Formability and failure mechanisms of AA2024 under hot forming conditions
L. Wang1, M. Strangwood2 D. Balint1, J. Lin*1 and T. A. Dean3
1
2
Department of Mechanical Engineering, Imperial College London
SW7 2AZ, UK
School of Metallurgy and Materials, University of Birmingham, Edgbaston,
Birmingham B15 2TT, UK
3
School of Mechanical Engineering, University of Birmingham, Edgbaston,
Birmingham B15 2TT, UK
Abstract
Aluminium alloy 2024 (AA2024) is extensively used as a structural material in the aircraft industry
because of its good combination of strength and fatigue resistance. However, complex shaped
components, particularly those made from sheet, are extremely difficult to form by traditional cold
forming due to its low ductility at room temperature. A possible solution of this problem is to form
sheet workpieces at elevated temperature. The aim of the work described in this paper is to
determine the relationship between formability and temperature for AA2024 by conducting a series
of isothermal tensile tests at elevated temperatures ranging from 350 to 493 °C. Ductility of
AA2024 was found to increase gradually with increasing temperature up to 450 °C, followed by a
sharp decrease with further increase in temperature. So-called cup tests confirmed that the
formability of AA2024 is very high at a temperature of about 450 °C. Fracture surfaces and
longitudinal sections of formed samples were examined by scanning electron microscope. It was
found that fracture occurred in three different modes depending upon the temperature, and the sharp
decrease in ductility when the temperature exceeds 450 °C was caused by softening of grain
boundaries by solute enrichment (at higher heating rates liquation may be involved) and softening
of the matrix around inclusion particles.
Keywords: AA2024; Fracture mechanisms, Hot forming, Formability
* Corresponding author. Tel.: +44 020 7594 7082; fax: +44 020 7594 7017;
Email address: jianguo.lin@imperial.ac.uk (J. Lin).
1
1. Introduction
Aluminium alloy 2024 (AA2024) is widely used for structural applications in the aerospace
industry, due to its good combination of strength and fatigue resistance. Age hardening markedly
affects the microstructure and mechanical behaviour of AA2024. After appropriate heat treatment,
finely dispersed precipitates are obtained and a high strength to weight ratio is achieved. Low
ductility, however, is typically a limitation of AA2024, making complex-shaped parts difficult to
form using traditional cold forming processes. Hence some thin-walled and complex-shaped parts
are machined from solid blocks of metal, which can incur up to 90% material wastage for some
applications, with corresponding cost and energy wastages.
In recent years, efforts have been made to improve the strength and ductility of AA2024 [1-9] by
developing advanced processing routes. For example, following severe plastic deformation (i.e.
ECAP [1, 4, 5], cryo-rolling [3] or high-pressure torsion processes [9]) and suitable ageing, highly
dispersed nanometre-sized precipitates can be obtained. Due to the low density of dislocations
present after ageing, and the strong pinning effect and aggregation of precipitates to dislocations, an
elongation of 18% can be achieved at room temperature [3]. However, the use of extreme strain
levels is limited to very small samples, and is currently not a suitable alternative for bulk materials
processing. On the other hand, pre-homogenisation treatment [7, 8] might be an effective way of
improving the ductility of AA2024, which, however, sacrifices the material strength. On the
contrary, the material strength could be enhanced by adding SiCp into AA2024, but the ductility of
the resulting composite has been found to decrease [6].
As for most metallic alloys, the ductility of AA2024 increases with increasing deformation
temperature [6]. However, hot stamping is seldom used for forming complex-shaped parts because
the desired microstructure (which has been fixed in the sheet prior to forming) tends to be destroyed
at high temperatures, thereby reducing the mechanical performance. Furthermore, if heat treatment
to restore strength is carried out after stamping, the parts tend to lose their shape due to thermal
distortion.
Recently, a new technique, Hot Forming and cold-die Quenching (HFQ), has been developed [10].
The basic idea of this novel process is to (i) heat the sheet metal to its solution heat treatment (SHT)
temperature, at which its ductility is expected to be maximal, and (ii) simultaneously form and
quench the sheet by using cold dies. After forming, the workpiece material is held within the dies
for roughly 5 to 6 seconds in order to reduce its temperature rapidly to approximately 100 °C and
freeze the microstructure as a supersaturated solid solution (SSSS). If a heat-treatable, aluminiumbased alloy is used, the part can then be aged to obtain full strength. The feasibility of this novel
2
process has been demonstrated for AA6XXX in [11]. This technique has gained favour over
traditional forming methods because it produces parts with high formability, negligible springback,
rapid processing and efficacious mechanical properties.
