Discontinuous cellular precipitation in a Cr–Mn–N steel with niobium and vanadium additions

Discontinuous cellular precipitation in a Cr–Mn–N steelwith niobium and vanadium additionsR.D.Knutsen *,ng,J.A.Basson1Department of Mechanical Engineering,Centre for Materials Engineering,University of Cape Town,Private Bag Rondebosch,7701Cape Town,South Africa Received 6November 2003;received in revised form 22January 2004;accepted 26January 2004AbstractThe influence of niobium and vanadium additions on the precipitation behaviour in a 24Cr–18Mn–1N austenitic stainless steel is characterised during ageing at temperatures from 800to 1100°C.Niobium demonstrates a tendency to stabilise the cubic MX-type precipitates,whereas vanadium encourages the formation of hexagonal close-packed M 2X-type precipitates.Vanadium,further-more,promotes formation of M 2X-type precipitates by the discontinuous cellular precipitation (DCP)reaction.The presence of sigma phase and a high frequency of austenite twinning occur in association with the DCP reaction during ageing.This behaviour,together with the influence of niobium and vanadium,is used to understand the driving force for boundary migration during the DCP reaction.Ó2004Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Stainless steels;Ageing;Interface migration;Discontinuous cellular precipitation1.IntroductionHigh nitrogen austenitic stainless steels provide im-proved mechanical properties and corrosion resistance compared to conventional austenitic grades [1].As a result,these high nitrogen steels have found broader application in engineering,particularly where advantage is taken of their excellent combination of high strength and toughness.Modern developments in steel making have made possible the manufacture of steels with nitrogen levels up to 1wt%or even more [1,2].The processes employed to achieve these high nitrogen levels can vary:solid-state nitriding processes are easily adapted to thin sheet and wire stock,whereas thicker sections are limited to liquid state metallurgy involving either high or low (atmospheric)pressure melting prac-tices.For example,the manufacture of high nitrogen steels with up to 1wt%nitrogen is made possible at lowpressure by manipulating the solubility of the nitrogen in the melt with appropriate alloying.Central to this approach is the effective use of manganese in controlling the melt chemistry [3].Besides influencing the nitrogen levels in the melt,the manganese additions reinforce the high work hardening ability associated with nitrogen in solution,and consequently promote even greater strength and toughness.In addition to the structural advantage provided by good mechanical properties,the combination of high strength and toughness also pro-motes excellent wear resistance in these steels,which when coupled with good corrosion resistance,provides attractive tribological properties [4].Notwithstanding the good wear resistance provided by the high nitrogen austenitic steels,attempts have been made to further improve the wear resistance by promoting nitride dispersions in the austenite matrix.To this end the addition of strong nitride formers,niobium and vanadium,to casting alloys has been considered with a view to promoting desirable precipitate disper-sions under controlled ageing conditions.However,second phase particles can also be quite damaging to mechanical and corrosion properties [5].For example,*Corresponding author.Tel.:+27-216503172;fax:+27-216897571.E-mail address:rknutsen@eng.uct.ac.za (R.D.Knutsen).1Present address:180Degrees Engineering Solutions,Cape Town,South Africa.1359-6454/$30.00Ó2004Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.actamat.2004.01.031Acta Materialia 52(2004)2407–2417chromium nitrides(Cr2N)form discontinuously in high nitrogen stainless steels on ageing in the temperature range700–1000°C[6–8]and their cellular morphology significantly reduces toughness and corrosion resistance. During discontinuous cellular precipitation(DCP), Cr2N lamellae form behind a migrating boundary leaving precipitate cells consisting of alternate layers of Cr2N precipitate and less nitrogen-saturated austenite [6,9].The Cr2N precipitate is of the M2X-type,where M is usually Cr,but can be a mixture of other compatible elements.X usually represents nitrogen,but can also be a mixture of carbon and nitrogen[10].The