Third generation of AHSS with increased fraction of retained austenite for the automotive industry

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Original Research Article

Third generation of AHSS with increased fraction of retained austenite for the automotive industry

A.Grajcar a ,n ,R.Kuziak b ,W .Zalecki b

a

Silesian University of Technology,Faculty of Mechanical Engineering,Institute of Engineering Materials and Biomaterials,18a Konarskiego Street,44-100Gliwice,Poland b

Institute for Ferrous Metallurgy,12-14Karola Miarki Street,44-100Gliwice,Poland

a r t i c l e i n f o

Article history:

Received 30May 2012Accepted 17June 2012Available online 23June 2012Keywords:TRIP-aided steel AHSS

DCCT diagram Stress–strain curves Retained austenite

a b s t r a c t

Third generation of advanced high-strength steels for the automotive industry contains a high volume fraction of ?ne-grained ferrite,carbide-free bainite,martensite and retained austenite.The level of strength and ductility is highly dependent on the fraction and mechanical stability of austenitic phase.One of the methods to obtain an increased fraction of g phase is trough its chemical stabilization by Mn.T wo 0.17C–3Mn–1.5Al–0.2-Si–0.2Mo steels with and without Nb microaddition were developed in the present study.The steels were subjected to the thermomechanical processing designed on the basis of the DCCT diagram (deformation –continuous cooling transformation).The paper presents the results of the multi-stage compression tests and multiphase microstructures obtained as a result of the controlled multi-stage cooling.It was found that the hot workability of a new generation of AHSS is very challenging due to high values of ?ow stresses required.However,the thermomechanical processing enables to obtain very ?ne-grained bainite-based microstructures with a fraction of retained austenite up to 20%.The highest fraction of ?ne grains and interlath austenite was obtained for the temperature range between 400and 4501C.The effect of Nb results in higher ?ow stresses and better distribution of all the microstructural constituents.

&2012Politechnika Wroc?awska.Published by Elsevier Urban &Partner Sp.z o.o.All rights

reserved.

1.Introduction

Third generation of advanced high strength steels (AHSSs)for the automotive industry is a further step in development of modern high strength–ductility balance steel sheets.They utilize complex interaction of solid solution hardening,pre-cipitation hardening,microalloying,grain re?ning,strain aging and TRansformation Induced Plasticity (TRIP)or TWin-ning Induced Plasticity (TWIP)effects.Third generation AHSS

combine the advantages of multiphase structures (character-istic for 1st generation AHSS)[1–6]and austenitic phase,especially suitable to enhance different hardening mechan-isms,strain-induced martensitic transformation and mechan-ical twinning (characteristic for 2nd generation AHSS)[7–11].In spite of excellent mechanical properties of high-Mn austenitic steels belonging to the 2nd generation of AHSS their applica-tion will be probably limited only for most challenging auto-body elements,for example,specially geometrically-designed

1644-9665/$-see front matter &2012Politechnika Wroc?awska.Published by Elsevier Urban &Partner Sp.z o.o.All rights reserved.https://www.360docs.net/doc/c810591661.html,/10.1016/j.acme.2012.06.011

n

Corresponding author.

E-mail address:adam.grajcar@polsl.pl (A.Grajcar).

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sections absorbing energy during crash events[12].The rea-sons are various technological problems related to poor cast-ing,hot-working above11501C,corrosion resistance,Mn segregation and especially the high cost due to Mn(between 20and30wt%),Al and Si alloying concept[7–11].

