Stress-induced thickening of Ω phase in Al–Cu–Mg alloys containingvarious Ag additions

Song Bai a,b, Zhiyi Liu a,b,n, Xuanwei Zhou a,b, Peng Xia a,b, Meng Liu a,b

a Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Central South University, Changsha 410083,People's Republic of Chinab School of Material Science and Engineering, Central South University, Changsha 410083, People's Republic of China

a r t i c l e i n f o

Article history:Received 8 August 2013Received in revised form17 September 2013Accepted 19 September 2013Available online 25 September 2013

Keywords:Al–Cu–Mg–Ag alloyΩ phaseStress agingThickening

a b s t r a c t

The thickening of Ω phase in Al–Cu–Mg alloys containing various bulk Ag contents during stress aging at200 1C with a tensile stress of 240 MPa was investigated by a combination of transmission electronmicroscopy (TEM), high resolution transmission electron microscopy (HRTEM) and atom probetomography (APT). TEM characterization confirmed preferred orientation of Ω phase in all stress-agedsamples. Corresponding quantitative TEM calculations revealed the thickening kinetics of Ω phase wassignificantly accelerated during stress aging as compared to that during stress-free aging at 200 1C.HRTEM analysis on the α/Ω interfacial structure confirmed that the applied tensile stress facilitated therapid nucleation of the growth ledge on the broad face of Ω phase, thereby resulting in the acceleratedplate thickening during stress aging at 200 1C. Meanwhile, quantitative TEM analysis highlighted thestress-induced thickening of Ω phase at 200 1C was affected by the bulk Ag content. This was consistentwith the HRTEM observation as the ledge nucleation was found to be suppressed with increasing Agaddition. Our APT analysis on different stress-aged samples further suggested the progressive enrich-ment of Ag atoms in the segregation layer helped to stabilize the interfacial structure and wasresponsible for the lowest nucleation rate of the ledge in 1.77Ag alloy as compared to that in 0.46Ag alloy.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

Trace addition of Ag was well known to accelerate the agehardening response of Al–Cu–Mg alloys by producing a fine anduniform distribution of Ω phase that nucleated and grew as thinhexagonal-shaped plates on the {111}α plates [1–3]. Compared tothe common strengthening precipitates (e.g. θ′, S′) in aluminumalloys, Ω phase exhibited excellent thickening resistance duringprolonged aging at elevated temperatures. Therefore, particularinterest focused on the origin of the high thickening resistance ofΩ phase [4–6]. The thickening of plate-like precipitates, e.g. θ′/θ inAl–Cu alloys [7,8], was generally explained by a ledge mechanism.Similarly, HRTEM investigation by Fonda et al. [5] suggested thatthe thickening of Ω phase was accomplished by the propagation ofthe large ledge containing a misfit-compensating dislocation.Later, Hutchinson et al. [6] developed this idea to clarify thenature of the thickening resistance of Ω phase. In terms of theunrelaxed misfit normal to the Ω plate, their atomic resolutionZ-contrast microscopy analysis confirmed that a lack of the growth

ledges produced by the strong vacancy misfits was responsible forthe excellent thickening resistance of Ω phase at 200 1C. Based onthe above studies, the thickening of Ω phase in Al–Cu–Mg–Agalloys was now considered to be dominated by a ledge mechanismduring conventional aging process.

Due to the combination of the attractive strength properties [9]and excellent creep performance [10], Al–Cu–Mg–Ag alloys wereused as promising materials for aerospace applications. Therefore,it was important to clarify the microstructural evolution of Al–Cu–Mg–Ag alloys, particularly in the thickening behavior of Ω phaseduring stress aging (similar to the real service condition). In somealloy systems, the preferred orientation of the strengtheningprecipitates with various morphologies was induced by an exter-nally applied stress. For example, previous studies confirmed anoticeable orientation effect on the alignment of θ′ phase in binaryAl–Cu alloys by stress aging [11,12]. In the case of Al–Cu–Mg–Agquaternary alloy, the remarkable effect of the applied stress on theΩ precipitation in Al–Cu–Mg–Ag alloys was also well investigatedin previous studies. For example, TEM study by Muraishi et al. [13]proposed that the stress effect on the precipitation of Ω phase waspredominant in its initial nucleation stage, implying stress agingfollowing the initial stress-free aging failed to cause the preferredprecipitation of Ω phase. Later, their work concerning the compe-titive precipitation between Ω and θ′ phases during two-step aging

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0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.msea.2013.09.065

n Corresponding author at: School of Material Science and Engineering,Central South University, Changsha 410083, People's Republic of China.Tel.: þ86 731 88836011; fax: þ86 731 88876692.

E-mail address: liuzhiyi@mail.csu.edu.cn (Z. Liu).

