Aging Behavior of Intercritically Quenched Ductile Iron

metals

Article Aging Behavior of Intercritically Quenched Ductile Iron

Ali Abdelmonem 1, Mohamed Soliman 2,3,*, Heinz Palkowski 2 and Ahmed Elsabbagh 1

1 Design and Production Engineering Department, Faculty of Engineering, Ain Shams University, Cairo 11517, Egypt; [email protected] (A.A.); [email protected] (A.E.) 2 Institute of Metallurgy, Clausthal University of Technology, 38678 Clausthal-Zellerfeld, Germany; [email protected] 3 Faculty of Engineering, Galala University, Galala City 43511, Egypt * Correspondence: [email protected]

Abstract: Although extensive aging and strain aging (bake hardening, BH) studies have been carried out on dual-phase steels, the aging behavior of the dual matrix structure (DMS) ductile iron (DI), as a potential way to improve its mechanical properties, has not been addressed until now. This research was designed to study the aging behavior of DI with a ferrite-martensite matrix structure. DMS- DI with a martensite volume fraction of 30% was produced by intercritical austenitizing at 785 ◦C followed by quenching in water to room temperature. Aging treatments were carried out without pre- straining at aging temperatures of 140, 170, and 220 ◦C for 2–10,000 min. DMS-DI was investigated by light optical microscopy (LOM) for unaged samples and scanning electron microscopy (SEM) for selected samples after aging treatments. The effect of aging conditions on the mechanical properties were investigated. Microhardness measurements for ferrite and martensite were also examined as a function of aging conditions. The increase in yield strength due to aging was determined. The



 results indicate that the aging conditions have a small effect on the ultimate tensile strength UTS. It is shown that the yield strength increased to a maximum value of 45 MPa (~11% increase) after Citation: Abdelmonem, A.; Soliman, aging for particular time, which is found to be dependent on the aging temperature. The peak aging M.; Palkowski, H.; Elsabbagh, A. response is followed by a decrease in yield strength, which is observed to be attributed to martensite Aging Behavior of Intercritically tempering as confirmed by microhardness measurements. Quenched Ductile Iron. Metals 2021, 11, 897. https://doi.org/ Keywords: dual matrix structure ductile iron; aging; precipitation; microhardness. 10.3390/met11060897

Academic Editor: Anders E. W. Jarfors 1. Introduction Received: 6 May 2021 Due to the combination of good mechanical properties and low cost of ductile iron (DI), Accepted: 27 May 2021 it is widely used in many critical engineering applications [1]. The mechanical properties Published: 31 May 2021 of DI depend on both nodule characteristics and the constitution of the matrix. The former develop during the solidification stage and the latter results from the transformation of Publisher’s Note: MDPI stays neutral the austenite at high temperature [2]. The matrix structure of DI in the as-cast condition with regard to jurisdictional claims in can be controlled by a number of variables such as pouring temperature, spheroidization published maps and institutional affil- treatment, inoculation process, chemical compositions, and cooling rate [3]. Since the 1960s, iations. various methods have been carried out to improve the mechanical properties of DI. These methods relate to all known heat-treatments such as austenitizing, normalizing, annealing, quenching, and tempering [4]. Ductile iron with dual phase or dual matrix structure (DMS-DI) was first introduced in Copyright: © 2021 by the authors. 1980 [5]. From that time, the DMS-DI has gained the attention of scientists and researchers Licensee MDPI, Basel, Switzerland. for its excellent combination between strength, elongation, and its improved machinability This article is an open access article compared to other conventional DI [6,7]. The mechanical properties and physical character- distributed under the terms and istics of dual-phase DI depend mainly on the kind of microstructures (ferrite + martensite conditions of the Creative Commons or ferrite + bainite or ausferrite), which are obtained by special heat treatment cycles. Attribution (CC BY) license (https:// These involve DI in an as-cast condition being subjected to a partial austenitization step at creativecommons.org/licenses/by/ different elevated temperatures within the intercritical (α + γ + graphite) region, followed 4.0/).

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by a quenching step (a rapid cooling step) to transform γ austenite into martensite or ausferrite [8–14]. There is another schedule for obtaining DMS-DI, in which the DI in an as-cast condition is heated to above the upper critical transformation temperature (γ region) of Fe-C-Si equilibrium ternary phase diagram and then soaked for a certain time before slow cooling to the intercritical region to form a certain amount of ferrite before quenching to room temperature [6]. Although there is limited previous work on aging or precipitation hardening in the cast iron field, the work in certain nonferrous alloys and steel is more exhaustive. Hence, to illustrate the aging behavior in DMS-DI, it is necessary to review the aging behavior in these alloys. The phenomenon of aging (precipitation hardening) was first discovered in an aluminum alloy by the Alfred Wilm in 1906 [15,16]. The aging (precipitation hardening) mechanism in nonferrous alloys is described in detail in the literature [17–19]. The strain aging behavior or Bake Hardening (BH) in steel and multiphase steels (DP, TRIP, etc.) is explained in detail in the literature [20]. According to Elsen and Hougardy [21], the increase in yield strength due to BH in steels occurs in two stages. In the first stage, the increase in yield strength is caused by the formation of Cottrell atmosphere around dislocations. In the second stage, the strengthening is caused by the precipitation of fine carbides (coherent) at dislocations. The first work referring to aging behavior in cast iron (malleable cast iron) was published by G.E. Kempka in 1955 [22,23]. In 1963, R. Ebner published the first study on aging in grey cast iron. Two published studies also addressed aging behavior of cast iron, one of them on malleable iron by Burgess P.B. in 1969, and the other on grey cast iron by Novichkov P.V in 1970 [22]. From that time, no other work on this topic has been published until interest in it was revived again with a study by Gokhale, Saumitra P. in 1995 [24]. This study was performed to explain the aging characteristics of ADI. In 1996, another study was conducted to describe the precipitation hardening of ferritic ductile iron by Rezvani, M. [25]. This study showed that vanadium additions (up to 0.15%) to ferritic ductile iron increased the yield strength by up to 60% as well as hardness and the modulus of elasticity but reduced the ductility. From 1997 to 2016, the American Foundry Society (AFS) supported numerous studies on age strengthening of cast iron. All these studies are documented in the literature [26–40]. Additionally, Bayati H. and Elliott R. [41] studied the effect of aging temperature and aging time on the microstructure and mechanical properties in austempered ductile iron ADI. The results also showed that lowering the aging temperature increases strength and re- duces ADI ductility. V. L. Richards. et al. [37] studied the aging behavior of DI. In this study, three grades of DI were investigated, two in as-cast conditions and one heat treated. The as-cast and heat-treated grades were 65-45-12, 85-55-06, and 100-70-03 respectively. Data for as cast conditions showed higher strain hardening rates for aged specimens, probably due to precipitates that would act as barriers to the dislocation movement and that would assist the hypothesis of aging precipitation in cast iron. Those that were heat treated (quenched and tempered) showed no signs of aging characteristics. V. L. Richards. et al. [29] studied the natural aging characteristics and machinability effects of ductile iron and compacted iron. Machinability test articles and tensile test samples from ductile iron and compacted iron were cast in two foundries. Natural aging in ductile iron and compacted graphite iron was statistically verified by evaluating hardness and tensile strength. As mentioned above, many studies have dealt with aging in ductile iron; nonetheless, none of them have treated the aging behavior in DMS-DI. There are no available reports about the aging behavior of this important material category. The aim of the present study is to address the aging behavior of DMS-DI with a ferrite-martensite matrix structure, by applying varied aging conditions (aging-temperatures and times) and investigating the microstructure characteristics, tensile behavior, macro, and microhardness of the aged DMS-DI. MetalsMetals2021 2021, ,11 11,, 897 x FOR PEER REVIEW 33 of of 17 17

