The temporal evolution of the decomposition of a concentrated multicomponent Fe Cu-based steel

Transcription

1 Avalable onlne at Acta Materala 56 (2008) The temporal evoluton of the decomposton of a concentrated multcomponent Fe Cu-based steel R. Prakash Koll *, Davd N. Sedman Department of Materals Scence and Engneerng, Northwestern Unversty, Northwestern Unversty Center for Atom-Probe Tomography (NUCAPT), Evanston, IL 60208, USA Receved 7 August 2007; receved n revsed form 28 December 2007; accepted 31 December 2007 Avalable onlne 4 March 2008 Abstract The nucleaton (to a lmted extent), growth and coarsenng behavor of Cu-rch precptates n a concentrated multcomponent Fe Cu-based steel aged at 500 C from 0.25 to 1024 h s nvestgated. The temporal evoluton of the precptates, heterophase nterfaces, matrx compostons and precptate morphologes are presented. Wth ncreasng tme, Cu parttons to the precptates, N, Al and Mn partton to the nterfacal regon, whereas Fe and S partton to the matrx. Coarsenng tme exponents are determned for the mean radus, hr(t), number densty, N V (t), and supersaturatons, whch are compared to the Lfshtz Slyzov Wagner (LSW) model for coarsenng, modfed for concentrated multcomponent alloys by Umantsev and Olson (UO). The expermental results ndcate that the alloy does not strctly follow UO model behavor. Addtonally, we delneate the formaton of a N Al Mn shell wth a stochometrc rato of 0.51:0.41:0.08 at 1024 h, whch reduces the nterfacal free energy between the precptates and the matrx. Ó 2008 Acta Materala Inc. Publshed by Elsever Ltd. All rghts reserved. Keywords: Coarsenng; Fe Cu alloy; Phase separaton; Atom-probe tomography; Interfacal segregaton 1. Introducton Copper precptaton-strengthened structural steels are of consderable commercal mportance due to ther hgh strength, good mpact toughness, excellent weldablty wthout preheat or postheat, and corroson resstance [1 4]. Ths desrable combnaton of propertes s derved from Cu precptates formed wthn these steels after solutonzng and thermal agng. The sequence of precptaton, structure and composton has been studed by Mössbauer spectroscopy [5], feld-on mcroscopy (FIM) [6,7], atomprobe feld-on mcroscopy (APFIM) [8 11], atom-probe tomography (APT) [12], conventonal and hgh-resoluton electron mcroscopes (CTEM and HREM) [11,13 20], small-angle neutron scatterng (SANS) [21 24], extended X-ray absorpton fne structure (EXAFS) [25], X-ray * Correspondng author. Tel.: ; fax: E-mal addresses: rpkoll@nalco.com (R. Prakash Koll), d-sedman@ northwestern.edu (D.N. Sedman). absorpton spectroscopy (XAS) [26], small-angle X-ray scatterng (SAXS) [18,20], and computer smulatons [27 35]. Most of these studes have been on model bnary, ternary and quaternary Fe Cu-based carbon-free alloys and low-carbon steels. Fewer studes exst on Cu precptaton n multcomponent steels wth ncreased alloyng concentratons and mcrostructural complexty. The recent studes of Vaynman et al. [1] and Ishem et al. [36 38] characterzed the Cu-rch precptates found wthn multcomponent Fe Cu-based steels. The precptates were studed at nearpeak hardness and n the slghtly over-aged condton. Studes of growth and coarsenng (Ostwald rpenng) of the Cu precptates are lkewse few n number. Spech and Oran [14] studed coarsenng n bnary Fe Cu alloys contanng wt.% Cu, at agng temperatures between 730 and 830 C. These authors reported that the knetcs obey the t 1/3 power-law law for mean radus, hr(t), predcted by the Lfshtz Slyzov Wagner (LSW) [39,40] model for coarsenng, see below. Monzen et al. [41,42] studed an Fe 1.5 wt.% Cu alloy at temperatures between 600 and /$34.00 Ó 2008 Acta Materala Inc. Publshed by Elsever Ltd. All rghts reserved. do: /j.actamat

2 2074 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) C and also reported good agreement wth the t 1/3 power-law for hr(t). These authors also reported that the knetcs of the matrx supersaturaton, followed usng electrcal resstvty measurements, obey the t 1/3 powerlaw, as predcted by the LSW model. Sosson et al. [27] studed, by lattce knetc Monte Carlo smulatons, a bnary Fe 1.34 at.% Cu alloy aged between 300 and 500 C; they reported a t 1/3 dependency for hr(t). Recently, nvestgatons by Ishem et al. [37] descrbed the temporal evoluton of hr(t), number densty, N V (t), and concentraton profles for a multcomponent Fe Cu-based steel. The LSW model descrbes mean-feld dffuson-lmted coarsenng of precptates n dlute bnary alloys [43,44]. The asymptotc solutons of the LSW model [39,40] predct: hrðtþ hrðt 0 Þ / t 1=3 ; ð1þ where t s tme and hr(t 0 ) s the mean radus at the onset of quas-statonary coarsenng, where t 0 s greater than zero. The model also predcts: N V ðtþ N V ðt 0 Þ/t 1 ; ð2þ where N V (t 0 ) s the number densty at the onset of coarsenng, and: DC mat: ðtþ ¼ j½hc mat:;ff ðtþ C mat:;eq: ð1þšj / t 1=3 ; ð3þ where DC mat: ðtþ s the matrx (mat.) supersaturaton of an element at tme t between the far-feld concentraton of element, hc mat:;ff ðtþ, and equlbrum concentraton, C mat:;eq: ð1þ. The assumptons of the LSW model are: (1) the lnearzed Gbbs Thompson equaton s vald; (2) the precptate volume fracton s essentally zero; (3) the dffuson felds of the precptates do not overlap; (4) dlute soluton theory apples; (5) no elastc nteractons occur among the precptates; (6) the precptates have a sphercal morphology; (7) the precptates form and coarsen wth the composton gven by the equlbrum phase dagram; (8) the evaporaton condensaton mechansm s mplct n the model; and (9) the system s n the statonary state, whch s obtaned from the asymptotc solutons at nfnte tme. The Umantsev Olson (UO) model [45] extends the LSW model from bnary to multcomponent alloys, employng the same assumptons excludng that the alloy s a dlute soluton and the precptatng phase has a volume fracton of zero. It predcts the same tme exponents as the LSW model albet wth dfferent rate constants. And, fnally, Kuehmann and Voorhees (KV) [46] have analyzed ternary alloys n sgnfcantly greater detal. The focus of ths nvestgaton s the nucleaton (to a lmted extent), growth and coarsenng of the Cu-rch precptates n a concentrated multcomponent Fe Cu-based steel contanng 1.82 at.