Time-dependent synchrotron X-ray diffraction on the austenite … synchrotron X... · 2013. 2....

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Acta Materialia xxx (2012) xxx–xxx

Time-dependent synchrotron X-ray diffraction on theaustenite decomposition kinetics in SAE 52100 bearing steel

at elevated temperatures under tensile stress

E. Jimenez-Melero a,b,⇑, R. Blonde a,c, M.Y. Sherif d, V. Honkimaki e, N.H. van Dijk a

a Fundamental Aspects of Materials and Energy, Faculty of Applied Sciences, Delft University of Technology, Mekelweg 15, 2629 JB Delft, The Netherlandsb Dalton Cumbrian Facility, University of Manchester, Westlakes Science & Technology Park, Moor Row, Cumbria CA24 3HA, UK

c Materials Innovation Institute, Mekelweg 2, 2628 CD Delft, The Netherlandsd SKF Engineering & Research Centre, Kelvinbaan 16, 3439 MT Nieuwegein, The Netherlands

e European Synchrotron Radiation Facility, 6 Rue Jules Horowitz, BP 220, 38043 Grenoble Cedex, France

Received 25 April 2012; accepted 26 October 2012

Abstract

We have studied the decomposition kinetics of the metastable austenite phase present in quenched-and-tempered SAE 52100 steel byin situ high-energy synchrotron X-ray diffraction experiments at elevated temperatures of 200–235 �C under a constant tensile stress. Wehave observed a continuous decomposition of austenite into ferrite and cementite. The decomposition kinetics is controlled by the long-range diffusion of carbon atoms into the austenite ahead of the moving austenite/ferrite interface. The presence of a tensile stress of295 MPa favours the carbon diffusion in the remaining austenite, so that the activation energy for the overall process decreases from138–148 to 82–104 kJ mol�1. Before the austenite starts to decompose, a significant amount of carbon atoms partition from the sur-rounding martensite phase into the metastable austenite grains. This carbon partitioning takes place simultaneously with the carbideprecipitation due to the over-tempering of the martensite phase. As the austenite decomposition proceeds gradually at a constant tem-perature and stress, the elastic strain in the remaining austenite grains continuously decreases. Consequently, the remaining austenitegrains act as a reinforcement of the ferritic matrix at longer isothermal holding times. The texture evolution in the constituent phasesreflects both significant grain rotations and crystal orientation relationships between the parent austenite phase and the newly formedferritic grains.� 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: SAE 52100 steel; Austenite; Phase transformation kinetics; Carbon diffusion; Synchrotron diffraction

1. Introduction

Advanced engineering components require an accuratecontrol of their shape and dimensions during the manufac-turing process and their service life [1,2]. At the end of theindustrial production chain, many components are sub-

1359-6454/$36.00 � 2012 Acta Materialia Inc. Published by Elsevier Ltd. All

http://dx.doi.org/10.1016/j.actamat.2012.10.025

⇑ Corresponding author at: Dalton Cumbrian Facility, University ofManchester, Westlakes Science & Technology Park, Moor Row, CumbriaCA24 3HA, UK. Tel.: +44 (0)1946 508872.

E-mail address: [email protected] (E.Jimenez-Melero).

Please cite this article in press as: Jimenez-Melero E et al. Time-dependent sSAE 52100 bearing steel at elevated temperatures under tensile stress. Acta

jected to a multi-step heat treatment that aims to attainthe desired materials properties while still meeting the strictdimensional criteria imposed by the target application [3,4].However, their resultant mechanical and dimensional char-acteristics after production will gradually deteriorate overtime during the service life, due to the susceptibility ofthe microstructure to external thermal and mechanicalstimuli [5–7]. This is the case in ball and roller bearings thatare nowadays used as key precision components in a broadscope of machinery and gearboxes [8]. During their servicelife, bearings are subjected to complex multi-axial stressstates that change periodically in combination with

rights reserved.

ynchrotron X-ray diffraction on the austenite decomposition kinetics inMater (2012), http://dx.doi.org/10.1016/j.actamat.2012.10.025


mailto:[email protected]



2 E. Jimenez-Melero et al. / Acta Materialia xxx (2012) xxx–xxx

relatively high operating temperatures. A high accuracy ofall the internal bearing dimensions and a good surface fin-ish quality are mandatory requirements in order to mini-mize rotational friction effects and the associated heatdissipation [9,10].

Through-hardened steel SAE 52100 (DIN 100Cr6) iswidely used in bearing applications due to its high strengthand resistance to rolling contact fatigue. This material isrelatively easy to machine in the spheroidize-annealed con-dition. A subsequent heat treatment process of austenitiza-tion, quenching and tempering leads to a complexmultiphase microstructure containing either bainitic-ferriteor martensite, finely dispersed e/g transition carbides,cementite and retained austenite [11]. A significant amountof austenite (>10 vol.%) is typically retained after temper-ing in the room-temperature microstructure due to a com-bined effect of the carbon content, grain size and residual(compressive) stresses [12,13]. An additional carbonitridingprocess may be added during the austenitization stage inorder to increase the amount of retained austenite(>20 vol.%) in the (sub-) surface region of the bearing com-ponent to attain an extended service life [14]. During oper-ation, the bearing must sustain relatively large contactstresses (>3 GPa) during the extremely high number ofcycles (>108) that are ascribed to the phenomenon of roll-ing contact fatigue [13,15]. As a consequence, relativelycomplex microstructural changes develop in the bearingover time. These changes involve the decay of martensiteinto ferrite, the decomposition of the metastable austeniteand changes in texture and in the residual carbides [9]. Inthe case of good clean lubrication conditions and smoothsurfaces, the main failure mechanism in bearings is the sub-surface initiated fatigue spall that results in metallic parti-cles flaking from the raceway surface [16,17]. Cracksinitiate in the sub-surface region at inclusions or carbideclusters and propagate to the surface of the raceway. Crackinitiation is preceded by microplastic deformation thatleads to localized damage. The ability to maintain an elas-tic response during cyclic loading seems to be critical toextend the bearing life. The development of a radial tensilestress and sharp changes in texture will promote the crackpropagation parallel to the rolling surface [18].