2.2. ExperimentalExperimental ProcedureProcedure 2.1.2.1. MaterialsMaterials andand PreparationPreparation ofof SamplesSamples TheThe charging material used used to to produce produce ductile ductile iron iron consisted consisted of foundry of foundry castings castings and andsteel steel scrap. scrap. Melting Melting was carried was carried out in out the in induction the induction furnace. furnace. Magnesium Magnesium treatment treatment (Sphe- (Spheroidizationroidization treatment) treatment) was performed was performed using using the ladle the ladle transfer transfer method. method. In this In treatment, this treat- ment,the master the master alloy containing alloy containing 5.0–6.0 wt.% 5.0–6.0 Mg, wt.% 44.0–50.0 Mg, 44.0–50.0wt.% Si, 0.75–1.25 wt.% Si, wt.% 0.75–1.25 Ca, 0.8–1.6 wt.% Ca,wt.% 0.8–1.6 Re and wt.% up to Re 1 andwt.% up Al to was 1 wt.% placed Al in was the placed bottom in of the the bottom ladle and of the has ladle been andcovered has beento prevent covered magnesium to prevent evaporation magnesium by evaporation the sandwich by thetechnique. sandwich Inoculation technique. treatment Inoculation was treatmentcarried out was using carried a 75% out ferrosilicon using a 75% alloy. ferrosilicon The melt alloy. was The cast melt into was a Y cast shape into sand a Y shape mold sandprepared mold according prepared accordingto ISO/DIN to 1083. ISO/DIN The dimensions 1083. The dimensions of the Y block of theand Y the block Y shape and the sand Y shapemold sandare shown mold in are Figure shown 1. in Figure1.

FigureFigure 1.1. ((aa)) DimensionsDimensions ofof YY blockblock and,and, ((bb)) Y-shape Y-shape sand sand mold. mold.

TheThe chemicalchemical composition composition analysis analysis of theof the as-cast as-cast ductile ductile iron usediron inused the in current the current study wasstudy performed was performed using atomic using atomic emission emission spectroscopy spectroscopy (AES). (AES). 2.2. Microstructure Characterization 2.2. Microstructure Characterization An optical microscopic investigation was performed on the as-cast structure and on An optical microscopic investigation was performed on the as-cast structure and on the DMS-DI obtained after treatment cycles. The conventional preparation procedures were performedthe DMS-DI on obtained samples accordingafter treatment to standard cycles metallographic. The conventional technique preparation [42]. The procedures samples werewere mechanicallyperformed on grinding samples by according silicon carbide to standard papers withmetallographic 80, 220, 400, technique 600, 800, 1200, [42]. andThe 2500samples grit were paper. mechanically The polishing grinding was performed by silicon using carbide 1 µ papersm alumina with suspension.80, 220, 400, Metallo-600, 800, graphic1200, and samples 2500 grit were paper. examined The polishing by FEROX was PL performed optical microscopy using 1 µm (Medline, alumina Oxfordshire, suspension. UK)Metallographic after etching samples with 2% were Nital examined etchant. Theby FEROX volume PL fraction optical of microscopy each phase (Medline, in the as-cast Ox- andfordshire, in dual UK) matrix after structure etching with has been 2% Nital carried etch outant. using The volume the point fraction counting of each technique phase inin accordancethe as-cast withand in ASTM dual E562 matrix [43 ].structure Moreover, has some been samples carried wereout using selected the for point examination counting usingtechnique the CAMSCAN in accordance 44 with (Cambridge ASTM E562 scanning [43]. CompanyMoreover,Ltd., some Cambridge, samples were UK) selected scanning for electronexamination microscopy using the (SEM). CAMSCAN For SEM 44 observations, (Cambridge a 0.3scanningµm oxide-polish Company silicaLtd., Cambridge, suspension wasUK) used scanning for the electron final polishing microscopy and etching (SEM). withFor 2%SEM Nital observations, for revealing a 0.3 the µm different oxide-polish phases insilica the microstructure.suspension was Afterused etching,for the final the samplespolishing were and rinsedetching with with ethyl 2% Nital alcohol for and revealing dried underthe different a warm phases air drier. in the microstructure. After etching, the samples were rinsed with ethyl alcohol and dried under a warm air drier. 2.3. Defining the Intercritical Region 2.3. DefiningIn this study, the Intercritical the DI with Region a ferrite-martensite matrix was investigated. Therefore, it is necessaryIn this to study, specify the the DI intercritical with a ferrite-martens region in orderite tomatrix obtain was the investigated. ferrite-martensite Therefore, matrix. it Fromis necessary the previous to specify studies, the intercritical it was found region that thein order intercritical to obtain region the ferrite-martensite can be identified ma- by threetrix. From different the methods,previous namelystudies, dilatometry,it was found thermodynamic that the intercritical calculations, region can and be intercritical identified austenitizingby three different heat treatmentmethods, cycles.namely dilatometry, thermodynamic calculations, and inter- critical austenitizing heat treatment cycles. 2.3.1. Dilatometry

For dilatometric analysis to specify the intercritical region, a LINSEIS DIL L75 PT H (differential horizontal pushrod dilatometer) (LINSEIS Company, Robbinsville, NJ, USA) Metals 2021, 11, x FOR PEER REVIEW 4 of 17

