The origins of room temperature hardening of Al–Cu–Mg alloys

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Materials Science and Engineering A 387–389 (2004) 222–226

The origins of room temperature hardening of Al–Cu–Mg alloys

M.J. Starink ∗, N. Gao, J.L. Yan

Materials Research Group, School of Engineering Sciences, University of Southampton, Southampton, SO17 1BJ, UK

Received 26 August 2003; received in revised form 14 January 2004

Abstract

The most commonly cited mechanisms for the rapid age hardening of Al–Cu–Mg alloys at about 100–200 ◦C are hardening by Guinier–

Preston B zone formation, formation of Cu–Mg co-clusters and a dislocation–solute interaction mechanism. New experiments on ageing–deformation–ageing cycles at room temperature indicate that no substantial additional age hardening occurs with the addition of deformation to

the cycle, and hence, a dislocation–solute interaction mechanism appears unlikely. Instead, strengthening due to modulus hardening generated

by the difference in shear modulus of Cu–Mg co-clusters and matrix is proposed as the main strengthening mechanism for room temperature

hardening.

© 2004 Elsevier B.V. All rights reserved.

Keywords: Clusters; Calorimetry; Precipitation; Al–Cu–Mg alloys; Natural ageing

1. Introduction

During the rapid hardening and the room temperature,

hardening in Al–Cu–Mg based alloys [1–3], no distinct pre-cipitate phase can be detected by conventional transmission

electron microscopy (TEM) [4], but differential scanning

calorimetry (DSC) experiments clearly show a dissolu-

tion effect evidencing that a metastable pre-precipitate has

formed [5]. High-resolution electron microscopy (HREM)

images fail to show any ordered structures, which indicates

the pre-precipitate in this stage has a random arrangement.

Results from several atom probe field ion microscopy (AP-

FIM) and three-dimensional atom probe (3DAP) studies

reveal Cu–Mg co-clusters of typically 1 nm diameter (10–40

atoms) which by some authors are held responsible for the

rapid hardening reaction [6,7]. Information on the com-

position of Cu–Mg clusters is limited. Ringer et al. [6,7]

observed that in a high purity Al–1.7 at.% Mg–1.1 at.% Cu

aged for 5 min at 150◦C clusters with high Cu:Mg ratios

existed, whilst Gao et al. [8] found that in an Al–2.78 wt.%

Cu–0.44 wt.% Mg and in an Al–2.78 wt.% Cu–1.05 wt.%

Mg (both alloys are commercial purity with Mn additions)

the Mg:Cu ratio of the clusters varied from one cluster to

another. They also contain vacancies [4].

∗ Corresponding author. Tel.: +44-238-059-5094;

fax: +44-238-059-3016.

E-mail address: [email protected] (M.J. Starink).

Reich et al. [9] proposed an alternative mechanism for the

rapid hardening in which Cu and Mg solute atoms segre-

gate to the dislocations (especially dislocation loops), lock-

ing dislocations and increasing the hardness. This suggestionwas based on three main observations from experiments on

an Al–1.7 at.% Mg–1.1 at.% Cu alloy and related alloys: (i)

no clusters were detected by 3DAP during the rapid harden-

ing stage in Al–1.7 at.% Mg–1.1 at.% Cu; (ii) further, rapid

age-hardening occurred following deformation after ageing;

(iii) in quenched and room temperature aged Al–Cu–Mg

based alloys many dislocation loops are present [4]. In sup-

port of this, positron spectroscopy data [4,10] was inter-

preted to suggest that very rapidly either vacancy–Mg com-

plexes [4] or vacancy clusters [10] f orm, which is followed

by the formation of the hardening dislocation–vacancy ag-

gregates.

Ratchev et al. [11,12] suggested that the initial rapid

hardening in a commercial purity alloy with low Cu con-

tent (Al–4.7 at.% Mg–0.3 at.% Cu) is due to a combination

of cluster formation and heterogeneous formation of S

phase on dislocation helices but in most pre-1990s work

early hardening is attributed to Cu and Mg containing GPB

zones [13], which would have a rod-like shape [10]. It has

been suggested that the distinction between Cu–Mg clus-

ters (or vacancy–Cu–Mg complexes) and GPB zones can

be made on the basis of size, shape, composition, degree of

order, orientation and structure [6]. The vacancy–Cu–Mg

complexes were considered as precursor of GPB zones

0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved.

doi:10.1016/j.msea.2004.01.085



M.J. Starink et al. / Materials Science and Engineering A 387–389 (2004) 222–226 223

[4,10]. However, 3DAP work shows no difference between

zones and clusters except different sizes corresponding

to the ageing temperatures or times [8]. Hence, on bal-

ance, the evidence for the existence of GPB zones that

are responsible for the early hardening reaction and have

internal order and/or a distinct shape is weak and follow-

ing other work [4,6,7,9], we will not use the term GPBzones for the structures involved in the early hardening

reaction.

