On low temperature bainite transformation characteristics usingin-situ neutron diffraction and atom probe tomography

Khushboo Rakha a,n, Hossein Beladi a, Ilana Timokhina a, Xiangyuan Xiong b,Saurabh Kabra c, Klaus-Dieter Liss c, Peter Hodgson a

a Institute for Frontier Materials, Deakin University, Geelong, VIC 3216, Australiab Department of Materials Engineering, Monash University, Victoria 3800, Australiac Australian Nuclear Science and Technology Organisation, The Bragg Institute, New Illawarra Road, Lucas Heights, NSW 2234, Australia

a r t i c l e i n f o

Article history:Received 7 September 2013Accepted 16 September 2013Available online 20 September 2013

Keywords:Atom probe tomographyNeutron diffractionBainitePhase transformation

a b s t r a c t

In-situ neutron diffraction was employed to monitor the evolution of nano-bainitic ferrite during lowtemperature isothermal heat treatment of austenite. The first 10 peaks (austenite, γ and ferrite, α) weremonitored during austenization, homogenisation, rapid cooling and isothermal holding at 573 K.Changes in the α-110 and γ-111 peaks were analysed to determine the volume fraction changes andhence the kinetics of the phase transformation. Asymmetry and broadening in the α-200 and γ-200peaks were quantified to lattice parameter changes due to carbon redistribution as well as the effects ofsize and dislocation density. Atom Probe Tomography was then used to confirm that, despite thepresence of 1.5 mass% Si, carbide formation was evident. This carbide formation is the cause of poorductility, which is lower than expected in such steels.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

The ongoing industry quest for high performance and fuelefficiency encourages scientists to explore novel microstructuresin steels. The potential of new generation nanostructured bainiticsteels offering notable strength (2.5 GPa), toughness (40 MPa m1/2)and ductility (30%) has been widely debated [1–4]. These excellentproperties are attributed to the refinement of the microstructureleading to the formation of very fine bainitic ferrite laths with anaverage thickness of �50 nm and retained austenite filmsbetween these laths [5]. The Transformation Induced Plasticity(TRIP) effect arising from retained austenite films requires anoptimum amount of carbon for the stability of retained austenite,thus, carbide formation is suppressed during isothermal bainiteformation using Si. However, recently it has been made evidentthat Fe3C carbides do form in nanostructured bainite despite highSi content (�1.5 mass%) [6,7].

The mechanical properties, especially the excellent balance ofstrength/toughness/ductility are very sensitive to transformationparameters such as isothermal temperature and time. Unfavourably,decreasing the isothermal temperature reduces the carbon diffusionrate tremendously, thus, increasing the time for transformation upto 2 weeks. At such slow rates, control of carbon diffusion for an

optimum stability of retained austenite and the commercial pro-duction are restricted. While the effect of transformation tempera-ture on microstructure and variant selection has been wellestablished [5,8,9], there is an acute need to establish the in-situtransformation mechanism for further development of such steels.The stipulated mechanism could be different from the conventionalbainitic transformation as the size of bainitic ferrite laths is close tothe simulated critical bainitic nuclei size [10].

The isothermal bainitic transformation of nanostructured steelshas been previously investigated using in-situ X-ray and neutrondiffraction techniques, though the results have not been consistent[11–13]. Babu et al. [11] showed that the austenite diffractionpeaks split before the onset of bainitic transformation, suggestingcarbon partitioning in the austenite phase prior to the onset ofbainitic transformation. However, Stone et al. [12] did not observethe splitting of diffraction patterns before the onset of bainitictransformation suggesting homogenous distribution of carbon inaustenite before the start of bainitic transformation. Later, Kooet al. [13] also observed the peak broadening and peak shiftfollowing the onset of bainitic transformation using in-situ neu-tron diffraction, however, the peak split could not be identifieddue to the low resolution of the employed technique.

In the present study, in-situ neutron diffraction was employedusing the high-intensity powder diffractometer WOMBAT [14] atthe OPAL reactor of the Australian Nuclear Science and TechnologyOrganisation (ANSTO) to monitor the evolution of nano-bainitic

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0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.msea.2013.09.055

n Corresponding author. Tel.: þ61 405057290; fax: þ61 352271103.E-mail addresses: krakha@deakin.edu.au, krakha@gmail.com (K. Rakha).