Metals 2021, 11, x FOR PEER REVIEW 4 of 17

2.3.1. Dilatometry For dilatometric analysis to specify the intercritical region, a LINSEIS DIL L75 PT H 2.3.1. Dilatometry Metals 2021, 11, 897 (differential horizontal pushrod dilatometer) (LINSEIS Company, Robbinsville, NJ, USA)4 of 17 shownFor in dilatometricFigure 2 was analysisused. This to specifyanalysis the was in performedtercritical region, on three a LINSEIScylindrical DIL specimens L75 PT H of(differential 5 mm diameter horizontal and 20 pushrod mm length. dilatometer) A quartz (LINSEIS push rod Company, was used Robbinsville,to transmit the NJ, length USA) changeshown toin Figurethe linear 2 was variable used. Thisdifferential analysis transformer. was performed The on dilatometric three cylindrical measurements specimens ofshown 5 mm in diameter Figure2 wasand used.20 mm This length. analysis A quartz was performedpush rod was on threeused cylindricalto transmit specimens the length wereof 5 mmperformed diameter under and 240 20 mm mN length.contact A force. quartz During push rodthe wasthermal used cycle, to transmit the dilatometric the length curvechange was to recorded the linear to determinevariable differential the intercritical transformer. region. The dilatometric measurements werechange performed to the linear under variable 240 mN differential contact transformer.force. During The the dilatometric thermal cycle, measurements the dilatometric were performed under 240 mN contact force. During the thermal cycle, the dilatometric curve curve was recorded to determine the intercritical region. was recorded to determine the intercritical region.

Figure 2. Schematic drawing of the dilatometer DIL L75 PT [44].

2.3.2.Figure Thermodynamic 2. Schematic drawing Software of the dilatometer DIL L75 PT [44]. [44]. 2.3.2.Thermo-Calc Thermodynamic software Software with the TCFE9 database was used to specify the intercritical 2.3.2. Thermodynamic Software region.Thermo-Calc This software software enables with calculating the TCFE9 the databaseamount and was composition used to specify of each the intercriticalphase de- Thermo-Calc software with the TCFE9 database was used to specify the intercritical pendingregion. on This the software annealing enables temperature. calculating the amount and composition of each phase region. This software enables calculating the amount and composition of each phase de- depending on the annealing temperature. 2.3.3.pending Intercritical on the annealing Austenitizing temperature. Heat Treatment 2.3.3.The Intercritical intercritical Austenitizing region for the Heat as-cast Treatment ductile iron was established by the practical 2.3.3. Intercritical Austenitizing Heat Treatment methodologyThe intercritical described region in the for previous the as-cast work ductile [45,46]. iron According was established to this knowledge, by the practical the samplesmethodologyThe with intercritical dimensions described region inof the15 for mm previous the ×as-cast 15 mm work ductile × [545 mm,46 iron]. machined According was established from to thisY-block knowledge,by the were practical sub- the jectedmethodologysamples to thermal with described dimensions heat treatment in the of 15 previous cycles mm ×invo work15lving mm [45,46]. ×austenitizing,5 mm According machined holding, to this from andknowledge, Y-block quenching were the stage.samplessubjected Austenitizing with to thermal dimensions stage heat treatmentwas of 15 performed mm cycles × 15 atmm involving te mperatures× 5 mm austenitizing, machined ranging from holding,from Y-block 720 and to 860 quenchingwere °C sub- by ◦ 20jectedstage. °C step. to Austenitizing thermal At each heataustenitizing stage treatment was performedtemperature, cycles invo atlving temperaturesthe samples austenitizing, are ranging soaked holding, from for 20 720 and min to quenching and 860 thenC by ◦ quenchedstage.20 C step.Austenitizing in Atwater. each austenitizingFigure stage was3 shows performed temperature, a summary at te themperatures of samples the thermal areranging soaked heat from treatment for 720 20 min to cycles860 and °C then toby determine20quenched °C step. the inAt water.intercriticaleach austenitizing Figure region.3 shows temperature, a summary the of samples the thermal are soaked heat treatment for 20 min cycles and then to quencheddetermine in the water. intercritical Figure region. 3 shows a summary of the thermal heat treatment cycles to determine the intercritical region.

Figure 3. Summary of the intercritical austenitizing heat treatment cycles to produce DMS-DI; RT (room temperature), LOM (light optical microscopy), and PVF (phase volume fraction). Metals 2021, 11, x FOR PEER REVIEW 5 of 17

Figure 3. Summary of the intercritical austenitizing heat treatment cycles to produce DMS-DI; RT Metals 2021, 11, 897 (room temperature), LOM (light optical microscopy), and PVF (phase volume fraction). 5 of 17

2.4. Aging Treatment The aging treatment was performed on the DMS-DI with the ferrite-martensite ma- 2.4. Aging Treatment trix. The specimens were soaked for nine different aging times (2, 10, 20, 50, 100, 500, 1000, 5000, 10,000 min)The at aging treatmenttemperatures was of performed 140, 170, and on the 220 DMS-DI °C. After with aging the treatment, ferrite-martensite the matrix. specimens Thewere specimens quenched werein water soaked (to terminat for ninee carbon different diffusion) aging times and then (2, 10, stored 20, 50, frozen 100, 500, 1000, 5000, 10,000 min) at aging temperatures of 140, 170, and 220 ◦C. After aging treatment, the at 18 °C (to avoid natural aging) until applying the tensile test. The aging temperatures specimens were quenched in water (to terminate carbon diffusion) and then stored frozen and times were selected based on the common aging conditions applied on dual-phase at 18 ◦C (to avoid natural aging) until applying the tensile test. The aging temperatures steels [47,48]. and times were selected based on the common aging conditions applied on dual-phase steels [47,48]. 2.5. Tensile Test Cylindrical2.5. Tensile tensile Test test samples with a gauge length of 25 mm and an original diam- eter of 5 mm wereCylindrical machined tensile from test the samples lower with part a of gauge Y-blocks length according of 25 mm andto ISO an original6892- diameter 1:2009E [49].of The 5 mm samples were machined were machined from the prior lower to heat part oftreatment Y-blocks cycles according to avoid to ISO develop- 6892-1:2009E [49]. ing of any residualThe samples stresses were during machined machinin priorg process. to heat treatmentThe tensile cycles testing to was avoid carried developing out of any at room temperatureresidual stresses using a duringLLOYD machining universal testing process. machine The tensile (LLOYD testing Instruments, was carried UK) out at room with a 300 temperaturekN loading capacity using a LLOYDand the universalloading speed testing (cross machine head (LLOYD speed) of Instruments, 1 mm/min. UK) with a For each aging300 kNcondition, loading three capacity samples and thewere loading tested. speed (cross head speed) of 1 mm/min. For each aging condition, three samples were tested. 2.6. Hardness Test In this2.6. study, Hardness both Testmacrohardness (surface hardness) and microhardness measure- ments were performedIn this study, on as-cast, both macrohardnessunaged, and aged (surface DI. For hardness) macrohardness and microhardness measure- measure- ments, the mentsWilson were Rockwell performed hardness on as-cast, test with unaged, a Rockwell and aged B DI.scale For (100 macrohardness kg load—1/16” measurements, tungsten carbidethe Wilson ball) Rockwell(Instron GmbH, hardness Darmstadt, test with aGermany) Rockwell was B scale used. (100 For kg each load—1/16” sample, tungsten five indentationscarbide were ball) done (Instron and GmbH, the averag Darmstadt,e hardness Germany) HRB was was calculated. used. For For each micro- sample, five in- hardness measurements,dentations were the done Mitutoyo and the microhar averagedness hardness tester HRB (Mitutoyo, was calculated. Sakado, For Japan) microhardness with an opticalmeasurements, microscope thewas Mitutoyo used. The microhardness hardness of all tester phases (Mitutoyo, in unaged Sakado, (ferrite-mar- Japan) with an tensite) andoptical after aging microscope (ferrite-martensite) was used. The was hardness measured of all with phases a Vickers in unaged (10(ferrite-martensite) g load, and and 15 s durationafter time) aging hardness (ferrite-martensite) technique. Fi wasve measuredVickers microhardness with a Vickers indentations (10 g load, and were 15 s duration applied in eachtime) phase. hardness Figure technique. 4 depicts Five the Vickersexperimental microhardness procedure indentations performed during were applied the in each current work.phase. Figure4 depicts the experimental procedure performed during the current work.

