a hrtem study of metastable phase formation in al–mg–cu alloys during artificial aging

Acta Materialia 52 (2004) 2509–2520

www.actamat-journals.com

A HRTEM study of metastable phase formation in Al–Mg–Cualloys during artificial aging

L. Kovarik a,*, P.I. Gouma b, C. Kisielowski c, S.A. Court d, M.J. Mills a

a Department of Materials Science and Engineering, The Ohio State University, 2041 College Road, Watts Hall 477, Columbus, OH 43201, USAb Department of Materials Science and Engineering, State University of New York at Stony Brook, Stony Brook, NY 11794, USA

c National Center for Electron Microscopy, Ernest Orlando Lawrence Berkeley National Laboratory, One Cyclotron Road, Berkeley, CA 94720, USAd Alcan Technology and Management Ltd., CH-8212 Neuhausen, Switzerland

Received 10 October 2003; received in revised form 29 January 2004; accepted 30 January 2004

Available online 5 March 2004

Abstract

Microstructure evolution of an age hardenable Al–3Mg–0.4Cu–0.12Si (wt%) alloy has been studied during artificial aging at

180 �C prior to the formation of the stable S-phase. The primary investigation method used in this study was high-resolution

transmission electron microscopy (HRTEM), coupled with image processing and image simulation. After 1 h of aging, the presence

of super-lattice reflections was detected in the Fourier spectra of the HRTEM images, suggesting an L10 type ordering of Mg and Cu

atoms in the Al matrix. After 4 and 8 h of aging, coherent particles were observed in the microstructure. These particles give rise to

diffraction spots that in previous literature have been considered to be characteristic of the S00-phase in the ‘‘Cu-lean’’ Al–Mg–Cu

alloys. It is shown that these diffraction spots can be indexed in terms of a crystal structure that is closely related to the L10 ordering

formed at the shorter aging times. The crystal structure is orthorhombic with lattice parameters a ¼ 1:2 nm, b ¼ 0:4 nm, c ¼ 0:4 nm

and space group Cmmm. We propose to identify these coherent particles as GPB-II zones, and the ordering that precedes them as

GPB zones.

� 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Al–Mg–Cu alloys; Age hardening; GPB zones; HRTEM; Image reconstruction

1. Introduction

The solid solution decomposition of age hardenable

Al–Cu–Mg alloys has been frequently studied in recent

years. Aside from the more numerous studies of Al–Cu–Mg alloys, some attention has also been paid to the Al–

Mg–Cu alloys (closely related to 5000 series aluminum

alloys) that represent a medium strength alloy, possibly

being suitable for car body applications. The notation

‘‘Al–Mg–Cu’’ is meant to imply that Mg is present in

much higher concentrations than Cu. In the work of

Ratchev et al. [1,2], it was shown that the age-hardening

characteristics of these Al–Mg–Cu alloys are in severalrespects similar to those in the Al–Cu–Mg alloys. The

* Corresponding author. Tel.: +1-614-688-3409; fax: +1-614-292-

1537.

E-mail address: [email protected] (L. Kovarik).

1359-6454/$30.00 � 2004 Acta Materialia Inc. Published by Elsevier Ltd. A

doi:10.1016/j.actamat.2004.01.041

similarities include two-stage hardening, with the first

stage proceeding very rapidly, after only several minutes

of aging. The differences are associated with the alloy

behavior following the rapid-hardening stage. In case of

the Al–Mg–Cu system, the alloys continue to harden ata relatively low rate, while the Al–Cu–Mg alloys develop

a plateau on the age hardening curve for up to several

hours of artificial aging.

One of the few studies that has addressed micro-

structure development of the Al–Mg–Cu alloys at vari-

ous stages of aging was published by Ratchev et al. [1].

Based on the authors findings, the following precipita-

tion sequence was considered: a-(supersaturated solidsolution))GPB) S00 ) (S0)S, which is identical to the

sequence proposed by several authors for the Al–Cu–

Mg [3,4]. In the work of Ratchev et al. [1,2], however,

the existence of the GPB zones was surmised only

indirectly, based on the strength response to various

ll rights reserved.

mail to: [email protected]

2510 L. Kovarik et al. / Acta Materialia 52 (2004) 2509–2520

thermal treatments. No direct evidence confirming the

existence of GPB zones, such as via the presence of

characteristic diffraction spots, was presented. Studies of

the Al–Cu–Mg alloys indicate that GPB zones have a

rod-like morphology, being 1–2 nm in diameter and 4–8nm in length. Depending on the aging temperature,

these rod-like GPB zones become detectable only after

several hours of aging [5]. In electron diffraction pat-

terns taken along the h001iAl zone, the presence of GPB

zones is evidenced as diffraction streaks forming crosses

around the forbidden {1 1 0}Al spots [5]. Although

studied by several authors, the crystal structure of the

GPB zones is still the subject of debate. From the pre-cipitation sequence, it is apparent that the GPB zone is

the first metastable phase to form, and has therefore

been linked to the rapid hardening phenomenon.

However, since it has not been detected at the very early

stages of aging, some recent work suggests that such

rod-like GPB zones are not reasonable for the rapid

hardening. Instead, alternative interpretations including

solute cluster hardening [5,6] or dislocation–solute in-teraction have been proposed [7].

