Materials and Design 178 (2019) 107862

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Materials and Design

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Austenite decomposition and carbon partitioning during quenching andpartitioning heat treatments studied via in-situ X-ray diffraction

Sandra Ebner a,⁎, Clemens Suppan b, Andreas Stark c, Ronald Schnitzer a, Christina Hofer a

a Department of Materials Science, Montanuniversität Leoben, Franz-Josef-Strasse 18, 8700 Leoben, Austriab voestalpine Stahl GmbH, voestalpine-Strasse 3, 4020 Linz, Austriac Institute of Materials Research, Helmholtz-Zentrum Geesthacht, Max-Planck-Straße 1, 21502 Geesthacht, Germany

H I G H L I G H T S G R A P H I C A L A B S T R A C T

• Transformation kinetics during Q&Pheat treatments were studied byHEXRD and correlated with the me-chanical properties.

• Effective carbon partitioning into theaustenite during 2-step processing en-hanced the strain hardening behavior.

• Poor austenite stabilization after 1-stepprocessing resulted in low yieldstrength and high initial strain harden-ing.

• More pronounced bainite formation athigher quenching temperature did notaffect the mechanical properties.

⁎ Corresponding author.E-mail address: sandra.ebner@unileoben.ac.at (S. Ebne

https://doi.org/10.1016/j.matdes.2019.1078620264-1275/© 2019 The Authors. Published by Elsevier Ltd

a b s t r a c t

a r t i c l e i n f o

Article history:Received 15 March 2019Received in revised form 30 April 2019Accepted 15 May 2019Available online 19 May 2019

High strength combined with excellent ductility can be achieved by quenching and partitioning (Q&P) micro-structures containing martensite and a considerable amount of retained austenite. Since the mechanical proper-ties are inherited from the microstructure, a thorough understanding of this relationship is indispensable. In thepresent work, in-situ synchrotron X-ray diffraction was used to investigate the transformation kinetics duringQ&P processing. The effect of different heat treatment conditions on themicrostructural evolutionwas examinedand correlated to the mechanical properties obtained by tensile testing. The results showed that austenite de-composition occurred for all Q&P cycles, especially at the beginning of partitioning. The extent of this decompo-sition was affected by a change of the quenching temperature, while the partitioning temperature showed nosignificant influence. Regardless of the heat treatment parameters, carbon partitioningwas clearly visible duringthe 2-step cycles, which led to enhanced work hardening with increasing strain. In contrast, this was not ob-served in the case of 1-step processing due to negligible carbon diffusion, and thus insufficient chemical stabili-zation of the austenite.© 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://

creativecommons.org/licenses/by-nc-nd/4.0/).

Keywords:Advanced high strength steelsQuenching and partitioningMicrostructure evolutionHigh-energy X-ray diffractionMechanical properties

r).

. This is an open access article under

1. Introduction

Quenching and partitioning is one of the concepts proposed for the3rd generation of advanced high strength steels (AHSS) to provide themechanical properties required for automotive applications [1–3]. The

the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).

2 S. Ebner et al. / Materials and Design 178 (2019) 107862

core of this heat treatment comprises a so-called quenching step, whereaustenite partially transforms into martensite, followed by apartitioning step, with the aim to stabilize the remaining austenite bycarbon enrichment. This partitioning step is either carried out at thequenching temperature (1-step Q&P) or at elevated temperatures (2-step Q&P). Insufficiently stabilized austenite may transform to freshmartensite upon final cooling. Apart from this, the final microstructureideally consists of tempered martensite and retained austenite, whichresults in a beneficial combination of strength and ductility [4–7]. Thelatter is achieved by enhanced strain hardening due to the transforma-tion induced plasticity (TRIP) effect [7–9]. Since the contribution of theTRIP-effect to the work hardening behavior is not only affected by theamount of retained austenite but also by its stability, much attentionhas been given to study the influence of different factors (carbon con-tent, grain size, morphology, surrounding matrix, etc.) on the austenitestability inQ&P steels [9–12]. For example, deKnijf et al. [10] studied themechanically-induced martensite to austenite transformation in a0.25C-1.5Si-3Mn steel subjected to a 2-step Q&P heat treatment. Thetransformation stability was found to be mainly influenced by grainsize, morphology and crystallographic orientation. The importance ofthe surroundingmicrostructure was revealed by Hidalgo et al. [12]. Mi-crostructures consisting of austenite and martensite with different de-grees of tempering were created by 2-step Q&P cycles followed byreheating steps to either 450 °C or 700 °C. A higher degree of martensitetempering lowered the mechanical stability of austenite.

The influence of the heat treatment parameters on the mechanicalproperties was extensively studied in several works [4,5,7,9,13–17]. Ingeneral, the quenching temperature should be high enough to obtainsufficient austenite fractions, but too high temperatures may lead tothe formation of fresh martensite upon final cooling due to insufficientstabilization. Partitioning parameters should be adopted in order to sta-bilize austenite while simultaneously preventing competing reactionslike carbide formation or isothermal decomposition of austenite.

It was shown by Toji et al. [18] and Thomas et al. [19] that the carbondiffusion from the supersaturated martensite into the austenite is themain reaction taking place during partitioning. In both studies, atomprobe tomography (APT) revealed clear carbon enrichment in the aus-tenite during partitioning. However, carbon trapping and carbide for-mation were found to compete with carbon partitioning. Based on thelocal chemistry obtained from APT, these carbides were assigned as ce-mentite by Toji et al. [18,20]. Thomas et al. [19] identified their carbidesas transition carbides (ε or η) by transmission electron microscopy(TEM). Precipitation of ε-carbides was also observed by HajyAkbaryet al. [21]. Pierce et al. [22] applied a combination of Mössbauer spec-troscopy and TEM to investigate and quantify η-carbides during Q&Pprocessing. Furthermore, cementite was found to occur as a conse-quence of austenite decomposition preferentially at higher partitioningtemperatures and prolonged holding [13,23].