addition of niobium and vanadium,on the other hand,is expected to result in the formation of MX precipitates that form in the temperature range1000–1200°C[10–12].This study reports on the influence of niobium and vanadium additions in the range0–1wt%on the pre-cipitation reactions in high nitrogen steel with an ap-proximate base composition of24Cr–18Mn–1N(all figures in wt%).The alloy compositions were chosen so that the individual and combined effects of the element additions could be determined.2.Experimental procedure2.1.MaterialsEight cast ingots of similar base composition,each weighing approximately20kg,were investigated.The ingots differed from each other with regard to niobium/ vanadium balance as shown in Table1.The base com-position is represented by alloy A,which contains low levels of niobium and vanadium.Alloys B and C contain approximately1wt%niobium and vanadium,respec-tively,whereas alloys D–H are composed to consider the influence of different Nb/V ratios on the microstructural evolution.2.2.Heat treatments and metallographyIn order systematically to study the precipitation behaviour of the eight alloys,the as cast alloys werefirst solution treated at1300°C for2h and later aged at temperatures ranging from800to1100°C for periods of 0.5,2,5,10and100h.All heat treatments were per-formed under argon partial pressure conditions(102 Torr)and specimens were water quenched after each heat treatment,including the initial solution treatment. Standard metallographic techniques were used to pre-pare specimens for microscopic investigation.Thefinal preparation stages for light microscopy included electro-polishing in a solution containing25g chromium trioxide,133ml acetic acid,and7ml distilled water, followed by electro-etching in the same solution.Both electrolytic procedures were performed at room tem-perature;electro-polishing at20V and electro-etching at 10V.The electro-etching stage was omitted in situations where orientation images were generated using electron backscattered diffraction in the scanning electron mi-croscope(SEM).The twin-jet polishing technique was utilised to prepare specimens for transmission electron microscope(TEM)studies.Suitable foils were prepared using a solution containing10%perchloric acid and20% glycerol in ethanol.The foils were electro-thinned by applying a potential difference of15V at temperatures ranging from)10to0°C.Semi-quantitative composi-tional analysis of the precipitate phases was performed using energy dispersive X-ray spectroscopy in the SEM. In most cases,elemental compositions were deduced from high resolution X-ray maps.The extent of DCP is reported by reference to the volume fraction of the precipitate cells as measured by areal analysis performed on a series of light micrographs.3.Results3.1.Solution treated conditionThe group of alloys showed varying levels of minor phase content in an austenite grain structure after the standard solution treatment at1300°C for2h.These secondary phases,also referred to as particles,were generally absent in the very low Nb-containing alloys(A and C),but increased in volume fraction with increasing Nb/V ratio in the remaining alloys.Although a rea-sonable level of homogenisation was expected after theTable1Composition of Alloys A–HElement wt%(at%)A B C D E F G HFe Balance Balance Balance Balance Balance Balance Balance Balance Cr24.124.423.523.123.924.723.023.7Mn17.517.518.318.818.818.318.418.4N 1.03(3.89) 1.02(3.87) 1.17(4.39)0.98(3.70) 1.02(3.86) 1.15(4.32) 1.13(4.26) 1.13(4.26) C0.08(0.35)0.11(0.49)0.13(0.57)0.23(1.01)0.11(0.49)0.14(0.61)0.12(0.53)0.14(0.62) Nb0.04(0.02) 1.17(0.67)0.04(0.02)0.92(0.52)0.62(0.35)0.29(0.16)0.29(0.16)0.40(0.23) V0.20(0.21)0.11(0.11) 1.18(1.21)0.94(0.97)0.45(0.47)0.75(0.78)0.47(0.49)0.22(0.23) Allfigures in weight percent except where atomic percent appears in parentheses.2408R.D.Knutsen et al./Acta Materialia52(2004)2407–2417solution heat treatment,the distribution of the particles highlighted the interdendritic regions of the as cast structure,as shown by the distribution pattern in alloy E in positional analysis has demonstrated that these particles are rich in niobium and vanadium when both elements are present in the alloy,but only niobium is present in the particles in the very low V-containing alloy (alloy B).Electron diffraction analysis in the TEM shows that the particles