In AA2024, copper is the most effective strengthening constituent. The age hardening effect can be
increased by increasing copper content up to 6 wt % [12]. Magnesium is used in combination with
copper to accelerate the age hardening effect at room temperature. The equilibrium precipitate
phases for this system are mainly CuMgAl2 (S phase) and CuAl2 (θ phase) [13], although CuAl2 is
less often observed in the alloy compared to CuMgAl2. Both of these phases are largely soluble
during SHT. In addition, manganese has a marked strengthening effect, because it can influence
material properties through the formation of intermetallics that provide Zener drag and limit grain
size. The intermetallic compounds (IMCs) in AA2024 are very complicated [14-16], and the shape,
size and chemical composition of the IMCs vary remarkably, which are dictated by the processing
route [16]. A range of IMC compositions have been reported for AA2024, such as
Al20(Cu,Fe,Mn)5Si(Al8Fe2Si), periphery, Al10(Cu,Mg), (Al,Cu)93(Fe,Mn)5(Mg,Si)2, Al3(Cu,Fe,Mn),
Al2Cu (θ phase), CuMgAl2 (S phase) and Al7Cu2Fe [15]. These are normally formed during the
casting or homogenization stages of processing, i.e. high temperature stages; these are normally not
soluble during SHT.
A number of studies have been carried out on the fracture mechanisms of AA2024 at room
temperature [17-20] and under superplastic or creep age forming conditions [8, 9, 19, 21], i.e. high
temperature and low strain rate conditions. However, little has been published on the fracture
mechanisms of AA2024 under hot forming conditions, i.e. high temperature and high strain rate
conditions. The objective of the present study is to address the above issue; to determine the
conditions for which maximum ductility arises, and to investigate the flow behaviour and fracture
mechanisms of AA2024 when subjected to HFQ.
2. Experimental procedure
2.1. Materials
The composition of the AA20241 used in the current research is listed in Table 1. The 2 mm thick
AA2024 sheets were in a T3 condition (solution heat treated, quenched and stress-relieved by cold
stretching).
1
Provided by Airbus S.A.S; contains small amounts of proprietary constituents.
3
Table 1. Chemical composition of AA2024 (wt %) [12].
Si
Fe
Cu
Mn
Mg
Zn
Al
0.5
0.5
3.8-4.9
0.3-0.9
1.2-1.8
0.25
balance
2.2. Testing programme
Gleeble thermomechanical testing
Isothermal tensile tests were conducted on a Gleeble 3800 thermomechanical simulator, which can
heat a specimen by direct resistance heating at a rate as high as 10,000 °C/s. On the other hand, the
specimen was held by two high thermal conductivity grips clamped by two jaws, each with an
embedded water-cooling system. This makes the Gleeble 3800 also capable of high cooling rates. In
addition, a pair of thermocouples was welded on the specimen to provide signals for the accurate
feedback control of the specimen temperatures. The combined effect of the efficient heating and
cooling system and the accurate thermal control system ensures that the temperature of the
specimen can be controlled accurately.
Due to the possibility of overheating, specimens were first heated to a temperature 25 °C lower than
the target temperature at a heating rate of 50 °C/s, then were further heated to the target temperature
at a rate of 5 °C/s. Isothermal tensile tests to failure were performed as soon as the target
temperature was reached, at a constant strain rate of 1 s-1, representative of typical strain rates in hot
forming processes. Table 2 lists the isothermal tensile tests that were conducted in the present study.
Additional tests were conducted at 350 and 450 °C to verify the repeatability of the tests; good
agreement from one test to another (mean deviation less than 5%) of both ductility and flow stress
was achieved.
Table 2. Isothermal Gleeble 3800 tensile tests (bold indicates select temperatures where
repeatability of the tests was verified).
