crystals
Article Effect of Annealing Process on Microstructure, Texture, and Mechanical Properties of a Fe-Si-Cr-Mo-C Deep Drawing Dual-Phase Steel
Hongbo Pan 1,2, Xiaohui Shen 3,*, Dongyang Li 2,*, Yonggang Liu 4, Jinghua Cao 3, Yaqiang Tian 5, Hua Zhan 4, Huiting Wang 3, Zhigang Wang 6 and Yangyang Xiao 4
1 Key Laboratory of Metallurgical Emission Reduction & Resources Recycling of Ministry of Education, Anhui University of Technology, Ma’anshan 243002, China; [email protected] 2 The Department of Chemical and Materials Engineering, University of Alberta, Edmonton, AB T6G1H9, Canada 3 School of Metallurgical Engineering, Anhui university of Technology, Ma’anshan 243032, China; [email protected] (J.C.); [email protected] (H.W.) 4 Technical center, Ma’anshan Iron and Steel Co., Ltd., Ma’anshan 243000, China; [email protected] (Y.L.); [email protected] (H.Z.); [email protected] (Y.X.) 5 Key Laboratory of the Ministry of Education for Modern Metallurgy Technology, North China University of Science and Technology, Tangshan 063210, China; [email protected] 6 Faculty of Materials Metallurgy and Chemistry, Jiangxi University of Science and Technology, Ganzhou 341001, China; [email protected] * Correspondence: [email protected] (X.S.); [email protected] (D.L.)
Received: 13 August 2020; Accepted: 31 August 2020; Published: 2 September 2020
Abstract: Dual phase steel generally has poor deep drawing property with a low r value less than 1.0, making it difficult to be used for deep drawing automotive parts. In order to improve the mechanical properties of the steel through heat treatment, effect of heat treatments with different conditions on a Fe-Si-Cr-Mo-C deep drawing dual-phase steel was investigated with the aim of identifying effective heat treatment parameters for effective modification towards optimal properties. Relevant thermal dilation and heat treatment experiments were performed. Corresponding characters were investigated. The results show that island martensite can be obtained at low cooling rate. With the increase of cooling rate, the formation of pearlite and bainite is favored. During annealing at low temperatures, recrystallization of the steel is incomplete with the presence of the shear bands. With the increase of annealing temperature, the recrystallization process is gradually complete, and the number of high angle grain boundaries increases significantly. The ratio of gamma orientation components to alpha orientation components decreases first and then increases with the increase of annealing temperature. The strain hardening exponent and r value show an upward trend with respect to annealing temperature, and the r value is as high as 1.15.
Keywords: deep-drawing dual-phase steel; mechanical properties; microstructure; recrystallization; texture
1. Introduction Low carbon gas emission, information, and intelligence are the main components for the future trend of automotive technological advance [1–4]. The automotive industry is one of strategically important industrial sectors, which affects the technological progress and social modernization, and plays a key role in economic growth [5,6]. The development of automotive industry is largely related to steel technology. The automobile manufacturing industry is the largest user of steel sheet.
Crystals 2020, 10, 777; doi:10.3390/cryst10090777 www.mdpi.com/journal/crystals Crystals 2020, 10, 777 2 of 16
In recent years, with the increasingly rigorous implementation of safety and fuel economy regulations, advanced steels with higher strength and toughness are demanded by the automotive industry [7]. Vehicle weight and engine power are two of the most important parameters that influence a vehicle’s fuel consumption and carbon dioxide (CO2) emissions, among which the most effective method is to reduce the vehicle weight [8,9]. The use of lightweight components is a fundamental requirement in material engineering. Owing to the economic and ecological considerations, the weight of automotive structure should be reduced. In order to reduce the structure’s weight while retaining required strength, materials for the structure must be stronger and tougher [10–12]. Profound progress in automobile steel manufacturing has been achieved through the development of advanced high-strength steels (AHSS), fueled by the conflicting demands on the automotive industry to simultaneously improve crash safety and fuel economy [13–15]. Dual-phase (DP) steel is the flagship of advanced high-strength steels, which was the first among various candidate alloy systems to find application in light-weight automotive components [16]. With relatively simple thermomechanical treatment and lean alloying, a hard martensitic or bainitic phase is dispersed in a ductile ferritic matrix [17–19], providing a series of excellent and industrially accessible mechanical properties [16,18]. The dual-phase steel exhibits an excellent combination of strength and ductility, which has higher ultimate tensile strength, lower initial yielding stress, higher initial stage strain hardening, high ductility,and high bake hardening [17,20–22]. These high-level mechanical properties obtained in the DP steels lend good energy-absorption capacity, making it suitable for use in structural parts and reinforcements [19,22]. The high strain hardening combined with a pronounced bake hardening effect gives the DP steel great potential for reducing the weight of structural parts and even skin parts [23,24]. Dual-phase steel is a low carbon or low carbon micro-alloying steel, which can be produced by hot rolling or cold rolling [25–27]. The rolling temperature and cooling process influence the microstructure of the steel, e.g., transformation of a ferrite-martensite structure from austenite during hot rolling [28,29]. For cold rolling, properties of the steel can also be modified on the continuous annealing lines where there is even greater control over thermal treatment [30–32]. However, due to the complex microstructure of the steel, from a scientific perspective, a full understanding and well guided control of microstructure evolution in DP steel for optimized properties have not been achieved yet [16]. Currently, traditional DP steel is not a good candidate for applications that require high drawability. The DP steel usually exhibits rather poor plastic strain ratio value (r value), rendering it difficult to be used as the skin parts by stamping as well as the structure parts with high deep drawing performance [33–38]. This drawback, however, could be eliminated by adding carbide forming elements that may induce precipitation at low temperatures or act as solute for solution hardening at high temperatures [34,36,37]. Alternatively, the volume fraction of martensite can be reduced, leading to decrease the texture density and fraction of α fiber [35,37]. However, there are still open questions, mainly related to the through-process microstructure and texture development as well as the structure–property relationships [16]. It is not easy to answer the questions, ascribed to the complexity involved in the effect of processing conditions on the mechanical behavior of DP steels because of the mutual influences among the processing- and composition-dependent microstructural parameters [16]. In this study, with the specific consideration for deep drawing capability and theoretical understanding of dual-phase steel, a low carbon 1.0Si-0.15Cr-0.5Mo deep drawing dual phase steel was designed and investigated. The main objective of study is to determine the process–structure–property relationships with the eventual goal to effectively control microstructure during continuous annealing and cooling processes for maximizing deep drawing capability (plastic strain ratio) alone with other optimized mechanical properties. Crystals 2020, 10, 777 3 of 16