Austenite decomposition is another competing process frequentlyobserved during partitioning, which was often indicated by an expan-sion visible in dilatometer curves [17,21,24–26]. In the case of 2-stepQ&P, partitioning is conducted above Ms in a temperature range typicalfor bainite formation. Thus, an observed decrease of the austenite frac-tion is likely caused by the formation of bainite. For example,HajyAkbary et al. [21] conducted a comprehensive study of the compet-ing reactions occurring during 2-step Q&P processing of a 0.3C-1.6Si-3.5Mn (wt%) steel. Bainite formation was found to decrease withlower quenching temperatures, whichwas accounted to the lower frac-tion of initially formed austenite and its faster stabilization duringpartitioning. Regarding 1-step Q&P, partitioning is conducted belowMs. Van Bohemen et al. [27] analyzed dilatometer curves and scanningelectron microscopy (SEM) images of a 0.66C-0.69Mn-0.3Si (wt%)steel partitioned at varying temperatures above and below Ms. Theyconcluded that isothermal bainite is formed below Ms. Similar resultswere obtained by Samanta et al. [28], who attributed the observedtransformation to a displacive growth of bainite. Contradictions to

these findings can be found by Somani et al. [25] and Kim et al. [29].Both observed an isothermal product with wavy boundaries andledge-like protrusions after isothermal treatments below Ms. Whilethis phase was identified as martensite by Somani et al. [25], Kim et al.[29] concluded that the characteristics are neither fully martensiticnor bainitic. These findings have been discussed recently by Navarro-López et al. [30]. Their study of bainitic/martensitic structures obtainedafter heat treatments conducted above and below Ms showed that theobserved protrusions can be explained by the formation of bainitic fer-rite from nucleation sites of the athermal martensite resulting in wavyboundaries. This supports the occurrence of bainite below Ms in hypo-eutectoid steels.

Common methods for microstructural characterization are oftenlimited to the final heat-treated condition. Though phase transforma-tions can be observed by dilatometry [17,21,24–26], precise detailsabout the occurring phases and the transformation kinetics during theheat treatment are not accessible. This issue can be overcome by theuse of in-situ experiments. In this manner, time-resolved informationabout the qualitative and quantitative amounts of present phases canbe gathered during the entire process. Regarding Q&P steels, neutrondiffraction [31,32] and high-energy X-ray diffraction (HEXRD) [33–36]were used to study the transformation characteristics during Q&P pro-cessing. Both methods revealed carbon diffusion into the austenite bythe dilatation of the austenite lattice parameter and confirmed thatthe final carbon content in the austenite is lower than theoretically ex-pected due to carbon trapped in themartensite [31,34,35]. Moreover, itwas shownbyRieger et al. [33] that austenite decomposition takes placeduring partitioning, leading to similar final austenite fractions regard-less of the chosen quenching temperature. In contrast, a dependenceof the retained austenite fraction on the process parameters wasfound by Allain et al. [35], who conducted a comprehensive study ofthe temporal evolution of austenite fraction, lattice parameters and car-bonmass balance including three different Q&P cycles. A decrease of thequenching temperature led to less retained austenite with higher car-bon content. With higher partitioning temperature, the amount ofretained austenite increased and the carbon content decreased. It wasfurther pointed out in thiswork that the estimation of the austenite car-bon content strongly depends on the applied method. This was attrib-uted to internal stresses that are generated in the austenite duringprocessing and studied in-depth in [36].

Although the microstructural evolution during Q&P processing hasbeen extensively investigated using in-situ experiments, these resultshave not been correlated to the mechanical properties so far. Thereexist, however, a few publications in which HEXRD experiments servedas a basis for numerical modeling of mechanical properties. For exam-ple, Hu et al. [37,38] studied the individual phase properties of a Q&Psteel by combining in-situ tensile tests and HEXRD. The experimentallyobtained results served as input for computationmodeling using differ-ent approaches. Building up on this work, Cheng et al. [39] demon-strated computational material design of Q&P steels based on aparametric study using the plastic instability criterion to estimate ulti-mate tensile strength and uniform elongation. Furthermore, 2-stepQ&P heat treatments with and without a previous hot stamping stepwere compared by a combination of numerical modeling and differentexperimental techniques, including in-situ HEXRD to gather informa-tion on lattice parameter changes during partitioning [40].

Since the correlation between process parameters, microstructureand mechanical properties is of high importance for industrial applica-tions, the present work aims to link the information from in-situ exper-iments with those attained from mechanical testing. A combination ofdilatometry and high-energy X-ray diffraction was applied for an in-situ study of five different Q&P cycles. By this means, the influence ofvarying the quenching or partitioning temperature on the extent of aus-tenite decomposition was quantified. An identification of the decompo-sition product as tempered martensite, bainite or fresh martensite wasaccomplished by taking into account the heat treatment stage at

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which the transformation was observed. Furthermore, the carbonpartitioning kinetics and the final carbon content in the austenitecould be compared by the dilatation of the austenite lattice parameter.Based on the microstructural information obtained by HEXRD, the me-chanical properties derived from tensile testing of samples subjectedto the investigated heat treatments are discussed. This enabled an en-hanced understanding of the relationship between process parametersand resulting material properties, which finally helps to define guide-lines for the implementation of Q&P heat treatments on an industrialscale.