Temperature (°C)
Strain Rate 350 360 370 380 390 400 410 420
1 s-1
430 440 450 460 470 480 487 493
4
Formability testing
Formability tests (so-called cup tests) were performed on a 25 tonne ESH high-speed (ram speeds
up to 5 m/s) press. The formability test rig was designed to be a portable, integrated structure onto
which either a spherical head or flat head punch could be mounted. A high speed camera and prism
were employed during the tests to record the diameter evolution of the central hole. Fig. 1 shows the
formability test rig in place on the ESH high-speed press. In Table 3, the dimensions of the test
piece and punches, and main process parameters for the forming tests are listed. A central hole, 16
mm in diameter, was cut in the centre of the test pieces, for the qualitative evaluation of plastic
deformation and verification of modelling results.
Microstructure examination
Precipitate and inclusion populations have a strong effect on the ductility of AA2024. A Hitachi
S3400-N scanning electron microscope (SEM) equipped with a Gatan H1002 heating stage (750 °C
peak temperature) was used to examine the microstructure of AA2024; the area close to the fracture
surface was examined in order to reveal the fracture mechanisms at elevated temperatures. The
operating voltage was 15 kV. Longitudinal sections, as well as the fracture surfaces, of the samples
deformed at 350, 450 and 493 °C were examined. Precipitate evolution and inclusion distribution
were examined at a temperature of 493 °C using a heating stage installed in the SEM for in-situ
observations. All the samples were polished prior to SEM examination.
Table 3. Dimensions of the test piece and HFQ punches, and the main process parameters.
Test piece width × length × thickness (mm)
170 × 170 × 2
Central hole diameter (mm)
16
Hemispherical head punch (HH) diameter (mm)
80
Flat head punch (FH) diameter (mm)
80
Punch temperature (°C)
20 (room temperature)
Ram speed (mm/s)
170
486
486
HH Punch displacement (mm)
26 ± 3
26 ± 3
36 ± 5
FH Punch displacement (mm)
21 ± 2
21 ± 2
-
Initial test piece temperature (°C)
493 ± 5
500 ± 5
Soaking time (hours)
0
1
Forming temperature (°C)
450 ± 10
493 ± 10
5
3. Results
3.1. Dependence of ductility on temperature
As the alloy is hot formed in the HFQ process, the evolution of the tensile ductility at elevated
temperatures, especially around the SHT temperature, is of particular importance. In the present
research, a series of tensile tests, shown in Table 2, were performed on a Gleeble 3800
thermomechanical simulator at temperatures ranging from 350 to 493 °C. Fig. 2 shows the results
obtained from the tensile tests. It was found that the maximum failure strain of AA2024 is in excess
of 1.1 (equivalent to an elongation of 200%), occurring at 450 °C, hence suggesting a temperature
window for high formability of AA2024. This behaviour is consistent with increased thermal
activation leading to greater atomic mobility, dislocation motion and recovery. However, at 493 °C,
a typical SHT temperature for AA2024, the ductility was very poor, only about one-tenth of the
peak value. Such a steep decrease in ductility is quite different from other heat treatable aluminium
alloys, such as AA6XXX [22]. Usually, the ductility of aluminium alloys at the SHT temperature
should be improved, as the amount of precipitates in the matrix is significantly decreased relative to
lower temperatures, hence the chance of void nucleation should be significantly diminished [17]
which would improve the ductility of the material. In the Gleeble tests, melting of precipitates and
inclusions (see section 3.4) appears to cause the drop in ductility at temperatures in excess of
450 °C. Fig. 3 shows the stress versus strain relations of AA2024 at different temperatures.
Significant softening due to increasing temperature can be observed, and the flow stress decreases
with rising temperature from approximately 200 MPa at 350 °C to about 50 MPa at 493 °C.
3.2. Dependence of formability on temperature
Temperature played a dominant role in the formability tests of AA2024, as it did in the ductility
tests. Fig. 4 shows the effect of temperature on formability. Formability at 450 °C was high (see Fig.
4a), as expected from the high ductility observed at that temperature in tensile testing. Significant
thinning was observed, especially in the central hole area (Table 4), indicating that severe plastic
deformation occurred during the tests. As the forming temperature approached the SHT temperature
(493 °C), the AA2024 samples showed extremely poor formability and exhibited cracking with low
levels of ductility (Fig. 4b). Thickness measurements showed little or no localised plastic
deformation or thinning (average thickness is 1.95 mm), suggesting that cracking occurred early in
the forming operation. The lack of plasticity associated with cracking is consistent with the low
ductility noted for tensile tests at these temperatures.
