Design of an Effective Heat Treatment Involving Intercritical Hardening for High Strength/High Elongation of 0.2C–3Al

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Article Design of an Eﬀective Heat Treatment Involving Intercritical Hardening for High Strength/High Elongation of 0.2C–3Al–(6–8.5)Mn–Fe TRIP Steels: Microstructural Evolution and Deformation Behavior

1, 1, , 1 2 1 Yanjie Mou †, Zhichao Li * †, Xiaoteng Zhang , Devesh Misra , Lianfang He and Huiping Li 1,* 1 School of Materials Science and Engineering, Shandong University of Science and Technology, Qingdao 266590, China; [email protected] (Y.M.); [email protected] (X.Z.); [email protected] (L.H.) 2 Laboratory for Excellence in Advanced Steel Research, Department of Metallurgical, Materials and Biomedical Engineering, University of Texas at El Paso, El Paso, TX 79968, USA; [email protected] * Correspondence: [email protected] (Z.L.); [email protected] (H.L.); Tel.: +86-0532-8068-1145 (Z.L.); +86-0532-8605-7920 (H.L.) These two authors contributed equally to this work. † Received: 19 October 2019; Accepted: 26 November 2019; Published: 28 November 2019

Abstract: High strength/high elongation continues to be the primary challenge and focus for medium-Mn steels. It is elucidated herein via critical experimental analysis that the cumulative contribution of transformation-induced plasticity (TRIP) and microstructural constituents governs high strength/high elongation in 0.2C–3Al–(6–8.5)Mn–Fe steels. This was enabled by an effective heat treatment involving a combination of intercritical hardening and tempering to obtain high strength/high ductility. An excellent combination of high ultimate tensile strength of 935–1112 MPa and total elongation of 35–40% was obtained when the steels were subjected to intercritical hardening in the temperature range of 700–750 ◦C and low tempering at 200 ◦C. The intercritical hardening impacted the coexistence of austenite, ferrite, and martensite, such that the deformation behavior varied with the Mn content. The excellent obtained properties of the steels are attributed to the cumulative contribution of the enhanced TRIP effect of austenite and the microstructural constituents, ferrite and martensite. The discontinuous TRIP effect during deformation involved stress relaxation, which was responsible for the high ductility. Lamellar austenite, unlike the equiaxed microstructure, is envisaged to induce stress relaxation during martensitic transformation, resulting in the discontinuous TRIP effect.

Keywords: microstructure; deformation; mechanical properties; heat treatment; deformation behavior; discontinuous transformation during deformation

1. Introduction Transformation-induced plasticity (TRIP) steels have good crashworthiness, superior ductility, and high strength. The steels are potential candidates for automotive applications. The enhanced work hardening rate and energy absorption behavior are attributed to the TRIP eﬀect, which is the transformation of retained austenite to martensite [1,2]. The conventional low-alloyed TRIP steels with Mn content of less than ~2.5% and a product of ultimate tensile strength and total elongation (PSE) of ~15 10 GPa% consist of ferrite, retained austenite, a small amount of martensite, and occasionally ± bainite [3–6]. The high-manganese twinning-induced plasticity (TWIP) steels with PSE of ~60 ± 10 GPa% are generally completely austenite [7,8].

Metals 2019, 9, 1275; doi:10.3390/met9121275 www.mdpi.com/journal/metals Metals 2019, 9, 1275 2 of 13

Recently, a number of studies on medium-manganese (4–10%) steels have been conducted [9–16], The medium-Mn steels with PSE 30 GPa% are considered promising automobile steels. The ≥ composition of austenite greatly influences the austenite stability. It has been suggested [12,13] that austenite stability can be increased by the distribution of manganese between the austenite and ferrite during annealing in the intercritical region with sufficient holding time for the microstructure to reach phase equilibrium. Miller [14] reported that the austenite fraction in 6.0% Mn steels could be as high as 30%. More recently, retained austenite fractions of ~16.8% in 5.2Mn–0.10C steel, ~27.8% in 5.8Mn–0.10C steel, and ~38% in 7.1Mn–0.10C steel were obtained by Merwin [11]. The volume fraction of retained austenite was significantly increased with increasing Mn content, which increases with PSE. In conventional TRIP steel, a two-stage heat treatment process is adopted to stabilize austenite during austempering [17]. As regards medium-Mn steels, the steel is first heated in the austenitic region and then quenched to obtain a complete martensite microstructure. Subsequently, annealing followed by air cooling stabilizes a high retained austenite fraction. The quenching and partitioning (Q&P) process basically consists of complete austenization followed by rapid cooling to the region between Ms and Mf. Annealing at the quenching temperature or above is subsequently conducted to facilitate the partitioning of carbon from martensite to austenite [18–20]. The process of obtaining austenite from the initial martensite is referred to as “austenite reverted transformation” (ART) [10,14,18–22]. The heat treatments discussed above require cold rolling and a long annealing time (6–12 h). A long annealing time renders the austenite too stable and leads to the TRIP effect being significantly weakened. Herein, for medium-manganese steel, two different traditional heat treatments (Q&P and ART processes) are described. The shortcomings of the two traditional heat treatments are summarized. Our purpose in this is to explain the necessity for putting forward a new heat treatment. Thus, a novel heat treatment is proposed for the experimental medium-Mn steels to obtain high strength/high elongation, constituting the objective of the study.

