[25]the Mechanism of Acicular Ferrite in Weld Deposits

8/9/2019 [25]the Mechanism of Acicular Ferrite in Weld Deposits

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The mechanism of acicular ferrite in weld deposits

Sudarsanam Suresh Babu *

Metals and Ceramics Division, Bldg 4508, Mail Stop 6096, Oak Ridge National Laboratory, Oak Ridge, TN 37831, United States

Abstract

Research has shown that the acicular ferrite microstructure in steel weld metal, which provides an optimum combination of strength and toughness, is indeed intragranularly nucleated bainite. It is possible to maximize the content of acicular ferrite byincreasing the intragranular nucleation sites while maintaining a critical weld metal cooling rate and the steel hardenability. Thispaper highlights recent research related to nucleation and growth of acicular ferrite during decomposition of austenite.Published by Elsevier Ltd.

Keywords: Steels; Weld microstructure; Acicular ferrite; Bainite and phase transformations

1. Introduction

In the late 1970s, the importance of an acicular ferritemicrostructure towards optimizing strength and tough-ness was identied [1]. Since then, extensive researchhas been done on acicular ferrite formation and its rela-tion to oxide inclusions, weld metal hardenability andcooling conditions [**2,**3] . An overview of the high-temperature phase changes that occur in steel weld metalis presented (see Fig. 1 ) to describe acicular ferrite for-mation in the context of the overall microstructure evo-lution. The weld metal microstructure is affected bymelting, gas dissolution, solidication and solid-statetransformations. Since the weld pool region is heatedto temperatures as high as 2500 K, the liquid steel dis-solves oxygen. The extent of oxygen dissolution depends

upon the thermodynamic properties of liquid metal, gas,and slag phases [4]. As the liquid weld metal cools fromthis temperature ( Fig. 1 a), in the temperature range2000–1700 C, the dissolved oxygen and deoxidizingelements in liquid steel react to form complex oxideinclusions in the range of 0.1–1 l m size range. In the

temperature range 1700–1600 C (Fig. 1 b), solidicationto d-ferrite (body centered cubic phase) starts and envel-ops these oxide inclusions; and this d-ferrite transformsto austenite ( c-face centered cubic phase). In the temper-ature range from 1600 to 800 C (Fig. 1 c), austenitegrain growth may occur. In the temperature range800–300 C ( Fig. 1 d–g), the austenite decomposes to dif-ferent ferrite (body centered cubic phase) morphologies.The austenite to ferrite decomposition ( Fig. 1 d and e)starts with the formation of allotriomorphic ferrite ( a )at prior c – c boundaries and eventual coverage of theseboundaries. With continued cooling ( Fig. 1 f), the Wid-mansta¨ tten ( a W ) ferrite nucleates at the a /c boundariesand extends into the untransformed austenite–graininteriors. Further cooling to low temperatures, ( Fig.1g) the acicular ferrite ( a a ) would nucleate on the inclu-

sion. If there are no potent inclusions, bainitic ( a b ) fer-rite might form instead of acicular ferrite, from theremaining austenite. On further cooling to temperaturesclose to room temperature, any remaining austenite maycompletely or partially transform to martensite. Thismixture of martensite–austenite phases is referred to asM–A constituent. These phase transformation se-quences are important while discussing the mechanismof acicular ferrite formation because each of the abovereactions modies the nucleation and growth kinetics

1359-0286/$ - see front matter Published by Elsevier Ltd.doi:10.1016/j.cossms.2004.10.001

* Tel.: +1 865 574 4806; fax: +1 865 574 4928.E-mail address: [email protected]

Current Opinion in Solid State and Materials Science 8 (2004) 267–278

mailto:[email protected]:[email protected]


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of acicular ferrite. The importance of considering theoverall microstructure evolution was recently stressedeven for identication of microstructural constituents[*5].

A typical weld microstructure containing acicular fer-rite ( a a ) is shown in Fig. 2 [6]. This morphology is usu-ally referred as a ‘‘basket-weave structure’’ [7]. Theoptical microstructure (see Fig. 2 a) shows the interlock-ing nature of acicular ferrite ( a a ) co-existing with otherferrite morphologies such as Widmansta ¨ tten ferrite(a W ) and allotriomorphic ferrite ( a ). The acicular ferriteplates have an aspect ratio of 0.1–0.2 in any randomcross sectional area. High-magnication analyses of the same microstructure in a transmission electronmicroscope show primary plates (marked as P) thatnucleate on inclusions and the secondary ferrite pates(marked as S) that nucleate on preexisting ferrite plates(see Fig. 2 b). This chaotic arrangement of lenticular fer-rite plates within an austenite grain is considered to beone of the mechanisms by which the toughness is im-proved in the welds [**8]. This arrangement is promotedby the nucleation of acicular ferrite on oxide inclusions.The concept of intragranular nucleation on inclusionshas also been extended to structural steels also by inten-

tionally adding titanium oxide during steelmaking be-fore any other thermomechanical processing [*9,*10].