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Table 4. Wall thickness distributions for the samples shown in Figs. 5aI and 5bI.
Distance from the centre
10
12
14
18
20
24
30
34
35
38
40
42
(mm)
Wall thickness (mm) of
1.67
1.84
1.83
1.88
1.8
1.77
1.28
1.24
1.2
1.44
1.5
1.71
16
18.5
20.6
26.6
32.5
36
40
42
1.32
1.37
1.44
1.53
1.7
1.83
1.86
1.9
sample Fig. 5aI
Distance from the centre
(mm)
Wall thickness (mm) of
sample Fig. 5bI or 4a
3.3. Dependence of formability on forming velocity
Fig. 5 shows the effect of deformation speed on formability. For the hemispherical head punch, the
material flow was influenced significantly by forming velocity. For low punch speeds, localised
plastic deformation occurred circumferentially, lower on the test piece relative to the hole, with a
mild necking zone visible on the side-wall (see Fig. 5aI). On the other hand, at high speeds, a
considerable increase in the diameter of the central hole was observed (Fig. 5bI), i.e. plastic
deformation predominated in the vicinity of the central hole, rather than around the circumference,
which indicates a change in the location of localised plastic deformation during the high speed HFQ
tests.
Figs. 5aII and 5bII show the deformation characteristics using a flat head punch instead of a
hemispherical punch. In both the low and high speed tests, an enlarged central hole can be observed,
suggesting that plastic deformation took place primarily in the vicinity of the central hole in both
tests. Again, the high-speed test exhibited a larger final hole diameter (D = 28.3 mm; Fig. 5bII) than
the low speed test (D = 27.2 mm; Fig. 5aII), indicating that more plastic deformation took place
around the central hole area under high speed forming. The implication is that an optimum speed
exists between the low and high speed limits such that plastic deformation is most evenly
distributed, and localises minimally as a neck or a concentration around open features, both of
which could ultimately lead to ductile failure during forming.
3.4. Damage mechanisms
Fig. 6aI shows the fracture surface of the tensile sample deformed at 350 °C. It has a dimpled cupand-cone appearance, indicating that microvoid nucleation, growth and coalescence was the
dominant fracture mechanism, which is similar to the fracture mechanism of AA2024 at room
temperature [17, 19, 20]. Fig. 6aII indicates the region adjacent to the fracture area on a longitudinal
7
section of the deformed sample showing the path of the fracture and intermetallic particles (white).
This section shows that secondary crack branching from the primary fracture surface to nearby large
inclusions occurred, but inclusions further away from the primary fracture surface did not show
evidence of decohering or voiding. This suggests that the primary crack nucleated perpendicular to
the loading direction and propagated across the section without initiation of further cracks from
inclusions or other sites occurring readily.
Fig. 6bI shows the fracture surface of the sample deformed at 450 °C. The features of this fracture
surface are much different than those of the sample deformed at 350 C, which appears as fine
ductile features superimposed on a background of elongated surface features and resembles an
intergranular type of failure. Heating the sample to 450 C will have resulted in an initial tendency
for the original T3 condition to age with ”, ’ and S’ forming prior to the equilibrium  and S
(these stages will also have taken place on heating to 350 C). Rapid heating will restrict the
formation of strengthening precipitates until temperatures when dissolution again becomes
favourable. This results in most of the Cu and Mg being in solid solution during deformation at 450
C. During elevated temperature deformation, thermally activated diffusion of alloying elements
will be taking place, enhanced by the fast diffusion paths offered by the excess dislocations present
from the initial T3 treatment and those introduced during deformation. Segregation of alloying
element atoms to defect sites such as dislocations and, more importantly, grain boundaries would be
expected through this treatment although high-resolution transmission electron microscopy (TEM)
would be needed to confirm this behaviour. Any precipitates formed would be too fine for SEM
resolution and so the bright particles observed in Figs. 6aII and 6bII are from Mn-, Fe- and Si-rich
intermetallic phases formed largely in the melt. These will be present with a range of compositions,
sizes and, hence, thermal stabilities, but would generally be hard and undeformable within the
aluminium matrix. As seen from a comparison of Figs 6aII and 6bII, there is an apparent decrease
in the volume fraction of inclusions from 6aII to 6bII, but, given the inhomogeneous distribution of
inclusion particles, this falls within the range of anticipated scatter in inclusion volume fraction.