The main idea of3rd generation of AHSS is to obtain the mechanical properties regime between1st and2nd generation of AHSS at cost slightly higher compared to1st generation multiphase steels(DP,TRIP,CP)[13–15].New microstructure concepts consist in increasing volume fraction of hard consti-tuents and retained austenite.It is related to replacing poly-gonal ferrite by acicular or bainitic ferrite,non-carbide bainite, martensite and stabilization of austenitic phase in different ways,i.e.,chemically or mechanically.Replacing the polygonal ferrite matrix by non-carbide bainite/martensite improves also the stretch-?angeability of multiphase steel sheets,which is one of the most important technological parameter character-izing their practical application for the auto-body structure [13,16].One of the chemical strategy concepts to obtain a bainitic matrix containing a high volume fraction of metastable retained austenite is Mn alloying but to a much lesser extent than in high-Mn steels of the2nd generation AHSS.Manganese is a main austenite stabilizer and its content reaches up to 12wt%in recently investigated C–Mn–Al–Si steels[13–15]. Increased hardenability of steel due to Mn alloying leads to a considerable decrease in the ferrite fraction as a result of shifting the g-a transformation region to the right side on CCT diagrams[17,18].A heat treatment for steels with increased Mn content is similar to that used for DP and TRIP steels,i.e.,intercritical annealing after cold rolling or intercritical annealing and isothermal holding at a bainitic transformation range.The thermomechanical processing, microalloying with Nb,Ti,V,reverse martensite transforma-tion and quenching and partitioning processing are the examples of ideas to obtained?ne-grained complex micro-structures with a high fraction of austenitic phase of optimal stability for strain-induced martensitic transformation dur-ing drawing,stretching,bending,etc.[7,13–15].

In our previous works[19,20]the stress–strain curves and softening kinetics of austenite under conditions of hot-compression for3Mn–1.5Al and5Mn–Al steels were investi-gated.Elaboration of the thermomechanical rolling requires also the knowledge of hot-working behavior of steels during multi-stage deformation.Additionally,especially important for hot-rolled multiphase steels are CCT(Continuous Cooling Transformation)diagrams,which are necessary for proper design of multi-stage cooling after?nishing rolling.Despite many DP and TRIP steels investigated,very few CCT diagrams can be found in the literature[17,18,21–24].Silicon and aluminum increase A c3temperature of steel.In addition,Si shifts pearlitic transformation to the left,however C,Mn,Al, Mo,Cr,Ni and Nb act in the opposite direction.The effect of Nb usually manifests by slight increase of the g-a transfor-mation temperature and shortening time to initiate the transformation[17,18].

2.Experimental procedure

The paper addresses the thermomechanical processing of new-developed high-Mn high-Al TRIP steels with and without Nb microaddition.Special attention was paid to the effect of Nb on the hot-working behavior and multiphase microstruc-ture formed at various cooling conditions of the thermome-chanical treatment.The chemical compositions of the investigated steels are given in Table1.

The chemical composition strategy was designed with the focus on the maximizing of retained austenite volume frac-tion(increased Mn content)and obtaining carbide-free bai-nite by low-Si high-Al concept[2,3,23].Special attention was paid to the effect of Nb on the hot deformation resistance and ?nal multiphase microstructures after the thermomechanical processing.The steels were produced by vacuum induction melting in the Balzers VSG-50furnace.Liquid metal was cast in the Ar atmosphere into a cast iron mould.Ingots with a mass of about25kg were forged at temperature range from1200to 9001C to a thickness of22mm.Then,cylindrical samples |10?12mm for hot compression tests and|4?7mm speci-mens for dilatometric analyses were machined.

Parameters of the thermomechanical processing are shown in Fig.1.The experiments were carried out using the DSI Gleeble3800thermomechanical simulator.The specimens were inserted in a vacuum chamber,where they were resis-tance-heated to a temperature of12001C.After austenitizing for30s the specimens were cooled to a temperature of?rst deformation.The thermomechanical processing consisted of four deformation steps(1150,1050,950,8501C)and multi-stage cooling according to Fig.1.The logarithmic strain value was equal to0.25at the strain rate of10sà1for each deformation step.After the?nal deformation at8501C the specimens were cooled to the temperature of7001C.The essential step of the thermomechanical treatment consisted in the various time of slow cooling in a temperature range700–6501C.The time was equal to10,25or50s depending on the cooling rate applied (Fig.1).Next,the samples were cooled at the rate of401C/s (similar to the laminar cooling facilities in industrial condi-tions)to the holding temperature in the bainite transformation range(T B).The isothermal holding temperature of450,400or 3501C was the second variable parameter of the thermome-chanical processing.The holding time(t B)was constant,300s. Finally,the specimens were colled with a rate of0.51C/s to room temperature.The whole cooling route was designed on the basis of the DCCT diagram.

