Interaction of Precipitation With Recrystallisation and Phase Transformation In

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REVIEW

Interaction of precipitation withrecrystallisation and phase transformation inlow alloy steels

D P Dunne

The processes of precipitation restoration and phase transformation can interact in complex

ways during thermomechanical processing of microalloyed steels to profoundly alter their

structures and properties Precipitation in austenite during hot deformation can strongly modify

the kinetics of recovery and recrystallisation subsequently affecting the nucleation and growth of

ferrite during cooling For steels containing strong carbidenitride formers interphase precipita-

tion (IP) can occur in ferrite at the austeniteferrite interface conferring significant coherency

strengthening Much of what is known about this phenomenon is attributable to the impressive

research efforts of Robert Honeycombe and his colleagues at Cambridge

Keywords Interphase precipitation Alloy carbonitrides Microalloyed steels Hot deformation Phase transformation

This paper is part of a special issue in memory of Professor Sir Robert Honeycombe

PreambleIntensive research has been conducted since the 1960s onthermomechanical controlled processing (TMCP) ofcommercial microalloyed steels The results of this workhave been extensively published in the internationalliterature and have addressed many of the problems inalloy design and process control Many excellent reviewshave been written on the role of alloy carbonitrideprecipitation on the structure and properties of hotrolled microalloyed steels1ndash4 and although these sourcesare referred to where relevant it was not my intention toproduce another comprehensive review Rather I haveattempted to highlight some of the important contribu-tions to the understanding of the science underlying theprecipitation process that were made by HoneycombersquosAlloy Steels Group as well as referring in some detail torelated Australian work based on research projectsconducted by the University of Wollongong in colla-boration with BHP Steel at Port Kembla The outcomesof this Antipodean work are not well known in theinternational arena because of only limited publicationmostly in Australian journals and national and interna-tional conference proceedings This research effort wasconcentrated largely on specific commercial microal-loyed steels produced by BHP but it resulted in thediscovery and elucidation of some novel generalphenomena I gratefully acknowledge that the seeds ofthis work originated in the experience knowledge andstimulation that I gained during a study leave atCambridge in 1977ndash1778 working with RobertHoneycombe and the Alloy Steels Group This study

leave served as an opportunity for me to re-focus myresearch pursuits on my return to Australia

I was generously supported and encouraged atCambridge by Robert Honeycombe to conduct researchon an aspect of interphase precipitation (IP) in alloysteels I had the privilege of working in a research groupthat included David Edmonds Barry Muddle PaulHowell and John Bee Harry Bhadeshia and PeterSouthwick were PhD research students at this timeworking respectively and enthusiastically on bainitictransformations and duplex stainless steels LindsayGreer was also a PhD student and I recall that he assistedme in learning to operate a DSC as a tool for anotherresearch strand that I pursued with Mike Stobbs onthermoelastic martensitic transformation in FendashPt alloys

My research activities leading up to my year inCambridge started with PhD candidature from 1964 to1968 at the University of New South Wales supervisedby John Bowles who in collaboration with JockMackenzie developed the phenomenological crystal-lographic theory of martensite transformation My PhDresearch topic was shape strain measurement in (225)F

and (3 10 15)F plate martensites in FendashC FendashMnndashC andFendashNindashC alloys This research was followed by nearlythree years at the University of Illinois working withMarvin Wayman on the crystallography of ferrousmartensites and thermoelastic martensitic transforma-tion in FendashPt alloys

My return to Australia to an academic position at theUniversity of Wollongong in 1970 required a period ofadjustment to the lsquoteaching tradersquo and allowed onlylimited research on ferrous martensitic transformationsand shape memory alloys For these reasons the invita-tion from Robert Honeycombe to conduct research atCambridge presented an excellent opportunity to takestock and to diversify my research portfolio by

Faculty of Engineering University of Wollongong Northfields RoadWollongong NSW 2522 Australia

Corresponding author email druceuoweduau

410

2010 Institute of Materials Minerals and MiningPublished by Maney on behalf of the InstituteReceived 26 May 2009 accepted 29 August 2009DOI 101179026708309X12526555493350 Materials Science and Technology 2010 VOL 26 NO 4

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branching out into research fields that were potentiallymore funding-friendly than martensite crystallography

A significant part of the research conducted in theAlloy Steels Group was focused on clarifying funda-mental aspects of alloy carbide and nitride precipitationin laboratory prepared ternary alloys with the ultimateaim of fuller understanding of the influence of structureon the mechanical properties of commercial microal-loyed steels The project that I proposed to RobertHoneycombe fitted with this research emphasis andconsisted of an experimental study of the restorationbehaviour of cold rolled FendashVndashC alloy that had beenheat treated before cold rolling to produce IP of V4C3

Precipitation and recrystallisationI had previously published a paper on lsquopancakersquo grainformation in batch annealed Al killed deep drawingsteels5 The development of grain shape anisotropy in anannealed steel fascinated me and I set out to map theevolution in ferrite grain shape through the cold rollingand recrystallisation processes used for this type of steelWhen lsquofinished hot and coiled coldrsquo these steels aresupersaturated with AlN which subsequently precipi-tates out during batch annealing and interacts with therecrystallisation process Recrystallisation is severelyimpeded and is nucleation limited resulting in largerecrystallised grains with a preferred orientation that isfavourable for sheet forming by deep drawing The workthat I carried out showed that the description of therecrystallised grain shape as lsquopancakedrsquo is incorrect assome linear anisotropy is retained effectively mimick-ing albeit to a reduced extent the slab shape presentafter cold rolling Figure 1 shows that the recrystallisedgrains have similar planarndashlinear anisotropy to that ofthe cold rolled ferrite grains Cross-rolling instead ofunidirectional rolling followed by simulated batchannealing resulted in no linear anisotropy and the grainshape was truly lsquopancakedrsquo Recrystallised grain shape

anisotropy arises because the AlN precipitates on theboundaries of the subgrains and grains of the unidir-ectionally cold rolled steel pinning these boundaries andrestricting the lsquonucleationrsquo of recrystallised domains bygrain boundary bulging or subgrain growthcoalescenceOnce nucleated the growth of recrystallised grainsproceeds within a forest of impeding particles distrib-uted in planarndashlinear layers in the deformed matrixTherefore growth is anisotropic

My interest in and curiosity about the interactionbetween precipitates and restoration (andor poly-morphic transformation) led naturally to the projectthat I undertook at Cambridge ferrite restoration ofcold rolled ferrite in a Fendash105Vndash023C (wt-) alloycontaining networks of fine coherent interphase pre-cipitate particles that were present before cold rolling7

