Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

THE METALLOGRAPHY EVALUATION OF THE RELATIONSHIPBETWEEN THE HOT DUCTILITY SHORTNESS AND THESUSCEPTIBILITY TO TRANSVERSE CRACKING IN CC PRODUCTS

Jiří Pavliska, Zdeněk Jonšta, Karel Mazaneca)

Eva Mazancováb)

a) Institute of Mater. Eng., Technical University, 17. listopadu 15, 708 33Ostrava-Poruba, CZ

b) NOVÁ HUŤ a. s., Research and Testing Institute, 707 02 Ostrava, CZ

AbstractThe work is devoted to the study of low-carbon steel microstructure on the plastic

behaviour during the hot deformation at the critical temperature region. Under theseconditions the steel microstructure consists of a ferrite and austenite mixture. The inducedplastic deformation is concentrated in the intergranular ferrite zones having theallotriomorphic feature. The levels of applied strain rate and the heterogeneity effect ofdeformation in soft ferrite zones are analyzed. Simultaneously, the physical metallurgycharacteriscics of the trough of ductility shortness found in dependence: R in A value - testingtemperature are evaluated from point of view of the developing recrystallization process andof the possible limitation of this process, respectively. Further, the influences of theallotriomorphic ferrite morphology on the type of subsequent austenite decompositionsproducts are investigated. The behaviour of allotriomorphic ferrite – austenite is veryimportant on the subsequent decomposition of austenite (inert/active interface). The effects ofdifferent microstructure types on the susceptibility of CC products to the transverse crackingare discussed.

1. INTRODUCTIONIn recent years, the hot ductility of steel at low strain rates has become important because

of its relationship to the problem of transverse cracking observed during continuous casting(CC). These cracks are believed to form when the strand, usually cast in curved mould, isstraightened in the temperature range: 700 ° to 1100°C. This corresponds to the range inwhich steels, when tested in a tensile test, show a ductility trough due to the intergranularweakness in austenite [1]. The formation of these cracks running along the austenite grain boundaries can be causedby two mechanisms: namely by grain boundary sliding in the austenite matrix and/or by phasetransformation leading to the intergranular failure. In the former case, the failure is associatedwith the following successively realized processes: grain boundary sliding, cavitation andconcluding cavity linkage. This process is much encouraged by having particles at grainboundaries. Further, this effect can be enhanced due to matrix strengthening by very fineprecipitates contributing to the localization of deformation along austenitic grain boundaries. In the later case, failure can be again associated with grain boundary sliding in theaustenite. It can occur by simultaneous phase transformation when a narrow ferrite film formsaround the austenite grain boundaries. The formation of this film is often strain-induced andcan be realized in a wide temperature range [2]. Usually this behavior is evaluated using the isothermal tensile tests that are performed atstrain rate around 10-3÷10-4 s–1. The effects of temperature dynamics realized during thecooling of CC products complicate the easy evaluation of mechanical response as it resultsfrom crack/expert system presented by University of BC [3,4].

Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

2. BASIC APPROACHESThe hot ductility shortness by tensile testing and the proposed susceptibility to the crack

formation in CC products have the relation to the morphology of ferrite formed in theinvestigated steel type. The microstructure characteristics are linked to the realized phasetransformation process of austenite into ferrite that is observed in cast steel during itsstraightening operation [5,6]. Besides the above mentioned microstructure effects, thepresence of some residual elements in steel and their influence on a higher occurrence ofcracks resulting from the detrimental action of these elements on the plastic behaviour atgiven critical temperature range, are taken into consideration [7].

As sufficient ferrite will be present at the lowtemperature end of the trough, ductility will alwaysimprove because ferrite has higher ductility level[8]. The ferrite is formed either via normaltransformation prior to deformation, or by thedeformation itself. The narrow ferrite zones found atthe austenite grain boundaries are generally theproduct of deformation induced phasetransformation realized at temperatures situated upto Ae3 [1]. The detected recovery in ductility at lowtemperature end of the trough of plastic propertiesdepends for the main on the amounts of ferrite beingpresent in the microstructure (above 30% whatprevents strain concentration and gives highervalues of R in A than 40%). This plasticity level isindependent of whether it is produced by normal orstrain induced transformation [8].

At the high temperature end of the trough, the ductility improves when dynamicrecrystallization occurs (Fig.1). Any cracks that are formed at the grain boundaries areisolated if the grain boundary moves away from this cracks during the development ofrecrystallzation. The new grains are formed and in this consequence crack growth is halted.However, dynamic recrystallization is not possible during CC process because thedeformation induced in strand is too small (around 2-3%). In addition to this effect, it is alsonecessary to take into account the coarse grained microstructure of CC products. Althoughstrains can be greater in the thin slabs casting when predeformation is used and the received

DynamicRecrystalliz.

No dynamicRecrystallizat.

