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Materials Science and Engineering A 528 (2010) 706–714

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

recipitation process along dislocations in Al–Cu–Mg alloy during artificial aging

ongqiang Fenga,∗, Yanqing Yanga, Bin Huanga, Ming Hana, Xian Luoa, Jigang Rub

School of Materials, Northwestern Polytechnical University, 127 West Youyi Road, Xi’an 710072, PR ChinaBeijing Institute of Aviation Materials, Beijing 100095, PR China

r t i c l e i n f o

rticle history:eceived 5 July 2010eceived in revised form0 September 2010ccepted 23 September 2010

a b s t r a c t

The heterogeneous nucleation and growth of precipitates at dislocations in Al–Cu–Mg alloy were inves-tigated by examining sample aged at 195 ◦C for various times from 10 min to 9 h. High resolutiontransmission electron microscopy (HRTEM) observations and electron energy loss spectrometers (EELS)analysis show that the precipitation sequence of the S (Al2CuMg) phase along dislocations should beSSS → GPB zones → S (Type I) → S (Type I) + S (Type II). Type II S precipitate can nucleate and grow sepa-

eywords:l–Cu–Mg alloysislocationsrecipitation sequence

rately at some dislocation parts which are unfavorable for Type I S precipitate formation, or can transitfrom Type I S precipitate during lateral growth along dislocations by means of continuous or steppedchanging in lattice orientations. Both of these modes are accomplished under the action of dislocationstrain field. The deviation extent of Type II S precipitate is considered not changing with time but closelyrelated with the interactions between the transformation strain during nucleation and the dislocation

rientation relationship strain field.

. Introduction

Al–Cu–Mg alloy is one of the most commonly used age-ardening aluminum alloys for aircraft applications due to itsxcellent properties of strength, fracture toughness and fatiguend corrosion resistance [1–3]. The homogeneous precipitation of(Al2CuMg) phase, together with its precursors S′ ′ and S′, is con-

idered as the main reason for aging hardening [2]. Besides, theeterogeneous formation of precipitates along dislocations waslso revealed to have effect on the hardness of Al–Cu–Mg alloyuring early aging [4]. Generally, the precipitation of S phase alongislocations will experience the following precipitation sequence4]:

SS (supersaturated solid solution) → S′′ → S′/S (SQ1)

he S phase, which has a face-centered orthorhombic structureith the lattice parameters a = 0.400 nm, b = 0.923 nm, c = 0.714 nm

5], always forms on {0 1 2}Al habit planes and grows mainlylong 〈1 0 0〉Al directions [1,2,6]. Concerning the S′ ′ phase, or calledPB2/GPBII [7,8] (GPB stands for Guinier–Preston–Bagaryatsky

9]), it has been reported by many researchers using differential

canning calorimeter (DSC) [10,11] and high resolution transmis-ion electron microscopy (HRTEM) [4,11,12]. However, its crystaltructure, chemical composition and orientation relationship withl matrix are still not very clear. For the S′ phase, although DSC

∗ Corresponding author. Tel.: +86 29 88460499; fax: +86 29 88460499.E-mail address: qiang733@163.com (Z. Feng).

921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2010.09.069

© 2010 Elsevier B.V. All rights reserved.

results given by Ratchev et al. [4] may prove its existence, otherstudies show that the S′ phase has the same formation enthalpies[13] and almost identical structures observed by HRTEM [14] as itsequilibrium phase S, and it is difficult to distinguish S′ from S exceptthe misfit between precipitates and Al matrix [15,16]. Given Per-litz and Westgren model [5] as the actual and exact structure ofS phase (it has been proved by many other authors [17,18]), thenthe precipitate phase, which has the similar phase structures buta little different lattice parameters with the S precipitate, could bedefined as the S′ phase or the slightly distorted version of the Sphase [1,7,13].

Based on the research in Al–2.5Cu–1.2Mg (wt%) alloy, Wilsonand Partridge [6] suggested that the S′ phase is not preceded by S′ ′

and it nucleates heterogeneously on dislocations, which can releasepartially the misfit between the precipitates and the matrix. Thenthe precipitation sequence should be

SSS → SSS → S′/S (SQ2)

Using Coincidence Doppler Broadening (CDB) technique, Nagaiet al. [19] revealed that vacancy–Mg–Cu complexes form along dis-locations in Al–1.3%Mg–1.7%Cu (at.%) alloy during aging at 150 ◦Cfor 1 min. Wang and Starink [1] considered that the S′ phase formsmostly on dislocations and Cu–Mg solute clusters during agingin Al–Cu–Mg based alloys, then another precipitation sequence

should be

SSS → Cu − Mg clusters → S′/S (SQ3)

