Heat treatment, thickness and chemical composition of tested aluminium alloys
For a long time, aluminium wrought alloys have been used as the predominant material for the sheet metal processing industry. Aluminium alloys not only show a high affinity for passivation which results in corrosion resistance in many environments , but also indicate high specific strengths (strength versus density) and rigidity (Young´s modulus versus density). The most commonly used aluminium alloys in aircraft applications of the non-precipitation hardenable groups are EN AW-3xxx, EN AW-5xxx and of the precipitation hardenable groups EN AW-2xxx, EN AW-6xxx and EN AW-7xxx. The advantages of the EN AW-2xxx group are good mechanical properties which remain at higher temperatures, whereas the EN AW-7xxx groups show good mechanical properties at lower temperatures and high stress corrosion resistivity. A negative aspect is, however, that aluminium alloys have a very low formability when compared with modern steels. Figure 1 illustrates the fracture elongation versus yield strength of aluminium wrought alloys and deep drawing steels in optimised hardened condition. Aluminium wrought alloys with the designation EN AW-7xxx series in peak-aged condition show yield strength values comparable with the micro-alloyed steels; although the fracture elongation differs slightly in favour of the steels.
Yield strength values from the wrought alloys of EN AW-6xxx group reach only about 50% of the strength which can be obtained from the EN AW-7xxx group at ambient temperatures; both aluminium wrought alloys are precipitation hardenable in peak-aged condition. The fracture elongation of the 6xxx and 7xxx group can be identified to be approximately half of the value which can be obtained using bake-hardening steels. Yield strength values from non-heat-treatable, and therefore only strain-hardenable, aluminium alloys of the EN AW-5xxx group are slightly higher when present in highly strain-hardened condition (e.g. H18). But their fracture elongation values in an annealed condition are equivalent to those exhibited by the EN AW-6xxx group. So, ways of enhancing the fracture elongation of wrought aluminium alloys to levels that are at least comparable with those of deep-drawing steels and additionally possess corresponding strength values are sought. Forming aluminium alloys below room temperature would be an approach, where strain hardening losses and precipitation processes are not an issue. In this regard there is no danger of a reduction in strength through the mechanism of recovery, recrystallization and aging which is in most cases induced by forming at higher temperatures. Furthermore, in previous investigations during tensile testing it was observed that in certain aluminium alloys, which were present in various heat treatment conditions, the Lüders and the Portevin-LeChatelier (PLC) effect occurs at room temperature  . This PLC effect not only causes a bad surface quality on sheet metals, but also creates a reduction of the uniform elongation and fracture elongation . It is expected that in tensile testing at low temperatures the negative impact caused by the PLC effect can be reduced and therefore the values of uniform elongation and fracture elongation can be increased.
2. Portevin-LeChatelier (PLC) Effect
2.1. Phenomena of the PLC effect
Considering a tensile testing specimen, consisting of an isotropic ductile material, it can be observed that straining in the direction of the principal normal stress is accompanied by a cross-sectional reduction in width and thickness. In the region of uniform elongation the prismatic geometry of the specimen remains, since a
In this deformation state it can be observed that only localised necking occurs () and there is no reduction in width over the whole specimen length () which would be denoted as diffuse necking . This is
let σ1≠ 0 and σ2=σ3=0 then
If only dεl and dεw are considered (because of a small sheet thickness), the expression can be rewritten as
Since dεw< 0 a Mohr´s circle of strain can be developed as illustrated in Figure 3. Applying a coordinate transformation it can be shown that α=70°32´if dεw´=0.
Thus, the angle between the direction of the localised neck and the direction of the principal normal stress is ideally 54°44´ which is in accordance with the hypothesis of Hill . Eventual variations can be attributed to textures of the material.