The presence of metastable austenite in the sub-surfaceregion of the bearing and its evolution with time are con-sidered as useful local probes of the microstructuralchanges occurring during rolling contact fatigue. It is there-fore a key parameter to assess the fatigue life of the bearing[19]. The decomposition of austenite is normally monitoredin situ: (a) during cyclic loading at room temperature usingconventional X-ray diffraction. In-depth studies normallyimply successive removal of surface layers by electropolish-

Table 1Chemical composition (in wt.%) of the SAE 52100 steel used in this study.

Cr C Mn Cu Si

1.44 0.96 0.34 0.22 0.20

Please cite this article in press as: Jimenez-Melero E et al. Time-dependent sSAE 52100 bearing steel at elevated temperatures under tensile stress. Act

ing or chemical etching [9,13,20]; (b) during heating at zerostress using a wide variety of experimental techniquesincluding dilatometry [21], calorimetric analysis [22], mag-netic [23] and thermoelectric power measurements [24]. Theaim of the present study is to assess the combined effect oftemperature and applied stress on the retained austenitedecomposition kinetics in SAE 52100 steels.

To achieve this goal, we have monitored in situ theaustenite decomposition process by time-dependent X-raydiffraction experiments at a synchrotron source. The useof X-rays of high energy (>50 keV) has allowed us to probethe austenite phase present in the bulk microstructure with-out any prior removal of the surface layer. The high flux ofthe synchrotron X-ray beam, together with the use of a fasttwo-dimensional (2-D) detector, provided the key to mon-itor the decomposition kinetics in real time. We haveselected temperatures that resemble the operation tempera-tures of bearings and the local temperature peaks generatedunder rolling conditions. We consider the application of aconstant tensile stress as a first step in understanding thematerial response during cyclic loading. Equivalent diffrac-tion experiments were carried out with and without anapplied stress in order to decouple the temperature andstress effect on the austenite decomposition kinetics.

2. Experimental

2.1. Sample preparation and metallographic investigations

The starting material consisted of a spheroidize-annealed bar with a diameter of 60 mm. The chemicalcomposition is given in Table 1. In this condition themicrostructure can be described as spheroidized cementiteparticles embedded in a ferritic matrix. Cylindrical dog-bone-shaped tensile specimens with a gauge length of10 mm and a diameter of 1 mm were machined from thestarting bar. The cylindrical axis of the samples wasselected to be parallel to the axis of the bar. In order tokeep track of the sample orientation a mark parallel tothe diameter of the original bar was made on the top partof the cylindrical samples. The samples were austenitizedfor 20 min in a salt bath furnace. Three different austeniti-zation temperatures (Taust = 860, 900 and 950 �C) wereused in order to generate microstructures with differentcharacteristics of the metastable austenite phase. The sam-ples were subsequently quenched in oil to 60 �C and held atthis temperature for 15 min. During this process, a signifi-cant amount of austenite is transformed into martensite.Finally, the samples were rinsed in cold water and tem-pered at 160 �C for 90 min in a hot air stream furnace.Additionally, one disc of spheroidize-annealed material

Mo Al P S Fe

0.05 0.011 0.007 0.002 Bal.

ynchrotron X-ray diffraction on the austenite decomposition kinetics ina Mater (2012), http://dx.doi.org/10.1016/j.actamat.2012.10.025


E. Jimenez-Melero et al. / Acta Materialia xxx (2012) xxx–xxx 3

was heat-treated at the same time as the tensile specimensfor each of the selected austenitization temperatures usingthe process described above. These discs were used to char-acterize the resultant microstructures using visible lightoptical microscopy. The metallographic samples werecold-mounted in a resin, mechanically ground and polisheddown to 1 lm with a diamond suspension, and finallyetched with a nital 1.5 vol.% solution.

2.2. High-energy X-ray diffraction experiment

We have studied the decomposition kinetics above roomtemperature for samples with Taust = 860, 900 and 950 �Cby means of in situ high-energy X-ray diffraction. Theexperiment was performed at the beam line ID15A of theEuropean Synchrotron Radiation Facility (Grenoble,France). Technical details about the beam line can befound in Ref. [25]. The samples were mounted on a 2 kNmicro-tensile rig placed on a Huber translation table thatallowed translations in three dimensions in space, togetherwith x-rotations around the sample cylindrical axis. Thesamples were aligned with the diameter of the originalbar parallel to the frame of the rig. The loading direction(LD) coincided with the cylindrical axis of the sampleand the raw bar material. The central part of the samplewas illuminated by an intense monochromatic X-ray beamwith an energy of 69.95 keV (k = 0.1772 A). We focusedthe beam down to 5 � 25 lm2 using parabolic focusinglenses. The diffracted intensity was collected on a 2-D Pix-ium 4700 detector placed behind the sample [26]. Thisdetector has a useful pixel array of 2640 � 1920 pixelsand a pixel size of 154 � 154 lm2. The sample was contin-uously rotated along its cylindrical axis in steps ofDx = 0.5�, covering a total angular rotation range from45 to 135� on both sides of the tensile rig. During eachx-step the diffracted intensity was continuously recordedon the Pixium detector while the sample was rotated at aconstant angular velocity. We selected an exposure timeof texp = 0.4 s for each diffraction pattern in order to avoidsaturation of the detector pixels. The diffraction patternswere collected for the mentioned x-scan at five sampleheights. The summation of the 2-D patterns recorded atfive different heights for a given step in x yielded an effec-tive beam size of 25 � 25 lm2. Each sample was character-ized at room temperature by recording the correspondingdiffraction patterns at zero stress and at a constant tensilestress of r = 295 MPa.