Table1–Chemical composition of the investigated steels,wt%.

Designation C Mn Al Si Mo S P Nb

3Mnà1.5Alà0.2Sià0.2Mo0.17 3.30 1.700.220.230.0140.010–

3Mnà1.5Alà0.2Sià0.2MoàNb0.17 3.10 1.600.220.220.0050.0080.04

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Critical temperatures A c 3and A c 1of steels and continuous cooling transformation diagrams after hot deformation (DCCT diagrams)were determined using the DIL 805A/D dilatometer with the LVDT gauge head (Baehr Thermoanalyse GmbH).The cylindrical specimens were heated at a rate of 31C/s to the austenitizing temperature of 11001C and held for 300s.Subsequently,the samples were cooled at a rate of 51C/s to the deformation temperature of https://www.360docs.net/doc/c810591661.html,pression of cylindrical samples was conducted to the logarithmic strain of 0.5and at the strain rate of 1s à1.Vickers hardness of thermo-mechanically processed samples and these cooled in a range from 100to 0.51C/s was measured at the load of 98N.Microstructure observations of nital-etched specimens were carried out at magni?cation of 400?.For the best identi?ca-tion of retained austenite,etching in 10%aqueous solution of sodium metabisul?te at room temperature was applied for the thermomechanically processed specimens.Metallographic investigations were carried out by the use of LEICA MEF4A optical https://www.360docs.net/doc/c810591661.html,puter image analysis was used to evaluate metallographically a fraction of retained austenite on the basis of ?ve binary maps for each variant.

3.Results and discussion

A result of the four-step deformation at the temperature range between 11501C and 8501C are the stress–strain curves in Fig.2.The stress values increase signi?cantly with decreasing deformation temperature.The registered temperature increase due to applied deformation was very small (i.e.up to 41C),so stress correction was not applied.It appears on the basis of the curves that dynamic recovery is the process controlling work-hardening for the whole range of applied hot-working tem-perature.These results are in accordance with earlier tests of continuous compression,where initiation of dynamic recrys-tallization occurred only at the temperature of 11501C after the true strain about 0.5[19].The highest growth of stress occurs for the deformation temperature of 8501C,what is probably caused by a limited course of static recrystallization in the interval between third and fourth deformation.The time of 10s is enough to get only about 30%of recrystallized austenite for the applied temperature and strain rate conditions ($9001C,10s à1).The strain accumulation enhances the rate of a diffusive transformation from a pancake austenite during

subsequent cooling.Due to complex Mn–Al–Mo–Nb alloying the values of ?ow stress are slightly higher compared to conventional C–Mn–Si and C–Mn–Si–Al and C–Mn–Al TRIP steels [5,25–27].

The stresses of Nb-microalloyed steel are initially slightly higher compared to Nb-free steel and then the difference is much higher for the deformation temperature of 950and 8501C (Fig.2).The effect of Nb microaddition is known to retard the dynamic recrystallization and to increase the ?ow stress.The in?uence of Nb is rather lower than expected,especially for the two initial deformation steps.A reason is probably a masked effect of Nb due to the presence of high contents of alloying elements.Al alloying suppresses pre-cipitation of Nb(C,N)during hot-working due to strong af?-nity of Al to nitrogen [19,26].Additionally,it is well known [20,27]that increasing the Mn content in Nb-microalloyed steels delays the precipitation of Nb(C,N).