In this review, recent progress in understanding thedetails of nucleation and growth of acicular ferrite, theparameters that affect the formation of acicular ferriteduring weld cooling, overall transformation kinetics of acicular ferrite formation and the transitions from acic-ular ferrite to bainite microstructure are reviewed. Theconcepts presented are also applicable to acicular ferritethat forms in oxide-inoculated structural steels.

2. Mechanisms of austenite to acicular ferritetransformation

Past research has revealed that acicular ferrite platesare crystallographically related to the parent austenitegrain approximately either by the Nishiyama–Wasser-man (NW) orientation relationship or by Kurdjumov– Sachs (KS) orientation relationship [**8,**11,12] . Inaddition, the acicular ferrite microstructure exhibits sur-face relief phenomena [13] indicating the formation of these ferrite plates from austenite involves an invariant

Fig. 2. Typical acicular ferrite ( a a ) microstructure in a low alloy steel(Fe–0.06C–0.51Si–1.11Mn–0.48Crwt.%) weld metal with small addi-tions of aluminum, titanium and oxygen: (a) optical micrograph showsthe presence of grain boundary ( a ) and Widmanstätten ferrite ( a W )coexisting with acicular ferrite, (b) transmission electron micrographshows the apparent nucleation of primary acicular ferrite plate(marked as P) on an oxide inclusion and a secondary acicular ferriteplate (marked as S) on prexisting ferrite plate.

Fig. 1. Schematic illustration of a weld metal cooling curve and ahypothetical continuous cooling transformation diagram shows differ-ent phase transformations that may occur as the weld metal cools toroom temperature.

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plane strain with a large shear component. Elementalanalyses across acicular ferrite and austenite phaseshave also revealed no substitutional alloying element(i.e., Si, Mn, Cr, etc.) partitioning [14,6]. In addition,atom probe eld ion microscopy indicated that inter-stitial alloying elements (i.e., carbon) could partition

between austenite and ferrite [15]. The austenite to acic-ular ferrite transformation kinetic data exhibit incom-plete reaction phenomena, i.e., the transformationstops when the carbon concentration in austenitereaches a critical value ( T 0 condition) above which thediffusionless transformation is thermodynamicallyimpossible [**8,12–14] . Furthermore, the ‘‘basket-weave’’ morphology of acicular ferrite plates can bechanged to a non-random orientation by applyingelastic stresses during transformation [16]. Duringtransformation under stress, evidence for transforma-tion-induced plasticity has also been observed. Basedon the above-mentioned research, one can conclude thatthe transformation mechanisms of acicular ferrite andbainite are similar, except for the requirement of potentinclusions within the austenite grain for acicular ferriteformation. Although, the bainite and acicular ferrite for-mation mechanisms are similar, due to the presence of many crystallographic variants of plates in acicular fer-rite and resulting chaotic microstructure, the toughnessof the acicular ferrite microstructure in welds is betterthan the toughness of a bainitic microstructure [17,*10].

2.1. Nucleation of acicular ferrite

To understand the mechanism of acicular ferrite for-mation, one needs to understand why it nucleates oninclusions? Based on the published literature, four majormechanisms can be put forward to explain the

nucleation of acicular ferrite [*18,19–22,*23,*24,25– 27,*28,29,*30,**31,*32,33,34,*35,36] . These mecha-nisms are illustrated schematically in Fig. 3 . (1) In therst mechanism (see Fig. 3 a), the inclusions act as inertsurfaces leading to a reduction of activation energy andferrite nucleation is promoted. The calculations showthat the ratio of the activation energy for nucleationon inclusions to that for homogeneous nucleation willdecrease with an increase in inclusion diameter. (2)Good lattice matching between inclusions and ferritecan also reduce the activation energy for nucleation.Due to the constraint of a reproducible orientation rela-tionship between austenite and ferrite, the probability of achieving suitable orientation relationships betweeninclusion and ferrite and also between ferrite and aus-tenite may be difficult. (3) Inclusions may deplete the ele-ments such as carbon, manganese and silicon from theaustenite. This depletion may lead to a local increasein the driving force for the nucleation of ferrite fromaustenite at the inclusion surface [*28]. (4) Due to a dif-ference in thermal expansion coefficients ( De) of austen-ite and inclusions, thermal strains may develop near the

