U.S. patent application number 10/451217 was filed with the patent office on 2004-05-13 for age-hardenable aluminium alloys.
Invention is credited to Mahon, Gary John.
Application Number | 20040089379 10/451217 |
Document ID | / |
Family ID | 9905491 |
Filed Date | 2004-05-13 |
United States Patent
Application |
20040089379 |
Kind Code |
A1 |
Mahon, Gary John |
May 13, 2004 |
Age-hardenable aluminium alloys
Abstract
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. There is described 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, Cooper : 0.1 to 0.6,
Manganese : up to 0.8, Iron : up to 0.5, Silicon : up to 0.3,
Chromium : up to 0.15, Titanium : up to 0.15, Balance : Aluminium
with incidental impurities b) optionally homogenising the cast
alloy, c) hot working the casting at an initial temperature of at
least 400 DEG 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 solvus temperature of the alloy, d) cooling
the intermediate product either during hot working or in a
subsequent step at a rate such that at least a partially recovered
or recrystallised structure is formed ant 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. The
described method is particularly suited to the production of can
end stock and sheet for automotive applications.
Inventors: |
Mahon, Gary John; (Oxon,
GB) |
Correspondence
Address: |
BROWDY AND NEIMARK, P.L.L.C.
624 NINTH STREET, NW
SUITE 300
WASHINGTON
DC
20001-5303
US
|
Family ID: |
9905491 |
Appl. No.: |
10/451217 |
Filed: |
December 10, 2003 |
PCT Filed: |
December 20, 2001 |
PCT NO: |
PCT/GB01/05686 |
Current U.S.
Class: |
148/552 |
Current CPC
Class: |
C22C 21/06 20130101;
C22F 1/047 20130101 |
Class at
Publication: |
148/552 |
International
Class: |
C22F 001/04 |
Foreign Application Data
Date |
Code |
Application Number |
Dec 20, 2000 |
GB |
0031104.3 |
Claims
1. 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:
9 Magnesium 1.0 to 4.0 Copper 0.1 to 0.6 Manganese up to 0.8 Iron
up to 0.5 Silicon up to 0.3 Chromium up to 0.15 Titanium up to 0.15
Balance Aluminium with incidental impurities
b) optionally homongenising the cast alloy, c) hot working the
casting at an initial temperature of at least 400.degree. 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 solvus temperature of the alloy, d) cooling the intermediate
product either during hot working or in a subsequent step at a rate
of less than 5.degree. 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.
2. A method as claimed in claim 1 wherein the alloy has the
following composition expressed in weight percent:
10 Magnesium 2.0 to 4.0 Copper 0.2 to 0.5 Manganese up to 0.6,
preferably up to 0.5 Iron 0.1 to 0.3 Silicon up to 0.2 Chromium up
to 0.15 Titanium up to 0.05 Boron or Carbon up to 0.01 Balance
Aluminium with incidental impurities
3. A method as claimed in claim 2 wherein the magnesium content is
2.5 to 4.0%.
4. A method as claimed in any one of claims 1 to 3 wherein the
intermediate product has a substantially fully recovered or
recrystallised structure.
5. A method as claimed in any one of the preceding claims wherein
the casting is homogenised before hot working at a temperature of
at least 480.degree. C., preferably 500 to 600.degree. C., so that
substantially all of the magnesium and copper in the casting are in
solid solution.
6. A method as claimed in any one of the preceding claims wherein
the casting is hot worked, optionally with re-heating of the
casting to above the alloy's solvus temperature, and preferably at
least 450.degree. C., to take substantially all of the magnesium
and copper present into solid solution.
7. A method as claimed in any one of the preceding claims wherein
the hot working step is carried out when the casting has an initial
temperature of from 450.degree. C. to 580.degree. C.
8. A method as claimed in any one of the preceding claims wherein
the alloy is DC cast.
9. A method as claimed in any one of the preceding claims including
the step of cold rolling the hot worked casting, optionally with
coiling.
10. A method as claimed in any one of claims 1 to 9 wherein the hot
working is effected by extrusion.
11. A method as claimed in any one of claims 1 to 9 wherein the hot
working is effected by hot rolling.
12. A method as claimed in any one of the preceding claims wherein
the hot worked casting is cooled at a rate of less than 1.degree.