% Cu. Ths steel s beng studed as part of a program to develop an exploson-resstant steel for the US Navy [1,47,48]. The temporal evoluton of the morphologes and compostons of the Cu-rch precptates, from the as-quenched condton to 1024 h, when aged at 500 C, s studed utlzng APT [49]. We dscuss the temporal evoluton of the precptate core and precptate/a-fe matrx heterophase nterface compostons n detal. The power-law tme exponents are expermentally determned for hr(t), N V (t) and DC mat: ðtþ. Addtonally, we determne the exponents for the precptate (ppt.) core, DC ppt: ðtþ ¼ j½hc ppt: ðtþ C ppt:;eq: ð1þšj, and heterophase nterfacal regon (nt.), DC nt: ðtþ ¼ j½hc nt: ðtþ C nt:;eq: ð1þšj, supersaturatons of an element at tme t, where C j ðtþ s the concentraton as a functon of tme and C j:eq: ð1þ s the equlbrum concentraton for phase j (j = mat., ppt., or nt.). The results obtaned are compared to predctons of the UO model [45] and to earler nvestgatons for coarsenng n bnary Fe Cu alloys [14,27,41,42]. Furthermore, we dscuss the equlbrum morphology of the precptates and delneate the formaton of a N Al Mn phase at the Cu-rch precptate and a-fe matrx nterfaces. 2. Expermental methods 2.1. Materal detals A 45.5 kg (100 lb) steel heat was vacuum nducton melted and cast at ArcelorMttal Steel Global Research & Development, East Chcago, IN. The heat was reheated to approxmately 1150 C and hot-rolled n multple passes to a thckness of 12.5 mm and then ar-cooled to room temperature. The fnal hot-rollng temperature was approxmately 900 C. The bulk composton of the steel, determned by spectrographc analyss at ArcelorMttal Steel Global Research & Development, s presented n Table 1. We denote ths steel NUCu (170 desgnates the yeld strength n ks); detals regardng the development of the NUCu seres of steels can be found n Ref. [1]. The plates were trmmed and cut nto rods mm 250 mm and solutonzed at 900 C for 1 h and then quenched nto water at 25 C. Materal (12.5 mm 12.5 mm 25 mm blocks) for hardness testng and APT were aged at 500 C for 0.25, 1, 4, 16, 64, 256 or 1024 h and subsequently quenched nto water at 25 C Vckers mcrohardness Sectons were cold mounted n acrylc and polshed utlzng standard metallographc procedures to a fnal surface fnsh of 1 lm for hardness testng. Hardness testng was performed n accordance wth ASTM standards (ASTM E ) usng a Buehler Mcromet II mcrohardness tester wth a Vckers mcrohardness ndenter, a load of 500 g and a testng tme of 15 s. The reported mcrohardness values are derved from 10 measurements on each specmen. 1 The term NUCu stands for Northwestern Unversty copper alloyed steel.

3 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Table 1 Nomnal composton of hgh-strength low-carbon Fe Cu-based steel Cu C Al N S Mn Nb P S wt.% at.% Balance s Fe Atom-probe tomography The steel blocks were further reduced to 0.3 mm 12.5 mm 25 mm coupons, cut from the center, utlzng an abrasve saw. The APT tp blanks (0.3 mm 0.3 mm 25 mm) were mechancally cut from the coupons and electropolshed usng standard technques [49,50]. Intal polshng was performed wth a soluton of 10 vol.% perchlorc acd n acetc acd at V DC at room temperature. Ths was followed by a manually controlled pulsed fnal-polshng step usng a soluton of 2 vol.% perchlorc acd n butoxyethanol at 10 5 V DC at room temperature, producng a tp wth a radus <50 nm. LEAP TM tomography [51,52] was performed at a specmen temperature of 50 K under ultrahgh vacuum (UHV) condtons of Pa ( torr). The pulse repetton rate was Hz and the pulse-voltageto-standng-dc voltage rato (pulse fracton) was 15 20%. Vsualzaton and reconstructon of the APT data s performed usng the Imago Vsualzaton and Analyss Software (IVAS TM ) package. The precptates are dentfed utlzng an soconcentraton surface methodology [53] wth the threshold concentraton set at 10 at.% Cu, whch gves morphologcally and compostonally stable results. The parameters chosen to obtan nose-free soconcentraton surfaces are a voxel sze of 1 nm, a delocalzaton dstance of 4 nm, a sample count threshold of 5%, and a polygon flter level of Concentraton profles wth respect to dstance from the reference soconcentraton surfaces are obtaned utlzng the proxmty hstogram (proxgram for short) wth a bn wdth of 0.25 nm [54]. The ±2r standard devatons for the reported concentratons are gven by countng statstcs: sffffffffffffffffffffffffffffffffffffff c ð1 c Þ r ¼ ; ð4þ N TOT where c s the atomc fracton of an element n a bn and N TOT s the total number of ons n a bn. The sphercal volume equvalent radus, R, of a precptate [55] s gven by: R ¼ 3 1 N atoms 3; ð5þ 4p q th g where N atoms s the number of atoms detected wthn a delneated precptate, q th, s the theoretcal atomc densty, whch s equal to 84.3 atoms nm 3 for ths steel, and g s the estmated detecton effcency of 0.5 of the multchannel plate detector. The value of N atoms belongng to a precptate s determned from the envelope method [49,56], based on a maxmum separaton dstance, d max, of nm, a mnmum number of solute atoms, N mn, of Cu atoms (dependng on the agng tme), and a grd spacng of 0.2 nm. The detals regardng selecton of values for the parameters utlzed n the envelope method for ths steel can be found n Ref. [55]. The radal dstrbuton functon (RDF) [57], wth respect to Cu for ths steel, s defned as the average concentraton of an element at a dstance r from the Cu atoms, hc Cu ðrþ, normalzed to the overall concentraton of the element, C o, n the alloy: RDF ¼ hccu ðrþ C o ¼ 1 C o X N Cu k¼1 N k ðrþ N k tot ðrþ ; where N k ðrþ s the number of atoms n a radal shell around the kth Cu atom at a dstance r, and N k totðrþ s the total number of atoms wthn the shell at a dstance r. The partal RDFs are calculated usng 0.01 nm ncrements from 0.2 to 1.0 nm. Only the partal RDFs for r P 0.2 nm are dsplayed, as the physcal nterpretaton at smaller r values s dffcult due to possble on trajectory effects. The results are smoothed by a weghted movng average usng a Gaussan-type splne functon [53]. Radal dstrbuton functon values >1 mply a postve correlaton wth Cu atoms, whereas values <1 mply a negatve correlaton wth Cu atoms. A value of unty ndcates a random dstrbuton of element wth respect to Cu. Greater detal regardng the RDF can be found n Refs. [57,58]. The partal RDF s utlzed to evaluate any potental clusterng, n the as-quenched specmen, wth a greater degree of senstvty than the soconcentraton surface [53] or envelope [49,56] methodologes Transmsson electron mcroscopy Specmens for conventonal transmsson electron mcroscopy (TEM) were mechancally cut from the 0.3 mm 12.3 mm 25 mm coupons aged for 1024 h. The 3 mm fols were then mechancally ground to a thckness of approxmately 150 lm and subsequently twn-jet electropolshed n a soluton of 5 vol.% perchlorc acd n methanol at 60 C and 15 V. The specmen perforatons were further thnned utlzng on-beam thnnng (IBT) at 3.0 kv and 3.5 ma. The specmens were examned n a Htach H-8100 transmsson electron mcroscope operatng at 200 kv. 3. Results 3.1. Temporal evoluton of the morphology Fg. 1 dsplays the Vckers mcrohardness (VHN) as functon of agng tme. Ths fgure demonstrates that a sgnfcant ncrease n hardness s attanable when the steel s ð6þ