Materials Science & Engineering A 589 (2014) 89–96

(stress aging for 1 h followed by stress-free aging) further showedthe preferentially nucleated GP zones against {111}α precipitatesduring the initial stress aging and therefore assisted the followingnucleation and precipitation of θ′ phase [14]. The study of Skrotzkiet al. [15] suggested a critical stress must be exceeded to inducethe noticeable preferential orientation of Ω phase. Hargarter et al.[16] investigated the relationship between the preferentiallyaligned precipitates and plastic anisotropy in one stress-agedAl–Cu–Mg–Ag alloy. Their study revealed the similar preferentialalignment of Ω phase paralleled to the tensile stress of 120 MPaafter aging at 160 1C for 50 h. According to the published works[14–16], the preferred precipitation of Ω phase by stress agingattracted too much attention. The thickening of Ω phase at a hightemperature of 200 1C with a high tensile stress, however, waspoorly investigated. Although the accelerated thickening kineticsof Ω phase during stress aging was implied in Ref. [15], it was stilldifficult to confirm the exact connection between the appliedtensile stress and the increase of the plate thickness due to HRTEMstudy on the α/Ω interfacial structure in stress-aged alloys waslack. Thus, one aspect, or an as-yet ambiguous question, was theintrinsic mechanism that contributed to the accelerated thickeningprocess of Ω phase during stress aging at certain temperatures,below which Ω phase exhibited excellent thickening resistance inconventional stress-free aging.

Till present there was little published information concerningthe thickening behavior of Ω phase during stress-aging in Al–Cu–Mg alloys containing various Ag additions. Thus, the present workgave a close examination on the plate thickening in stress-agedAl–Cu–Mg alloys containing various Ag contents (above 0.46%) at200 1C. The purpose of this paper was firstly to clarify themechanism that contributed to the accelerated thickening of Ωphase during stress aging by a nano-scaled characterization of theα/Ω interfacial structure. Secondly, the effect of Ag variations onthe plate thickening during stress aging was also investigated byAPT analysis.

2. Experimental procedures

The compositions of the Al–Cu–Mg–Ag alloys used in presentstudy are given in Table 1. 2 mm thick strips were solution treatedat 515 1C for 6 h, cold water quenched and immediately aged at165 1C for 2 h. After that, three underaged (165 1C�2 h) samplesfor each alloy were prepared to experience subsequent stressaging under a constant tensile stress of 240 MPa in a creepmachine. The temperature of subsequent stress aging was inten-tionally set to 200 1C arising from the fact that Ω phase rarelycontained the growth ledge at this temperature. According to therupture time of the stress-aged samples, the underaged samplesalso experienced stress-free aging at 200 1C to certain periods.Thus, it provided a good chance to examine the intrinsic mechan-ism of the tensile stress on the thickening kinetics of Ω phase byquantitative TEM characterization. Samples for stress aging wereprepared vertically to the longitudinal direction of the strips witha thickness of 2 mm and a gauge section of 30 mm.

Microstructural characterization was performed on a TecnaiG2 20 ST TEM machine operating at 200 KV. Samples suitable for

TEM observation were thin discs with 3 mm in diameter andelectro-polished by using twin-jet equipment with a voltage of15 V in an 80% methanol and 20% nitric acid at approximately�20 1C. For stress-aged samples, the stress direction was markedbefore discs were punched out from the gauge section. These discswere obtained from the locations that were far away from thefracture surfaces of the stress-aged samples if fracture occurred.All quantitative characterization was performed on TEM imagestaken near ⟨110⟩α orientations. For each sample, a minimum of1000 precipitates were measured from at least three differentgrains. The calculations of the volume fraction and number densityof the Ω plates followed the approach of Underwood [17]. The foilthickness was determined by convergent beam electron diffraction(CBED) using Kossel–Mollenstedt fringe spacing [18]. Furthermore,JEOL 3010 and JEOL F20 high resolution transmission electronmicroscope were used to examine the differences in the ledgenucleation at the α/Ω interfaces in stress-aged samples containingvarious Ag additions.

APT analysis of stress-aged samples was carried out on a LEAP3000HR instrument to clarify the solutes segregation and its effecton the stress-induced thickening of Ω phase. Very sharp needle-shaped samples for APT analysis were prepared using an electro-polishing procedure with small rods (0.6�0.6�12 mm) cut frombulk materials. APT experiments were performed at a specimentemperature of about 25–35 K in an ultrahigh vacuum of�1.2�10�10 Torr with a pulse fraction of 15% and a pulserepetition rate of 200 kHz. Reconstruction and visualization ofAPT data was conducted by Integrated Visualization & AnalysisSoftware (IVAS) 3.6 package.

3. Results

3.1. Stress rupture behavior

The studied alloys containing various Ag additions exhibiteddifferent stress rupture behavior. As shown in Table 2, the averagestress–rupture time for the underaged samples at 200 1C graduallyincreased from 91 h to 145 h, and finally to 288 h as the bulk Agcontent increased from 0.46 to 1.77. The average elongation ofstress-aged samples, inversely, followed a decrease trend from9.3% to 5.8% with the increasing bulk Ag content. This downwardtrend of the elongation, to some extent, implied the increasing Agcontent was effective to improve the creep performance of Al–Cu–Mg–Ag alloys at 200 1C.