Taking into account retarding effects of Mn and Al on the precipitation of Nb(C,N)it can be expected that Nb is almost entirely dissolved in austenite above 10001C and rises the ?ow stress only by its solute-drag effect.This is probably why the stresses for the ?rst two deformation steps are just slightly higher for Nb-microalloyed steel compared to the basic steel.The difference in stresses between two steels for two successive deformation steps is much higher and reaches 20MPa (Fig.2).It is probably related to the initiation of Nb(C,N)precipitation because of a higher accumulated strain and longer total time for precipitation between successive deformation stages.

The diagram of phase transformations of supercooled plas-tically deformed austenite presented in Fig.3a is necessary to design multi-stage cooling of steels from the temperature of a ?nal deformation.The determined A c 1temperatures are equal 633and 6291C –respectively for the steel without Nb and for the Nb-microalloyed steel.The A c 3temperatures were not achieved because of a high aluminum content in the steels.It was also con?rmed by the calculations using JMatPro software where the maximum fraction of g phase under conditions of thermodynamic equilibrium occurs at the temperature of 1170and 11851C for Nb-free and Nb-microalloyed steels.For the austenitizing temperature of 11001C the fraction of ‘‘primary’’non-transformed ferrite amounts to 7%for 3Mn à1.5Al à0.2Si à0.2Mo steel and about 2%for the steel containing Nb.The micrographs in Fig.3b-e con?rm also the presence of some fraction of large round ‘‘primary’’ferrite grains.Martensite

start

Fig.1–Parameters of the thermomechanical processing for axisymetrical specimens.

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Fig.2–Changes of stress as a function of strain for successive deformation

stages.

Fig.3–DCCT diagram of the Nb-free steel (a)and microstructures obtained after cooling the specimens from the deformation temperature of 9001C with a rate of 1001C/s (b),151C/s (c),21C/s (d)and 0.51C/s (e).

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temperature of the Nb-microalloyed steel for the deformed austenite is equal to3531C and it is121C higher for the steel without Nb microaddition.

The investigated steels have high hardenability mainly due to the increased Mn content.In DCCT diagram(Fig.3a)it manifests by relatively low martensite start-and bainite start tempera-tures,a bainite region extended to short time and a right-displaced ferritic bay.There is not a pearlitic transformation for the cooling rates applied–what is especially important for TRIP steels.A lack of pearlite is caused by the effect of Mn,Al,Si and Mo.The DCCT diagrams are very similar for both steels,so the detailed analysis is limited only to3Mnà1.5Alà0.2Sià0.2Mo steel.The analysis of the DCCT diagram in Fig.3a and the microstructures obtained for various cooling rates(Fig.3b–e) indicate that the dominate constituent of the steels for the applied range of cooling rates is a mixture of martensite and bainite.Cooling the steel at a rate of approximately1001C/s results in obtaining the martensitic microstructure with hard-ness of about460HV(Fig.3b).The fraction of bainite increases with decreasing the cooling rate(Fig.3c,d)but the hardness decreases only a little amounting to about420HV for the cooling rate of11C/s.A fraction of ferrite(excluding‘‘primary’’ferrite)is very low,therefore a ferritic region in Fig.3a is plotted by the dashed line.Very small ferrite grains can be observed at boundaries of primary austenite for the cooling rates lower than101C/s but their volume fraction is marginal even for the lowest cooling rate applied(Fig.3e).A more detailed analysis of ferritic transformation using dilatograms and SEM deserves to be undertaken elsewhere.

Obtaining multiphase structures containing a high volume fraction of retained austenite is highly dependent on a ferrite fraction and a bainite morphology.These factors in?uence the redistribution of carbon during cooling and decide about austenite stability at room temperature.Cooling curves for ?ve developed variants of processing(Fig.1)were super-imposed on the DCCT diagram of the Nb-free steel in Fig.4. It can be seen that there are different conditions for ferrite forming depending on the time of slow cooling from7001C.It should be also noted that two isothermal holding tempera-tures are situated at a bainitic transformation region whereas the specimens cooled down to3501C are held below marten-site start temperature of steels.