C./min.
13. A method as claimed in any one of the preceding claims wherein
if after the said hot working step the temperature of the
intermediate product exceeds the solvus temperature of the alloy
then cooling of the intermediate product to a temperature below the
alloy's solvus temperature is effected at a rate less than
2.degree. C./sec.
14. A method as claimed in claim 12 or claim 13 wherein the cooling
is effected by forced air cooling.
15. A method as claimed in any one of the prceeding claims wherein
no separate annealing step is carried out after the hot working
step (c)and before the cooling step (d).
16. A method as claimed in any one of the preceding claims wherein
the product of the method is can end stock.
Description
[0001] 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.
[0002] In the art AA5000 series alloys are usually regarded as
non-heat treatable alloys i.e. they are not regarded 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.
[0003] 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.
[0004] 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.
[0005] 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.
[0006] WO-A-98/01593 describes a process for producing aluminium
alloy can body stock but again a separate annealing step is
required.
[0007] 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.
[0008] U.S. Pat. No. 5,655,593 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.
[0009] U.S. Pat. No. 3,464,866 describes a process for obtaining
aluminium alloy conductors but again teaches rapid cooling.
[0010] In accordance with the present invention there is provided a
method of producing an age-hardenable aluminium alloy comprising
the steps of:
[0011] (a) casting an alloy of a composition comprising the
following expressed in weight percent:
1 Magnesium: 1.0 to 4.0 Copper: 0.1 to 0.6 Manganese: up to 0.8
Iron: up to 0.5 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
[0012] (b) optionally homogenising the cast alloy,
[0013] (c) hot working the casting at an initial temperature of at
least 400.degree. 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 solvus temperature of the alloy,
[0014] (d) cooling the intermediate product during hot working or
in a subsequent step at a rate of less than 5.degree. 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
[0015] (e) optionally allowing or arranging for phase precipitation
to occur in the alloy.
[0016] Preferably after the said hot working step the intermediate
product is generally maintained at a temperature below the solvus
temperature of the alloy, provided that if the intermediate product
is heated above the alloy's solvus temperature then cooling thereof
is effected at a rate less than 2.degree. C./sec.
[0017] 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.
[0018] The precipitation phase if formed is believed to be S phase
(an AlCuMg phase) or its metastable precursors.
[0019] 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.
[0020] 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:
[0021] optionally homogenising the casting at a temperature of at
least 480.degree. C., and preferably 500 to 600.degree. C., so that
substantially all of the magnesium and copper in the casting are in
solid solution,
[0022] optionally hot rolling the casting, optionally with
re-heating of the casting to above the alloy's solvus temperature,
preferably at least 450.degree. C., to take substantially all of
the magnesium and copper present into solid solution,
[0023] hot rolling the casting with a rolling mill entry
temperature of the casting of at least 400, and preferably from
450.degree. to 580.degree. C.,
[0024] 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 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 is carried out,
[0025] 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.
[0026] During cold rolling, the metal temperature generally rises
to about 100-200.degree. 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.degree. C./sec
to avoid distortion or to avoid the need for a separate quench
stage.
[0027] 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 be produced of a thickness of
generally as low as 5 mm, and sometimes as low as 2 mm. 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.
[0028] 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.degree. C. for extrusion. Alternatively the casting can be
cooled directly from its homogenisation temperature to the desired
extrusion temperature.
[0029] The extrudate is cooled preferably with still air or with
forced air. If desired, the extrudate can be reheated to above the
solvus temperature of the alloy and then cooled at a rate of less
than 2.degree. 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 1/2to 2% and then aged.
[0030] 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 0.150
mm.
[0031] For the production of can end stock, the preferred method is
to cast an ingot, homogenise it, and hot roll to, say, 2 mm 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.
[0032] The range (in weight percent) for the principal elements
over which this invention is operable is:
2 Magnesium: 1.0-4.0 wt. %, preferably 2.0-4.0, still more
preferably 2.5 to 4.0% Copper: 0.1-0.6 wt. %, preferably 0.2-0.5,
still more preferably 0.2 to 0.4% Manganese: up to 0.8 wt. %,
preferably up to 0.6, more preferably up to 0.5, still more
preferably up to 0.4%. 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 0.2% Chromium: up to 0.15%,
preferably trace Titanium: up to 0.15, preferably up to 0.05%
Boron: up to 0.05, preferably up to 0.01% Carbon: up to 0.05,
preferably up to 0.01%
[0033] For grain refining of the casting either TiB.sub.2 or TiC
can be used, but generally not together.