3.2. Stress-aged TEM microstructures

For a better understanding of the mechanism that dominates thethickening process of Ω phase during stress aging at 200 1C, the Ωplates perpendicular to the applied stress direction are highlypreferred. However, it was extremely difficult to obtain the same

Table 1Chemical compositions of the studied Al–Cu–Mg–Ag alloys (wt%).

Cu Mg Ag Mn Ti Zr Fe Si Al

Alloy 1 6.45 0.65 0.46 0.29 0.09 0.14 0.04 0.01 Bal.Alloy 2 6.46 0.67 0.88 0.24 0.04 0.13 0.03 0.02 Bal.Alloy 3 6.31 0.65 1.77 0.26 0.04 0.15 0.03 0.02 Bal.

Table 2Stress rupture behavior of different underaged samples at 200 1C with a constanttensile stress of 240 MPa.

Alloys Temperature Time to rupture (h) Elongation after rupture (%)

0.46Ag alloy 165 1C�2 h 86 91 9.9 9.391 9.895 8.2

0.88Ag alloy 165 1C�2 h 136 145 6.3 7.2147 7.1153 8.3

1.77Ag alloy 165 1C�2 h 281 288 6.4 5.8291 5.0292 6.1

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orientation of Ω phase with respect to the stress direction for allstress-aged samples. Therefore, the Ω variants approximately per-pendicular to the applied stress (the angle between the plates andthe stress direction was within a range of 80–90 deg) were selectedintentionally for TEM analysis when viewed along ⟨110⟩α zone axis.

Fig. 1 showed the typical stress-aged microstructures of differentalloys following initial underaging treatment. Corresponding quanti-tative calculation results were given in Table 3. In Fig. 1, the stress-orienting effect on the Ω precipitation was clearly observed in allstress-aged samples. Contrasted with the previous studies [15,16],the high density Ω plates were found to preferentially orientedperpendicular to the applied stress axis rather than parallel to it. TEMimages also indicated that the average diameter and number densityof Ω phase perpendicular to the stress axis were higher than thoseapproximately paralleled to the stress axis. After stress aging at200 1C for 91 h, it was undoubted that the Ω plates possessed adecrease of the average diameter from 122.4 to 103.0 nm and anenhanced precipitation from 2800 to 7000 precipitates/μm3 with theincreasing Ag addition. Coarsen θ′ particles (arrowed) were evidentin Fig. 1(a and b) but was seldom observed in Fig. 1(c). This changewas consistent with the SAED patterns. No extra diffractions pro-duced by θ′ phase could be clearly distinguished from corresponding

SAED pattern in Fig. 1(c), whereas the reflections at the 1/2 {022}αpositions were evident in Fig. 1(a and b). In fact, small amount of θ′particles in 1.77Ag alloy after stress aging for 91 h could also be seenin ⟨100⟩α beam orientations. Since the precipitation of θ′ phase wasgenerally promoted during stress aging as proved in Refs. [14,15], itwas considered important to underline that the increase of Agprobably delayed the positive effect of the tensile stress on theaccelerated precipitation of θ′ phase. As seen in Fig. 1(e and f), it wasstill hard to confirm the θ′ precipitation in 1.77Ag alloy as stress agingcontinued. But our TEM examination viewed along ⟨100⟩α zone axis(not present here) showed the presence of θ′ particles with a finerdimensions as compared to the coarsen ones in Fig. 1(d). In this case,the low number density of θ′ particles in stress-aged 1.77Ag alloy wasassociated with the only variable, the bulk Ag content.

3.3. Stress-aged HRTEM microstructures

As seen in Table 3, also noteworthy was the gradual increase ofthe plate thickness in stress-aged samples as compared to that ofthe stress-free aged ones. This trend was found to be independentof the bulk Ag content because it occurred in all stress-agedsamples. The apparent increasing trend of the plate thickness

Fig. 1. Bright field TEM images showing the typical microstructures of the underaged (165 1C�2 h) samples subjected to stress aging at 200 1C under a constant tensilestress of 240 MPa for various times. The direction of the applied tensile stress was indicated by white arrows. All the images were taken near the ⟨110⟩α zone axis. (a) 0.46Agalloy, 91 h, (b) 0.88Ag alloy, 91 h, (c) 1.77Ag alloy, 91 h, (d) 0.88Ag alloy, 147 h, (e) 1.77Ag alloy, 147 h and (f) 1.77Ag alloy, 291 h.

Table 3Summary of the quantitative calculation results of Ω plates in underaged (165 1C�2 h) samples after subsequent stress aging (200 1C�240 MP) for various times.