Once the starting microstructures were characterized atroom temperature, the samples were heated by cartridgeheaters mounted in the clamps to a selected temperature(Texp) in the range of 200–235 �C, and kept at a constanttemperature for 2–6 h to complete the decompositionprocess. During this process we collected the diffractionpatterns for the five sample heights continuously, coveringa total angular range from 85 to 95� in steps of Dx = 0.5�on only one side of the tensile rig. Each data point onthe austenite decomposition kinetics was obtained


approximately every 2 min and was the result of combining100 2-D diffraction patterns (20 patterns per sampleheight). The exposure time for each diffraction patternwas again 0.4 s. Once the Debye rings originating fromthe austenite phase have disappeared from the detector,the sample was cooled back to room temperature. For eachof the selected decomposition temperatures, two equivalentsamples were measured: one at zero stress and one with atensile stress of r = 295 MPa. In order to avoid oxidationof the sample surface, the tensile rig was covered with adome made of borosilicate glass (chemical composition ofSiO2: 80.6%, B2O3: 13.0%, Na2O: 4.0% and Al2O3: 2.3%)with a total thickness of 15 mm. The inner space was evac-uated to reach a vacuum of �10�5 mbar before we startedheating the sample. The scattering signal from the borosil-icate dome was also recorded before the steel sample wasmounted on the tensile rig.

2.3. Analysis of the diffraction data

The recorded dataset at room temperature comprises aseries of 2-D diffraction patterns as a function of the x-rotation and sample height, for selected combinations ofaustenitization temperature and applied stress for the stud-ied SAE 52100 steel samples. In addition to this, we alsocollected the diffraction patterns of CeO2 powder calibrant(NIST Standard Reference Material SRM-674b) insertedinto a quartz tube and placed at the same position in thetensile rig as the studied steel samples. We summed allthe recorded 2-D patterns for the calibrant and also foreach of the samples. The resultant CeO2 pattern was usedto determine the sample to detector distance, beam centrecoordinates and the detector tilt angles with respect tothe incoming X-ray beam using the program Fit2D [27].Afterwards, we integrated the summed 2-D patterns alongthe azimuth angle for constant scattering angles to obtainone-dimensional (1-D) patterns (intensity vs. scatteringangle). The resultant 1-D patterns were analysed via theRietveld method implemented in the Fullprof Suite soft-ware package [28]. The instrumental resolution was deter-mined by fitting the CeO2 pattern to a pseudo-Voigtfunction. The experimental values of the instrumentalparameters were subsequently used in the full profile anal-ysis of the 1-D patterns of the studied steel samples. Thisanalysis yielded the phase fraction and lattice parameterof the constituent phases in the starting microstructures.

The lattice parameter and phase fractions obtained atroom temperature were then used as starting values inthe Rietveld analysis of the diffraction patterns collectedduring heating. For this, the 2-D patterns were summedin bunches of a hundred (20 patterns for each x-scan � 5sample heights) and integrated to obtain the corresponding1-D patterns. In order to carry out a whole profile fitting ofthe series of 1-D patterns, we set up Fullprof in batch pro-cessing mode. Additionally, for the samples studied undertensile stress, we integrated each summed 2-D pattern forconstant scattering angles over sections of the detector



Fig. 1. Visible light optical micrograph of the SAE 52100 steel austen-itized at a temperature of (a) 860 �C, (b) 900 �C and (c) 950 �C. The steelsamples were subsequently quenched in oil to 60 �C for 15 min, rinsed incold water and tempered at 160 �C for 90 min. The microstructure hasbeen revealed using a nital 1.5 vol.% etching solution.


covering an azimuth angular range of 10�. When the sam-ple is under stress, the diffraction rings adopt an ellipsoidalsymmetry that reflects the orientation of the diffractinggrains with respect to the applied stress for different partsof the detector. Each 2-D pattern then yielded 36 1-D pat-terns. All the resultant 1-D patterns were fitted using Full-prof in batch processing mode. The variation of the latticeparameter as a function of the azimuth angle for each timestep reveals the evolution of the elastic properties of theconstituent phases with temperature and time. For selectedtime steps, the sample texture was inferred from acombined Rietveld analysis of the 1-D patterns using theE-WIMV algorithm [29] implemented in the softwarepackage MAUD [30].

3. Results and discussion

3.1. Microstructures at room temperature

Fig. 1 shows the visible light optical micrograph of theroom-temperature microstructure as a function of theaustenitization temperature. The microstructure consistsof tempered martensite, retained austenite and globularcarbides. These carbides can be clearly identified in the860 �C sample as white-etched particles embedded in a veryfine martensite/austenite matrix. Higher austenitizationtemperatures cause a coarsening of the austenite grains,which results in an increase in the length of the martensiticplates that can be easily revealed by optical microscopy.These plates constitute a brown-etched network whereirregular-shaped austenite grains are intermixed with glob-ular carbide particles.