Nb contain. steel

C-Mn aC-Mn-Alsteels

Temperature [°C]

R in

A [%

]

Fig. 1 Schematic diagram illustrating theductility levels that can beachieved with and withoutdynamic recrystallization.

Temperature [°C]

Dynamic.recrystallization

εf

εc

a)

R in

A [%

]

b)

εc2

εc1

εf2

εf1

TD1 TD2

Temperature [°C]

R in

A [%

]

Fig. 2 Schematic diagram showing: a) how the width of ductility could be controlled by dynamicrecrystallization; b) how increasing the strain rate reduces the depth and width of thetrough [9].

Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

grain size is finer, the microstructure is still too coarse and dynamic recrystallization does nottake place.

Schematic diagrams presented in Figs. 2 a, b show how the width of the ductility trough iscontrolled by dynamic recrystallization (a) and how increasing the strain rate reduces thedepth and the width of the trough (b). The temperature TD corresponds to the onset ofdynamic recrystallization while εc and εf express the critical strain for dynamicrecrystallization and the total strain to failure in the absence of dynamic recrystallization,

respectively [9]. Figure 3 demonstrates the effect ofgrain size refinement.

Increasing the strain rate work - hardens theferrite more rapidly. This in turn displaces thedeformation toward the centers of grains, therebyincreasing the amount of material participating inthe phase transformation and the deformationprocess is more homogeneous [2]. The trough isshallow, the width is smaller and the lower value ofR in A is shifted to higher ductility level. In sum,we can hold the following parameters that controlductility values: strain rate, grain size, precipitationand inclusion content, for the most important variables [1,10]. Concerning the precipitation, thefiner particles, the worse is the ductility. The grainboundary precipitation is particularly deleterious[1]. Strain induced precipitation is always finer and

more detrimental to ductility than the precipitation process that is present before strain [11]. Inthe case of Nb and V containing steels, a larger part of the precipitation comes outdynamically during unbending operation and so is very detrimental to ductility [1].

Before starting the metallography analysis, it is important to study the relevance of the hotductility tensile tests to the problems of transverse crack formation in CC products. As itresults from the above-mentioned crack/expert system [3], the application of thecorresponding software is very useful to predict the changes in the casting speed and waterflow rate on the thermal history. For this reason, it is necessary to subject the tensilespecimens to the laboratory simulation based on the predicted time/temperature schedules andto apply the tensile strain to fracture at temperature corresponding to straightening [3]. Thesimple isothermal tensile tests can be held for useful because contribute to the informativeassessment steel susceptibility to cracking. However the immediate technical application ofobtained results is not possible without the additional analysis based on the finding of therelationship between microstructure characteristics and obtained mechanical properties.

For this reason, the aim of present work is contribute to the elucidation of themicrostructure effect (the morphology of ferrite) on the hot ductility shortness of steels whatwe hold for very important information.

3. RELATION BETWEEN FERRITE MORPHOLOGY AND HOT DUCTILITYSHORTNESS.The relation between the ferrite morphology and the susceptibility of CC product to the

transverse cracking were discussed in whole range of papers [12-14]. The laboratorymeasurements show the transverse crack formation is linked to the low hot ductility valuesfound on tensile specimens performed at temperatures corresponding to the critical hotshortness interval. The microstructure of tested specimens is austenitic and/or mixedconsisting of ferrite and austenite. The chemical composition of investigated steels plays an

Fig. 3 Schematic diagram showing howrefining the grain size reduces thedepth and the width of the trough.

εc2 fine

εc1 croase

TD1 TD2

εf1 croase

Temperature [°C]

R in

A [%

]

εc2 fine

Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

important role in particular from viewpoint of austenite decomposition process and ferriteformation before loading and/or after loading by deformation induced mechanisms [9,10]. Incarbon and/or in C-Mn commercial steels, the formation of the low ductility trough (independence of R in A values on testing temperature) is related to the austenite decompositionusually.

The ferrite grains are softer than austenitic at given temperature what results from a higherstacking fault energy of ferrite in comparison with austenite. Therefore, the ferrite recoversmore rapidly. The result of this behavior is that when small amounts of ferrite are present inaustenite matrix, any strain applied to the material is concentrated in softer ferrite regions.The ferrite is therefore subjected to a deformation substantially higher than is the global levelimposed on the steel and exceeds this value several times [15]. It is proposed, the localizeddeformation is up to 15-20 times increased, in comparison with the global deformation level,approximately.

The analysis of these conditions shows, the minimal value of hot ductility can be found forferrite volume fraction of about 5%. It should be emphasized that the above-consideredconception is only acceptable if the ferrite film covers the entire austenite grain boundaries. Ifthe ferrite volume faction is smaller than the mentioned limit of 5% approximately, thediscontinuities of intergranular ferrite layers are found and the considered detrimental effectof ferrite films is suppressed partially. On the contrary, if the ferrite volume fraction is higherthan 30% then the strain concentrated in ferrite is less significant and this ferrite volumefraction becomes less detrimental and the attained hot ductility level is improved.

