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Ironmaking & SteelmakingProcesses, Products and Applications

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Structure and properties of κ-carbides in duplexlightweight steels

W. J. Lu, X. F. Zhang & R. S. Qin

To cite this article: W. J. Lu, X. F. Zhang & R. S. Qin (2015) Structure and properties of κ-carbides in duplex lightweight steels, Ironmaking & Steelmaking, 42:8, 626-631, DOI:10.1179/1743281215Y.0000000007

To link to this article: http://dx.doi.org/10.1179/1743281215Y.0000000007

Published online: 24 Apr 2015.

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Structure and properties of k-carbidesin duplex lightweight steels

W. J. Lu, X. F. Zhang and R. S. Qin*

k-Carbide in lightweight steel has been studied. Its thermal stability, crystal morphology,

orientation relationship, degree of lattice misfit and mechanical properties have been measured

experimentally. The mechanisms for the microstructural evolution of k-carbide are considered

based on the crystal structure, lattice misfit, element diffusivity and solute partition. The hard c/k

grain and ductile a þ k grain provide an opportunity to develop lightweight steels with a desirable

combination of mechanical properties.

Keywords: Lightweight steel, k-Carbide, Thermal stability, Lattice misfit, Morphology

IntroductionAdding aluminium and manganese to steels helps toreduce their mass density, and the microstructural andmechanical properties of lightweight steels change with Aland Mn compositions accordingly.1,2 The formation ofk-(Fe,Mn)3AlC carbide (so called k-carbide) is promotedby the addition of Al.3,4 Lightweight steels are found withferrite based,3,5 austenite based3,6,7 and ferrite–austeniteduplex basedmicrostructures.3,8,9 k-Carbide is observed inevery type of high Al lightweight steels,3,10 and it plays animportant role in affecting themechanical properties of thesteels.11 (Fe,Mn)3AlC type carbide has a cubicperovskite crystal structure in which Al occupies cornerpositions, Fe and Mn occupy face centred positions, andcarbon is at the body centre interstitial octahedral site.1

This is different from that of c-austenite face centred cubic(fcc) crystal and a-ferrite body centred cubic (bcc) crystalphase in steels in terms of atom configuration.

The growth of the same crystals from different crystal-lographic matrix provides an ideal situation to investigatethe physics of pattern formation in patterned back-grounds.12 Therefore, the matrix/k interface is stronglydependent on the matrix structure. Normally, the c/kinterface is coherent, whereas the a/k interface is semi-coherent, suggesting that lattice mismatch between matrixand k-carbide depends on their interface. Interface aniso-tropy plays an important role inmicrostructural evolution,but the derivation of interface anisotropy from surfaceanisotropy is not a trivial task.13 For lattice misfit strain, itcan be measured by X-ray diffraction or selected areadiffraction of transmission electronmicroscope. Thus, onecan investigate the microstructural evolution based onmeasured lattice misfit strain. The study helps tounderstand transformation kinetics, interface propertiesand microstructure–property relationship. k-Carbide canform in bcc ferrite phase and fcc austenite phase in duplex

lightweight steel simultaneously. Systematic andquantitative studyof thek-carbide structure andpropertiesis interesting to academia and engineering.

ExperimentalA steel (800 g) with chemical composition of 9 wt-% Al,34 wt-% Mn and 0.65 wt-% C was produced using aninduction melting furnace. The size of steel ingot wasnominally 80|50|50 mm (length|width|height).A number of plates of 10 mm thickness were cut from themiddle of casts and then heat treated at 1200uC for 4 h inorder to homogenise the chemical compositions beforequenching in cold water. The plates were subsequentlyannealed at different temperatures (600, 700 and 800uC)for 6 h to allow the development of microstructure beforequenching into coldwater. Specimensweremade from thequenched plates, ground, polished and etched with 2%Nital. The grain morphology and chemical compositionwere analysed using scanning electron microscope (SEM)equipped with energy dispersive X-ray spectroscopy(EDX). Phases were identified by X-ray diffraction(XRD)with a scan rate of 3u perminute and a scan step of0.02u. The volume fraction of each phase was calculatedby integrationof all the diffractionpeaksof the phase. Themicrostructure and crystal structure of the phases in thespecimens were further characterised using the trans-mission electron microscope (TEM) with an accelerationvoltage of 200 kV. The TEM specimens were prepared bypunching a 30 mm thick steel plate into 3 mm diameterdisc and then polished using a twin jet electropolisherwitha solution containing 10% perchloric acid and 90% aceticacid at 15uC. The Vickers hardness was measured using1 kg load. The crystallographic structures and the latticemisfit between phases were calculated and plotted usingthe commercial software Crystal Maker.