Cu–Mg clusters, which own sub-nanometer size and have spe-cial ‘cluster hardening’ effect on Al matrix, were observed after

Z. Feng et al. / Materials Science and Engineering A 528 (2010) 706–714 707

ure of

s(Gpaodgdbta

c[drtweto[ttBi

Fig. 1. (a) TEM image observed from [1 0 0]Al showing the microstruct

hort time aging by means of atom probe field ion microscopyAPFIM) [20]. Besides, another similar precipitate structure, i.e.PB zone, was also observed in Al matrix during early agingrocess. According to Silcock [9], GPB zones are small cylinderslong 〈1 0 0〉Al directions with a diameter of 1–2 nm and lengthf 4–8 nm. Kovarik et al. [8] suggested that GPB zones are smalliscrete ordered regions, which form through a process analo-ous to spinodal decomposition/ordering and do not have clearlyefined interfaces with Al matrix. However, GPB zones have noteen reported to form along dislocations during aging process, andhe relationship among Cu–Mg clusters, GPB zones, S precipitatend its precursors are still not clear.

The classic orientation relationship (OR) between S pre-ipitate and Al matrix is [1 0 0]Al//[1 0 0]S′/S, [0 2 1̄]Al//[0 1 0]S′/S,0 1 2]Al//[0 0 1]S′/S [1,6,15]. On the other hand, S precipitate with ORifferent from the classic one was also reported. Radmilovic et al. [2]evealed an S precipitate which differs by a rotation of about 5◦ fromhe classic OR, and termed this S precipitate as Type II S precipitate,hile that with the classic OR was Type I S precipitate. Majimel

t al. [21] observed that the orientation differences between thewo types of S precipitate varied from particle to particle, mostf them falling into the degree range of about 4–6◦. Kovarik et al.12] confirmed the existence of Type II S precipitate but referred

o the Type II S precipitate as S′ ′ phase due to the orientation rela-ionship satisfied the OR proposed by Li and Yan [22] for S′ ′ phase.ased on the measurement of the orientation of over three hundred

ndividual particles of the S phase, Winkelman et al. [23] revealed

Fig. 2. Zero loss image (a) and element mappings for Cu (b) and Mg (c) imaged b

2024 alloy aged at 195 ◦C for 10 min. (b) Corresponding SAD pattern.

that all orientations of S phase may differ by a rotation of up to ∼7◦

about the axes [1 0 0]S//[1 0 0]Al between two rational limits definedby the relationships: (0 0 1)S//(0 2 1)Al, [1 0 0]S//[1 0 0]Al (ROR1) and(0 2̄ 1)S//(0 1 4)Al, [1 0 0]S//[1 0 0]Al (ROR2) and individual S particlesmay adopt a continuous or near-continuous range of orientationsbetween these limits, accompanied by accommodating changes ininterface orientation and lattice parameters. The driving force forthe formation of Type II S precipitate was explained to result froman invariant line transformation strain [2], competitions betweentwo coherent (0 0 1)S//(0 2 1)Al and (0 2 1)S//(0 1 4)Al interfaces [12],or the interactions between the transformation strain during nucle-ation and the dislocation strain field [23]. The formation of Type IIS precipitate was found to be strongly dependent on the conditionsof solution treatment and cold work [3], or related with the localvariations in the environment for nucleation of the S phase such aspre-existing lattice defects and/or solute or defect/solute clusters[23].

Many researchers [1,4,6,7,19,24] have referred to the precipita-tion of the S phase along dislocations, but few of them paid muchattention to its precipitation sequence and distribution character-istics. Meantime, the formation process and mechanism of TypeII S precipitate along dislocations are still not clear. In this paper,therefore, the microstructure evolution during heterogeneous pre-

cipitation of the S phase was investigated and the formationmechanisms of Type II S precipitates along dislocations wasdiscussed.

y EELS showing no enrichment of Cu and Mg elements along dislocations.

708 Z. Feng et al. / Materials Science and Engineering A 528 (2010) 706–714

Table 1Nominal chemical composition of 2024 alloy (wt%).

2

it242qdebebGeaa

3

3

1Tphsp1tiEhcfiSbitiscftGztl

3

vhcdpt

Element Cu Mg Mn Si Fe Al

Amount 4.2 1.5 0.6 0.5 0.5 Bal.

