Formability and failure mechanisms of AA2024 under hot forming conditions

L. Wang1, M. Strangwood2 D. Balint1, J. Lin*1 and T. A. Dean3

1Department of Mechanical Engineering, Imperial College LondonSW7 2AZ, UK

2School of Metallurgy and Materials, University of Birmingham, Edgbaston, Birmingham B15 2TT, UK

3School 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).

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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, aluminium-

based alloy is used, the part can then be aged to obtain full strength. The feasibility of this novel

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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).

Table 1. Chemical composition of AA2024 (wt %) [12].1 Provided by Airbus S.A.S; contains small amounts of proprietary constituents.

3

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 Rate1 s-1

35

0

36

0

37

0

38

0

39

0

40

0

41

0

42

043

0

44

0

45

0

46

0

47

0

48

0

48

7

49

3

Formability testing4

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

3. Results

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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

(mm)10 12 14 18 20 24 30 34 35 38 40 42

Wall thickness (mm) of

sample Fig. 5aI1.67 1.84 1.83 1.88 1.8 1.77 1.28 1.24 1.2 1.44 1.5 1.71

Distance from the centre

(mm)16 18.5 20.6 26.6 32.5 36 40 42

Wall thickness (mm) of

sample Fig. 5bI or 4a1.32 1.37 1.44 1.53 1.7 1.83 1.86 1.9

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 cup-

and-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

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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

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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 Mn-

rich 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

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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 solute-

enriched 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 solute-

enriched 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

Fig. 1. Formability test rig equipped on a 25 tonne ESH high-speed press.

17

High speed camera

Prism

Fig. 2. Ductility test results for AA2024 at different temperatures.

18

Fig. 3. Stress versus strain relations of AA2024 at different temperatures.

19

(a) (b)

Fig. 4. HFQ formability test results at (a) 450 ± 10 °C and (b) 493 ± 10 °C.

20

50mm 50mm

(a) (b)

(I)

(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

50mm 50mm

50mm 50mm

(I) (II)

(a) 350 °C

(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

X

(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