Likewise, there is no major difference in the range of sizes and aspect ratios between 350 and 450
C, indicating neither dissolution nor plastic deformation of these particles took place. The
continued non-deformable nature of the inclusion particles would explain the greater tendency to
form voids (associated with large particles or clumps of particles) further from the main crack (Fig.
6bII). The greater thermal softening of the matrix at 450 C compared with 350 C would result in
greater strain concentration in the matrix around the inclusion particles and so localised necking
would occur, giving rise to ductile voids. With a greater difference in strength between the particle
and the matrix, smaller particles would give rise to voiding resulting in the observed larger number
8
of smaller ductile dimples seen in Fig. 6bI compared with Fig. 6aI. The presence of the inclusions
explains the fine-scale ductile voiding, but does not explain the background surface features in Fig.
6bI. However, a secondary crack (marked ‘X’) is seen in Fig. 6bII, which is not associated with
voiding around particles and follows a curved boundary-like path. This suggests a role for grain
boundary failure processes at elevated temperatures.
Fig. 6cI shows the fracture surface at 493 °C, which is characterised by the absence of ductile
dimples and an entirely intergranular fracture path. Although increasing temperature would
facilitate voiding, changes at the matrix grain boundaries result in a much lower resistance to
cracking at these sites. Hence, at this temperature extensive secondary cracking occurs below the
main crack (see Fig. 6cII). Fig. 6cII appears to show the absence of inclusions, but this is just a
spatial variation; although some of the less stable Fe- and Si-bearing inclusions would be expected
to dissolve (limited due to the low time of exposure to elevated temperatures), the majority of Mnrich inclusions do not start to dissolve below 500 C. The behaviour exhibited in the fractographic
study necessitated a more detailed, in-situ SEM examination of the behaviour of inclusions and
grain boundaries.
3.5. In-situ SEM studies
Fig. 7 shows in-situ SEM observations of the evolution of inclusions and coarse precipitates after
heating to different temperatures using the Gatan H1002 heating stage. The irregularly shaped white
particles in Fig. 7 are inclusions, which are generated during material manufacture. The particles are
large (up to 15 m in size) and distributed inhomogeneously. The region shown in Fig. 7 also
exhibits several clusters of inclusions. The inclusions in AA2024 were identified by EDX as
compounds of iron, magnesium, manganese, silicon, and copper, but the composition varied from
one particle to another. It should be noted that surface preparation resulted in some residual
scratches, running from top left to bottom right in the figure; these, and their associated strain fields,
may have contributed a small amount stored energy that could have accelerated the annealing
process somewhat during in-situ heating. No visible changes were observed in the microstructure
until the specimen reached a temperature of 320 °C, when grain boundaries started to become
visible, as indicated by the arrows in Fig. 7b, which was probably due to the precipitation of second
phase particles on the grain boundaries. The relatively slow heating rate used (6 C/min) means that
more substitutional alloying element diffusion is possible than during the HFQ process and the
forming trials, so that greater development of precipitate structures should be expected. As the
specimen temperature rose to 400 °C, more continuous grain boundaries can be clearly observed
(Fig. 7c), indicating that more second phase particles precipitated on the grain boundaries. When
9
the specimen temperature was close to 450 °C, as shown in Fig. 7d, dispersed precipitates could be
observed not only on the grain boundaries, but also within the grains. The size and location of the
precipitates would suggest that they are the equilibrium  and S phases. The second phase particles
are larger in size compared with those present at lower temperatures. When the temperature was
further increased to 480 °C, most of the precipitates had re-dissolved (see Fig. 7e). Fig. 7e shows
the removal of precipitate phases from the grain boundaries by dissolution, which will leave the
alloying element atoms (Cu and Mg) in solution. Loss of these atoms from the grain boundary
regions is diffusion-limited and so it is likely that the grain boundary regions remain solute-enriched
at this temperature.
In Fig. 7e, a grain boundary crack is opening up in the region within a cluster of inclusions. The
width of the grain boundary defect is far greater than would be seen through thermal grooving. This
suggests that the crack was softened whilst under surface traction forces. As the sample was just
heated in the SEM and not strained, these forces must arise from residual stresses produced by the
T3 state (quenched and plastically deformed). These residual stresses would be concentrated in
regions between hard, non-deformable particles consistent with the features shown in Figs. 6 and 7.