. Experiment procedure

The 2024 alloy was chosen for the investigation, and the chem-cal composition of this alloy is given in Table 1. The cast ingot ofhe alloy was homogenized at 460 ◦C for 16 h and hot rolled to amm thin sheet, then solution treated at temperature 505 ◦C for5 min, water quenching, followed by aging at 195 ◦C for 10 min,0 min, 1 h, 4 h and 9 h, respectively. TEM samples were subse-uently prepared by mechanical grinding and punching to 3 mmisks in diameter. The disks were finally thinned using a twin jetlectropolishing with a solution of 30% nitric acid and 70% methanolelow −25 ◦C at 15 V. TEM and HRTEM observations as well aslectron energy loss spectrometry (EELS) analysis were carried outy means of a transmission electron microscope of Tecnai F302 equipped with an energy filter of Gatan Tridiem and an X-raynergy dispersive spectrometer (EDS). All the HRTEM observationsre performed from 〈1 0 0〉Al direction, since they are the most suit-ble conditions for GPB zones, S′ ′, S′ and S precipitates visualization.

. Results

.1. Aging at 195 ◦C for 10 min

Fig. 1 shows the microstructure of the sample aged at 195 ◦C for0 min and corresponding selected area diffraction (SAD) pattern.here are no clear precipitates except some rod-like Al20Cu2Mn3articles [1,25]. Lots of dislocations, especially dislocation loops andelices, can be observed in Al matrix (Fig. 1a). SAD pattern (Fig. 1b)hows no obvious diffraction characteristics for the S phase and itsrecursors S′ ′ and S′, but weak crosses (see arrows), centered at/2 {0 2 2}Al and formed by diffraction streaks along 〈0 1 0〉Al direc-ions, which are indication of GPB zones [11,13,26–28]. Detailednvestigations are focused on the dislocations and adjacent regions.ELS element mappings (Fig. 2) and EDS analysis results (not shownere) indicate no enrichment of Cu and Mg elements along dislo-ations. HRTEM observations on several dislocation parts cannotnd the atomic arrangement feature of GPB zones as suggested byilcock [9]. However, the dark contrast regions, probably causedy strain field contrast, are observed along dislocation as shown

n Fig. 3. The corresponding fast Fourier transformation (FFT) spec-rums obtained from these regions show the same characteristics,.e. besides the Al diffraction spots, there are sharp diffraction spotsituated at the position of 0 1 1Al (inset). (Note that the whiterosses, located at the transmission and diffraction spots, are arti-act due to the FFT transformation.) Kovarik et al. [8] suggestedhat these sharp 1 1 0Al reflections indicate the existence of thePB zones (or small discrete ordered regions). However, these GPBones do not have clearly defined interfaces with Al matrix, thushey are hard to be detected although may have formed along dis-ocations.

.2. Aging at 195 ◦C for 20 min and 1 h

The microstructure of the sample aged at 195 ◦C for 20 min isery similar to that for 10 min. Many dislocations, especially the

elical ones, are distributed in Al matrix (Fig. 4a). However, theontrast of these dislocations in bright field image seems mucharker than that of its previous state. In the corresponding SADattern (Fig. 4b), no diffraction characteristics proving the exis-ence of the intermediate S′ ′ phase as considered by Cuisiat et

Fig. 3. HRTEM image observed from [1 0 0]Al showing regions near the dislocationline. The inset is the corresponding FFT spectrum of the white rectangular region.

al. [29] can be discerned. Nonetheless, it should be noted herethat extra diffraction spots, arrayed approximately along [0 1 2]Aldirection (see arrow), can be obviously observed. This diffractioncharacteristic is very similar to that of S precipitate proposed byWang and Starink [1] as shown in Fig. 4c. These discrete diffrac-tion spots, with some of them deviating from [0 1 2]Al direction,might be caused by the S′/S precipitate formed during initial agingwith the characteristics of minor size, different lattice parametersand apparent periodic structure in one dimension. This specula-tion can be further testified by the HRTEM observation on themicrostructure of this state. Fig. 5a illustrates the HRTEM imageof a dislocation segment. There is no obvious lattice distortion inAl matrix but along the dislocation. Moreover, the lattice distortionat some parts (labeled by arrows) seems larger than those at otherparts. Detailed observation is given to these parts and the enlargedimage of the white rectangular region is shown in Fig. 5b. The Allattice in this region bears high distortion and shows dark contrast.Moreover, equidistant white atoms which deviate from their nor-mal Al lattice sites can be found arraying orderly on rows of (0 2 1)Alplanes. The row and atomic distances are measured approximatelyto be 0.712 nm and 0.920 nm, which are close to the interplanarspacings of (0 0 1)S′/S and (0 1 0)S′/S as ever proposed S′/S structure[2,4,5,16,23,30] (Table 2). Meantime, in the corresponding FFT spec-trum (Fig. 5c), diffraction streaks (see arrow) can be observed along[0 2 1]Al direction. All these features fully demonstrate that the S′/Sphase has formed at this part of dislocation. Since the S′ phase couldbe defined as the slightly distorted version of the S phase [1,7,13],the term ‘S phase’ is used hereafter instead of ‘S′ phase’ to describethe precipitates formed at dislocations.