Figure 4. SchematicFigure diagram 4. Schematic of the experimental diagram of the procedures experimental followed proced in thisures work: followed Ti (intercritical” in this work: annealing Ti (intercriti- temperature), cal” annealing temperature), F (ferrite), M (martensite), RT (room temperature). F (ferrite), M (martensite), RT (room temperature).

3. Results and Discussion 3.1. Characteristics of the As-Cast The chemical composition of the as-cast condition is given in Table1. The Si and Mn contents are the most significant elements that influence the range of intercritical region. These contents are conforming to the recommended values in the literature [50] for Metals 2021, 11, x FOR PEER REVIEW 6 of 17

3. Results and Discussion 3.1. Characteristics of the As-Cast Metals 2021, 11, 897 6 of 17 The chemical composition of the as-cast condition is given in Table 1. The Si and Mn contents are the most significant elements that influence the range of intercritical region. These contents are conforming to the recommended values in the literature [50] for char- characterizingacterizing the theintercritical intercritical region. region. Figure Figure 5 illustrates5 illustrates the themicrostructure microstructure of the of theas-cast as-cast con- condition,dition, which which consists consists of a of ferritic-pearlitic a ferritic-pearlitic matrix. matrix. The Theas cast as castmicrostructure microstructure was wascom- composedposed of 64% of 64% ferrite ferrite and and 36% 36% pearlite. pearlite. Addi Additionally,tionally, the the as-cast as-cast morphology morphology displayed displayed a anodularity nodularity of of about about 66% 66% according according to the requirementsrequirements ofof ASTMASTM E2567E2567standard standard [ 51[51],], the the µ 2 nodulenodule size size was was 26.4 26.4 µm,m, and and the the nodule nodule count count was was 227 227 nodules/mm nodules/mm2..

TableTable 1. 1.Chemical Chemical composition composition of of as-cast as-cast ductile ductile iron iron (wt.%). (wt.%). 1 CC SiSi MnMn MgMg SS PP FeFe CECE1 3.393.39 2.632.63 0.330.33 0.030.03 0.010.01 0.050.05 Bal.Bal. 4.274.27 CE1 = %C + 1 1(% Si + %P). CE =% C+3 % Si + % P. 3

FigureFigure 5. 5.Microstructure Microstructure of as-cast of as-cast with wi ferritic-pearliticth ferritic-pearlitic matrix. matrix. 3.2. The Intercritical Region 3.2. The Intercritical Region For the empirical equations, the lower and upper critical transformation temperatures For the empirical equations, the lower and upper critical transformation tempera- Ae1 and Ae3 were estimated using empirical equations [52]. Accordingly, the intercritical regiontures rangeA and lays A 782.7 were◦C toestimated 813.3 ◦C. using empirical equations [52]. Accordingly, the in- tercriticalFrom theregion Thermo-Calc range lays calculations, 782.7 °C to it813.3 was °C. found that the formation of austenite begins From◦ the Thermo-Calc calculations, it was found that the formation of austenite be- at 787.8 C, which is known as the lower critical transformation temperature Ae1. It was alsogins observed at 787.8 °C, that which the end is known of ferrite as formationthe lower critical occurred transformation at 821.3 ◦C, whichtemperature is known A as. It was also observed that the end of ferrite formation occurred at 821.3 °C, which is known the upper critical transformation temperature Ae3. The intercritical region is characterized ◦ byas the upper presence critical of γ tr+ansformationα + graphite temperature together, and A this. The is intercritical only achieved region from is character- 787.8 C toized 821.3 by◦ theC. presence of γ + α + graphite together, and this is only achieved from 787.8 °C to 821.3The dilatometric°C. measurements were carried out by heating the samples up to 1000 ◦C. For confirmingThe dilatometric the equilibrium measurements measurements, were carried the out samples by heating were heatedthe samples with up a heating to 1000 rate°C. For of 0.08 confirming K/s and the cooled equilibrium with a measurements, cooling rate of the 0.17 samples K/s to were detect heated the dilatometric with a heat- (Temperature—lengthing rate of 0.08 K/s and change) cooled curve with shown a cooling in Figure rate of6. 0.17 K/s to detect the dilatometric (Temperature—lengthFrom the dilatometry change) measurements, curve shown we in find Figure that 6. the curve begins to deviate from linearity at ~755 ◦C during heating, which indicates the (austenite formation) start of the ferrite to austenite transformation. The curve restores linearity at ~890 ◦C, indicating the transformation is complete. During cooling, the dilatometric curve deviates from linearity at ~843 ◦C, indicating the beginning of ferrite formation, and then returns to linearity again at ~714 ◦C, which corresponds to the end of ferrite formation. Within these temperature ranges (755 to 890 ◦C during heating and 714 to 843 ◦C during cooling), the ferrite, austenite, and graphite coexist, and these temperature ranges correspond to the intercritical region. Metals 2021, 11, 897 7 of 17 Metals 2021, 11, x FOR PEER REVIEW 7 of 17

Figure 6. DilatationDilatation temperature temperature curve during heating with 0.08 (K/s) and and cooling cooling with with 0.17 0.17 (K/s). (K/s).