are face-centred cubic with lat-tice parameters in the range 4.24–4.38 A,which is con-sistent with the mixed (Nb,V)N phase.Accordingly,the interdendritic phases are identified as MX-type particles,where the M component can be a combination of nio-bium and vanadium,with increasing dominance in niobium content as the niobium level in the alloy increases.3.2.Ageing treatments at 1100°CWith the exception of the base alloy (alloy A),ageing at 1100°C resulted in either the formation of lamellar precipitates or fine globular precipitates,depending on the Nb/V ratio in the steel.No additional second phases were observed in alloy A after ageing at 1100°C.Fig.2shows the lamellar-type precipitates that were observed in alloys C,D,F and G after ageing for 2h.These precipitates formed by DCP and indications of the mi-grating front normally associated with DCP are marked by arrows in both micrographs.The volume fraction oflamellar precipitates is greatest in alloy C,where the total cell volume per grain ranges from 0.2to 0.3.The cell volume fraction for alloys D,F and G is estimated at 0.08,0.15and 0.05,respectively.Examination of the same precipitate distributions in the TEM reveals fur-ther distinction in that the precipitate lamellae are often fragmented in alloys F and G compared to the more continuous nature of the individual lamellae in alloys C and D.Electron diffraction analysis has shown that the lamellae are hexagonal close-packed and that they de-velop a close packed orientation relationship with the host austenite phase such that (111)c //(0001)M 2X and [)110]c //[1)210]M 2X.Fig.3shows the [011]diffrac-tion pattern for precipitate lamellae formed in alloy C.Although the diffraction analysis is consistent with the identification of Cr 2N type precipitates,compositional analysis by EDS has shown that the lamellae are rich in both chromium and vanadium.Furthermore,the lattice parameters for the lamellar precipitates vary accordingto a ¼2:57–2.84 Aand c ¼4:57–4.61 A.Given this situation,it is reasonable to conclude that the lamellae are of the mixed M 2X-type precipitate of the form (Cr,V)2N.After extended ageing (up to 100h)at 1100°C,the DCP process ceases,and instead the same phase forms continuously as fine needle-like precipitates (Fig.4).Discontinuous cellular precipitation is notably absent in alloys B,E and H,but instead fine globular precipi-tates form at grain boundaries and within the large austenite grains.This precipitate form is depicted in a TEM micrograph in Fig.5after ageing alloy E at 1100°C for 2h.The [112]zone axis pattern onceagainFig.1.Occurrence of MX-type particles in alloy E after solution treatment at 1300°C.mellar precipitates after ageing at 1100°C for 2h,(a)alloy C and (b)alloyF.Fig.3.[011]Zone axis diffraction pattern for austenite matrix.Extra spots correspond to the [1)210]zone axis for the HCP system.R.D.Knutsen et al./Acta Materialia 52(2004)2407–24172409confirms the MX-type precipitate and the calculatedlattice parameter (4.2–4.4 A)is consistent with the pre-cipitate form that was already present after the solution treatment process.DCP did not occur in alloys B,E and H even after extended ageing up to 100h at 1100°C.3.3.Ageing treatments at 1000°CDiscontinuous cellular precipitation occurred in all but alloy B after ageing at 1000°C for up to 2h.The niobium and vanadium levels influenced the extent of lamellar precipitation and the greatest volume fraction of precipitate cells was observed in alloy C,which has maximum vanadium level and lowest Nb/V ratio.The cell volume fraction per grain is nearly 1.0.Alloy A had the next greatest lamellar precipitate volume (cell volume ¼0.15),with the remaining alloys D–H having much lower precipitate volumes (0.05–0.10).The con-trast in lamellar precipitate distribution is exhibited in Fig.6where alloys C and F are compared after ageing at 1000°C for 2h.The entire austenite grain structure is transformed to lamellar precipitate cells in alloy C (Fig.6(a)).After extended ageing at 1000°C for up to 100h,needle-like precipitates formed in the remaining untransformed austenite matrix in alloys D–H,as shown in Fig.7(a).The DCP reaction ceased inalloysmellar precipitation distributions after ageing at 1000°C for 2h,(a)alloy C and (b)alloyF.Fig.7.Extended ageing for 100h at 1000°C,(a)continuous precipitation in alloy F and (b)discontinuous