Alloys Temperature Average plate diameter (nm) Average plate thickness (nm) Aspect ratio Volume fraction (%) Number density (no. plates/μm3)

0.46Ag alloy 200 1C�91 h 122.4 3.6 (2.0)a 34 (65) 0.1218 28000.88Ag alloy 200 1C�91 h 116.5 3.0 (1.8) 39 (68) 0.1302 4100

200 1C�147 h 120.1 3.7 (2.0) 32 (64) 0.1362 33001.77Ag alloy 200 1C�91 h 103.0 2.8 (1.3) 37 (86) 0.1628 7000

200 1C�147 h 111.5 3.2 (1.5) 34 (78) 0.1966 6300200 1C�291 h 119.6 3.9 (1.8) 31 (72) 0.2665 6000

a Data in the brackets are the quantitative calculation results of Ω phase after corresponding stress-free aging.

S. Bai et al. / Materials Science & Engineering A 589 (2014) 89–96 91

could be observed in 1.77Ag alloy during stress aging from 91 h to291 h at 200 1C. The thickening of Ω phase was controlled by aledge nucleation and the ledge propagation that associated withthe elastic strain field [5]. Generally, the plate thickening rate wasrelatively low at 200 1C as revealed in Refs. [4,6]. The acceleratedthickening kinetics observed in Table 3 indicated that the appliedtensile stress lead to the modification of the elastic strain fieldaround the Ω plates. The mechanism governed the plate thicken-ing was believed to be not as simple as that proposed in Ref. [6].Therefore, HRTEM was used to examine the α/Ω interfacialstructures in different stress-aged samples.

Fig. 2(a) showed an end of a typical Ω plate in 0.46Ag alloy afterstress aging for 91 h. This plate varied from 2.5 to 4 unit cells inthickness due to the presence of small ledges. As indicated byblack arrows, two small ledges on the broad face of Ω phase were1/2 unit cell high and coherent with the matrix. Both of themproduced an additional vacancy misfit normal to the Ω plate.Despite the complex strain field introduced by the tensile stresscovered up the details of the interfacial structure of this Ω plate, itcould also be confirmed that the plate end contained a misfitdislocation (arrowed) with an extra half plane in Ω phase. More-over, special interest was the right part of this plate (if divided byM–N white line) migrated integrally along the direction normal tothe Ω plate by a displacement of 1/4 unit cell, leaving a minorledge (white arrowed) of 1/4 unit cell high at its top and bottomfaces, respectively. As expected, these minor ledges were alsocoherent with the matrix. In Fig. 2(b), the coherent 1/2 unit cellhigh ledge (black arrowed) was evident at the middle of the broadface of a 3 unit cell thick Ω plate. Again, the integral migration ofthis plate was confirmed as a 1/4 unit cell high ledge (whitearrowed) was present at the top face of Ω phase. Contrasted withthe results of Fig. 2(a and b), small ledges were seldom observedon the broad face of Ω phase in 1.77Ag alloy after stress aging for91 h. Fig. 2(c) showed the end of a 1.5 unit cell thick Ω plate. Thisplate exhibited an interstitial type strain field with an extra halfplane in Ω phase. A 3 unit cell thick Ω plate, which was notperpendicular to the stress direction in stress-aged 0.46Ag sample,was given in Fig. 2(d). Apparently, a misfit dislocation (arrowed)could be found at the end of this plate but no ledge was observed.According to our HRTEM examination, the quantity of small ledgeson the plate broad face in stress-aged samples at 200 1C wasconsiderably increased as compared to that in stress-free agingcondition. Small ledges were much more common in the Ω platesnormal to the stress direction rather than parallel to it. Therefore,the rapid nucleation of the growth ledges was ascribed to theapplied tensile stress and resulted in the observed acceleratedthickening of Ω phase at 200 1C.

The underaged 1.77Ag alloy exhibited the longest time to stressrupture at 200 1C and small ledge was seldom observed after

stress aging for 91 h. Particular interest was to understand by whatmechanism the thickness of Ω phase in 1.77Ag alloy increasedafter stress aging for 291 h as seen in Table 3. Fig. 3 showed somedetails of two different Ω plates in 1.77Ag alloy after stressaging for 291 h. The end of a 6 unit cell thick Ω plate was givenin Fig. 3(a). The presence of a small coherent ledge (arrowed) onthe broad face of this plate was evident. Similarly, four 1/2 unit cellhigh ledges (arrowed) were clearly observed on the broad face of a7 unit cell thick plate in Fig. 3(b). Therefore, the presence of smallledges was considered to be the representative feature thatexpected for the accelerated thickening of the isolated Ω platesduring stress aging at 200 1C. Based on the HRTEM observation inFigs. 2 and 3, the observed accelerated thickening kinetics of Ωphase in stress-aged samples was closely related to the enhancednucleation and continuous growth of small ledges.