As a result of the conducted thermomechanical processing bainitic–martensitic–ferritic microstructures were obtained with a high fraction of retained austenite.The realization of phase transformation under conditions of accumulated strain in austenite results in a signi?cant re?nement of all the microstructural constituents.It refers to bainitic–martensitic packets and to retained austenite as well.Additionally,a few percent fraction of ferrite can be identi?ed in Fig.4b–e.It has a morphology of very?ne equiaxial grains but more often it is acicular ferrite.A volume fraction of ferrite is higher for specimens cooled slowly in the g-a transformation range for 50s(Fig.4d and e).A few ferrite grains occurring at the microstructures of specimens cooled slowly for10s is a result of strain-induced transformation.Under conditions of repeated straining the ferritic region is probably left-displaced and ferrite forms easier and faster compared to deformation con-ditions applied for the determination of DCCT diagram.How-ever,a maximum portion of ferritic phase does not exceed5%,hence the distribution of carbon into austenite at this tem-perature range is marginal.The‘‘primary’’ferrite disappeared because of the higher austenitizing temperature.

A fraction of retained austenite varies from17.3to20.4%for the specimens isothermally held at4501C(Table2).Due to minor changes in the ferrite fraction there is also no distinct difference in retained austenite fraction for samples cooled slowly for10,25 and50s.Retained austenite is uniformly distributed in a bainitic–martensitic matrix and two main morphologies of this phase can be distinguished in Fig.4b–d–?ne granules and interlath retained austenite.A majority of?ne grains has a diameter of about1m m and they are usually located on ferrite–bainite phase boundaries or are a constituent of bainitic–martensitic–austenitic(B–M–A)regions.Interlath retained austenite belongs to bainitic regions and its growth is consistent with a growth of ferritic bainite laths.This bainitic morphology is very similar to the degenerate upper bainite, where cementite is replaced by thin austenite plates[28].Some larger austenite regions with a size up to8m m transformed into martensite during a?nal slow cooling.These martensite regions are indicated by arrows in Fig.4b–e.The M s temperature of this type of g phase is higher than room temperature and higher compared to?ne granules and interlath austenite which are thermally stable.Sometimes,there is a halo of untransformed austenite around martensite because outer regions of g phase are easier enriched in C(have lower M s temperature)than a middle of the grains due to a longer diffusion path of carbon. The difference in a retained austenite fraction in Fig.5is only dependent on the isothermal holding temperature in a bainitic transformation range.The distribution,volume fraction,size and morphology of retained austenite for the specimens held at 4001C(Fig.5c and d)are similar to those obtained for the highest holding temperature applied(Fig.5a and b).A higher fraction of interlath austenite compared to?ne granules can be only observed.Moreover,the bainitic regions have more distinct lath character.With further decrasing in isothermal holding tem-perature the morphology of a bainitic–martensitic matrix changes considerebly into clear laths(Fig.5e and f).At this temperature interlath austenite dominates but it is impossible to determine its volume fraction metallographically.A preliminary X-ray examination allowed to evaluate the fraction of austenitic phase as9%.However,the more detailed quantitative X-ray analysis and microstructural investigations using the higher resolution equipment are in progress.It is a very challenging problem due to very small austenite grain size and its dispersion.

A characteristic feature of bright laths of bainitic ferrite in Fig.5e and f are some precipitates inside the laths.They are present in both steels,so it is not related to the precipitation of Nb carbides. It is also low probable these are Mo carbides because of a hampering effect of Al and Si on carbide precipitation.The observed morphology at this temperature region is similar to the degenerate lower bainite containing intralath martensite–austenite islands instead of cementite plates[20,28].

Niobium microalloying does not contribute considerably to re?nement of the microstructure compared to corresponding microstructures of the steel without Nb.However,the evident effect consists in the uniform distribution and re?nement of all the structural constituents as distinct from the Nb-free steel where a few large bainitic–martensitic islands(B–M)can be found(Fig.4b and d and Fig.5a).Hardness of the

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Table 2–Fraction of retained austenite and hardness of the thermomechanically processed samples.