The heat treatment used to generate the startingstructure was annealing under argon at 1160uC for15 min followed by isothermal transformation in a saltbath for 15 min at 747uC According to Batte andHoneycombe8 this treatment results in a volumefraction of 00123 of V4C3 precipitates of 3ndash5 nm radiusCold rolling of 75 mm strip was used to reduce thethickness by 01 mm per pass and samples were takenafter each 10 reduction up to 60 in order toinvestigate recrystallisation kinetics at 706uC

Cold rolling resulted in a dense and uniform distribu-tion of dislocations as well as the development of shearbands for reductions greater than 20 The shear bandswere typically 02ndash4 mm thick and oriented at an angleof about iexcl30u to the rolling direction The shear bandsubstructure consisted of elongated subgrains 01 mmthick and provided potent sites for nucleation ofrecrystallisation Although recovery and recrystallisa-tion occurred rapidly within the bands on annealing of60 cold rolled samples at 706uC restoration ofsurrounding material was severely retarded and recrys-tallisation was incomplete after 1000 h The rate

1 Linear VLIN and planar VPL anisotropy factors for ferrite grains after a cold rolling and b cold rolling and batch anneal-

ing as a function of percent cold reduction Following Underwood6 VPL is defined as the ratio SV(PL)SV(TOT)

VLIN5SV(LIN)SV(TOT) and VPLndashLIN5[SV(PL)zSV(LIN))SV(TOT)] where SV is grain boundary surface area per unit volume and

SV(TOT) includes the planar linear and isometric values which are determined from grain boundary intercept densities

in the rolling direction the transverse direction and the direction of the rolling plane normal56

Dunne Precipitation recrystallisation and phase transformation in low alloy steels

Materials Science and Technology 2010 VOL 26 NO 4 411

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controlling step in migration of recrystallising interfaceswas the rate of coarsening of pinning particlesFigures 2a and b shows coarsened V4C3 particles inthe wake of the advancing recrystallising interfacePrecipitates in the unrecrystallised regions remainedcoherent longer and coarsened more slowly than thoseincorporated into the recrystallised grains

Grain shape anisotropy was exhibited by the recrys-tallised grains because of the distribution of precipitatesafter cold rolling However the grain elongation(average length to average width in a transversedirection section) decreased with grain growth fromabout 15 to 12 on annealing for 1000 h at 706uC(Fig 3) Owing to the extended time required fordissipation of the stored strain energy the three classicalrestoration processes (recovery recrystallisation andgrain growth) operated simultaneously and competi-tively in reducing system energy

This experimental alloy provides a relatively highvolume fraction of carbide precipitates facilitating thestudy of their morphology and crystallography Incontrast commercial microalloyed steels have lowvolume fractions of precipitates and are subject to thecomplicating interactive effects of multiple elements aswell as more complex thermomechanical treatmentNevertheless commercial alloys are at the production(paydirt) end of the metallurgical spectrum and it was

there post-Cambridge and in collaboration with BHPSteel that I decided to invest a major research effort

Precipitation and phase transformationin low alloy steelsThe significant interaction between precipitation andphase transformation is well illustrated by the evolu-tionary advances since the 1970s in procedures for thehot rolling of plate and strip steels9 These advanceshave been underpinned by developments in both steeldesign and TMCP Microalloyed steels containing smallamounts of the strong carbidenitride formers Nb Tiand V have allowed the production of strong and toughhot rolled steels with improved weldability due toreduced carbon content

Fundamental aspects of IPThe carbides and nitrides of the alloying elements usedin microalloyed steels have fcc crystal structures withclosely similar lattice spacings Isomorphism is thereforepossible allowing mixing of both the metallic elementsand carbon and nitrogen In general the precipitatespecies in microalloyed steels is a carbonitride X(CN)where X is Ti Nb or V or combinations of theseelements depending on the steel composition

Honeycombersquos Alloy Steels Group took up thechallenge of advancing fundamental understanding ofthe physical metallurgy of microalloyed steels by study-ing more highly alloyed ternary or quaternary labora-tory steels that are less complex than commercial steelsIn particular the group clarified the nature andmechanism of IP establishing that various types of IPare possible depending on the temperature of isothermaltransformation andor the nature of the austeniteferriteinterface The most well characterised form of IPconsists of planar layers of particles with a uniforminterlayer spacing that form at periodically staticsemicoherent interfaces before being incorporated intothe ferrite by lateral propagation of incoherent steps orledges [interphase precipitation (planar) ndash IPP]1011

In addition curved layers of precipitates have beenreported in steels containing Cr12 and V13 These layersare related to curved incoherent boundaries between theaustenite and ferrite and have been described asinterphase precipitation (curved) (IPC)1213 In this casesolute accumulation at a moving ac interface can arrestits motion long enough for precipitation to occur and

a b

2 a Optical micrograph of the FendashVndashC alloy showing partially recrystallised structure after annealing of a 60 cold

reduced sample at 706uC for 905 h3 b TEM image showing a boundary between a recrystallised grain and the

deformed matrix of a 60 cold rolled sample of the FendashVndashC alloy annealed at 706uC for 1008 h7 Note interface bow-

ing around pinning particles

3 Average diameter and elongation ratio of recrystallised

grains in transverse direction section as function of

annealing time at 706uC (log scale) elongation ratio is

average mean intercept length to average mean inter-

cept width of ferrite grains7


412 Materials Science and Technology 2010 VOL 26 NO 4

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induce more effective interfacial pinning Eventualunpinning allows further advancement of the interfacebefore the process repeats itself and the ferrite grain isdecorated with curved parallel layers of fine precipitatesRicks et al12ndash14 proposed an additional subdivision ofregular (Reg) and irregular (Irreg) layer spacings It washypothesised that regular spacing can arise from theformation of lsquoquasi-ledgesrsquo due to localised but restrictedbreak-out of the interface that is followed by lateralpropagation of the ledges to incorporate another layer ofparticles into the ferrite This mechanism is similar to thatoccurring at a semicoherent interface to give rise to IPPIn the irregular case the interfacial pinning processoccurs more erratically leading to layers of variablespacing In all of these cases surprisingly the plateshaped fcc carbide precipitates are characterised by asingle variant of the BakerndashNutting orientation relation-ship15 with the bcc ferrite 100fcc||100bcc andn011mfcc||n001mbcc It was concluded by Honey-combe1116 that of the three possible variants the oneselected makes the smallest angle with the interface inorder to maximise growth kinetics by interfacial diffusionand to reduce interfacial energy