At this magnification it is not possible to determine whether the residual stresses caused separation
of a liquated phase or ductile failure of grain boundaries weakened by segregation bringing the local
solidus and liquidus temperatures closer to the sample temperature. The slow heating rates would
suggest that the latter is more likely, but segregation and inclusion clustering is necessary for grain
boundary separation. In either case the observations suggest considerable weakening of the grain
boundaries above 450 C, which upon application of mechanical loading will greatly limit the
ductility and formability. Raising the sample temperature to 493 C (see Fig. 7f) causes some
rounding of the inclusions, but no dissolution. In addition, the grain boundaries were clearly
delineated throughout the field of view, but not opened up as seen in Fig. 7e (the particular
boundary clearly visible in Fig. 7e appeared to have suffered some contamination upon heating
beyond 480 C). This behaviour is consistent with greater vaporisation losses from the edges of the
inclusions and the grain boundaries (thermal etching or grooving).
4. Discussion
4.1. Effect of punch velocity on the material flow during HFQ tests
The combined effects of material rate dependence and a forming-speed-dependent, radially varying
distribution of temperature in the workpiece during forming led to the difference in plastic
localisation observed in the low- and high-speed tests (Section 3.3). During the hemispherical
punch tests, the material making good contact with the cold punch, the central hole area, was cooled
10
rapidly by conduction with an additional, smaller contribution to cooling by radiation and
convection of heat into the surrounding, room temperature air; on the other hand, the material with
little or no contact with the cold punch, the circumferential area, was cooled mainly by the air and
hence at a much lower rate. The radial variation in the temperature of the workpiece due to
conduction of heat into the cold die is accentuated at a low punch speed, which provides greater
time for heat transfer to take place. Decreasing temperature below 450 C increases the flow stress
and decreases the ductility of AA2024, and decreasing the strain rate (lower punch speed) decreases
the flow stress; in the vicinity of the hole, the two effects oppose each other for the low speed punch
tests, but the effect of cooling dominates leading to an overall higher flow stress and lower ductility
in the hole area relative to the circumferential area. In the circumferencial area, accentuation of the
punch cooling effect by lower punch speed is mitigated or absent altogether, hence the lower strain
rate acts alone to decrease the flow stress in that area, causing localisation of plastic deformation
and neck formation. On the other hand, at a high punch speed, the duration of contact between the
punch and workpiece is shorter, hence less heat transfer takes place in the vicinity of the hole
keeping the material relatively warm and easy to deform. However, the higher strain rate (high
punch speed) acts to increase the flow stress, opposing the effect of reduced cooling; again, the
effect of cooling, or relative lack thereof in the high punch speed case, dominates the effect of
changing the strain rate, and the material in the vicinity of the hole in the high punch speed case
deforms more easily relative to the circumferential area (the opposite of the low punch speed case),
causing thinning around the hole and a propensity for necks radiating from the hole that lead to
fracture in the case of excessive punch displacement. It is important to reiterate that the difference
observed between low and high punch speed suggests that an optimum punch velocity exists at
which near uniformity of strain can be achieved, which will lead to higher formability by delaying
neck formation.
In the flat head punch tests, good contact between the punch and work piece occurred only around
the punch face circumference, which mitigated the effect of punch cooling in these tests. The
variation observed in the flat punch tests with punch speed is attributable primarily to the
corresponding difference in strain rate experienced by the material, and to a lesser extent, the
cooling effect of the cold punch. At high punch speed, the material contacting the circumference of
the punch face was warmer than in the low speed punch tests, but overall relatively strong due to
the higher applied strain rate. The effect of the higher strain rate dominated the relative lack of
cooling in this case, and as a consequence, the workpiece material deformed at a higher flow stress
at the punch circumference in the high punch speed test and overcame the effect of friction at the
11
punch circumference to a greater extent, thereby decreasing circumferential localisation and
spreading deformation to the top face and the central hole.