Results and discussionThe SEM images of two different resolutions for thespecimen annealed at 600, 700 and 800uC are presentedin Fig. 1a–f respectively. The precipitation of Al,

Department of Materials, Imperial College London, Exhibition Road,London SW7 2AZ, UK

*Corresponding author, email r.qin@imperial.ac.uk

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� 2015 Institute of Materials, Minerals and MiningPublished by Maney on behalf of the InstituteReceived 20 November 2014; accepted 1 February 2015DOI 10.1179/1743281215Y.0000000007 Ironmaking and Steelmaking 2015 VOL 42 NO 8

Mn and C from the bulk phase at the temperature causesthe formation of k-carbide. In the steel annealed at600uC, as illustrated in Fig. 1a and b, k-carbide can befound in a phase with an average size of 0.17 mm, butnot in c phase. The reason for this is thought due to thelow diffusivity of the elements in c phase than that in aphase. For example, it is reported in literature that theratios of diffusivities of carbon in a to that in c phasesare 629 at 730uC, 5493 at 750uC and 30 795 at 800uC.14

The high element mobility in a phase enables theformation of nanoscale k-carbide at 600uC. However,the annealing time is not sufficient long to allow theformation of k-carbide in c phase. In the steel annealed

at 700uC (Fig. 1c and d), k-carbides form in both a and cphases. The k-carbide in a phase is coarsening to*2.62 mm in thickness. In some of the c phasegrains, nanoscale k-carbides are observed. Figure 1eand f shows the microstructure of steels annealed at800uC. k-Carbides in a phase are further coarsening to*17.2 mm in thickness. However, no k-carbide isfound in c phase. This shows that the thermal stability ofk-carbide in a phase is better than that in c phase inthis steel.

High resolution SEM and TEM were used to furtheranalyse the structure and properties of k-carbide in thesteel annealed at 700uC. The results are presented in

1 SEM images for steels annealed at a, b 6008C, c, d 7008C, and e, f 8008C for 6 h

Lu et al. k-Carbides in duplex lightweight steels

Ironmaking and Steelmaking 2015 VOL 42 NO 8 627

Fig. 2. Figure 2a and b shows the SEM and TEM brightfield images for the k-carbide in c phase. It is atypical microstructure formed in spinodal decompo-sition (e.g. the decomposition is determined solely byuphill diffusion but does not experience the nucleationstage). Figure 2c and d shows the SEM and TEM brightfield images for the k-carbide in a phase. The micro-structure is typical of nucleation/growth controlled bydiffusion. TEM diffraction pattern from the selected c/karea is demonstrated in Fig. 2e. The calculations show alattice distance of k-carbide of 3.75 A and that of cphase of 3.66 A. k-Carbide has a cube on cube orien-tation relationship with the c matrix with (001)k//(001)cand [001]k //[001]c. The relationships are identical to thatbetween c and c9 in Ni based superalloys.15,16 Fig. 2f

presents the TEM diffraction pattern from the selecteda þ k area. The lattice distances of k-carbide and aphases are 3.76 and 2.87 A respectively. The orientationrelationships between k-carbide and a matrix grain are(111)k //(110)a and [011]k //001]a. This obeys theNishiyama–Wasserman relationship and is in agreementwith report in literature.17

Figure 3apresentsXRDresults for all the steels annealedat different temperatures. The volume fractions of c, a andk phases are calculated according toFig. 3a and are plottedin Fig. 3b. It should be pointed out that the accuracy givenin Fig. 3b was based on image analysis, while the accuracyofusingXRDtoget the illustrated imageswas not includedin this analysis. There should be some fluctuation betweenmeasurements. The fraction of c phase is found to decrease

2 Experimental results for steel annealed at 7008C for 6 h: a SEM image of fine k-carbide in c phase; b TEM bright field image of

fine k-carbide in c phase; c SEM image of lamellar k-carbide in a phase; d TEM bright field image of lamellar k-carbide in a

phase; e TEM diffraction pattern from selected c/k area; f TEM diffraction pattern from selected a 1 k area

Lu et al. k-Carbides in duplex lightweight steels

628 Ironmaking and Steelmaking 2015 VOL 42 NO 8

from 52.9% at 600uC to 45.6% at 700uC. The fraction of aphase changes from 40.1% to 31.3%, and that of kphase changes from 7.0% to 23.3% in the temperature.k-Carbide presents in both a and c grains at 700uC. TheVickers hardness of each samplewasmeasured 10 repeated

times on both a and c phase areas. The average values andtheir standard deviation are illustrated in Fig. 3c. The areawith c phase containing significant amount of k-carbidehas the highest hardness (704+23 HV). The a phaseannealed at 700uC has a Vickers hardness of 526+20 HV.

3 a XRD spectrum for samples annealed at different temperatures; b calculated volume fractions of a, c and k phases;

c Vickers hardness

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Ironmaking and Steelmaking 2015 VOL 42 NO 8 629

For the steel annealedat 800uC,k-carbide freecphase hasaVickers hardness of 328+16 HV only, which is lower thanthat of a phase containing coarsening k-carbide of348+14 HV.

