Scandinavian Journal of Metallurgy 2002; 31: 272–280 Copyright C© Blackwell Munksgaard 2002Printed in Denmark. All rights reserved SCANDINAVIAN

JOURNAL OF METALLURGYISSN 0371-0459

Evolution of microstructure in an ultra-lowcarbon Cu bearing HSLA forging

Samar Das1, A. Ghosh2, S. Chatterjee2 and P. Ramachandra Rao1

1National Metallurgical Laboratory, Jamshedpur, India; 2Department of Metallurgy, Bengal Engineering College (D.U.), Howrah, W.B., India

Over the last decade low carbon age hardenable steelshave become a subject of interest for naval applica-tions. In the present study an ultra-low carbon Ni-bearingsteel alloyed with Cu was cast and processed by con-trolled forging and subsequently cooled in air, sand andwater. A variety of multiphase structures have beenobserved because of different cooling rates. The slowestcooling rate has produced bainitic ferrite by ledge mech-anism whereas lath- or plate-type bainitic ferrite was pre-dominant at the faster cooling rates. The volume fractionof retained austenite and/or martensite austenite constituentsincreased with decreasing cooling rate. Precipitation of

finer micro-alloying carbides and carbonitrides was sup-pressed by faster cooling. Slow cooling rates also resultedin the formation of ε-Cu precipitates of varying sizes andshapes.

Key words: bainitic ferrite, Cu precipitates, MA constituents,Moire fringes, polygonal ferrite, retained austenite, strain-induced precipitation.

C© Blackwell Munksgaard, 2002

Accepted for publication 23 November 2001

Since the 1960s, forging has emerged as an alternateroute to rolling for producing structural steels. Isotropicdeformation accompanied by a high rate of strain hard-ening during forging results in improved toughness andextra soundness in the material. Thermomechanical con-trolled processing played a major role in engenderingvarieties of microstructure in HSLA steels. Developmentof HSLA-80 and HSLA-100 steels has opened a newarena for structural applications. Higher strength andtoughness values were achieved as a result of disloca-tion substructures and precipitation hardening in thesesteels. Decrease in carbon content in HSLA steels pro-motes bainitic microstructures of different morpholo-gies in a wide range of continuous cooling transfor-mations [1–3]. Lowering the carbon content gives anextra benefit in terms of weldability, corrosion resistanceand subambient impact toughness. Moreover, use of Cuin HSLA steel envisages an extra advantage of precip-itation hardening on ageing [4–6]. Extensive researchefforts have already been directed to developing Cu-bearing HSLA steels but the influence of Cu on ultra-lowcarbon steel is yet to be explored. Ultra-low carbon steelsare very often referred to as steels containing less than orequal to 0.02 wt% carbon [7]. The achievement of 100 KSistrength in ultra-low carbon steel has been reported else-where [8] but reports concerning in-depth microstruc-

tural characterization are scanty. Moreover, precipita-tion behaviour of ε-Cu on continuous cooling has be-come a subject of interest. Hence, in the present inves-tigation an ultra-low carbon HSLA-100 grade steel wascast and controlled forged followed by cooling in differ-ent media. The microstructural constituents as well asmicro-alloying carbide/carbonitride and Cu precipita-tion because of thermomechanical controlled processingand variation in cooling rates, have been studied.

Experimental procedure

In this investigation, a (3.2 wt%) Ni-bearing steel with1.6 wt% Cu and micro-alloyed with Nb and Ti was pre-pared by air induction melting. The chemical analysiswas carried out in Quantovac (Model no. ARL80) andC, S and P contents were measured by LECO analysis.The composition is given in Table 1. The 50 mm × 50 mmingots were soaked at 1200◦C for 2 h. Controlled forgingwas done in two stages. The schematic representation ofcontrolled forging is shown in Fig. 1. 50% deformationtook place in the temperature range of 1100–1050◦C inStage I and in Stage II another 50% deformation tookplace in the temperature range of 850–800◦C. Subse-quently the forged slabs were subjected to sand, air andwater cooling. Cooling rates were measured from 800◦C