2. Materials and Experimental Procedure E. De Moor et al. [23] proposed a model which is used to design the composition of steels and predict the heat treatment temperature. The model suggests that the amount of austenite stabilized at room temperature depends on the content of carbon, aluminum, and manganese alloying elements enriched in austenite [15,16,19]. It further suggests that a temperature exists in the intercritical region that yields the maximum fraction of retained austenite at room temperature. In Thermo-Calc software, we input the composition of steel and set the temperature range. The prediction of the phase volume fraction of the experimental steels and composition in austenite could then be calculated. The cast ingots were refined by Central Iron and Steel Research Institute in Beijing using a vacuum melting furnace. The real composition of the experimental steels was measured by atomic absorption spectrophotometry (SP-3803, Spectrum, Shanghai, China) and is given in Table1. Then, the ingots were heated at 1250 ◦C for 3 h and hot forged between 1200 ◦C and 900 ◦C into billets of section size 100 mm 30 mm. Subsequently, the billets were cooled in air to room temperature (RT). Finally, the × billets were heated at 1250 ◦C for 5 h and hot rolled to sheets with 4 mm thickness after 7 passes of hot rolling; the primary rolling temperature was 1150 ◦C, the final rolling temperature was not lower than 850 ◦C, and the hot-rolled steel was finally cooled to room temperature at a cooling rate of 15 ◦C/s.

Table 1. Chemical composition (wt.%) of the two steels.

Steel Mn Al C Fe 6Mn 6.00 3.20 0.20 90.60 8.5Mn 8.40 3.11 0.19 88.30

The traditional heat treatment for medium-Mn steels mentioned above requires a long holding time in the intercritical region, and it was proved to be inapplicable to the experimental steels [18–22]. Also, the long annealing time renders the retained austenite too stable. As a result, the TRIP effect is Metals 2019, 9, 1275 3 of 13 weakened.Metals Thus, 2019, 9 an, x FOR effective PEER REVIEW and simple two-stage heat treatment process was adopted (as3 of 13 shown in Figure1): (1) Intercritical hardening—the as-hot-rolled steels were heated in the intercritical is weakened. Thus, an effective and simple two-stage heat treatment process was adopted (as shown temperaturein Figure region 1): 600–850 (1) Intercritical◦C for 1hardening—the h and then immediately as-hot-rolled quenchedsteels were in heated water in to the room intercritical temperature. (2) Tempering—thetemperature region quenched 600–850 steels °C for were 1 h and tempered then immediately at 200 quenched◦C for 15 in min, water followedto room temperature. by air cooling to room temperature(2) Tempering—the to obtain quenched a good steels balance were tempered between at 200 ductility °C for 15 and min, strength. followed by In air the cooling processes to of annealingroom and temperature tempering, to carbon obtain dia ffgooduses balance from ferritebetween to ductility austenite, and whichstrength. enhances In the processes the stability of of austenite,annealing leading toand superior tempering, ductility carbon [diffuses6]. from ferrite to austenite, which enhances the stability of austenite, leading to superior ductility [6].

Figure 1. FigureA heat 1. A heat treatment treatment schedule schedule for for steels steels using using an effective an e ffandective simple and two-stage simple heat two-stage treatment heat process. treatment process. Tensile specimens of a gauge width 12.5 mm and length 50 mm were taken parallel to the rolling Tensiledirection specimens and ofwere a gauge machined width from 12.5 mmthe heat-t and lengthreated 50plate. mm A were universal taken paralleltesting machine to the rolling direction and(SANSCMT5000, were machined MTS, fromSaint thePaul, heat-treated MN, USA) wa plate.s used A at universal a constant testingcrosshead machine speed of (SANSCMT5000, 3 mm·min−1 MTS, Saintfor Paul, tensile MN, testing USA) at room was temperature. used at a constant The micros crossheadtructure was speed studied of3 using mm anmin optical1 for microscope tensile testing · − at room temperature.(OM, OLYMPUS, The OLYMPUS-GSX500, microstructure was Tokyo, studied Japan), using a field-emission an optical scanning microscope electron (OM, microscope OLYMPUS, OLYMPUS-GSX500,(FE-SEM, SSX-550, Tokyo, Supra, Japan), Shimadzu, a field-emission Tokyo, Japan) scanning, and a electron field-emission microscope transmission (FE-SEM, electron SSX-550, microscope (FE-TEM, TECNAI G2-20, FEI, operated at 200 kV, Hillsboro, MO, USA). The OM and Supra, Shimadzu, Tokyo, Japan), and a field-emission transmission electron microscope (FE-TEM, SEM samples were etched in an aqueous solution of 75% H2O + 25% NaHSO3. The TEM samples were TECNAI G2-20,twin-jet FEI,polished operated (Struers, at 200Tenupol-5, kV, Hillsboro, Copenhagen, MO, Denmark) USA). The in OMa solution and SEMof 95% samples alcohol and were 5% etched in an aqueousperchloric solution acid at of ~20 75% °C. H Using2O + 25%a direct NaHSO comparison3. The method TEM samples[24], the austenite were twin-jet volume polished fraction was (Struers, Tenupol-5,determined Copenhagen, by X-ray Denmark) diffraction in (XRD, a solution D/Max of2250/PC, 95% alcohol Rigaku, and Tokyo, 5% perchloricJapan) with acidCu K atα radiation. ~20 ◦C. Using a direct comparisonThe volume fraction method of [austenite24], the was austenite measured volume using the fraction integrated was intensities determined of the by(311) X-rayγ and di(220)ffractionγ (XRD, D/Max2250peaks of austenite/PC, Rigaku, and the Tokyo, (211)α and Japan) (200) withα peaks Cu of K ferrite.α radiation. The austenite The volume volume fractionfraction, V ofA, austenitewas calculated using the equation [25] was measured using the integrated intensities of the (311)γ and (220)γ peaks of austenite and the (211)α and (200)α peaks of ferrite. The austenite volume1 fraction, , VA, was calculated using the equation [25] ∑ , (1) Iγ,i 1 1,PN 1 , ∑N i=1 R∑ , γ,i , VA = (1) 1 PN Iγ,i 1 PM Iα,i N i=1 R + M i=1 R where Iγ and Iα are the integrated intensity valuesγ,i of the austeniteα,i and α-phase, respectively; N and M are the numbers of peaks of austenite and peaks of ferrite, respectively; and Rγ and Rα are the where I and I are the integrated intensity values of the austenite and α-phase, respectively; N and γ standardizationα constant of austenite and ferrite, respectively, such that Rγ, (311) = 2.282 m−1, Rγ, (220) = M are the1.796 numbers m−1, Rα,of (211)peaks = 2.932 ofm−1 austenite, and Rα, (200)and = 1.269 peaks m−1. of ferrite, respectively; and Rγ and Rα are the 1 standardization constant of austenite and ferrite, respectively, such that Rγ, (311) = 2.282 m− , Rγ, (220) = 13. Results and Discussion1 1 1.796 m− , Rα, (211) = 2.932 m− , and Rα, (200) = 1.269 m− . 3.1. Design of the Chemical Composition of Steels 3. Results and Discussion