Fig. 2 displays the 2-D diffraction pattern recorded atroom temperature and the corresponding 1-D pattern,before (2a and c) and after (2b and d) having heated thesample to 216 �C and held at that temperature for 4.5 h.The recorded diffraction patterns at room temperatureafter tempering at 160 �C (Fig. 2a and c) consist of a num-ber of diffraction rings that can be indexed with the threeaforementioned crystallographic phases: martensite(body-centred cubic (bcc), a), austenite (face-centred cubic(fcc), c) and cementite (orthorhombic, Fe3C). The resultsof the Rietveld analysis of the corresponding 1-D patternsare displayed in Fig. 3a, c and e for the three studied micro-structures. During austenitization the starting spheroidize-annealed carbide/ferrite microstructure gives rise to anaustenitic matrix containing a small fraction of undissolvedcarbide particles. The ferrite phase is expected to have fullytransformed at a temperature of 740 �C, and a furtherincrease in temperature will promote the dissolution ofthe remaining particles [31]. We experimentally observe adecrease in the amount of cementite (Fe3C) present atroom temperature from �3% to �1% when increasingthe austenitization temperature from 860 to 950 �C. Forthe studied experimental conditions, we haven’t foundany indication for the presence of other chromium-rich car-bides such as M7C3 or M23C6 (M = Fe, Cr) [31] in the


recorded diffraction patterns for this material. Thermody-namic calculations predict the complete dissolution of thecarbide particles at temperatures above 870 �C [32]. Anaustenitization time of 20 min does not suffice to reachthe expected equilibrium fractions. During the progressivecarbide dissolution, part of the chromium atoms will dif-fuse into the undissolved carbide particles so that theybecome more stable over time. The carbide particles pres-ent an initial chromium content of �7 wt.% [33] andcementite can host up to 18.4 wt.% Cr [31]. The favouredcarbide dissolution at higher temperatures leads to anincrease in both the carbon content and the grain size ofthe austenite phase. The carbon enrichment induces a



X

Y

0 500 1000 1500

0

500

1000

1500

2000

2500

γγα α

γ

α α

X

Y

0 500 1000 1500

0

500

1000

1500

2000

2500 Before heating

Before heating After heating

After heating

(a) (b)

(c) (d)

Fig. 2. 2-D X-ray diffraction pattern of the SAE 52100 steel sample austenitized at 900 �C. The patterns have been recorded at room temperature (a)before and (b) after having heated the sample to 216 �C and held at that temperature for 4.5 h. The presence of the {200}, {220} and {311} austenitediffraction rings before heating the sample are denoted as c200, c220 and c311. The Rietveld refinement of the corresponding 1-D patterns is shown:(c) before and (d) after heating. The position of the ferrite (a), austenite (c) and cementite (Fe3C) reflections is indicated by vertical lines. The insetsillustrate the disappearance of the {111}c reflection due to the austenite decomposition during heating.


larger amount of austenite retained in a metastable condi-tion at room temperature when the sample has beenquenched from a higher austenitization temperature (Fig. 3).

The lattice parameter of austenite at room temperaturein the as-tempered microstructure takes the value of3.5892(2) A (Taust = 860 �C), 3.5951(2) A (Taust = 900 �C)and 3.5976(1) A (Taust = 950 �C). The carbon content inaustenite can be derived from the experimental value ofthe lattice parameter according to [34,35]:

ac ¼ 3:556þ 0:0453xC þ 0:00095xMn þ 0:056xAl

þ 0:0006xCr þ 0:0015xCu ð1Þ


where ac is in A and xC, xMn, xAl, xCr and xCu are in wt.%.The chromium content in austenite (xCr = 1.10 wt.%) is ta-ken from reported electron probe micro-analysis (EPMA)studies [36]. The carbon content in austenite derived fromEq. (1) amounts to 0.69 wt.% (Taust = 860 �C), 0.82 wt.%(Taust = 900 �C) and 0.88 wt.% (Taust = 950 �C). The mar-tensite that is formed during quenching from high temper-ature will also present an increase in its carbon contentwith the austenitization temperature. However, the mar-tensite fraction decreases due to the higher stability of thecarbon-enriched parent austenite phase. The as-quenchedmartensite is expected to present a tetragonal symmetry



850 900 950 1000

σ = 295 MPa

α-Fe

Austenitization temperature (ºC)


Fe3C

σ = 295 MPa


σ = 295 MPa

Before heatingα-Feγ-Fe

After heatingα-Feγ-Fe

α-Fe

γ-Fe

0

20

40

60

80

100

α-FeFr

actio

n (%

)

Before heatingα-Feγ-Fe

After heatingα-Feγ-Fe

σ = 0 MPa

γ-Fe

0

1

2

3

4

Fe3C

σ = 0 MPa

Frac

tion

(%)


800 850 900 950 1000

2.866

2.868

2.870

2.872

2.874 σ = 0 MPa

Austenitization temperature (ºC)

Before heating After heating α-Fe

(a) (b)

(c) (d)

(e) (f)

Fig. 3. Variation in the fraction of austenite, ferrite and cementite, together with the variation in ferrite lattice parameter with the austenitizationtemperature. These results have been obtained via a Rietveld analysis of the 1-D diffraction patterns at room temperature, before (blue) and after (red)having heated the sample to a temperature in the range of 200–235 �C and held at that temperature for a period of 2–8 h depending on the austenitedecomposition kinetics. The experiments have been performed both at zero stress and in the presence of an applied stress of r = 295 MPa. The phasespresent in the microstructure have been labelled as follows: ferritic matrix (a-Fe), metastable austenite (c-Fe) and cementite (Fe3C). (For interpretation ofthe references to colour in this figure legend, the reader is referred to the web version of this article.)