As shown in Fig. 5b, during early aging, the S precipitate formsas a thin-layer on (0 2 1)Al plane along dislocation. This charac-teristic will be observed more obviously as prolonging the agingtime to 1 h (Fig. 6). The S precipitates distribute continuously alongthe dislocation and take on stepped and/or corrugated appearancesfrom part to part. The insets in Fig. 6 are magnified HRTEM imagesand corresponding FFT spectrums of the two rectangular regions.It can be found that both of these two S precipitates keep flaton {0 1 2}Al (or (0 1 0)S) habit planes and show thin-layer charac-

teristic in their cross-sections. According to Bagaryatsky [15], theclassic OR between S precipitate and Al matrix is [1 0 0]Al//[1 0 0]S,[0 2 1̄]Al//[0 1 0]S, [0 1 2]Al//[0 0 1]S and the angle between [0 0 1]Sand [0 0 1]Al (or termed as �OR) is about 26.56◦. Table 3 shows four

Z. Feng et al. / Materials Science and Engineering A 528 (2010) 706–714 709

Fig. 4. (a) TEM image observed from [1 0 0]Al showing the microstructure of 2024 alloy aged at 195 ◦C for 20 min. (b) Corresponding SAD pattern. (c) The simulated SADpattern of S precipitate and Al matrix [1].

Fig. 5. (a) HRTEM image observed from [1 0 0]Al showing the nucleation of S precipitates at discrete parts of dislocation (see arrows). (b) The magnification of the whiterectangular region in (a) and the corresponding FFT spectrum (c).

Table 2Proposed structures for S′/S phase and the corresponding interplanar spacings of (0 1 0)S′ /S (or 2(0 1 0)S′ /S) and (0 0 1)S′ /S.

Precipitate Crystal structures Interplanar spacings of (0 1 0)S′ /S and (0 0 1)S′ /S References

S′ Orthorhombic, Cmcm, a = 0.405 nm, b = 0.906 nm, c = 0.724 nm. 0.9060 nm and 0.724 nm [16]Orthorhombic, Cmcm, a = 0.404 nm, b = 0.925 nm, c = 0.718 nm. 0.9250 nm and 0.718 nm [4]

S Orthorhombic, Cmcm, a = 0.400 nm, b = 0.923 nm, c = 0.714 nm. 0.9230 nm and 0.714 nm [5]Orthorhombic, Cmcm, a = 0.401 nm, b = 0.925 nm, c = 0.715 nm. 0.9250 nm and 0.715 nm [16]Orthorhombic, Pmm2, a = 0.400 nm, b = 0.461 nm, c = 0.718 nm. 0.9220 nm (2(0 1 0)S′ /S) and 0.718 nm [30]Orthorhombic, Cmcm, a = 0.403 nm, b = 0.930 nm, c = 0.708 nm. 0.9300 nm and 0.708 nm [2]Orthorhombic, Cmcm, a = 0.400 nm, b = 0.9056–0.9321 nm, c = 0.7040–0.7245 nm. 0.9056–0.9321 nm and 0.7040–0.7245 nm [23]

Table 3Four equivalent orientation relationships between S and Al matrix along [1 0 0]Al//[1 0 0]S [1].

Variant Equivalent orientation relationship Two directions to determine �OR

1 [1 0 0]Al//[1 0 0]S, [0 2 1̄]Al//[0 1 0]S, [0 1 2]Al//[0 0 1]S [0 0 1]S &[0 0 1]Al

2 [1̄ 0 0]Al//[1 0 0]S, [0 2 1]Al//[0 1 0]S, [0 1 2̄]Al//[0 0 1]S [0 0 1̄]S & [0 0 1]Al

3 [1 0 0]Al//[1 0 0]S, [0 1̄ 2̄]Al//[0 1 0]S, [0 2 1̄]Al//[0 0 1]S [0 0 1]S &[0 1 0]Al

4 [1̄ 0 0]Al//[1 0 0]S, [0 1 2̄]Al//[0 1 0]S, [0 2̄ 1̄]Al//[0 0 1]S [0 0 1̄]S & [0 1 0]Al

710 Z. Feng et al. / Materials Science and Engineering A 528 (2010) 706–714

F ipitatec

edit3nnctodacds

3

catcacff

ig. 6. HRTEM image observed from [1 0 0]Al showing the development of S precorresponding FFT spectrums.