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Table 1. Chemical composition of the steel (wt%)

C S P Si Mn Mo Ni Cu Nb Ti Al N (ppm)

0.01 0.008 0.008 0.21 1.21 0.55 3.2 1.60 0.05 0.105 0.02 34

Fig. 1. Schematic diagram of controlled forging operation.

to 500◦C by embedding thermocouples into the steelcoupons. Standard grinding and polishing techniqueswere used for the metallographic samples. Specimenswere etched in 2% nital solution and examined by lightmicroscopy. For the transmission electron microscopy(TEM), thin slices were cut by a low-speed diamond sawand subjected to hand grinding before punching 3-mmdiameter discs and electrolytically thinning in a twin-jetpolisher using 90:10, acetic acid: perchloric acid as elec-trolyte at 12◦C. Samples were observed in TEM (PhilipsCM200) at 200-kV operating voltage. Energy dispersiveanalysis by x-ray (EDAX) was carried out to identify Cuprecipitates.

Results and discussion

Microstructure of the sample cooled at 0.68◦C/sThe optical microstructure (Fig. 2) shows polygonal fer-rite with a bright lustre and white coloured appearance.The light grey polygonal constituents along with darkdiscrete islands are often designated as granular bainite[1, 2]. Formations of faint subgrains in globular bainitehave been noticed. This is because of the presence of low-angle lath boundaries, which are not resolved properlyin light microscopy but are clearly observed in the trans-mission electron microscope. The bright field transmis-sion electron images taken from different areas are dis-played in Figs. 3a–c. Bainitic ferrite laths with low-anglelath boundaries are clearly observed and the dark filmsof retained austenite have been identified in the interlath

regions. When viewed at higher magnification (Figs. 4and 5), decomposition of austenite at very slow coolingbecomes prominent. The general sequence of austenitedecomposition with a decrease in transformation tem-perature is polygonal and acicular ferrite followed byupper and lower bainite and martensite. The austenitedecomposition starting temperature, i.e Ar3 and bainiticstarting temperature, i.e Bs can be determined from thefollowing equations [9, 10]:

Ar3 = 910–310%C-80%Mn-20%Cu-15%Cr

-80%Mo-55%Ni (1)

Bs = 830–270%C-90%Mn-37%Ni-70%Cr-83%Mo. (2)

From the above equation Ar3 and Bs temperatures forthis particular steel were found to be 550◦C and 546◦Crespectively. It is evident that austenite stabilization hasbeen increased for this steel chemistry and Ar3 tem-perature is just above Bs temperature. On very slowcooling when the steel attains the Ar3 temperature,austenite starts to decompose into polygonal ferrite by adiffusional mechanism. Figure 4 shows the existence ofpolygonal ferrite with its characteristic features of low-dislocation density. The low-dislocation density becauseof the transformation of austenite into polygonal ferritehas been reported by Thompson et al. [2]. They mea-sured the dislocation density in polygonal ferrite, whichwas found to be less by a factor of two or more thanthat of acicular/granular ferrite at the cooling rate of10◦C/s. As the temperature drops to the Bs temperature,the mode of austenite decomposition shifts from diffu-sional to mixed mode, i.e. bainitic transformation. Thepresence of ledges (marked by arrows) is identified inthe bright field electron micrograph in Fig. 5. The growthof bainite by lateral movement of such ledges has beenstrongly advocated by Aaronson and his colleagues [11].They pointed out that the equilibrium dislocation struc-ture either generated by the transformation from f.c.cto b.c.c or by precipitation provides an effective bar-rier to the growth of newly formed phase. The bar-rier to growth can be surmounted by the creation ofledges or steps whose lateral movement is parallel to theplane of boundaries. This lateral movement of the ledgescauses disordered boundaries, to migrate along the di-rections normal to the phase boundaries. These ledgeshave been observed to vary from nanometer to micronin size. It is envisaged that during controlled forging

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Fig. 2. Optical micrograph of the sample cooled at 0.68oC/s.