that would easily be distinguished in the diffraction patterndue to the clear separation of the {002} reflection from the{20 0}/{020} doublet [37]. However, the tempering stepafter quenching causes the formation of nano-sized e/gtransition carbides in the martensite phase [38], so thatthe reservoir of carbon atoms in solid solution in martens-ite is significantly reduced. As a consequence, the symmetryof the martensite lattice shifts from tetragonal to cubic. Themartensite lattice parameter at room temperature takes thevalue of 2.8706(2) A (Taust = 860 �C), 2.8718(2) A(Taust = 900 �C) and 2.8723(1) A (Taust = 950 �C). Theremaining carbon content in solution in the martensite


phase can be estimated from the measured values of its cu-bic lattice parameter using the equation [39,40]:

aa ¼ 2:8664þ 0:00055xMn þ 0:020xC ð2Þwhere aa is in A and xC and xMn are in wt.%. The remain-ing carbon content in martensite after quenching and tem-pering amounts to 0.20 wt.% (Taust = 860 �C), 0.26 wt.%(Taust = 900 �C) and 0.29 wt.% (Taust = 950 �C). The exper-imental values of the cementite lattice parameters are:a = 5.077(3) A, b = 6.708(4) A and c = 4.513(2) A. Theapplication of an external stress of r = 295 MPa at roomtemperature does not trigger the decomposition of the




metastable austenite phase in any of the studied samples.The Rietveld results of the 1-D diffraction patterns mea-sured in the presence of the applied stress are collected inFig. 3b, d and f. Similar trends in the phase fraction andlattice parameter of the constituent phases are observedwhen comparing the results for r = 0 and r = 295 MPa.The main difference between the results from Fig. 4e andf stems from the increase in the lattice parameter whenthe stress is applied due to the occurrence of elastic strains.We have not observed a significant broadening of the dif-fraction peaks when the load is applied.

The diffraction rings from the metastable austenitephase have completely disappeared after having heatedthe sample at 216 �C and held at that temperature for4.5 h (Fig. 2b and d). The 1-D diffraction pattern has beenanalysed taking into account only two crystallographicphases: the bcc phase (a) for the ferrite matrix and theorthorhombic cementite phase (Fe3C). The results of theanalysis are displayed in red in Fig. 3. The microstructureis composed of a ferritic matrix with an increased amountof cementite. The bcc lattice parameter has decreased con-siderably with respect to its initial value before heating thesample. This decrease in lattice parameter translates into areduction in the carbon content of the ferritic matrix. Theobtained values of the bcc lattice parameter as a functionof the austenitization temperature are: 2.86713(5) A(Taust = 860 �C), 2.8673(1) A (Taust = 900 �C) and2.86707(6) A (Taust = 950 �C). The derived carbon contentin the ferritic matrix from the experimental values of thelattice parameter using Eq. (2) amounts to 0.03 wt.%(Taust = 860 �C), 0.03 wt.% (Taust = 900 �C) and 0.02 wt.%(Taust = 950 �C).

{110

} α

{200

} α

{111

} γ

{200

} γ

{220

} γ2θ (°)

time

(h)

3.5 4.5 5.5 6.5 7.5

1

2

3

4

5

6

{121

} Fe 3

C+

{210

} Fe 3

C

{022

} Fe 3

C+

{131

} Fe 3

C

{301

} Fe 3

C+

{231

} Fe 3

C

{312

} Fe 3

C+

{123

} Fe 3

C

Fig. 4. 1-D X-ray diffraction patterns collected for the Taust = 900 �C samplestress of r = 295 MPa. The temperature profile during heating is shown in thindicated in the figure. The additional weaker reflections come from the Fe3C


3.2. Evolution of phase fractions during heating

The time evolution of the 1-D diffraction patternrecorded while heating the Taust = 900 �C sample aboveroom temperature is shown as an example in Fig. 4. Thestarting 1-D pattern is composed of relatively strong fer-ritic (a) and austenitic (c) reflections, together with a num-ber of weaker reflections coming from the Fe3C phase.During the initial heating stage the recorded reflectionsshift to lower scattering angles due to the thermal expan-sion of the corresponding phases. At �30 min the intensityof the austenite reflections starts to decrease, and simulta-neously the ferrite reflections become more intense. Oncethe austenite decomposition has started, it progresses overtime until the austenite phase has disappeared completely.During this process there is also an increase in intensity ofthe Fe3C reflections. At the end of the process only the fer-rite and Fe3C reflections can be observed in the diffractionpattern. The observed time evolution of the diffraction pat-tern reveals that when the necessary conditions to destabi-lize the austenite phase have been reached, it continuouslytransforms into ferrite until all the austenite has disap-peared. The excess of carbon atoms present in the startingaustenite is expelled in the form of carbides during the aus-tenite decomposition process.

In order to monitor the austenite decomposition kineticsabove room temperature, we have derived the time evolu-tion of the austenite fraction by performing a Rietveldanalysis of the 1-D diffraction patterns. The resultsobtained for the studied microstructures and selected com-binations of temperature and stress are shown in Fig. 5.Significant differences in both the starting fraction and

0501001502000

1

2

3

4

5

6

temperature (°C)

time

(h)

{211

} α

{220

} α

{311

} γ

{222

} γ

8.5 9.5

as a function of time during heating in the presence of a constant appliede right part of the figure. The ferritic (a) and austenitic (c) reflections arephase. Selected Fe3C reflections have also been indicated in the figure.




the time to complete the transformation can clearly beobserved when varying the austenitization temperature.For example, the metastable austenite has transformedafter 3 h for Taust = 900 �C, while a holding time of 6 h isrequired for the austenite to transform in the case ofTaust = 950 �C. The presence of an applied stress ofr = 295 MPa leads to faster decomposition kinetics. After1 h the remaining austenite fraction amounts to 15%(r = 0 MPa) and 10% (r = 295 MPa) for Taust = 900 �C.A significant acceleration of the kinetics is also achievedby an increase in the transformation temperature (Texp)from 203 to 216 �C. Therefore, three different parametersseem to influence the decomposition kinetics of the meta-stable austenite phase: (1) the starting microstructure (viathe austenitization temperature, Taust), (2) the temperatureat which the isothermal holding is carried out (Texp), and(3) the presence of an applied stress (r).

