Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
1
AGE-HARDENABLE ALUMINIUM ALLOYS
This invention concerns AA5000 series alloys with
the addition of Cu that can be retained in a solution
treated condition after hot working, for example by hot
rolling on a hot mill or by hot extruding.
In the art AA5000 series alloys are usually regarded
as non-heat treatable alloys i.e. they are not regarded
l0 as age hardenable. The addition of Cu to these alloys
renders them age hardenable, as described in
EP-A-0773303, EP-0616044 and EP-A-0645655. However these
known methods also require a formal solution treatment.
The novel feature of this invention is the discovery
that for certain Cu - containing AA5000 series alloys
sufficient solution treatment occurs during hot working,
for example hot rolling, to render the alloys age
hardenable without a further expensive solution treating
step. This gives a very significant economic advantage
especially for commodity products such as can end stock,
automotive sheet products, or extruded products such as
structural sections.
EP-A-0605947 describes manufacturing can body sheet
using two sequences of continuous operations. The
described additional steps of uncoiling the hot coiled
sheet, quenching the sheet without intermediate cooling,
cold rolling and re-coiling the sheet are required, but
these additional steps are not needed in the method of
the present invention.
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
2
WO-A-99/39019 describes a method for making can end
and tab stock but annealing of the sheet is required as a
separate operation after hot rolling which is not needed
in the method of the present invention.
WO-A-98/01593 describes a process for producing
aluminium alloy can body stock but again a separate
annealing step is required.
JP-A-100121179 describes aluminium alloy sheet for
carbonated beverage can lids but a formal solution heat
treatment is required, which is not needed in the method
of the present invention.
US-A-5655593 describes aluminium alloy sheet
manufacture in which the hot strip is cooled rapidly to
minimise the precipitation of the alloying elements.
This teaching of rapid cooling is contrary to that of the
present invention.
US-A-3464866 describes a process for obtaining
aluminium alloy conductors but again teaches rapid
cooling.
In accordance with the present invention there is
provided a method of producing an age-hardenable
aluminium alloy comprising the steps of:
(a) casting an alloy of a composition comprising the
following expressed in weight percent:
Magnesium: 1.0 to 4.0
Copper: 0.1 to 0.6
Manganese: up to 0.8
Iron: up to 0.5
CA 02431029 2003-06-11 GB0105686
.15-03-2003 __-__ _ _ ___
3
Silicon: up to 0.3
Chromium: up to 0.15%
Titanium: up to 0.15%, preferably up to
_. 0.05%
Boron: from 0 up to 0.05, preferably up
to 0.01
Balance: Aluminium with incidental
impurities
(b) optionally homogenising the cast alloy,
(c) hot working the casting at an initial temperature of
at least 400°C, to form an intermediate product, wherein
at Least part of the hot working is carried out whilst
the casting is at a temperature above the solves
temperature of the alloy,
(d) cooling the intermediate product during hot working
or in a subsequent step at a rate of less than 5°C/min
such that at least a partially recovered or
recrystallised structure is formed and that sufficient
copper is retained in solid solution in the alloy to
cause an age hardening effect on the alloy if phase
precipitation takes place during the alloy's subsequent
thermal history, and
(e) optionally allowing or arranging for phase
precipitation to occur in the alloy.
Preferably after the said hot working step the
intermediate product is generally maintained at a
temperature below the solves temperature of the alloy,
provided that if the intermediate product is heated above
3o the alloy's solves temperature then cooling thereof is
effected at a rate less than 2°C/sec.
AMENDED SHEET
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
4
By the term "the solvus temperature of the alloy" is
meant the temperature below which under equilibrium
conditions the copper begins' to be removed from solid
solution to form a precipitate. However, as to the rate
of copper removed that will depend on the kinetics of the
reaction.
The precipitation phase if formed is believed to be
S phase (an AlCuMg phase) or its metastable precursors.
l0
The alloy may be cast by DC casting to form an ingot
or by continuous casting, for example in a belt caster or
a twin roll casting machine, to form a sheet.