2. Experimental Material and Procedure
2.1. Material Design The material under study is selected mainly based on its deep drawing property and proper hardenability to form a certain amount of martensite, which could be adjusted by alloying with selected elements. Carbon can act as an austenite stabilizer and affect the formation of martensite at practical cooling rates. However, higher carbon contents would inhibit the formation of γ fiber texture, and then deteriorates the deep drawing property of the steel [39]. As discussed later, texture can induce plastic anisotropy and influence the formability of a metallic material. Mn can improve not only the stability of austenite but also strengthen ferrite by solution-hardening as well as retard ferrite formation [40,41]. Mo and Cr could be used to partially replace C and Mn for sufficient hardening ability, since they help retard the formation of pearlite or bainite and thus keep a sufficient amount of martensite as well as help enhance the solute-hardening effect. Mo addition is more effective than Cr and Mn to obtain the required volume fraction of martensite and higher mechanical strength [36,42,43]. Silicon promotes ferrite transformation and has a negligible solubility in cementite, thus suppressing the precipitation of cementite [44–46]. Furthermore, instead of cementite precipitation, carbon partitions from ferrite into untransformed austenite. The carbon enrichment increases the stability of the austenite and thus allows most of the austenite to transform into martensite in the temperature range of about 300 ◦C even under the slow cooling condition, due to the increased hardenability attained by the enriched carbon and the presence of molybdenum, chromium, manganese, etc., in solution. Nb can react with C in steel to form niobium carbide, which can dissolve and stabilize austenite at high temperature [47]. Based on the above consideration, a novel steel is designed which is denoted as Fe-0.025C-1.0Si-1.0Mn-0.15Cr-0.50Mo (wt %). Details of chemical composition of the steel are given in Table1. The A e1 and Ae3 temperatures calculated by Thermo-calc software are 740 and 943 ◦C, respectively.
Table 1. Chemical composition of the experimental steels (mass fraction, %).
C Si Mn Nb Cr Mo Al P S N 0.026 0.98 1.09 0.041 0.15 0.50 0.08 0.007 0.003 60 ppm ≤
2.2. Experimental Procedure The designed steel was fabricated by vacuum induction melting. The ingot of 100 kg was forged to make slabs of 30 mm in thickness, which were reheated to 1200 ◦C and held for 2 h, followed by hot-rolling at a temperature between 850 and 1100 ◦C to reduce their thickness to 3 mm. After water cooling to 600 ◦C, the plates were placed in a box resistance furnace for 2 h, and then cooled in the furnace to simulate coiling. The microstructure of the as-rolled plates was a mixture of ferrite and pearlite. The plates were pickled in a 10% hydrochloric acid solution and further cold-rolled to reduce their thickness to 1.0 mm using a four high reversing mill. ϕ 4 10 mm small cylinders were made from the forged slabs by wire cutting and machining. × The critical temperatures were measured, and the continuous cooling experiments were performed on a Bähr DIL805 A quenching dilatometer (Baehr-Thermo, New Castle, Germany). The continuous cooling experiments were arranged as follows. For austenitization and homogenization, the samples were first heated to 1000 ◦C at a heating rate of 10 ◦C/s and held for 3 min, followed by cooling to the room temperature at cooling rates of 0.5, 1, 3, 5, 10, 15, 20, and 30 ◦C/s, respectively. Samples were annealed in a salt bath at 790, 820, 850, 880, and 900 ◦C inter-critical temperature, respectively, for 600 s, then transferred within 3 s to a salt bath at 300 ◦C and held there for 420 s, and then water-cooled to room temperature. The entire process is illustrated in Figure1. Crystals 2020, 10, 777 4 of 16 Crystals 2020, 10, x FOR PEER REVIEW 4 of 16
Figure 1. Schematic illustration of annealing process. Figure 1. Schematic illustration of annealing process. Samples were sequentially ground with 100, 200, 400, 600, and 800 grit SiC papers, polished Samples were sequentially ground with 100, 200, 400, 600, and 800 grit SiC papers, polished with with particle diameter of 5 and 1.5 µm diamond grinding paste and then etched in a 4% nital particle diameter of 5 and 1.5 μm diamond grinding paste and then etched in a 4% nital solution. solution. Microstructure characterization was carried out using an Olympus BX51 optical microscope Microstructure characterization was carried out using an Olympus BX51 optical microscope (Olympus, Taoyuan, Japan) and FEI Nova NanoSEM 430 field emission scanning electron microscopy (Olympus, Taoyuan, Japan) and FEI Nova NanoSEM 430 field emission scanning electron microscopy (FEI, Eindhoven, Netherlands). Electron backscatter diffraction (EBSD) analysis was conducted using a (FEI, Eindhoven, Netherlands). Electron backscatter diffraction (EBSD) analysis was conducted using Zeiss-SIGMA field-emission scanning electron microscope (Carl Zeiss AG, Cambridge, UK) with a a Zeiss-SIGMA field-emission scanning electron microscope (Carl Zeiss AG, Cambridge, UK) with a step size of 0.6 µm. Prior to the EBSD analysis, the samples were sequentially ground with 400, 800, step size of 0.6 μm. Prior to the EBSD analysis, the samples were sequentially ground with 400, 800, 1000, and 2000 grit SiC papers and electropolished at room temperature in a solution of 8% perchloric 1000, and 2000 grit SiC papers and electropolished at room temperature in a solution of 8% perchloric acid alcohol solution at an operating voltage of 32 V. For tensile tests, specimens with 50 mm gauge acid alcohol solution at an operating voltage of 32 V. For tensile tests, specimens with 50 mm gauge length and 12.5 mm gauge width were machined from the sheet with the gauge length parallel to length and 12.5 mm gauge width were machined from the sheet with the gauge length parallel to the the rolling direction. The tensile tests with a tensile speed of 2 mm/min were performed using a rolling direction. The tensile tests with a tensile speed of 2 mm/min were performed using a Zwick/Roell Z150 universal electronic tensile testing machine (ZwickRoell, Ulm, Germany) at the room Zwick/Roell Z150 universal electronic tensile testing machine (ZwickRoell, Ulm, Germany) at the temperature. Vickers hardness of samples was measured using a HVT-1000 A digital microhardness room temperature. Vickers hardness of samples was measured using a HVT-1000 A digital tester (Huayin, Laizhou, China). Five different positions were measured for each sample, and the microhardness tester (Huayin, Laizhou, China). Five different positions were measured for each reported value is an average of the results from the five measurements. sample, and the reported value is an average of the results from the five measurements. 