4 2076 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Fg. 1. Vckers mcrohardness (VHN) as a functon of agng tme (h) when the NUCu-170 steel s aged at 500 C. A plateau of 395 VHN s seen between 2 and 16 h. aged at 500 C. Pror to 2 h of agng the steel s under-aged. From 2 to 16 h of agng the mcrohardness reaches a plateau of approxmately 395 VHN. The observed hardness plateau s a result of dfferent nucleaton and growth rates of the Cu-rch and NbC precptates [59,60] present wthn the steel and s consstent wth earler observatons on NUCu steels contanng at.% Cu [1]. At 64 h of agng the mcrohardness decreases to ± 8.7 VHN, ndcatng the steel s over-aged. Agng for tmes greater than 64 h results n further reducton of hardness, reachng a value of ± 6.1 VHN at 1024 h. Other researchers have reported agng tmes rangng from 1.94 to 10 h to acheve peak hardness n model Fe Cu-based alloys contanng between 1.1 and 1.5 at.%. Cu, when aged at 500 C [7,10,12,13,20,24,30,61,62]. Evoluton of the Cu-rch precptate morphology s seen qualtatvely n Fg. 2a g, whch depct nm 3 (3000 nm 3 ) subsets of an analyzed volume for each agng tme studed, contanng approxmately 130,000 atoms. These fgures llustrate that the precptates, whch are delneated wth 10 at.% Cu soconcentraton surfaces, grow and coarsen wth ncreasng agng tme. The precptates are ntally spherodal and become more ellpsodal or rod-lke at 1024 h of agng, although, as dscussed below, not all precptates at 1024 h have a rod-lke morphology. A representatve Cu-rch precptate for a 1 h agng tme s presented n Fg. 3a. In Fg. 3a c the Cu atoms are orange, 2 the N atoms are green, the Al atoms are teal, the Mn atoms are mustard, and the Fe atoms are blue, 2 For nterpretaton to color n Fg. 3, the reader s referred to the web verson of ths artcle. Fg. 2. Copper-rch precptates as delneated by 10 at.% Cu soconcentraton surfaces, when the steel s aged at 500 C for: (a) 0.25 h; (b) 1 h; (c) 4 h; (d) 16 h; (e) 64 h; (f) 256 h; (g) 1024 h. Each reconstructon, nm 3 (3000 nm 3 ), s a subset of an analyzed volume and contans approxmately 130,000 atoms. and the S atoms are gray. The dstrbuton of N, Al and Mn atoms adjacent to the Cu atoms, at the precptate/a- Fe matrx heterophase nterface, s llustrated. The segregaton s more pronounced on one sde of the precptate. Other precptates for ths agng tme exhbt smlar segregaton and ths observaton s consstent wth prevously reported results n smlar concentrated multcomponent steels [36 38]. A representatve precptate for the 4 h agng tme s dsplayed n Fg. 3b. The observed N, Al and Mn segregaton s smlar to that observed at 1 h, whch s dscussed n detal below. We observe smlar segregaton behavor for N, Al and Mn n other precptates at ths

5 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Fg. 3. Three-dmensonal atom-probe tomographc reconstructons of representatve precptates n the (a) 1 h, (b) 4 h and (c) 1024 h aged condtons. The precptate presented n (c) s of rod-lke morphology. The Cu, N, Al and Mn atoms are shown as spheres (not to scale), allowng vsualzaton of the precptates and heterophase nterfaces. Only 20% of the Fe and 50% of the Cu, N, Al, and Mn atoms are shown n (c) for clarty. agng state. The precptate dsplayed n Fg. 2g, 1024 h of agng, s larger than the cross-sectonal area employed, and extends beyond the boundares of the volume dsplayed. The exact same precptate s presented n Fg. 3c. The precptate has a rod-lke morphology consstng almost entrely of Cu atoms n ts core and s partally surrounded by N, Al and Mn atoms, where the observed segregaton s greater qualtatvely than for the 4 h aged state, see below. The segregaton of the elements N, Al and Mn s not unform around the Cu atoms n the core of the precptate. Unlke for the 1 and 4 h agng condtons, the observed segregaton occurs predomnantly on two sdes of ths precptate. Not all precptates at ths agng tme possess the same morphology presented n Fgs. 2g and 3c. Other precptates exhbt a more spherodal morphology wth N, Al and Mn segregatng toward one sde of the Cu-rch core (Fg. 4) Nucleaton and growth We do not fnd any evdence of precptaton wthn the matrx n the as-quenched state utlzng the soconcentraton surface methodology [53] wth the desgnated parameters. Reducton of the threshold concentraton to values as small as the bulk Cu concentraton does not affect the observed result. We also do not fnd any evdence of precptaton utlzng the envelope method [49,56] wth the desgnated parameters and the methodology descrbed n Ref. [55]. The steel, when n the supersaturated state, exhbts some evdence of formng metastable clusters (embryos). Utlzng the same value for d max and settng N mn =3Cu atoms gves 813 clusters (embryos) contanng 3 20 Cu atoms. The majorty of clusters (embryos), 801, contan less than 11 Cu atoms, suggestng they are metastable clusters (embryos). The larger agglomeratons are possbly stable clusters,.e. early-stage stable nucle. To further evaluate ths steel n the as-quenched state, an expermental partal RDF, determned by Eq. (6) and presented n Fg. 5a f, s utlzed. The 1st 6th nearest-neghbor (NN) dstances and NN numbers for body-centered cubc (bcc) a-fe are gven on the top abscssa. These fgures llustrate that strong overall postve correlatons exst for Cu Fg. 4. Three-dmensonal atom-probe tomographc reconstructon of a spherodal shaped precptate found n the 1024 h aged condton. The Cu, N, Al and Mn atoms are shown as spheres (not to scale) allowng vsualzaton of the precptates and heterophase nterfaces. Only 20% of the Fe and 50% of the Cu, N, Al and Mn atoms are shown for clarty. Cu (Fg. 5a), Cu Al (Fg. 5d) and Cu S (Fg. 5f) atoms, whereas a strong negatve correlaton exsts for Cu Mn (Fg. 5e) atoms. The strong postve oscllaton for Cu Cu between the second and thrd NN postons along wth the results of the envelope method suggests that the onset of phase decomposton may occur as early as after solutonzng and quenchng to room temperature. The slght negatve oscllaton for Cu Fe (Fg. 5b) s a mnmum between the frst and second NN postons, suggestng that Cu and Fe are negatvely correlated at ths locaton. However, the closeness of the partal RDF to unty ndcates that Fe s almost randomly dstrbuted relatve to Cu. The Cu N partal RDF (Fg. 5c) demonstrates a weak overall negatve correlaton but the proxmty of the partal RDF to unty between the frst and second NN postons suggests a random dstrbuton of N wth respect to Cu.