The cast and preferably homogenised alloy can be
extruded but for the production of can end stock it is
generally hot rolled. After casting the preferred steps
are:
optionally homogenising the casting at a temperature
of at least 480°C, and preferably 500 to 600°C, so that
substantially all of the magnesium and copper in the
casting are in solid solution,
optionally hot rolling the casting, optionally with
re-heating of the casting to above the alloy's solvus
temperature, preferably at least 450°C, to take
substantially all of the magnesium and copper present
into solid solution,
hot rolling the casting with a rolling mill entry
temperature of the casting of at least 400, and
preferably from 450° to 580°C,
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
continuing rolling the casting to the desired
thickness to form a sheet so that at least part of the
rolling reduction is carried out above the solvus
temperature of the alloy and cooling the alloy, either
5 while rolling or subsequently, slow enough so as to form
at least a partially recovered or recrystallised
structure but fast enough to ensure that sufficient of
the Cu is retained in solid solution to provide an age
hardening effect if a subsequent precipitation treatment
l0 is carried out,
optionally cold rolling the hot rolled sheet, and
optionally age hardening the cold rolled alloy, wherein
preferably after the essential hot rolling step the
rolled ingot is always maintained at a temperature below
its solvus temperature.
During cold rolling, the metal temperature generally
rises to about 100-200°C as it is passed through the mill.
Conventionally after cold rolling, the metal is coiled
and being so massive the coiled metal takes a long time
to cool down to room temperature. Phase precipitation
and hardening can occur during this cooling down period
without the need forcibly to cool the coil. Additional
cooling can, however, be used if required. If desired
after cold rolling re-heating can be effected if desired,
for example to control the amount of cold work in the
alloy. If this re-heating takes the alloy above its
solvus temperature then cooling is preferably effected at
a rate less than 2°C/sec to avoid distortion or to avoid
the need for a separate quench stage.
As an alternative to batch DC casting, the alloy
could be cast continuously by for example belt casting or
twin roll casting. These techniques allow thin strip to
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
6
be produced of a thickness of generally as low as 5mm,
and sometimes as low as 2mm. Such thin cast strip may or
may not require homogenisation before hot rolling since
it tends to cool so quickly that the Cu and Mg present
are likely to remain in solid solution.
The casting could be extruded using direct or
indirect extrusion. Preferably the casting is
homogenised as described above and then cooled to room
temperature before being re-heated to 400 to 500°C for
extrusion. Alternatively the casting can be cooled
directly from its homogenisation temperature to the
desired extrusion temperature.
The extrudate is cooled preferably with still air or
with forced air. If desired, the extrudate can be re-
heated to above the solvus temperature of the alloy and
then cooled at a rate of less than 2°C/sec. This re-
heating treatment may be needed for texture and/or grain
size control. After extrusion the extrudate is generally
stretched by about ~ to 2% and then aged.
The present invention has particular applicability
for the production of can stock, especially can end stock
(CES) which possesses a combination of high strength and
formability. The combination of composition and process
of the present invention overcomes many of the
manufacturing difficulties of the conventional AA5182
sheet currently in use and is capable of producing CES at
lower cost. It also improves the subsequent performance
of the can end, most notably its scoreline corrosion
resistance. The invention is particularly suitable for
downgauging to produce lighter weight can ends, i.e.
gauges down to say 0.150mm.
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
7
For the production of can end stock, the preferred
method is to cast an ingot, homogenise it, and hot roll
to, say, 2mm to form strip. A key aspect of the
invention is that the strip does not need an additional
solution heat treatment step. Furthermore, even if it
does, the material does not need to be rapidly cooled,
e.g. does not need to be quenched into water; the cooling
is generally air cooled (possible forced air). The coil
is then cold rolled to final gauge and lacquered.
The range (in weight percent) for the principal elements
over which this invention is operable is:
Magnesium: 1.0 - 4.0 wt. o, preferably 2.0 - 4.0,
still more preferably 2.5 to 4.0o
Copper: 0.1 - 0.6 wt. o, preferably 0.2 - 0.5,
still more preferably 0.2 to 0.4%
Manganese: up to 0.8 wt. o, preferably up to 0.6, more
preferably up to 0.5, still more
preferably up to 0 . 4 0 . For some alloys a
minimum Mn content of 0.1% is preferred.