FromThe intercritical the dilatometry region measurements, was determined we according find that the to the curve methodology begins to deviate described from in linearitythe experimental at ~755 °C procedure during heating, (Section which2.3.3). Partiallyindicates austenitizing the (austenite at formation) different temperatures start of the ferritewithin to the austenite intercritical transformation. region (α + Theγ) enables curve restores controlling linearity as well at ~890 as obtaining °C, indicating different the transformationamounts of ferrite is complete. and martensite. During Figure cooling,7 demonstrates the dilatometric the microstructures curve deviates offrom ductile linearity iron atobtained ~843 °C, after indicating austenitizing the beginning and water of ferrite quenching formation, steps. and Generally, then returns these to microstructures linearity again atconsist ~714 °C, of ferrite which (labeled corresponds F), martensite to the end (labeled of ferrite M), formation. and graphite Within (spheroidal these temperature graphite). Through this metallographic examination, it can be observed that the martensite was ranges (755 to 890 °C during heating and 714 to 843 °C during cooling), the ferrite, austen- completely formed at 860 ◦C, indicating that 20 min (holding time) at this austenitizing ite, and graphite coexist, and these temperature ranges correspond to the intercritical re- temperature is sufficient to reach equilibrium. gion. Figure8 depicts the measured matrix volume fractions of ferrite (FVF) and martensite The intercritical region was determined according to the methodology described in (MVF) as a function of austenitizing temperature. It is evident that as the partial austenitiz- the experimental procedure (Section 2.3.3.). Partially austenitizing at different tempera- ing temperature increases, the MVF increases and the FVF decreases. The A is established tures within the intercritical region (α + γ) enables controlling as well as obtaininge1 differ- at 740 ◦C, due to the presence of a martensite, produced by quenched austenite. Likewise, ent amounts of ferrite and martensite. Figure 7 demonstrates the microstructures of duc- the A is defined as 860 ◦C, due to the martensite is fully formed. tile irone3 obtained after austenitizing and water quenching steps. Generally, these micro- Table2 lists the results of characterizing of the intercritical region by the different structures consist of ferrite (labeled F), martensite (labeled M), and graphite (spheroidal techniques as described above together with the empirically calculated ones. Based on the graphite). Through this metallographic examination, it can be observed that the marten- results shown in this table, it can be concluded that estimating the range of the intercrit- Metals 2021, 11, x FOR PEER REVIEWsite was completely formed at 860 °C, indicating that 20 min (holding time) at this austen-8 of 17 ical region using dilatometric method by continuous heating is more accurate than the itizingother method. temperature is sufficient to reach equilibrium. Figure 8 depicts the measured matrix volume fractions of ferrite (FVF) and martensite (MVF) as a function of austenitizing temperature. It is evident that as the partial austen- itizing temperature increases, the MVF increases and the FVF decreases. The A is estab- lished at 740 °C, due to the presence of a martensite, produced by quenched austenite. Likewise, the A is defined as 860 °C, due to the martensite is fully formed. Table 2 lists the results of characterizing of the intercritical region by the different techniques as described above together with the empirically calculated ones. Based on the results shown in this table, it can be concluded that estimating the range of the intercritical region using dilatometric method by continuous heating is more accurate than the other method.

FigureFigure 7.7. MicrostructuresMicrostructures ofof ductileductile ironiron obtainedobtained afterafter austenitizingaustenitizing andand waterwater quenched:quenched: ((aa)) 720720◦ C, (°C,b) 740 (b)◦ 740C, ( c°C,) 760 (c)◦ 760C, ( d°C,) 780 (d)◦ 780C, ( e°C,) 800 (e)◦ 800C, ( f°C,) 820 (f)◦ C,820 (g °C,) 840 (g)◦ C,840 (h °C,) 860 (h◦) C,860 F (ferrite),°C, F (ferrite), M (martensite). M (mar- tensite).

Figure 8. Measured phase fractions (respecting the DI matrix) as a function of austenitizing tem- perature, FVF (ferrite volume fraction), MVF (martensite volume fraction).

Table 2. Results of the intercritical region by four different techniques.

Dilatometry Technique Empirical Equa-ThermoCalc Soft-Heating Cooling Partial Austenitizing Heat tions ware 𝐀 𝐀 Treatment Temperature 𝐜𝟏 𝐫𝟑 𝐀𝐜𝟑 𝐀𝐫𝟏 Lower critical transformation tem- 782.7 787.8 755 714 740 perature [°C] Upper critical transformation tem- 813.3 821.3 890 843 860 perature [°C]

3.3. Microstructural Analysis Based on the results obtained from the determination of the intercritical region, the austenitizing temperature of 785 °C within this range was selected to produce the ferrite- martensite DMS-DI. Figure 9 shows the microstructures of the unaged samples after aus- tenitization at 785 °C for 20 min and quenching in water to room temperature. As shown Metals 2021, 11, x FOR PEER REVIEW 8 of 17

Figure 7. Microstructures of ductile iron obtained after austenitizing and water quenched: (a) 720 Metals 2021, 11, 897 8 of 17 °C, (b) 740 °C, (c) 760 °C, (d) 780 °C, (e) 800 °C, (f) 820 °C, (g) 840 °C, (h) 860 °C, F (ferrite), M (mar- tensite).

Figure 8. Measured phase fractions (respecting the DI matrix) as a function of austenitizing tempera- ture,Figure FVF 8. (ferrite Measured volume phase fraction), fractions MVF (martensite(respecting volume the DI fraction). matrix) as a function of austenitizing tem- perature, FVF (ferrite volume fraction), MVF (martensite volume fraction). Table 2. Results of the intercritical region by four different techniques. Table 2. Results of the intercritical region by four different techniques. Technique Dilatometry Partial Empirical ThermoCalc Heating Cooling Austenitizing Equations Software Dilatometry TemperatureTechnique Ac1−Ac3 Ar3−Ar1 Heat Treatment Empirical Equa-ThermoCalc Soft-Heating Cooling Partial Austenitizing Heat Lower critical transformation ◦ 782.7tions 787.8ware 755𝐀 𝐀 714 740Treatment temperatureTemperature [ C] 𝐜𝟏 𝐫𝟑 𝐀 𝐀 Upper critical transformation 𝐜𝟑 𝐫𝟏 813.3 821.3 890 843 860 Lower criticaltemperature transformation [◦C] tem- 782.7 787.8 755 714 740 perature [°C] Upper critical transformation3.3. tem- Microstructural Analysis 813.3 821.3 890 843 860 perature [°C] Based on the results obtained from the determination of the intercritical region, the austenitizing temperature of 785 ◦C within this range was selected to produce the ferrite- martensite3.3. Microstructural DMS-DI. FigureAnalysis9 shows the microstructures of the unaged samples after austenitization at 785 ◦C for 20 min and quenching in water to room temperature. As shownBased in Figure on9 thea, the results microstructure obtained observed from the by lightdetermination optical microscope of the (LOM) intercritical shows region, the aaustenitizing matrix consisting temperature of ferrite (bright of 785 color) °C andwithin martensite this range (dark color).was selected There may to beproduce phases the ferrite- thatmartensite are not definedDMS-DI. under Figure the LOM9 shows due the to their micr smallostructures size. For of this the reason, unaged a scanning samples after aus- electrontenitization microscope at 785 (SEM) °C for was 20 used min to and elucidate quenching these phases. in water The to SEM room micrograph temperature. of the As shown unaged samples is presented in Figure9b. It is also observed that there is an indistinct phase, such as the area denoted by red rectangle in the same figure. Referring to the corresponding micrographs with higher magnification shown in Figure9c,d, these revealed that the plate martensite morphology was formed by quenching ductile iron to room temperature. These features correspond to the typical morphology of quenched ductile iron [53], but without retained austenite due to austenitization at 785 ◦C within the intercritical region. The SEM micrographs of Figure9 illustrate that the formed martensite is chiefly a plate-type martensite. However, pearlite is formed at the interface with the ferrite phase as shown in Figure9. The holding time and temperature were not sufficient for dissolving the whole pearlite existing in the as-cast condition. About 1.5% of this phase is retained at the ferrite-martensite interface and appears as a black-phase in the LOM micrographs (Figure9a) and revealed a fine lamellar structure in the SEM micrographs (Figure9c,d). Metals 2021, 11, x FOR PEER REVIEW 9 of 17