Fig. 3. Schematic illustrations of different mechanisms for nucleation of acicular ferrite on inclusions: (a) Inclusion surface acts as an inert surface fornucleation and therefore large inclusions are potent sites. (b) Matching between ferrite and inclusion lattice will decrease the interfacial energybetween inclusion and ferrite in comparison to inclusion–austenite interface. (c) Local depletion of hardening elements such as carbon and/ormanganese which might lead to an increase in driving force for nucleation of ferrite from austenite. (d) Strain energy near the inclusion may increasedue to differences in thermal expansion between austenite and inclusion, which might reduce the activation energy for nucleation.

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inclusion austenite interface, which may reduce the acti-vation energy for the formation of a ferrite nucleus [36].

In this section, the above mechanisms are evaluatedbased on two methodologies. In the rst methodology,the formation of acicular ferrite involves thermal activa-tion. Based on this assumption, one can dene the nucle-

ation rate per unit volume per second ( I V ) by thefollowing equation [**8]:

I V ¼ mexp DG kT ; ð1Þ

where m is the attempt frequency, k is the Boltzmannconstant, T is the temperature in Kelvin and DG * isthe activation energy for nucleation per unit volume of the matrix. Classical textbook analyses have shown thatDG * is inversely proportional to the square of the chem-ical driving force. Recent research has focused on relat-ing the activation energy for nucleation on inclusion todifferent metallurgical variables. In a classical work onacicular ferrite nucleation, Ricks et al. [*18] used classi-cal nucleation theory to describe the nucleation potencyof inclusion (see Fig. 3 a). By assuming a spherical geom-etry, the activation energy for nucleation on an inclusionwas related to the driving force for the austenite to fer-rite transformation ( DG c! amax jvol ), interfacial energies of austenite-inclusion ( r c /inc ), ferrite–inclusion ( r a/inc ) andaustenite–ferrite ( r c /a ) [see Fig. 3 a]. As per these calcula-tions, the nucleation potency of inclusions will increasewith an increase in their radius, i.e. activation energy de-creases as inclusion diameter increases. Furthermore,the calculations also showed that the activation energy

nucleation of ferrite at the grain boundary is alwayslower than at inclusions. It is also known that not allinclusions provide preferential nucleation sites. Thismay be related to possible changes in the interfacial fer-rite–inclusion ( r a/inc ) with different inclusion types. Re-cent publications have used the model by Ricks et al.has been used to explain acicular ferrite transformationboth in welds and structural steels by calculating areduction in activation energy for nucleation due toone or more phenomena listed: (i) increase in the inclu-sion radius [21]; (ii) reduction of interfacial energy bylattice matching (see Fig. 3 b) [*23,*24] , (iii) increase inthe DG c! a

maxjvol

by local depletion of substitutional alloy-ing elements including manganese [37] or (iv) localdepletion of carbon [27,*28] near the inclusion.

In the alternative methodology, the nucleation of acicular ferrite can be understood by a theory basedon bainite nucleation [**8,*38] . In this theoretical treat-ment, the activation energy is assumed to be propor-tional to the driving force for transformation, as givenby the following equation:

DG ¼ bT ; ð2Þ

where b is assumed to be an constant. This theory isbased on martensite nucleation theory and assumes that

the ferrite nuclei are forming by displacive mechanismswith the partitioning of carbon to the austenite. As perthis theory, the austenite already contains preexistingferrite embryos and the nucleation of ferrite occurs assoon as the chemical driving force increases above thebarrier energy for interface motion. In addition, the the-

ory also requires the presence of a glissile interface be-tween the ferrite and austenite. Rees and Bhadeshia[*38] found experimental evidence for Eq. (2) by analyz-ing the acicular ferrite start temperature as a function of weld metal composition. With this concept, one can de-rive an expression for nucleation rate of acicular ferrite[**8,39,40] ,

I V ¼ C 3 exp C 4

RT

C 4DG c! amaxC 2 RT ; ð3Þ

where C 3 is related to the number of nucleation sites andC 2 and C 4 are constants based on experimental measure-ments. For acicular ferrite transformation, C 3 would berelated to number density of potent inclusions and forbainite transformation C 3 is related to grain boundarynucleation sites.