3.4. APT analysis

As revealed in Figs. 2 and 3, stress aging at 200 1C lead to thesignificant differences in the interfacial structures of Ω phase invarious alloys. Special effort was to clarify whether thesestructural differences were dependent on the change of thechemical compositions and the solute distribution of Ω phase.Fig. 4(a) illustrated APT elemental maps of an analyzed volumecontaining a segment of a coarsen Ω plate in stress-aged(200 1C�91 h) 0.46Ag alloy. The absence of an entire Ω plate inFig. 4(a) was ascribed to the existence of high density defectsproduced by stress aging. These defects facilitated the destructionof the needle-like specimen under the pressure of electric fieldforce during APT data acquisition, thereby failing to obtain an

Fig. 2. HRTEM images showing the interfacial structures of several Ω plates in different alloys after stress aging at 200 1C for 91 h. (a) and (b) the Ω plates in 0.46Ag alloyrevealed the presence of small ledges, (c) the end of an Ω plate in 1.77Ag alloy without ledge, (d) the Ω plate in 0.46Ag alloy approximately parallel to the applied tensilestress. The black and white arrows indicated the small ledges with different height. The white arrows at the corner in each image indicated the direction of the tensile stress.All the images were taken along the ⟨110⟩α zone axis in the same magnification.

Fig. 3. HRTEM images showing the ends of (a) 6 unit cell thick and (b) 7 unit cellthick Ω plates in 1.77Ag alloy after stress aging at 200 1C for 291 h. The black arrowsindicated the small ledges. The white arrows indicated the direction of the tensilestress. Both images were taken along the ⟨110⟩α zone axis in the samemagnification.

S. Bai et al. / Materials Science & Engineering A 589 (2014) 89–9692

entire plate for analysis. It was unambiguous that no Mg and Agsegregation was observed at the α/Ω interface in Fig. 4(a). From theconcentration profile across the Ω plate, it turned out that Cu, Mgand Ag were all within this plate. The concentration peaks of Cu,Mg and Ag were 31.9%, 7.3% and 6.4%, respectively. In the case ofstress-aged 1.77Ag alloy, two different types of Ω phase wereconfirmed in Fig. 4(b–c). Similarly, the interfacial segregation ofMg and Ag in the crossing Ω plates was absent in Fig. 4(b).Compared to 0.46Ag alloy, the concentration profile in Fig. 4(b) indicated that Cu concentration dramatically decreased to22.9% but the concentration peaks of Mg and Ag increased to12.1% and 13.7% in the crossing Ω plate, respectively. Furthermore,the interfacial segregation of Mg and Ag was revealed in Fig. 4(c),supported by the concentration profile possessing a two-peakcharacter. The concentration peak of Ag was still higher than thatin Fig. 4(a). Arrowed in Fig. 4(c) was probably a small ledge at thebroad face of Ω phase because the plate thickness varied slightly.

4. Discussion

4.1. Stress-induced thickening of the oriented Ω plates at 200 1C

According to the present work, it was surprising that thepreferred orientation of Ω phase was confirmed for the initialstress-free aging followed by stress aging at 200 1C. This con-tracted with the explanation of the applied stress on the oriented

nucleation and precipitation between {111}Ω and {001}θ′ byMuraishi et al. [13,14]. They reported the stress effect on theprecipitation kinetics of the competitive phases was absent as ifthe stress aging (150 MPa, 10 h) was applied after the initial stress-free aging (177 1C, 5 min). Inversely, another two-step aging intheir work (stress aging for 1 h followed by stress-free aging for10 h at 177 1C) revealed a noticeable change in the number densityof the oriented precipitates. On the basis of the quantitativeanalysis by a classic nucleation theory, they proposed that theexternal stress primarily induced the preferred nucleation of G.Pzones at the initial aging stage. Therefore, the successive precipi-tation of θ′ phase was assisted by the applied stress, whereas theprecipitation of Ω phase was suppressed during the subsequentstress-free aging [13,14]. According to their theory, further stressaging should fail to cause the orientation effect as if the compe-titive nucleation kinetics between Ω and θ′ was not changed at theearly stage of aging. This suggested no orientation effect should beconfirmed in our study because all the alloys experienced initialstress-free aging at 165 1C for 2 h. Apparently, their explanationcontrasted with our TEM characterization results in Fig. 1. Theobserved discrepancy suggested the preferred precipitation andgrowth of Ω phase was not merely determined by the competitivenucleation during the initial stress aging.