The activation energy for the decomposition process at agiven combination of Taust and r can be obtained from thetime elapsed between two fixed transformation stagesmeasured at different temperatures (Texp). In this study,we have chosen the time when 1=4 and ½ of the starting aus-tenite fraction remains (t1=4 and t½). These two stages werechosen because the temperature had reached a constantvalue at this part of the transformation. The activationenergy can be estimated by the following equation [41]:

lnðt1=2 � t1=4Þ ¼ ln Aþ QRT

ð3Þ

5

10

15

20

25

30

σ = 0 MPa

Taust

860ºC 900ºC 950ºC

aust

enite

frac

tion

(%)

Texp = 216ºC

0 1 2 3 4 50

5

10

15

20

25

30

aust

enite

frac

tion

(%)

time (h)

Taust

860ºC 900ºC950ºC

σ = 295 MPaTexp = 216ºC

(a)

(c)

Fig. 5. Time evolution of the austenite phase fraction for selected combinatisynchrotron X-ray diffraction experiment (Texp) and applied stress (r).


where A is a constant, Q the activation energy, R the gasconstant and T the temperature. Since the mechanism ofthe austenite decomposition remains the same in the tem-perature range used in this experiment, Eq. (3) can be usedto estimate Q. The experimental values for the activationenergy are collected in Table 2. When r = 0 MPa, the acti-vation energy Q takes a value of 138(4) kJ mol�1 forTaust = 900 �C and 143(6) kJ mol�1 for Taust = 950 �C.These values are in good agreement with the values of138–148 kJ mol�1 reported in the literature for the diffu-sion of carbon in austenite [42–44]. This observation indi-cates that the carbon diffusion in the remaining austeniteis the limiting step in the decomposition kinetics. Duringthe decomposition of the metastable austenite into ferrite,the carbon atoms will diffuse out of the newly formed fer-rite grains into the ferrite/austenite interface. The activa-tion energy for the diffusion of carbon in ferrite, 78–84 kJ mol�1 [45,46], is lower than in austenite, so that thecarbon atoms will tend to pile up at the interface. Thelong-range diffusion of carbon into the austenite ahead ofthe moving austenite/ferrite interface will control the over-all transformation kinetics. This diffusion of carbon aheadof the interface will generate local carbon enrichments thatwill trigger the precipitation of cementite in the austenitephase. This will act as a barrier for the moving austenite/ferrite interface, so that the transformation will stop untila new successful nucleation event takes place in the remain-ing parts of the original austenite grain.

1 4 5 6

σ = 295 MPaTaust = 900ºC

time (h)

Texp

203ºC216ºC

Texp

203ºC216ºC

Taust = 900ºCσ = 0 MPa

2 36

(b)

(d)

ons of austenitization temperature (Taust), temperature during the in situ



Table 2Experimental values of the activation energy for the austenite decompo-sition process (see text).

Taust (�C) Activation energy (kJ/mol)

r = 0 MPa r = 295 MPa

900 138(4) 104(4)950 143(6) 82(4)


In the presence of an applied stress of r = 295 MPa, theactivation energy is significantly reduced and takes a valueof 104(4) kJ mol�1 for an austenitization temperature of900 �C and a value of 82(4) kJ mol�1 for an austenitizationtemperature of 950 �C. A finer surrounding martensiticphase (lower Taust) seems to exert a larger shielding effecton the stress felt by the austenite grains, leading to a higheractivation energy for the austenite decomposition. Anapplied tensile stress will (1) increase the driving force forthe short-range diffusion of iron atoms across the austen-ite/ferrite interface [47], and (2) accelerate the carbon diffu-sion kinetics in both austenite [47] and ferrite [48]. Duringthe carbon diffusion ahead of the interface, the carbonatoms are expected to interact with the point-like defectspresent in the austenite. The binding energy of nearest-neighbour carbon–vacancy pairs is estimated at 35–40 kJ mol�1 [49]. We have experimentally observed areduction in the activation energy of the austenite decom-position of the same magnitude as the carbon–vacancybinding energy. This suggests that the applied stress accel-erates the austenite decomposition by perturbing the car-bon–vacancy binding that slows down the carbon diffusion.

3.3. Temperature dependence of the lattice parameter

The evolution of the austenite and ferrite lattice param-eter is shown in Fig. 6a and b for an applied stress ofr = 0 MPa and r = 295 MPa. At temperatures lower thanthe tempering temperature of 160 �C (represented by a ver-tical dashed line in Fig. 6), the temperature dependence ofthe lattice parameter is governed by the coefficient of linearexpansion a = (1/a)(da/dT) for each phase, as no phasetransformation takes place in this temperature range. Fromroom temperature up to the maximum temperature reachedin this experiment of 235 �C, the thermal expansion of bothphases is expected to behave linearly with temperature[34,50]. We have therefore fitted the experimental valuesof the lattice parameter up to 160 �C to a quadratic temper-ature dependence of the form: ac,a = A + BT + CT2. Theobtained values for the fitting parameters are collected inTable 3. The zero-stress values at 20 �C of ac = 18.1� 10�6 K�1 and aa = 12.6 � 10�6 K�1 are in good agree-ment with the reported room-temperature values ofac = 16.0 � 10�6 K�1 and aa = 11.7 � 10�6 K�1 [34].

We have extrapolated the obtained thermal expansionof the phases to temperatures above 160 �C. Above thistemperature, we observe changes in phase fractions dueto the tempering of the martensite and the decompositionof the metastable austenite. The difference between the


experimental values of the lattice parameter and the pre-dicted values from the thermal expansion coefficient isattributed to a change in the carbon content due to theaforementioned processes. The change in carbon content(DxC) with temperature, as derived using Eqs. (1) and (2),is shown in Fig. 6c and d for austenite and ferrite, respec-tively. Above 160 �C, both phases undergo a continuouschange in their carbon content. In the ferrite phase, the car-bon content decreases gradually during heating, so that atthe end of the process only 0.03 wt.% C is left in solution inthe bcc lattice. In contrast, the austenite experiences a sig-nificant increase in carbon content above 160 �C, attaininga maximum of DxC = 0.07 wt.% C. Beyond this maximum,its carbon content decreases continuously until all themetastable austenite grains have transformed.