quivalent variants of S precipitate observed along [1 0 0]Al//[1 0 0]Sirection [1] and the corresponding two directions for determin-

ng �OR. Accordingly, S precipitate in Fig. 5b could be attributedo Variant 4, while those in Fig. 6 are attributed to Variant 2 and. The �ORs for these S precipitates are all approximate to 26.56◦,o apparent OR changes could be detected for them. However,ot all of these S precipitates form on {0 1 2}Al habit planes whenompared with the closest (0 1̄ 2)Al or (0 2̄ 1)Al habit plane illus-rated as white lines shown in Fig. 6. At some parts such as thenes labeled by arrows the orientation of S precipitates seemseviating from (0 1̄ 2)Al or (0 2̄ 1)Al habit plane of about smallngles. Unfortunately, these deviations cannot be observed morelearly at the atomic level in higher magnification observationsue to the minor size in the precipitate cross-sections in the initialtages.

.3. Aging at 195 ◦C for 4 h and 9 h

With increasing of aging time to 4 h, the average size of S pre-ipitates along dislocations became larger than that during earlyging process. Fig. 7 shows the HRTEM image obtained in closeo [1 0 0]Al at a dislocation segment. S precipitates A–D arrange

ontinuously along the dislocation line and the corresponding FFTnalyses are shown bellow the HRTEM image. The �ORs for S pre-ipitates A and C (Variant 4) are very close to 26.56◦, while thoseor S precipitates B and D (Variant 1) differed about 4.2◦ and 6.4◦

rom this angle, respectively. As ORs differs from the classic one,

s along dislocation. The insets are magnifications of the rectangular regions and

S precipitate tends to be equiaxed in the cross section, accom-panied by the morphology changing from lath to rod. Accordingto the previous researchers [2,3,12,21,23], S precipitates B andD can be termed as Type II S precipitates. Attentions should bepaid to the sites for Type II S precipitate formation, which arevery similar to those labeled by arrows as shown in Fig. 6. TheOR deviations are both observed at these sites and indicate that Sprecipitate with different OR may have formed during early agingprocess.

Fig. 8 shows part of dislocation decorated by S precipitates afteraging at 195 ◦C for 9 h. S precipitates of the same S variant contin-uously formed at this dislocation part. However, the ORs for theseS precipitates vary from part to part, i.e. S precipitate near regionA can be termed as Type I S precipitate since (0 0 1)S is parallel to(0 2 1)Al, while that near region C can be termed as Type II S precip-itate because this precipitate possesses an different OR (about 6.5◦

deviation from the classic one) and (0 0 1)S is very close to (0 3 1)Al(about 1.63◦ offset). The orientation transition from Type I to TypeII is accomplished by two or more stepped changes in lattice orien-tation within narrow regions such as the white rectangular regionslabeled by ‘OR transition region 1 and 2′. For instance, the orien-tation of region B rotates clockwise about 2.6◦ from that of region

A, while the orientation of region C further rotates about 3.9◦ fromthat of region B. Within these regions two orientations of S latticescoexist and/or lie out of their normal sites. All these characteristicsmentioned above may indicate another formation mechanism ofType II S precipitate.

Z. Feng et al. / Materials Science and Engineering A 528 (2010) 706–714 711

F ates alA

4

4

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ig. 7. HRTEM image observed from [1 0 0]Al showing the configuration of S precipit–D.

. Discussion

.1. Precipitation sequence at dislocations

Since the diffusion of vacancies and solute atoms along ‘dislo-ation pipes’ is much easier than that in the matrix [31], and the Alattice around dislocations often bear high distortions and possessarge elastic strain energy, the dislocations are usually preferredor new precipitate nucleation and growth [32–34]. According tohe previous researchers [1,4,6], the precipitation sequence of thephase at dislocations seems different with that in Al matrix. In

onsideration of the characteristics of heterogeneous precipitation,his precipitation sequence cannot be studied by means of thermalnalysis (such as DSC) but by analyses on the changes in HRTEMtructures, diffraction characteristics and chemical compositions.n the present work, these analyses are given to the microstructuref 2024 alloy aged at 195 ◦C for periods from 10 min to 9 h. Afterging for 10 min, no obvious precipitates except many dislocationsan be observed in Al matrix. The clustering of Cu and Mg elementslong dislocations is not enough to be detected by EELS (and EDS)Fig. 2). However, weak crosses centered at 0 1 1Al in the SAD pat-ern (Fig. 1b) as well as sharp diffraction spots at 0 1 1Al in the FFTpectrum corresponding to regions near the dislocation (Fig. 3),re proved to be the indication of GPB zones [8,11,13,26–28].hen aging time increases to 20 min, extra diffraction spots can be

bserved approximately along [0 1 2]Al direction (Fig. 4b). More-ver, white atoms line up in the small regions near dislocationnd diffraction streaks form along [0 1 2]Al directions in the corre-

ponding FFT spectrums (Fig. 5c). All the characteristics of electroniffraction and atomic arrangement are not in accord with the cor-esponding features of the S′ ′ phase given by Ratchev et al. [4]nd Cuisiat et al. [29], but consistent with that of the S phase2,4,5,16,23,30]. The initial formed S precipitate is observed on

ong part of dislocation, the FFT spectrums below are corresponding to S precipitates