Our investigation revealed that the stress-orienting effectoccurred not only during the initial nucleation stage but alsoduring the growth stage of Ω phase. In stress-free aging, the strongelastic interaction between the Ω variants on different habit planes

Fig. 4. APT maps showing the distributions of Cu (yellow), Mg (blue) and Ag (purple) atoms of the underaged (165 1C�2 h) samples after stress aging (200 1C�240 MPa).(a) A segment of a coarse Ω plate in 0.46Ag alloy after stress aging for 91 h, (b) two crossing Ω plates in 1.77Ag alloy without interfacial segregation of Mg and Ag atoms and(c) a segment of an Ω plate revealing the distinct interfacial segregation of Mg and Ag atoms in 1.77Ag alloy after stress aging for 291 h. The concentration profiles wereobtained along the black lines. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

S. Bai et al. / Materials Science & Engineering A 589 (2014) 89–96 93

tended to achieve a self-accommodative uniform distribution tominimize the total elastic energy. Naturally, the application of ahigh stress level (240 MP) at a high temperature of 200 1Cdisturbed the primary balance between the distribution of differ-ent Ω variants. It served as an exterior inducement of the observedorienting precipitation. The preferentially oriented Ω plates wereenergetically favored in contrast to those with other orientations.The alignment of the oriented Ω plates in the studied Al–Cu–Mg–Ag alloys failed to reveal the same stress-orienting effect in Refs.[15,16] but involved with further selective growth and thickeningby subsequent aging under a tensile stress of high magnitude.As shown in Figs. 2 and 3, small ledges were much more commonin the Ω plates normal to the stress axis. Rapid nucleation of thesesmall ledges, therefore, facilitated the continuous growth of Ωplates in the favored orientation as stress aging prolonged. Thisphenomenon was believed to be a result of the change of theinstantaneous elastic strain fields surrounding each precipitateduring stress aging.

The mechanism that accounted for the accelerated thickeningkinetics of Ω phase in stress aging was seldom investigated inprevious studies as the stress-induced nucleation of the growthledge was unconsciously ignored [15]. In general, two types of theΩ plates could be identified in Al–Cu–Mg–Ag alloys during thestress-free aging at 200 1C. One type is the plates intersect witheach other by rapid lengthening, the other is the isolated thickplates. The intersection or impingement between the precipitateson different habit planes was considered to play an important roleon the accelerated thickening kinetics of the plate-like precipitatesby providing a potential source for the nucleation of the growthledges, such as γ/γ′ in Al-Ag alloys [19,20]. This mechanismdominated the thickening kinetics of the plate-like precipitateswith large aspect ratios. In our work, the applied tensile stress leadto the preferred Ω precipitation and a significant reduction of theplate aspect ratio, as revealed in Fig. 1 and Table 3. Therefore, theimpingement frequency of different Ω plates was considerablylimited. Despite a small amount of Ω plates was observed tointeract with each other, the accelerated thickening by precipitateimpingement was not predominant in present study. The isolatedΩ plates present in Fig. 1 were most likely to be responsible for theaccelerated plate thickening during stress aging at 200 1C ascompared to those in stress-free aging conditions.

The growth and thickening of the precipitates in simple binaryalloys, such as θ′ [7,8] and γ′ [20], was widely discussed by a ledgemechanism. This mechanism seems to be overwhelming inexplaining the thickening kinetics of the plate-like precipitates.Turn to Ω phase, HRTEM investigations by Fonda et al. [5] andHutchinson et al. [6] indicated that the plate thickening wasassociated with the nucleation and growth of the ledges withrespect to both positive and negative misfit that normal to thebroad face of Ω phase. Accordingly, the elastic strain field aroundthe Ω plate was found to determine the ledge propagation. Thenucleation of small ledges on the broad face of Ω phase was hardto reveal during stress-free aging at 200 1C, as suggested byprevious investigations [4,6,21,22]. Therefore, comparisons of theplate thickness in Table 3 indicated the applied tensile stress wasdetrimental to the stability of Ω phase at 200 1C and lead to thedestruction of Mg–Ag segregation layer to satisfy the conditionrequired for ledge nucleation. The increasing trend of the averageplate thickness in Table 3 implied a substantial change in theunrelaxed misfit normal to the Ω plates during stress aging. It wasundoubtedly that the change of the strain type of Ω phase (fromvacancy to interstitial type and vice versa) during stress aging wasfacilitated. Thus, it was quite natural that the thickening kinetics ofΩ phase in stress aging at 200 1C was comparable to that of thestress-free aging at temperatures above 250 1C. During stress-freeaging at 200 1C, the plate thickening occurred at a rate much lower

than that of diffusion control due to the lack of a potent source ofgrowth ledges. In Figs. 2 and 3, the accelerated kinetic conse-quence of the growth ledge in stress-aged samples was readilyrevealed, highlighting the ledge nucleation and migration werefacilitated at 200 1C. This striking result at 200 1C indicated a highdriving force for the rapid ledge nucleation and could be rationa-lized on the basis of the interaction between the pre-existingelastic strain field of Ω phase and the tensile stress. The instanta-neous elastic strain field around the Ω plate, which was deter-mined by the interfacial misfit for the externally unloadedcondition, was highly influenced by the applied stress. The tensilestress of high magnitude should make the ledge growth easier bylowering the activation energy barrier for ledge nucleation, asimplied in Figs. 2 and 3.