6 2078 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Fg. 5. Expermental partal radal dstrbuton functon (RDF) n the asquenched state for: (a) Cu Cu atoms; (b) Cu Fe atoms; (c) Cu N atoms; (d) Cu Al atoms; (e) Cu Mn atoms; and (f) Cu S atoms. Values of RDF >1 ndcate a postve correlaton between atoms, whereas values of RDF < 1 ndcate a negatve correlaton between atoms; when the RDF = 1 ths ndcates a random dstrbuton. The 1st 6th nearest neghbor (NN) dstances and NN numbers for bcc Fe are ndcated on the top abscssa. By 0.25 h, Cu-rch precptates have formed wthn the matrx (Fg. 2a) wth hr = 1.2 ± 0.1 nm (Fg. 6a), where the error s gven by standard error of the mean. At 0.25 h, N V = (5.2 ± 1.8) m 3 (Fg. 6b) and the volume fracton, /, s equal to 0.3 ± 0.01%, where the reported error s based on countng statstcs for both quanttes. Further agng to 1 h results n a sgnfcant ncrease of N V (t) to (4.2 ± 0.5) m 3, wth hr = 1.5 ± 0.05 nm, and / equal to 3.6 ± 0.06%. The ncrease n N V (t) occurs wth a temporal dependency of t 2.7, n conjuncton wth a tme exponent for hr(t) equal to 0.16, ndcatng that nucleaton s occurrng. The slowly ncreasng value of hr(t) ndcates, however, that growth s also occurrng smultaneously Growth and coarsenng The quantty N V (t) reaches a maxmum at 1 h of agng and subsequently decreases ndcatng the onset of coarsenng. The value of hr(t) ncreases from 1.5 ± 0.05 nm at 1 h to 6.5 ± 0.7 nm at 1024 h (Fg. 6a) and the temporal dependences for hr(t) are t 0.16±0.01 from 1 to 64 h and t 0.34±0.09 Fg. 6. The (a) mean precptate radus, hr(t); and (b) number densty, N V (t), as a functon of agng tme for an agng temperature of 500 C. from 64 to 1024 h. Only from 64 h onwards s the tme exponent equal to that predcted by Eq. (1). The quantty N V (t) decreases by a factor of 56 to (7.4 ± 0.5) m 3 at 1024 h wth / equal to 4.3 ± 0.06%. The temporal dependences for N V (t) are t 0.45±0.03 from 1 to 64 h and t 0.63±0.07 from 64 to 1024 h, whch ndcates coarsenng s occurrng slower than that predcted by Eq. (2) Temporal evoluton of composton The temporal evoluton of the concentraton profles s presented n Fgs. 7 and 8a c. The fgures represent the temporal evoluton of the Cu, Fe, N, Al, Mn and S profles from the under-aged condton, 0.25 h (Fg. 7) and 1 h (Fg. 8a), to the peak yeld strength aged condton, 4h (Fg. 8b), and fnally to the over-aged condton at 1024 h (Fg. 8c). These fgures demonstrate clearly that the composton of the matrx, precptates and precptate/a-fe matrx heterophase nterfaces are evolvng temporally. The observed segregaton of N, Al and Mn s non-monotonc (confned), whereas the observed segregaton of Cu, Fe and S s monotonc (non-confned) Far-feld matrx compostons In ths nvestgaton the plateau ponts [63] wthn the matrx (far-feld) yeld the a-fe matrx concentratons (Table 2), where the plateau regon of the proxgram s

7 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) delneated by utlzng the Fe concentraton profles as a fducal marker. Only data ponts wthn the flat regon of the profle, a mnmum of 1.5 nm away from the heterophase nterface and wth ±2r < 0.4 at.%, are ncluded. The concentraton of Cu decreases n the a-fe matrx, reachng a value of 0.2 ± 0.01 at.% at 1024 h of agng. The concentratons of N, Al and Mn n the a-fe matrx also decrease wth prolonged agng tme. The concentratons of S and Fe ncrease n the a-fe matrx wth ncreasng agng tme, whch becomes enrched n these elements at 1024 h. Nobum s not detected wthn the a-fe matrx, wth the excepton of the 1 h aged state, whle C s detected at a reduced concentraton when compared to ts nomnal value. In the NUCu steels Nb s not normally detected by APT due to the exstence of NbC precptates at a smaller number densty ( m 3 ) than the Cu-rch precptates [1,59,60]. In the 1 h aged state the presence of a sngle NbC precptate at the boundary of an analyzed volume gves the measured Nb concentraton. Smlarly, C s detected at lower concentratons wthn the matrx due to the presence of NbC precptates, cementte (Fe 3 C) [37] and segregaton at gran boundares [37]. Fg. 7. Proxgram concentraton profles (at.%) for Cu, Fe, Al, S and Mn, when the steel s aged at 500 C for 0.25 h. The dotted vertcal lne corresponds to the a-fe matrx/precptate heterophase nterface Composton of the Cu-rch precptates At 0.25 h of agng the precptate cores are Cu-rch (46.7 ± 4.3 at.%) but also contan sgnfcant concentratons of Fe, N and Al, and smaller quanttes of S and Mn (Table 3). At 1 h of agng the precptate cores are also Cu-rch (44.6 ± 3.1 at.%) but stll contan sgnfcant Fg. 8. Proxgram concentraton profles (at.%) for Cu, Fe, N, Al, S and Mn, when the steel s aged at 500 C, for (a) 1 h, (b) 4 h and (c) 1024 h. The dotted vertcal lnes ndcate the a-fe matrx/precptate heterophase nterface.