3. Results and Discussion 3. Results and Discussion 3.1. The Phase Transformation Behavior 3.1. The Phase Transformation Behavior The dilatometer can be used to accurately measure dilatation of a sample with temperature, and theThe phase dilatometer transformation can be used temperatures to accurately can measur be determinede dilatation from of the a sample dilatation-temperature with temperature, curves. and Thethe phase dilatation transformation curve of the temperatures sample heated can at be 10 dete◦C/rmineds is illustrated from the in dilatati Figureon-temperature2a. From the dilation curves. curve,The dilatation one may curve determine of the sample the transformation heated at 10 temperatures,°C/s is illustrated Ac1 inand Figure Ac3 ,2a. which From are the 850 dilation and 973 curve,◦C, respectively.one may determine A partial the equilibrium transformation phase diagramtemperatures, calculated Ac1 byand Thermo-calc Ac3, which software are 850 (Thermo-Calc and 973 °C, Softwarerespectively. AB, A Solna, partial Sweden) equilibrium is shown phase indiagram Figure ca2b,lculated from which by Thermo-calc it is concluded software that (Thermo-Calc the critical temperaturesSoftware AB, of Solna, Ae1 and Sweden) Ae3 are is 740 sh andown 943 in◦ FigureC, respectively. 2b, from The which results it is show concluded that the heatingthat the ratecritical has atemperatures large influence of onAe1 the and transformation Ae3 are 740 and temperature 943 °C, respectively. of pearlite but The little results influence show on that the the transformation heating rate temperaturehas a large influence of ferrite. on This the is transformation mainly ascribed temperature to the fact thatof pearlite the diff usionbut little of alloyinginfluence elements on the istransformation slow at low temperatures, temperature of thus ferrite. requiring This is amainly long timeascribed for theto the phase fact transformationthat the diffusion to of complete. alloying Forelements martensitic is slow transformation, at low temperatures, its starting thus temperature requiringMs a long can betime estimated for the usingphase an transformation empirical model, to whichcomplete. is influenced For martensitic by the concentrationstransformation, of its alloying starting elements temperature in wt %Ms [48 can]: be estimated using an empirical model, which is influenced by the concentrations of alloying elements in wt % [48]: Ms[◦C] = 561 423 C 30.4 Mn 17.7 Ni 12.1 Cr 11 Si 7 Mo Ms[°C] = 561− − 423·C· − − 30.4·Mn· − − 17.7·Ni· −− 12.1·Cr· − 11·Si· −− 7·Mo·
ForFor thethe designeddesigned steel steel (see (see Table Table1), 1), its its martensite martensite starting starting temperature temperature is calculated is calculated to be to 500.8 be 500.8◦C. °C.
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FigureFigure 2.2. DilatationDilatation curve curve of aof sample a sample heated heated at 10 ◦atC /10s ( a°C/s) and ( ana)) equilibriumandand anan equilibriumequilibrium phase diagram phasephase calculated diagramdiagram bycalculated Thermo-calc by Thermo-calc software (b software). (b).).
3.2.3.2. Effect Effect of Cooling Rate on MicrostructureMicrostructure EvolutionEvolution andand HardnessHardness MicrostructuresMicrostructures of samplessamples cooled at didifferentfferent coolingcooling ratesrates areare depicteddepicted inin FigureFigure3 .3. AtAt lowlow coolingcooling ratesrates ofof 0.5 and 1 °C/s,◦C/s, microstructures microstructures of the the samples samples are are composed composed of of ferrite ferrite and and M/A MM/A/A islands islands asas FigureFigure3 3a,ba,b illustrate. illustrate. PearlitePearlite andand bainitebainite areare notnott observed observedobserved in inin the thethesamples. samples.samples. InInIn orderorderorder to toto reveal revealrevealthe thethe MM/A/A islandsislands more clearly,clearly, thethe areasareas markedmarked byby thethe redred dotteddotted boxesboxes inin FigureFigure3 3a,ba,b are are enlarged enlarged and and shownshown in the upper right corner of the figures, figures, wh whereere the M/A M/A islands are convex and clearly clearly visible. visible. WhenWhen thethe coolingcooling raterate increases increases to to 3 3◦ C°C/s,/s, pearlite pearlite appears appears in in the the microstructure, microstructure, which which is composedis composed of ferrite,of ferrite, pearlite, pearlite, and and M/A M/A islands islands as Figure as Figure3c shows. 3c shows. At cooling At cooling rates of 5rates and of 10 5◦ Cand/s, bainite10 °C/s, begins bainite to appearbegins to in appear the microstructures, in the microstructures, which consist which of ferrite,consist pearlite,of ferrite, bainite, pearlite, and bainite, M/A islands.and M/A However, islands. theHowever, size and the content size and of pearlitecontent andof pearlite M/A decrease and M/A at decrease the higher at cooling the higher rates cooling as shown rates in as Figure shown3d,e. in AsFigure the cooling3d,e. As rate the is cooling equal to rate or higheris equal than to or 15 high◦C/s,er the than M /15A islands°C/s, the and M/A pearlite islands in and the microstructure pearlite in the disappear.microstructure In this disappear. case, the microstructuresIn this case, the aremicrostructures composed of are a small composed amount of of a fine small ferrite amount and aof large fine amountferriteferrite andand of bainite. aa largelarge With amountamount increasing ofof bainite.bainite. the cooling WithWith rate,increasingincreasing the content thethe cooling ofcooling bainite rate,rate, continuously thethe contentcontent increases ofof bainitebainite and thecontinuously microstructure increases becomes and the finer microstructure as Figure3f–h becomes illustrate. finer as Figure 3f–h illustrate.
FigureFigure 3. SEMSEM micrographs micrographs of of samples samples cooled cooled at at (a ())a 0.5,0.5,) 0.5, ((b ())b 1,1,) 1,((c)) ( c3,3,) 3,((d) () d5,5,) ( 5,(e)) ( 10,e10,) 10, ((ff)) (15,15,f) 15, ((g)) (20,20,g) 20,andand and ((h)) (30h) °C/s, 30 ◦C respectively./s, respectively. M/A: M /martensite-austeniteA: martensite-austenite island; island; P: pearlite; P: pearlite; B: bainite. B: bainite.