In the work by Ratchev et al. [1] the existence of the

S00-phase at an early stage of aging was also surmised

based on carrying out thermal treatments, while in the

later stages S00-phase became directly detectable. Al-

though not discussed by the authors, the diffraction

spots from this S00-phase are somewhat different than

those considered to be characteristic of the S00-phase inAl–Cu–Mg [8,9]. In the Al–Cu–Mg alloys, moreover,

the existence and characteristics of the S00-phase are

controversial. Some authors support its existence [8,10],

and envision the S00-phase to have a fundamentally dif-

ferent crystal structure from the S-phase while contrary

to this view, it has been shown that the diffraction spots

that were considered characteristic of the S00-phase ac-

tually can be understood in terms of the S-phase [11,12].Several important issues remain to be answered re-

garding the precipitation processes in the Al–Mg–Cu

alloys. The first and the most important is the crystal

structure and morphology of the transition phases. In

addition to that, conclusive evidence supporting the

proposition that the precipitation behavior of Al–Mg–

Cu (Cu-lean) and Al–Cu–Mg (Cu-rich) alloys are the

same has not been provided. The present uncertainty iscompounded by the confusion in the literature con-

cerning the nature of the ‘‘characteristic’’ S00-phase dif-

fraction spots in these two systems. With respect to the

Al–Mg–Cu alloys, it has been suggested moreover that

Table 1

Composition of the alloy studied (in wt%)

Al Mg Cu Si

Wt% Bal. 2.96 0.42 0.12

diffraction spots considered as due to the S00-phase are

actually an artifact caused by a surface contamination,

probably an oxide layer [13]. The goal of this work has

been to revisit the microstructural study of Al–Mg–Cu

alloys using high-resolution transmission electron mi-croscopy (HRTEM) in order to provide more direct

structural information regarding the nature of the

metastable phases as a function of aging time.

2. Experimental

The alloy studied was laboratory cast and supplied(Alcan International Ltd.) in the form of a 1 mm rolled

sheet. The composition of the alloy is given in Table 1.

The alloy was solution heat treated in a salt bath at the

550 �C for 30 min. Aging was carried out at 180 �C for a

range of times between 1 and 8 h. Samples for the

HRTEM observations were first ground to a thickness

of 200 lm. The grinding was followed by jet polishing

(Twin-Jet Polisher, Fischione, Model-110) in a 25% ni-tric acid and 75% methanol solution. The polishing

conditions were )30 �C, and a voltage �8 V. The per-

forated thin foils were subsequently thinned using a low

energy ion mill (LINDA) operated at 500 V. Finally, an

amorphous �2 nm thin film was deposited on the back

side of the foil to prevent the knock-off damage during

the observation.

The high-resolution TEM investigations were per-formed on Philips CM 300 FEG microscopes at both the

National Center for Electron Microscopy and at the

Ohio State University. An accelerating voltage of 300 kV

was used. In addition to recording individual images,

focal-series reconstructions were also performed. In the

latter cases, a series of 20 images were recorded for each

area studied that was later used for the image recon-

struction. The image reconstruction was performed withthe Philips/Brite-Euram software by Coene et al. [14,15].

The result is an exit wave function that contains ampli-

tude and phase information. Diffraction patterns and the

HRTEM images of the proposed crystal structure were

simulated using EMS� software package. In the

HRTEM simulations, the microscope parameters were:

accelerating voltage¼ 300 kV, spherical aberration co-

efficient¼ 0.6 mm, beam convergence¼ 0.2 mrad, spreadof defocus¼ 4.9 nm. The HRTEM simulation was per-

formed on an Al crystal supercell (25 · 25 · 15 or

25 · 25 · 65 unit cells in size) into which the proposed

GPB zones were incorporated.

Mn Fe Zn Ti

0.25 0.21 0.007 0.002

Fig. 2. The schematic representation of the morphology of GPB zones

and their crystal structure.

L. Kovarik et al. / Acta Materialia 52 (2004) 2509–2520 2511

3. Results

The results of this TEM study are organized to show

the progression in the microstructure as a function of

aging time. Section 3.1 treats the observations after 1 hof aging, and presents the proposed nature of the GPB

zones at short aging times. An analysis of the size of the

GPB zones is then presented in Section 3.2. Section 3.3

presents the microstructure after 4 h of aging for which

two distinct types of zones are observed. In Section 3.4,

the microstructure after 8 h is presented in which distinct

particles are observed using bright field imaging, and the

structure of the particles is deduced from HRTEM ob-servations along several zone axes.

3.1. HRTEM studies of samples aged for 1 h

HRTEM observations of samples aged for 1 h failed

to reveal any obvious contrast relevant to the presence

of GPB zones in the Al matrix. For the most part, ob-

servations such as shown in Fig. 1(a) were made. Allrecorded HRTEM images were fast Fourier trans-

formed (FFT) and it was noticed that some of the FFT

spectra showed the presence of faint {1 1 0} super-lattice

reflections; an example is shown in Fig. 1(b). These

{1 1 0} reflections are quite faint and relatively sharp,

nevertheless rising above the diffuse background inten-

sity. The appearance of the {1 1 0} super-lattice reflec-

tions is a new observation for the Al–Mg–Cu system.While preliminary reporting of this observation has been

published elsewhere [16], the following is a much more

detailed account of the nature of the zones giving rise to

these reflections, and of how these zones are related to

the subsequent precipitation sequence. Note also that

the apparent streaking of the transmitted and funda-

mental matrix diffraction spots, aligned perpendicular to

the edges of the image, is an artifact due to the FFTtransformation.