3.1. Design of the Chemical Composition of Steels The fraction of retained austenite as a function of the annealing temperature was predicted using the model proposed by De Moor et al. The modeling steps of the calculation process are as follows [23]: Metals 2019, 9, x FOR PEER REVIEW 4 of 13

The fraction of retained austenite as a function of the annealing temperature was predicted using the model proposed by De Moor et al. The modeling steps of the calculation process are as follows Metals 2019, 9, 1275 4 of 13 [23]: (1) The phase fraction of each temperature in the equilibrium state is calculated using Thermo-Calc. (1) TheShown phase in Figure fraction 2a,d, of each the three temperature curves are in thethe equilibriumtrend charts stateof temperature is calculated change using of Thermo-Calc. austenite, Shownferrite, inand Figure cementite.2a,d, the three curves are the trend charts of temperature change of austenite,

(2) ferrite,The contents and cementite. of Mn, C, and Al elements in the austenite at different temperatures are calculated by Thermo-Calc, as shown in Figure 2b,e. (2) The contents of Mn, C, and Al elements in the austenite at diﬀerent temperatures are calculated (3) According to the calculated content of each element in the austenite, the Ms point of the by Thermo-Calc, as shown in Figure2b,e. experimental steel can be calculated using Equation (2) [26]: (3) According to the calculated content of each element in the austenite, the Ms point of the experimental steel can be calculatedMs = 550 using − 350 EquationC − 40Mn (2) + 30 [26Al.]: (2) (4) According to the Koistinen–Marburger (K–M) equation, the amount of transformation from Ms = 550 350C 40Mn + 30Al. (2) austenite to martensite during the cooling −process− is calculated using Equation (3): (4) According to the Koistinen–Marburgerfm = 1 − exp[ (K–M)−0.11 equation, (Ms − T)] the amount of transformation(3) from whereaustenite T is the to room martensite temperature, during which the cooling is taken process as 20 °C. is calculated using Equation (3): According to the initial austenite fraction (as shown in Figure 2a,d) and the newly formed martensite during the cooling processf mcalculated= 1 exp[ by 0.11Equations (Ms (2)T)] and (3), the fraction of retained (3) − − − austenite at room temperature can be obtained, as shown in Figure 2c,f. For 8.5Mn steel, a pronounced peakwhere (~65.2 T isvol.%) the room was observed temperature, at ~755 which °C, isand taken for 6Mn as 20 steel,◦C. the maximum fraction of retained austenite reached ~37.7 vol.% at about 690 °C.

FigureFigure 2. 2.Schematic Schematic illustrations illustrations ofof thethe predictive model model for for austenite austenite stabilization stabilization as asa function a function of of temperaturetemperature for for 6Mn 6Mn steels steels andand 8.5Mn8.5Mn steels:steels: ( (a,da,d)) phase phase fractions; fractions; (b,e (b,e) C,) C, Al, Al, and and Mn Mn content content in in austenite;austenite; and and (c,f (c,f) calculated) calculated volume volume fractionfraction of retained austenite. austenite.

3.2.According Microstructure to theEvolution initial and austenite Mechanical fraction Properties (as shown in Figure2a,d) and the newly formed martensiteSEM duringmicrographs the cooling of the 8.5Mn process and calculated 6Mn intercritical by Equations hardened (2) and and tempered (3), the fraction steels are of shown retained austenitein Figure at 3. room The microstructures temperature can of be the obtained, 8.5Mn and as 6Mn shown steels in Figurecomprised2c,f. Forintercritical 8.5Mn steel, ferrite a ( pronouncedα-ferrite), peakδ-ferrite, (~65.2 retained vol.%) wasaustenite, observed and martensite, at ~755 ◦C, and and the for morphology 6Mn steel, and the maximumcontent of each fraction phase of varied retained austenitesignificantly reached with ~37.7 change vol.% in the at aboutintercritical 690 ◦ C.hardening temperature. The amount of stripelike ferrite decreased with increasing temperature, while the thin lath martensite gradually thickened with 3.2.increasing Microstructure temperature. Evolution Similar andMechanical morphological Properties changes were observed for martensite. The thin lath martensiteSEM micrographs was dispersed of the in 8.5Mnpackets and (marked 6Mn intercritical rectangle in hardened Figure 3) and in the tempered parent austenite steels are matrix, shown in Figurewhich3. was The broadly microstructures divided into of thegranular 8.5Mn and and thin 6Mn layers steels in comprisedthe temperature intercritical range of ferrite 750–800 ( α °C-ferrite), for δ -ferrite, retained austenite, and martensite, and the morphology and content of each phase varied signiﬁcantly with change in the intercritical hardening temperature. The amount of stripelike ferrite decreased with increasing temperature, while the thin lath martensite gradually thickened with increasing temperature. Similar morphological changes were observed for martensite. The thin lath martensite was dispersed in packets (marked rectangle in Figure3) in the parent austenite matrix, Metals 2019, 9, 1275 5 of 13