Iron: up to 0.5 wt.%,, preferably 0.1 - 0.3%
Silicon: up to 0.3wt.%, preferably up to 0.20
Chromium: up to 0.15%, preferably trace
Titanium: up to 0.15, preferably up to 0.050
Boron: up to 0.05, preferably up to 0.010
Carbon: up to 0.05, preferably up to 0.010
For grain refining of the casting either TiB2 or TiC can
be used, but generally not together.
The present invention will now be described in more
detail with reference to the accompanying drawings in
which:
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
8
Figure 1 shows a thermodynamic calculation of the solvus
temperature for S-phase precipitation in A1-xoMg-yoCu-
0.25Mn-0.2Fe-0.12Si,
Figure 2 shows the conductivity changes (MACS) during
isothermal annealing of an A1-3Mg-0.4Cu-0.25Mn-0.2Fe-
0.12Si alloy after solution heat treatment and cold water
quenching,
l0 Figure 3 shows the conductivity changes (oIACS) during
isothermal annealing of an A1-3Mg-0.4Cu-0.25Mn-0.2Fe-
0.12Si alloy after solution heat treatment, cold water
quenching and cold rolling, and
Figure 4 are curves showing the effect of time and
temperature on the extent of recrystallisation during
isothermal annealing of an Al-3Mg-0.4Cu-0.25Mn-0.2Fe-
O.I2Si alloy after solution heat treatment, cold water
quenching and cold rolling.
The theoretical basis for the present invention is as
follows:
The basic premise is to select an alloy composition
which will enable solute to be kept in solid solution
during cooling from hot rolling temperatures (250°C to
400°C, say) . The strip is then processed to bring out a
precipitation hardening phase which provides extra
strength. This precipitation forms preferentially on the
dislocation structure introduced during cold deformation
In the case of CES this cold deformation is cold rolling,
for extrusions it is stretching, and for sheet it is
during forming of the sheet when it is fabricated into a
component.
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
9
Although there is a thermodynamic driving force for
the solute to be removed from solid solution during hot
working and subsequent cooling, the nucleation and
diffusion effects are such to keep a substantial amount
of solute in solution, i.e. 'missing the nose of the c-
curve'. Accompanying Figure 1 shows a calculation of the
solvus temperatures for a range of A1-Cu-Mg alloys. This
shows that the solute will stay in solid solution above
the temperatures indicated. Thus, the solute can not
l0 start to come out of solid solution until the strip is at
or below this temperature. It should be noted, that even
if the solute does start to come out of solid solution,
there may still be sufficient solute available to provide
an appreciable strengthening effect during subsequent
processing.
The conductivity has been determined for a
3Mg-0.4Cu-0.25Mn-0.2Fe-0.12Si alloy (wt.o) to demonstrate
that for this type of alloy there is a barrier to
nucleation and growth of the precipitates which can be
commercially exploited to provide an improved balance of
strength and formability. Accompanying Figure 2 shows
the effect of isothermal ageing on the conductivity of a
full solution heat treated and cold water quenched
material subject to isothermal ageing. This shows that
at temperatures below the solvus the conductivity
increases (indicating Cu along with Mg removed from solid
solution), but that at lower temperatures the
precipitation becomes difficult. Thus, the solute can be
kept in solid solution if the strip can be cooled to
these temperatures sufficiently rapidly.
If there are dislocations present then the
conductivity rise is more rapid, since the precipitating
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
phase is believed to be S-phase (an AlCuMg phase), or its
metastable precursors which is well-known to nucleate
preferentially on dislocations. To demonstrate this, a
further set of isothermal ageing experiments have been
5 performed on the same alloy, but after solution heat
treatment, cold water quenching and cold rolling. This is
shown in accompanying Figure 3. In this case the
conductivity drop starts to occur after a few seconds.