The presence of super-lattice {1 1 0} reflections is

presumed to arise from the ordering of Mg and Cu at-

oms on the FCC Al lattice. By the reasoning to be

Fig. 1. (a) The HRTEM image obtained from a sample aged for 1

presented below, we have deduced this ordering to be of

the L10 type, as illustrated in Fig. 2. First, consider that

only the variant that has its c-axis aligned along the

[0 0 1]Al would contribute to the {1 1 0} super-lattice re-

flections for a given HRTEM image recorded on the[0 0 1]Al zone axis. The two other variants, which have

the c-axis perpendicular to the [0 0 1]Al zone axis, should

give rise to {1 0 0} reflections. However, no observations

of these {1 0 0} reflections have been made. A justifica-

tion for the absence of these {1 0 0} reflections in L10structure is nevertheless possible by considering that the

L10 ordered regions have a thin rod-like morphology, as

illustrated in Fig. 2. Under these conditions, the twovariants with the c-axes perpendicular to the viewing

direction cause an insufficient change in projected po-

tential in the TEM foil. Therefore, the {1 0 0} reflection

should be practically invisible. Additionally, there can

be a canceling effect if antiphase boundary (APB)-re-

lated domains are superimposed on top of each other.

With the rod-like morphology of the L10 order, it is thus

possible to explain the presence of the {1 1 0} reflections,the only kind of super-lattice reflections observed in the

spectra. Moreover, with this morphology, it is also

possible to explain why the strength of the {1 1 0} re-

flections can vary between different regions. Those re-

gions for which the super-lattice reflections were weak or

absent may imply that the variants are present with their

c-axis perpendicular to the viewing direction.

h at 180 �C. (b) FFT spectrum of the corresponding image.


In the above discussion, the variation in the intensity

of the {1 1 0} reflections was assumed to reflect the

evolution of the microstructure. However, the experi-

mental conditions used for recording the HRTEM im-

ages might also be the source of the variation in {1 1 0}intensities. For instance, using certain defocus values,

the {1 1 0} spatial frequencies can be substantially at-

tenuated. In order to eliminate possible imaging arti-

facts, an image reconstruction procedure was performed

by recording a focal series of images for each region, and

then extracting the wave function at the exit plane of the

sample. The image reconstruction process yielded some

images with stronger {1 1 0} reflections, while in othercases no {1 1 0} intensity was observed. This behavior is

similar to that observed in individual images.

3.2. Size of the GPB zones

The reconstructed HRTEM ‘‘phase images’’ were

subsequently used in an image processing procedure that

was intended to locate the regions that gives rise to the{1 1 0} reflections. This procedure consists of applying a

ring mask on the FFT spectrum of the phase image. The

ring diameter is chosen such that only the spatial fre-

quencies of a magnitude close to the super-lattice re-

flections {1 1 0}Al are included. The filtered FFT

spectrum is then inverted, resulting in an intermediate

image. This image was subsequently sampled digitally

for pattern matching of a motif that gives rise to {1 1 0}reflections. A similar technique has been applied to the

study of partial order in Au4Cr and Au4Mn alloys [17].

The phase image used in this analysis is shown in

Fig. 3(a), and the result after ring filtering and inverse

Fourier transformation (IFT) are given in Fig. 3(b).

Cursory inspection of peak positions, marked with

crosses in the image, suggests that they are randomly

located instead of having the square pattern as expectedfrom {1 1 0} reflections. However, digital sampling

(pattern matching) showed that there exist areas with

patterns that give rise to {1 1 0} super-lattice reflections.

Fig. 3. (a) The phase image of the reconstructed wave function. (b) The image

these that give rise to (1 1 0) diffraction spots.

These areas are circled, as seen in Fig. 3(c). The enclosed

gray areas represent the positions where the circles are

located on/or just a few pixels from the atomic column

positions. The other encircled areas are not located on

the atomic columns and therefore their validity can bequestioned [17]. The image processing procedure en-

abled us to determine that the ordered areas are very

small, of the order of only 1–2 nm in diameter. To

validate this procedure, it was first ‘‘calibrated’’ on

simulated images without any local order, and no such

ordered zones were detected.

HRTEM image simulations were subsequently per-

formed using the EMS� software package to provideinformation related to the size of the L10 ordering of Cu

and Mg atoms on the FCC lattice. The images were

simulated along the [0 0 1] zone axis for two different

cases: (a) the whole crystal was considered to have a L10crystal structure (maintaining the overall alloy compo-

sition) and (b) small L10 zones (1 nm in diameter and 4

nm in length), modeled on the structure shown in Fig. 2,

were introduced into a pure Al matrix. In the FFT ofthese simulated images, very diffuse {1 1 0} super-lattice

reflections were seen in the case of small zones in the

matrix, as opposed to sharp {1 1 0} reflections for the

homogeneous L10 structure.