There are two other types of IP that have beenidentified by Honeycombe and his research collea-gues213 One is a fibrous form (IPF) observed typicallyin higher carbon higher alloy steels Fibrous carbideprecipitation has been reported in Mo17 and V8 steelsunder conditions favouring slow coupled growth at anincoherent interface Interphase precipitation (fibrousform) is essentially a eutectoid type transformationproduct similar to lamellar pearlite but with finercarbide particles and a smaller volume fraction ofcarbide The second type is referred to as interphaseprecipitation (random) (IPR) because the single variantparticles are randomly dispersed rather than distributedin layers If a migrating incoherent ac interface isundulating due to transient pinning by precipitates IPRcould arise Unlike a planar or smoothly curved inter-face precipitation at such a boundary occurs in anuncoordinated way (Fig 4a) However it should benoted that an apparently random distribution ofparticles may occur for planar precipitate arrays becauseof the orientation of the foil surface plane relative to thelayer plane The Honeycombe group reported layerswith spacings typically between 5 and 30 nm for themore concentrated alloys that they investigated whereas

Smith et al18ndash22 found that layer spacings were between15 and 150 nm for the more dilute commercial steelsthat they studied For an untilted 200 nm thick TEMfoil IPP with layer spacings of 20 and 75 nm will onlyappear to be in the form of distinct layers in the image ifthe layers are respectively aligned closer than 6 and 22uto the foil normal1819 For small spacings thereforeparticles in layers can be easily construed as randomlydispersed particles The layer morphology may also beindistinct because the precipitates have formed with acoarse and irregular spacing by the IPC (Irreg) mode

The type of IP is temperature dependent because thenature of the transforming ac interface is temperaturesensitive Incoherent boundaries are favoured by hightransformation temperatures and semicoherent bound-aries provide the most energetically frugal and kineti-cally preferred boundaries at lower temperaturesTherefore the dominant form of the precipitate layerswould be expected to shift from IPC (Irreg) to IPC(Reg) to IPP with decreasing temperature However thevariable nature of the transforming interface at a giventemperature will ensure that there is considerableoverlap of precipitate types Furthermore it has beenestablished that different types of IP can occur withinthe same ferrite grain associated with a change in thenature of the interface or interfacial boundary segmentsof different character2819 It should be noted that thereis a finite temperature window over which IP can occureven if carbonitride precipitation is thermodynamicallypossible over a wider temperature range Precipitationcan be suppressed on transformation of austenite bothat high and low temperatures The resulting super-saturated ferrite is amenable to subsequent precipitationof carbonitride during continuous cooling isothermalholding or subsequent aging In such cases all threevariants of the BN relationship are equally feasibleleading to multivariant precipitation typically nucleat-ing at dislocation sites

Interphase precipitation in commercial steelsInvestigations of structurendashproperty relationships inmicroalloyed steels accelerated sharply in the 1960snot only in universities such as Sheffield whereHoneycombe and his colleagues were active but alsowithin research laboratories of major steel producersand materials resource industries eg British SteelUnited States Steel Great Lakes Steel Corp and UnionCarbide This decade was an especially fruitful one inwhich the first electron micrographs of row precipitationof NbC were reported initially by Morrison in the UK25

and then by Gray and Yeo26 in the USA Theobservation of these precipitates confirmed the inferencedrawn in earlier reports27ndash30 that precipitation in ferriteresulting from the addition of small concentrations ofNb V or Ti could produce significant strengthening oflow carbon steels Initial problems with toughness wereovercome by the development of controlled rolling torefine the prior austenite grain size and effect transfor-mation to fine grained ferrite3132 Curiosity about themechanism of precipitation and how it can be optimisedto improve strength and toughness of low carbonstructural steels drove much of Honeycombersquos researchinitially at Sheffield and then at Cambridge followinghis appointment as Goldsmith Professor of Metallurgyin 1966

4 a Schematic diagram illustrating how IPR can arise by

uncoordinated precipitation of alloy carbidenitride par-

ticles at advancing incoherent ac interface and b dia-

gram showing how coordinated precipitation can result

in curved layers of particles in ferrite (IPC)18



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The modes of IP identified and characterised byHoneycombersquos group using laboratory ternary alloyshave largely been confirmed in more dilute multicom-ponent microalloyed steels in the isothermally trans-formed and hot rolled conditions For example researchat the University of Wollongong18ndash22 on commercial lowcarbon V Nb and Ti microalloyed steels has demon-strated that apart from IPF isothermal transformationof these steels resulted in the same types of IP of alloycarbonitride particles that were documented byHoneycombersquos group for higher alloy experimentalsteels Moreover continuously cooled hot rolled stripsamples showed the same types of IP as for isothermallytransformed samples

Early work was concentrated on a commercially hotrolled 007C 014V alloy18ndash20 This alloy has aboutone-third of the C and one-seventh of the V for theternary alloy studied at Cambridge Nevertheless re-austenitising at 1200uC followed by isothermal trans-formation in molten salt at selected temperatures in therange 820ndash650uC showed that IPC (Irreg) was dominantfor temperatures 810uC (Fig 5) with IPC (Reg) andIPP being prominent for temperatures 810uC (Fig 6)Interphase precipitation (planar) increased in frequencywith decreasing temperature

Hot deformation by up to 50 by rolling in thetemperature 920ndash800uC resulted in profound refinementof the precipitate size and spacing on subsequent iso-thermal transformation at a lower temperature Figure 7shows IPP in a sample of the V steel which had been

reduced 50 by rolling at 800uC and quenched after a2 min hold

The precipitate layers are typical of IPP and the layerspacing (y80 nm) is close to that shown in Fig 6 for thetransformation from undeformed austenite at a tem-perature 50uC lower22 Progressive transformation toferrite on holding at 800uC resulted in ferrite grainrefinement well in excess of that expected from theincrease in austenite grain surface areaunit volume SV

by deformation22 For gt50 reduction of austenite witha coarse starting grain size (150 mm) the transformationproduct consisted of a series of layers of fine ferritegrains22 produced by a process termed lsquocascadersquonucleation23 This behaviour is different to that typicallyfound for Nb steels in which continued ferrite formationtends to proceed by selective growth of preferredgrains24 It was inferred that as a result of highersolubility of V(CN) in austenite solute drag andcopious IP during transformation limit boundarymobility allowing nucleation of new grains in successivelayers ahead of the advancing transformation frontAnother important conclusion of this work was thattransformation to ferrite from deformed austenite is ineffect a surrogate recrystallisation process driven byboth stored strain energy and chemical free energy

As a result of continuous cooling during transforma-tion of either recrystallised or deformed austenite incommercial processing the nature of the ac interface islikely to be highly variable Therefore IP is not expectedto be as extensive nor as well developed or readilyobservable as in isothermally transformed samplesNevertheless both IPP and IPC have been observed incommercially hot rolled steels Figure 8 shows IPP in thecommercially rolled V steel In addition carbonitrideprecipitation has been found on sub-boundaries of thedeformed austenite These particles do not exhibit theBN orientation relationship and have clearly formed indeformed austenite during finish rolling before beingincorporated into the ferrite formed during continuouscooling The important role of carbonitride precipitationin austenite in controlled non-recrystallisation hot roll-ing is discussed further in the section on lsquoPrecipitationand retardation of austenite recrystallisationrsquo It isapparent that carbonitride that is preprecipitated inaustenite is lost to precipitation in and strengthening of