In the following paragraphs, possible mechanisms forthe modication of parameters in Eq. (3) by the presenceof inclusions within the austenite will be discussed. Thenucleation rate of acicular ferrite can be associated witha change in the magnitude of C 2 , C 3 , C 4 or DG c! amax . SinceC 2 is associated with the slope of the ‘‘universal nucle-ation function’’ for acicular ferrite or bainite, one canassume that this does not change with the inclusionstype. The presence of inclusions can increase the C 3 by

providing an inert surface [**41] within the austeniteand thereby promoting the nucleation of acicular ferrite.However, this assumes that the interface structure of inclusion–austenite is more or less similar to that of anaustenite–austenite grain boundary. Therefore, it is notclear how changing the inclusion–austenite crystallo-graphic orientations will lead to a different interface dis-location structure, which might inuence the nucleationof acicular ferrite.

In the original derivation of Eq. (3), the C 4 is also re-lated to number of array dislocations in the fault whichacts as embryo for nucleation, shear resistance of the lat-tice to dislocations, magnitude of Burger s vector, en-ergy barrier between adjacent equilibrium positions of dislocations and the interfacial energy between austeniteand ferrite. Rees and Bhadeshia [*38] speculated thatembryos might exist at the dislocation debris creatednear the inclusions due to thermal strains induced bydifferences in the thermal expansion between inclusionsand austenite. Furthermore, it is suggested that in caseof large inclusions, these stresses cannot be accommo-dated and may lead to the generation of a dislocation ar-ray [**8] by dislocation motion. In this mechanism, it isimportant to note that the stress build up is proportionalto both De and the degree of undercooling ( DT ) below



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the austenite to ferrite transformation temperature[33,34,35*,36] . These treatments indicate that a plasticzone may develop near the inclusion interface. Panet al. [36] extended the treatment of Brooksbank andAndrews [*35] and indicated that the strain energy valuemight not be very high in comparison to the chemical

driving force because of the plastic deformation of aus-tenite close to the inclusion. The plastic deformationmay help in nucleating the acicular ferrite by creatingdislocation defect structures that would be suitable fordisplacive nucleation of ferrite. At the same time, exces-sive plastic deformation may not work in favor of acic-ular ferrite growth, which will be described in the nextsection. In real weld cooling conditions, the stresses nearthe inclusion–austenite interface can also be relaxeddue to stress relaxation phenomena. This implies thatthe acicular ferrite formation at the inclusion will befavored only under certain thermomechanical coolingconditions.

The depletion of alloying elements near the inclusioninterface may increase the driving force for transforma-tion ( DG c! amax ), which will lead to an increase in the nucle-ation rate. It is noteworthy that the concentration of alloying elements such as manganese [*32,37] may notbe homogenized after repeated thermal cycling betweenthe austenite and ferrite phase regions during weldingdue to the low diffusivity in austenite. At the same time,the manganese-depleted regions are not always observedin welds [21]. Decarburization of austenite close to theinclusion–austenite interface may also promote nucle-ation, as proposed by Gregg and Bhadeshia [27,*28].

However, several questions exist with regard to thismechanism. How can this carbon depletion be presentduring repeated thermal cycling? Will the inclusion po-tency decrease with repeated thermal cycling betweenaustenite and ferrite phase eld? Will the inclusion actas a carbon pump, i.e., deplete the carbon while coolingand reject it back to the austenite while heating to hightemperature? Further work is needed to elucidate thismechanism by careful experimentation with repeatedthermal cycling.

Recently, He and Edmonds [42] speculated that theformation of Fe–V clusters could act as nucleation sitesfor acicular ferrite, in addition to inclusions. In recentelegant work, Furahara et al. [**31] showed that addi-tion of V and N to C–Mn steel containing MnSinclusions led to the formation of intragranular idiomor-phic ferrite. Although, this ferrite forms above the bain-itic start temperature for this alloy, i.e. at and above873 K, the discussions in this work are very relevant tothe overall consideration of nucleation potency. Fura-hara et al. [**31] concluded that the reduced interfacialenergy between the VC and/or V(CN) and ferrite mightcontribute to an increase in the nucleation rate of intra-granular ferrite. Furahara et al. [**31] also found thatthe strains around the inclusion might promote acicular

ferrite at 823 K, a lower transformation temperature.Previous to this work, Guo et al. [*43] showed that theaddition of vanadium could lead to precipitation of VC on preexisting MnS sulde inclusions, which maylead to the nucleation of pearlite colonies in a hypereu-tectoid Fe–Mn–C alloy. In addition, they also found