Since the experimental observations indicate quite

sharp {1 1 0} reflections, as shown in Fig. 1(b), the zones

may not be narrow ordered regions with sharp and

discrete borders. To comply with the results from boththe image processing analysis of Fig. 3, which suggests

that the ordering originates from small discrete areas,

and the HRTEM image simulations, we propose that

the ordered regions are thin rod-like GPB zones that do

not have clearly defined interfaces. The ordering mod-

ulation in the GPB zone must vary smoothly, with a

maximum ordering near the center of the zones, and

diminishing order with radial distance from the center ofthe zones. In fact, such a view of these GPB zones im-

plies that they form through a process analogous to

spinodal decomposition/ordering [18].

obtained after the IFFT. (c) The image with the areas corresponding to


The rod-like morphology of GPB zones is consistent

with the geometry of the GPB zones proposed by Sil-

cock [4] in the Al–Cu–Mg alloys. Despite the geomet-

rical similarities, the diffraction spots that Silcock [4]

attributed to the GPB zones are very different. Silcock,observed half-crosses located around the forbidden

{1 1 0} reflections, streaked along the [0 1 0]Al direction,

which is unlike the sharp {1 1 0} reflections observed in

our work. Nevertheless, it will be shown later that such a

streaking observed by Silcock is very similar to what we

observe in the present work in the latter stages of aging.

3.3. HRTEM studies of samples aged for 4 h

The majority of the HRTEM observations from

samples aged for 4 h were similar to those seen for the 1

h aged sample. However, a few exceptions revealed the

presence of small particles in the microstructure. An

example of one such observation is shown in Fig. 4(a).

The particle is extended along [1 0 0]Al and is visible only

due to a slight, systematic change in contrast of adjacentatomic columns – brighter and darker atomic columns

are observed than in the matrix. This image was taken at

a relatively large defocus value (Df � �298 nm) and

given the microscope parameters, the bright spots

should represent rows of atomic columns. Considering

an ordering of Mg and Cu on the Al lattice, this rela-

tively strong contrast between the adjacent atomic col-

umns suggests that the particle is also extended in theviewing direction. In view of the fact that the particle is

extended along [1 0 0]Al, and through the depth of the

foil along [0 0 1]Al, this suggests that the zone has a thin,

plate-like morphology lying parallel to the (0 1 0)Al.

The FFT transformation of this image provides in-

sight into the type of ordering and the morphology of

the detected particle, as shown in Fig. 4(b). The intensity

at {1 1 0} positions is observed with distinct streakingalong the [0 1 0]Al. Since these diffraction streaks lay at

{1 1 0} positions, the particle can be considered as a

GPB zone that has grown along the [1 0 0] direction. In

fact, close examination of these streaks reveals that they

Fig. 4. (a) The HRTEM image of the developed G

consist of three distinct spots. It will be shown in Section

3.4 that similar streaking is expected from one variant of

small GPB-II zones that are found after 8 h of aging.

Another observation of a metastable particle in the

matrix after 4 h of aging was obtained via image re-construction, as shown in Fig. 5(a). Employing image

reconstruction was crucial to visualize the particle, as it

was essentially invisible in the directly acquired images

(probably caused by strong delocalization effects due to

the large defocus values (Df ¼ �295 to )262 nm) used

to record the images). The result of the image recon-

struction showed that the particle is extended along the

[1 0 0]Al direction and again it probably exists in theform of a plate lying on the (0 1 0)Al, similar to the ob-

servation in Fig. 4(a). However, examination of the FFT

spectra reveals that the diffraction spots do not corre-

spond to those found in Fig. 4(b). In fact, we find close

similarities with diffraction pattern considered by Rat-

chev et al. [1] as characteristic of the S00-phase in ‘‘Cu-

lean’’ Al–Mg–Cu alloys. Thus, after the 4-h age, we see

the presence of two distinct types of particles: one typethat exhibits diffracted intensity at {1 1 0} reflections (the

same or slightly more developed as what was seen after

1 h of aging) and a second type exhibiting distinctly

different diffracted intensities. The latter type of particle

is observed in more abundance after 8 h of aging, as

discussed in the next section.

3.4. TEM observations of particles in samples aged for 8 h

Distinct particles were more easily found in the mi-

crostructure after 8 h of aging. These particles were

detectable using conventional TEM, being visible even

without the use of an objective aperture. In order to

demonstrate the particle morphology, bright-field im-

ages from different zone axes are shown in Fig. 6. The

image taken on the [0 0 1]Al is shown in Fig. 6(a) and itshows the presence of two types of particles in the mi-

crostructure. The first type, circled and labeled as A,

shows relatively strong contrast while the second type,

circled and labeled as B, shows weaker contrast. The

PB zones. (b) Corresponding FFT spectrum.

Fig. 5. (a) Phase image showing the presence of a small particle in the matrix. (b) Corresponding FFT spectrum.