HardnessHardness valuesvalues ofof thethe samples samples cooled cooled at at di differentfferent cooling cooling rates rate weres were measured measured and and are are shown shown in Figureinin FigureFigure4. As 4.4. AsAs illustrated, illustrated,illustrated, the thethe hardness hardnesshardness does doesdoes not notnot change change much much when when the the samples samples are are cooled cooled at at cooling cooling rates of 0.5 and 1 °C/s, which show similar microstructures consisting of coarse ferrite and a small amount of martensitic islands. With increasing the cooling rate, hardness of the steel increases due to
Crystals 2020, 10, 777 6 of 16
rates of 0.5 and 1 ◦C/s, which show similar microstructures consisting of coarse ferrite and a small amountCrystals 2020 of, martensitic10, x FOR PEER islands. REVIEW With increasing the cooling rate, hardness of the steel increases due6 of to16 the microstructure refinement and the generation of a certain amount of bainite. The hardness increases the microstructure refinement and the generation of a certain amount of bainite. The hardness slowly within the cooling rate range of 3–10 ◦C/s, shown by the small slope of the hardness-cooling increases slowly within the cooling rate range of 3–10 °C/s, shown by the small slope of the hardness- rate curve, because the samples have similar microstructures with only slight changes in grain size and cooling rate curve, because the samples have similar microstructures with only slight changes in grain bainite content. After the cooling rate is further increased to and above 10 ◦C/s, the slope of the curve size and bainite content. After the cooling rate is further increased to and above 10 °C/s, the slope of increases significantly, accompanied with obvious changes in microstructure as Figure3 illustrates. the curve increases significantly, accompanied with obvious changes in microstructure as Figure 3 The microstructures corresponding to the high cooling rates are composed mainly of bainite with a illustrates. The microstructures corresponding to the high cooling rates are composed mainly of small amount of fine ferrite. Further increasing the cooling rate does not change the slope of the curve bainite with a small amount of fine ferrite. Further increasing the cooling rate does not change the much and corresponding microstructures are rather similar. slope of the curve much and corresponding microstructures are rather similar.
Figure 4. Hardness versus the cooling rate.
3.3. Effect Effect of Annealing Temperature onon Microstructure,Microstructure, Texture,Texture, and and Corresponding Corresponding Mechanical Mechanical Properties Properties 3.3.1. Microstructure 3.3.1. Microstructure As only a very small fraction of carbon can be dissolved in the ferrite, a substantial fraction of As only a very small fraction of carbon can be dissolved in the ferrite, a substantial fraction of carbon is available for stabilizing austenite in the low alloyed steel during inter-critical annealing [48–50]. carbon is available for stabilizing austenite in the low alloyed steel during inter-critical annealing [48– During inter-critical annealing, a higher carbon content is obtained in the austenite and the 50]. During inter-critical annealing, a higher carbon content is obtained in the austenite and the carbon-rich austenite transforms into martensite during subsequent cooling [49]. The metallographic carbon-rich austenite transforms into martensite during subsequent cooling [49]. The metallographic microstructures of the samples at different inter-critical annealing temperatures are shown in Figure5. microstructures of the samples at different inter-critical annealing temperatures are shown in Figure It can be seen from Figure5 that there is basically no di fference in microstructure among samples 5. It can be seen from Figure 5 that there is basically no difference in microstructure among samples annealed at 850 C and below except there is a little deformed ferrite in samples annealed at 820 C and annealed at 850◦ °C and below except there is a little deformed ferrite in samples annealed at ◦820 °C below. That is to say, when annealed at 820 C and below, the samples are not completely recrystallized. and below. That is to say, when annealed◦ at 820 °C and below, the samples are not completely The microstructures are all composed of predominant ferrite and carbide, among which is the result recrystallized. The microstructures are all composed of predominant ferrite and carbide, among that stable chromium carbides are not completely dissolved due to the low inter-critical annealing [51]. which is the result that stable chromium carbides are not completely dissolved due to the low inter- Such a low degree of austenitizing results in little martensite as shown in Figure5a–c. When the critical annealing [51]. Such a low degree of austenitizing results in little martensite as shown in annealing temperature is increased to 880 C, the ferrite is fully recrystallized and the carbide is Figure 5a–c. When the annealing temperature◦ is increased to 880 °C, the ferrite is fully recrystallized completely dissolved for stabling austenite, resulting in a microstructure consisting of polygonal and the carbide is completely dissolved for stabling austenite, resulting in a microstructure consisting ferrite and island martensite. Island martensite having a higher carbon content is formed at the grain of polygonal ferrite and island martensite. Island martensite having a higher carbon content is formed boundary between ferritic grains and/or at the triple-points. This indicates that the full potential for at the grain boundary between ferritic grains and/or at the triple-points. This indicates that the full stabilizing the austenite has been unlocked at the employed annealing temperature (see Figure5d). potential for stabilizing the austenite has been unlocked at the employed annealing temperature (see When the annealing temperature is further increased to 900 C, the microstructure is the same as that Figure 5d). When the annealing temperature is further increased◦ to 900 °C, the microstructure is the same as that of 880 °C, but the content of island martensite is largely reduced (see Figure 5e). This is attributed to the fact that at the high annealing temperature, the carbon content in austenite is not enough to stabilize it and form martensite during subsequent cooling.
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of 880 ◦C, but the content of island martensite is largely reduced (see Figure5e). This is attributed to the fact that at the high annealing temperature, the carbon content in austenite is not enough to stabilizeCrystals 2020 it, and10, x formFOR PEER martensite REVIEW during subsequent cooling. 7 of 16
Figure 5. 5. OpticalOptical micrographs micrographs of specimens of specimens annealed annealed at (a) 790, at ( a(b)) 790,820, ( (cb)) 850, 820, (d ()c 880,) 850, and ( d(e)) 880,900
and°C, respectively. (e) 900 ◦C, respectively.