that there is no preferred crystallographic orientationrelationship between MnS and VC, as well as betweenVC and austenite. Based on this observation, theauthors concluded that the kinetics of transformationis accelerated, possibly due to depletion of carbon inthe austenite as a result of precipitation of carbide.Based on these results, the nucleation of acicular ferriteon Fe–V clusters during weld cooling suggested by Heand Edmonds [42] appears to be improbable. The nucle-ation very well could have been accelerated due to theprecipitation of VC or V(CN). A simple thermodynamiccalculation of their alloy system shows that is possible toform V(CN) in their steels. However, the evidence forthe precipitation of VC or V(CN) was not present, be-cause the precipitates were very small. While theabove-mentioned models, in many ways, assume sim-plistic and uniform properties for inclusions, the realinclusions in the weld are very heterogeneous in compo-sition and often exhibit amorphous phases [*32,44].Therefore, a critical experiment is needed to add well-controlled particles that induce only one of the mecha-nisms in the austenite grain. In summary, for displaciveacicular ferrite nucleation on inclusions, the possibilityof thermal strains and local depletion of carbon and/or manganese appears to be the most plausible

mechanisms.In this section, the nucleation potency of inclusions

for nucleation of a single ferrite plate was discussed.Experimental observations show that there is no one-to-one correspondence of inclusion with every acicularferrite plate because acicular ferrite is also promotedby autocatalytic nucleation on a pre-existing ferrite plate(see Fig. 2 ). Recent work using electron back scattereddiffraction methods have related morphology to thecrystallographic misorientation and showed the impor-tance of autocatalytic nucleation of secondary ferriteplates on pre-existing ferrite plates [*10,**11] . Theimportance of autocatalysis will be discussed later inthe context of overall transformation kinetics of acicularferrite.

2.2. Growth of acicular ferrite

Acicular ferrite plates grow by repeated nucleation of ferrite subunits by displacive transformation with nopartitioning of carbon [**8]. Subsequent to the forma-tion of one subunit, the carbon escapes to austenite.With continued growth of acicular ferrite, the carbonin the austenite enriches and reaches a critical limit be-yond which the displacive formation of acicular ferrite



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becomes impossible. This critical carbon concentrationcorresponds to a condition ( T 00 condition) where the freeenergy of the austenite equals that of ferrite with astored energy of 400 J/mole. This cessation of the trans-formation at this point is referred as incomplete reac-tion. Experimental evidence for such incomplete

reaction during acicular ferrite transformation in weldshas been published [**8,45]. As mentioned earlier, theacicular ferrite growth also leads to invariant planestrain surface relief [**8]. The formation of sheaves sim-ilar to bainite are stied due to hard impingement of oneset of plates with another set of plates nucleated fromnearby inclusions. However, if the number density of potent nucleation sites is less, it may be possible to formsheaves even within the austenite grain due to autocata-lytic nucleation on primary acicular ferrite plates. Simi-lar to bainite, by applying an elastic stress, the selectionof a particular growth variant among the 12 differentNW- or 24 different KS orientational relationships canbe forced, leading to alignment of acicular ferrite plates(see Fig. 4 ). Interestingly, this result was observed in

conjunction with transformation-induced plasticity also[16]. It is important to note that this effect was observedwith the presence of an elastic stress.

An interesting question arises: what will happen if theaustenite deforms plastically and leads to extensive dis-location network? As per the displacive transformation

theory, the formation of stable dislocation networksmust retard the glissile interface and therefore must slowdown the acicular ferrite transformation kinetics. Thisphenomenon is referred to as mechanical stabilizationand it occurs in bainitic steels [**8]. A denitive proof for this phenomenon during acicular ferrite transforma-tion was obtained by Lee et al. [*46]. In their work, theaustenite grain boundaries were made ineffective by dec-orating them with allotriomorphic ferrite [45,47] therebyforcing the decomposition of austenite to acicular fer-rite. The austenite was then strained to 10%, 25% and45% before isothermally heat-treating at 390 C and360 C. The transformation kinetics of austenite to acic-ular ferrite were measured at both temperatures. Themeasurements from 390 C showed the retardation of acicular ferrite transformation kinetics with an increasein percentage strain, which is in agreement with the ear-lier hypothesis. Interestingly, the measurements at360 C failed to show any difference between strainedsamples. This was attributed to an increase in chemicaldriving force for transformation due to the increase inundercooling. This result shows that the growth of acic-ular ferrite in welds will be affected by complex thermo– mechanical–chemical conditions that may prevail withinthe austenite during weld cooling.

At this juncture, it is noteworthy for certain weld me-tal compositions under certain cooling conditions it maybe possible to form intragranularly nucleated Wid-manstätten ferrite above the B s temperature [*48]. Thisintragranular Widmansta ¨ tten ferrite will be coarser thanacicular ferrite that forms below the bainitic start tem-perature. All the mechanisms we have discussed will beapplicable except for the reduction in undercooling forthis transformation since it occurs at a highertemperature.