coarsening in alloyC.Fig.5.Globular MX-type precipitates in alloy E after ageing at 1100°C.Inset:[112]zone axis pattern forprecipitate.Fig.4.Continuous precipitation in alloy C after extended ageing for 100h at 1100°C.2410R.D.Knutsen et al./Acta Materialia 52(2004)2407–2417D,F and G after 5h,but continued for longer periods in alloys E and H,although it is obvious that extensive plate-like precipitates occurred during ageing up to 100h.Diffraction analysis of the needle-like precipitates yielded the same result as the lamellar phase.As an-ticipated,both precipitate forms were found to contain chromium and vanadium.On the other hand,needle-like precipitates did not form in alloys A and C after prolonged ageing,but instead the DCP reaction was regenerated by the formation of new cells on old cell boundaries,as shown in Fig.7(b).The new pre-cipitate cells have a much greater interlamellar spacing than the Ôold Õprecipitate cells and the reaction has been described as discontinuous coarsening [13].A highincidence of austenite twinning was also detected in the DCP cells in all the alloys and an example after ageing at 1000°C for 10h is shown for alloy A in Fig.8.The high Nb-containing alloy B appeared much the same after ageing at 1000°C compared to the ageing treatments at 1100°C,even for periods of up to 100h at temperature.DCP was notably absent and instead TEM studies identified prolific globular precipitate popula-tions on the austenite grain boundaries.An example of this precipitate form after ageing for 2h is exhibited in Fig.9.Diffraction and compositional analysis indicates that these precipitates are the same MX-type precipi-tates that were identified at 1100°C,although not sur-prisingly,they are much smaller in size than those depicted in Fig.5.3.4.Ageing treatments at 800°CAgeing treatments at 800°C resulted in the formation of lamellar precipitates in all of the alloys including alloy B.The volume fraction of lamellar precipitation after ageing for 2h was greatest in alloy C,followed by alloy A and then the alloys with significant niobium addi-tions.The volume fraction of lamellar precipitation in alloy B was the lowest after the same ageing time.Fig.10(a)and (b)show the lamellar precipitation in alloys A and E,respectively.Alloy E is representative of the microstructures that evolved in Nb-containing al-loys.Of particular importance is the observation that the interlamellar spacing of the precipitates within the precipitate cells is much finer than those that formed after ageing at 1000°C and that extensive austenite twinning continued to occur in the cell structures.Al-though twinning was also observed in the DCP cell structures at the higher temperatures (Fig.8),the fre-quency of twinning is much greater at 800°C.Diffrac-tion analysis in the TEM once again showed that these are M 2X-type precipitates with the same orientation relationship with the austenite as those that form at 1000°C.The host austenite in alloys A and CwereFig.8.Parallel twin boundaries in alloy A after ageing at 1000°C for 10h.Fig.9.Globular MX-type precipitation in alloy B after ageing at 1000°C for 2h.Fig.10.Distribution of DCP volumes after ageing at 800°C for 2h,(a)alloy A and (b)alloy E.R.D.Knutsen et al./Acta Materialia 52(2004)2407–24172411completely transformed to DCP cells after ageing at 800°C for 100h.Sigma phase (r -FeCr)also appears after ageing at 800°C and is shown in bright contrast between the DCP cells in alloy A in Fig.11.The body-centred tetragonal r -phase is known to occur in the Fe–Cr system at temperatures below 820°C,although the reaction is quite sluggish [14].Nevertheless,r -phase can dissolve large amounts of other elements,including vanadium and manganese [14–16],and it has been shown that manganese in particular increases the sta-bility of the r -phase [15,16].The high manganese levels,as well as the additions of vanadium,have clearly assisted in the formation of r -phase in these alloys.4.DiscussionThe precipitation behaviour in the present range of alloys during ageing between 800and 1100°C has been shown to be quite variable.Two predominant precipi-tate types occur,namely MX and M 2X.The substitu-tional element composition of these precipitates is complex,although niobium dominates the MX-type,whereas chromium dominates the M 2X-type.Notwith-standing the variation in total alloy