Figure6 6 presents presents SEM SEM images images of of samples samples annealed annealed at di atff differenterent temperatures, temperatures, followed followed by salt by bath salt quenching,bath quenching, which which provide provide more details more aboutdetails microstructure about microstructure evolution. evolution. As shown, As at shown, the annealing at the temperatureannealing temperature of 790 ◦C, recrystallizationof 790 °C, recrystallization of ferrite precedesof ferrite theprecedes formation the formation of austenite. of austenite. There are There a lot ofare non-recrystallized a lot of non-recrystallized ferrites in ferrites the microstructure, in the microstructure, which contain which a contain lot of shear a lot bands, of shear as shownbands, inas Figureshown6 a.in AtFigure the annealing6a. At the temperature annealing temperature of 820 ◦C, recrystallization of 820 °C, recrystallization of ferrite and of the ferrite formation and the of austeniteformation proceed of austenite simultaneously. proceed si Theremultaneously. are still non-recrystallized There are still non-recrystallized ferrites in the microstructure ferrites in with the amicrostructure certain amount with of sheara certain bands, amount and aof small shear amount bands, ofand island a small martensite amount isof formed, island martensite as shown inis Figureformed,6b. as When shown the in annealing Figure 6b. temperature When the is annealin further raisedg temperature to 850 ◦C, is the further recrystallization raised to 850 of ferrite°C, the is basicallyrecrystallization completed of ferrite with only is basically a small amountcompleted of non-recrystallized with only a small ferrite amount left inof thenon-recrystallized microstructure, andferrite the left austenitizing in the microstructure, process is slightly and promotedthe austenitizing as Figure process6c illustrates. is slightly As the promoted annealing temperatureas Figure 6c isillustrates. further increased As the annealing to 880 ◦ C,temperature the recrystallization is further increased of ferrite to and 880 formation °C, the recrystallization of austenite have of ferrite been fullyand formation implemented, of austenite the ferrite have grains been growfully implemented, up obviously withthe ferrite more grains island grow martensite up obviously as Figure with6d illustrates.more island The martensite recrystallized as Figure grains grow6d illustrates further and. The the recrystallized austenitizing processgrains isgrow more further complete and at the annealingaustenitizing temperature process is of more 900 ◦ Ccomplete but the contentat the a ofnnealing island martensite temperature appears of 900 to °C be reduced,but the content as shown of inisland Figure martensite6e. appears to be reduced, as shown in Figure 6e.
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Figure 6.6. SEM micrographs of specimens annealed at ( a) 790, ( b) 820, ( c) 850, (d) 880, and (e) 900 ◦°C,C, respectively.respectively. MM/A:/A: martensite-austenitemartensite-austenite island.island. 3.3.2. Mechanical Properties 3.3.2. Mechanical Properties Mechanical properties including initial stress–strain curves, ultimate tensile strength, yield strength, Mechanical properties including initial stress–strain curves, ultimate tensile strength, yield elongation, yield ratio, strain hardening exponent (n value), and plastic strain ratio (r value) of the strength, elongation, yield ratio, strain hardening exponent (n value), and plastic strain ratio (r value) samples were determined, results of which are illustrated in Figure7. For all the specimens, their yield of the samples were determined, results of which are illustrated in Figure 7. For all the specimens, strengths and yield ratios decrease with increasing the annealing temperature, while the ultimate tensile their yield strengths and yield ratios decrease with increasing the annealing temperature, while the strength decreases with increasing annealing temperature up to 880 C and then it rises. The elongation ultimate tensile strength decreases with increasing annealing temperature◦ up to 880 °C and then it increases by 41.3% with increasing the annealing temperature up to 850 C, then decreases by 20.6% rises. The elongation increases by 41.3% with increasing the annealing temperature◦ up to 850 °C, then at the annealing temperature of 880 C, and finally increases rapidly. Strain hardening exponent decreases by 20.6% at the annealing◦ temperature of 880 °C, and finally increases rapidly. Strain (n value) rises rapidly with increasing the annealing temperature, which is in the range between 0.12 hardening exponent (n value) rises rapidly with increasing the annealing temperature, which is in and 0.20. The plastic strain ratio (r value) increases quickly with increasing the annealing temperature the range between 0.12 and 0.20. The plastic strain ratio (r value) increases quickly with increasing up to 850 C, then decreases by 1.4% at 880 C, and finally increases again but slowly. The r value the annealing◦ temperature up to 850 °C, then◦ decreases by 1.4% at 880 °C, and finally increases again is in the range between 0.65 and 1.14. As mentioned earlier, when the sample is annealed at a lower but slowly. The r value is in the range between 0.65 and 1.14. As mentioned earlier, when the sample temperature, its recrystallization process is incomplete, resulting in high strength, low elongation, is annealed at a lower temperature, its recrystallization process is incomplete, resulting in high low n value, and low r value. With increasing the annealing temperature, the recrystallization process strength, low elongation, low n value, and low r value. With increasing the annealing temperature, and the dissolution of carbides become more and more complete, resulting in the decrease in strength the recrystallization process and the dissolution of carbides become more and more complete, but increase in elongation, n and r values. A large number of dissolved carbides can stabilize austenite resulting in the decrease in strength but increase in elongation, n and r values. A large number of which is supposed to transform to martensite during the following cooling process, resulting in dissolved carbides can stabilize austenite which is supposed to transform to martensite during the slight decrease in elongation and r value [52,53]. As the annealing temperature increases to 900 C, following cooling process, resulting in slight decrease in elongation and r value [52,53]. As the◦ the stability of carbon-rich austenite decreases, thus reducing the content of martensite and leading to annealing temperature increases to 900 °C, the stability of carbon-rich austenite decreases, thus a consequent increase in elongation and r value. reducing the content of martensite and leading to a consequent increase in elongation and r value.