2.3. Transition from bainite to acicular ferrite

The discussion above shows that the acicular ferriteand bainite are similar in nature. An obvious questionarises: is it possible to switch the microstructure in thesame steel or weld by modifying the ratio of intragranu-lar to intergranular nucleation sites? Based on the pub-lished literature the answer is yes. It is possible toswitch between bainite and acicular ferrite microstruc-tures in three ways, as schematically shown in Fig. 5 .

The rst method relies on increasing the austenitegrain size, which in turn reduces the c – c grain boundaryarea per unit volume. Since bainite nucleation initiatesat the c – c grain boundaries, coarse austenite grains re-

Fig. 4. Optical micrographs from Fe–Cr–C weld metal samplessubjected to isothermal acicular ferrite transformation with an appliedelastic stress of (a) 12 MPa and (b) 174 MPa during isothermaltransformation of austenite to acicular ferrite [6].



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duce the bainite kinetics and favor the intragranularnucleation sites. This condition will lead to a predomi-nant acicular ferrite [49,**8] microstructure.

The next method relies on the removal of active c – cgrain boundary by decorating it with ne allotriomor-phic ferrite [45,47]. It is noteworthy that a favorably ori-ented a – c interface may also develop bainite sheaves[50]; however, it can be stied by the carbon enrichmentnear the a – c interface. This was proven by reducing thecarbon enrichment prole ahead of the a – c interface,which leads to the emanation of bainitic sheaves frompreexisting the a – c interfaces [6].

The third method is related to the presence of potentinclusions within the austenite grain. If there are noinclusions, acicular ferrite formation will not occur,and this was proved by Harrison and Farrar in their ele-gant experiment [51]. In addition, even if the inclusionsare present, if they are ineffective, bainite formation willbe favored once again. This was recently observed inmany cases by varying the titanium addition to the weldmetal [52–54]. At very low levels of titanium, the inclu-sions are predominantly manganese silicates, which areineffective acicular ferrite nucleants, and the weld metalmicrostructure is predominantly bainitic. However, onaddition of small amounts of titanium, the inclusioncharacteristics are modied and the microstructure ispredominantly acicular ferrite. This illustrates theimportance of inclusion engineering during the design

of weld metal consumable and inoculated structuralsteel development.

Elimination of c – c grain boundaries as potential

nucleation sites for bainite by boron segregation is alsobeen postulated by Peng et al. [55]. However, grainboundary allotriomorphic ferrite can be seen in theirmicrograph and therefore the mechanism for acicularferrite enhancement may be similar to the one shownin Fig. 5 . Snieder and Kerr have observed more than90% acicular ferrite without any grain boundary allot-riomorphs by adding titanium and boron [56]. Theseresearchers argued that the boron segregates to the aus-tenite grain boundary and sties the nucleation of allo-triomorphic ferrite at c – c grain boundary and thetitanium forms potent inclusions within the austenitegrain to promote acicular ferrite. These researchers didnot report any presence of bainitic sheaves originatingfrom the c – c grain boundary. This suggests that theboron may be retarding the bainitic nucleation too.

Another interesting suggestion for making the c – cgrain boundaries ineffective nucleation sites for bainitehas been attributed to possible segregation at c – c grainboundaries [57]. These authors found that, in a Fe-C-Mn structural steel inoculated with titanium oxide, theacicular ferrite microstructure can be promoted byincreasing the manganese concentrations from1.4 wt.% to 2.46 wt.%. At high concentrations of man-ganese, the authors observed no allotriomorphic ferrite

Fig. 5. Schematic illustrations of different mechanisms by which bainite transformation can be stied to obtain predominantly acicular ferrite: (a)Austenite grain size effect: with large grain size, the c – c grain boundary area per unit volume decreases for a given inclusion density. (b) Poisonedaustenite grain boundary: by forming a thin layer of allotriomorphic ferrite, the c – c grain boundary can be made ineffective due to reduced area of favorably oriented a /c interface, as well as, carbon enrichment near the interface (c) Presence of potent inclusions: for a given austenite grain if weincrease the number of potent inclusions, then acicular ferrite nucleation will be favored [6].



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and still observed enhanced acicular ferrite formation.Even with the presence of free c – c grain boundaries,there was no evidence for bainitic sheave formation.Therefore, these authors, based on the work of Liuand Zhang [58], suggested that the Mn segregates toc – c grain boundaries and reduces the driving force for

bainite nucleation at the interface. Liu and Zhang [58],concluded that this segregation is of a nonequilibriumtype and will depend upon the austenitization tempera-ture. Although, such segregation can occur even in alloywelds, changes in the austenite-to-ferrite driving forceunder paraequilibrium conditions needs to be quantiedto evaluate the relative effects of segregation.