5 Dark field electron micrograph using 002VC electron

diffraction spot for V steel isothermally transformed for

180 min at 810uC18 irregularly spaced curved layers of

V(CN) are present ie IPC (Irreg) bar represents 1 mm

6 Bright field electron micrograph of V steel isothermally

transformed at 750uC for 15 min1819 precipitate layers

are typical of IPP bar represents 05 mm

7 Bright field electron micrograph of V steel deformed

50 at 800uC then isothermally transformed at 800uCfor 2 min before quenching18 bar represents 05 mm



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ferrite that forms on cooling Nevertheless the fullpotential for precipitation of the remaining carbonitridein solution in ferrite is unlikely to be realised duringcommercial TMCP of strip steels

Effect of precipitation on ferrite strength andtoughnessThe effect of isothermal holding time on the hardness offerrite produced in the temperature range 600ndash750uC fora commercial NbzV microalloyed steel is shown inFig 9a In general the hardness of the first formedferrite grains increased with decreasing temperature inthe range 750ndash650uC Interphase precipitation is a majormicrostructural contributor to the hardness (strength)and its effect intensifies with decreasing temperature offormation because of decreasing particle size and layerspacing933 Holding at the isothermal transformationtemperature resulted in a decrease in hardness mainlydue to particle coarsening Another more minor soften-ing factor is the recovery of the dislocation substructureresulting from the polymorphic transformation Anexceptional result occurred for 600uC The initialhardness was relatively low because the ferrite formedwithout IP However isothermal holding resulted in ahardness peak due to multivariant precipitation of VC

from the supersaturated ferrite These hardness changesmirror trends in yield strength which indicate thatprecipitation hardening can increase the yield stress ofNb microalloyed steels by up to 200 MPa dependingon the particle size and volume fraction33

Experiments on aging of commercially hot rolled VNb and VndashNb strip steels demonstrated that they hadadditional precipitation hardening capacity because theferritic transformation product formed on continuouscooling remained supersaturated with alloy carbonitride(Fig 9b) This additional precipitation is clearly distin-guishable from IP through its multivariant nature2021

Ferrite grain refinement is a strong toughness enhan-cing factor for structural steels and carbonitride pre-cipitation in both austenite and ferrite can affect the asrolled ferrite grain size The solubilities of Ti Nb and Vcarbonitrides in austenite have been thoroughly exam-ined2 and it has been concluded that Ti carbonitridesare least soluble and V carbonitrides have the highestsolubility Furthermore the nitrides are more stablethan carbides and coarsen more slowly on isothermalholding in the austenitic state2 Undissolved carboni-trides can limit austenite grain size during reheating forhot rolling and thereby enhance refinement throughrecrystallisation during successive hot rolling passesPrecipitation of fine carbonitrides during the finishrolling stage can strongly inhibit recrystallisation andensure that stored strain energy catalyses ferrite nuclea-tion during cooling below the Ar3 temperature as well asreducing growth rate by IP The net result is a finegrained ferrite structure with impact transition tempera-tures as low as 250uC combined with yield strengthsabove 400 MPa33

Interphase precipitation in Cu bearing steelsRicks et al14 established that the IP mode in steels is notunique to carbonitrides by demonstrating the presenceof IP in FendashCundashNi alloys Interphase precipitation(planar) of e-Cu was obtained by isothermal transfor-mation of Fendash2Cundash2Ni alloy at 720uC This form ofprecipitation has also been reported for a commercialASTM A710 type steel produced by BHP Steel with thecomposition of 0055Cndash140Mnndash025Sindash085Nindash110Cundash002Nbndash0013Tindash0012Pndash0003Sndash00075N34ndash37 This typeof steel is being increasingly used as a more weldable

8 Bright field electron micrograph of V steel in as hot

rolled condition IPP of V(CN) is evident1819 layer spa-

cing is 15 nm and bar represents 02 mm

9 a Effect of isothermal holding temperature and time on hardness of first formed ferrite in commercial NbzV steel re-

austenitised at 1200uC for 15 min then quenched into molten salt at each of the indicated temperatures21 (steel com-

position 009Cndash100Mnndash0051Nbndash0057Vndash005Alndash0012N) and b aging responses at 500uC of three commercial TMCP

strip steels containing 008C and 014V 0057Vz0051Nb and 0049Nb (Ref 21)



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substitute for HY80 as structural plate for naval andoffshore structural applications Although normallyused in the quenched and tempered condition (ASTMA710 grade A class 3) BHP Steel developed theabove more alloy lean steel for production by a TMCProute The processing schedule involved recrystallisationrolling to produce fine recrystallised austenite non-recrystallisation finish rolling to ensure that ferritetransformation occurred from deformed austenite con-trolled cooling to 500ndash550uC to minimise e-Cu pre-cipitation and finally aging at 550uC for 30 min toinduce e-Cu precipitation37 The specified minimumyield stress of this steel designated CR HSLA 80 is550 MPa

The e-Cu precipitates that were observed in hot rolledand aged steel were often in arrays of a single variantThe orientation relationship between the fcc e-Cu andbcc a-ferrite is KurdjumovndashSachs 111e||110a andn110me||n112ma Although there are 24 possible variantssingle variant planar arrays were observed both inisothermally transformed samples (see Fig 10) and alsoin rolled and aged material In many cases the pre-cipitates of e-Cu were apparently in a random distribu-tion but with a single orientation variant Theseobservations suggest that the precipitate arrays form inan uncoordinated way during cooling in associationwith the motion of an incoherent interface (see Fig 4a)Subsequent aging may have served to coarsen theparticles without substantial new multivariant precipita-tion The apparently random distribution of particlesmay arise from a layered distribution that is not obviousbecause of the foil section relative to the layers asmentioned previously or because they have formed bythe IPC (Irreg) mode

After solution treatment and cooling slowly enough togenerate a ferrite plus pearlite structure samples wereaged at 500uC The hardness increased above that of theunaged condition (185 HV) due to precipitation hard-ening before falling due to particle coarsening andoveraging (Fig 11) Alternative heat treatments werecarried out that involved re-austenitising at 1200uC for15 min followed by quenching to ambient to formmartensite and isothermal transformation to bainite at440uC for 5 min The bainitic and martensitic sampleswere also aged at 500uC resulting in significant

secondary hardening because of multivariant precipita-tion e-Cu36 The peak hardnesses shown in Fig 11 are225 HV (ferrite) 245 HV (bainite) and 292 HV (mar-tensite) but the hardness increment relative to the basehardness is similar in each case (in the range 25ndash40 HVpoints) These results and those shown in Fig 9 indicatethat alternative processing routes to conventionalTMCP can produce significantly higher strength levelsby taking full advantage of the precipitation hardeningpotential of the microalloyed steel as well as the form ofthe matrix phase