Metals 2019, 9, x FOR PEER REVIEW 5 of 13 which was broadly divided into granular and thin layers in the temperature range of 750–800 ◦C for 6Mn steel and at 800 ◦°CC for 8.5Mn steel, respectively.respectively. The content of retained austenite in 8.5Mn steel waswas higherhigher thanthan that that in in 6Mn 6Mn steel, steel, while while the the ferrite ferrit contente content in 8.5Mnin 8.5Mn steel steel was was less thanless than that inthat 6Mn in steel.6Mn Accordingsteel. According to the to Fe–C the phaseFe–C phase diagram, diagram, with increasing with increasing annealing annealing temperature, temperature, the amount the amount of ferrite of decreased,ferrite decreased, the amount the amount of austenite of austenite increased, increased, and the austeniteand the austen grainsite grew. grains With grew. increased With increased austenite grainaustenite size, grain the stability size, the of stabilit austenitey of decreases. austenite Duringdecreases. quenching, During morequenching, austenite more was austenite transformed was intotransformed martensite. into Thus,martensite. with increasingThus, with annealing increasing temperature, annealing temperature, the amount the of amount martensite of martensite increased afterincreased quenching. after quenching. Austenite Austenite in the two in steels the two comprised steels comprised a high fraction a high in fraction the temperature in the temperature ranges of 650–700ranges of◦ C650–700 for 6Mn °C steel for 6Mn and steel 650–750 and◦ 650–750C for 8.5Mn °C fo steel,r 8.5Mn but steel, when but they when were they treated were attreated 750 ◦ Cat and750 800°C and◦C, respectively,800 °C, respectively, the amount the of amount austenite of decreasedaustenite markedlydecreased becausemarkedly of extensivebecause of martensitic extensive transformation.martensitic transformation. Moreover, we performedMoreover, nanoindentation we performed (Keysight nanoindentation G200XT nanoindenter (Keysight system)G200XT of thenanoindenter 0.20C–8.65Mn–4.12Al–Fe system) of the medium-manganese0.20C–8.65Mn–4.12Al–Fe steels medium-manganese after quenching from steels 750 ◦C after and quenching tempering fromfrom 200750 °C◦C and (as showntempering in Figure from 2004), which °C (as couldshown also in Figure explain 4),the which hardness could also di ﬀ erencesexplain the between hardness the austenite,differencesδ -ferrite,betweenα the-ferrite, austenite, and martensite. δ-ferrite, α-ferrite, and martensite.

Figure 3.3. SEMSEM micrographs micrographs of of hot-rolled 6Mn and 8.5M 8.5Mnn steels after quenching from didifferentﬀerent temperatures.temperatures. ( a) 6Mn, 650 ◦°C;C; ( b) 6Mn, 700 ◦°C;C; ( c) 6Mn, 750 ◦°C;C; ( d) 6Mn, 800 ◦°C;C; (e) 8.5Mn, 650 ◦°C;C; ((ff)) 8.5Mn,8.5Mn, 700700 ◦°C;C; (g)) 8.5Mn,8.5Mn, 750750 ◦°C;C; and (h)) 8.5Mn,8.5Mn, 800800 ◦°C.C.

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FigureFigure 4. (4.a) ( Aa) SEMA SEM image image showing showing the the indentation indentation impressionsimpressions on austenite austenite grains, grains, δ-ferriteδ-ferrite grains, grains, α-ferriteα-ferrite grains, grains, and and martensite martensite grains grains in in hot-rolled hot-rolled 0.20C–8.65Mn–4.12Al–Fe 0.20C–8.65Mn–4.12Al–Fe steels steels after after quenching quenching

fromfrom 750 750◦C °C and and tempering tempering from from 200 200◦ C.°C. ( b(b)) The The nanohardness nanohardness valuesvalues of of different diﬀerent phase phase grains. grains. Each Each errorerror bar bar represents represents a standard a standard deviation. deviation.