Fig. 6. Presence of small particles in the microstructure. (a) Image from the [0 0 1] zone. (b) Image from the [0 1)1] zone. (c) Diffraction pattern from

the [0 0 1]Al.


particles labeled as B are about twice the length of those

labeled as A. The image shown in Fig. 6(b) was taken onthe [1 1 0] zone axis and shows only one particle mor-

phology, indicating that both A and B are in fact dif-

ferent variants of the same type of particles. Based on

the observations, and assuming that the particles are

visible only along a direction in which they are sub-

stantially extended through the foil thickness, it appears

that the particles have a lath shaped morphology. The

lath edges are oriented along the h100i matrix direc-tions. Based on the cursory measurements from such

images, the laths have a thickness of approximately

1 nm, width of 2–4 nm and length of 6–8 nm.

Complementary to the direct microstructure obser-

vations, diffraction patterns taken on the [0 0 1]Al, such

as that in Fig. 6(c), show the presence of two types of

extra diffraction spots. The first type of reflection can be

thought of as diffraction spots located close to {1 2 0}Al

positions. These reflections are identical to those de-

tected in Fig. 5(b) (see the circled reflection), and are

considered characteristic of the S00-phase in Cu-lean Al–

Mg–Cu alloys [1]. The second type of reflection are

streaks which form ‘‘crosses’’ centered at (1 1 0)Al.

It must be noted here, however, that the diffraction

spots of the first type have been found in various Al

alloy systems and interpreted as due to surface con-

tamination [19,20]. The interpretation given by Park andArdel [20] assumes that fine polycrystalline c-Al203forms on the surface of the foil and gives rise to the

diffraction rings. Then, as a result of double diffraction,

the rings duplicate themselves around the fundamental

matrix reflections. At the intersection of the rings, the

visible diffraction spots are formed. This interpretation

provides a very good qualitative explanation for the

presence of contamination diffraction spots. In fact, thisinterpretation also holds for other major low index

zones. Indeed, when samples were viewed multiple times

without plasma cleaning, these contamination diffrac-

tion spots were observed. However, the extra spots

shown in Fig. 6(c), and also those shown later in

Fig. 7(b), are due to the presence of particles in the

microstructure as these samples were plasma cleaned

immediately prior to TEM examination, which effec-tively suppresses diffracted intensities due to the pres-

ence of contamination.

An example of an HRTEM observation of a type A

particle is shown in Fig. 7(a). The particle can be iden-

tified mainly due to a change in the intensity of the

bright spots in the central part of the image. Again, the

particle is extended along the [1 0 0]Al direction, consis-

Fig. 7. (a) Presence of a particle identified after 8 h of aging. (b) Corresponding FFT spectra. (c) FFT spectrum from an area next to the identified

particle (indicated by the white box).


tent with the expected plate-like morphology discussed

above. The FFT, shown in Fig. 7(b), reveals that the

particle gives rise to diffraction spots that are identical to

the diffraction spots of the ‘‘first type’’, described with

reference to Fig. 6(c). In Fig. 7(b), only one of the twopossible diffraction spot variants is present, and the

streaking of the spots is consistent with the narrow,

plate-like geometry of the particle in Fig. 7(a). It is also

important to mention that FFTs obtained from the vi-

cinity of the particles (e.g. Fig. 7(c)) exhibit no extra

diffraction spots in the spectrum. This result is typical of

all conditions for which particles are resolved in the

matrix, and suggests that the surrounding matrix may bedepleted of solute. Chemical analysis to confirm this

supposition is presently underway.

An example of an HRTEM observation of a particle

of type B is shown in Fig. 8(a). Again, the particle can be

identified mainly due to a change in the intensity of the

bright spots in the central part of the image. At close

inspection, it is seen that the particle gives rise to con-

trast in which brighter and darker columns alternatealong the [1 0 0]Al. The contrast in the image is in fact

very similar to the observation shown in Fig. 4. How-

ever, after 8 h of aging the particle has grown to larger

Fig. 8. (a) Presence of a particle identified after 8 h of aging. The particle

indicate location of the particle. (b) Corresponding FFT spectra.

size, now being about 8.0 nm in length and 3–4 atomic

layers thick (or about 0.6 nm). The particle width

measured from the HRTEM image is somewhat smaller

than the measured value of 1.0 nm from the diffraction

contrast image of Fig. 6(a). The FFT of the image re-veals the character of the extra diffraction spots associ-

ated with the particle; split streaks positioned on the

{1 1 0} forbidden reflection are observed.

The streaking observed in Figs. 6(c), 7(b) and 8(b) is

due to the small size of the plate-like particles. For all

the particles considered here, the strongest elongation of

diffraction spots is parallel to the normal of the platelets

and thus the streaking is always along one of the cubedirections. Therefore, the streaks observed in Figs. 6(c),

7(b) and 8(b) are lying in the plane of the Ewald sphere

and hence the maxima of the diffraction spots should be

positioned at the middle of the observed streaks.

