipitates.

4. Conclusions

The evolution of microstructure, precipitates and mechanical properties of the micro-alloyed medium Mn steel under different thermal deformation temperatures were investigated by microstructure observation and tensil tests. The following conclusions can be made:

(1) By analyzing the influence of thermal deformation temperature on the quenched medium Mn steel, it can be found that the lath martensite size was smaller as the deformation temperature became lower. And after the quenched specimens were annealed in the intercritical temperature range, the martensite laths became thicker and shorter and the distortion of the original laths disappeared. A large number of dislocations accumulate to form stable dislocation cell structures, and fine precipitates are scattered between the laths.

(2) In comparison with the undeformed and quenched specimen, it can be concluded that the acicular cementite was precipitated during the cooling process, while the granular NbC was deformation induced precipitate. As the deformation temperature increased, the average size of the second phase particles got decreased firstly from 30 nm to 20 nm at the deformation temperature range from 750 to 950 °C and then increased to 36 nm at the deformation temperature of 1050 °C.

(3) The tensile strength and yield strength of the annealed specimens got decreased, while the elongation and the product of strength and elongation were improved at the same trend along with the raise of the deformation temperature. As the deformation temperature increased, the product of strength and elongation first increased and then decreased. Compared with the high temperature deformation (1050 °C), the strength and plasticity of low temperature deformation have apparent advantages. For the tested medium Mn steel, the smallest average size of second phase particles (20 nm) and the best mechanical properties (32.5 GPa.%) can be obtained at the deformation temperature of 950 °C.

Acknowledgement

This work was supported by the National Key Research and Development Program [Grant No. 2018YFA0702900], the National Natural Science Foundation of China [Grant No. U1508215, 51774265], the National Science and Technology Major Project of China [Grant No. 2019ZX06004010], the Key Program of the Chinese Academy of Sciences [Grant No. ZDRW-CN-2017-1], the Key Program of Natural Science Foundation of Hebei Province of China [Grant No. E2017203161], and the CAS Interdisciplinary Innovation Team.

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