FigureFigure5 shows 5 shows the XRDthe XRD pattern pattern and and measured measured fraction fraction of austenite of austenite in heat-treated in heat-treated 8.5Mn 8.5Mn and and 6Mn steels6Mn at steels room at temperature. room temperature. For the Fo 8.5Mnr the 8.5Mn steel, steel, the retainedthe retained austenite austenite volume volume fraction fraction increased increased to 58.5to vol.% 58.5 vol.% at 750 at◦C, 750 followed °C, followed by a significant by a significant decrease decrease to 24% to at 80024%◦ atC. 800 A similar °C. A variationsimilar variation in austenite in austenite content as a function of temperature occurred for the 6Mn steel. It attained a peak of 33 content as a function of temperature occurred for the 6Mn steel. It attained a peak of 33 vol.% at 700 ◦C. Thus,vol.% Mn at promoted 700 °C. Thus, an increase Mn promoted in the an retained increase austenite in the retained volume austenite fraction. volume Comparing fraction. Figure Comparing2c,f with FigureFigure5c, we2c,f note with that Figure the theoretical5c, we note predictionsthat the theo areretical in good predictions agreement are within good the experimentalagreement with results. the experimental results. With increasing annealing temperature, the austenite volume fraction increased and the volume fraction of ferrite decreased. Thus, the austenite volume fraction increased with increasing temperature when 8.5Mn steel and 6Mn steel were intercritically annealed at 600–750 ◦C and 650–700 ◦C, respectively. Compared with the 6Mn steel, the 8.5Mn steel had a larger austenite content. With increasing temperature, the grain size of the austenite increased gradually, which resulted in a decrease in the austenite stability [27,28]. The higher the annealing temperature, the more austenite transformed to martensite during quenching when 6Mn steel was intercritically annealed in the temperature range of 750–850 ◦C. Because of martensitic transformation, the volume fraction of austenite decreased for 8.5Mn steel intercritically annealed at 800 ◦C. The mechanical properties are presented in Figure6. At least five samples for each temperature were examined in the experiment. The standard deviation is also presented in the results in Figure6. Figures6a and5b show that ultimate tensile strength (UTS) increased continuously with increasing temperature, whereas the total elongation (TEL) almost decreased with increasing temperature after attaining peak values of 35.6% for 6Mn steel at 700 ◦C (6Mn-700 steel) and 39.5% for 8.5Mn steel at 750 ◦C (8.5Mn-750 steel). From the plots of TEL, UTS, and product of UTS and TEL in Figure6, it can be seen that the mechanical properties of the 8.5Mn steel were obviously superior to those of the 6Mn steel, which mainly resulted from a greater austenite amount in the 8.5Mn steel, as apparent from Figure5c. A comparison between Figure6a,b and Figure5c show that the total elongation had a trend similar to the austenite volume fraction. As shown in Figure6, the 6Mn-700 steel (when 6Mn steel was heat treated at 700 ◦C) exhibited a TEL of 35.2%, UTS of 935 MPa, and PSE of 32.9 GPa%, while 8.5Mn-750 steel (when 8.5Mn steel was heat treated at 750 ◦C) was characterized by a TEL of 39.4%, UTS of 1113 MPa, and PSE of 43.9 GPa%. In contrast, as shown in Table2[ 9,11,13,29], other medium-Mn TRIP steels had lower or similar ductility with similar strength but required additional cold rolling work and a prolonged annealing time. Therefore, 8.5Mn and 6Mn steels heat treated at Figure 5. XRD patterns and measured austenite fractions of 6Mn and 8.5Mn steels heat treated at 750 ◦C and 700 ◦C, respectively, and tempered at 200 ◦C are significantly important. In particular, the different temperatures: (a) XRD patterns for 6Mn steels; (b) XRD patterns for 8.5Mn steels; and (c) two samples had a superior product of ultimate tensile strength and total elongation. measured austenite fractions.

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Figure 4. (a) A SEM image showing the indentation impressions on austenite grains, δ-ferrite grains, α-ferrite grains, and martensite grains in hot-rolled 0.20C–8.65Mn–4.12Al–Fe steels after quenching from 750 °C and tempering from 200 °C. (b) The nanohardness values of different phase grains. Each error bar represents a standard deviation.

Figure 5 shows the XRD pattern and measured fraction of austenite in heat-treated 8.5Mn and 6Mn steels at room temperature. For the 8.5Mn steel, the retained austenite volume fraction increased to 58.5 vol.% at 750 °C, followed by a significant decrease to 24% at 800 °C. A similar variation in austenite content as a function of temperature occurred for the 6Mn steel. It attained a peak of 33 vol.% at 700 °C. Thus, Mn promoted an increase in the retained austenite volume fraction. Comparing FigureMetals 2c,f2019 with, 9, 1275 Figure 5c, we note that the theoretical predictions are in good agreement7 ofwith 13 the experimental results.

Figure 5. XRD patterns and measured austenite fractions of 6Mn and 8.5Mn steels heat treated at Figure 5. XRD patterns and measured austenite fractions of 6Mn and 8.5Mn steels heat treated at diﬀerent temperatures: (a) XRD patterns for 6Mn steels; (b) XRD patterns for 8.5Mn steels; and different temperatures: (a) XRD patterns for 6Mn steels; (b) XRD patterns for 8.5Mn steels; and (c) (c) measured austenite fractions. measured austenite fractions. Table 2. Comparisons of medium-Mn transformation-induced plasticity (TRIP) steels.

Initial Composition UTS (MPa) TEL (%) PSE (GPa%) Reference Condition 0.13C–5Mn–0.6Si–1.1Al CR/annealing 988/1461 16.7/8.6 16.5/12.6 9 0.13C–6Mn–0.5Si–3.1Al CR/annealing 854/1161 21.7/12.4 18.5/14.4 11 0.2C–5Mn HR/ART (12 h) 890 33 29.4 13 0.15C–4Mn CR/annealing 1084/1187 19.1/15.9 20.7/18.9 29 0.2C–6Mn–3Al HR/QT (1 h) 935 35.2 32.9 present work 0.2C–8.5Mn–3Al HR/QT (1 h) 1113 39.4 43.9 present work UTS, ultimate tensile strength; TEL, total elongation; PSE, product of strength and elongation; HR, hot rolled; CR, cold rolled; QT, quenching and tempering; ART, austenite reverted transformation.