This shows the importance of passing through this
10 temperature regime without large numbers of dislocation
present, since if the phase nucleates at these high
temperatures it is likely to be relatively coarse and
provide little strengthening. The example shown is an
extreme example since the strip was cold rolled to
introduce a high dislocation density prior to ageing. In
hot deformation the dislocation density is lower for a
fixed level of macroscopic strain, thus providing fewer
sites for nucleation of the precipitates.
For the production of CES the hot rolling conditions
are selected to ensure that the hot rolled sheet
recrystallises on or before coiling or very shortly
thereafter. Preferably the sheet is fully recrystallised
resulting in a low dislocation density.
Recrystallisation is encouraged by arranging for the
minimum temperature of the sheet as it exits from the
rolling mill to be 250°C, preferably 270°C and more
preferably 300°C and/or arranging for thewcooling rate of
the sheet to be sufficiently slow to allow time for the
sheet to recrystallise when in its coiled form or during
coiling. In a conventional mill the coiling temperature
is approximately the same as the exit rolling mill
temperature. Where additional cooling means are provided
after the mill the minimum coiling temperature should be
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
11
in the range of minimum mill exit temperatures mentioned
above. In practice acceptable cooling rates are found to
be of the order of 0.1 to 10°Clminute and preferably 0.2
to 5°C/minute over the temperature range of 400-200°C.
There is no need to uncoil the sheet during cooling in
order, for example, to quench it.
An indication of the time required to achieve
recrystallisation has been determined for a 3Mg-0.4 Cu-
0.25Mn-0.2Fe-0.12Si alloy (wt.o). This material was
solution heat treated, cold water quenched and cold
rolled 500. Isothermal heat treatments were performed to
determine the extent of recrystallisation, as shown in
Figure 4. This shows that after this deformation, full
recrystallisation is possible within a few minutes at
temperatures in excess of around 320°C. It should be.
noted that the precise details of the recrystallisation.
kinetics will depend on the deformation conditions and
the material microstructure.
A high rolling mill exit temperature encourages
precipitation of S phase or its precursors while the
strip or coil is cooling. Cooling more quickly can
counter this and prevent precipitation but if the exit
temperature becomes too high, the cooling rate required
is too fast to be practically useful. To take maximum
advantage of the rapid cooling during hot rolling, the
upper limit to the mill exit temperature, especially for
the alloys richer in Cu and Mg, should preferably be
lower than the solvus temperature of the alloy. Figure 1
gives an indication of the solvus temperature as a
function of the Mg and Cu contents. Preferably the
maximum exit temperature should be between 340°C and 360°,
although up to 380°C is possible for some alloys.
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
12
It is important to note that the location of the
nose of the c-curve for these alloys when recrystallised
varies with the composition of the alloy. For example,
for the alloy referred to in Figure 2, the nose of the
curve is located at a time of around 100 to 1000 seconds.
For dilute alloys the nose is moved to longer times
whilst for more concentrated alloys the nose is moved to
shorter times. The time indicated in Figure 2 compares
with times of between 1 and 100 seconds for conventional
age hardening systems such as AA7075, AA2017, AA6061 and
AA6063. For the alloys described in the present
invention, this provides longer times at temperatures
below the solvus temperature in which to cool the strip
and still maintain the Cu (and Mg) in solid solution.
For this preferred alloy of Figure 2 it has been found
that a cooling rate of 1°C/min and preferably 5°C/min is
sufficient substantially to miss the nose of the c-curve
and provide a substantial age hardening response during
subsequent processing. This cooling rate can be achieved
by, for example, forced air cooling of a coil. Previous
art regarding solution heat treatment of these Al-Mg-Cu
alloys teaches that, not only is a separate solution heat
treatment stage required, but that the strip must be
quenched with a cooling rate of 2°C/second or faster. For
the present invention it has been found that neither of
these steps need to be used, thereby providing a lower
cost manufacturing route for these alloys. Likewise no
separate annealing step is needed after the hot working
step and before the cooling step.