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Figure 7. Mechanical properties of specimens annealed at different temperatures. (a) Stress–strain Figurecurves; 7. (b )Mechanical relation between properties strength of specimens and temperature; annealed (c) relationshipat different temperatures. between elongation, (a) Stress–strain yield ratio, curves;and temperature; (b) relation (d )between relationship strength between andn temperature;value, r value (c and) relationship temperature. between elongation, yield ratio, and temperature; (d) relationship between n value, r value and temperature. 3.3.3. Texture 3.3.3. Texture Figure8 displays the key BCC (body center cubic) texture components in the ϕ2 = 45◦ section of the EulerFigure space 8 displays and the the calculated key BCC Orientation (body center Distribution cubic) texture Function components (ODF) in at theϕ2 =φ245 = ◦45°for section different of theannealing Euler space temperatures and the [calculated54,55]. The Orientation textures of Distribution the samples afterFunction annealing (ODF) show at φ2 alpha = 45° for {uvw} different<110> annealingand gamma temperatures {111}
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Figure 8. ϕ2 = 45° sections of Orientation Distribution Function (ODF) at surface layers in the
Figurespecimens 8.8. ϕϕ 2annealed2= =45 45°◦ sections sections at different of Orientationof Orientation temperatures. Distribution Distribution (a) Main Function BCCFunction (ODF)texture (ODF) at components surface at layerssurface shown in thelayers specimensin ϕin2 =the 45° ϕ specimensannealedsection of at theannealed di ffEulererent space; temperatures.at different (b) 850 temperatures. °C; (a) ( Mainc) 880 BCC °C; ( a( texture)d Main) 900 °C. componentsBCC texture shown components in 2 = shown45◦ section in ϕ2 of = 45° the sectionEuler space; of the ( bEuler) 850 space;◦C; (c) ( 880b) 850◦C; °C; (d) ( 900c) 880◦C. °C; (d) 900 °C. 4. Discussion 4. Discussion 4. Discussion 4.1. Alloy Elements Diffusion 4.1. Alloy Elements Diffusion 4.1. AlloyIf cooling Elements rate Diffusion is slow after soaking at full austenitizing or the α + γ region high temperature, the If cooling rate is slow after soaking at full austenitizing or the α + γ region high temperature, alloy elements in the transformed ferrite will have enough time to diffuse to untransformed austenite the alloyIf cooling elements rate is in slow the after transformed soaking at ferrite full austenitizing will have enough or the timeα + γ toregion diffuse high to temperature, untransformed the domains and stabilize them, so that this part of austenite will finally transform into island martensite. alloyaustenite elements domains in the and transformed stabilize them, ferrite so will that have this enough part of austenitetime to diffuse will finally to untransformed transform into austenite island However, if the cooling rate is high, there is no sufficient time for diffusion during the phase domainsmartensite. and However, stabilize them, if the so cooling that this rate part is high,of austenite there is will no finally sufficient transform time for into diff islandusion martensite. during the transformation process [56], so the original austenite will eventually transform into lath bainite as phaseHowever, transformation if the cooling process rate [ 56is ],high, so the there original is no austenite sufficient will time eventually for diffusion transform during into laththe bainitephase illustrated in Figure 3. It is a complex process combining recrystallization, carbide dissolution, and transformationas illustrated in process Figure 3[56],. It isso a the complex original process austenit combininge will eventually recrystallization, transform carbide into lath dissolution, bainite as austenite transformation during inter-critical annealing [16]. Schematic illustrations of phase illustratedand austenite in Figure transformation 3. It is a complex during inter-criticalprocess combining annealing recrystallization, [16]. Schematic carbide illustrations dissolution, of phase and transformation during inter-critical annealing are presented in Figure 9. In the case of annealing at austenitetransformation transformation during inter-critical during inter-critical annealing are annealing presented [16]. in Figure Schematic9. In the illustrations case of annealing of phase at low temperatures, a large amount of carbide is not dissolved and only a small amount of austenite is transformationlow temperatures, during a large inter-critical amount of annealing carbide is are not presented dissolved in and Figure only 9. a In small the amountcase of annealing of austenite at formed, so the content of martensite in the microstructure after cooling is very small as shown in lowis formed, temperatures, so the content a large amount of martensite of carbide in the is not microstructure dissolved and after only cooling a small is amount very small of austenite as shown is Figures 9a and 6b,c. When the annealing temperature is moderate, the carbide dissolves and stabilizes formed,in Figure so6b,c the and content Figure of9 a.martensite When the in annealingthe micros temperaturetructure after is cooling moderate, is very the carbidesmall as dissolves shown in the austenite, which transforms into martensite during the subsequent cooling process. A higher Figuresand stabilizes 9a and the 6b,c. austenite, When the which annealing transforms temperatur into martensitee is moderate, during the thecarbide subsequent dissolves cooling and stabilizes process. content of martensite can be consequently obtained as shown in Figures 9b and 6d. At higher Athe higher austenite, content which of martensite transforms can into beconsequently martensite during obtained the as subsequent shown in Figures cooling6d process. and9b. At A higherhigher annealing temperatures, the increase in the austenite content reduces the content of alloying elements contentannealing of temperatures,martensite can the be increase consequently in the austenite obtained content as shown reduces in theFigures content 9b ofand alloying 6d. At elements higher and lowers stability, leading to the decrease in the martensite content during the subsequent cooling annealingand lowers temperatures, stability, leading the increase to the decrease in the austenite in the martensite content reduces content the during content the of subsequent alloying elements cooling process, as shown in Figures 9c and 6e. process,and lowers as shownstability, in leading Figures 6toe andthe 9decreasec. in the martensite content during the subsequent cooling process, as shown in Figures 9c and 6e.
Figure 9. Schematic illustration of phase transformation during inter-critical annealing. Figure 9. Schematic illustration of phase transformation during inter-critical annealing. (a) Low (a) Low annealing temperature; (b) moderate annealing temperature; (c) high annealing temperature. Figureannealing 9. Schematic temperature; illustration (b) moderate of phase annealing transforma temperature;tion during (c) highinter-critical annealing annealing. temperature. (a) Low(The (The color in the figure represents the content of alloy elements in austenite schematically.) annealingcolor in the temperature; figure represents (b) moderate the content annealing of alloy temperature; elements in au (cstenite) high schematically.)annealing temperature. (The 4.2. Ferritecolor in Recrystallization the figure represents and Growththe content of alloy elements in austenite schematically.) 4.2. Ferrite Recrystallization and Growth The recrystallization and grain growth of ferrite during annealing play an important role in 4.2. Ferrite Recrystallization and Growth affectingThe mechanicalrecrystallization properties and grain of the growth steel under of ferrite study. during In the annealing case of incomplete play an important recrystallization role in ofaffecting ferrite,The recrystallizationmechanical the sample willproperties haveand grain high of the strength,growth steel under of high ferrite study. yield during In ratio, the lowannealingcase elongation,of incomplete play lowan recrystallizationimportantn value, and role low in of affectingferrite, the mechanical sample will properties have high of thestrength, steel under high study.yield ratio, In the low case elongation, of incomplete low recrystallizationn value, and low of r ferrite, the sample will have high strength, high yield ratio, low elongation, low n value, and low r