2. Experimental

The investigated steel is a 0.2C/1-1.5Si/2.2-2.7Mn (wt%) alloy with amartensite start temperature (Ms) of 339 °C as determined by dilatom-etry. The material was provided in the form of cold rolled sheets. Sam-ples of 10 × 4 × 1.19 mm3 were machined with the longest axesperpendicular to the rolling direction for dilatometer experiments.

In-situ HEXRD measurements were carried out at the P07 high-energy materials science (HEMS) beamline of HZG at PETRA III, DESY,Germany [41]. The experiments were performed in transmission geom-etry using monochromatic synchrotron X-ray radiation with a photonenergy of 100 keV (λ=0.124 Å). LaB6 was used as calibration standard.Diffraction patterns were acquired continuously during the in-situ heattreatments by a Perkin Elmer XRD1621 flat panel detector positioned ata distance of 1.5 m behind the sample. This allowed the detection of fullDebye-Scherrer rings up to amaximum 2Theta angle of 7.5°. An exposi-tion time of 1 s for each frame was used during austenitization and theend of the heat treatments, and 0.5 s for the quenching section and thebeginning of partitioning.

For the heat treatments, a dilatometer DIL 805A/D from TA Instru-mentswas placed in the beamusingAl2O3 rods and a type S thermocou-ple for temperature control. Both 1-step and 2-step Q&P heattreatments were performed. All samples were slowly heated to 850 °C,cooled to 750 °C, and subsequently quenched to the quenching temper-ature Tq. At the beginning of the quenching step, the reached tempera-tures were higher than the target values (up to 10 °C) and successivelyapproached the given Tq. By the end of the quenching step, the obtainedtemperatures were no higher than 4 °C compared to the given Tq. Afterthe quenching step of 3 s, the samples were reheated and partitioningwas executed for 300 s at varying partitioning temperatures Tp. All inall, five different heat treatments were investigated with different heattreatment parameters summarized in Table 1.

The recorded Debye-Scherrer rings were azimuthally integratedusing the DAWN software package [42,43]. The obtained 1D diffractiondata was analyzed by a Rietveld refinement procedure implemented inthe commercial software TOPAS from Bruker AXS using a fundamentalparameters approach. The instrumental contribution to the line profileshapes was determined by fitting the pattern of the LaB6 reference sam-ple, which ideally does not contribute to peak broadening. The obtainedparameters were then fixed and used for the instrumental function ofthe actual fitting procedure. The refinement further included back-ground, zero displacement, scale factors, unit cell parameters, tempera-ture factors and texture parameters. Strain and size broadening effectswere considered by the double-Voigt approach [44]. The quality of thefit was evaluated from the resulting difference curve. The fit of the 1D

Table 1Overview of the heat treatments conducted in this work.

Sample ID Tq/°C tq/s Tp/°C tp/s

Reference 260/360260

3

260

300Variation Tp

260/260 360260/400 400

Variation Tq230/360 230

360290/360 290

diffraction pattern recorded at the end of the reference cycle (260/360) is exemplary shown in Fig. 1.

Peaks corresponding to austenitic or ferritic (i.e. ferrite/martensite/bainite) phase were unequivocally identified. There was no evidenceof any additional phases occurring during the heat treatments. Austen-ite was fitted by a face-centered cubic cell (fcc, Fm-3m space group). Ithas to be noted that possible peak asymmetries caused by different car-bon contents in the austenitewere not taken into account during refine-ment. For fitting the ferritic phase, a body-centered tetragonal unit cell(bct, I4/mmm space group) was considered. This is reasonable, sincethe vast amount of this phase fraction is martensitic.

The evolution of the phase fractions as well as the change of latticeparameters were investigated. Considering possible inaccuracies of therefinement process, themaximumerrors of the received phase fractionsare assumed to be about 1%, which corresponds to typical values ob-tained for Rietveld refinements. The determination of the lattice param-eters is more accurate and the standard deviation calculated by theTOPAS software is not higher than 8 × 10−5 Å.

The carbon content in austenite can be estimated from the austenitelattice parameter aγ. This requires a prior subtraction of the thermalcontribution, since themeasurements were conducted at elevated tem-peratures. The coefficient of thermal expansion (CTE) was determinedusing the exponential temperature dependence proposed by vanBohemen [45]:

CTE ¼ B � e−θT; ð1Þ

where B is denoted as the CTE in the high temperature limit and θ as theDebye temperature of austenite. Allain et al. [36] calibrated CTE by a fitof the linear regime of aγ during cooling from 900 to 300 °C, where thealloy was fully austenitic. Additional ex-situ dilatometric experimentswere conducted in the present work, because the formation of smallamounts of a ferritic phase prior to the martensitic formation couldnot be entirely excluded and might affect the calibration results. Sam-ples were slowly heated (1 °C/s) to 950 °C and rapidly quenched toroom temperature with a cooling rate of 70 °C/s in a dilatometerDIL805 A from TA Instruments using SiO2 rods and a type S thermocou-ple for temperature control. The coefficients (B= 2.44 × 10−5 K−1, θ=230 K)were obtained by a fit of the dilatation curves in the range of 400to 800 °C. Subsequently, these coefficients were used for the calculationof the austenite lattice parameter (termed aγtherm), which was thensubtracted from the measured aγ. The resulting dilatation of the ther-mally corrected aγ at the end of the partitioning stepwas used for an es-timation of the final carbon enrichment in the austenite via the integralform of the correlation between aγ and carbon content suggested by

Fig. 1. Comparison of the measured (black dots) and calculated (blue line) diffractionpattern after Rietveld refinement of the diffraction pattern recorded at the end of thereference cycle (260/360). The resulting differential curve is plotted as red line at thebottom. Peaks corresponding to fcc (γ) and bct (α′) phase are indicated.