Meanwhile, the plate thickening was highly diffusion related.The rate of diffusion was known to depend on the strain field aswell as upon compositional gradients. The substantial change inthe chemical concentrations of Ω phase during growth andthickening in stress aging was mediated by diffusion. Hence, onepotential problem with this process must be emphasized. Thedissolution of the Ω plates (e.g. those not perpendicular to thestress axis or with small dimensions) inevitably transportedsufficient solute atoms to the favored plates. This might beexpected to involve with the accelerated thickening of Ω phase.Analysis of the thickening kinetics of the plate-like precipitatesshowed the ledge thickening was interface controlled but therelated migration of the growth ledge was diffusion controlled [6].As seen in Table 3, the trend of the plate aspect ratios in 0.88 and1.77Ag alloys with the increasing stress-aging time was down-ward. This was a reasonable change because the aspect ratiosignificantly decreased when the plate-like precipitate evolvedinto the equilibrium phase with a low aspect ratio. The high valuesof aspect ratio possessed an important effect of restricting thediffusion flux to the broad face of the plate, thereby giving a slowrate of the plate thickening. The apparent decreasing trend of theplate aspect ratio in Table 3, therefore, suggested that the platethickening was intended to achieve the rapid solute diffusion tothe rim of the new developed ledge during stress aging. Thisoutstanding difference underlined that the rapid growth of smallledges, as well as the effective solute diffusion, was responsible forthe high thickening rate of Ω phase during stress aging at 200 1C.As revealed in Figs. 2 and 3, the continuous stress aging provided asupply of the growth ledges at a sufficient rate so that thecondition required for the diffusion-controlled mechanism tendedto be gradually established. As a consequence, the thickeningkinetics of Ω phase was not only determined by a ledge-controlled mechanism but also involved with a diffusion-assistedmechanism. The solute-rich zones surrounding the Ω plates couldbe associated with a substantial overall supersaturation of soluteatoms that transferred from the dissolution of the unfavorable Ωplates. Once a single ledge was generated at the broad face of Ωphase under the high driving force of stress aging, the platethickening was highly driven by the rapid diffusion of soluteatoms to the terrace of the newly formed ledge. In such way, themechanism governed the thickening of Ω phase was believed toexperience a gradual transition from a pure ledge mechanism instress-free aging to a complicated mechanism consisted of thecombined contributions of the ledge nucleation and sufficientsolute diffusion during stress aging. Thus, the plate thickening in1.77Ag alloy was considered to be dominated by a ledge mechan-ism after stress aging for 91 h due to the lack of small ledges. Butthe plate thickening in stress-aged 0.46Ag alloy was probablycontrolled by the combined contributions of the ledge nucleationand rapid solute diffusion since the nucleation kinetics of smallledges was significantly accelerated, as supported by Fig. 2. Even ifthe ledge nucleation rate was greatly accelerated during stress

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aging, it should be recognized that in present work the platethickening was far from that dominated by the individualdiffusion-controlled mechanism. It was arising from the fact thatthe stable diffusion-controlled mechanism could be established asif the aspect ratio of the plate-like precipitate was predicted toremain constant [23,24]. In the view of the plate aspect ratio, itcould be concluded that the tendency of the diffusion-controlledthickening at 200 1C was intensified by stress aging, as supportedby an apparent downward trend of the plate aspect ratio inTable 3.