(a)

(b)

(c)

Fig. 3. (a)–(c) Bright field transmission electron micrograph of the samplecooled at 0.68◦C/s.

Fig. 4. Bright field transmission electron micrograph of the sample cooledat 0.68◦C/s shows the microstructural constituents containing polygonalferrite (PF), bainitic ferrite (BF) and retained austenite (RA).

Fig. 5. Bright field transmission electron micrograph shows the presence ofledges marked by arrows.

pancaking of austenite below the recrystallization, stoptemperature has introduced substantial amount ofstrain into the matrix, which on subsequent transforma-tion of austenite into bainitic ferrite produces insertionof these ledges. Creation of such ledges might have suf-ficiently disordered the boundaries for the easy glide,necessary during the growth of these platelets. The greyappearance of the bainitic ferrite phase in the opticalmicrograph (Fig. 2) can be inferred with the help of the

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Fig. 6. Bright field transmission electron micrograph exhibits an interfacialdislocation network.

transmission electron micrograph (Fig. 5). The contrastdeveloped by chemical etching in light microscopy sam-ples is because of the anodic–cathodic potential differ-ence between the phases that is very dependent on thestructural constituents and/or the mode of transforma-tion. The polygonal ferrite having low dislocation den-sities is revealed as a bright phase whereas the bainiticferrite phases containing high dislocation density arerevealed as grey colour in Fig. 2. The retained austen-ite may be dark in the optical micrograph. A very closeinspection of the electron micrograph reveals interfa-cial dislocation tangles along with very tiny particles atthe dislocation network (Fig. 6). The ingress of interfa-cial dislocation as observed in Fig. 6 has been the subjectof considerable debate. Sankaran & Larid [12] conjec-tured that these types of dislocations of Burger vectora/2 〈110〉 may first nucleate at the edges of the plateletsand their subsequent glide into the broad face may leadto the formation of the arrays. Other possibilities of ad-sorption of matrix dislocations into the interfaces havebeen described by Larid & Aaronson [13]. However, aseries of perturbations has been generated at the γ /α in-terphase as observed in Fig. 5. Aaronson and coworkershave described these types of perturbations in the γ /α

interphase region as the initial stages of developmentof Widmanstatten saw teeth. Invariably, these perturba-tions in the interphase region are associated with dislo-cation networks along with fine strain-induced precipi-tates, which are micro-alloying carbides, carbo-nitridesand ε-Cu particles. The kinetics of strain-induced pre-cipitation of NbC/CN has been well studied by Sellarsand his colleagues [14]. They have observed that the pre-

Fig. 7. Bright field transmission electron micrograph displays an enlargedview of dislocation precipitation interaction.

cipitation mechanism is enhanced by the formation ofa dislocation network. In the bright field electron mi-crograph of Fig. 7, a magnified view of the dislocationprecipitation interaction is exhibited. It is very interest-ing to note that the dislocations have been arrested bythe particles having interspacings ranging from 20 to50 nm.

Microstructure of the sample cooled at 35◦C/sThe optical microstructure of the water-quenchedsample is displayed in Fig. 8. Fine bainitic ferrite lathsare predominant along with some equiaxed/globular-shaped second-phase particles. The volume fraction ofdeeply etched second-phase particles has decreasedremarkably at the faster cooling rate. Bright fieldtransmission electron micrographs of this sample havebeen taken from three different positions and are shownin Figs. 9a–c. Fig. 9a shows the parallel arrays of bainiticferrite laths at the corner of the micrograph. Some block-like bainitic ferrites are distributed with dark discreteislands of MA constituents (marked by arrows). Theserrated edges of the laths are revealed in Fig. 9b. Highdislocation densities along with fine micro-alloyingcarbide/carbonitride precipitates predominate in themicrostructure. Interlath-retained austenite films arealso noticed in Fig. 9c. It is interesting to observe thevariation of bainitic ferrite morphology with variationin cooling rate. At the slowest cooling rate, the bainiticmicrostructure reveals formations of ledges at the in-terphase whereas for fastest cooling, lath- or plate-typebainitic ferrite has been engendered as the predominantphase. Fig. 10 shows an enlarged view of a bainitic sheaf

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Fig. 8. Optical micrograph of the sample cooled at 35◦C/s.