4 S. Ebner et al. / Materials and Design 178 (2019) 107862

Allain et al. [34]. By this means, the effect of internal stresses on aγ thatare found to arise during the final cooling is ruled out [35].

Additionally, steel stripswere subjected to theheat treatments givenin Table 1 on an annealing simulator. For these samples, the fraction ofretained austenite was evaluated magnetically with an accuracy of0.1% [46]. Flat tensile test samples with a gauge length of 25 mmweremachined from the heat treated strips and tested on a tensile testingmachine BETA 250 from Messphysik. Two samples for each heat treat-ment were tested. The procedure was similar to the one described in[14].

Fig. 3.Comparison of the temporal evolution of the bct phase fraction for thedifferent heattreatments after the quenching step (t=0 s). Themarked increase of the 260/260 curve atthe end of the heat treatment indicates the formation of fresh martensite.

3. Results

3.1. Phase fractions

The temperature profile of the 260/360 reference heat treatmentafter the onset of cooling from 850 °C can be seen in the upper part ofFig. 2. The lower part of Fig. 2 shows the corresponding evolution ofthe fcc and bct phase. The different heat treatment sections are dividedby vertical lines. A small amount of ferritic phase (b2%) remains at 850°C that slightly increases during the cooling section (I–III), beforereaching Ms. This amount does not exceed 5% for any heat treatment.Below Ms, the fraction of bct phase increases rapidly due to martensitetransformation. By the end of the quenching step (IV), values of 72%,83% and 87% are obtained for the bct phase after quenching to the re-spective Tq of 290, 260 and 230 °C. Since this bct phase fraction corre-sponds to athermal martensite, which is tempered during the ongoingheat treatment, it is designated as temperedmartensite. A further, albeitsmaller increase of the bct phase is observed during the subsequentreheating (V) and partitioning (VI) step. Due to the strong evidence ofthe bainitic nature of this phase mentioned in the introduction part,this fraction is stated as bainite. The bct phase that possibly formsduringthe final cooling (VII) is defined as fresh martensite.

Fig. 3 shows the temporal evolution of the bct volume fraction afterthe quenching step for all heat treatments. A fast increase in the first25 s is observed for all heat treatments that levels off with prolongedholding. The amount of bct phase increases with higher Tq from 5%(230/360) to a significantly higher value of 17% (290/360). By elevatingTp to 400 °C (260/400), a steeper increase is observed at the beginning,but the curveflattens faster and reaches slightly lower values (6%) com-pared to the reference cycle. During partitioning of the 1-step cycle(260/260), about 5% volume fraction bainite is formed. This is also theonly cycle were the formation of a small amount of fresh martensite

Fig. 2.Overview of the reference Q&P heat treatment (260/360) starting from the onset ofcooling after austenitization. The temperature profile is shown in the upper part and theevolution of fcc and bct phase fraction in the lower part of the graph. The different heattreatment sections are divided by solid vertical lines. The beginning of martensiteformation is marked by a dashed line.

was detected during the final cooling step, which can be seen as in-crease of the 260/260 curve (encircled area in Fig. 3).

For a better comparability of the extent of austenite decompositionafter quenching, the austenite phase fractions fγ at specific points (i.e.end of quenching, start and end of partitioning, and end of heat treat-ment) are listed in Table 2. The austenite phase fraction decreases byabout 2% for all cycles during reheating. Concerning partitioning andfinal cooling, the most significant decrease of austenite fraction from28% to 11% is observed for the highest Tq (290/360). This correspondsto a ‘loss’ of austenite of 61%. In comparison, the ‘loss’ of austenite forthe other heat treatments ranges between 28 and 47%.

The final phase fractions are summarized in Fig. 4. As mentioned be-fore, the bct phasewas divided into temperedmartensite formed duringthe initial quench, bainite formed during reheating and partitioning,and fresh martensite formed during the final quench. Depending onthe heat treatment, the achieved austenite fractions range between 8and 13% and bainite fractions from 5 to 17%. As can be seen, higher Tqresults in a lower amount of temperedmartensite and a higher amountof bainite, while the austenite fraction marginally increases. Comparedto the reference heat treatment, both higher and lower Tp lean towardsmore austenite and slightly less bainite.

3.2. Lattice parameters and austenite carbon content

To calculate the carbon content, the measured austenite lattice pa-rameter aγ was corrected for the contribution of thermal expansion ac-cording to the procedure described in the experimental part. For the260/360 cycle, a comparison of aγ with the calculated lattice parameteraγtherm can be seen in Fig. 5a together with the resulting differentialvalue adiff. The enlarged section, marked by a red rectangle, is shownin Fig. 5b. A good agreement of aγ and aγtherm above Ms is obtained.Below Ms, aγ deviates from the calculated line as can be seen as asmall increase of adiff (encircled area in Fig. 5a). Subsequently, aγ is

Table 2Volume fractions of austenite (%) determined at specific points during the heattreatments.