concentration over the range of alloys,the influence of niobium and vana-dium on the morphology and mechanisms of precipita-tion over this temperature range is significant and distinctions can be made even in the solution treated cast structures.Although it is unfortunate that the ele-ment additions,excluding niobium and vanadium,are not highly consistent,there is sufficient excess interstitial element to encourage the formation of the respective precipitates,and thus greater influence on the progress of precipitation is partitioned to the substitutional nio-bium and vanadium additions.The sequence of precip-itation events will be explained in terms of the influence of niobium and vanadium on precipitate stability as well as the influence that these elements have on controlling the progress of DCP.4.1.Influence of niobium and vanadium during solution treatmentThe role of niobium in promoting nitride formation in the interdendritic regions in the cast structure is es-tablished by the abundance of MX-type particles that are still visible after the solution treatment at 1300°C,which is a mere 50–70°C below the solidus temperature in these alloys.Although vanadium also occurs in solid solution in these particles,there is no evidence to suggest that vanadium promotes their formation.The formation of large,blocky particles of the MX-type and face-cen-tred cubic structure has been reported in mild steels [17],high strength low alloy steels [18],stainless steels [19]as well as high nitrogen steels [20].Further,it has been shown that niobium additions strongly favour this re-action,especially in the presence of nitrogen [17,21,22].Clearly the level of 1wt%niobium is more than suffi-cient to promote the strong presence of Nb-rich nitrides in the interdendritic liquid during solidification,whereas these particles are absent in the two low Nb-containing alloys (alloys A and C).The formation of the stable MX-type particles during solidification will reduce the nitrogen level in solution at 1300°C,although the im-pact is quite small.Despite the fact that the volume fractions of the MX-type particles have not been quantified,even complete consumption of the 1wt%Nb in alloy B in the form of NbN particles will still leave 0.89wt%nitrogen in solution.Since it is also likely that some niobium will still remain in solution,the level of nitrogen in solution in the austenite phase after the solution treatment at 1300°C will be above at least 0.9wt%.4.2.Discontinuous cellular precipitation in the base alloy In an attempt to explain the respective roles of the niobium and vanadium additions in affecting precipita-tion behaviour at the different ageing temperatures,it is appropriate to first discuss the precipitation behaviour in the base alloy (alloy A).No precipitation was ob-served at 1100°C;precipitation only occurred in earnest during the 1000°C ageing treatment.At this tempera-ture profuse DCP occurred to produce Cr-rich nitride (M 2X)lamellae behind migrating austenite grain boundaries.Similar precipitate reactions have been re-ported by Vanderschaeve et al.[7]and Kikuchi et al.[6]in high nitrogen stainless steels.The DCP reaction continued until the entire austenite matrix was trans-formed to the cellular structure,and upon continued ageing after the saturation point,the DCP reaction was regenerated and discontinuous coarsening occurred.At 800°C,the same DCP reaction occurs,although the interlamellar spacing is much finer.In addition to the precipitation of MX 2lamellae at 800°C,r -phase also formed at the DCP cell boundaries.The formationofFig.11.Occurrence of r -phase at DCP cell boundaries (alloy A aged at 800°C).2412R.D.Knutsen et al./Acta Materialia 52(2004)2407–2417r-phase is not unexpected and has been predicted by thermodynamic analysis to occur in these steels[23]. However,the distribution of r-phase in relation to the DCP cells is interesting(Fig.11)and suggests that a point is reached during the migration of the cell boundaries when the chromium content becomes suffi-ciently enriched at the migrating interface to allow r-phase to form.After a certain volume of r-phase has formed,the DCP process is again initiated and the boundary migration continues.Therefore,the sequence of events portrayed by the morphological arrangement of the DCP cells and r-phase indicates that chromium becomes enriched at the migrating grain boundary,in-stead of being depleted if one places emphasis on the growth of the Cr-rich lamellae.This deduction is at odds with the conclusions reached by Vanderschaeve et al.