Materials Science & Engineering A 589 (2014) 303–309

ferrite during low temperature isothermal heat treatment. Thiswork was inspired by the need to verify prior observations [11–13]using an instrument with higher temporal resolution to cope upwith subtle changes during the isothermal heat treatment, espe-cially just before and after reaching the isothermal heat treatmenttemperature. Since the neutrons penetrate and probe a largervolume of material, the measurement is less sensitive to thesurface decarburization during heating, as compared to the X-raydiffraction technique. Neutron diffraction also provides improvedcrystallites statistics. Transmission electron microscopy (TEM) wasperformed complementary to evaluate the effect of lath size anddislocation density on the peak broadening of both bainitic ferriteand austenite phases observed in the neutron diffraction results.Interesting outcomes from carbon content calculations in thetransformation product phases led to the use of Atom ProbeTomography (APT) to investigate in more detail the microstruc-tural features and their effect on the mechanical properties ofnanostructured bainite.

2. Experimental procedure

The chemical composition of the alloy used in this study isgiven in Table 1. The transformation temperatures Bs and Ms forbainitic and martensitic formation were estimated to be 658 K and428 K, respectively [15]. A slab, with an initial thickness of 40 mm,was hot-rolled in a deformation temperature range of 1473–1273 K through successive passes to obtain a final thickness of�12 mm. The samples were then sealed into a stainless steel bagand homogenised at 1523 K for 480 min in a vacuum of 100 Pa,followed by furnace cooling to room temperature. Cylindricalsamples with a height of 15 mm and a diameter of 10 mm weremachined from the hot rolled material along the rolling direction.

The experiment was performed using in-situ neutron diffrac-tion with the high intensity powder diffractometer, WOMBAT [14],equipped with a heating system and a rapid quencher. The latterprototype was critically important to avoid any phase transforma-tion to equilibrium ferrite during cooling from the austenitizingtemperature (1093 K) to the bainitic transformation temperature(573 K). The rapid quencher was designed and built by the sampleenvironment team at Bragg Institute of ANSTO providing coolingrates over 27 K/s. The heating occurred via halogen lamps and thequench was performed through a controlled flow of liquid nitro-gen. This was a first-case ever of a rapid quencher in the neutronsample environment community.

The temperature was monitored throughout the experimentusing two thermocouples embedded at each end of the sample.The specimen was hung using a thermocouple wire in betweenthe two heating lamps. The thermocouple was attached to aEurotherm temperature regulator, which was later controlledremotely through the WOMBAT system software. A germaniumsingle-crystal was employed as a monochromator for WOMBAT,using the Ge-115 reflection in symmetric Bragg geometry and 1201take-off angle. Delivered wavelength and wavenumber wereλ¼1.49 Å and k¼3.31 Å�1, respectively. Calibration of the instru-ment was undertaken using the Al2O3 and LaB6 standard speci-mens. All diffraction data has been calibrated and expressed inreciprocal space Q¼4π/λsin((2θ)/2), where 2θ is the scatteringangle measured by the instrument.

During in-situ neutron diffraction, a cylindrical sample wasaustenitized by heating to 1093 K at the rate of 5 K/s and retainedat the same temperature isothermally for 30 min, followed byrapid cooling to 573 K at a rate of 25 K/s under liquid nitrogenflow. The sample was then held isothermally at 573 K for 12 h(750 min). It is worth mentioning here that the temperaturefluctuations did not influence the result as the temperature dropwas still higher than the Ms temperature, and the time fortemperature stabilisation was much less than the incubation timefor the bainitic phase transformation. The time resolution used foracquisition of data was 13 s/scan, which was much faster thanprevious studies [11–13].

Diffraction data was obtained from where the austenite phaseformed until the end of transformation. The diffraction data was fittedby a voigt function to obtain the position, intensity (area) and width(Full width half maximum, FWHM) for each peak. The integratedintensities of α-110 and γ-111 were used in direct comparison toobtain the phase fractions. The peak position Ghkl in reciprocal space isa valuable parameter denoting total strain ε¼�ΔG/G includingthermal expansion εT, chemical shifts εc (Vegard's law) and elasticstrain offset εe. At a constant temperature (573 K), carbon concentra-tion (χc) is a linear function of εc [16,17], given by

Δχc¼76.25 εc (mass%)

A comparison between the FWHM and the shapes of α-200 andγ-200 was made to calculate the carbon distribution, straindistribution and coherent lattice volume size.

The microstructure of the sample was characterised using TEM.Thin foils were prepared by twin-jet electro-polishing usingsolution of 5% of perchloric acid in methanol at 248 K andoperating voltage of 50 V. Bright and dark field images andselected area electron diffraction patterns were obtained using aPhilips CM 20 microscope operated at 200 kV.