3.5. Proposed structure of the GPB-II zones

Based on the work of Ratchev et al. [1], the obviousoption is to consider that the detected particles after 4

and 8 h of aging are in fact S00-phase. However, this

terminology seems to be inappropriate as we believe that

is most prominent when viewed along the [1 0 1]Al directions. Arrows


a phase termed as the S00-phase should be consistent with

the crystal structure originally proposed by Bagaryatsky

[3]. According to Bagaryatsky, the S00-phase is a mono-

clinic version of the S-phase with a slightly modified

orientation relationship compared with that of theS-phase. This orientation relationship can be under-

stood from the S-phase/matrix orientation relationship

[1 0 0]S//[1 0 0]Al, [0 1 0]S//[0 2 1]Al, but with the [0 1 0]Srotated about 4�–5.5� towards the [0 1 0]Al: As discussed

by Bagaryatsky, the S00-phase can homogeneously nu-

cleate in the matrix, which accounts for the modified

orientation relationship. In fact, very recent work by

Radmilovich et al. [21] and Majimel et al. [22] shows theexistence of ‘‘S-phase’’ with the modified orientation

relationship as proposed by Bagaryatsky, although it

was termed the S-phase with type-II orientation rela-

tionship instead of S00-phase. The work by Majimel et al.

[22] also shows that with increasing size of the S00-phase,the orientation relationships gradually change into that

known as the S-phase/matrix orientation relationship.

This change in the orientation relationship was previ-ously suggested by Shchegoleva and Shpektor [23].

Instead, we propose that the particles detected in the

present study be identified as GPB-II zones, rather than

suggest that the particles are S00-phase. We have shown

that the extra diffraction spots associated with the GPB-

II zones can be indexed in terms of a crystal structure

which is similar to the L10 ordering detected in the

earlier stages of aging. This GPB-II structure can besimply understood by considering three Al FCC unit

cells attached to each other. The lattice sites are popu-

lated by Mg and Cu atoms in such a way that layers of

alternating Mg rich and Cu rich (0 0 2) planes are

formed, just as they would to form the L10 structure.

However, along the x-direction at position a=2, the Mg

and Cu (0 2 0) layers are interchanged, resulting in a

creation of an APB. A schematic of this structure isshown in Fig. 9. A unit cell of such an ordered crystal

structure with incorporated APB along one lattice di-

rection is often referred to as a one dimensional super-

Fig. 9. Crystal structure of the GPB-II zone.

lattice. Another example of such a structure is that of

AuCuII [24]. Crystallographically, the GPB-II structure

has an orthorhombic C-centered lattice, with parame-

ters: a ¼ 1:212, b ¼ 0:404, c ¼ 0:404. The space group

is Cmmm and the atom positions, assuming perfectordering of the structure, are given in Table 2.

The orientation relationship of the GPB-II zone

with the matrix is [1 0 0]Al//[1 0 0]GPBII and [0 1 0]Al//

[0 1 0]GPBII. This orientation relationship and the crystal

symmetry results in six variants relative to the matrix.

As previously discussed, and shown in Fig. 6, four of

these variants are detectable along the [0 0 1]Al zone. To

fully describe the envisioned GPB-II zone, it is alsoimportant to mention that the [0 0 1]GPBII is aligned

along the long axis of the lath-like GPB-II zone, while

the [1 0 0]GPBII is parallel with the short edge. (It is in-

teresting to note that the measured width of the lath is

only about 1.0 nm, which is actually smaller than the

a – lattice parameter.) Based on these morphological

considerations, it is then possible to deduce that the

particles labeled A in Fig. 6(a), and the HRTEM ob-servations shown in Fig. 7, represent the variant viewed

along the [0 0 1]GPBII: The particles labeled as B in

Fig. 6(a), and the HRTEM observations shown in

Fig. 8, on the other hand, represent the variant viewed

along the [0 1 0]GPBII.

Based on the similarity in ordering and orientation of

the GPB-II and GPB zones, we may presume that the

GPB-II zones develop continuously from the GPBzones. Therefore, the GPB-II zones may not have well-

defined phase boundaries, unlike the GP/GPII zones in

the Al–Cu system, which can be ideally considered as

sharply bounded, mono-layer/double layers of Cu at-

oms. Nevertheless, the fact that we observe streaking in

the FFT spectra from the detected particles and strong

contrast in TEM observations indicates that the order-

ing is very localized, and thus the phase boundary isnarrow.

Simulated electron diffraction patterns for the crystal

structure of the GPB-II zones described in Fig. 9 show

very good agreement with the experimental observa-

tions. To show the similarity, results from the simula-

tions along three orthogonal directions are presented in

Fig. 10. The weaker spots in the patterns are solely due

to the ordering of Cu and Mg atoms while the strongerspots, on the other hand, arise due to the underlying

‘‘FCC arrangements’’ of all the atoms. Fig. 10(a) is a

Table 2

Atom positions in the GPB-II zone

X Y Z

Cu 0 0 0

Cu 1/6 0 0.5

Mg 0 0.5 0.5

Mg 1/6 0.5 0

Fig. 10. Simulated electron diffraction pattern of the GPB-II zone along three orthogonal directions. (a) On [0 0 1] zone. (b) Comparison of the [0 0 1]

simulation with the experimental observations. (c) On [0 1 0] zone. (d) Comparison of the [0 1 0] simulation with the experimental observations. (e) On

[1 0 0] zone.


diffraction pattern from the [0 0 1]GPBII zone. As seen

from the simulation, the spots due to the ordering of the

Cu and Mg atoms are very similar to the characteristic

spots of the first type, mentioned with reference to

Fig. 6(c). An overlay of observed and simulated patterns(magnified) is shown in Fig. 10(b). The simulation of

diffraction pattern from [0 1 0]GPBII is shown in

Fig. 10(c). In this case, the ‘‘ordering’’ diffraction spots

correspond very well with the diffraction spots described

as the second type in Fig. 6(c). The differences between

the simulation and experiment are due to the streaking

of the spots, as seen in the overlay shown in Fig. 10(d).