2.2. Microscopic background of the PLC-effect
The microscopic background of the PLC effect is still a subject of discussion. This is because the PLC effect appears at temperatures where interstitial or substitutional diffusion of solute alloying elements do not offer sufficient mobility by bulk diffusion. A brief introduction of the development of ideas concerning the PLC effect is illustrated in  . The first model developed by Cottrell and Bilby  assumed that the mobility of solute atoms is actually high enough to follow mobile dislocations. Subsequently, McCormick recognised that the mobility of solute atoms is barely sufficient to follow mobile dislocations, and he postulated that clustering mobile dislocations only occurs if they are arrested by obstacles such as by forest dislocations, for example . Van den Beukel presupposed in his model, which was based on McCormick’s work, that arresting mobile dislocations at forest dislocations is only caused by volume diffusion of solute atoms . Following a proposal of Sleeswyk , the model created by Mulford  not only postulates the arresting of the stopped dislocation by volume diffusion, but additionally by a simultaneously running pipe diffusion (see Figure 4).
A requirement for such a condition is a prior formation of a Cottrell-atmosphere at stationary forest dislocations. The pipe diffusion occurs without the acceleration caused by the vacancy mechanism. Modelling the pipe diffusion of the substitutionally dissolved magnesium in the aluminium bulk showed that the effectiveness of this diffusion process is not high enough to explain the experimental results . As a consequence of this model the diffusion velocity of magnesium atoms is too slow to form Cottrell-atmospheres around the core of the arrested dislocation during representative times. Considering these results Picu  developed a new model which is based on the estimation that a forest dislocation embedded in a Cottrell-atmosphere reacts with a mobile dislocation by forming
3. Experimental methods
A set of experimental tests were conducted based around a tensile test machine. Four types of wrought aluminium alloys were examined which are detailed in Table 1. In this table the sheet thickness and detailed composition are shown.
strain-hardened to a half hard condition - no thermal treatment
slightly strain-hardened (less than H11) – no thermal treatment
solution heat-treated at 465°C for 20 minutes and naturally aged to a substantially stable condition
solution heat-treated at 465°C for 20 minutes and tested within 30 minutes
For the purpose of evaluating potential strain rate dependencies the uniaxial tensile tests for the aluminium alloys EN AW-6016 and EN AW-7021 in the solid solution heat treated and naturally aged condition T4, and the EN AW-5182 aluminium alloy in the strain hardened condition H111 (as currently implemented in industrial applications), were carried out with the following three forming velocities; (maximum velocity of the tensile testing machine), and resulting in strain rates s-1, s-1 and s-1 respectively (see Table 2). By way of comparison, the highest strain rate illustrated, , approaches the forming velocity used in a typical industrial process. Determining the effect of precipitations for the 6000 and 7000 series a number of subsequent tests with the highest strain rate were carried out within 30 minutes after a solid solution heat treatment (W30 condition). As a reference the aluminium alloy EN AW-1050A was tested in half-hard condition (H14) with . To be thorough regarding a potential influence of the strain rates, the reference aluminium alloy was also tested with the lowest strain rate ()
|6.60 . 10-2 s-1||1.70 . 10-2 s-1||0.17 . 10-2 s-1|
Flat sheet specimens were CNC machined showing a test section width of 12.5mm in accordance with DIN EN ISO 6892-1 (A50) . Attention was given towards the rolling direction of the sheet metal which was perpendicular to the applied load. The length of the 50mm wide clamp sections of the specimens were designed to fit with the requirements of the cryogenic set-up.