4.2. The effect of Ag on the ledge nucleation during stress-agingat 200 1C

The stress-induced accelerated thickening kinetics was evidentfor Ω phase in all stress-aged alloys, regardless of bulk Ag contents.However, the accelerated thickening of Ω phase in stress-agedalloys containing various Ag additions was found to be quitedifferent. In Fig. 2(c), the presence of small ledge on the broadface of Ω phase in 1.77Ag alloy was rarely observed. This con-trasted with our HRTEM results of 0.46Ag alloy in Fig. 2(a and b),from which the small ledges could be clearly distinguished. Afterstress aging for 91 h, the pronounced difference between 0.46 and1.77Ag alloys was the nucleation density of the growth ledges.Concentration profiles of APT analysis in Fig. 4(a and b) revealedthat the concentration peaks of Mg (7.3%) and Ag (6.4%) for thick Ωplate in 0.46Ag alloy significantly increased to 12.1% and 13.7%for thin Ω plate in 1.77Ag alloy, respectively. Apparently, theconcentration of Ag in thin plate in 1.77Ag alloy doubled thatof the thick plate in 0.46Ag alloy. Considering the sample inFig. 4(b) experienced a much longer time to stress–rupture thanthat in Fig. 4(a), it was proposed that the enrichment of Ag atomswithin Ω phase was beneficial to enhance the plate thickeningresistance during stress aging as the average thickness (3.9 nm) ofΩ plates in stress-aged (200 1C�291 h) 1.77Ag alloy was compar-able with that (3.6 nm) of Ω plates in stress-aged (200 1C�91 h)0.46Ag alloy. Moreover, concentration profile in Fig. 4(c) revealedthe concentration peaks of Mg and Ag of Ω phase dropped to 6.7%and 7.1%, which was approximately one half of those in Fig. 4(b).This was a reasonable change because Mg and Ag atoms diffusedout from the interior of Ω phase to the α/Ω interface, ultimatelyleading to the formation of Mg–Ag segregation layer at both side ofthe plate broad face. Thus, the Mg and Ag concentrations of Ωphase with distinct segregation layers were considered to be muchlower than those of Ω phase without segregation layers. AlthoughAPT analysis in Fig. 4(a) failed to reveal the interfacial segregationof Mg and Ag in stress-aged 0.46Ag alloy, it was reasonable toassume that the interfacial concentrations of Mg and Ag in 0.46Agalloy were lower than those in 1.77Ag alloy, similar to the APTresults present in Fig. 4(a) and (b). HRTEM results in Figs. 2 and 3suggested stress aging at 200 1C lead to the generation of smallledges in Ω phase. The slowest nucleation rate of the growthledges in stress-aged 1.77Ag alloy was most likely to be the directresult of the enhanced stability of Ag-richer interfacial structure.As revealed by APT analysis, the progressive enrichment of Agwithin segregation layer would help to stabilize the structure of Ωphase and act as an effective obstacle to the nucleation andmigration of the growth ledge, finally lowering the thickeningrate of Ω phase with respect to the increasing Ag content duringstress-aging. This explanation was supported by quantitativecalculation results in Table 3 and was in agreement with the firstprinciples investigation by Sun et al. [25]. Their work on thechemical bonding of Mg and Ag atoms at the α/Ω interfaceindicated that the drop of Mg/Ag atomic ratio within the segrega-tion layer probably satisfy the lowest-energy interfacial structurerequirement to enhance the stability of Ω phase. Therefore, the

progressive enrichment of Ag lead to the drop of Mg/Ag atomicratio and was responsible for the observed difference in the ledgenucleation of various stress-aged samples. The segregation layerricher in Ag helped to enhance the structural stability of Ω phaseand limited the ledge migration during stress aging at 200 1C.Finally, the ledge nucleation rate and the plate thickness in 1.77Agalloy were lower than those in other alloys during stress aging.Based on HRTEM and APT results, the potential influence ofincreasing Ag on delaying the ledge nucleation during stress agingwas highlighted. But it should be noted that this delaying effectcould not totally balance the accelerated thickening of Ω phaseduring stress aging because the average thickness of Ω phase instress-aged 1.77Ag alloy, as seen in Table 3, also exhibited aconsiderably increase when compared to that of stress-freeaging alloy.

5. Conclusions

The accelerated thickening of Ω phase in underaged Al–Cu–Mgalloys containing various Ag additions during stress aging at200 1C was well investigated by a combination of TEM, HRTEMand APT analysis. Results are summarized as follows:

(1) The time to rupture of the stress-aged Al–Cu–Mg–Ag alloy wassignificantly increased from 91 h to 288 h with an increase ofthe bulk Ag addition from 0.46 to 1.77.

(2) The preferred orientation of Ω phase was revealed in all thestudied alloys during stress aging at 200 1C following theinitial aging at 165 1C for 2 h. Therefore, the stress-orientingeffect was also confirmed to be available during the growthstage of Ω phase.

(3) Compared to the stress-free aging at 200 1C, quantitative TEMcalculation revealed the average thickness of Ω phase pos-sessed an apparent increase trend during stress aging, sug-gesting the thickening resistance of Ω phase in all stress-agedalloys was considerably degraded. This stress-induced accel-erated thickening of Ω phase also showed its close connectionwith the bulk Ag content as the average plate thicknessdecreased with increasing Ag addition after stress aging forthe same time.

(4) HRTEM observation on the stress-aged microstructures indi-cated the applied tensile stress facilitated the rapid nucleationof the growth ledges on the broad face of Ω phase at 200 1C,thereby leading to the accelerated thickening kinetics of Ωphase observed by quantitative TEM characterization.

(5) APT analysis proposed that the α⧸Ω interfacial structure richerin Ag helped to stabilize the structure of Ω phase, therebysuppressing the rapid nucleation of small ledge during stressaging at 200 1C. This was a reasonable explanation on thedifferences in the ledge nucleation between stress-aged 0.46and 1.77Ag alloys as the Ag concentration of Ω phase in 1.77Agalloy was remarkably higher than that in 0.46Ag alloy.

Acknowledgments

The authors are grateful for the financial support from theNatural Science Foundation of China (Grant no. 51171209), theNational Basic Research Program of China (973 Program) and 2011Program. Special thanks are given to Prof. Wenqin Liu of Instru-mental Analysis & Research Center at Shanghai University for hisassistance with APT analysis.

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