[7] who argue that the DCP reaction is prematurely termi-nated in their steels due to the development of chro-mium impoverished zones ahead of the migrating interface.Their support for the incomplete reaction is based on the fact that the untransformed austenite ahead of the migrating DCP cell still contains sufficient nitrogen(about75%of the original level)to cause su-persaturation in the host,and therefore drive the for-mation of Cr2N.Although they have elegantly illustrated that the DCP reaction in high nitrogen au-stenitic stainless steels does not occur under steady state conditions,they have based their argument on the comparison of the measured growth rate of the precip-itate cells with the theoretical growth rate of lamellar structures predicted by grain boundary diffusion and volume diffusion of chromium,respectively.They con-cluded that,whilst the chemical driving force for the lamellar precipitation is governed by the level of nitro-gen supersaturation,the rate of cell advancement is controlled by the supply of chromium to the growing lamellae.In the early stages of the DCP reaction they report that the growth rate is consistent with grain boundary diffusion of chromium,but that the growth rate steadily retards thereafter towards the slower bulk diffusion rate and then prematurely terminates as men-tioned above.In our deliberations,we wish to place different emphasis on the issues affecting the non-steady state growth conditions of the DCP cells,particularly with regard to the migration of the austenite/austenite grain boundary.If the formation of the lamellar precipitates is ignored for the moment,then it is possible to develop an argu-ment for the migration of the austenite/austenite grain boundaries based on the existence of a pulling force for the migrating boundary.To establish this,it isfirst necessary to consider the widely accepted conditions for normal grain growth in a single phase polycrystalline solid.In the case of two adjacent grains separated by a curved boundary,the effect of the pressure difference caused by the curved boundary is to create a difference in free energy(D G)that drives the atoms across the boundary,and the boundary migrates towards the centre of curvature.Consequently,the free energy dif-ference is thought of as a force pulling the grain boundary towards the grain with the higher energy.This pulling force per unit area of boundary is given byF¼D G=V m;ð1Þwhere V m is the molar volume.The above situation not only applies to grain growth,but also to recrystallisation for example,where the boundaries between the new strain-free grains and the original deformed grains are acted on by a force D G=V m,where in this case,D G is due to the difference in dislocation strain energy between the two grains.In the present case,it is proposed that the energy difference is provided by the difference in nitro-gen saturation across the moving interface,and hence D G can be referred to as the chemical potential differ-ence,D l.Accordingly,the transformed austenite grain will migrate into the untransformed,and still supersat-urated,neighbouring grain.However,just as the mi-gration of a recrystallisation front is made possible by the annihilation of dislocations as the boundary sweeps into the deformed grain,so is the present situation made possible by the consumption of nitrogen by the ad-vancement of the lamellae behind the migrating grain boundary.The similarity to the process of recrystalli-sation is further suggested by grain orientation mea-surements depicted in Fig.12.In this case(alloy A), precipitate cells have formed at either side of the original grain boundary separating the austenite grains1and2 (Fig.12(a)).The grey-scale orientation image for aus-tenite in Fig.12(b)demonstrates that the cell volume(1) migrating into grain2has the same orientation as grain 1,whereas the cell volume(2)migrating into grain1has the same orientation as grain2.The same orientation reversal can arise during recrystallisation when large subgrains on one side of a grain boundary can grow into the adjacent grain and vice versa,and the process is often referred to as strain induced grainboundary Fig.12.Parent