Tq End Tp Start Tp End End

260/360 17 15 9 9260/260 18 – 13 13260/400 16 14 10 10230/360 13 11 8 8290/360 28 26 11 11

Fig. 4. Summary of the phase fractions (%) obtained for the different Q&P cycles in thepresent work. Note that “Bainite” indicates the amount of bct phase formed duringreheating and partitioning.

Fig. 6. Evolution of adiff (measured austenite lattice parameter aγ minus calculatedaustenite lattice parameter aγtherm) after the end of the quenching step (t = 0 s). Notethat the final values cannot be compared directly, since the curves do not have the samestarting point.

5S. Ebner et al. / Materials and Design 178 (2019) 107862

lower than aγtherm at Tq and during the reheating step. A slight increaseof the slope during reheating is observed above ~300 °C, which is indi-cated by the red lines in Fig. 5b. During partitioning, aγ significantly in-creases while aytherm stays constant. A similar behavior was alsoobserved for the other heat treatments.

For a comparison of the transformation kinetics, the determined adifffor all heat treatments is shown in Fig. 6. For the 1-step cycle (260/260),the change of adiff is negligible compared to the significant increase thatis observed for the other heat treatments. By the end of the partitioningstep, an increase between 2.10 and 2.37 × 10−2 Å is attained for all 2-step cycles, where the increase is the fastest for 260/400. During finalcooling, adiff further increases by about 0.26 × 10−2 Å.

The increase of adiff from reheating to the end of the partitioning step(Δaγ) was used for the estimation of the total carbon enrichmentΔCγ inthe austenite. The results are summarized in Table 3. Except for the 260/260 cycle with a value below 0.1 wt%, there is a clear increase of carboncontent around 0.66 wt% for the other heat treatments.

It was possible to investigate the reduction of tetragonality duringthe heat treatments by using a bct phase for the fit of the ferriticpeaks. It was found that the lattice parameter aα remains constant,

Fig. 5. (a) Comparison of the measured austenite lattice parameter aγ with the calculated audifference adiff is presented in the lower part of the graph. The end of austenitization corresponred lines serve to improve the comparability of the different slopes.

while the lattice parameter cα constantly decreases after quenching.This behavior is shown as c/a ratio in Fig. 7. The decrease of c/a ismost pronounced in the first 25 s and for higher Tp.

3.3. Mechanical properties

A summary of themechanical properties and retained austenite frac-tions is given in Table 4, and the corresponding engineering stress-strain curves can be seen in Fig. 8a. For the 2-step cycles, yield strength(YS) is around 1100 MPa and ultimate tensile strength (UTS) is around1250 MPa. The 260/260 engineering stress-strain curve clearly differsfrom the others with lower YS (879 MPa) and higher UTS (1474 MPa).The achieved uniform elongation (UE) varies between 4.5 and 7.6%and the total elongation (TE) between 12.3 and 15.1%. Except for the260/260 cycle, the trends regarding the austenite fraction determinedmagnetically after the heat treatments are consistent with the resultsobtained by HEXRD (Fig. 4). With increasing Tq, higher amounts ofretained austenite are achieved. At the same time, YS and UTS slightlydecrease and both UE and TE increase. Higher Tp also result in higheraustenite fractions. UTS decreases, while YS, UE and TE increase withhigher Tp.

stenite lattice parameter aγtherm for the 260/360 reference heat treatment. The resultingds to t = 0 s. (b) Enlarged section of (a) showing the quenching and reheating step. The

Table 3Summary of the increase of adiff between reheating and the end of partitioning, and theresulting increase of carbon content in the austenite.

Δadiff/10−2 Å ΔCγ/wt%

260/360 2.13 0.65260/260 0.25 0.08260/400 2.32 0.70230/360 2.10 0.64290/360 2.19 0.66

Table 4Summary of the mechanical properties and the corresponding retained austenite fractionfγ (%). Note that the given retained austenite fractionswere determinedmagnetically afterthe heat treatments. The values of the mechanical properties correspond to the mean oftwo samples tested.

fγ/% YS/MPa UTS/MPa UE/% TE/%

260/360 10.5 1101 ± 9.2 1283 ± 1.4 6.1 ± 0.0 13.8 ± 0.1260/260 9.6 879 ± 3.5 1474 ± 8.5 6.7 ± 0.4 12.3 ± 0.1260/400 11.6 1113 ± 9.2 1245 ± 3.5 7.6 ± 0.4 15.1 ± 0.1230/360 9.4 1155 ± 33.2 1317 ± 0.0 4.5 ± 0.4 12.3 ± 0.6290/360 11.1 1055 ± 7.1 1270 ± 3.5 6.7 ± 0.4 14.1 ± 0.2

6 S. Ebner et al. / Materials and Design 178 (2019) 107862

The instantaneouswork hardening exponent ndiff is shown in Fig. 8b.The discontinuities seen at 0.015 logarithmic strain are due to the strainrate change during tensile testing. Significant work hardening at higherstrains is observed for all 2-step heat treatments. In contrast, higher ini-tial n-values are obtained for the 1-step cycle (260/260) that rapidly de-cline with increasing strain.