This solute is then used to give a significant
precipitation hardening effect during subsequent
thermomechanical processing. During subsequent cold (or
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
13
warm) deformation of the strip an increased dislocation
density is introduced giving enhanced nucleation sites
for the strengthening phase. This deformation may not be
needed for all applications of this invention, since for
these compositions it is known that the precipitation can
also occur in the absence of dislocations, albeit at
slower rates. The precipitating phase is believed to be
S-phase which can form as needles or rods on the
dislocation structure. In the case of CES this
l0 precipitation could occur during a separate ageing step
or during the thermal history which the material would
experience during deformation in, for example, strip
rolling.
As shown above, it may be important to achieve rapid
recrystallisation in order to remove the dislocations
from the material as it cools. Mn can be added as a
strengthening element and to control grain size and is
therefore desirably kept as high as possible. However,
Mn inhibits recrystallisation after hot rolling or during
annealing, and so a maximum Mn content of 0.4% may have
to be set in order to achieve full recrystallisation for
some alloys under certain conditions. For many of the
alloys to assist in controlling the grain size of the
recrystallised sheet, it may be desirable to have a
minimum of at least 0.05%Mn and preferably at least
0.loMn present in the alloy. Recrystallisation may also
be important for crystallographic texture control in CES,
but this may not be necessary if, the can end tooling is
modified to take significantly higher levels of Baring
into account. Crystallographic texture control can also
be important for automotive sheet formability; another
potential application of this invention.
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
14
Another feature of the composition used in the
present invention is the importance of having low Fe and
Si in the alloy, since this will prevent the presence of
excessive numbers of coarse constituent particles in the
sheet. These form during solidification and cannot be
fully dissolved during homogenisation of the ingot.
Although they break up during rolling, their presence is
sufficient adversely to affect formability. Since this
invention has been found to produce improved formability
over existing AA5182 CES, the strip may be able to
tolerate higher levels of these elements, thus reducing
cost. Tolerance of higher levels of 5i and Fe may allow
greater use of recycled aluminium scrap and this is
another important aspect of this invention. Up to 0.5oFe
may be tolerated in the alloy and preferably up to
0.3%Fe. The minimum amount of Fe present will be
dictated by cost and there is unlikely to be less than
0.lFe. Silicon up to 0.3o may be present, preferably up
to 0.2~.
Another advantage over conventional AA5182 CES is
that the lower Mg content will also make the can end less
susceptible to stress corrosion cracking (SCC), which can
lead to catastrophic failure of the end under the
stressed conditions which are encountered in the
pressurised can. The invention described here will make
the end less sensitive to these conditions, since the
lower Mg content reduces beta-phase precipitation, which
has been linked to SCC. Avoidance of SCC is also
important in many other applications including car body
sheet.
CES is currently made from AA5182 and gets its
strength predominantly from a combination of solute
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
hardening and strain hardening. This makes it difficult
to roll and gives a relatively high manufacturing cost.
The alloy used in the present invention has lower
5 strength during the rolling operations, but develops its
strength during subsequent thermal exposure during
fabrication. Thus there is the benefit of rolling a
lower strength sheet, but still enabling the desired
sheet properties to be obtained ultimately. It is also
10 possible to produce a higher strength sheet suitable for
downgauging without a reduction in rollability (higher
rolling loads, more difficulties in performing the
rolling operation) encountered in higher Mg containing
alloys such as AA5182 and AA5019A.
The present invention is also applicable to
production of low cost automotive sheet, where the
material could be used in the hot rolled condition
(Direct Hot Roll to Gauge), thereby potentially avoiding
2o the need to solution heat treat the sheet.
Alternatively, the sheet could be cold rolled to gauge,
as for CES, with a final continuous anneal to impart the
formability required for this application and to take the
solute into solution. Cooling after annealing should be
sufficiently rapid to retain substantially all of the
solute in solution. Ageing could be carried out in a
separate operation before or after forming, for example
during the paint bake stowing of the automotive part.