A specially designed cryogenic dewar (Figure 5) was used to obtain and maintain the sub-zero temperatures required in these tests. A variation of testing temperatures was reached by way of a
Figure 6 a) through f) shows for the aluminium alloys tested the temperature dependency of the yield strength
It can be seen that with one exception and for all materials the yield strength, ultimate tensile strength and elongation rises consistently when temperature is decreased. The strength values of the aluminium alloys in W30 condition are lower than in T4 condition, which is as expected. For the aluminium alloys EN AW-5182 and EN AW-7021 in W30 condition a uniform elongation of over 40% and a fracture elongation of considerably more than 50% can be reached at 77K. A similar trend can be observed for the aluminium alloy EN AW-6016 in both heat treatment conditions, where the elongation in T4 condition is less than in W30 condition. For a decrease in temperature the aluminium alloy EN AW-1050A shows a rather shallow rise of the ultimate and yield strengths over the whole temperature range; whereas curves of uniform and fracture percentage elongation of this alloy indicate significant increases in the magnitudes of their gradients. The strength and elongation of the aluminium alloys EN AW-7021 also show a growing trend if temperature is reduced. For both heat treatment conditions, T4 and W30, a gradual elongation rise of 25% at 152 K can be seen. If the temperature is lower than 152 K the elongation of the peak aged aluminium (T4) alloy decreases again. In turn, when present in W30 condition, the uniform elongation rises to more than 40% and the fracture elongation approaches 50% at 77K.
The influence of strain rates regarding strength values can be considered as small (see Figure 7 a), c), e) & g)). A slightly more distinct effect can be observed for the elongation values, especially for EN AW-5182 and EN AW-6016 T4. Nevertheless, those variations converge as the temperature approaches 77K (see Figure 7 b), d), f)&g).
Some overall observations where made as a result of the tests from stress-strain data. Firstly aluminium alloy EN AW-1050A-H14 showed no pronounced yielding for all strain rates and temperatures. The same effect was observed for EN AW-6016 in both heat treatment conditions, T4 and W30. In contrast the aluminium alloy EN AW-5182-H111 showed pronounced yielding at all strain rates and temperatures. Aluminium alloy EN AW-7021 in W30 condition indicated pronounced yielding only at 298K and 248K. When the temperature was reduced this effect was not evident. On the other hand, EN AW-7021 in T4 condition showed similar yielding effects for all strain rates as was also observed for EN AW-6016. If one now considers Figure 8, the results of increasing the uniaxial load on an EN AW-5182-H111 specimen can be seen over the elastic and plastic regions. It is clear that there is a pronounced yielding behaviour accompanied by a region of stress oscillations. For the results, shown at T=298K, the amplitude is as high as ± 2 MPa; whereas for T ≤ 248K the magnitude reduces to ± 0.2 MPa.
The PLC-effect, caused by Cottrell-atmospheres, offers a plausible mechanism to explain the resulting reduction of fracture elongation
Ludwik, Hollomon and Backofen identified that the flow stress
A relationship which was empirically developed by Ludwik modified by Hollomon and extended by Backofen is
Using this relationship most stress-strain curves can be described accurately in the region of uniform elongation a
The strain hardening coefficient
Figure 9 a) through f) shows only a slight strain rate dependency. Apart from the temperature range between 298K and 248K and from EN AW-7021 in the precipitation hardened condition (T4), an increase of the strain hardening coefficient and uniform elongation as temperature is reduced can be seen. Except for the aluminium alloy 1050A, (Figure 9 a)) which is consistent with the Ludwik curve, tested aluminium alloys indicate a parallel offset. It is assumed that this shift is due to the PLC-effect. Therefore Eq. 6, which takes only the strain hardening process into account, seems not to be applicable for all aluminium alloys. A modification to this equation would seem appropriate to allow for the description of materials which exhibit a PLC-effect. Eq. (8) illustrates a
The modified Ludwik curve is illustrated on Figure 9 b) through f) which incorporates a value of
The plastic behaviour of commercial aluminium wrought alloys in annealed or solution annealed condition can be positively enhanced if the forming temperature is below 123K. This advantage can be reached without any interference of higher strain rates.
Figure 10 demonstrates the fracture elongation values which are currently attainable from aluminium alloys compared with values achievable from deep-drawing steels. A competitive advantage of cryogenic forming of annealed or solution annealed aluminium alloys is clearly shown. Flow curves showing PLC-effects can be described more precisely with addition of a constant p which represents an extension of the initial Ludwik relationship.
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