Steel type

Variant

designation Isothermal holding temperature,1C Cooling time between 700and 6501C,s Austenite fraction,%HV10

3Mn–1.5Al–0.2Si–0.2Mo–Nb

700–10s à650–4504501017.372.8375712700–25s à650–4504502517.872.4369725700–50s à650–4504505019.772.1345713700–50s à650–4004005018.872.2384723700–50s à650–35035050nd.

4177113Mn–1.5Al–0.2Si–0.2Mo–Nb

700–10s à650–4504501018.872.7414729700–50s à650–4504505020.472.7376715700–50s à650–4004005019.571.8435710700–50s à650–350

350

50

nd.

469721

nd.–not

determined.

Fig.4–DCCT diagram of the Nb-free steel with superimposed cooling paths (a)and multiphase microstructures obtained after applying various cooling time of specimens between 700and 6501C:10s (b and c)and 50s (d and e);3Mn–1.5Al–0.2Si–0.2Mo (b and d),3Mn–1.5Al–0.2Si–0.2Mo–Nb (c and e).

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339

thermomechanically processed specimens correlates good to the microstructures in Figs.4and 5.The highest hardness has the samples held at the temperature of 3501C with a distinct lath martensitic–bainitic microstructure (Table 2).The hard-ness values decreases with increasing the isothermal bainitic transformation temperature and with increasing the time of slow cooling during the g -a diffusive transformation.

Nb microalloying increases the hardness considerably from 20to 50HV for corresponding treatment variants.It is a combined effect of Nb(C,N)probably precipitated in a ?nal stage of hot-working and solid solution hardening as well as increased hardenability of austenite caused by dissolved niobium.

4.Conclusions

Third generation of the advanced high-strength steels for the automotive industry requires a high volume fraction of hard carbides-free microstructural constituents and retained

austenite of optimal mechanical stability for strain-induced martensitic transformation utilizing the TRIP effect.A simple solution for the increase of a volume fraction of austenitic phase is to add a higher manganese content compared to classical DP or TRIP steels (about 1.5wt%).The increased Mn concentration results in higher hardenability of steel what leads to the delay of ferritic transformation to long times on DCCT diagrams and favours martensite forming.The ?rst consequence is favourable for high-strength steels because of replacing polygonal ferrite by bainitic ferrite but the second one can be a reason of reduced ductility of steel sheets.It requires decreasing the carbon content and to keep a good balance between C and Mn contents.The obtained multi-phase microstructures contain some fraction of ferrite due to acceleration of a diffusive transformation under conditions of multi-stage deformation but the fraction of this phase is lower than 5%.Increasing the time of slow cooling from 10to 50s in the g -a transformation range does not in?uence the microstructure in a meaningful way.The major effect

for

Fig.5–Multiphase microstructures containing a various fraction of retained austenite obtained after applying a different isothermal holding temperature:4501C (a and b),4001C (c and d)and 3501C (e and f);3Mn–1.5Al–0.2Si–0.2Mo (a,c and e),3Mn–1.5Al–0.2Si–0.2Mo–Nb (b,d and f).

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stabilization of retained austenite has isothermal holding of steel in a range of bainitic transformation.The best tempera-ture for obtaining up to20%of g phase occurs for a range between400and4501C,where?ne granules and interlath retained austenite were identi?ed.The?ne austenite regions are stable whereas these ones with a size above3m m trans-form into martensite,what con?rms the importance of grain re?nement for the stabilization of retained austenite.Both steels have?ne-grained microstructures due to the thermo-mechanical processing applied but the distribution of all the microstructural constituents is more uniform for the Nb-microalloyed steel.

Acknowledgment

The research work was?nancially supported by the Polish Ministry of Science and Higher Education in a period of 2010–2012in the framework of project no.N N508590039.

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