in Figure 9a, the microstructure observed by light optical microscope (LOM) shows a ma- trix consisting of ferrite (bright color) and martensite (dark color). There may be phases that are not defined under the LOM due to their small size. For this reason, a scanning electron microscope (SEM) was used to elucidate these phases. The SEM micrograph of the unaged samples is presented in Figure 9b. It is also observed that there is an indistinct phase, such as the area denoted by red rectangle in the same figure. Referring to the cor- responding micrographs with higher magnification shown in Figure 9c,d, these revealed Metalsthat2021 the, 11, 897plate martensite morphology was formed by quenching ductile iron to room tem- 9 of 17 perature. These features correspond to the typical morphology of quenched ductile iron [53], but without retained austenite due to austenitization at 785 °C within the intercritical region. The untempered martensite appears as a featureless structure with a higher level compared to the ferrite phase.

Figure 9. Microstructure of ferrite-martensite matrix (unaged), (a) LOM, (b) SEM, (c) and (d) SEM micrographs with high Figure 9. Microstructure of ferrite-martensite matrix (unaged), (a) LOM, (b) SEM, (c) and (d) SEM magnifications of area indicated by red rectangle in (b). micrographs with high magnifications of area indicated by red rectangle in (b). The SEM micrographs of specimens aged at a temperature of 170 ◦C for a 1000 and The SEM micrographs10,000 min of areFigure presented 9 illustrate in Figure that10. It the was observedformed thatmartensite after 1000 is min, chiefly there area visible plate-type martensite.small However, needle precipitatespearlite is in formed the ferrite. at Extendingthe interface of the with aging the time ferrite up to 10,000 phase min led to increased needle precipitations. These precipitates are believed to be formed during aging, as shown in Figure 9. Theas they holding were absent time in and the unagedtemperature condition were as shown not sufficient in Figure9c,d. for As dissolving clearly observable the whole pearlite existingin Figure in the10b,c,e, as-cast the carbide condit precipitatesion. About are 1.5% rarely of evident this phase within is a ferriteretained margin at around the ferrite-martensite theinterface ferrite-martensite and appears interface. as Thisa black-phase observation canin bethe correlated LOM micrographs to the lower saturation (Figure 9a) and revealedwith a carbon fine lamellar at these zones stru comparedcture in tothe the SEM regions micrographs that showed large(Figure precipitates 9c,d). [48]. It is predicted using thermodynamic calculations using Thermo-Calc (Thermo-Calc Software The untempered martensiteAB, Solna, appears Sweden) as that a ~0.02featureless wt.% of Cstructure dissolves inwith ferrite a higher during intercritical level com- annealing. pared to the ferrite phase.Throughout quenching from the intercritical region to RT, the solubility of carbon in ferrite The SEM micrographsdecreases of byspecimens decreasing aged the temperature, at a temperature which results of 170 in diffusion °C for ofa carbon1000 and from ferrite 10,000 min are presentedto the in austenite/martensite Figure 10. It was ob phaseserved during that the after course 1000 of quenching. min, there However, are visible this diffusion small needle precipitatesprocess in the is limited ferrite. by Extending the short time of ofthe quenching, aging time therefore up to resulting10,000 min in lower led carbon content at the neighborhood of the martensite phase compared to the that far away from it, to increased needle precipitations.which require long-range These prec diffusionipitates between are phasesbelieved [48 ].to be formed during aging, as they were absent in the unaged condition as shown in Figure 9c,d. As clearly observable in Figure 10b,c,e, the carbide precipitates are rarely evident within a ferrite margin around the ferrite-martensite interface. This observation can be correlated to the lower saturation with carbon at these zones compared to the regions that showed large precipitates [48]. It is predicted using thermodynamic calculations using Thermo-Calc (Thermo-Calc Software AB, Solna, Sweden) that ~0.02 wt.% of C dissolves in ferrite during Metals 2021, 11, x FOR PEER REVIEW 10 of 17

intercritical annealing. Throughout quenching from the intercritical region to RT, the sol- ubility of carbon in ferrite decreases by decreasing the temperature, which results in dif- fusion of carbon from ferrite to the austenite/martensite phase during the course of quenching. However, this diffusion process is limited by the short time of quenching, Metals 2021, 11, 897 therefore resulting in lower carbon content at the neighborhood of the martensite phase 10 of 17 compared to the that far away from it, which require long-range diffusion between phases [48].