[0034] The present invention will now be described in more detail
with reference to the accompanying drawings in which:
[0035] FIG. 1 shows a thermodynamic calculation of the solvus
temperature for S-phase precipitation in
Al-x%Mg-y%Cu-0.25Mn-0.2Fe-0.12Si,
[0036] FIG. 2 shows the conductivity changes (%IACS) during
isothermal annealing of an Al-3Mg-0.4Cu-0.25Mn-0.2Fe-0.12Si alloy
after solution heat treatment and cold water quenching,
[0037] FIG. 3 shows the conductivity changes (%IACS) during
isothermal annealing of an Al-3Mg-0.4Cu-0.25Mn-0.2Fe-0.12Si alloy
after solution heat treatment, cold water quenching and cold
rolling, and
[0038] FIG. 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-0.12Si alloy after solution heat
treatment, cold water quenching and cold rolling.
[0039] The theoretical basis for the present invention is as
follows:
[0040] 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.degree. C. to 400.degree. 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.
[0041] 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 FIG. 1 shows a calculation
of the solvus temperatures for a range of Al--Cu--Mg alloys. This
shows that the solute will stay in solid solution above the
temperatures indicated. Thus, the solute can not 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.
[0042] The conductivity has been determined for a
3Mg-0.4Cu-0.25Mn-0.2Fe-0- .12Si alloy (wt. %) 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 FIG.
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.
[0043] If there are dislocations present then the conductivity rise
is more rapid, since the precipitating 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 performed on the same alloy, but after solution heat
treatment, cold water quenching and cold rolling. This is shown in
accompanying FIG. 3. In this case the conductivity drop starts to
occur after a few seconds. This shows the importance of passing
through this 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.
[0044] 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.degree. C., preferably 270.degree. C. and more preferably
300.degree. C. and/or arranging for the cooling 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 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.degree. C./minute and preferably 0.2 to
5.degree. C./minute over the temperature range of 400-200.degree.
C. There is no need to uncoil the sheet during cooling in order,
for example, to quench it.
[0045] 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. %). This material was solution
heat treated, cold water quenched and cold rolled 50%. Isothermal
heat treatments were performed to determine the extent of
recrystallisation, as shown in FIG. 4. This shows that after this
deformation, full recrystallisation is possible within a few
minutes at temperatures in excess of around 320.degree. C. It
should be noted that the precise details of the recrystallisation
kinetics will depend on the deformation conditions and the material
microstructure.
[0046] 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. FIG. 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.degree. C. and
360.degree., although up to 380.degree. C. is possible for some
alloys.
[0047] 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
FIG. 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 FIG. 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 FIG. 2 it has been found that a cooling rate of 1.degree.
C./min and preferably 5.degree. 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.degree. 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.
[0048] This solute is then used to give a significant precipitation
hardening effect during subsequent thermomechanical processing.
During subsequent cold (or 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 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.
[0049] 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 0.1% Mn 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 earing into account.
Crystallographic texture control can also be important for
automotive sheet formability; another potential application of this
invention.
[0050] 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 Si and
Fe may allow greater use of recycled aluminium scrap and this is
another important aspect of this invention. Up to 0.5% Fe 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 0.1Fe. Silicon up to 0.3% may be present,
preferably up to 0.2%.
[0051] 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.
[0052] CES is currently made from AA5182 and gets its strength
predominantly from a combination of solute hardening and strain
hardening. This makes it difficult to roll and gives a relatively
high manufacturing cost.
[0053] The alloy used in the present invention has lower 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
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.
[0054] 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 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 stoving of the automotive part.
[0055] Some embodiments of the present invention will now be
described by way of example:
EXAMPLE 1
[0056] An alloy of the following composition was cast as a 225
mm.times.75 mm cross section DC ingot;
3 Magnesium 3.0 wt. % Copper 0.4 wt. % Manganese 0.25 wt. % Iron
0.20 wt. % Silicon 0.12 wt. %
[0057] Balance aluminium with incidental impurities. The ingot was
not grained refined during casting and as a consequence the Ti
level was 0.0018% and B was less than 0.0001%.