(a)

(b)

(c)

Fig. 9. (a)–(c) Bright field transmission electron micrographs of the samplecooled at 35◦C/s.

Fig. 10. Bright field transmission electron micrograph displays a magnifiedview of bainitic sheaf with interlath-retained austenite (RA) film.

that consists of a number of plate-type subunits with re-tained austenite films at the lath boundaries. The drivingforce for the transformation can be enhanced by increas-ing the cooling rate, which will result in the maximumnumber of sites for nucleation of bainitic ferrite subunits.The subunits grow in diffusion-less fashion but the ex-cess carbon in ferrite is soon partitioned into the resid-ual austenite. The reaction will proceed upto the pointwhen the free energy of austenite and ferrite of identi-cal composition, will attain the same value. As a con-sequence, decomposition of austenite will not proceedfurther and this will remain as retained austenite. If thecarbon concentration in the residual austenite exceedsthe equilibrium value, then it will reject carbon and asa result carbide will precipitate at the interface. In thisultra-low carbon steel, the initial carbon concentrationwas very low and, hence, carbide precipitation was notobserved very frequently, but the interfacial-retainedaustenite/martensite constituent has been identified, asmentioned earlier. An enlarged view of the interfacialMA constituent is shown in the bright field electron im-age of Fig. 11. The presence of these MA constituents isnot very desirable as they degrade the impact toughness

Fig. 11. Bright field transmission electron micrograph engenders the pres-ence of MA constituents.

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Fig. 12. Optical micrograph of the sample cooled at 1.15◦C/s.

property especially at subambient temperature. Tianet al. have elegantly reported that MA constituents areharmful kink formers and are responsible for poor sub-ambient impact toughness property of steel [15].

Microstructure of the sample cooled at 1.15◦C/sThe optical microstructure (Fig. 12) shows prior austen-ite grain boundaries along with second-phase particleswith an intermediate size compared to the previoustwo categories. The lath width has been increased com-pared to that of the water-quenched sample and thelath boundaries are not revealed prominently. The trans-mission electron micrograph, in Fig. 13a, shows arraysof parallel laths with discontinuous lath boundaries.Within the laths, bainitic ferrites of blocklike or equiaxedshape are visible along with randomly distributed smalldark regions of retained austenite that have been con-firmed by selected area diffraction. Fig. 13b shows theSADP taken from the black region of photograph andits interpretation is given in Fig. 13c.

Precipitation

Precipitation of Nb and Ti carbides andcarbonitridesThe austenizing temperature of 1200◦C was employedin the present study for allowing all carbides and car-bonitrides to go into the solution. The solubility of Tiand Nb in steel has been computed using the thermo-dynamic equations [16, 17]

log10 [Ti][C] = −7000/T + 2.75 (3)

log10 [Ti][N] = −16586/T + 5.90 (4)

log10 [Nb][C + 12/14N] = −6770/T + 2.26 (5)

log10 [Nb][C] = −6770/T + 2.26 (6)

(all T in K).

(a)

(b)

(c)

Fig. 13. (a) Bright field transmission electron micrograph of the samplecooled at 1.15◦C/s. (b) The marked SADP from the black discrete regionof Fig. 13(a). (c) The schematic illustration of SADP.

It can be seen from the above equations that the solu-tionizing temperature of 1200◦C was high enough totake Nb- and Ti-carbides and carbonitride into solution.Faster quenching from the finish forging temperatureof 850◦C suppressed their precipitation. Bright fieldelectron micrographs in Fig. 14a, show a very low

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(a) (b)

(c) (d)

(e)

Fig. 14. Bright field transmission electron micrograph exhibits precipitation behaviour of micro-alloying carbides and carbonitrideat various cooling rates: (a)35◦C/s, (b) 1.15◦C/s and (c) 0.68◦C/s. (d) The SADP taken from the cuboid precipitation of Fig. 14(c). (e) The schematic illustration of Fig. 14(d).