In fact, the streaking in the experimental observationcan be anticipated, as the lath-like GPB zones are only

1 nm wide and about 6–8 nm in length in the projected

direction. Nevertheless, the positions of the spots agree

well. Finally, the simulated pattern from the [1 0 0]GPBII

zone is shown in Fig. 10(e). The fact that the ‘‘ordering’’

spots from this zone have not been observed is probablydue to the fact that extra reflections are relatively weak,

even for these simulations based on an ‘‘infinite crystal’’,

and also because of the shape of the GPB-II zones

relative to the viewing direction.

The HRTEM observations of the GPB-II zones were

also compared with image simulations. For the image

simulations, we built an aluminum FCC crystal of a size

25� 25� 65 (unit cells) with the GPB-II zone insertedinside this crystal. The sample thickness of 65 unit cells

Fig. 11. HRTEM image simulation of the GPB-II incorporated into the Al matrix and viewed along the [0 0 1]GPBII. (a) At the Scherzer defocus.

(b) At defocus corresponding to the first indirect transfer condition. (c) FFT corresponding to (a).


represents about 26 nm, which is greater than half of theextinction distance of the (2 0 0)Al on the h001iAl zone.

Thus, at the Scherzer defocus, the bright spots corre-

spond to the atomic columns, whereas at the first indi-

rect transfer conditions, the bright spots correspond to

the spaces between atomic columns.

For the simulations, the GPB-II zone was situated

approximately in the middle of the matrix and oriented

such that either the [0 0 1]GPBII or [0 1 0]GPBII zone axiswas aligned along the viewing direction. The size of the

zone was 1:0� 2:4� 6:0 nm, which is approximately

equal to the experimentally determined size. If simulated

along the [0 0 1]GPBII direction, the projected area was

thus 1.0 · 2.4 nm, while if simulated along the

[0 1 0]GPBII, the projected area was 1.0 · 6.0 nm. It has to

be mentioned that the simulated GPB-II zone was not

built as a compositional modulation, but rather ashaving a well-defined order and sharply bounded inter-

faces. This assumption may represent a shortcoming of

this exercise, as described below.

The HRTEM image simulations shown in Fig. 11

represent the case when the GPB-II zone was oriented

along the [0 0 1]GPBII. The image shown in Fig. 11(a)

corresponds to the Scherzer defocus conditions, whereas

the image shown in Fig. 11(b) corresponds to the de-focus value at the first indirect transfer condition. One

Fig. 12. HRTEM image simulation of the GPB-II zone incorporated into the

(b) At defocus corresponding to the first indirect transfer condition. (c) FFT

can clearly see that the Scherzer defocus condition givesrise to a stronger contrast from the embedded GPB-II

zone. It must be noted, however, that such a strong

contrast at Scherzer defocus is seen only for images with

thickness greater than a half of the extinction distance of

the (2 0 0)Al, which is the thickness in the present simu-

lations (26 nm). On the other hand, at a thickness below

half of the extinction distance, the first indirect transfer

condition gives much stronger contrast from the em-bedded GPB zone. For the simulated condition, the

GPB-II zone is seen as a pattern of alternating brighter/

darker atomic columns as shown in Fig. 11(a). Indeed

such an alternating pattern of bright spots was seen in

the experimental images, an example of which is shown

in Fig. 7(a) (indexing of the crystallographic directions

in Fig. 11 can be directly applied to Fig. 7(a)). The

Fourier spectrum of the simulated image in Fig. 11(a)shows the presence of extra diffraction spots, as shown

in Fig. 11(c). This spectrum compares very well with the

FFT spectrum shown in Fig. 7(b).

The HRTEM image simulations shown in Fig. 12

represent the case when the GPB-II zone is oriented

along the [0 1 0]GPBII. Two simulated images, the first

corresponding to the Scherzer defocus and the second

corresponding to the first indirect transfer defocus, areshown in Fig. 12(a) and (b). As already seen in the

Al matrix and viewed along the [0 1 0]GBPII. (a) At the Scherzer defocus.

corresponding to (a).


previous simulations, the Scherzer defocus conditions

give rise to a stronger contrast for the given thickness. In

fact, at the indirect transfer conditions, the particle is

virtually invisible (such changes in the particle visibility

with a change of the defocus were confirmed experi-mentally). Comparing the contrast from the simulations

at Scherzer defocus and experimental observation we see

good correlation with respect to the presence of alter-

nating brighter/darker spots. Nevertheless, the experi-

mental image shown in Fig. 8(a) does not exhibit the

obvious presence of an APB, unlike the simulated im-

ages of Fig. 12(a) (the APB can be seen by observing at

the glancing angle along the [1 0 1]Al). This discrepancymay be related to the GPB-II zones having rather diffuse

phase boundaries, as discussed above, together with the

fact that the strongest ordering/segregation may be

confined to the central region of the GPB zone, which

was not accounted for in the simulations. The FFT

transformation of the simulated image from the Scher-

zer defocus condition is shown in Fig. 12(c). It is seen

that the character of the streaks and also the degree ofstreaking is quite similar to the experimental observa-

tions, shown in Fig. 8(b).