4.2. Failure mechanisms of AA2024 under hot forming conditions
It is well known that fracture of AA2024 is strongly affected by second phase particles, as they
provide fracture nucleation sites, with intermetallic inclusions being dominant [17, 18]. The
strength of the alloy is controlled by the precipitate distribution developed during the ageing
treatment. The precipitation sequence of an Al-Cu-Mg alloy can be summarised as [23]:
SSSS  Cu, Mg co-clusters  S’  S-phase precipitates (Al2CuMg)
Initial strengthening by second phase particles
For AA2024 in the T3 state, initial strengthening is caused by spherical or linear Cu-Mg co-clusters
[24]. These zones form during room temperature ageing, and although their coherent nature makes
their nucleation rate the fastest, their thermal stability is low so that dissolution or transformation to
S’ and S occurs (or ”, ’ and  if segregation changes the Cu/Mg ratio). In the segregated regions,
when the Mg content is very low, the system behaves more like an Al-Cu alloy, thus the precipitate
sequence SSSS  Cu, Mg (mainly Cu) co-clusters  ”  ’   (Al2Cu) occurs [25-27]. This
initially occurs in the temperature range from about 150 °C to 200 °C [1, 3, 23, 28]. At this stage,
the maximum density of precipitates can be obtained, which are normally very small in size (about
10 - 15 nm), and are needle or plate shaped with a large aspect ratio ranging from 10 to 40 [29, 30].
In general, these precipitates are considered to be unshearable at room temperature, regardless of
their size, and the presence of such highly dispersed precipitates leads to a considerable
enhancement of strength. The strengthening effect is determined by the average spacing between
the randomly distributed particles, as the stress required to bow a dislocation between particles
increases with decreasing spacing. Therefore, the precipitation hardening effect can be enhanced by
increasing the volume fraction of precipitates formed, and reducing the average size of the
precipitates [30]. Precipitation hardening is usually natural (room temperature) or artificial (120 –
170 C), but the development of precipitates is controlled by substitutional alloying element
diffusion with peak strength being achieved after 3 to 8 hrs. Continued holding at, or heating above,
the ageing temperature leads to overageing and coarsening [23, 28, 31, 32]. Transformation to the
equilibrium  and S phases occurs coupled with Ostwald ripening (growth of larger particles at the
expense of finer ones), reducing strength. The rapid heating rates of the HFQ and in-situ SEM tests
means that the precipitation and coarsening reactions will not go to completion, although the extent
of these processes will be much greater for the in-situ SEM than for the HFQ process at the same
12
temperature. Thus, SEM examination reveals extensive grain boundary precipitation of the
equilibrium phases on heating to 450 °C (see Fig. 7d), followed by dissolution and grain boundary
separation within clusters of inclusions (see Fig. 7e). During HFQ tests, less grain boundary
precipitation would occur, but also less time is available for de-segregation so that solute
enrichment, softening and reduced solidus and liquidus at the grain boundaries would be expected
to be more extensive at grain boundaries for the HFQ condition than in the in-situ SEM trials.
Fracture mechanisms during hot forming
As demonstrated earlier, the nature of the grain boundary structures and temperature affected the
failure mode observed during the HFQ tests. Several different fracture modes were observed.
Fracture mode 1: Ductile fracture (see Fig. 8a). At low temperatures (350 °C ≤ T < 450 °C), voids
initiated by debonding or cracking of inclusions in the grains and ductile fracture was the dominant
mechanism, and is probably the same as that for cold forming [17, 19]. Fracture mode 2: Mixed
ductile and intergranular fracture (see Fig. 8b). Under medium temperature conditions (450 °C ≤ T
< 480 °C), a large amount of the precipitate phases have been dissolved, but a range of inclusion
sizes remains. The lower matrix strength gives rise to earlier void formation resulting in a finer
ductile dimple population. Precipitate dissolution results in solute enrichment of grain boundaries
and their weakening compared to the grain bulk regions at the forming temperature due to the lower
local solidus and liquidus temperatures at the grain boundaries (grain boundaries are being
deformed at a higher homologous temperature than the grain bulk). During HFQ deformation, the
grain boundaries separate so that the finer ductile dimples are seen on a background of intergranular
failure features. Theoretically, the precipitates are considered to be unshearable by the dislocations
at room temperature; but the strength of precipitates decreases with increasing temperature, and at a
higher softening rate than that of the matrix material. At 450 °C, the maximum ductility was
observed, which is most likely because both precipitates and matrix were very ductile at 450 °C and
their strengths were also about the same. Fracture mode 3: Intergranular fracture (see Fig. 8c). At
high temperature, 480 °C ≤ T ≤ 493 °C, which is achieved by a high heating rate, the homologous
temperature of the boundaries is still higher than that of the bulk as de-segregation from the soluteenriched boundaries will not have occurred. At these temperatures, the cohesive strength of the
boundary is so low that voiding around inclusions does not occur and thus intergranular fracture
was the dominant mechanism. The high heating rates may have caused some liquation of the soluteenriched boundaries, which would require greater SEM and TEM studies to confirm.