8 2080 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Table 2 Composton n at.% of the a-fe matrx of the specmens aged at 500 C as determned by atom-probe tomography Cu C Al N S Mn Fe Nb 0.25 h 1.6 ± ± ± ± ± ± ± 0.01 ND 1 h 0.6 ± ± ± ± ± ± ± ± h 0.4 ± 0.07 ND 0.8 ± ± ± ± ± 0.5 ND 16 h 0.5 ± ± ± ± ± ± ± 0.5 ND 64 h 0.3 ± ± ± ± ± ± ± 1.1 ND 256 h 0.3 ± ± ± ± ± ± ± 0.5 ND 1024 h 0.2 ± ± ± ± ± ± ± 0.04 ND ND = Not detected. Table 3 Composton n at.% of the Cu-rch precptate cores of the specmens aged at 500 C as determned by atom-probe tomography Cu C Al N S Mn Fe Nb 0.25 h 46.7 ± ± ± ± ± ± ± 4.2 ND 1 h 44.6 ± ± ± ± ± ± ± 3.1 ND 4 h 53.5 ± 2.3 ND 11.3 ± ± ± ± ± 2.1 ND 16 h 66.0 ± 1.8 ND 9.8 ± ± ± ± ± 1.4 ND 64 h 88.4 ± 1.7 ND 1.6 ± ± ± ± ± 1.3 ND 256 h 78.3 ± ± ± ± ± ± ± 1.6 ND 1024 h 97.1 ± 0.8 ND 1.2 ± ± ± ± 0.2 ND ND = Not detected. amounts of Fe, N, Al, Mn and S. As agng progresses, the Cu concentraton wthn the cores ncreases to 53.5 ± 2.3 at.% at 4 h and acheves a value of 97.1 ± 0.8 at.% at 1024 h. Concomtantly, the concentratons of Fe and S decrease, reachng a value of 0.3 ± 0.2 at.% and zero, respectvely, at 1024 h. The concentratons of N, Al and Mn wthn the cores have a more complcated behavor, whch s dscussed below. Pror to 1024 h, however, all three elements are found enhanced wthn the cores at the agng tmes studed. At 1024 h the cores are depleted n N (0.7 ± 0.4 at.%) and Al (1.2 ± 0.4 at.%), relatve to ther nomnal concentratons, whereas Mn s enrched (0.8 ± 0.4 at.%). Carbon and Nb are not found wthn the precptate cores, wthn the prescrbed expermental uncertanty Composton of precptate/a-fe matrx heterophase nterfaces The composton of the precptate/a-fe matrx heterophase nterfaces s also llustrated n Fgs. 7 and 8a c. The heterophase nterfacal regon, at as early as 0.25 h of agng, s enrched n Cu, N, Al and Mn but s depleted n Fe (Table 4). The N peak concentraton (8.3 ± 1.8 at.%) s located at a dstance of nm close to the Mn peak concentraton (1.3 ± 0.5 at.%), whch s found at a dstance of nm. The Al peak concentraton (13.5 ± 7.2 at.%) s located toward the center of the precptate at nm. At 1 h, the N (9.9 ± 0.5 at.%) and Mn (1.5 ± 0.2 at.%) peak concentratons are located at a dstance of nm, whereas the Al peak concentraton (7.1 ± 1.7 at.%) s found closer to the center of the precptate at a dstance of nm. At 4 h the N (10.7 ± 0.3 at.%) and Mn (1.7 ± 0.4 at.%) peak concentratons are co-located at a dstance of nm, whle the Al peak concentraton (12.2 ± 3.2 at.%) s at a dstance of nm. At 1024 h of agng a dstnct enhancement of N, Al and Mn s observed at the heterophase nterfaces, wth peak concentratons of 22.1 ± 0.8 at.% N and 21.2 ± 0.8 at.% Al at a dstance of nm and 3.3 ± 0.3 at.% Mn at a dstance of nm. The evoluton of the N, Al and Mn concentratons wthn the heterophase nterfaces s related to that wthn the precptate cores and s dscussed below. Table 4 Composton n at.% of the precptate/a-fe matrx heterophase nterfaces of the specmens aged at 500 C as determned by atom-probe tomography Cu C Al N S Mn Fe Nb 0.25 h 3.3 ± ± ± ± ± ± ± 0.04 ND 1 h 5.0 ± ± ± ± ± ± ± ± h 6.2 ± 0.2 ND 3.2 ± ± ± ± ± 0.3 ND 16 h 8.7 ± ± ± ± ± ± ± 0.07 ND 64 h 8.4 ± ± ± ± ± ± ± 0.09 ND 256 h 10.4 ± ± ± ± ± ± ± 0.2 ND 1024 h 10.3 ± ± ± ± ± ± ± 0.1 ND ND = Not detected.

9 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Slcon s also found wthn the heterophase nterfacal regon at a slghtly enhanced concentraton of approxmately at.%. Nobum s not detected, wth the excepton of the 1 h aged condton, whle C s detected at a reduced concentraton when compared to ts nomnal value Parttonng ratos The temporal evoluton of the parttonng rato [55], j ppt:=mat: ðtþ ¼C ppt: ðtþ=c mat:;ff ðtþ; ð7þ s dsplayed n Fg. 9. The standard error for j s determned by standard error propagaton methods for the concentraton errors [64]. Ths fgure llustrates clearly that Cu parttons to the precptates, whereas Fe and S partton to the matrx. The dashed lne, beyond 256 h, for the S rato, represents the zero S concentraton wthn the precptate core at 1024 h of agng. Nckel, Al and Mn exhbt a more complcated behavor but slghtly prefer the Cu-rch precptate phase from 0.25 to 256 h. At 1024 h Fg. 9 shows that Al and Mn have a slght preference for the precptate, whereas N has a slght preference for the a-fe matrx. The proxmty of all three elements, though, to the dashed lne separatng precptate and matrx phases ndcates that the three elements partton to the nterfacal regon. The parttonng ratos for N and Mn exhbt smlar trends from 0.25 to 16 h, where they dverge slghtly, whereas the profles of N and Al follow smlar trends from 64 to 1024 h. Fgs The equlbrum concentratons are obtaned by extrapolatng the concentraton of each element as a functon of (agng tme) 1/3 to nfnte tme, analogous to the procedure found n Ref. [65]. The expermentally determned equlbrum concentraton values, the concentratons measured at 1024 h of agng, and those predcted by Thermo-Calc are dsplayed n Table 5. The tme exponents for DC mat: ðtþ (Fg. 10a c) are 0.23 ± 0.07 for Cu, 0.32 ± 0.06 for Fe, 0.25 ± 0.03 for N, 0.22 ± 0.09 for Al, 0.14 ± 0.14 for Mn, and 0.25 ± 0.13 for S. The tme exponents for all elements, except S, are derved from agng tmes of 1 to 1024 h, whle that of S s from 1 to 64 h, snce DC mat: ðtþ ¼0 at.%, s reached at 256 h for S. The tme exponents for DC ppt: ðtþ (Fg. 11a c) are 0.26 ± 0.06 for Cu, 0.37 ± 0.04 for Fe and 0.41 ± 0.05 for S. The exponents for Cu and Fe are derved for agng tmes of h, whle the exponent for S s for agng tmes of h. The equlbrum concentraton of S s 0 at.%, whch s reached at 1024 h. The tme exponents for DC nt: ðtþ (Fg. 12a c) are 0.16 ± 0.12 for N, 0.30 ± 0.07 for Al and 0.20 ± 0.09 for Mn, where the exponents are derved for agng tmes of h Coarsenng knetcs The matrx, precptate core and precptate/a-fe matrx heterophase nterface supersaturatons are dsplayed n Fg. 9. Parttonng ratos, j ppt:=mat: ðtþ, of the concentrated multcomponent Fe Cu steel as a functon of agng tme, when aged at 500 C, from 0.25 to 1024 h. The horzontal dashed lne ndcates the dvson between the precptate and matrx phases. The dashed arrow for the S rato represents the zero S concentraton wthn the precptate at 1024 h. Fg. 10. Double logarthmc plots of matrx supersaturatons, DC mat: ðtþ,as a functon of agng tme, when aged at 500 C, for (a) Cu and Fe; (b) N and Al; and (c) Mn and S. The slopes of the plots yeld the coarsenng tme exponents for each element.