FigureFigure 6. 6. TensileTensile properties properties of of 6Mn 6Mn steel steel and and 8.5Mn 8.5Mn steel: steel: (a ()a )ultimate ultimate tensile tensile strength strength and and total total elongationelongation of of 6Mn steel;steel; ( b(b)) ultimate ultimate tensile tensile strength strength and and total tota elongationl elongation of 8.5Mn of 8.5Mn steel; andsteel; (c )and product (c) Metals 2019, 9, x FOR PEER REVIEW 9 of 13 productof ultimate of ultimate tensile strengthtensile strength and total and elongation total elongation comparisons comparisons between between 6Mn steel 6Mn and steel 8.5Mn and steel. 8.5Mn steel.

3.3. Deformation Behavior A comparison of the XRD patterns of undeformed 8.5Mn-750 and 6Mn-700 steels with fractured samples is shown in Figure 7. The results showed that the amount of austenite decreased to a small amount after tensile failure. The retained austenite volume fractions measured by XRD analysis of the 8.5Mn-750 and 6Mn-700 fractured samples were ~9.8 and ~7.7 vol.%, respectively, and the original austenite volume fractions of the 8.5Mn-750 and 6Mn-700 samples were 58.5 and 33 vol.%, respectively. Thus, it is indicated that a significant TRIP effect occurred during tensile deformation. Figure 8 shows the microstructures of the 6Mn-700 and 8.5Mn-750 steels in the vicinity of the fracture position after tensile deformation. Based on Figures 7 and 8, there was a marked decrease in austenite, while the variation in the martensite fraction was the opposite. Black and white lath martensite subdivided the blocky austenite into thin film and granules. Comparing the micrographs of the undeformed steels with the fractured surface, it can be inferred that the black lath martensite FigureFigure 7. 7.XRDXRD patterns patterns of thethe fracturedfractured and and undeformed undeformed steels: steels: (a) 6Mn-700 (a) 6Mn-700 steel andsteel (b and) 8.5Mn-750 (b) 8.5Mn-750 steel. is thesteel. martensite that was formed during heat treatment, while the white lath martensite with higher carbonFigure content8 shows corresponds the microstructures to strain-induced of the 6Mn-700deformation. and 8.5Mn-750During quen steelsching, in the the vicinity martensitic of the transformationfracture position occurs after preferentially tensile deformation. in low-carb Basedon-containing on Figures 7austenite, and8, there and was higher-carbon-content a marked decrease austenitein austenite, is retained. while the The variation reason infor the this martensite is that carbon fraction can was effectively the opposite. improve Black the and stability white lathof austenitemartensite [9]. subdivided Thus, the strain the blocky-induced austenite martensite into thin is the ﬁlm resi anddual granules. white lathy Comparing martensite. the micrographs of the undeformed steels with the fractured surface, it can be inferred that the black lath martensite is the martensite that was formed during heat treatment, while the white lath martensite with higher carbon content corresponds to strain-induced deformation. During quenching, the martensitic transformation occurs preferentially in low-carbon-containing austenite, and higher-carbon-content austenite is

Figure 8. Optical micrographs of the 6Mn-700 steel (a) and 8.5Mn-750 steel (b) after tensile deformation.

According to previous studies [10,30], the work hardening rate evolution presents three domains in steels containing 5–7% Mn content: (1) a rapid decrease of the work hardening rate (WH), (2) an increase of the WH, and (3) a final smooth decrease of the WH. Shi et al. [10] suggested that the three- stage WH exists only when the volume fraction of retained austenite is greater than ~15%. A majority of studies [10,30–32] have concluded that the first stage is mainly associated with ferrite deformation, the second intermediate stage is related to the occurrence of the TRIP effect, and the third stage may be related to straining of martensite and ferrite because martensite phase transition in this stage is inactive. The WH behavior for the two steels is presented in Figure 9. A similar three-stage WH was observed in 6Mn-700 steel, but in Stage III, there was a small difference: the WH decreased and there were fluctuations. Based on previous studies [33–37], this behavior is related to the discontinuous TRIP effect. As shown in Figure 9, the strain in Stage III from each peak point of the fluctuation can be divided into two stages (rise and drop of the WH curve). In the rising stage, when a certain critical stress is attained, martensitic transformation is activated, and retained austenite with similar stability continuously transforms to martensite. The WH increases rapidly in this stage because of the occurrence of the TRIP effect. As shown in Figure 10, the granular and lamellar grains located between adjacent martensite laths are austenite grains, as apparent from the diffraction pattern (inset in Figure 10b). It is known that the stability of granular and lamellar retained austenite is greater than that of equiaxed austenite [38]. Chiang et al. [39] showed that lamellar austenite exhibits high and sustained work hardening ability at high strain, while equiaxed austenite exhibits a relatively rapid transformation rate. It can be deduced that the fluctuations in Stage III are because of stress relaxation

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retained.Figure The 7. XRD reason patterns for this of isthe that fractured carbon and can undeformed eﬀectively improvesteels: (a) the 6Mn-700 stability steel of austeniteand (b) 8.5Mn-750 [9]. Thus, thesteel. strain-induced martensite is the residual white lathy martensite.