4. Discussion

4.1. Austenite decomposition during heat treatment

Austenite decomposition during subsequent reheating andpartitioning was found to occur for all heat treatments, as can be seenin Fig. 3. Since no additional peaks indicating carbide formation weredetected, this decomposition can be attributed to the formation of bai-nite. As mentioned earlier, there is strong evidence that the isothermaltransformation product below Ms is bainite [27,28]. Contradictory re-sults obtained by Somani et al. [25] and Kim et al. [29] were discussedin a very convincing way by Navarro-López et al. [30]: The observedprotrusions of themartensitic featureswere attributed to bainitic ferritegrowing from the athermal martensite and resulting in wavy bound-aries. The present results approve that austenite decomposition doesoccur for the 1-step (260/260) heat treatment, albeit to a smaller extentcompared to the reference cycle (Fig. 4).

In the case of 2-step cycles, where partitioningwas performed aboveMs, the transformation product is unambiguously identified as bainite.Considering the T0 concept that describes bainite formation [47], themain influences that determine the amount of bainite can be summa-rized as (1) the initial austenite fraction, (2) the carbon content in aus-tenite and (3) the partitioning parameters (time and temperature).Higher Tq results in a higher austenite fraction (see Table 2) thatneeds to be stabilized. This higher austenite content also gives rise to

Fig. 7. Evolution of the c/a ratio of the ferritic phase during heat treatment. The end of thequenching step is taken as t = 0 s.

more bainite formation, as can be seen in Fig. 4, which is consistentwith the results reported in [21]. Consequently, only a slightly higherfinal austenite fraction compared to the reference cycle was obtainedfor the 290/360 heat treatment, although the amount of austeniteafter the partitioning step was significantly higher (Table 2). Regardingdifferent Tp, carbon diffusion is faster with increasing temperature andthe critical carbon content in austenite, where bainite formation stops,is reached earlier. Thus, slightly less bainite is expected to form at 400°C compared to 360 °C. This trend is indeed reflected in Fig. 4, althoughthis difference is negligible when considering the inaccuracies of the re-finement process of about 1%.

During the final cooling, insufficiently stabilized austenite trans-forms into fresh martensite, as can be seen for the 260/260 cycle inFig. 3. At the end of the heat treatments, the detection of diffraction pat-terns ended before the specimen temperature reached room tempera-ture (30–50 °C). Therefore, a possible martensite formation uponcoolingwas not entirely recorded andmay be underestimated.Withouttaking into account the 1-step cycle, the final austenite fractions do notdiffer largely (Table 2), regardless of the chosen heat treatmentparameters.

4.2. Change of austenite lattice parameters and estimated carbon content

Several factors influence the change of the austenite lattice parame-ter, namely temperature, stress state and chemical composition. Thethermal contribution was subtracted considering an exponential tem-perature dependence. The influence of prior induced internal stressesas well as the effect of stress relaxation during partitioning are nottaken into account. The difference between aγ and aγtherm shown inFig. 5 is similar to that observed by Allain et al. [36]. According to theirwork, the onset of martensite transformation introduces hydrostaticstresses. This leads to a deviation of aγ from the calculated curvebelowMs and to values lower than predicted at Tq. During the followingreheating step, the calculation overestimates the increase of aγ, sincethe austenite expansion is constrained by the surrounding matrix,which is not considered in the CTE calculation. The following significantincrease of aγ is then accounted to the carbon enrichment of the austen-ite. The effect of stress relaxation is considered too small, and thus can-not be responsible for this increase. Finally, the deviation during cooling,which is better seen as increase of adiff in Fig. 6, is attributed to the intro-duction of final hydrostatic stresses in the austenite. The authors of thepresentwork agreewith these explanations for themost part. However,an alternative explanation for the deviation observed immediately aftermartensite transformation is offered. As reported by Epp et al. [48], car-bon diffusion starts immediately after martensite transformation. Thismay retard the decrease of aγ and lead to the observed deviation,which can be seen as small increase of adiff (encircled area in Fig. 5a). De-pending on Tq, the measured increase of adiff is between 0.18 and 0.35× 10−2 Å, where the lower values are attained at higher Tq. This changeof adiff corresponds to a carbon increase in the austenite of 0.05–0.11wt% immediately after martensite formation.

Fig. 8. (a) Engineering stress-strain curves and (b)work hardening behavior. Note that the discontinuities at 0.015 logarithmic strain seen in (b) are caused by the change of the strain rateduring tensile testing.

7S. Ebner et al. / Materials and Design 178 (2019) 107862

The negligible carbon diffusion in austenite at low temperaturescan be seen by the small changes during partitioning of the 1-stepQ&P heat treatment (Fig. 6) that lead to an estimated carbon enrich-ment of only 0.08 wt%. Thus, a poor austenite stabilization isexpected. This is also emphasized by the detected formation offresh martensite upon final cooling seen in Fig. 3. Although carbonenrichment is extremely low, a certain part of austenite remains bythe end of the heat treatment that is the highest of all cycles. Sincepattern recording stopped before room temperature was reached,this is mainly attributed to further, though not detected martensiteformation at the end. This is also approved by the distinct values ofthe retained austenite fraction determined by HEXRD (13%,Table 2) and by magnetic measurements (9.6%, Table 4) for the260/260 cycle.