Crystals 2020, 10, x FOR PEER REVIEW 11 of 16 Crystals 2020, 10, 777x FOR PEER REVIEW 11 of 16 value, as shown in Figures 6a,b and 7. When the recrystallization and grain growth of ferrite are fully value, as shown in Figures 6a,b and 7. When the recrystallization and grain growth of ferrite are fully developed, the sample will have low strength, low yield ratio, high elongation, high n value, and developed,r value, as shown the sample in Figure will6 a,bhave and low Figure strength,7. When low the yield recrystallization ratio, high elongation, and grain high growth n value, of ferrite and high r value, as Figures 6e and 7 illustrate. Figure 10 shows the distribution of recrystallized grains, highare fully r value, developed, as Figures the sample6e and 7 will illustrate. have low Figure strength, 10 shows lowyield the distribution ratio, high elongation,of recrystallized high n grains,value, sub-structured grains, and deformed grains at higher annealing temperatures. When annealed at 850, sub-structuredand high r value, grains, as Figures and deformed6e and7 grainsillustrate. at higher Figure annealing 10 shows temperatures. the distribution When of annealed recrystallized at 850, 880, and 900 °C, the recrystallized grains are 85.92%, 93.10%, and 97.22%, respectively; while the 880,grains, and sub-structured 900 °C, the recrystallized grains, and deformed grains are grains 85.92%, at higher 93.10%, annealing and 97.22%, temperatures. respectively; When while annealed the deformed grains are 11.85%, 4.75%, and 1.50%, respectively; the sub-grains are 2.23%, 2.15%, and deformedat 850, 880, grains and 900 are◦C, 11.85%, the recrystallized 4.75%, and grains 1.50%, are respectively; 85.92%, 93.10%, the andsub-grains 97.22%, are respectively; 2.23%, 2.15%, while and the 1.28% respectively. This indicates that with an increase in the annealing temperature, the 1.28%deformed respectively. grains are 11.85%,This indicates 4.75%, and that 1.50%, with respectively; an increase the sub-grainsin the annealing are 2.23%, temperature, 2.15%, and 1.28% the recrystallization process becomes gradually complete, corresponding to the variations in mechanical recrystallizationrespectively. This process indicates becomes that with gradually an increase complete in the, corresponding annealing temperature, to the variations the recrystallization in mechanical properties as shown in Figure 7 that the samples annealed at higher temperatures show higher n propertiesprocess becomes as shown gradually in Figure complete, 7 that the corresponding samples annealed to the at variations higher temperatures in mechanical show properties higher asn values and r values. Figure 11 illustrates the distributions of grain boundary misorientation angle for valuesshown and in Figure r values.7 that Figure the samples 11 illustrates annealed the at distribution higher temperaturess of grain boundary show higher misorientationn values and angler values. for different annealing temperatures. One may see that the percentage of low-angle grain boundaries differentFigure 11 annealingillustrates temperatures. the distributions One of may grain see boundary that the misorientation percentage of anglelow-angle for di grainfferent boundaries annealing (LAGBs, 2–15°) decreases and that of high-angle grain boundaries (HAGBs, >15°) increases. There is (LAGBs,temperatures. 2–15°) One decreases may see and that that the percentageof high-angle of low-anglegrain boundaries grain boundaries (HAGBs, (LAGBs,>15°) increases. 2–15◦) decreases There is a certain amount of recovery structure in the sample annealed at a lower temperature, while the aand certain that ofamount high-angle of recovery grain boundaries structure (HAGBs, in the sample>15◦) increases. annealed Thereat a lower is a certain temperature, amount of while recovery the recrystallization process is fully completed at a higher temperature, which further verifies the recrystallizationstructure in the sampleprocess annealed is fully atcompleted a lower temperature, at a higher whiletemperature, the recrystallization which further process verifies is fully the previous discussion. previouscompleted discussion. at a higher temperature, which further verifies the previous discussion.
Figure 10. Electron backscatter diffraction (EBSD) images showing recrystallization fraction of the Figure 10. 10. ElectronElectron backscatter backscatter diffraction diffraction (EBSD) (EBSD) images images showing showing recrystallization recrystallization fraction fraction of the of heat-treated sample versus the inter-critical temperature. (a) 850, (b) 880, and (c) 900 °C. Blue areas heat-treatedthe heat-treated sample sample versus versus the inter-critical the inter-critical temperature. temperature. (a) 850, ( (ba) 880, 850, and (b) ( 880,c) 900 and °C. (Bluec) 900 areasC. represent recrystallized grains, yellow areas are sub-structures, and red areas represent deformed◦ representBlue areas recrystallized represent recrystallized grains, yellow grains, areas yellow are sub-structures, areas are sub-structures, and red areas and represent red areas deformed represent structures. structures.deformed structures.
Figure 11. 11. GrainGrain boundary boundary misorientation misorientation distribution distribution histograms histograms of samples of samples annealed annealed at different at Figure 11. Grain boundary misorientation distribution histograms of samples annealed at different ditemperatures.fferent temperatures. (a) 850, (b (a) )880, 850, and (b) ( 880,c) 900 and °C. (c ) 900 ◦C. temperatures. (a) 850, (b) 880, and (c) 900 °C. 4.3. Texture Evolution 4.3. Texture Evolution 4.3. Texture Evolution Texture is an important parameter of steel sheet because it induces plastic anisotropy and Texture is an important parameter of steel sheet because it induces plastic anisotropy and appropriateTexture textureis an important can be beneficial parameter to formability of steel sheet [57– 60because]. A convenient it induces measure plastic ofanisotropy anisotropy and is appropriate texture can be beneficial to formability [57–60]. A convenient measure of anisotropy is appropriatethe r value, whichtexture is can the be ratio beneficial of width to strain formability to thickness [57–60]. strain, A convenient determined measure by a simple of anisotropy tensile test. is the r value, which is the ratio of width strain to thickness strain, determined by a simple tensile test. theGenerally, r value,a which higher isr thevalue ratio corresponds of width strain to a to higher thickness deep strain, drawability determined [60]. Aby favorablea simple tensile texture test. for Generally, a higher r value corresponds to a higher deep drawability [60]. A favorable texture for Generally,formability a is higher one in r which value a corresponds high proportion to a of higher the grains deep are drawability oriented with [60]. their A favorable {111} planes texture parallel for formability is one in which a high proportion of the grains are oriented with their {111} planes parallel formability is one in which a high proportion of the grains are oriented with their {111} planes parallel