FigureFigure 8. OpticalOptical micrographs micrographs of theof 6Mn-700the 6Mn-700 steel (a )steel and 8.5Mn-750 (a) and steel8.5Mn-750 (b) after steel tensile (b deformation.) after tensile deformation. According to previous studies [10,30], the work hardening rate evolution presents three domains in steels containing 5–7% Mn content: (1) a rapid decrease of the work hardening rate (WH), (2) an According to previous studies [10,30], the work hardening rate evolution presents three domains increase of the WH, and (3) a final smooth decrease of the WH. Shi et al. [10] suggested that the in steels containing 5–7% Mn content: (1) a rapid decrease of the work hardening rate (WH), (2) an three-stage WH exists only when the volume fraction of retained austenite is greater than ~15%. A increase of the WH, and (3) a final smooth decrease of the WH. Shi et al. [10] suggested that the three- majority of studies [10,30–32] have concluded that the first stage is mainly associated with ferrite stage WH exists only when the volume fraction of retained austenite is greater than ~15%. A majority deformation, the second intermediate stage is related to the occurrence of the TRIP effect, and the third of studies [10,30–32] have concluded that the first stage is mainly associated with ferrite deformation, stage may be related to straining of martensite and ferrite because martensite phase transition in this thestage second is inactive. intermediate stage is related to the occurrence of the TRIP effect, and the third stage may be relatedThe WHto straining behavior of for martensite the two steels and ferrite is presented because in Figuremartensite9. A phase similar transition three-stage in WHthis wasstage is inactive.observed in 6Mn-700 steel, but in Stage III, there was a small difference: the WH decreased and there wereThe fluctuations. WH behavior Based for on the previous two steels studies is presente [33–37], thisd in behaviorFigure 9. is A related similar to three-stage the discontinuous WH was observedTRIP eff ect.in 6Mn-700 As shown steel, in Figurebut in 9Stage, the III, strain there in Stagewas a IIIsmall from difference: each peak the point WH of decreased the fluctuation and there werecan befluctuations. divided into Based two stageson previous (rise and studies drop of[33–37 the WH], this curve). behavior In the is rising related stage, to the when discontinuous a certain TRIPcritical effect. stress As is shown attained, in martensiticFigure 9, the transformation strain in Stage is activated,III from each and peak retained point austenite of the fluctuation with similar can bestability divided continuously into two stages transforms (rise and to drop martensite. of the WH The curve). WH increases In the rising rapidly stage, in this when stage a certain because critical of stressthe occurrence is attained, of martensitic the TRIP e fftransformationect. As shown is in activated, Figure 10, and the granularretained austenite and lamellar with grains similar located stability continuouslybetween adjacent transforms martensite to lathsmartensite. are austenite The grains,WH increases as apparent rapidly from thein dithisffraction stage patternbecause (inset of the occurrencein Figure 10 ofb). the It isTRIP known effect. that As the shown stability in of Figure granular 10, andthe lamellargranular retained and lamellar austenite grains is greater located betweenthan that adjacent of equiaxed martensite austenite laths [38 ar].e Chiang austenite et al.grains, [39] showed as apparent that lamellarfrom the austenite diffraction exhibits pattern high (inset inand Figure sustained 10b). It work is known hardening that the ability stability at high of granular strain, while and equiaxedlamellar retained austenite austenite exhibits is a relativelygreater than thatrapid of transformationequiaxed austenite rate. [38]. It can Chiang be deduced et al. that[39] theshowed fluctuations that lamellar in Stage austenite III are because exhibits of high stress and sustainedrelaxation work in the hardening martensitic ability transformation at high strain, process, while which equiaxed is due to austenite the presence exhibits of austenite a relatively types rapid of transformationdifferent morphology. rate. It can be deduced that the fluctuations in Stage III are because of stress relaxation

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Metalsin the2019 martensitic, 9, 1275 transformation process, which is due to the presence of austenite types of different10 of 13 morphology.

MetalsFigure Figure2019, 9, 9. 9.x FORWork PEER hardening REVIEW rate and stress–st stress–strainrain curves of of 6Mn-700 6Mn-700 steel steel ( (aa)) and and 8.5Mn-750 8.5Mn-750 steel steel (b (b).). 11 of 13

In the case of the 8.5Mn-750 sample, as shown in Figure 9b, the WH behavior showed only two stages, which were similar to Stages II and III of the 6Mn-700 steel. The absence of the rapidly decreasing Stage І in the 8.5Mn-750 sample may be related to two main factors: first, a small amount of soft ferrite; second, a larger amount of austenite (Figures 3g and 5c) which is largely equiaxed and rarely transformed into martensite during the commencement of plastic deformation. It is reasonable to conclude that the equiaxed austenite first transformed into martensite intensely in Stage II, and lamellar austenite was retained. A majority of the austenite retained after tensile fracture was granular. Therefore, a majority of lamellar austenite was transformed into martensite in Stage III. Thus, the fluctuation in this stage results from the discontinuous TRIP effect associated with lamellar austenite, which relieves the local stress concentration produced by martensite laths. In contrast, Stage III of the 6Mn-700 sample fluctuated to a small degree because of the large amount of ferrite effectively relieving the internal stress induced by volume expansion associated with the martensite transformation process. Thus, the cracks marked with a rectangle in Figure 8a are in δ-ferrite grains. It can be seen from the results in Figure 9 that the discontinuous TRIP effect in the 8.5Mn-750 sample was stronger than that in the 6Mn-700 sample. Moreover, the products of the ultimate tensile strength and total elongation of the 8.5Mn-750 sample and 6Mn-700 sample were 43.9 GPa% and 32.9 GPa%, respectively. Thus, the strong discontinuous TRIP effect in the 8.5Mn-750 sample contributed to its superior mechanical properties.