The dislocation density of the bainitic ferrite was calculated using

^¼2NL/Lt

Where, NL is the number of intersections with dislocations, L isthe length of random lines and t is the foil thickness. The totaldislocation line length in a unit volume of crystal giving a parameterin terms of length (m) in m�3 was measured. The foil thickness (t)was determined from intensity oscillations in the two-beam con-vergent beam electron diffraction (CBED) patterns [7]. Four brightand dark field images were used at the magnification of 88,000.

APT analyses were carried out using the Oxford nano-Science 3DAPat the Monash Centre for Electron Microscopy in ultrahigh vacuum(10�8 Pa). A pulse repetition rate of 20 kHz and a pulse fraction of0.2 were used. The sample temperature was 80 K. Atom probe needlespecimens were cut using the wire-cut method. A standard two-stageelectro-polishing procedure was used to prepare the atom probe tipsusing 33% nitric acid in methanol for the first stage, followed by 2%perchloric acid in butoxy-ethanol at 16 V. Data sets from the experi-ment were reconstructed using PoSAP (Position Sensitive Atom Probe)software. The software displays both the mass spectra and a 3-Dvisualisation of the analysed volume. The maximum atomic mass thatcan be analysed by PoSAP is 300 amu giving the capability of analysinghigh mass complex ions [18].

3. Results

3.1. Neutron diffraction

Using reciprocal space Q as the independent variable, azimuth-ally integrated intensities of first 10 peaks were plotted against

Table 1Alloy steel composition.

C Si Mn Cr Mo Al Co

wt% 0.79 1.51 1.98 0.98 0.24 1.06 1.58at% 3.49 2.86 1.9 1.0 0.13 2.08 1.42

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time and temperature profile for an overall illustration of theevolution of different phases during the heat treatment cycle(Fig. 1). The initial pattern at room temperature consisted of ferritephase only. At the beginning of the heating, a shift in theα-110 peak position was observed towards a lower scatteringvector due to an increase in the lattice size because of the thermalexpansion. During heating to 1093 K, a narrow γ-111 peak initiallyappeared at around 1023 K and became stronger at the expense ofα-110, which gradually diminished. On rapid cooling, there was asudden shift in γ-111 to higher Q value due to the thermalcontraction (i.e. a reduction in the lattice size). Onset of thebainitic transformation was marked by the reoccurrence of theα-110 peak. It took 16.7 min for the reappearance of α-110 whichmarked the onset of bainitic transformation after reaching theisothermal holding temperature. The start of isothermal holdingtreatment at 573 K was taken as the zero time for reference.

The area under the peak, representing the intensity, for α-110was fitted using voigt function and the ratio taken over intensity of100% ferrite was plotted as volume fraction of bainitic ferrite overtime (Fig. 2). The equation for the kinetics of the transformationthrough volume fraction progression was further fitted to theAvrami Equation [1� f¼exp(�ktn)] and the parameter ‘n’ wasfound to be 1.7. Hence, the transformation kinetics was identifiedas nucleation-controlled as the parameter ‘n’o2. The volumefraction of bainitic ferrite at the end of transformation wascalculated to be 64% (Fig. 2). The maximum solubility of carbonin bainitic ferrite phase is 0.27 at% [19]. Thus, the remaining 36%

austenite should have 9.21 at% carbon, in the absence of cementite,by the law of mixtures.

During the low temperature isothermal heat treatment auste-nite peaks became wider and asymmetric with the averageposition shifting towards lower scattering vectors (Fig. 3a). Theshift in peak position of γ-200 was attributed to the increase inlattice size due to the rejection of carbon from bainitic ferrite andits redistribution in the remaining austenite (Fig. 3b). The asym-metry in γ-200 was attributed to the occurrence of more than onepopulation of austenite with different carbon contents duringtransformation. As evident in Fig. 3b, the asymmetry increasedwith time, with the increase in intensities of higher carbonaustenite with time. Towards the end of the transformation thepeak becomes symmetric with an evident shift in position. Carbonenrichment of austenite at the end of transformation, calculatedfrom peak position changes using Eq. (1) was 1.470.2 mass%(6.0370.9 at%). Comparing γ-200 and instrument LaB6 function,the FWHM of γ-200 was proportional to the scattering vector G, sothe strain broadening, ε¼�ΔG/G¼12�10�3 adds up to thebroadening due to carbon content gradient.