Failure in uniaxial tension tests performed in the dual-phase microstructure, simulating asclosely as possible the conditions pertaining to the unbending operation during CC process,occurs by void linkage along the thin ferrite film decorating the austenite grain boundaries.The basic effect causing this low ductility trough formation corresponds to the localization ofplastic deformation at austenite grain boundaries if the intergranular sliding mechanism isdominant in described deformation process. In addition, it is necessary to take intoconsideration the participation of precipitated non-metallic inclusions (e.g. sulphides) and thephysical metallurgy parameters of strain-induced precipitation of carbides, nitrides and/orcarbonitride particles [3].

The ferrite morphology influences the conditions for the realization of localizeddeformation in CC products and also the tendency to their transverse cracking. The ferritedecorating the austenite grain surfaces has the allotriomorphic feature usually, what furtherresults in the modification of following austenite decomposition process realized in adjoiningmatrix regions. The physical metallurgy behaviour of the ferrite – austenite interface plays animportant role as active and/or as inert interface in this connection [16]. The activeallotriomorphic ferrite (ATF) is defined as a microstructure component, which is able todevelop into other transformation products such as Widmanstätten ferrite (WF) or bainiticpackets (B) at given temperature. The ATF was said to be inert when the local reduction intransformation temperature at the ferrite/austenite interface due to the partitioning of carbonprevents the development of the secondary WF or B. The example of the inert interface ispresented in Fig. 4. Under this condition, the ATF can be rendered inert by the build up ofcarbon in the austenitic matrix ahead of the boundary of austenite and the ATF [16].

The corresponding controlling process is linked to the local carbon enrichment at one sideof interface. The observed asymmetric morphology represents a very interesting behaviour.The causes of this effect are not elucidated unambiguously up to this time. The followingpossibilities of the material response could be taken into consideration. For example, thedescribed behaviour can be attributed either to the localized grain boundary enrichment withmanganese what leads to a lower carbon activity or to the specific property of the interface.The interface has a „straight – line“ appearance without ascertainable deviations in its course

Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

(Fig. 5). This figure shows a section of the interface enriched of carbon. On the contrary, thesection without this local enrichment can be described as a continuously growing acicular WRimmediately connected with the zone of ATF.

The feature of observed cracks are also different. The cracks running along the „straight –line“ inert interface of austenite-ATF are not accompanied with the formation of somedeviations (Fig. 6a). These cracks can be held for smooth according their course. The cracksfound in the central part of the intergranularly precipitated ATF are accompanied with theformation of numerous deviations in their course. The observed deviations in crack course canbe interconnected with the non-metallic inclusions. (Fig.6b). In this case, the linear crackroughness parameter RL is 1,25, approximately, while in the first presented example (smoothcracks – Fig 6a) the roughness parameter of cracks RL is only about 1,04 [17].

Figure 7 shows a very interesting example of the ATF initiation at the non-metallicinclusions. The given initiation examples are usually classified as the „boxing incharacterization morphology” [18]. In their vicinity, the traces of ferrite deformation can bedetected, inclusive the micro-crack initiation running from the interface of the non-metallicinclusion and AFT matrix.

Figure 8 shows the regions of intragranular acicular ferrite (IAF) formation including theindications of the interlocking morphology of this microstructure compound [19]. This type ofmicrostructure is found in CC product comparatively seldom and its formation depends on thefulfillment of some conditions: a) coarse austenite grain size leading to the limitation of

Fig. 4 Inert interface of ATF Fig. 5 Section of the interface carbonenrichment

Fig. 6a Smooth crack running along the“straight-line” interface Fig 6b Crack initiated in the central

part of ATF

Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

nucleation probability of the B and/or of the WF at austenite grain boundary; b) the sufficientnumber of non-metallic inclusions having the dimension assuring the favourable condition forthe IAF heterogeneous nucleation; c) the presence of the nucleant inclusions in steel matrix[20].

4. CONCLUSIONSThe work describes the effect of microstructure found in CC products on the achieved level

of hot ductility in the critical temperature region if the deformation process is realized in theheterogeneous dual - matrix consisting of ferrite and austenite. The influence ofmicrostructure type on the response of CC products to straining process during theirunbending operation was determined. It was demonstrates, the microstructure types play avery important role in the formation of evaluation of susceptibility level to the transversecrack formation in CC steels.

AcknowledgementThe authors would like to acknowledge the Grant Agency of Czech Republic for the

financial support. The work was realized in the project GAČR No. 106/96/K032.

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Fig. 7 ATF initiation at the non-metallic inclusions

Fig. 8 Regions of IAF formationhawing the interlocking morphology

Metal 2001 15. – 17. 5. 2001, Ostrava, Czech Republic

[13] Mintz, B., Abushosha, Crowther, D. N.. Mater. Sci. Technol., 1995, vol. 11, pp. 474-481.

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