Finally, based on the above discussions, one can won-der, if acicular ferrite will form all the time in welds. Pastresearch and recent work [54,59] shows that, if the weldcooling rates are too fast, even with the presence of po-tent inclusions a mixed bainite and acicular ferritemicrostructure will be obtained. For example, two weldmetal samples with different titanium contents wereaustenitized at 1200 C for 10 min and were cooled ata rate of 80 c/s. The micrograph in Fig. 6 a is of a weldmetal sample with no potent inclusions and thereforeexhibits predominantly bainite and martensite. Themicrograph shown in Fig. 6 b is of a weld metal samplewith potent inclusions exhibits grain boundary allotrio-morphic ferrite, predominant acicular ferrite microstruc-ture and small fractions of bainitic microstructure.Detailed microstructural analyses showed that theallotriomorphic ferrite was not continuous along theaustenite–austenite grain boundaries due to rapid cool-

ing conditions and as a result, the formation of bainitesheaves ( a b ) was been favored in some regions. Thismixed microstructure formation is also related to themechanisms presented in Fig. 5 .

2.4. Overall transformation kinetics of acicular ferrite

If the acicular ferrite formation is similar to that of bainite, is it possible to apply already existing bainite ki-netic models to acicular ferrite? Such an exercise wascarried out recently by Babu and David [54]. They ap-plied the overall transformation kinetic model devel-

oped by Chester and Bhadeshia given by the followingequation [*38,39,40] :

dnd t

¼ uK 1

h ð1 nÞð1 þ bhn Þ exp

K 2 RT

1 þDG 0m

r þ K 2ðDG 0m

G N ÞrRT

n ;ð4Þ

where K 1 is a parameter related to austenite grainboundary surface area per unit volume, u is the ferritesubunit volume, n is the extent of transformation, h isthe maximum bainite volume fraction that can form at

that temperature, K 2 is a constant, R is the gas constant,T is temperature in Kelvin, DG 0m is the maximum drivingforce for the nucleation of ferrite, r is another constantgiven by 2540 J/mole, G N is the universal nucleationfunction given by G N (J/mole) = 3.6375 T r, b is theautocatalysis factor. The autocatalysis factor is givenby the following equation:

b ¼ k1ð1 k2 xÞ; ð5Þwhere x is the bulk carbon concentration in the units of mole fraction and k1 and k2 are constants. Using isother-mal kinetic data for bainite one can derive the constantsK 1 , K 2 , k1 and k2 [*38,39] and the variation of subunitvolume with temperature. The K 1 parameter representsthe nucleation rate of bainite at any boundary. If theaustenite grain boundaries are free, K 1 for bainite for-mation will be related to grain boundary area per unitvolume ( S c

– cV ) [40]. If the austenite grain boundaries

are decorated by the allotriomorphic ferrite, it will be re-lated to the a – c interfacial area per unit volume ðS c

– aV Þ

Fig. 6. Optical micrographs of weld samples continuously cooled from1200 C at a rate of 80 K/s. (a) A predominantly bainite and martensitemicrostructure was observed due to lack of titanium containing potentinclusions. (b) Predominantly acicular ferrite microstructure wasobtained due to presence of titanium containing inclusions. However,a small fraction of bainite was also found due to lack of continuousallotriomorphic ferrite coverage along the austenite grain boundary[54].



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that is capable of nucleating bainite [50]. In the case of acicular ferrite, the K 1 parameter is related to the inclu-sion–austenite interfacial area ( S c

– incV ) that is capable of

nucleating acicular ferrite subunits [**41,*28,*24] . Babuand David [54] evaluated experimental measurements of volume fraction as a function of temperature duringcontinuous cooling conditions (see Fig. 7 a and b) fortwo weld samples using Eq. (5). In the weld containingno titanium addition, the austenite transforms predom-inantly to bainite and in the weld containing titaniumthe austenite transforms predominantly to acicular fer-rite (see Fig. 7 c and d). The tted K 1 value for bainiticdata was (1.0 · 1017 ) higher than the one (4.4 · 1016 ) ob-tained for the acicular ferrite data. The autocatalyticfactor ( k1 ) for acicular ferrite (34) was found to be high-er than that of bainite (4). It is important to note thatthese values are achieved by purely tting the equationto experimentally observed kinetic data and are not de-rived from fundamental knowledge of nucleation po-tency. Nevertheless, it is interesting to compare thetted parameters.