In this paper, we will present new experiments on room

temperature ageing of two commercial purity Al–Cu–Mg–

Mn alloys which are aimed to test whether mechanisms pro-

posed in the literature can explain the room temperature

hardening in these alloys, and finally present the outline of

a model which can explain early hardening in a quantitative

manner.

2. Experimental procedures

Two commercial purity Al–Cu–Mg–Mn alloy plates with

Cu:Mg ratio close to 1 (atomic ratio) have been studied,

the compositions are reported in Table 1. For both alloys,

99.90wt.% aluminium was used as base to ensure compa-

rable impurity contents that reflect commercial alloys of

relatively high purity. Fe and Si contents of the alloys are

expected to be between 0.02 and 0.03 at.%. All samples

were freshly solution treated at 495 ◦C and subsequently

quenched into water at room temperature before the age-

ing treatments commenced. Room temperature ageing was

performed at about 25 ◦C.

Aged samples were studied in a Perkin-Elmer Pyris 1calorimeter. Samples were discs (5 mm diameter and ap-

proximately 1 mm thickness) that were machined prior to

solution treatment. DSC runs were started immediately af-

ter introduction of the sample at 0 ◦C. Scanning over the

temperature range 5–540 ◦C at a constant heating rate of

10 K/min was performed. Baseline correction procedures are

described elsewhere [14].

Vickers hardness tests were performed on quenched

and naturally aged specimens and samples that were sub-

sequently deformed and naturally aged. The hardness

values were obtained from surfaces ground with #1200

grade SiC-paper. Four indentations were made on each

specimen with a 20-kg or a 1-kg mass and a mean hard-

ness is reported. Tensile tests were carried out according

to ASTM standard E-8. Specimens were taken in the

longitudinal ( L) orientation (i.e. parallel to the rolling

direction).

Table 1

Compositions of the alloys (at.%)

Alloy Cu Mg Mn

Al–1.2 Cu–1.2 Mg 1.21 1.19 0.21

Al–1.9 Cu–1.6 Mg 1.89 1.56 0.21

3. Results

The hardness evolution during ageing at 25 ◦C of solu-

tion treated samples shows hardening within a few hours

(Fig. 1). (The hardness values in the freshly quenched con-

dition are 72 for the Al–1.2 Cu–1.2 Mg alloy and 91 for

the Al–1.9 Cu–1.6 Mg alloy.) Additional hardening duringartificial ageing after deformation has been suggested to be

indicative of a solute dislocation interaction being responsi-

ble for the rapid ageing in high purity Al–Cu–Mg [9]. We

performed an experiment that attempts to test this for age-

ing at room temperature for our alloys. We solution treated

and quenched our Al–1.9 Cu–1.6 Mg alloy and aged it at

room temperature for a week and measured the hardness.

Subsequently, we deformed the sample by 2.5% uniaxial

compression and measured the hardness on subsequent nat-

ural ageing for 10–10,000 min (∼1 week). The results show

that the deformation had increased the hardness by about

10 HV, which would be caused by the network of dislo-

cations introduced during deformation, but further ageingup to 1 week does not result in a further increase in hard-

ness (data not presented). Thus, for ageing at room tempera-

ture, these experiments provide no indication of a substantial

dislocation–solute interaction that causes age hardening.

To further investigate the latter point we performed cy-

cles of thermo-mechanical treatments on Al–1.9 Cu–1.6 Mg

tensile samples. Typical tests consisted of: solution treat-

ment and quench, 7 days ageing at room temperature, ten-

sile deformation to 1.5 or 3.5%, room temperature ageing

for times between 5 min and 16 h, and repeats of the last

two (tensile deformation, room temperature ageing) for up

to five cycles. An example is presented in Fig. 2. From thesecurves, we determined the stress at the end of the strain-

ing of a cycle when the stress is released σ (εr) as well

as, by extrapolation, σ (εr + 0.1%), the procedure is illus-

trated in Fig. 3. We compared this with the 0.1% proof

stress, σ 0.1%, determined in the subsequent straining cycle,

which was started after natural ageing for 4 or 16 h. The

data for two samples that underwent a very similar sequence

Fig. 1. Age hardening of the Al–1.2 Cu–1.2 Mg and the Al–1.9 Cu–1.6

Mg alloy at 25 ◦C.