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(a) (b)

(c)

Fig. 15. (a) Bright field transmission electron micrograph display ε-Cu precipitate of the sample cooled at 0.68◦C/s. (b) Dark field electron image of Fig. 15(a).(c) EDAX analysis of the precipitate marked at Fig. 15(b).

precipitation density in the matrix though some strain-induced precipitation has occurred near the disloca-tion network. The intermediate holding between the1st and 2nd stage deformation during controlled forg-ing has resulted in strain-induced precipitation. Precip-itation density has increased with reduction in coolingrate. An intermediate cooling rate of 1.15◦C/s producedfiner precipitates with a dense distribution in the matrix,

shown in Fig. 14b. When the cooling rate was very slow(0.68◦C/s) particle coarsening occurred giving rise tolarger size precipitates of regular geometric configura-tion. Figure 14c shows a bright field electron image of thesample cooled in sand medium where cuboid-shapedparticles of 70–75 nm size were observed in the bainiticferrite matrix. An SADP was taken from this precipitateas shown in Fig. 14d and its analysis is given in Fig. 14e.

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Fig. 16. Dark field electron micrograph shows various shapes and sizes ofε-Cu precipitates.

Precipitation of Cu. At the cooling rate of 0.68◦C/s, pre-cipitates of ε-Cu have been identified with various mor-phologies and size. The precipitation of ε-Cu is exhibitedin the bright field and the corresponding dark field elec-tron micrographs in Figs. 15a and b. The average sizes ofCu precipitates are very small and lie in the range 10–18nm. A large precipitate ∼50 nm size was observed at thecentral region of Fig. 15b and EDAX analysis was per-formed on this particle. The EDAX analysis shown inFig. 15c confirms the particle as Cu. Spherical particles∼28–30 nm size along with rod-shaped precipitates of43–45 nm length are observed in the dark field electronmicrograph of Fig. 16. The occurrence of rod-shapedprecipitate in Cu bearing steel on ageing has been re-ported by many workers [18]. In a closer observation ofFig. 16 parallel fringes are revealed in the Cu particlesand marked by arrows. Miglin et al. [19] have reportedthat the possibility that stacking faults may exist in thelarger-sized ε-Cu precipitate. During the growth of theseprecipitates coherency of the precipitates with matrixis gradually lost and they ultimately become incoher-ent with the matrix that may result in stacking faultsperpendicular to the most closely packed direction of〈111〉Cu. Yet another possibility of these parallel lines asMoire fringes cannot be ruled out. The fringes appearedin the large-size precipitates, and these precipitates havesevere possibilities for atmospheric oxidation just afterelectrolytic jet thinning. The fine oxide layer on the pre-cipitate may give rise to this type of fringe contrast inthe electron micrograph.

Conclusion

From the foregoing results and discussion it can beconjectured that bainitic ferrites of different morpholo-gies have been obtained as a result of variations in thecooling rate. Faster cooling produced bainitic ferritewith lath- or plate-shaped morphology together withretained austenite or MA constituent. With reductionin the cooling rate, formation of polygonal ferrite hasbeen facilitated along with bainitic ferrite and retainedaustenite. The bainitic ferrite at slower cooling rates ismainly formed by a ledge mechanism. Precipitation offine carbides and carbonitride of Ti and Nb was sup-pressed by faster cooling whereas very slow cooling pro-moted particle coarsening resulted in large-sized parti-cles with regular geometric configuration. Rod-shapedor spherical ε-Cu particles have also been found at theslowest cooling rate.

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Address:S. ChatterjeeDepartment of MetallurgyBengal Engineering College (D.U.)Howrah 711 103,West BengalIndiae-mail: Subrata Chatterjee@hotmail.com

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