In contrast, carbon enrichment was clearly visible for all 2-step Q&Pcycles. The change in slope during the reheating step above 300 °C(Fig. 5b) underlines the onset of effective carbon diffusion into the aus-tenite [49]. In accordance with faster diffusion kinetics, the increase ofadiff during partitioning is fastest for 260/400 and from the flatteningof the curve it can be assumed that carbon redistribution is finishedwithin 300 s of partitioning time. Regarding the estimated carbon con-tent of austenite at the end of the partitioning step, similar results areobtained for the 2-step cycles (Table 3). Allain et al. [36] pointed outthat the carbon content derived from the lattice parameter stronglydepends on the selected methodology, i.e. the choice of the temper-ature dependent CTE function and whether effects induced by inter-nal stresses are considered or not. Consequently, small variationsbetween the carbon content obtained for the different 2-step cycleswere indeed detected (Table 3), but precise conclusions concerningthe influence of the process parameters cannot be drawn. Apartfrom this, the estimated carbon increase of 0.64–0.70 wt% after the2-step Q&P cycles leads to a final carbon content of 0.84–0.90 wt%in the austenite for the investigated steel. These results representmean carbon values, since measurements by atom probe tomogra-phy showed that the carbon content can significantly vary betweendifferent austenite regions [50–52].

4.3. Effects on the mechanical properties

According to the HEXRD experiments, three different phases can beclearly identified during all heat treatments investigated in this work:(1) martensite formed during initial quenching, (2) austenite with suf-ficient stability for existence at room temperature, and (3) bainiteformed during reheating and partitioning. In the case of the 1-stepcycle (260/260), small amounts of fresh martensite formed during the

final quench were also detected. In addition to the phase fractions(Fig. 4), it was possible to estimate the decrease of tetragonality of themartensite by using a bct cell to fit the ferritic diffraction peaks, as canbe seen in Fig. 7. The loss of tetragonality with prolonged holding canbe attributed to tempering effects, which are the partitioning of carbonto energetically favorable sites (dislocations, grain boundaries, etc.) orinto the austenite and the decrease of dislocation density. Further infor-mation about the carbon enrichment into the austenite was also ob-tained by the dilatation of aγ (see e.g. Fig. 6). Based on these findings,the mechanical properties of the resulting multiphase microstructuresand the contribution of the individual phases are discussed in thefollowing.

Comparing the reference cycle (260/360) and the 1-step Q&P cycle(260/260), a similar amount of tempered martensite was observed, ascan be seen in Fig. 4. As expected, the degree of tempering is signifi-cantly lower for 260/260, which was revealed by the c/a ratio shownin Fig. 7. Since UTS is primarily affected by the amount and strength ofthe dominating phase, i.e. tempered martensite in this case, the about190MPa lower value obtained by 260/360 can be explained by strongermartensite tempering. The difference between the bainite fractionformed during partitioning is low (2%), and thus the effect on the me-chanical properties is assumed to be negligible. Furthermore, the mag-netic measurements showed a lower austenite fraction for 260/260. Asdiscussed in the previous section, this austenite fraction is assumed tohave a lower chemical stability compared to the 2-step cycles due tominimal carbon diffusion, which was derived from the minor increaseof adiff (see Fig. 6). Austenite stability is not only affected by the chemicalcomposition, but also by other factors, including size, morphology, andsurrounding matrix. Hidalgo et al. [12] showed the beneficial effect ofa low-tempered martensitic matrix on the stabilization of austenite.For 260/260, the less tempered martensite certainly accounts for a me-chanical stabilization of the austenite. However, the detected formationof freshmartensite showed that thismight be less effective compared tothe chemical stabilization by carbon enrichment observed for the 2-stepcycles. Based on these findings, it seems likely that the mechanical sta-bility for 260/260 might not be sufficient to enhance ductility by theTRIP-effect. This assumption is supported by a recent study of the pres-ent authors [14],where the influence of Q&P heat treatment parameterson themechanical properties of a steelwith similar compositionwas in-vestigated. For the 1-step cycle conducted at 230 °C, it was found thatthe retained austenite fraction continuously increased with prolongedholding time up to 600 s, but neither UE nor TEwere positively affectedby this increase. This was accounted to a poormechanical stability and afast austenite to martensite transformation in the early stages of exter-nal loading, which may also be the case in the present work. As a

8 S. Ebner et al. / Materials and Design 178 (2019) 107862

consequence of the poor austenite stability and the early formation offresh martensite, the obtained YS for the 260/260 condition is lowercompared to the 260/360 reference cycle. Similar results can be alsofound in literature [5,53].

As can be seen in Fig. 8b, the work hardening behavior of the 1-stepcycle clearly differs from the reference cycle, and also from the other 2-step cycles. An explanation of the work hardening behavior observed inQ&P steels can be found, for example, in a recent study by Findley et al.[54]. At lower strains, thework hardening rate is dominated by themar-tensitic matrix and its dislocation density, which can be controlled bytempering. With increasing strain, this contribution diminishes as thedislocation density reaches saturation, and consequently the workhardening rate decreases. This decrease can be mitigated by themechanically-induced austenite to martensite transformation. Strongertempering, as it was the case for the 2-step cycles, reduces the disloca-tion density of the temperedmartensite, which results in a smaller con-tribution to the initial work hardening rate and to the remarkabledifference between 1-step and 2-step Q&P. From the continuously in-creasing work hardening curve seen above 0.015 logarithmic strain forthe 260/360 cycle – and also for the other 2-step cycles – in Fig. 8b, itcan be concluded that the austenite to martensite transformation effec-tively contributes to an enhanced work hardening behavior at higherstrains. Hence, an adequate stabilization of the austenite was obtainedby the 2-step processing, which was mainly accomplished by a carbonenrichment into the austenite around 0.66 wt% (Table 3). Nevertheless,further studies, e.g. interrupted tensile tests, are needed in order to in-vestigate the austenite stability and the contribution of austenite andmartensite to the work hardening behavior, especially in the case ofthe 1-step Q&P cycle.