224 M.J. Starink et al. / Materials Science and Engineering A 387–389 (2004) 222–226

Fig. 2. Stress–strain curves of repeated straining interrupted for 4 or

16 h room temperature ageing for the solution treated and naturally aged

Al–1.9 Cu–1.6 Mg alloy. (See Table 2.) The flow stress increases with

each repeated straining.

Fig. 3. Illustration of determination of stress at permanent strain εi + 0.1%during a straining cycle (indicated by I), which was interrupted at per-

manent strain εi, and σ 0.1% at the subsequent straining (indicated by II).

of ageing–straining–ageing are presented in Table 2 (which

also contains the average difference of these two stresses,

σ ). This data indicates that σ is always very small. In

total five tensile samples were tested in this manner, with

tensile deformation in the cycles 1.5 or 3.5% and room tem-

perature ageing for times between 5 min and 16 h. The ab-

solute magnitude of σ was always less than 7 MPa and

on average 3 MPa. This shows that there is no substantial

Table 2

The stress at the end of the straining of a cycle extrapolated to σ (r

+ 0.1%) (as illustrated in Fig. 3) and the 0.1% proof stress, σ 0.1%,

determined in the subsequent straining cycle

I → II II → III III → IV IV → V V → VI

t a (h) 4 4 16 4 4

σ (εr + 0.1%) 350 386 407 415 421

σ 0.1% 356 388 405 415 418

σ 6 2 −2 0 −3

∆σ (average) 5 2 −2 −1 −3

t a is the natural ageing time applied between subsequent straining.

Fig. 4. DSC curves of the Al–1.2 Cu–1.2 Mg alloy after ageing for several

intervals at 25 ◦C.

age hardening after dislocations are introduced during de-

formation in these alloys when ageing–deformation–ageing

is conducted at room temperature, and this is in contrast

to results for ageing–deformation–ageing of a high purityAl–Cu–Mg alloy at 150 ◦C [9].

DSC curves for the Al–1.9 Cu–1.6 Mg alloy after ageing

at 25 ◦C are presented in Fig. 4. Four effects (two exothermic

and two endothermic), numbered I–IV, are evident. Effects

III and IV are discussed in detail elsewhere [8]; they can

be identified as due to S phase formation (Effect III, in the

range of 230–320 ◦C) and S phase dissolution (Effect IV,

in the range of 330–470 ◦C). These effects are unrelated to

rapid ageing. As discussed below, Effect I is thought to be

due to the formation of Cu–Mg co-clusters. Comparison of

Fig. 1 and Fig. 4 shows that the progress of this reaction

coincides with the increase in hardness. Effect II is thoughtto be due predominantly to dissolution of clusters.

For all DSC curves, the exothermic heat evolution Qcl

due to cluster formation was measured. In Fig. 5, the change

in Qcl with room temperature ageing is plotted as a func-

tion of the change in hardness during room temperature age-

ing. This figures reveals that that the increase in hardness

exactly coincides with heat release at room temperature.

4. Discussion

For alloys aged to the very start of the plateau strength

stage, Cu–Mg clusters have been evidenced by APFIM [6,7],

and in a subsequent paper [15], we will present APFIM

evidence showing the formation of co-clusters during the

age hardening and the concomitant heat release, whilst no

GPB zones are observed. Thus, whilst the present short paper

does not contain evidence that cluster formation rather than

zone formation is responsible for the age hardening (Fig. 1)

and the exothermic Effect I (Fig. 4), and this issue remains

controversial, we will here discuss the room temperature

age hardening in terms of cluster formation. (In alternative

interpretations/terminologies, the structures formed during

Effect I are termed GPB zones [13,16]).



M.J. Starink et al. / Materials Science and Engineering A 387–389 (2004) 222–226 225

Fig. 5. Increment in hardness during room temperature ageing vs. the change in heat content in the DSC thermogram for the cluster formation Effect I

compared to un-aged samples for the Al–1.2 Cu–1.2 Mg () and the Al–1.9 Cu–1.6 Mg (×) alloys.

For our commercial purity Al–Cu–Mg alloys.