Our results indicate that the precipitation of nano-sizede/g transition carbides within the martensite phase isresumed once the temperature of 160 �C is surpassed.These intermediate nano-carbides will eventually lead tothe formation of cementite particles [51]. The occurrenceof carbon enrichment in the metastable austenite takesplace at a stage in the heating process when the austenitedecomposition has not yet started. Therefore, the observedcarbon enrichment cannot be due to the diffusion of carbonout of the newly formed ferrite grains and into the remain-ing austenite ahead of the moving interface. The additionalcarbon enrichment will stem from the partition of carbonfrom martensite into austenite [52,53]. This carbon parti-tioning takes place simultaneously with the precipitationof transition carbides inside the martensite phase. At thetime when the maximum carbon enrichment in austeniteof DxC = 0.07 wt.% C has been reached, the decrease incarbon content in the martensite phase amounts toDxC = –0.18 wt.% C. The additional carbon atoms in theaustenite will further destabilize this phase due to anincrease in the chemical driving force for carbide nucle-ation [54]. The high-carbon austenite grains are expectedto transform first, and the excess carbon atoms in the resul-tant ferrite grains will diffuse into the austenite wherecementite precipitation will take place. Consequently, thecarbon content of the remaining austenite phase will grad-ually decrease. The newly formed ferrite grains will have acarbon content close to the equilibrium concentration andsignificantly lower than the carbon content in the supersat-urated martensite phase.

3.4. Elastic behaviour

The application of an external uniaxial stress causes thediffraction rings to adopt an elliptical shape. The value ofthe interplanar spacing for the grains contributing to thediffracted intensity in a specific region along the ringdepends on the grain orientation with respect to the appliedload [55]. Fig. 7 shows the value of the lattice parameter asa function of the azimuth angle (g) for the austenite andferrite phase at selected times. The obtained maximumand minimum values of the lattice parameter correspond



50 100 150 200 250 300-0.3

-0.2

-0.1

0.0

σ = 0 MPaσ = 295 MPa

Δ xC

(wt.%

)

temperature (°C)0 50 100 150 200 250 300

-0.3

-0.2

-0.1

0.0

σ = 0 MPaσ = 295 MPa

ΔxC

(wt.%

)

temperature (°C)

γ -Fe

γ -Fe3.58

3.59

3.60

3.61

3.62

3.63σ = 0 MPaσ = 295 MPa

2.871

2.874

2.877

2.880

2.883σ = 0 MPaσ = 295 MPa

α-Fe

α-Fe

(a) (b)

(c) (d)

Fig. 6. Temperature dependence of the lattice parameter of (a) austenite and (b) ferrite in the Taust = 900 �C sample during heating, at zero stress and in anapplied stress of r = 295 MPa. The continuous lines correspond to the thermal expansion obtained for temperatures below 160 �C (see Table 3). Theobtained thermal expansions have been extrapolated to higher temperatures. The change in carbon content (DxC) as a function of temperature, as derivedfrom the lattice parameter (see text), is shown for (c) austenite and (d) ferrite. The vertical dashed line in (a–d) indicates the tempering temperature of160 �C used in the prior material processing.

Table 3Values obtained by fitting the temperature dependence of the lattice parameter of austenite and ferrite up to 160 �C to a quadratic behaviourac,a = A + BT + CT2, both at zero stress and in the presence of an applied tensile stress of r = 295 MPa.

r (MPa) Austenite Ferrite

A (A) B (A/K) � 105 C (A/K2) � 108 A (A) B (A/K) � 105 C (A/K2) � 108

0 3.5786(13) 5.12(63) 2.37(94) 2.8634(4) 1.98(29) 2.78(42)295 3.5792(12) 5.11(67) 3.11(93) 2.8624(5) 2.98(30) 1.92(42)


to subsets of grains whose scattering vector (Q) is orientedapproximately parallel (g = 0� and 180�) or perpendicular(g = 90� and 270�) to the applied stress, respectively. Thisis due to the small values of the scattering angles for highenergy (>50 keV) X-ray diffraction experiments [56]. Thefitting procedure for the g-dependence of the lattice param-eter to a sin2g law yields the maximum value of the latticeparameter (amax) and the width of the sinusoidal function(Da). These values have been used to derive the averagephase strain (hei) and the elliptical distortion of the diffrac-tion rings (De) according to:

hei ¼ hai � a0

a0

ð4Þ

De ¼ Daa0

ð5Þ

where a0 is the lattice parameter obtained when no externalstress was applied to the sample and hai is the average lat-


tice parameter in the applied stress condition (hai =amax– (Da/2)). The experimental values of hei and De canbe used to obtain the Poisson’s ratio (m) of the constituentphases since:

hei ¼ 1� m2

ek ð6Þ

De ¼ ð1þ mÞek ð7Þ

where m = �e\/e||, with e|| and e\ denoting the axial andtransverse strains, respectively. By combining Eqs. (6)and (7), we obtained a value of m = 0.26(1) for austeniteand m = 0.29(1) for ferrite.