Despite the typically adverse effect of precipitationhardening on impact toughness33 this low carboncopper bearing steel exhibited Charpy V notch (CVN)values in the TMCP and aged condition that exceededthe minimum specified for ASTM A710-84 (27 J at245uC) as well as for HY80 (MIL-S US militaryspecification 47 J at 251uC) The minimum CVN valueswere also found to be surpassed for simulated heataffected zone (HAZ) microstructures equivalent towelding of 20 mm plate without preheat at a heat inputof 29 kJ mm2138 However after post-weld heat treat-ment at 550uC for 1 h the average CVN for thesimulated grain coarsened HAZ was 18 J lower thanthe MIL-S specified minimum of 47 J at 251uC Incontrast post-weld heat treatment at 450 or 650uC for1 h gave CVN values 120 J and so it was con-cluded that significant rehardening by precipitation ofe-Cu at 550uC compromised the impact toughness38 Thelow impact toughness is associated with maximumprecipitation strengthening by fine precipitate particlesor zones (5 nm in diameter) However for the solutiontreated and the overaged conditions exceptional tough-ness is exhibited In the latter case it is likely that thecoarsened e-Cu particles deform plastically absorbingimpact energy rather than cracking or decoheringfrom the matrix in the stress concentration zone aheadof a sharp crack Tomita et al39 proposed a similarinterpretation for the unexpectedly high toughness ofhigh strength Cu bearing steels They examined a rangeof low carbon Cu bearing steels designed for a 590 MPaminimum yield stress and suitable for medium to highheat input welding of 25ndash80 mm plate The steelsdeveloped showed excellent Charpy impact energies(260uC) and fracture toughness for the grain coarsenedHAZ at medium (5 kJ mm21) heat input as well as

10 Bright field electron micrograph of Cu steel isother-

mally transformed at 600uC for 160 min particles are

predominantly single variant and layer morphology is

consistent with IP3435 bar represents 05 mm

11 Aging response curves of Cu steel samples at 500uC

in ferritic bainitic and martensitic conditions36



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acceptable results for a high heat input (10 kJ mm21)They concluded that the weld thermal cycle dissolved theCu and it was not reprecipitated for weld cooling ratesthat exceeded 5uC min21 [008uC s21 or a cooling timebetween 800 and 500uC (Dt825) of y1 h]

Precipitation and retardation ofaustenite recrystallisationIt is well known that finish rolling of Nb steels atrelatively low temperatures significantly increases therolling load This increase in load is widely considered tobe due to NbC precipitation in austenite whicheffectively prevents austenite softening by recrystallisa-tion334041 Research on hot deformation of austenite inV and Ti bearing steels has also demonstrated that staticrecrystallisation is profoundly retarded by deformationbelow a specific temperature in the approximate range1000ndash850uC which for the purpose of industrial hotrolling is called the lsquonon-recrystallisationrsquo temperaturesince finishing below it ensures that ferrite forms fromdeformed austenite Strain in the austenite catalysesprecipitation and the precipitates then stabilise thedeformed austenite against restoration by recovery andrecrystallisation Subsequent cooling produces phasetransformation in a heterogeneous environment with ahigh density of potential nucleating sites Substantialferrite grain refinement occurs provided that nucleationrather than grain growth dominates in driving trans-formation The resulting fine grained hot rolled steeltypically exhibits a yield stress well in excess of350 MPa a minimum often used to define a low carbonstructural steel as a high strength low alloy steel Thetoughness is simultaneously increased because of grainrefinement

Despite the inference drawn from hot mechanicaltesting that alloy carbonitride precipitation in deformedaustenite is responsible for retarded recrystallisationdirect evidence remained sparse for many years becausethe polymorphic transformation of austenite to ferritedestroys the nexus between the precipitates and thedeformed austenite This uncertainty generated somecontroversy about whether it was Nb in solution or NbCprecipitation that was responsible for the retardedrestoration of deformed austenite in Nb microalloyedsteels This difference evaporated over time with therealisation that the issue was not lsquoeitherorrsquo and thatboth Nb in solution and NbC precipitation exert retard-ing effects on recovery and recrystallisation Retardationof austenite restoration through the presence of Nb Tior V in solution in steels with very low interstitialcontents (0002Cndash0002N wt-) has been demonstratedby Yamamoto et al42 The effectiveness of retardationdecreases in the order listed and is consistent withdecreasing atomic distortion by the solute atom2

Precipitation of fine alloy carbide in deformed austeniteis however much more effective than the solid solutionof the carbide forming element in retarding recovery andrecrystallisation of deformed austenite4041 Howeverthe polymorphic transformation makes it difficult todistinguish fine particles that formed in austenite andthose that formed in ferrite

Indirect evidence of precipitation in deformed auste-nite has been obtained using TEM replicas of samplesof Nb microalloyed steels quenched after austenite

deformation (see for example Ref 42) Similar obser-vations18 for thin foil samples of V microalloyed steelshowed that V4C3 particles were distributed along grainand subgrain boundaries of austenite deformed duringthe finishing rolling pass (Fig 12)

The sites and morphologies of the precipitates andtheir orientation relationships with the matrix can beused to infer particle formation in austenite as shownfor the subgrain network distribution of V4C3 particlesin Fig 12 However quenching was used in this case totry to preserve vestiges of the grain and subgrain struc-ture of the deformed austenite and it would obviouslybe preferable to retain the as deformed austenite atroom temperature This concept was pursued byAbdollazadeh et al4344 who designed an austeniticalloy (Fendash015Cndash307Nindash002Nb) as an analogue of alow carbon Nb microalloyed steel to allow directobservation at room temperature of NbC precipitatesin the retained hot deformed austenite Uniaxial com-pression was used for hot deformation

This work confirmed that both recovery and recrys-tallisation in the Nb steel were retarded relative to a Nbfree reference steel with similar contents of the otherelements Retardation of both recovery and recrystalli-sation was even stronger following strain stimulatedprecipitation of NbC on grain boundaries and thesubgrain network of the deformed austenite Theresearch established that carbide precipitation not onlyserved to maintain the work hardened condition of theaustenite but it also conferred a transient precipitationhardening effect (Fig 13) The orientation relationshipbetween austenite and the cubic carbides nitrides andcarbonitrides of Nb V and Ti is cubendashcube11 Since thelattice parameter of NbC is close to 044 nm and that ofaustenite in an Fendash30Ni alloy is about 036ndash037 nm(ignoring effects of thermal expansion) there is arelatively large atomic mismatch of about 20 alongdirections in the 100 planes Nevertheless age hard-ening does occur as demonstrated by the peak recordedfor the analogue steel