(i) natural ageing after extensive natural ageing followed

by deformation does not give rise to significant addi-

tional age hardening;

(ii) room temperature hardening of freshly solution treated

material coincides with a substantial heat release;

(iii) DSC shows only one exothermic reaction, and hence, it

appears that the reaction involved in room temperaturehardening is the same as the reaction involved in rapid

hardening at 100–200 ◦C.

Point (ii) would suggest that room temperature harden-

ing is a direct result of the formation of clusters, whilst

point i suggest that any interaction between solute and dis-

locations introduced by deformation is not the main cause

for this hardening. This interpretation is different from

the dislocation–solute hardening mechanism that was sug-

gested [9] to operate for the rapid hardening of high purity

Al–Cu–Mg at 150 ◦C. A further reasoning indicating that

this is not the main the cause of room temperature hardening

or rapid hardening is the following. TEM data shows that

in quenched and room temperature aged Al–Cu–Mg based

alloys the inter-loop spacing of dislocations is in the order

of 100–200 nm [4], which is almost an order of magnitude

larger than the precipitate spacing in artificially aged high

purity Al–Cu–Mg (see Fig. 3d in Ref. [9]), commercial

purity Al–Cu–Mg [14] and Al–Cu–Mg–Li (Fig. 4 in Ref.

[17]). The strengthening increment due to the formation of

these precipitates is similar to that of the rapid hardening,

and, hence, the loops (with or without solute associated

with it) are unlikely to provide a major contribution to

hardening. Non-shearable objects with spacing 100–200 nm

would provide a yield strength increase of 100–200 MPa

(40–80 HV) [18]. Hence, measured hardness increases could

only be caused by dislocation loop–solute complexes if

these would be non-shearable. However, this appears quite

unlikely. The strong composition dependency of plateau

hardness (e.g. Fig. 1, [2]) also suggests that non-shearable

dislocation loops can not be the main cause of hardening.

Fig. 5 strongly suggests a direct correlation between clus-ter formation and age hardening. As clusters have no internal

ordering, order strengthening and stacking fault strengthen-

ing will not play a role and only chemical hardening and

modulus hardening are expected to occur. As cluster–matrix

interfaces are diffuse, chemical hardening is expected to be

relatively small, and hence we concentrate on modulus hard-

ening only. Modulus hardening has been approximated from

the difference between the shear moduli of the matrix and

the clusters, G, and the volume fraction of the clusters f cl

according to [19]:

τ cl=

G

4π√ 2f

1/2

cl

(1)

where τ cl is the increment in critical resolved shear stress

(CRSS) due to clusters. The yield strength of the alloys is

related to the total CRSS by factor M (sometimes referred

to as the Taylor factor [20]):

σ y = σ i +M(τ qss +τ

q

cl)1/q (2)

where τ ss is the increment in CRSS due to solution

strengthening [14], and the superposition exponent q de-

pends on the relative strengths of the two types of obstacle

(1 < q< 2) [21,22].



226 M.J. Starink et al. / Materials Science and Engineering A 387–389 (2004) 222–226

To calculate strengthening increments on ageing we need

to estimate solution strengthening due to Cu and Mg atoms,

estimate the solvus of the clusters, estimate q, and estimate

the composition (especially the Al content) of the clusters

to obtain f cl. These data can be estimated from earlier work

[8,14,23], and it can be shown that the measured increments

in hardness of both alloys in Fig. 1 can be explained wellprovided G is about 3 GPa. Direct data on Gcl is virtually

impossible to obtain, and hence no direct verification of the

latter is possible. Nevertheless, this value for G is consid-

ered reasonable, and the present semi-quantitative analysis

indicates that modulus strengthening by Cu–Mg clusters is

the most probable cause of room temperature hardening.

5. Conclusions

Experiments on commercial purity Al–Cu–Mg alloys with

Cu:Mg ratios close to 1 have shown:• room temperature ageing after extensive natural ageing

followed by deformation does not give rise to significant

additional age hardening;

• room temperature age hardening coincides with a signif-

icant heat evolution.

An analysis of literature data and the new experiments in-

dicates that room temperature hardening can be due to mod-

ulus hardening due to the formation of Cu–Mg co-clusters.

Acknowledgements

The authors would like to acknowledge the financial sup-

port from EPSRC, Airbus UK and QinetiQ for parts of this

work.

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The origins of room temperature hardening of Al–Cu–Mg alloys

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