Some embodiments of the present invention will now
be described by way of example:
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
16
Example 1
An alloy of the following composition was cast as a
225mm x 75mm cross section DC ingot;
Magnesium 3.0 wt.%
Copper 0.4 wt.%
Manganese 0.25 wt.%
Iron 0.20 wt.o
Silicon 0.12 wt.o
Balance aluminium with incidental
l0 impurities. The ingot was not grained refined during
casting and as a consequence the Ti level was 0.0018% and
B was less than 0.00010.
This was homogenised for 2 hours at 540°C (50°C/hr
heating rate), followed by laboratory hot rolling to 6mm.
During this rolling stage the temperatures were only
about 100-200°C, so the strip was re-solution heat treated
to bring about full recrystallisation and to put the
solute back into solid solution. This reproduces solute
levels more like those which would be found during
rolling on an industrial hot line (but prior to coiling).
Different heat treatments were then applied at this
gauge. The strip was either solution heat treated (SHT)
(5 minutes at 550°C) and cold water quenched (CWQ) or it
was solution heat treated and then air cooled to
temperatures in the range 300 to 340°C and then cooled at
1°C/min. Conductivity was measured at this stage to
determine how much solute remained in solid solution.
These conditions were selected to simulate the conditions
which might be expected to exist during commercial use of
this invention. Until the strip temperature drops below
the solvus temperature for the alloy the S phase therein
cannot precipitate and therefore the Cu (and Mg) would be
substantially in solid solution. The strip could then be
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
17
quenched at the end of hot rolling or, preferably, cooled
after coiling. During this process the starting
temperature could be in the range 300 to 340°C and a
typical initial cooling rate would be 1°C/min. The
temperature range between the solvus temperature (about
390°C for the alloy) and the coiling temperature is passed
through very quickly since this is when the strip might
typically be in the hot tandem mill and, hence, there is
lubricant applied to the strip which acts as a coolant.
This phase was simulated using the air cool from the
solution heat treatment temperature.
The strip was then cold rolled to 0.24mm and given a
simulation of a coil cool down to -ambient temperature
from 150°C at 0.4°C/min. It was then given a simulation
of a lacquer curing cycle for 3 minutes at 205°C. Tensile
testing was performed at each stage of the treatment and
the results compared with results on conventional AA5182
CES materials processed in the laboratory.
The effect of strength development was also studied
at various stages of the laboratory simulation of the CES
production route. An example is given below for this
alloy which has been solution heat treated at 2mm and
rolled to 0.20mm gauge. This is compared with AA5182
rolled in the laboratory using a simulation of the
commercial route for that alloy. The 0.2% yield strength
is shown in Table 1 below. The as-rolled strength was
found to be lower than AA5182, indicating easier rolling,
and the strength drop during coiling and lacquer stoning
simulation was less, showing the benefits of
precipitation hardening. In addition, in AA5182 CES the
softest direction is usually at about 45° to the rolling
direction of the sheet (softer by about 10-20 MPa) and
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
18
this is believed to control the buckle pressure of the
sheet. In this invention the levels of cold reduction
needed to generate the desired strength level are lower
and thus the weakest direction is likely to be this
longitudinal value. Hence, at its best, the combination
of the composition and processing route of the present
invention is capable of producing a strength level
approximately 45 MPa stronger than existing AA5182.
Table 1: Comparison of properties with conventional CES
Condition 5182 CES This alloy
As-rolled at final gauge 430 MPa 399 MPa
As-rolled and coil annealed 358 MPa 386 MPa
As-lacquered 345 MPa 370 MPa
Conductivity results are shown in Table 2 below. This
shows that the conductivity at the solution heat
treatment stage is capable of being increased from 33.1
to 35.0 if the solute is allowed to be removed from solid
solution, but that if the material is cooled to ambient
temperature at 1°C/min from 300°C there is only a fraction
of the increase in the conductivity (0.3o versus 1.90).
This implies that a significant amount of the solute is
kept~in solid solution, even at these cooling rates.