austenite orientation reversal at grain boundary due to DCP,(a)cells migrating either side of austenite grain boundary and (b)grey-scale austenite orientation map of same area as(a).R.D.Knutsen et al./Acta Materialia52(2004)2407–24172413migration.In this way it is proposed that the boundary only indirectly migrates to seek solute for the precipitate growth,and that the principal force behind boundary migration is the difference in chemical potential associ-ated with the transformed and untransformed austenite.In order for the process of DCP to progress at some given rate,the mobility of the grain boundaries must be considered as well as the driving force.According to the discussion in the preceding paragraph,the rate of DCP is influenced by the rate of austenite/austenite boundary migration and hence it is common to write the progress, or velocity,asV¼M D l=V m;ð2Þwhere M is the mobility of the migrating grain bound-ary,and V m in this case is the molar volume of the transformed austenite phase.The mobility of the grain boundary will be determined by the mechanism of boundary migration and whether the migration is dif-fusion or interface controlled,as well as the influence ofsolute drag that is expected to occur in such a complex alloy.Returning to the precipitation of the M2X-type la-mellae,it is widely observed that the precursor to the DCP reaction is that precipitatesfirst form at a grain boundary and develop an orientation relationship with one of the grains(the host grain)to reduce the interfa-cial energy[9].After this initial stage,the host grain migrates into the neighbouring grain and the lamellar precipitates advance in the direction of the migrating grain boundary.But,if the grain boundary did not migrate,the process of precipitation would still continue and would occur by growth of the lamellae into the in-terior of the host grain to maintain the favourable ori-entation relationship;in other words,opposite to the direction in which the grain boundary would normally migrate.The fact that the grain boundary does migrate provides a much faster supply of solute for precipitate growth and,consequently,the precipitate growth fol-lows closely behind the migrating grain boundary.This is the preferred growth direction for the lamellae since it may be argued that the solute is provided by grain boundary diffusion along the advancing grain boundary rather than by volume diffusion.In fact,the precipita-tion of r-phase at the cell boundaries suggests that there is an oversupply of at least chromium at the migrating grain boundary.To summarise,the fact that the DCP reaction occurs at all is a result of the migration of the austenite/aus-tenite grain boundary and not that the boundary mi-grates due to the DCP reaction.The initial driving force for the grain boundary migration is provided by the chemical potential difference(D l)between the nitrogen depleted volume around the newly formed precipitates and the nitrogen supersaturated volume in the adjacent grain,which is shown schematically in Fig.13.As mentioned previously,this is analogous to the process of recrystallisation where the high angle boundary migrates in the direction of higher stored energy.In the case of recrystallisation,the dislocations are annihilated at the sweeping grain boundary and consequently the energy difference is maintained across the boundary.In the present situation,the advancement of the grain bound-ary into the supersaturated grain collects nitrogen at the boundary,which is then consumed by the growing M2X lamellae,and thus D l is more or less maintained be-tween the transformed austenite volume and the un-transformed austenite grain.It is understandable that the M2X lamellae will follow close behind the migrating boundary and the lamellae are often observed to be at-tached to the boundary.However,this is not necessarily always the case and an example is presented in Fig.14 where the migrating boundary has moved ahead of the precipitate lamellae.In the same way that a boundary would not be able to migrate during recrystallisation if the dislocations did not dissolve at the migrating boundary,it is imperative that the solute is removed by the advancing lamellae.Furthermore,it follows that theFig.14.Austenite/austenite boundary migrating ahead of lamellar precipitates.2414R.D.Knutsen et al./Acta Materialia52(2004)2407–2417。