(0 2 1)Al habit plane with only a few atoms thick and take on thin-layer morphology in its cross sections (Fig. 5b). As aging proceeds,more and more S precipitates became to form along dislocations(Fig. 6). Most of these S precipitates are observed on {0 1 2}Al habitplanes and consistent with the classic OR (Type I S precipitate)while S precipitates with different ORs are also observed at somedislocation parts labeled by arrows (Fig. 6). Further HRTEM obser-vations on the similar dislocation parts in the sample aged for 4 hprove the formation of Type II S precipitate, the ORs of which varyfrom particle to particle (Fig. 7). Upon further aging, these two typesof S precipitates grow and coexist up to 9 h (Fig. 8). Similarly, thecoexistence of different orientations of S phase was also observedin Al–0.2Cu–1.7Mg and Al–0.8Cu–1.7Mg (wt%) alloys aged at 150 ◦Cfor times in the range of 1–2000 h [23]. Accordingly, the precipita-tion sequence at dislocations during aging process might be

SSS → GPB zones → S (Type I) → S (Type I) + S (Type II)

4.2. The formation mechanism of Type II S precipitate

The orientation deviations from the classic OR proposedby Bagaryatsky [15] were reported by many researchers[2,3,12,21,23,35,36] and several theories were proposed to explainthe variation in ORs or the origin of Type II S precipitate. Kovarik etal. [12] suggested that Type II S precipitate forms due to the com-petition between two coherent interfaces, i.e. (0 0 1)S//(0 2 1)Al and(0 2 1)S//(0 1 4)Al, and will transform into Type I S precipitate uponfurther aging due to its metastable nature. On the contrary, based

on the differential scanning calorimetry (DSC) analyses at a heat-ing rate of 10 ◦C/min and the equivalent thermal calculation of theprevious reported formation conditions of the Type II S precipitate,Wang and Starink [3] concluded that Type II S precipitate can formfrom Type I S precipitate due to its superior stability in the thermo-

712 Z. Feng et al. / Materials Science and Engineering A 528 (2010) 706–714

Table 4Factors influencing S precipitate formation and the corresponding results.

Situation I Situation II Situation III Situation IV

FavorLowTypeDiffic

dawthtvmm[stf

ts[cart

otbaeai

Ff

Dislocation orientation FavorableStrain field energy HighS precipitate type Type IPrecipitation condition Easier

ynamic sense and result in an exothermic peak between 280 ◦Cnd 350 ◦C. The transformation from S precipitate Type I to Type IIas considered to be strongly dependent on the condition of solu-

ion treatment and cold work. After statistical analysis of over threeundred individual S precipitates, Winkelman et al. [23] suggestedhat the changes in the orientation of S precipitates will cause theariation in interface orientations as well as lattice parameters anday be related with the critical interactions between the transfor-ation strains during nucleation and the local elastic strain field

23]. Besides, an invariant line transformation strain was also con-idered as the driving force for Type II S precipitate [2]. However,his theory was found unable to explain the spread in the OR andor S precipitate with corrugated interfaces [3,12].

When the heterogeneous precipitation occurs along dislocation,he resistance to precipitate nucleation and growth will be loweredince extra elastic energy can be provided by dislocation strain field6]. Thus, the distribution characteristics of dislocation strain fieldan further affect the variant selection, nucleation and growth char-cteristics, and the final morphology of S precipitate. Two factorselated to the dislocation part are considered to be concerned withhe nucleation of S precipitate with or without OR changes:

(1) The strain field energy. The dislocation in Al matrix is moreften a mixed one rather than a pure edge or screw one. Generally,he total strain energy for this mixed dislocation can be representedy the elastic energy since the core energy is estimated to be only

small fraction of the elastic energy. According to the continuous

lastic media model for dislocation [33,34], the elastic distortionround an infinitely-long, straight dislocation can be representedn terms of a cylinder of elastic material. The elastic energy stored

ig. 8. HRTEM image observed from [1 0 0]Al showing the stepped changes in ORsor S precipitates near region A, B and C.