The carbidendashaustenite interface in these cases is likelyto be semicoherent producing a strain field that inhibitsdislocation motion However overaging occurred

12 Dark field electron micrograph using 002VC for V

steel deformed 50 at 840uC and held for 20 min at

840uC before quenching precipitates have formed on

deformation substructure of austenite18 bar repre-

sents 05 mm



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lishe

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rapidly at the relatively high aging temperature used togenerate the data in Fig 13 The precipitation of NbCinferred from the aging peak in Fig 13 is consistent withthe substantial delay evident in Fig 14 in the progressof static recrystallisation at 850uC

Figure 14 compares the times to 50 recrystallisationin the Nb and Nb free steels for two strain levels 025and 050 The delay in recrystallisation below 900uC forthe Nb steel reflects the stabilising effect of NbC pre-cipitation on the deformed austenite structure At highertemperatures recrystallisation is still retarded in the Nbsteel compared with the Nb free steel because of theeffect of solute Nb Figures 13 and 14 are consistentwith hot working results reported for many investiga-tions of austenite deformation and restoration inmicroalloyed steels by laboratory hot rolling torsionand uniaxial and plane strain compression The interac-tion of precipitation recrystallisation and phase trans-formation during the finishing stages of hot deformationcan be rationalised by considering the processes that arepossible over different temperature regimes At highertemperatures precipitation does not exert any effect onrecrystallisation if it occurs after recrystallisation iscomplete However when precipitation occurs on thedeformation substructure it can severely retard theprogress of recrystallisation as shown in Fig 14 A hotrolling investigation of a 0076 Ti microalloyed steel45

also showed that recrystallisation was strongly retardedon rolling below 1030uC Moreover a CndashMn referencealloy also showed retarded recrystallisation on rollingbelow 950uC and it was suggested that AlN precipita-tion on the substructure of the deformed austenite canover an appropriate temperature range exert a similareffect to alloy carbonitride precipitation especially if thestarting austenite grain size is large and solute dragcontributes to a delay in recrystallisation until precipita-tion commences These observations return the discus-sion full circle to the starting analysis of the effect ofAlN precipitation in deformed ferrite on recrystallisa-tion during batch annealing The general principle

emerges that precipitation of fine particles in a hot orcold deformed ferritic or austenitic structure can beexpected to significantly retard subsequent static recrys-tallisation

The age hardening effect shown in Fig 13 is con-sistent with many reports on hot mechanical testing ofaustenite in Nb microalloyed steels which show sub-stantial increases in flow stress with decreasing testtemperature (eg Ref 40) This austenite strengtheningis generally assumed to arise from NbC precipitationthat prevents static recrystallisation In the analoguealloy case the hardening of the austenite on isothermalholding was followed by direct structural analysis usingTEM43 Hardening started on holding at 850uC for12 s For shorter times the austenite in the Nb steelshowed a higher dislocation density and finer cells withmore tangled dislocation cell walls compared with theNb free steel at the same strain The higher hardness ofthe austenite in the Nb steel for times up to y10 s(Fig 13) is a result of the difference in dislocationsubstructure arising from the presence and absence ofNb in solution No precipitation of NbC was detecteduntil holding time exceeded about 20 s but onceprecipitation occurred the dislocation substructurewas strongly stabilised as illustrated by the compara-tive TEM images shown in Fig 15 for a 40 s holdingtime

The time to 5 static recrystallisation at 850uC after025 strain was found to be 10 s for the Nb free steeland 100 s for the Nb steel43 A TEM image of asample of the partially recrystallised Nb steel (450 s at850uC 025 strain) is given in Fig 16 and shows a grainboundary separating a recrystallised grain from thedeformed austenite structure As for recrystallisation offerrite in the FendashVndashC alloy (see Fig 2b) the precipitatesof NbC were significantly coarser after passage of therecrystallising interface as a result of transient pinningof the interface and coarsening by grain boundarydiffusion In both cases fine carbides were presentbefore and exerted a retarding effect on recrystallisa-tion of the deformed matrix Similar recrystallised grainstructures evolved in the two cases with a commonfeature of dispersed coarsened carbide particles withinthe matrix of the recrystallised grains

13 Hardness of austenite in Nb containing and Nb free

Fendash305Nindash015C alloys as function of isothermal hold-

ing time at 850uC after hot compression to strain of

025 (Ref 44)

14 Time to 50 recrystallisation as function of deforma-

tion temperature at strain levels of 025 and 05 for

Nb containing and Nb free austenitic steels43



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ConclusionsPrecipitation during finish rolling of microalloyed steelsstabilises the deformed austenite structure increasingthe rolling load and ensuring that substantial strain isretained to augment the driving force for subsequentpolymorphic transformation to ferrite A high density ofpotential nucleating sites for ferrite is generated by theincrease in SV by deformation and the structuralheterogeneities associated with the excess dislocationsThe result is transformation to ferrite with a smallaverage grain size

Despite pre-precipitation of alloy carbonitride inaustenite its lower solubility in ferrite generally ensuresthat precipitation also occurs in ferrite producing astrength increment by precipitation hardening The alloycomposition and the TMCP route are typically designedto balance the extents of precipitation in austenite andferrite to optimise the contributions of grain refinementand precipitation strengthening to the mechanical pro-perties of the ferrites Nevertheless the research con-ducted at the University of Wollongong has establishedthat the relatively fast cooling rates associated with striprolling normally ensure that the ferrite remains super-saturated with alloy carbonitride because of incompleteprecipitation in ferrite during continuous coolingSubsequent aging treatment has been shown to enableenhanced precipitation strengthening A similar

conclusion applies to the type of Cu bearing precipita-tion hardening microalloyed steel considered in thispaper Heat treatment to produce a martensitic orbainitic structure followed by a suitable aging treat-ment can result in significantly higher strength levelsthan those achieved by the TMCP and aging route Forboth alloy carbonitride and e-Cu precipitation in ferriterelatively small volume fractions of precipitate particlescan exert an unexpectedly profound effect on mechan-ical properties

My intention in writing this contribution was to showhow effectively the science underpinning the develop-ment of microalloyed steels was elucidated by Honey-combe and his research colleagues at Cambridge andhow profoundly it influenced international research inthis area This component of Robert Honeycombersquosresearch output stands both as a legacy to metallurgicalscience and engineering and a tribute to his outstandingabilities as a researcher and research leader