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Crystals 2020, 10, x FOR PEER REVIEW 12 of 16 to the sheet plane [61]. Other texture components are relatively unfavorable. Thus, to help the steel gainto the a sheet higher plane deep [61]. drawing Other capability, texture components one needs to ar develope relatively the {111} unfavorable. texture components Thus, to help as strongthe steel as possiblegain a higher and eliminatedeep drawing or weaken capability, other one texture needs components to develop to the amaximum {111} texture degree. components as strong as possibleTexture and can eliminate be induced or weaken by processes other texture involving components plastic deformation, to a maximum e.g., degree. rolling. In general, slip takesTexture place can onbe induced {110}, {112}, by processes and {123} involving planes forplastic BCC deformation, structure, withe.g., rolling. the first In two general, as more slip predominanttakes place on [ {110},62–64 ].{112}, During and rolling{123} planes and deformation, for BCC structure, the crystalline with the orientationfirst two as more aggregates predominant around {001}[62–64].<110 During>, then rolling spreads and to deformation, {112}<110> and the {111}crystalline<110> orientationalong α fiber aggregates [62,65]. Thearound internal {001}<110>, stored energythen spreads with dislocations to {112}<110> presented and {111}<110> in the deformed along α grains fiber of[62,65]. a cold-rolled The internal steel providesstored energy the driving with forcedislocations for recrystallization presented in the at temperatures deformed grains above of thea cold-rolled recrystallization steel provides temperature. the driving It is commonly force for believedrecrystallization that the orientationat temperatures dependence above ofthe a deformedrecrystallization structure temperature. is responsible It is for commonly the development believed of certainthat the texture orientation components dependence during of thea deformed recrystallization structure process is responsible [57]. The for stored the development energy of deformation of certain structurestexture components increases induring the sequence the recrystallization {001} < {112} Figure 12. Figure 12. Schematic sequence of nucleation rates of recrystallized grains during annealing of a cold rolled steel [57]. rolled steel [57]. Based on the above discussion, the production of α fiber components in Figure8b is likely related Based on the above discussion, the production of α fiber components in Figure 8b is likely related to the shear bands in the microstructure, as shown in Figure6c. However, the formation of α fiber to the shear bands in the microstructure, as shown in Figure 6c. However, the formation of α fiber components in Figure8c is related to the dislocations produced during the formation of martensite. components in Figure 8c is related to the dislocations produced during the formation of martensite. For BCC crystals, the slip usually occurs on {110} planes at medium temperatures. Since the martensitic For BCC crystals, the slip usually occurs on {110} planes at medium temperatures. Since the transformation temperature of the steel under study is close to 500 C, the local deformation from the martensitic transformation temperature of the steel under study◦ is close to 500 °C, the local transformation-induced volumetric dilation would lead to the increase in the {110} texture density. deformation from the transformation-induced volumetric dilation would lead to the increase in the The reason for the weaker γ fiber components shown in Figure8d should be related to the two {110} texture density. The reason for the weaker γ fiber components shown in Figure 8d should be transformations, α to γ and γ to α. The ratio of {111} component to other orientation components related to the two transformations, α to γ and γ to α. The ratio of {111} component to other orientation determines the r value of the material. The higher the ratio, the higher the r value. The results of components determines the r value of the material. The higher the ratio, the higher the r value. The orientation analysis shown in Figure8 correspond to the variations in r value illustrated in Figure7d. results of orientation analysis shown in Figure 8 correspond to the variations in r value illustrated in 5.Figure Conclusions 7d. 5. ConclusionsEffects of heat treatment on microstructure, texture, and mechanical properties of a designed Fe-Si-Cr-Mo-C deep drawing dual-phase steel were investigated. The following conclusions are drawn fromEffects the study. of heat treatment on microstructure, texture, and mechanical properties of a designed Fe- Si-Cr-Mo-C deep drawing dual-phase steel were investigated. The following conclusions are drawn from the study. (1) At high annealing temperatures, ferrite and island martensite are formed when the cooling rate is slow. With increasing the cooling rate, pearlite and bainite are gradually formed while the Crystals 2020, 10, 777 13 of 16 (1) At high annealing temperatures, ferrite and island martensite are formed when the cooling rate is slow. With increasing the cooling rate, pearlite and bainite are gradually formed while the island martensite gradually disappears. As the cooling rate reaches a certain level, i.e., 15 ◦C/s, the ferrite also disappears, and the microstructure is basically dominated by lower bainite. (2) When annealed at temperatures between 780 and 850 ◦C, recrystallization of ferrite cannot be fully implemented, and shear bands exist in the microstructure. The lower the annealing temperature, the more shear bands in the microstructure. However, when annealed at temperatures between 880 and 900 ◦C, the recrystallization process is basically completed and there are only a few non-recrystallized grains in the microstructure. When annealed at 820 ◦C, a small amount of fine island martensite begins to form; with increasing the annealing temperature, the content of island martensite increases; when annealed at 880 ◦C, the content of island martensite reaches the summit and then decreases when annealed at 900 ◦C. (3) With increasing the annealing temperature, the yield strength of the steel decreases. However, the tensile strength first decreases and then increases, leaving a trough at 850 ◦C. The elongation increases with increasing the annealing temperature, showing a trough at 850 ◦C and reaches a maximum of 30.1% at 900 ◦C; but the yield ratio decreases against the annealing temperature. The values of n and r increase with increasing the annealing temperature, and the range of n value is 0.12–0.20. When annealed above 850 ◦C, the r value is greater than 1.0, and reaches the maximum of 1.15 at 900 ◦C. (4) When annealed at 850–900 ◦C, the density of α-oriented components increases from 4.0 to 5.0 mrd first, then decreases to 1.0 mrd. The density of γ-oriented components decreases from 3.0 to 2.0 mrd. The corresponding ratio of γ-oriented components to α-oriented components decreases first and then increases. Within the current temperature range for annealing, the recrystallization grain fraction of the steel increases from 85.92% to 97.22%, while the proportion of high-angle grain boundaries increases significantly. Author Contributions: Conceptualization, H.P., X.S., Y.L. and H.W.; Data curation, H.P., X.S., J.C., Y.T., H.W. and Y.X.; Formal analysis, H.P., X.S., J.C. and Y.T.; Fuding acquisition, H.P. and Y.T.; Investigation, H.P. and J.C.; Methodology, H.P., H.W. and Z.W.; Project administration, X.S. and D.L.; Resources, Y.L., H.Z., Z.W. and Y.X.; Software, Z.W.; Supervision, D.L.; Validation, X.S., D.L. and H.Z.; Visualization, H.P. and D.L.; Writing-original draft, H.P.; Writing-review & editing, D.L. All authors have read and agreed to the published version of the manuscript. Funding: This research was funded by the National Natural Science Foundation of China, grant number 51774006 and NSFC-Iron and Steel Joint Fund of the National Natural Science Foundation of China, grant number U1860105; the major special projects of Anhui Province, grant number 18030901085 and the Liaoning Provincial Natural Science Foundation of China, grant number 2019-KF-25-01. 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