able Unfavorable UnfavorableHigh Low

I Type II Type IIult Easy More difficult

in the cylinder per unit length of a mixed dislocation is

Eel (mixed) = Gb2(1 − � cos2 �)4�(1 − �)

ln(

R

r0

)(1)

where G is shear modulus, r0 is the dislocation core radius, R is theouter radius of the cylindrical strain field, v is Poisson’s ratio, b isBurgers vector of the mixed dislocation, � is the angle between theBurgers vectors of the mixed dislocation and its screw component.For a dislocation with certain Burgers vector, the value of � alwaysvaries with the dislocation orientation, thus the fluctuation of elas-tic energy does exist along dislocation. The dislocation part, whichown high strain field energy, will be more easily for S precipitatenucleation and growth.

(2) The dislocation orientation. The changes in dislocation ori-entation will cause the variation of distribution characteristics ofthe dislocation strain field, and will further affect the variant selec-tion and the nucleation and growth characteristics of S precipitate.According to Wilson and Partridge [6], two factors may be relatedwith the variant selection of S precipitate along dislocation: (i) thelattice misfits between S precipitate and Al matrix, which are about−1.88% and +1.49% along bS and cS axes and producing tensile andcompressive stress in Al matrix, respectively, and (ii) the dislocationline tension, which restrain the dislocation to keep the minimumlength. When the orientation of a dislocation part is parallel tobS axis (or (0 0 1)S plane) of a certain S variant, the tensile stressstemming from the lattice misfit further prompt the effect of thedislocation line tension and ensure the dislocation with the mini-mum length. Meantime, the compressive stress along cS axis couldbe reduced by the dislocation lying parallel to the precipitate oneither the same or different surface [6]. Accordingly, the disloca-tion part, which is parallel or near parallel to (0 0 1)S plane, canbe termed as ‘orientation favorable’ site for Type I S precipitatenucleation, whereas those deviating from (0 0 1)S plane of about asmall angle, will not be preferred for Type I S precipitate any more.However, under the action of the dislocation strain field, a minuteadjustment may occur in OR during the nucleation of S precipitate,such as S precipitates labeled by arrows in Fig. 6, in order to obtainthe maximum coupling of the strain field induced by lattice misfitand dislocation. Attentions should be paid here that this OR adjust-ment for S precipitate is limited for the small orientation differencein minor regions. For large difference the dislocation part might benot appropriate for this S variant but for another. Thus the orienta-tion difference between dislocations and (0 0 1)S plane of a certainS variant is suggested to be the predominant factor for Type II S pre-cipitate formation. As the dislocation orientation varies from partto part, the OR adjustments for S precipitates formed along disloca-tions change from particle to particle. As a result, the deviation forType II S precipitate from the classic OR tend to be a variable withina small angle range (about 7◦) rather than a constant [21,23].

When the heterogeneous precipitation of the S phase occursat a certain dislocation part, with consideration of the two factorsmentioned above, i.e. the strain field energy and the dislocation ori-

entation, the formation of S precipitate tend to present four kindsof situations as shown in Table 4. For dislocation part with favor-able orientation and high strain field energy, Type I S precipitatewill form more easily such as the part shown in Fig. 5b, while thatwith unfavorable orientation but high strain field energy, Type II

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Z. Feng et al. / Materials Science a

precipitate will prefer to form such as the ones coexisting withype I S precipitate and labeled by arrows in Fig. 6.

Usually, Type II S precipitate could not be observed earlier thanype I S precipitate and was considered as the thermal stable vari-nt of Type I S precipitate upon further heat treatment [3]. However,ufficient observations have been made to confirm the coexistencef different orientation of S phase in Al–Cu–Mg alloys aged at 150 ◦Cor times in the range 1–2000 h [23]. The present work also showshat Type II S precipitate can coexist with Type I S precipitatet 195 ◦C for a period up to 9 h. These phenomena may be con-erned with the difference in nucleation conditions between twoypes of S precipitates. Generally, Type I S precipitates are boundy coherent planar interfaces parallel to (0 1 2)Al//(0 0 1)S, whileype II S precipitates have macroscopically planar facets parallelo (0 2 1)S//(0 1 4)Al comprising periodic steps, the spacing of whichecrease with the increasing lattice rotation [23]. As the effectivereas for the stepped interfaces are always larger than those for thelanar interfaces, for a certain precipitate volume, the total inter-acial energy for Type II S precipitate will be higher than that forype I S precipitate. Accordingly, the nucleation energy required forype II S precipitate seems larger than that for Type I S precipitate.ince the deviation for Type II S precipitate from the classic OR areithin a small angle range (about 7◦) [23] and most of them fall intorange of about 4–6◦ [21], a higher energy barrier may exist before