FigureFigure 10. Typical 10. Typical SEM SEM micrographs micrographs of of samples samples aged aged atat 170170 °C,◦C, ( (aa––cc) )for for 1000 1000 min, min, (d– (fd) –forf) for10,000 10,000 min. min. 3.4. Mechanical Properties of the Aged DMS-DI 3.4. Mechanical Properties of the Aged DMS-DI The mechanical properties of the DMS-DI before aging treatment are summarized The mechanical properties of the DMS-DI before aging treatment are summarized in Table 3. As mentionedin Table3 above,. As mentioned the specific aim above, of this the research specific is aimto study of this the effect research of aging is to study the effect of treatment on agingthe mechanical treatment properties on the mechanical of the ductile properties iron with ofthe the ferrite-martensite ductile iron with ma- the ferrite-martensite trix structure.matrix To assess structure. the influence To assess of ag theing influence on the mechanical of aging properties, on the mechanical the tensile properties, the tensile and hardnessand properties hardness were properties evaluated werebefore evaluated and after aging. before Figure and after11 shows aging. the Figureeffect 11 shows the effect of aging temperatureof aging and temperature time on tensile and time and onhardness tensile properties and hardness for samples properties intercriti- for samples intercritically cally annealedannealed at 785 °C at and 785 aged◦C andat temperatures aged at temperatures of 140, 170, and of 140,220 °C 170, for and aging 220 times◦C for aging times from from 2 to 10,0002 to min. 10,000 min. In unaged condition,In unaged the condition, samples had the samplesrelatively hadhigher relatively ultimate higher tensile ultimate strength tensile strength (UTS) (UTS) due todue the presence to the presence of the untempered of the untempered plate martensite. plate martensite. At the lowest At aging the lowest tem- aging temperature perature of 140of °C, 140 the◦C, (UTS) the (UTS)has not has shown not a shown significant a significant change with change increasing with the increasing aging the aging time as time as shown in Figure 11a. A significant decrease in the UTS for the samples aged at 170 shown in Figure 11a. A significant decrease in the UTS for the samples aged at 170 and and 220 °C is observed with prolonged aging time. The decreasing in UTS starts at about 220 ◦C is observed with prolonged aging time. The decreasing in UTS starts at about 500 min and 100 min of holding at 170 °C and 220 °C, respectively.◦ ◦ The aging500 response min and (BH 1000 level) min ofof holdingthe DMS-DI at 170 wasC calculated and 220 fromC, respectively. the difference between the 0.2% Theoffset aging yield strength response after (BH and0 level) before of aging the DMS-DIprocess. Figure was calculated11b shows from the difference the strengtheningbetween stages the of 0.2% aging offset process yield without strength pre-straining after and as before a function aging of process. aging Figure 11b shows time. In iron thecarbon strengthening alloys, it is a stages known of fact aging that process the mechanical without properties pre-straining may asbe a af- function of aging time. fected by soluteIn ironatoms carbon (carbon alloys, and nitrogen) it is a known diffusion fact according that the to mechanical the mechanisms: properties may be affected by 1. Stress-inducedsolute arrangement atoms (carbon of solute and nitrogen) atoms in interstitial diffusion sites according (Snoek toeffect). the mechanisms: 1. Stress-induced arrangement of solute atoms in interstitial sites (Snoek effect). 2. Segregation of solute atoms into dislocations to form and stabilize a Cottrell atmosphere. 3. Precipitation of carbides.

Table 3. The mechanical properties of the DMS-DI before aging treatment: UTS (ultimate tensile strength), YS (yield strength 0.2% offset stress), TEL (total elongation), HRB (Rockwell hardness B), FMH (ferrite microhardness, MMH (martensite microhardness), and E (elastic modulus).

FMH MMH UTS [MPa] YS [MPa] TEL [%] HRB E [GPa] [kgf/mm2] [kgf/mm2] 554.33 ± 14.14 402.33 ± 3.97 4.52 ± 2.11 85.8 ± 1.48 235.3 ± 2.59 1029 ± 15.01 154.6 ± 25.11 Metals 2021, 11, 897 11 of 17 Metals 2021, 11, x FOR PEER REVIEW 12 of 17

Figure 11. The effect of aging temperature and aging time on the tensile and hardness properties of aged DI with DMS: (a) Change in ultimate tensile strength UTS, (b) change in yield strength YS, and (c) hardness (HRB). Figure 11. The effect of aging temperature and aging time on the tensile and hardness properties of aged DI with DMS: (a) Change in ultimate Intensile general, strength the phenomenonUTS, (b) change of in aging yield isstrength based onYS, theand diffusion (c) hardness rate (HRB). of interstitial atoms towards dislocations. The increase in the yield strength (YS) goes through two plateaus as shown in3.5. Figure Microhardness 11b. The maximum level of BH0 is achieved in the samples aged at 140, ◦ 170, and 220 MicrohardnessC for an aging timetests were of 5000, carried 500, andout to 100 esti min,mate respectively. the changes Itoccurring is clear that in the ferrite the agingand temperature martensite plays phase a significant due to aging effect treatments. in the aging The kinetics; Vickers increasing microhardness the aging of the ferrite Metals 2021, 11, 897 12 of 17

temperature pronouncedly accelerates it. The effect of aging temperature on the achieved BH0 responses in the lower and upper plateaus is rather limited. Although there is no available research indicating that formation of a Cottrell atmo- sphere in DMS-DI, some research works [47,48] performed on dual-phase steel suggest that the first increase in yield strength (first plateau) is caused by Cottrell atmosphere. In this strengthening stage, the all-aged samples exhibited almost the same aging response during the short aging time. A subsequent strengthening stage is observed due to the formation of the precipitates (Figure 10b,c,e,f). At this stage, the BH0 value is mainly dependent on the aging temperature and aging time as a result of carbides precipitation. These precipitates were absent in the unaged condition (Figure9). ◦ A slightly decrease in BH0 level is observable for samples aged for 10 min at 140 C and at 170 ◦C. This behavior can be attributed to the weakening of dislocations locking by the Cottrell atmosphere. Prolonging the aging process, the effect of the over-aging treatment is observed, and the yield strength (BH0 level) decreases after reaching the maximum value. This effect can be observed at the early aging times of 1000, 500, and 100 min when aging at 140, 170, and 220 ◦C, respectively. The decrease in strength can be attributed to the martensite tempering as well, which will be shown in the next section. The measured Rockwell hardness B (HRB) as a function of aging time is shown in Figure 11c. Ductile iron with ferrite + martensite matrix does not exhibit the hardness plateau and over-aging effect associated with the traditional aging phenomenon; instead, the hardness fluctuated within a range of less than 5%, i.e., within the experimental error range.

3.5. Microhardness Microhardness tests were carried out to estimate the changes occurring in the ferrite and martensite phase due to aging treatments. The Vickers microhardness of the ferrite and martensite as a function of both the aging temperature and time are depicted in Figure 12. Figure 12 shows that the aging has no significant effect on the micro-hardness of the ferrite, whereas the martensite hardness declines with increasing the aging time. This observation is revealed in Figure 13, where the size microhardness indentation of the DMS- DI before aging and after aging for 1000 min and 10,000 min are almost similar. However, the indentation-size in martensite increases by increasing the aging time, indicating a decrease in hardness. Therefore, the formed precipitates in ferrite during aging (Figure 10) has not resulted in a significant increase in hardness. This might be due to the fact that the aging process is accompanied by decreasing the supersaturation of carbon in the ferrite as mentioned before. Additionally, Shan et al. showed analytically that the carbon saturating ferrite in the case of dual-phase steel decreases during the aging process [54]. A similar phenomenon is also expected in DMS-DI. Therefore, the increase in strength due to precipitates formation is counteracted by a decrease in solid solution strengthening of ferrite. This may justify the observation that aging the DMS-DI has not resulted in increasing the macrohardness. Furthermore, the ferrite regions that appeared under LOM and SEM as free of the needle precipitates showed similar microhardness as that with the precipitates. This can be also deduced from the similarity of indentation sizes within the former regions in Figure 13c,d with the latter regions in Figure 13b,d. Metals 2021, 11, x FOR PEER REVIEW 13 of 17

Metals 2021, 11, 897 13 of 17 and martensite as a function of both the aging temperature and time are depicted in Figure 12.