4. Discussion

The microstructural observations described above

show that a direct detection of the GPB zones in the Al–Mg–Cu alloys is possible only after several hours aging

at 180 �C. Such detection times are fairly consistent with

the published data on the Al–Cu–Mg alloys. However, if

the differences between the GPB zones in these two

systems are examined more closely, alloy dependent

characteristics can be identified. First is the morphology,

which has been reported numerously to be rod-like for

the Al–Cu–Mg alloys [5,25], while in the present ob-servations we find a lath-like morphology for the de-

tectable GPB-II zones. The character of the diffraction

spots associated with the presence of the GPB zones is

also different. While in both systems the GPB zones give

rise to diffraction streaks positioned around the forbid-

den {1 1 0} reflections, it is known from X-ray diffrac-

tion studies that half of the streaks in the electron

diffraction pattern of the Al–Cu–Mg system are due todouble diffraction. On the other hand, if the presently

proposed structure of the GPB zone for Al–Mg–Cu

system is correct, then no streaks would be due to

double diffraction. Another system-unique characteristic

is that the GPB zones appear to be coherent with the

matrix for the Al–Mg–Cu alloys, based on our HRTEM

observations, whereas observations made on the Al–Cu–

Mg indicate a tendency for faceting on the {1 2 0} and{1 1 0} planes [5].

Thus on the basis of these differences, it appears that

the GPB-II zones analyzed in this work are specific for

the Al–Mg–Cu alloy. These differences may be due to

the Cu/Mg ratio, as well as the presence of a small

amount of Si, which is known to have a significant in-

fluence on the hardening behavior in Al–Cu–Mg alloys.

From the published work by Wilson and Partridge [26]and Hutchinson and Ringer [27], it has been realized

that the presence of Si significantly increases the stability

of GPB zones in the Al–Cu–Mg alloys. This stabiliza-

tion effect may be more pronounced in the Al–Mg–Cu

alloys, and may eventually lead to the formation of

modified GPB zones, probably due to the lower Cu

levels in these alloys. Note that the location of the Si in

the GPB zones and its precise role in the Al–Mg–Cualloys remain to be determined.

The microstructure observations reported in this

study were made on samples aged for 1 h and up to 8 h.

The 1-h condition is significantly beyond the few min-

utes required for the rapid hardening response in these

alloys [1]. Thus, based on the present microstructure

observations, it is not yet possible to explain conclu-

sively the origin of the rapid hardening. Nevertheless, wecan speculate that the origin of the hardening is due to

formation of the L10 ordered GPB zones that were de-

tected after 1 h of aging. Immediately following the ra-

pid hardening, such zones would probably have a

smaller degree of ordering, thus making their detection

very difficult. The hardening at this stage of aging could

therefore be attributed to ‘‘destruction’’ of the ordering

associated with zones formed. In fact, since the GPBzones have L10 ordering, the dislocations with Burgers

vectors perpendicular to the c-axes of the zones (see

Fig. 2) would be influenced to lesser extent by the

presence of the GPB zones. In the later stages of aging,

when the GPB-II zones grow larger and develop APBs,

all 1/2h110i dislocations would have a pronounced ef-

fect on the ‘‘destruction’’ of the order in the zones. This

enhanced interaction between dislocations and the GPB-II zones may explain the gradual increase of the hard-

ness following the initial rapid hardening regime.

5. Conclusions

After 1 h of aging time, HRTEM observations and

focal series reconstructions revealed the presence of{1 1 0} super-lattice reflections in the FFT spectra, sug-

gesting the presence of L10 ordered, rod-like regions in

the matrix. These regions are considered to be GPB

zones. Image processing performed on images contain-

ing {1 1 0} reflections and also image simulations of

microstructures with rod-like GPB zones, suggested that

the GPB zones are not sharply bounded regions, but

instead are compositional modulations. Although pres-ently unsupported by chemical analysis, it is proposed

that the zones consist of alternating (2 0 0) layers of Mg

and Cu atoms. After 4 h of aging, some of the GPB


zones were found to grow to larger sizes and are iden-

tified as GPB-II zones. Based on the HRTEM obser-

vations and the diffraction information, it is proposed

that the crystal structure of the GPB-II zones is dis-

tinctly different from that of the GPB zones, with theformer having an orthorhombic lattice and space group

Cmmm. The GPB-II zone is fully coherent with the

matrix, and has the orientation relationship [1 0 0]Al//

[1 0 0]GPBII and [0 1 0]Al//[0 1 0]GPBII, and the morphology

is lath-like.

Acknowledgements

The authors acknowledge the financial support of

Alcan International Ltd. and the Center for the Accel-

erated Maturation of Materials (CAMM) at the Ohio

State University.

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a hrtem study of metastable phase formation in al–mg–cu alloys during artificial aging

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