Acknowledgements

I have already expressed my gratitude for the contribu-tion of Robert Honeycombe and I would also like toacknowledge Ross Smith Amir Abdollah-Zadeh andGhasemi Banadkouki who as PhD students under mysupervision generated many of the results referred to inthis review of work undertaken at the University ofWollongong In addition I would like to record myappreciation for the enthusiastic co-operation andsupport of colleagues at BHP Steel (now BlueScopeSteel) over many years In particular I would like tothank Jim Williams Chris Killmore and Frank Barbaro

References1 J H Woodhead and S R Keown lsquo A history of microalloyed

steelsrsquo in lsquoHSLA steels ndash metallurgy and applicationsrsquo Proc Int

Conf HSLA steels rsquo85 (ed J M Gray et al) 15 1986 Beijing

ASM International

2 R W K Honeycombe lsquoFundamental aspects of precipitation in

microalloyed steelsrsquo in lsquoHSLA Steels ndash metallurgy and applica-

tionsrsquo Proc Int Conf HSLA steels rsquo85 (ed J M Gray et al) 243

1986 Beijing ASM International

3 A J DeArdo lsquoNew concepts in the design and processing of high

performance steelsrsquo in Proc HSLA steels rsquo95 (ed G Liu et al)

99ndash112 1995 Beijing China Science and Technology Press

4 T N Baker lsquoMicroalloyed steels ndash an invited reviewrsquo in

lsquoDevelopments in metals and ceramicsrsquo Proc Conf on lsquo70th

birthday meeting of Sir Robert Honeycombersquo (ed J A Charles

ET AL) 76ndash119 1992 London Institute of Materials

5 D P Dunne and R L Dunlea Metall Forum 1978 12 156ndash164

a b

a Nb free steel b Nb containing steel15 Images (TEM) showing deformation substructures in austenite after holding for 40 s at 850uC following 025 strain in

uniaxial compression at same temperature4344 bar represents 2 mm for both images

16 Image (TEM) showing grain boundary separating

recrystallised grain from deformed austenite structure

in FendashNindashNbndashC alloy sample uniaxially compressed to

05 strain at 850uC and held for 450 s at 850uC43 bar

represents 05 mm



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6 E E Underwood lsquoQuantitative Stereologyrsquo 1970 Reading MA

Addison Wesley

7 D Dunne Met Sci 1982 16 259ndash267

8 A D Batte and R W K Honeycombe J Iron Steel Inst 1973

211 284

9 R W K Honeycombe and H K D H Bhadeshia lsquoSteels ndash

microstructure and propertiesrsquo 2nd edn Chap 10 1995 London

Metallurgy and Materials Science

10 A T Davenport F G Berry and R W K Honeycombe Met

Sci J 1968 2 104

11 R W K Honeycombe Metall Trans A 1976 7A 915

12 R A Ricks and P R Howell Met Sci 1982 16 317

13 R A Ricks and P R Howell Acta Met 1983 31 853

14 R A Ricks P A Howell and R W K Honeycombe Metall

Trans A 1979 10A 1049

15 R G Baker and J Nutting ISI special report no 64 1959 1

16 R W K Honeycombe lsquoFerritersquo Hatfield Memorial Lecture 1979

Met Sci 1980 14 201ndash214

17 F G Berry A T Davenport and R W K Honeycombe in lsquoThe

mechanism of phase transformation in crystalline solidsrsquo Institute

of Metals Monograph no 33 328 1969

18 R M Smith lsquoStructural aspects of the hot working of vanadium

and niobium high strength low alloy steelsrsquo PhD thesis University

of Wollongong Wollongong NSW Australia 1987

19 R M Smith and D P Dunne Mater Forum 1988 11166ndash181

20 R M Smith D P Dunne and J G Williams Met Forum 1982 5

(2) 109

21 D P Dunne R M Smith and T Chandra Proc Thermec rsquo88

Tokyo Japan June 1988 Iron Steel Institute 275

22 R M Smith D P Dunne and T Chandra in lsquoHigh strength low

alloy steelsrsquo (ed D P Dunne and T Chandra) 188 1984 Port

Kembla NSW South Coast Printers

23 R K Amin and F B Pickering in lsquoThermomechanical processing

of microalloyed austenitersquo (ed A J de Ardo et al) 377 1982

Pittsburgh PA AIME

24 W Roberts Scand J Metall 1980 9 13

25 W B Morrison J Iron Steel Inst 1963 201 317

26 J M Gray and R B G Yeo Trans ASM 1968 61 225

27 Great Lakes Steel Corp Mech Eng 1959 81 53

28 C A Beiser ASM preprint no 138 1959

29 W C Leslie NPL Conf Proc 334 1963 London HMSO

30 W B Morrison and J H Woodhead J Iron Steel Inst 1963 201

43

31 E R Morgan T E Dancy and M Korchynsky J Met 1965

829

32 R Phillips and J A Chapman J Iron Steel Inst 1966 204

615

33 F B Pickering lsquoPhysical metallurgy and the design of steelsrsquo 1978

Essex Applied Science Publishers Ltd

34 S S Ghasemi Banadkouki lsquoTransformation characteristics and

structure-property relationships for a copper-bearing HSLA steelrsquo

PhD thesis University of Wollongong Wollongong NSW

Australia 1996

35 D P Dunne S S Ghasemi Banadkouki and D Yu ISIJ Int

1996 36 324

36 S S Ghasemi Banadkouki D Yu and D P Dunne ISIJ Int

1996 36 61

37 C R Killmore J F Barrett A Cabello I F Squires and J G

Williams Proc Conf on lsquoOffshore mechanics and arctic engineer-

ingrsquo The Hague The Netherlands March 1989 ASME

38 DP Dunne lsquoWeldable copper strengthened low carbon steelsrsquo

Proc HSLA steels rsquo95 (ed G Liu H Stuart et al) 99ndash112 1995

Beijing China Science and Technology Press

39 Y Tomita T Haze N Saito T Tsuzuki Y Tokunaga and

K Okamoto ISIJ Int 1994 34 836

40 C M Sellars lsquoThe influence of particles on recrystallisation during

thermomechanical processingrsquo Proc 1st RISO Conference on

Metallurgy and Materials Science Recrystall ndash 1 section and Grain

Growth in Multiphase and Particle Containg Materials (ed by

N Hansen A R Jones and T Leffars) Raskilde Denmark 1980

291

41 I Weiss and J J Jonas Metall Trans A 1979 10A 831

42 S Yamamoto C Ouchi and T Otsuka in lsquoThermomechanical

processing of microalloyed austenitersquo (ed A J de Ardo et al) 613

1982 Pittsburgh AIME

43 A Abdollah-Zadeh lsquoInvestigation of deformation recrystallisa-

tion and precipitation in austenitic HSLA steel analogue alloysrsquo


Australia 1996

44 A Abdollah-Zadeh and D P Dunne lsquoRecovery and recrystallisa-

tion in steelrsquo Proc 37th MWSP Conf Hamilton Ont Canada

1996 ISS-AIME 74

45 D P Dunne T Chandra and S Misra lsquoHSLA steels ndash metallurgy

and applicationsrsquo in Proc Conf Microalloying rsquo85 (ed J M Gray

et al) 207 1985 Beijing ASM



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branching out into research fields that were potentiallymore funding-friendly than martensite crystallography