Figure 12. The effect of temperature and aging-time on the microhardness: (a) Ferrite microhardness, Figure (12.b) martensiteThe effect of microhardness. temperature and aging-time on the microhardness: (a) Ferrite microhard- ness, (b) martensite microhardness. The decline in the hardness of the martensite (Figure 12b) is also revealed in Figure 13a–c. FigureFigure 12 12 showsb shows that thatthe aging this decrease has no signif startsicant to beeffect more on significantthe micro-hardness after different of the times ferrite,depending whereas the on martensite the aging temperature.hardness declin Thesees with martensite increasing tempering the aging times time. correspond This ob- to the decay in the BH0 values (compare Figure 11b with Figure 12b). Therefore, the over-aging servation is revealed in Figure 13, where the size microhardness indentation of the DMS- phenomenon is closely related to the martensite tempering in DMS-DI. DI before aging and after aging for 1000 min and 10,000 min are almost similar. However, The high Si content of the investigated ductile iron has a direct effect on the activation the indentation-size in martensite increases by increasing the aging time, indicating a de- energy of aging process, which affects the yielding behavior; the silicon in solid solution crease in hardness. Therefore, the formed precipitates in ferrite during aging (Figure 10) reduces the diffusivity of carbon atoms in the ferrite phase [55]. The investigated ductile has not resulted in a significant increase in hardness. This might be due to the fact that the iron in this work has about 2.63 wt.% silicon, which is very high compared to most steels. aging process is accompanied by decreasing the supersaturation of carbon in the ferrite as Finally, in order to be able to develop a basic understanding of the aging process in mentioned before. Additionally, Shan et al. showed analytically that the carbon saturating ferrite + martensite ductile iron, an extensive research effort is still required. ferrite in the case of dual-phase steel decreases during the aging process [54]. A similar phenomenon is also expected in DMS-DI. Therefore, the increase in strength due to pre- cipitates formation is counteracted by a decrease in solid solution strengthening of ferrite. This may justify the observation that aging the DMS-DI has not resulted in increasing the macrohardness. Furthermore, the ferrite regions that appeared under LOM and SEM as Metals 2021, 11, x FOR PEER REVIEW 14 of 17

free of the needle precipitates showed similar microhardness as that with the precipitates. This can be also deduced from the similarity of indentation sizes within the former regions in Figure 13c,d with the latter regions in Figure 13b,d. The decline in the hardness of the martensite (Figure 12b) is also revealed in Figure 13a–c. Figure 12b shows that this decrease starts to be more significant after different times depending on the aging temperature. These martensite tempering times correspond to the decay in the BH0 values (compare Figure 11b with Figure 12b). Therefore, the over-aging phenomenon is closely related to the martensite tempering in DMS-DI. The high Si content of the investigated ductile iron has a direct effect on the activation energy of aging process, which affects the yielding behavior; the silicon in solid solution reduces the diffusivity of carbon atoms in the ferrite phase [55]. The investigated ductile iron in this work has about 2.63 wt.% silicon, which is very high compared to most steels. Metals 2021, 11, 897 14 of 17 Finally, in order to be able to develop a basic understanding of the aging process in ferrite + martensite ductile iron, an extensive research effort is still required.

Figure 13. MicrohardnessMicrohardness indentation indentation (denoted (denoted by by red red circles); circles); (a)( unaged,a) unaged, (b) (170b) 170 °C—1000◦C—1000 min, min, ((cc)) andand (d) 170 °C—10,000◦C—10,000 min min aging aging time. time.

4. Conclusions This study study is is an an attempt attempt to tounderstand understand the theaging aging behavior behavior of DMS-DI. of DMS-DI. DMS-DI DMS-DI with withMVF MVFof ~30% of ~30% was wasproduced produced by austenitizing by austenitizing in the in the intercritical intercritical region region followed followed by by quenching in water to room temperature. temperature. The The effect effect of of aging aging process process on on the the mechanical mechanical properties waswas investigated.investigated. The The increase increase in in yield yield strength strength (BH (BH0 response)0 response) was was investigated investi- forgated aged for samples. aged samples. Based onBased the experimentalon the experime resultsntal in results this study, in this the study, following the conclusionsfollowing areconclusions obtained: are obtained: 1. The estimating range of the intercritical region by dilatometric method applying con- tinuous heating transformation is more accurate than that estimated by ThermoCalc software and empirical equations. The latter two methods overestimate the lower and underestimate the upper critical transformation temperatures by about +45 ◦C and −43 ◦C, respectively. 2. For all investigated aged samples, the aging conditions have no major effect on the ultimate tensile strength UTS and macro hardness. 3. The increase in yield strength BH0 value (~11% increase) associated with the aging process is observed to go through two strengthening stages followed by a decrease in BH0 value (over-aging). 4. The maximum BH0 is found to be achieved after aging for 100, 500, and 5000 min when aging at 220 ◦C, 170 ◦C, and 140 ◦C, respectively. 5. The over-aging in DMS-DI is found to be closely related to the martensite tempering, which was revealed by microhardness measurements. The decrease in the martensite hardness starts to be more significant after aging times corresponding to the over- aging start-times. Metals 2021, 11, 897 15 of 17

6. The aging process has no significant effect on the microhardness of the ferrite phase in DMS-DI, which recorded a HV0.01 value of about 235 kgf/mm2.

Author Contributions: Conceptualization, M.S. and A.E.; data curation, A.A.; formal analysis, A.A. and M.S.; investigation, A.A.; methodology, A.A., M.S. and A.E.; project administration, A.A. and A.E.; resources, M.S., H.P. and A.E.; software, M.S.; supervision, M.S. and A.E.; validation, A.A.; visualization, A.A. and M.S.; writing—original draft, A.A. and M.S.; writing—review and editing, H.P. and A.E. All authors have read and agreed to the published version of the manuscript. Funding: This research received no external funding. Institutional Review Board Statement: Not applicable. Informed Consent Statement: Not applicable. Data Availability Statement: The data presented in this study are available on request from the corresponding author. The data are not publicly available at this time as they form part of an ongoing study. Conflicts of Interest: The author declares no conflict of interest.

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