The ferrite peaks α-110 and α-200 were observed to reoccurafter an incubation period of 16.7 min. They were, however, muchbroader than those observed at room temperature (Fig. 3c). Thepeak intensities and shape of ferrite peaks became almost constantat around 750 min from the start of the experiment. The Lorent-zian component of the Voigt function, which depends on themicrostructure refinement and heterogeneous strain gives rise tothe peak broadening. Austenite and ferrite peak broadening havebeen observed in previous studies [11–13] and attributed to therefinement in the microstructure. We have made an attempt toquantify the peak broadening in terms of size and strain effects.FWHM from LaB6 calibration sample were taken as instrumentalfunction. For the ferrite phase, there was a constant offset of α-200with respect to the instrumental function suggesting a dominatingsize effect on the FWHM. The peak broadening was quantified to avery fine structure of 2π/ΔGE182 nm and further verified usingtransmission electron microscopy.

3.2. Microstructural characterisation using transmission electronmicroscopy

The microstructure after 2 days at 573 k mainly consisted oflamella structure with the layers of bainitic ferrite and retainedaustenite (Fig. 4). The bainitic ferrite and retained austenite layersformed the colonies or packs. The layers within each pack wereoriented in the same direction. The bainite colonies grew from theprior austenite grain boundaries and had different orientationswithin the prior austenite grain. The average thickness of thebainitic ferrite and retained austenite layers were 118740 and

Fig. 1. Integrated Peak Intensities: Azimuthally integrated peak intensities displaying heating of the sample to austenite phase (1093 K), rapid quenching to 573 K andisothermal holding at 573 K.

10001001010

10

20

30

40

50

60

70

80

Bain

ite v

olum

e fra

ctio

n (%

)

log time (min)

Fig. 2. Kinetics of phase transformation calculated from changes in intensity ofα-110: Volume fraction of bainitic ferrite calculated from in-situ neutron diffraction.The plot follows the Avrami equation with the constant ‘n’ as 1.7 depicting anucleation controlled transformation mechanism.

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Fig. 3. Phase evolution: Phase evolution displayed by peak characteristics at different time intervals starting at the beginning of isothermal holding: (a) characteristics of first4 peaks changing with time, (b) peak shift and peak asymmetry evident in austenite γ-200 with time, (c) some peak asymmetry but no peak shift observed in ferrite α-200.

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60720 nm respectively. Some carbide particles were alsoobserved in the microstructure by TEM.

3.3. Atom probe tomography characterisation after heat treatment

The bainitic ferrite and retained austenite phases at the end ofthe heat treatment were also identified by APT on the basis ofcompositional analysis. A volume of 11.13�11.13�113.40 nm3 wasanalysed using PoSAP software (Fig. 5a). The composition ofretained austenite was calculated using the matrix calculationmethod for a selected box in the austenite phase. An averagecarbon content of 6.9570.1 at% in the austenite film was directlycomparable to the neutron diffraction results for carbon contentcalculation. The average carbon content in the super-saturatedbainitic ferrite phase was found to be approximately 0.25 at% andthe carbon atoms were distributed unevenly in this acicularmicrostructure. The presence of scarce carbides and Fe–C clusterswere detected inside the bainitic ferrite laths. Fig. 5a shows plateshaped Fe3C carbide containing �656 ions in the bainitic ferritephase with a thickness of �2 nm and diameter 3.71 nm. Fig. 5dshows the carbon composition profile of the selected box with theparticle. The carbon content of retained austenite depends on boththe volume fraction of each phase and the presence of carbide andclusters.

4. Discussion

Firstly, in the present neutron diffraction study, the accelerateddata acquisition rate of 13 s/scan helped to minutely detect anyprecursor events taking place in the austenite phase before theonset of the bainitic phase transformation. No peak shift or peakbroadening was noticed in γ-111 before initiation of the bainiticphase transformation, suggesting no carbon partitioning takingplace during the incubation time in the current experiment, whichcontradicts the results reported by Babu et al. [11]. Further thechanges in the γ-200 peak with time were a good evidence of themechanism of bainitic phase transformation, especially in terms ofcarbon redistribution.