A significant part of the research conducted in theAlloy Steels Group was focused on clarifying funda-mental aspects of alloy carbide and nitride precipitationin laboratory prepared ternary alloys with the ultimateaim of fuller understanding of the influence of structureon the mechanical properties of commercial microal-loyed steels The project that I proposed to RobertHoneycombe fitted with this research emphasis andconsisted of an experimental study of the restorationbehaviour of cold rolled FendashVndashC alloy that had beenheat treated before cold rolling to produce IP of V4C3

Precipitation and recrystallisationI had previously published a paper on lsquopancakersquo grainformation in batch annealed Al killed deep drawingsteels5 The development of grain shape anisotropy in anannealed steel fascinated me and I set out to map theevolution in ferrite grain shape through the cold rollingand recrystallisation processes used for this type of steelWhen lsquofinished hot and coiled coldrsquo these steels aresupersaturated with AlN which subsequently precipi-tates out during batch annealing and interacts with therecrystallisation process Recrystallisation is severelyimpeded and is nucleation limited resulting in largerecrystallised grains with a preferred orientation that isfavourable for sheet forming by deep drawing The workthat I carried out showed that the description of therecrystallised grain shape as lsquopancakedrsquo is incorrect assome linear anisotropy is retained effectively mimick-ing albeit to a reduced extent the slab shape presentafter cold rolling Figure 1 shows that the recrystallisedgrains have similar planarndashlinear anisotropy to that ofthe cold rolled ferrite grains Cross-rolling instead ofunidirectional rolling followed by simulated batchannealing resulted in no linear anisotropy and the grainshape was truly lsquopancakedrsquo Recrystallised grain shape

anisotropy arises because the AlN precipitates on theboundaries of the subgrains and grains of the unidir-ectionally cold rolled steel pinning these boundaries andrestricting the lsquonucleationrsquo of recrystallised domains bygrain boundary bulging or subgrain growthcoalescenceOnce nucleated the growth of recrystallised grainsproceeds within a forest of impeding particles distrib-uted in planarndashlinear layers in the deformed matrixTherefore growth is anisotropic

My interest in and curiosity about the interactionbetween precipitates and restoration (andor poly-morphic transformation) led naturally to the projectthat I undertook at Cambridge ferrite restoration ofcold rolled ferrite in a Fendash105Vndash023C (wt-) alloycontaining networks of fine coherent interphase pre-cipitate particles that were present before cold rolling7

The heat treatment used to generate the startingstructure was annealing under argon at 1160uC for15 min followed by isothermal transformation in a saltbath for 15 min at 747uC According to Batte andHoneycombe8 this treatment results in a volumefraction of 00123 of V4C3 precipitates of 3ndash5 nm radiusCold rolling of 75 mm strip was used to reduce thethickness by 01 mm per pass and samples were takenafter each 10 reduction up to 60 in order toinvestigate recrystallisation kinetics at 706uC

Cold rolling resulted in a dense and uniform distribu-tion of dislocations as well as the development of shearbands for reductions greater than 20 The shear bandswere typically 02ndash4 mm thick and oriented at an angleof about iexcl30u to the rolling direction The shear bandsubstructure consisted of elongated subgrains 01 mmthick and provided potent sites for nucleation ofrecrystallisation Although recovery and recrystallisa-tion occurred rapidly within the bands on annealing of60 cold rolled samples at 706uC restoration ofsurrounding material was severely retarded and recrys-tallisation was incomplete after 1000 h The rate

1 Linear VLIN and planar VPL anisotropy factors for ferrite grains after a cold rolling and b cold rolling and batch anneal-

ing as a function of percent cold reduction Following Underwood6 VPL is defined as the ratio SV(PL)SV(TOT)

VLIN5SV(LIN)SV(TOT) and VPLndashLIN5[SV(PL)zSV(LIN))SV(TOT)] where SV is grain boundary surface area per unit volume and

SV(TOT) includes the planar linear and isometric values which are determined from grain boundary intercept densities

in the rolling direction the transverse direction and the direction of the rolling plane normal56



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a b












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(c)

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Pub

lishe

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ey P

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ey P

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ions

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Pub

lishe

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ey P

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(c)

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sents 05 mm



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ey P

ublis

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(c)

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mun

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ions

Ltd










025 (Ref 44)






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ey P

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mun

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ions

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Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

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mun

icat

ions

Ltd









a b












Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

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mun

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ions

Ltd














Pub

lishe

d by

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ey P

ublis

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(c)

IOM

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mun

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ions

Ltd



















Pub

lishe

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ey P

ublis

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(c)

IOM

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mun

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ions

Ltd
















Pub

lishe

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ey P

ublis

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(c)

IOM

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mun

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ions

Ltd














Pub

lishe

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ey P

ublis

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(c)

IOM

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mun

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ions

Ltd














sents 05 mm



Pub

lishe

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Man

ey P

ublis

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(c)

IOM

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mun

icat

ions

Ltd










025 (Ref 44)






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lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

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mun

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ions

Ltd



















Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

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mun

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ions

Ltd
















Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

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mun

icat

ions

Ltd














Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

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mun

icat

ions

Ltd














sents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd










025 (Ref 44)






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd



















Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd
















Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














sents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd










025 (Ref 44)






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd
















Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














sents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd










025 (Ref 44)






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














sents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd










025 (Ref 44)






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd














sents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd










025 (Ref 44)






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd










025 (Ref 44)






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd





Acknowledgements





ASM International













a b







represents 05 mm



Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Pub

lishe

d by

Man

ey P

ublis

hing

(c)

IOM

Com

mun

icat

ions

Ltd


Addison Wesley



211 284





Sci J 1968 2 104





Trans A 1979 10A 1049












(2) 109








Pittsburgh PA AIME








43


829


615






Australia 1996


1996 36 324


1996 36 61














291








Australia 1996



1996 ISS-AIME 74






Interaction of Precipitation With Recrystallisation and Phase Transformation In

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Transcript of Interaction of Precipitation With Recrystallisation and Phase Transformation In