10 2082 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Fg. 11. Double logarthmc plots of the supersaturatons n the Cu-rch precptates cores, DC ppt: ðtþ, as a functon of agng tme, when aged at 500 C, for (a) Cu, (b) Fe and (c) S. The slopes of the plots yeld the coarsenng tme exponents for each element. 4. Dscusson 4.1. Morphology Fg. 12. Double logarthmc plots of the supersaturatons n the heterophase nterfaces, DC nt: ðtþ, as a functon of agng tme, when aged at 500 C, for (a) N, (b) Al and (c) Mn. The slopes of the plots yeld the coarsenng tme exponents for each element. The morphologes observed by APT n ths nvestgaton are consstent wth earler TEM studes, whch show that the bcc Cu? 9R Cu? 3R Cu? fcc e-cu phase changes are accompaned by a change from a spherodal morphology to an ellpsodal or rod-lke morphology [13 16]. Our observaton of a rod-lke morphology, by APT, n our over-aged Fe Cu alloy system s, to the best of our knowledge, the frst report of such a structure. Transmsson electron mcroscopy observatons of the 1024 h aged condton confrm the rod-lke morphology of the precptates (Fg. 13). The precptates orent along the [110] Cu k[111] a Fe (111) Cu k(110) a Fe drecton, whch s the Kurdjumov Sachs relatonshp, and s consstent wth earler observatons for bnary Fe Cu alloys [13,14,16,20, 41,42]. The equlbrum morphology of a precptate s determned by the balance between elastc and nterfacal ener- Table 5 Extrapolated to nfnte tme equlbrum compostons of the Cu-rch precptates, a-fe matrx and heterophase nterfaces compared to measured overall composton at 1024 h and also calculated from Thermo-Calc Cu Al N S Mn Fe Extrapolated from APT data to nfnte tme Measured from specmens aged 1024 h Calculated from Thermo-Calc and SGTE soluton database Precptate 100 ± ± ± ± ± Interface 10.2 ± ± ± ± ± ± 0.3 Matrx 0.2 ± ± ± ± ± ± Precptate 97.1 ± ± ± ± ± 0.2 Interface 10.3 ± ± ± ± ± ± 0.1 Matrx 0.2 ± ± ± ± ± ± 0.04 Precptate (bcc) Precptate (fcc) Interface Matrx

11 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) Fg. 13. Brght-feld TEM mcrograph of the steel when aged at 500 C for 1024 h. The precptates are algned along the Kurdjumov Sachs orentaton: [110] Cu k[111] a Fe (111) Cu k(110) a Fe. ges [43]. Thompson et al. [66] have ntroduced the L parameter, whch represents the equlbrum energy state of a precptate, and s gven by: L ¼ d2 C 44 l r ; ð8þ ppt:=mat: where d s the lattce msft, C 44 s an elastc shear modulus, l s a characterstc dmenson of the precptate, and r ppt./mat. s the precptate/matrx nterfacal free energy. For the purpose of calculatng the L parameter we utlze ntally values for d, C 44 and r ppt./mat. for a bnary Fe Cu alloy. The quantty d fcc-cu/bcc-fe s equal to where the lattce parameter, a 0, s equal to nm for Cu (fcc) and nm for a-fe (bcc) [67]. The remanng quanttes are C fcc-cu 44 ¼ 75 Gpa [67], whch s also approxmately the md-pont for the range of reported expermental and smulated values [31,67,68], l = 13 nm, and r ppt./mat. 600 mj m 2 [8], where we assume r ppt./mat. s sotropc. Eq. (8) yelds L = 87, whch suggests an equlbrum morphology controlled solely by the elastc energy. Other reported values for r ppt./mat. yeld smlar results. Ludwg et al. [33] report a value of r ppt./mat. = 245 mj m 2 for the sdes of a cylndrcal precptate, whch gves L =212. Spech and Oran [14] estmate a value of r ppt./mat. = 466 mj m 2 for the ends of a rod-shaped precptate, whch gves L = 112. Monzen et al. [41] report a value of r ppt./mat. = 1100 mj m 2 for the ends of a rod-shaped precptate, whch gves L = 47. Lee et al. [69] have shown, however, that the equlbrum morphology of precptates that are elastcally softer than the matrx, such as Cu (fcc) precptates wthn an a-fe (bcc) matrx (C bcc-fe 44 ¼ 116 GPa [67] or C bcc-fe 44 ¼ 101 GPa [31]), s a platelet, rrespectve of elastc ansotropy and orentaton relatonshp. The expermentally observed precptates possess the twofold symmetry of platelets but also have rounded edges and ellpsodal bodes, whch suggests that both elastc and nterfacal energy affect the equlbrum morphology. The dfferences between our observatons and the results of Eq. (8) are partally attrbuted to an uncertan nterfacal energy, whch s greater than the reported values, and s also affected by the formaton of a N Al Mn shell, whch s dscussed below. Agng beyond 1024 h or at hgher temperatures may result n a platelet morphology. Other researchers, however, dd not observe platelet morphologes n bnary Fe Cu alloys when agng at C for 1000 h [13], or at C for 300 h [14], and at 550 C for 1000 h [16]. An addtonal consderaton s that the L parameter assumes the same elastc constants for the precptate and matrx phases, whch s not the case for a Cu (fcc) precptate n an a-fe (bcc) matrx. Snce C bcc-fe 44 s greater than C fcc-cu 44 [67], the dfference n elastc constants cannot alone account for the dsparty between our observatons and the results of Eq. (8). The precptates, however, are more complcated than a pure bnary phase nterface. The formaton of a N Al Mn shell also affects the elastc constants and lattce msft of the Cu-rch precptates. For the purpose of calculatng L we utlze the values of d, C 44 and r ppt./mat. for a bnary equatomc NAl (B2 structure). For NAl (B2 structure), a 0 = nm [70], yeldng d B2-NAl/bcc-Fe = The quantty C B2-NAl 44 ¼ 114:7 Gpa [71], whch s also approxmately the md-pont for the range of reported expermental values [72,73]. Snce the nterfacal energy of NAl n a-fe s unknown we utlze ntally the surface energy of NAl, r ppt./mat. ff 1850 mj m 2 [70,74], where we assume r ppt./mat. s sotropc. Snce approxmately 50% of the N and Al bonds at a surface are free, ths estmate represents an upper bound for r ppt./mat.. Eq. (8) yelds L = 0.056, whch suggests an equlbrum morphology controlled solely by nterfacal energy. A smaller value for r ppt./mat. would result n a larger L value, but assumng lnearty, a 50% reducton n r