he formation of Type II S precipitate. During artificial aging, Type Ind Type II S precipitate mainly form in two different stages, whichay result in two exothermic peaks in DSC curve as reported byang and Starink [3]. Due to the differences in chemical composi-

ions and nucleation energies, some Type II S precipitates may alsoorm in the first stage while some Type I S precipitates in the secondtage. The transition between these two stages seems continuousnd may be not a process of dissolution and re-precipitation sinceo obvious endothermic peak can be observed between the twoxothermic peaks [3]. Thus during aging process the coexistence ofype I and Type II S precipitate can be observed for a long time.

As shown in Fig. 8, S precipitates formed at this dislocationart belong to the same S variant, and the orientation transitionrom region A (Type I) to region C (Type II) is accomplished bywo stepped changes in lattice orientation within two OR transitionegions 1 and 2. Similar phenomenon was also observed by Winkel-an et al. [23] and the change in orientation of S phase is about 4.2◦

nd accomplished within a narrow region in which lattice continu-ty is apparently continuously preserved and there is no evidence ofdiscrete interface between S phase domains. As discussed above,

or dislocation part with favorable orientation, it will be preferredor Type I S precipitate nucleation. Generally, Type I S precipitateas a characteristic lath shape and tends to be flat on (0 0 1)S planeue to the almost identical atomic arrangement of {0 1 2}Al and0 0 1)S plane which meet face to face [2]. As the dislocation orien-ations often changes from part to part, two kinds of results may benduced, i.e. (i) the most favorable variant changes and the S pre-ipitate of a new variant need to nucleate separately, and (ii) theost favorable variant dose not change but the lateral growth along

0 1 0]S is affected by the variation of dislocation strain field, i.e. theR for S precipitate gradually deviates from the previous one ofsmall angle and towards the dislocation orientation. For the for-er, a comparative large variation in dislocation orientation should

ccur within a small scale and much more energy is needed for newrecipitate nucleation. However, for the latter, the OR deviationsill be accomplished more easily within narrow regions throughcontinuous or stepped process since larger coincidence site lat-

ice (CSL) will be acquired between the adjacent S lattices due toittle changes in the lattice orientation and low resistance in lateralrowth. As the lattice deviation of S precipitate increases, the nucle-tion condition of the growth frontier, including lattice match anducleation energy, will gradually become unfavorable and the pref-

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gineering A 528 (2010) 706–714 713

erential growth direction will turn to the dislocation orientation,accompanied by accommodating changes in interface orientationas shown in Fig. 8. Accordingly, under the action of the dislocationstrain field, Type II S precipitate may also form through the lat-eral growth with corresponding OR deviations in order to get morenucleation energy from the dislocation strain field.

5. Conclusions

The precipitation process along dislocations in Al–Cu–Mg alloywas investigated in detail by TEM and HRTEM. The results showthat the heterogeneous precipitation of the S phase experiences adifferent sequence with the previous ones. GPB zones instead ofCu–Mg clusters are detected near dislocations in the initial stage ofthe aging process. As aging proceeds, no indication of the S′ ′ phasecould be discerned by SAD and HRTEM. However, S precipitates,which are only a few atoms thick, tend to form on {0 1 2}Al habitplanes. As more S precipitates form along dislocations, S precip-itates with different ORs from the classic one could be found atsome dislocation parts. Further HRTEM observations focusing onthe similar parts prove that these S precipitates are Type II S pre-cipitates, and can coexist with Type I S precipitate up to 9 h. Thusduring this aging process the precipitation sequence at dislocationsis suggested as SSS → GPB zones → S (Type I) → S (Type I) + S (TypeII).

Two formation modes of Type II S precipitate are concluded inthis work. First, Type II S precipitate can nucleate separately at somedislocation parts unfavorable for Type I S precipitate formation.Second, Type II S precipitate can be formed during lateral growthalong dislocations by means of continuous or stepped changing inlattice orientations. Both of these modes are accomplished underthe action of dislocation strain field. The deviation extent of TypeII S precipitate is considered not changing with time but closelyrelated with the interactions between the transformation strainduring nucleation and the dislocation strain field.

Acknowledgements

Thanks are given to the 111 Project (B08040) of China for sup-porting the research, One of the authors (ZQF) is grateful to Prof.J.G. Zhang and Dr. R.J. Zhang for many valuable discussions on someconcepts in dislocation theory and phase transformations. Thanksare also given to senior engineer L.F. Li and X.C. Li for TEM samplepreparation.

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