It has been noticed that some of the c phase grainshave none but others have k-carbides for the steelannealed at 700uC. For an area arranged with a þ kgrain, c/k grain and c in a row, as demonstrated inFig. 4a, EDX measurements were implemented to scanacross the region linearly to characterise the elementcompositions, as demonstrated in Fig. 4b. a Phasecontains high Al and low Mn than does c phase. The cgrain containing k-carbides has higher Al and lowMn than the c grain with free k-carbides. Al promotesk-carbide formation and stabilises k-carbide presen-tation.4 C is another element that promotes and stabil-ises k-carbide.18 High C diffusivity in a phasecontributes to k-carbide nucleation at low tempera-ture.14 Low C solubility in a phase may help to stabilisethe k-carbide at high temperature. k-Carbide containsvery high concentrations of C and Al. The diffusion of Cand Al from the matrix to the k-carbide determines itsgrowth rate. Mn and Fe have low diffusivity than C andAl. It should play a control role in the k growthaccording to the conventional diffusion limited growthmechanism. However, the compositions of Mn and Feare sufficiently high in the present samples, and hence,

their diffusions are not crucially important. Sato et al.observed that fine k-carbide precipitation required thediffusion of both C and Al in austenite during anisothermal aging treatment.18 Sutou et al. revealed thatthe fine k-carbide can be formed by rapid cooling fromsolution treatment via the diffusion of Al and C in theaustenite.22 Seol et al. presented that the fast diffusion ofC and high content of Al partition are the key factors toform lamellar k-carbide.23

The Crystal Maker software was implemented to plotthe unit cells of c phase, a phase and k-carbide. A modelis applicable to predict the lattice misfit of thesephases.19 According to the orientation relationshipsbetween phases revealed by TEM (Fig. 3e and f),a 2|2|2 c unit cells in (001) plane along [001] direc-tion has been constructed. This can be overlapped with a2|2|2 k-carbide unit cells in (001) plane along [001]direction. The results are plotted in Fig. 4c, where theblack solid circles represent c phase and red open circlesrepresent k-carbide. The lattice misfit along [001]k//[001]cis calculated to be 2.70%. Similarly, a 2|2|2 a unitcells in (110) plane along [001] direction and a 2|2|2k-carbide unit cells in (111) plane along [011] directionhad been constructed and overlapped, as demonstratedin Fig. 4d. The blue solid circles represent a phase, andred open circles represent the k-carbide. The calculatedlattice misfit along [011]k//[001]a is 6.67%.

4 a SEM image for area consisting of k 1 a, k/c and c phases; b chemical compositions across area; c planar lattice

matching diagrams for 26262 cells between k-carbide and c phase; d planar lattice matching diagrams for 26262 cells

between k-carbide and a phase

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630 Ironmaking and Steelmaking 2015 VOL 42 NO 8

The low lattice misfit between k-carbide and c grain isable to produce excellent mechanical properties.1,20,21

On the other side, the low lattice misfit means smallelastic strain energy. The strain energy is the kineticbarrier in first order phase transition. The small strainenergy enables the spinodal decomposition at smallundercooling. Thus, nanoscale phase can be formedunder low lattice misfit strain in spinodal decomposition,as shown in Fig. 2a. The strength of c/k is therefore high.On the other hand, the high lattice misfit betweenk-carbides and a phase made the strain energysignificant. The large misfit brings down the mechanicalproperty of k þ a phase. The high strain energy is a bigbarrier for k-carbide nucleation. Spinodal decompo-sition is therefore not possible in such a case. However,the high diffusivity in a phase accelerates the diffusionalgrowth. The observed k-carbide microstructure in aphase is therefore coarse at high temperature.

Choo et al. reported an overall phase transformation ofcRco þ k9Ra þ k in the aging of austenitic Fe–30Mn–7.8Al–1.3C steel at 550uC for 4 months,1 where k9 meansthe fine k-carbide, which is just called k-carbide in manyliteratures.15,22 The second transformation from co þ k9 toa þ k requestsmuch longer annealing time than that in thepresent work.

ConclusionsThe microstructure of Fe–34Mn–9Al–0.65C steelsannealed at 60, 700 and 800uC has been studied.The volume fractions of a, c and k phases at thetemperature range have been characterised. k-Carbides aremore stable in a phase than in c phase. The latticemisfit between k-carbide and c is low (2.70%). This leadsto low strain energy and hence enables spinodaldecomposition, which leads to the formation ofnanoscale phase. The mechanical property of c/k phase ishigh due to the fine microstructure and low misfit. Thelattice misfit between k-carbide and a is high (6.67%). Thegrowth mechanism for k-carbide in a phase is via nuclea-tion rather than spinodal decomposition. The diffusivity ofC inaphase causes significant coarsening ofk-carbide. Theoverall mechanical property of duplex lightweight steel isdetermined by the microstructure of a, c and k phases.Designing the thermomechanical processing in order togenerate the suitable microstructure can lead to the

manufacture of lightweight steels with desirable combi-nation of mechanical properties.

AcknowledgementsThe authors gratefully acknowledge the financial supportfrom TATA Steel, Royal Academy of Engineering andEPSRC (grant no. EP/J011460/1) at United Kingdom.

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