ppt./mat. would stll yeld an equlbrum morphology controlled by nterfacal energy. Clearly, both elastc and nterfacal energy contrbute to the equlbrum morphology of the precptates. The ends of the rod-lke precptates are domnated by elastc energy, whereas the sdes are controlled by nterfacal energy, whch s where the N, Al and Mn segregaton s predomnantly found. Other researchers have attrbuted the rod-lke morphology of Cu precptates n Fe Cu-based alloys to mnmzaton of elastc stran energy [13] or ansotropy of nterfacal energy [14]. Ths morphology has been attrbuted to shear assocated wth partal dslocatons producng an nvarant plane-stran transformaton to the 3R structure from the 9R structure, followed by relaxaton to the fcc structure by dffusonal growth to mnmze nterfacal energy [16] Coarsenng knetcs Our expermental results ndcate nucleaton s occurrng pror to 0.25 h, nucleaton and growth from 0.25 to 1 h,

12 2084 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) followed by a transton, and subsequently, after 64 h, an ncreasng proporton of growth and coarsenng. A tmelaw exponent of 1/2, for dffuson-controlled growth [43,75], however, s not observed, contrary to earler nvestgatons by TEM [13] and APFIM [6] n model bnary Fe Cu alloys. Assumng coarsenng starts at 64 h gves a temporal dependency of t 0.34±0.09 for hr(t), whch s n approxmate agreement wth the UO model value of 1/3 and ndcates a dffuson-lmted coarsenng mechansm. Assumng coarsenng starts at 64 h gves a temporal dependency for N V (t) oft 0.63±0.07, whch s not n agreement wth the UO model predcton of 1 and ndcates a slower coarsenng rate and an admxture of growth and coarsenng. Our results dffer from those reported by Spech and Oran [14] and Monzen et al. [41,42], who report good agreement for the t 1/3 power-law for hr(t) to agng tmes of 300 h. The alloys n ther studes are bnary Fe Cu and ternary Fe Cu N carbon free alloys, respectvely, whch can account for the dfferences wth our results. Addtonally, the agng temperatures n Ref. [14] were between 730 and 830 C, and those n Refs. [41,42] were between 600 and 750 C, whch s sgnfcantly greater than the 500 C we employed. At the reported temperatures quas-statonary state coarsenng would be acheved more rapdly than at 500 C. Our results also dffer from the lattce knetc Monte Carlo smulatons of Sosson et al. [27] who obtaned a t 1/3 dependency for hr(t) n a bnary Fe 1.34 at.% Cu alloy aged at 300 and 500 C. Ther smulatons, however, dd not nclude elastc stran effects, whch can have a sgnfcant effect on growth and coarsenng at longer agng tmes n a Fe Cu system. The tme exponents for hr(t) from 1 to 64 h are closer to the 1/6 to 1/4 values predcted for the cluster dffuson coagulaton coarsenng mechansm [43]. The recent observatons of Sudbrack et al. [76] and Mao et al. [77] on a model N Al Cr superalloy ndcate that precptates followng such a mechansm would form nterconnected necks wth clearly delneated lattce planes. We do not, however, observe nterconnected necks wth lattce planes employng LEAP TM tomography, ndcatng that the cluster dffuson coagulaton mechansm s noperatve. The expermentally determned tme exponents for DC mat: ðtþ, wth the excepton of Fe, do not satsfy the UO model predcton of 1/3. The devaton from the model value of the exponent for S s possbly due to the rapd approach of the S matrx concentraton to ts equlbrum value; S s a comparatvely fast dffusng element n a-fe (Table 6). The DC mat: ðtþ tme exponents for Cu, N, Al and Mn are affected by the bcc Cu? 9R Cu? 3R Cu? fcc e- Cu [13 16] phase changes of the precptates and the formaton of the N Al Mn shells. The tme exponents for DC ppt: ðtþ and DC nt: ðtþ are also affected, but more sgnfcantly, by the phase changes and formaton of the N Al Mn shell adjacent to the Cu-rch precptate cores. Wthn the heterophase nterfaces (Fg. 12a c) only the tme exponent for Al s close to, but not exactly dentcal wth, the Table 6 Tracer dffusvtes of Cu, N, Mn, S and Al n a-fe at 500 C D (m 2 s 1 ) Ref. Cu [78] N [79] Mn [80] S [81] Al [82] model predcton, whch can possbly be ascrbed to the relatvely large dffusvty of Al n a-fe (Table 6). The temporal dependency for Mn s less than that of Al but greater than that of N, whch corresponds to the dffusvty of Mn, whch s slower than that of Al but faster than that of N (Table 6). The phase changes affect the sold solublty of N, Al and Mn wthn the Cu-rch precptates. For example, when the precptates are <2 nm n radus, whch corresponds to the bcc structure, the Mn concentraton ntally ncreases from 0.6 ± 0.6 at.% at 0.25 h to 1.5 ± 0.6 at.% at 4 h. The Mn concentraton subsequently decreases to 1.1 ± 0.6 at.% at 64 h of agng and then ncreases to 2.3 ± 1.1 at.% at 256 h, before decreasng to 0.8 ± 0.4 at.% at 1024 h. At an agng tme of 1024 h the observed admxture of spherodal and rod-lke precptates ndcates a 9R (spherodal) and 3R or fcc (ellpsodal) precptates. Our observatons support the recent phase-feld smulatons of Koyama et al. [29]. Koyama et al. demonstrated that the Mn concentraton ncreases wthn the Cu-rch bcc precptates, but upon transformaton to the fcc phase the Mn concentraton decreases. The authors dd not report the effect of the 9R and 3R phases on Mn concentraton. Our results dffer from those of Monzen et al. [41,42], who report agreement wth the t 1/3 power-law for the knetcs of depleton of supersaturaton wthn the a-fe matrx. The alloys n ther studes are bnary Fe Cu and ternary Fe Cu N carbon-free alloys and were aged between 600 and 750 C; these alloys should reach quasstatonary state coarsenng more rapdly than our more complex alloy Composton of the Cu-rch precptates Precptates less than 2 nm n radus We measure a Fe concentraton between 39.9 ± 4.2 and 27.6 ± 2.1 at.% when the Cu-rch precptates are smaller than 2 nm n radus (Fg. 6a), whch corresponds to the bcc structure [5,16,25]. Ths s consstent wth observatons n a smlar steel, wth 1.17 at.% Cu, denoted NUCu (140 desgnates the yeld strength n ks and 1 desgnates the expermental heat number), where the Fe concentraton of the precptates was measured to be 25 at.% near-peak hardness [1]. Earler atom-probe studes of Fe Cu-based alloys, thermally aged at 500 C, have also detected sgnfcant Fe wthn Cu-rch precptates smaller than 2 nm n radus n both bnary [8,9,30] and multcomponent [9,10,37] alloys. In contrast to atom-probe measurements, SANS measurements nfer, based on several strong assumptons, an Fe concentraton