In Fig. 8 the variation in De for ferrite and austenite isshown during heating as a function of time and austenitefraction. De constitutes a very sensitive probe of thechanges in the elastic behaviour during the heating process.During heating to 216 �C, the elastic strain of both phasesincreases due to the softening of the lattice caused by the



0 50 100 150 200 250 300 3502.868

2.870

2.872

2.874

2.876

2.878

2.880

α-Fe

t = 0.5 h

t = 0.7 h

t = 0.1 h

azimuth angle (°)

t = 0 h

3.590

3.595

3.600

3.605

3.610

3.615

3.620

γ -Fe

t = 0.7 h

t = 0.5 h

t = 0.1 h

t = 0 h

(a)

(b)

Fig. 7. Variation in lattice parameter for (a) austenite (c) and (b) ferrite(a) as a function of the azimuth angle (g) for selected times during heating.Each data point corresponds to the Rietveld analysis of the 1-D diffractionpattern (intensity vs. scattering angle) originating from a region in thedetector covering 10� along the azimuth angle. The lines represent a fit to asin2g law (see text).

0.0 0.4 0.8 1.2 1.60.00

0.05

0.10

0.15

0.20

0.25

σ = 295 MPa

α-Fe

γ -Fe

Aus

teni

te e

last

ic s

trai

n (%

)

time (h)

0.05

0.10

0.15

0.20

0.25

0.30

0.35

Fer

rite

elas

tic s

trai

n (%

)

0 5 10 15 20 25 300.00

0.05

0.10

0.15

0.20

0.25

γ -Fe

Aus

teni

te e

last

ic s

trai

n (%

)

Austenite fraction (%)

α-Fe

0.05

0.10

0.15

0.20

0.25

0.30

0.35

Fer

rite

elas

tic s

trai

n (%

)

σ = 295 MPa

(a)

(b)

Fig. 8. Evolution of the elastic strain (De) in the austenite (c) and ferrite(a) phase with (a) time and (b) austenite fraction, of the Taust = 900 �Csample during heating and holding at a temperature of 216 �C in thepresence of an applied stress of r = 295 MPa.


rise in temperature. Once the constant temperature of216 �C has been reached, the strain produced in the ferriticphase continues to increase. This further increase in straincan be ascribed to the newly formed ferrite grains resultingfrom the austenite decomposition. These new grains aresofter than the martensite phase already present in themicrostructure. However, as the austenite decompositionproceeds at a constant temperature, the elastic strain inthe remaining austenite grains continuously decreases. Attimes longer than �0.6 h the strain in the remaining austen-ite grains becomes smaller than the overall strain observedin the ferrite phase. This implies that at longer holdingtimes the austenite grains act as a reinforcement of themajority ferritic matrix. The observed material softeningabove 200 �C during prolonged service of the bearing [9]originates in the ferritic matrix whose fraction increasesover time due to the austenite decomposition.

3.5. Texture evolution above room temperature

The evolution of the texture in the austenite and ferritephase is shown in Fig. 9 for selected steps during heating.The texture is expressed in terms of inverse pole figuresalong the three principal sample directions: rolling (RD),normal (ND) and transverse (TD) directions. In this study


RD corresponds to the loading direction, and also to thecylindrical axis of the sample and of the raw bar material.TD and ND are parallel and perpendicular to the diameterof the raw bar material, respectively. Both phases present aweak starting texture that is mainly concentrated alongRD. The metastable austenite shows a predominanth210i component along RD. In the case of the ferriticphase, three weak texture components are observed atroom temperature along RD: h210i, h5 11i and h111i[57]. During heating the decomposition of the metastablegrains contributing to the h21 0i component in the austen-ite phase gives rise to a significant enhancement of the fer-rite h21 0i component [58]. Furthermore, at longer holdingtimes grain rotations are observed within the austenitephase. This leads to the appearance of a h511i componentalong RD at the expense of the h210i component. The latertransformation of the h511i oriented austenite grains pro-duces an increase in the ferrite h5 11i component. Theh111i component remains unaltered during the heatingprocess.

4. Conclusions

We have monitored the microstructure evolution in SAE52100 steel during heating under tensile stress by in situ



t = 0 h(fγγ = 26 %)

t = 0.85 h(fγ = 13 %)

t = 1.10 h(fγ = 6.5 %)

austenite ferriteRD ND TD RD ND TD

100 110

111

100 110

111 111

100 110

111

100 110

111

100 110

111

100 110

111

100 110

111 111

100 110

111

100 110

111 111

100 110

111

100 110

111 111

100 110

111

100 110

111

100 110

100 110

100 110

100 110

100 110

111

Fig. 9. Texture evolution of the austenite and ferrite phase in the Taust = 900 �C sample for selected steps during heating in the presence of an appliedstress of r = 295 MPa. The austenite phase fraction (fc) is indicated in the figure. The texture is expressed in terms of the inverse pole figures recalculatedfrom the Rietveld analysis of the diffraction data. The scale is in multiples of random distribution (mrd).


high-energy X-ray diffraction. The main conclusions are asfollows. (1) The metastable austenite continuously decom-poses into ferrite and cementite until all the austenite hasdisappeared. This process is controlled by the carbon diffu-sion into the austenite ahead of the moving interface. (2)The acceleration of the kinetics by the applied stress hasbeen ascribed to the perturbation of the carbon–vacancybinding in austenite. (3) A lower austenitization tempera-ture results in a finer surrounding martensite phase and ahigher activation energy for the austenite decomposition.(4) A significant migration of carbon atoms from martens-ite into austenite takes place before the austenite starts todecompose. (5) The elastic strain in the remaining austenitecontinuously decreases at a constant temperature. (6) Thetexture evolution reflects grain rotations and crystal orien-tation relationships during the austenite decompositioninto ferrite.

Acknowledgements

We acknowledge the European Synchrotron RadiationFacility for provision of synchrotron radiation facilities.This research is supported by the Dutch Technology Foun-dation STW, applied science division of NWO and theTechnology Program of the Ministry of Economic Affairs.Romain Blonde acknowledges the support of the Materialsinnovation institute M2i (www.m2i.nl) under the ProjectNumber M41.5.08313 of the research programme.

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