FigureFigure 10. 10.TEM TEM micrographsmicrographs ofof undeformedundeformed 6Mn-7006Mn-700 steelsteel andand 8.5Mn-7508.5Mn-750 steel:steel: ((aa,c,c)) bright-ﬁeldbright-field (BF)(BF) imagesimages and and ( b(b,d,d)) dark-ﬁeld dark-field (DF) (DF) images. images.

4. ConclusionsIn the case of the 8.5Mn-750 sample, as shown in Figure9b, the WH behavior showed only two stages, which were similar to Stages II and III of the 6Mn-700 steel. The absence of the rapidly decreasing(1) With Stagechange I inin thethe 8.5Mn-750intercritical sample temperature, may be th relatede maximum to two mainvolume factors: fraction ﬁrst, of aretained small amount austenite of soft ferrite;stabilized second, at room a larger temperature, amount of calculated austenite base (Figuresd on 3theg and model5c) whichproposed is largely by De equiaxedMoor [23], and was rarelyin transformed good agreement into martensite with the experimental during the commencement data. The predicted of plastic results deformation. provided It a is basis reasonable for the to concludedesign thatof a heat the equiaxed treatment austenite scheme. The ﬁrst higher transformed Mn content into martensitein 8.5Mn steel intensely was responsible in Stage II, for and its lamellarhigher austenite austenite was content. retained. A majority of the austenite retained after tensile fracture was granular.

Therefore,(2) Optimal a majority mechanical of lamellar properties austenite were wasobtained transformed in 8.5Mn into and martensite 6Mn steels in Stage when III. they Thus, were the subjected to intercritical hardening at 750 °C and 700 °C, respectively, and identical tempering at 200 °C. The 8.5Mn-750 steel was characterized by an excellent combination of a TEL of 39.4%, UTS of 1113 MPa, and UTS × TEL of 43.9 GPa%. The mechanical properties of the two experimental steels were superior to those of other medium-Mn TRIP steels, with the advantage of reduced cold rolling work and a shorter annealing time. (3) Lamellar austenite, unlike the equiaxed microstructure, is more likely to cause stress relaxation during martensitic transformation, resulting in a discontinuous TRIP effect. The superior product of ultimate tensile strength and total elongation of the 8.5Mn steel is significantly related to a strong discontinuous TRIP effect.

Funding: This research was funded by the Natural Science Foundation of Shandong Province, grant number ZR2019BEE034; Natural Science Foundation of Shandong Province, grant number 2019GGX104009; Scientific

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fluctuation in this stage results from the discontinuous TRIP effect associated with lamellar austenite, which relieves the local stress concentration produced by martensite laths. In contrast, Stage III of the 6Mn-700 sample fluctuated to a small degree because of the large amount of ferrite effectively relieving the internal stress induced by volume expansion associated with the martensite transformation process. Thus, the cracks marked with a rectangle in Figure8a are in δ-ferrite grains. It can be seen from the results in Figure9 that the discontinuous TRIP e ffect in the 8.5Mn-750 sample was stronger than that in the 6Mn-700 sample. Moreover, the products of the ultimate tensile strength and total elongation of the 8.5Mn-750 sample and 6Mn-700 sample were 43.9 GPa% and 32.9 GPa%, respectively. Thus, the strong discontinuous TRIP effect in the 8.5Mn-750 sample contributed to its superior mechanical properties.

4. Conclusions (1) With change in the intercritical temperature, the maximum volume fraction of retained austenite stabilized at room temperature, calculated based on the model proposed by De Moor [23], was in good agreement with the experimental data. The predicted results provided a basis for the design of a heat treatment scheme. The higher Mn content in 8.5Mn steel was responsible for its higher austenite content. (2) Optimal mechanical properties were obtained in 8.5Mn and 6Mn steels when they were subjected to intercritical hardening at 750 ◦C and 700 ◦C, respectively, and identical tempering at 200 ◦C. The 8.5Mn-750 steel was characterized by an excellent combination of a TEL of 39.4%, UTS of 1113 MPa, and UTS TEL of 43.9 GPa%. The mechanical properties of the two experimental × steels were superior to those of other medium-Mn TRIP steels, with the advantage of reduced cold rolling work and a shorter annealing time. (3) Lamellar austenite, unlike the equiaxed microstructure, is more likely to cause stress relaxation during martensitic transformation, resulting in a discontinuous TRIP effect. The superior product of ultimate tensile strength and total elongation of the 8.5Mn steel is significantly related to a strong discontinuous TRIP effect.

Author Contributions: Data curation, Y.M.; formal analysis, Y.M.; funding acquisition, Z.L.; investigation, Y.M.; methodology, X.Z. and L.H.; project administration, Z.L.; resources, H.L.; software, L.H. and H.L.; supervision, Z.L.; validation, Z.L.; writing—original draft, Z.L.; writing—review and editing, D.M. Funding: This research was funded by the Natural Science Foundation of Shandong Province, grant number ZR2019BEE034; Natural Science Foundation of Shandong Province, grant number 2019GGX104009; Scientific Research Foundation of Shandong University of Science and Technology for Recruited Talents, grant number 2017RCJJ017; National Natural Science Foundation of China, grant number 51575324; and Shandong Province Key Laboratory of Mine Mechanical Engineering, China, grant number 2019KLMM104. Acknowledgments: R Devesh K Misra gratefully acknowledges collaboration with the former student and now Zhichao Li. Conflicts of Interest: The authors declare no conflict of interest.

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