13 R. Prakash Koll, D.N. Sedman / Acta Materala 56 (2008) of <10 at.% [21,22]; Fne et al. [83] recently dscuss the potental sources of the dfferences between atom-probe and SANS measurements n bnary Fe Cu alloys. Furthermore, recent frst-prncples calculatons by Lu et al. [31] predct that n a bnary Fe Cu alloy, bcc precptates wth Cu concentratons greater than 50% are mechancally unstable at 0 K. Although, strctly speakng, the argument of Ref. [31,83] only apples to bnary alloys, we also measure a Cu concentraton rangng between 44.6 ± 3.1 and 53.5 ± 2.3 at.% when the precptates are <2 nm n radus. Ths s consstent wth earler observatons for bnary [8] and multcomponent [10,37] Fe Cu-based alloys. When the precptates are <2 nm n radus they also contan quanttes of N, Al, Mn and S, whch s consstent wth earler atom-probe measurements of model ternary Fe Cu N [9,10] and multcomponent [37] Fe Cu-based alloys. The recent phase-feld smulatons of model quaternary Fe Cu N Mn alloys also llustrate Cu-rch precptates contanng N and Mn at early (non-dmensonless tme) stages of phase decomposton [28,29]. Furthermore, the EXAFS results of Pzzn et al. [25] demonstrate that the precptates retan a bcc structure longer n a model ternary alloy contanng N, when compared to the bnary Fe Cu alloy. Therefore, the presence of Fe and other alloyng elements may be mechancally stablzng the bcc precptates. In the over-aged condton the precptate cores are enrched n Cu (97.1 ± 0.8 at.%), achevng an almost elemental concentraton at 1024 h. Ths s consstent wth earler atom-probe and SANS nvestgatons of both bnary [8,21,84] and ternary [9,10] Fe Cu-based alloys Formaton of N Al Mn shells Nckel, Al and Mn are found enrched at the precptate/ a-fe matrx heterophase nterfaces as early as 0.25 h of agng (Fg. 7 and Table 4). The observed nterfacal segregaton becomes more pronounced wth ncreasng agng tme, thereby formng a dstnct shell contanng N, Al and Mn adjacent to the Cu-rch precptate cores at 1024 h. Smlar segregaton has been observed expermentally and n smulatons n over-aged Fe Cu N [9,10,28,29,35], Fe Cu Mn [11,28,29,35], Fe Cu N Al [19] and Fe Cu N Mn [23,28,29] alloys. Ishem et al. [36 38] have reported comparable segregaton effects n smlar concentrated multcomponent Fe Cu steels. At 1024 h the shell has a N:Al:Mn stochometrc rato of 0.51:0.41:0.08 (Table 7), whch suggests a NAl-type B2 phase wth Mn substtutng for Al on the latter s sublattce. The concluson that t has a B2 structure cannot, however, be made by stochometry alone. We have recently performed synchrotron radaton studes at the Advanced Photon Source (APS), Argonne, IL, that demonstrate that ths shell does ndeed have the B2 structure [85]. Furthermore, our recent frst-prncples calculatons confrm that Mn substtutes at Al sublattce stes rather than at N sublattce stes [85]. The expermental results llustrate that the shells form by segregaton of N at the heterophase nterfaces wth a Table 7 Temporal evoluton of N:Al:Mn stochometrc rato wthn the heterophase nterfaces as a functon of agng tme N Al Mn 0.25 h h h h h h h smaller quantty of Mn and Al also present. The N and Mn concentraton peaks are ntally assocated wth each other, whereas the Al concentraton peak s found closer to the center of the precptate n possble assocaton wth Cu. Further agng results, n ncreasng nterfacal segregaton of all three elements, from the matrx and precptate cores, to the heterophase nterfaces; concomtantly the Al concentraton peak moves toward the N and Mn peaks. As seen from Table 7 the stochometrc rato of the heterophase nterfaces also evolves temporally, as Al ncreasngly parttons to the heterophase nterfaces. Our observatons are n agreement wth those of Ishem et al. [37] and Vaynman et al. [1] for the NuCu steel and recent phasefeld smulatons of model quaternary Fe Cu Mn N alloys [28,29]. The shells are not homogenous around the Cu-rch cores; as Fgs. 3a c and 4 demonstrate that the nterfacal segregaton s nonunform. The Cu-rch precptates and a-fe matrx nterfaces may be actng as nucleaton stes for a N 0.5 (Al 0.5 x Mn x ) phase. The observed compostonal heterogenety at the nterfaces may also be due to possble trajectory aberratons. Comparable heterogeneous nucleaton on Cu precptates, wth nonunform nterfacal segregaton, however, has also been reported n amorphous FINEMET alloys [86] and maragng stanless steels [87]. Our observatons, however, do not support the three-phase Cu-rch precptate core, fcc homogeneous shell, and a-fe matrx model proposed n Ref. [23]. The presence of a N 0.5 (Al 0.5 x Mn x ) phase at the precptate/ a-fe matrx heterophase nterfaces s consstent wth Cahn s local phase rule for heterophase nterfaces [88], whch predcts that a sngle phase at ths heterophase nterface s permtted thermodynamcally Equlbrum compostons Table 5 compares the equlbrum compostons of the Cu-rch precptate cores, heterophase nterfaces and a-fe matrx derved by three methodologes: (a) extrapolated to nfnte tme from APT data as a functon of (agng tme) 1/3 ; (b) as measured from specmens aged for 1024 h; and (c) as calculated from Thermo-Calc utlzng the Scentfc Group Thermodata Europe (SGTE) solutons database [89]. Reasonable agreement exsts for the frst two methodologes but sgnfcant dfferences exst, especally for the Cu-rch precptates, wth the Thermo-Calc calculatons. Thermo-Calc predcts the exstence of both bcc and