It has been speculated that, as the volume fraction of bainiticferrite increased, two populations of austenite started to coexist inthe microstructure: one adjoining the ferrite phase with highercarbon content and the other far from the transformed ferrite withthe carbon content close to the parent austenite. Asymmetry in the

austenite peaks is due to these two populations of austenite.Because of a wider range of carbon content in the austenite, thepeak also becomes wider increasing the FWHM (Fig. 3b). Theincrease in FWHM can also be attributed to the refined size of thenew population of austenite and the strain induced in austenitebecause of the shear transformation in bainite. It is not peculiar tonote here that at around 744 min from the start of the transformationγ-200 becomes symmetric again, suggesting a uniformity in carboncontent and hence completion of the bainitic transformation.

A very valuable inference from this study which can be used forfurther neutron diffraction studies is the quantification of theFWHM values of the α-200 peak to the refinement of themicrostructure. When compared to the instrument function, thelath size was expected to be 182 nm, assuming that the broad-ening was only due to refinement while internal stress did notplay any role. With the TEM images revealing the true size of thebainitic laths to be �120 nm, this difference can be directlyattributed to strain due to the dislocation density, as the bainitictransformation is primarily a displacive transformation. There is agreat scope for the quantification of peak asymmetry and broad-ening directly to the internal stress in such steels.

It was validated though neutron diffraction results that thecarbon content calculated in retained austenite at 744 min, usingchanges in the peak position was 6.0370.9 at% which is muchlower than expected in a 36% volume fraction of retained austenitei.e. 9.21 at%. To account for the missing carbon and followingprevious studies [6,7], APT was done for a nano-bainite structure,formed under similar conditions. From the results of APT it wasconfirmed that despite the Si level in the steel, carbide formationcould not be completely suppressed in the bainitic ferrite phase.This carbide can form concurrently with the transformation whichcan be termed as auto-tempering or through ageing over time.Initiation of carbide formation might not be proven yet but theunfavourable effects are quite evident in the mechanical proper-ties. It is important to bear in mind that an optimal austenitestabilisation is the basis for the TRIP effect, leading to a desirableductility in the steel. The carbon content of retained austenitewere also compared to APT and an average value of 6.95 at%matched the result obtained by neutron diffraction, further sup-porting the existence of carbon in cluster and carbide form. Thekey factor affecting the stability of austenite is the carbon contentof retained austenite and, unless we can control the carboncontent of austenite, it would be very difficult to control theproperties of nanostructured bainitic steel.

Fig. 4. TEM image: Bright TEM images of the bainitic ferrite packs and bainitic ferrite and retained austenite layers within the pack.

K. Rakha et al. / Materials Science & Engineering A 589 (2014) 303–309 307

It has been well established now through various studies thatpresence of Si does not invalidate the formation of carbides andclusters in low temperature bainitic steels. It is now important tofind the true value of carbon content in different phases to be ableto design and exploit this promising new nanostructured material.

5. Conclusion

In-situ neutron diffraction was employed to monitor theevolution of nano-bainitic ferrite during low temperature isother-mal heat treatment of austenite. This was done to throw morelight on the much argued transformation mechanism of nanos-tructured bainite at relatively low temperature. The peak char-acteristics were quantified to changes in lattice parameter, volumefraction and carbon content for evidence of carbon partitioningduring transformation. A former debate of [11–13] on the parti-tioning of carbon in austenite during the incubation period wassupported in favour of no such evidence in the present study.

Mechanism of nanostructured bainite formation was discussedand Neutron Diffraction was used for the first time as evidence forthe formation of carbides despite the high level of Silicon in thecomposition. APT was further used to verify the above. Peakbroadening in α-200 was also quantified in terms of size effectand compared with the TEM characterisation. In-situ NeutronDiffraction has been established as a valuable characterisationtechnique for further studies.

Acknowledgement

The authors would like to acknowledge the support of theBragg Institute, Australian Nuclear Science and TechnologyOrganisation, in providing the high-intensity neutron powderdiffractometer WOMBAT used in this work. The sample environ-ment team, especially Stewart Pullen and Sean McTrusty togetherwith the team build the Rapid Sample Quencher which was criticalto our experiments. The authors would also like to thank AINSE

Fig. 5. APT: APT characterisation of austenite and bainitic ferrite phases: (a) carbon atommap showing carbon distribution along the analyses needle, (b) corresponding 6.95 at%isoconcentration surface, (c) carbon concentration profile along the z-axis of the needle analysed, (d) carbon concentration profile along the z-axis of the selected area in (a).

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Ltd. for providing financial assistance (Proposal ID 1748) to enablework to be conducted at ANSTO. This work was financiallysupported by grants through the Australian Research Councilincluding an ARC Laureate Fellowship (P.D.H.).

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