The increase in k1 for acicular ferrite may be possiblyrelated to the exibility to choose many variants fromthe inclusion, which in turn may allow for many morevariants due to autocatalysis effect. Moreover, in thecase of bainitic microstructure, the maximum experi-mental ferrite fraction achieved at 737 K is 0.81. Below737 K at 680 K, the transformation rate of austenite is

accelerated further due to formation of martensite. Incontrast, the maximum experimental ferrite fractionachieved at 737 K is 0.97 for predominant acicular fer-

rite microstructure. This result shows that the rate of transformation to acicular ferrite higher than bainitictransformation. This change may be related to largercarbon trapping between the acicular ferrite subunitsthan that between the bainitic ferrite subunits [39]. Thisdiscussion shows it is possible to describe the acicularferrite transformation kinetics by considering changesin S c

– cV (austenite–austenite grain boundary area per unit

volume), S c – a

V (austenite–ferrite interface boundary areaper unit volume), and S c

– incV (inclusion–austenite inter-

face boundary area per unit volume) on K 1 and k1 as afunction of weld metal composition and austenite grainsize. Further work is needed to evaluate these valueswith large number of experimental data set on acicularferrite transformation.

3. Future work

In the beginning of this review, it was stressed thatacicular ferrite formation depends on various high-tem-perature phase transformation sequences (see Fig. 1 ).Therefore, these reactions must follow a particulardirection to allow for acicular ferrite formation atlow temperature. This points to the obvious need for

Fig. 7. Comparisons of overall transformation kinetic model (black lines) predictions with the experimentally measured isothermal bainite andacicular ferrite kinetics in Fe–C–Mn steel welds (open circles) with two different inclusion characteristics during continuous cooling at 50 K/s. (a)Data from a sample with no intentional titanium addition shows sluggish transformation kinetics compared to (b) rapid kinetics measured from thesample with 17 wt. PPM titanium addition. (c) Corresponding optical micrograph from the sample with kinetic data shown in (a) showspredominantly bainite microstructure. (d) Optical micrograph from the samples with kinetic data shown in (b) shows predominantly acicular ferritemicrostructure [54].



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an integrated model that describes inclusion formation,primary solidication to d-ferrite, d-ferrite to austenitetransformation, austenite grain growth, and austenitedecomposition to different ferrite morphologies. Thereare recent developments to describe inclusion formationduring solidication [37,60,61*,62,63*,64–71] . There ex-

ist dendrite-tip growth and diffusion-controlled-growthmodels that can be used to describe both phase selectionand kinetics of solidication, respectively. These modelshave been extended to weld solidication [72,73*,74] .Diffusion controlled growth models can be used to de-scribe peritectic solidication [75]. The same modelscan be used to describe diffusion-controlled growth of austenite into d-ferrite [74,75]. With further develop-ment, these models can be coupled with Monte Carlotype grain growth models [76] to describe the nal stateof an austenite grain before the onset of ferrite forma-tion. Finally, the allotriomorphic-, Widmansta ¨ tten-,bainitic- and acicular-ferrite kinetic models [**77] mustbe linked through a simultaneous transformation kineticframework [78,79].

4. Conclusions

High-temperature transformation sequences in weldsthat affect acicular ferrite formation were highlighted.Evidence for the similarity between bainite and acicularferrite transformations is overwhelming. Using nucle-ation on inclusions and autocatalytic nucleation on pre-existing ferrite plates, the nucleation of acicular ferrite

can be described. The most probable mechanisms fornucleation on inclusions are either due to the generationof thermal strains in the austenite or the formation of solute-depleted regions. The experimental measurementsof acicular ferrite growth under externally imposed elas-tic and plastic strains illustrated the complex interplaybetween thermal–mechanical–chemical–crystallographicparameters. Now, it is possible to model the acicular fer-rite formation by using bainitic kinetic models. Theimportance of computational models that describe theoverall microstructure evolution in the context of pre-dicting acicular ferrite formation is stressed.

Acknowledgments

Research sponsored by the Division of Materials Sci-ences and Engineering, and Assistant Secretary for En-ergy Efficiency and Renewable Energy, IndustrialTechnologies Program, Industrial Materials for the Fu-ture, US Department of Energy, under contract DE-AC05-00OR22725 with UT-Battelle, LLC. The authoralso thanks Dr. Z. Feng of ORNL for translating oneof the research paper published in Chinese language.The author thanks Drs. J. M. Vitek and R. L. Klueh

of Oak Ridge National Laboratory for helpful discus-sions on this paper.

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