Table 2: Conductivity after different heat treatments at
2mm gauge
Condition Conductivity (o IACS)
Solution heat treated and CWQ 33.1
Solution heat treated and Fast Air Cooled 33.1
Solution heat treated and cooled from 340°C 33.9
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
19
Solution heat treated and cooled from 320°C 33.8
Solution heat treated and cooled from 300°C 33.4
SHT, cold worked and aged 14 hours at 320° 35.0
The strength developed in these materials at final CES
gauge after lacquer stoning is shown in Table 3 below.
In this case the sheet has been rolled to 0.24mm. This
shows that sufficient solute remains in solid solution
still to give an appreciable strength CES. Bend testing
has also been performed and indicates an improvement in
the amount of bending which can be performed prior to
failure when compared with conventional AA5182 CES.
Table 3: Strength developed after processing to 0.24mm
after various thermal treatments at 2mm 'hotband' gauge.
condition 0.2% Proof Stress (MPa
Solution heat treated and CWQ 350 MPa
Solution heat treated and cooled from 340C 327 MPa
Solution heat treated and cooled from 320C 329 MPa
Example 2
An alloy of the following composition was DC cast for
processing within an industrial plant:
Magnesium 2.9 wt.%
Copper 0.4 wt.o
Manganese ' 0.1 wt.o
Iron 0.20 wt.%
Silicon 0.08 wt.o
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
Balance aluminium with incidental
impurities. The ingots were cast with additional grain
refiner.
5 The ingot were homogenised at 540°C and hot rolled on a
single stand reversing mill to a thickness of 38mm at
which point the temperature was around 480°C. The strip
was then hot rolled through a 3-stand hot tandem mill to
a gauge of 2.5mm. The conditions were adjusted to give
10 two different coiling temperatures in order to show the
effects at opposite extremes of this invention. In both
cases the coils were forced-air cooled, giving a cooling
rate measured on the outer laps of the coil of around
0 . 7°C/min.
The cooler coil was processed to give a sidewall
temperature of 280-290°C. In this instance the
microstructure of the strip was largely unrecrystallised.
As a consequence the solute was easily removed from solid
solution on the pre-existing dislocation structure from
the hot deformation. The conductivity of this strip is
shown in Table 4, showing that the oIACS value is similar
to that in which all of the precipitation has been
allowed to occur. Also in Table 4 is presented the
conductivity obtained by using a still-air cool on strips
of the 2.5mm thick metal at the end of the hot rolling
(approximately 60°C per minute), showing that at these
cooling rates a significant amount of the solute can be
kept in solid solution.
The hotter coil was processed to give a coil sidewall
temperature of 330-340°C. Table 4 shows that in this case
the forced air cooling leaves more solute in solid
solution as a consequence of the fully recrystallised
CA 02431029 2003-06-11
WO 02/50329 PCT/GBO1/05686
21
grain structure achieved with the higher coiling
temperature. The amount of solute in solid solution with
the faster cool is even higher and approaches that of the
conventional solution heat treated (SHT) and cold water
quenched (CWQ) material. This shows that a cooling rate
of 0.7°C/min is able to keep some of the copper in solid
solution, but that more rapid cooling leaves more copper
in solid solution and yet is still fully recrystallised.
Thus, cooling the coil with forced-air from a temperature
lower than 330°C will achieve a similar effect (i.e. more
solute in solid solution), since the c-curve will be
substantially missed in that case too. The forced-air
cooled coil was cold rolled to 0.216mm and the as-rolled
tensile yield strength measured as 347 MPa.
Between these two limits of cooling temperature there
will be even more solute in solid solution at the end of
hot rolling and thus even higher strength sheet can be
produced.
Table 4: Conductivity after different thermomechanical
treatments in an industrial plant
Condition Conductivity (oIACS
Solution Heat Treated and CWQ 35.4
SHT, CWQ + 24 hrs. at 310°C 36.8
Forced-air cooled coil from 280-290°C 36.9
air cooled strip from 280-290°C 36.1
Forced-air cooled coil from 330-340°C 36.4
air cooled strip from 330-340°C ~ 35.9