13
5. Conclusions

The results of tensile tests showed that the maximum ductility of AA2024 under hot forming
conditions was achieved at 450 °C. Moreover, a steep decrease in ductility was observed as the
temperature ranged from 450 to 493 °C.

Temperature has a strong influence on the formability tests. Cup testing confirmed that the
formability of AA2024 could be improved considerably by forming in the vicinity, but not in
excess of 450 °C. (It should be noted that ductility decreases gradually as the temperature drops
below 450 °C, which provides an attainable temperature window for optimum forming.)

Punch velocity has a strong influence on formability. Punch velocity strongly affects the
strength of workpiece material due to its effects on temperature and strain rate distribution, and
hence flow stress and ductility. The forming trials suggest an optimum punch velocity exists
that would distribute plastic deformation most uniformly over the workpiece, hence delaying
necking and allowing for greater formability.

Coarsening/growth of second phase particles was clearly observed by in-situ SEM examination
at temperatures up to 450 C. It was found that the solute atoms preferably precipitated onto the
grain boundaries, and the presence of clustered inclusions combined with boundaries weakened
by solute enrichment caused grain boundary separation that may contribute to the low ductility
of AA2024 under HFQ hot forming conditions.

Under HFQ hot forming conditions, a transition of fracture mechanisms was revealed. At low
temperature conditions (350 °C ≤ T < 450 °C), ductile fracture due to debonding or fracture of
second phase particles was the dominant failure mode. At medium temperatures (450 °C ≤ T <
480 °C), solute enrichment of grain boundaries reducing grain boundary cohesion coupled with
softening of the matrix to make voiding around finer inclusions easier led to a mixed ductile and
intergranular fracture. At the highest forming temperatures (480 °C ≤ T ≤ 493 °C), intergranular
fracture due to very low strength grain boundaries (that may include liquation) became the
dominant fracture mode.
14
Acknowledgement
The support from EPSRC, grant No: EP/E00573X/2, is greatly appreciated. The authors also wish
to thank the support from Professor Peter Lee, Department of Materials, Imperial College London,
on microstructural analysis and Jingqi Cai, Department of Mechanical Engineering, Imperial
College for help with testing.
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16
High speed
camera
Prism
Fig. 1. Formability test rig equipped on a 25 tonne ESH high-speed press.
17
1.2
True strain to failure
1.0
0.8
0.6
0.4
0.2
0.0
360
380
400
420
440
460
480
500
o
Temperature ( C)
Fig. 2. Ductility test results for AA2024 at different temperatures.
18
350oC
Stress (MPa)
200
150
400oC
100
450oC
50
493oC
0
0.0
0.2
0.4
0.6
0.8
1.0
1.2
True strain
Fig. 3. Stress versus strain relations of AA2024 at different temperatures.
19
50mm
50mm
(a)
(b)
Fig. 4. HFQ formability test results at (a) 450 ± 10 °C and (b) 493 ± 10 °C.
20
(a)
(b)
(I)
50mm
50mm
50mm
50mm
(II)
Fig. 5. HFQ formability test results using (I) hemispherical head punch and (II) flat head punch at different
punch velocities of (a) 170 mm/s and (b) 486 mm/s.
21
(I)
(II)
(a) 350 °C
X
(b) 450 °C
(c) 493 °C
Fig. 6. SEM images for (I) fracture surfaces and (II) longitudinal (arrow direction) sections close to
failure areas of the samples deformed at (a) 350 °C, (b) 450 °C and (c) 493 °C
22
(a) 50 °C, initial structure
(b) 320 °C
(c) 400 °C
(d) 450 °C
(e) 480 °C
(f) 493 °C
Fig. 7. In-situ examination of the inclusions and precipitates at different temperatures. The
heating rate is 6 °C/min
23
(a)
(b)
(c)
Fig. 8. Fracture modes under hot forming conditions: (a) Ductile fracture, (b) mixed ductile
and intergranular fracture and (c) intergranular fracture.
24