[0058] This was homogenised for 2 hours at 540.degree. C.
(50.degree. C./hr heating rate), followed by laboratory hot rolling
to 6 mm. During this rolling stage the temperatures were only about
100-200.degree. 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).
[0059] Different heat treatments were then applied at this gauge.
The strip was either solution heat treated (SHT) (5 minutes at
550.degree. C.) and cold water quenched (CWQ) or it was solution
heat treated and then air cooled to temperatures in the range 300
to 340.degree. C. and then cooled at 1.degree. 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 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.degree.
C. and a typical initial cooling rate would be 1.degree. C./min.
The temperature range between the solvus temperature (about
390.degree. 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.
[0060] The strip was then cold rolled to 0.24 mm and given a
simulation of a coil cool down to ambient temperature from
150.degree. C. at 0.4.degree. C./min. It was then given a
simulation of a lacquer curing cycle for 3 minutes at 205.degree.
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.
[0061] 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 2 mm and rolled to 0.20 mm 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 stoving simulation was less, showing the
benefits of precipitation hardening. In addition, in AA5182 CES the
softest direction is usually at about 45.degree. to the rolling
direction of the sheet (softer by about 10-20 MPa) and 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.
4TABLE 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
[0062] 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.degree. C./min from
300.degree. C. there is only a fraction of the increase in the
conductivity (0.3% versus 1.9%). This implies that a significant
amount of the solute is kept in solid solution, even at these
cooling rates.
5TABLE 2 Conductivity after different heat treatments at 2 mm gauge
Condition Conductivity (% IACS) Solution heat treated and CWQ 33.1
Solution heat treated and Fast Air Cooled 33.1 Solution heat
treated and cooled from 340.degree. C. 33.9 Solution heat treated
and cooled from 320.degree. C. 33.8 Solution heat treated and
cooled from 300.degree. C. 33.4 SHT, cold worked and aged 14 hours
at 320.degree. 35.0
[0063] The strength developed in these materials at final CES gauge
after lacquer stoving is shown in Table 3 below. In this case the
sheet has been rolled to 0.24 mm. 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.
6TABLE 3 Strength developed after processing to 0.24 mm after
various thermal treatments at 2 mm `hotband` gauge. 0.2% Proof
Condition Stress (MPa) Solution heat treated and CWQ 350 MPa
Solution heat treated and cooled from 340.degree. C. 327 MPa
Solution heat treated and cooled from 320.degree. C. 329 MPa
EXAMPLE 2
[0064] An alloy of the following composition was DC cast for
processing within an industrial plant:
7 Magnesium 2.9 wt. % Copper 0.4 wt. % Manganese 0.1 wt. % Iron
0.20 wt. % Silicon 0.08 wt. % Balance aluminium with incidental
impurities. The ingots were cast with additional grain refiner.
[0065] The ingot were homogenised at 540.degree. C. and hot rolled
on a single stand reversing mill to a thickness of 38 mm at which
point the temperature was around 480.degree. C. The strip was then
hot rolled through a 3-stand hot tandem mill to a gauge of 2.5 mm.
The conditions were adjusted to give 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.degree. C./min.
[0066] The cooler coil was processed to give a sidewall temperature
of 280-290.degree. 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 %IACS 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 2.5 mm thick metal at the end of
the hot rolling (approximately 60.degree. C. per minute), showing.
that at these cooling rates a significant amount of the solute can
be kept in solid solution.
[0067] The hotter coil was processed to give a coil sidewall
temperature of 330-340.degree. 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 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.degree. 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.degree. 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 0.216 mm and the
as-rolled tensile yield strength measured as 347 MPa.
[0068] 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.
8TABLE 4 Conductivity after different thermomechanical treatments
in an industrial plant Conductivity Condition (% IACS) Solution
Heat Treated and CWQ 35.4 SHT, CWQ + 24 hrs. at 310.degree. C. 36.8
Forced-air cooled coil from 280-290.degree. C. 36.9 air cooled
strip from 280-290.degree. C. 36.1 Forced-air cooled coil from
330-340.degree. C. 36.4 air cooled strip from 330-340.degree. C.
35.9
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