Both the 260/400 and the 230/360 cycles can be compared to the260/360 reference cycle in a similar way. For 260/400, a fraction of tem-pered martensite and bainite similar to the reference cycle (260/360)was obtained (Fig. 4), and the amount of retained austenite marginallyincreased with increasing Tp. The higher degree of tempering at 400 °C,that can be seen from the stronger decline of the c/a ratio in Fig. 7, re-sulted in a slight decrease of UTS, while YS is only marginally affected.Furthermore, this resulted in a decline of the initial work hardeningrate seen in Fig. 8. The increase of UE and TE can be explained by thehigher austenite fraction. In the case of the 230/360 cycle, the initialquench to a lower Tq resulted in a higher amount of temperedmartens-ite (Fig. 4) with a similar degree of tempering (Fig. 7). Both bainite andaustenite fraction are slightly lower. As can be seen in Table 4, the higheramount of tempered martensite led to an increase of UTS and – to alesser extent – YS. The lower austenite fraction resulted in lower UEand TE.

In a comparison of the 260/360 reference cycle with 260/260, 260/400 and 230/360, the influence of bainite on the mechanical propertiesis supposed to be negligible, since the amount of bainite formed duringpartitioning is low (5–7%, Fig. 4). In the case of 290/360, 17% bainitewasdetected and the influence of bainite formation can no longer beneglected. The amount of austenite is comparable to the values obtainedby 260/360, and also by 260/400. Hence, the desired effect of higher Tqto increase the austenite fraction, and thus the ductility, was nullified bythe increasing formation of bainite. Interestingly, the ductility does notseem to be affected, although the proportion of tempered martensitewas partially replaced by bainite. As can be seen from Table 4, the in-crease of UE and TE correlates with the increase of austenite. There isalso no clear impact on UTS and YS. Obviously, the mechanical proper-ties of the bainite that forms form the carbon-enriched austenite duringpartitioning are similar to that of the carbon-depleted tempered mar-tensite. Hence, increasing bainite formation at higher Tq showed nobeneficial or detrimental effect on the tensile behavior. Nevertheless,it is unwanted, because austenite is consumed and themaximum reach-able austenite fraction is limited. The issue of bainite formation at higherTq might be avoided by increasing Tp to fasten the carbon diffusion andstabilization of the austenite.

5. Conclusions

The transformation kinetics offive different Q&Pheat treatments, in-cluding four 2-step and one 1-step cycle, were investigated by in-situHEXRD experiments, and the mechanical properties were determinedby tensile testing. The main findings can be summarized as follows:

• For all heat treatments, austenite decomposition during partitioningoccurred and was attributed to bainite formation. The highest bai-nite fraction of 17% was attained after quenching to 290 °C andpartitioning at 360 °C. The fractions of the other heat treatmentsranged from 5 to 7%. Independent of the process parameters, theobtained austenite fractions vary in a narrow range between 9and 13%.

• Carbon partitioning into the austenite was found to start immedi-ately after martensite transformation. For the 2-step cycles, thedilatation of the austenite lattice parameter during partitioningclearly revealed carbon diffusion into the austenite and similarcarbon contents around 0.66 wt% at the end of partitioning. Incontrast, the diffusion kinetics were too slow in the case of the 1-step heat treatment to achieve a comparable carbon enrichment.

• The dependence of the mechanical properties on the heat treat-ment parameters was primarily explained by the tempering de-gree of initially formed martensite and the amount and stabilityof austenite. As a consequence of the effective chemical stabiliza-tion of the austenite, enhanced work hardening at higher strainswas attained for the 2-step Q&P conditions. In contrast, the lowtempered martensite after 1-step processing resulted in a high ini-tial work hardening rate and it is assumed that there is no contri-bution of the austenite to the work hardening at higher strainsdue to a poor mechanical stabilization.

• At higher Tq, the relative amount of bainite clearly increased, whilethe austenite fraction was only slightly higher. Despite differentbainite fractions, the mechanical properties of the heat treatmentswith similar austenite fractions did not differ greatly. This was ex-plained by comparable mechanical characteristics of temperedmartensite and bainite.

CRediT authorship contribution statement

Sandra Ebner: Conceptualization,Methodology, Software, Validation,Investigation, Data curation, Writing - original draft, Visualization. Clem-ens Suppan: Conceptualization, Investigation, Resources, Writing -review & editing. Andreas Stark:Methodology, Software, Validation, In-vestigation, Resources, Data curation, Writing - review & editing. RonaldSchnitzer: Conceptualization, Writing - review & editing, Supervision.Christina Hofer: Conceptualization, Investigation, Writing - review &editing, Supervision.

Declaration of Competing Interest

None.

Acknowledgments

Funding of the Austrian BMVIT (846933) in the framework of theprogram “Production of the future” and the “BMVIT Professorship for In-dustry” is gratefully acknowledged. We acknowledge DESY (Hamburg,Germany), a member of the Helmholtz Association HGF, for the provi-sion of the experimental facilities. Parts of this research were carriedout at PETRA III and we would like to thank Dr. Nobert Schell for assis-tance at the P07 high-energy materials science (HEMS) beamline. Spe-cial thanks also to Dr. Andreas Landefeld for assistance in conductingthe experiments. In addition, the authors thank Dr. Michael Tkadletz,Dr. Christian Saringer and DI Dominik Nöger for the support regardingdata evaluation.

9S. Ebner et al. / Materials and Design 178 (2019) 107862

Data availability

The raw/processed data required to reproduce these findings cannotbe shared at this time as the data also forms part of an ongoing study.

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