Note: Descriptions are shown in the official language in which they were submitted.
2t85216
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Wo 95/25825 ~ r~"~i~ sl~
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AL~MINIU~ FOIL
This invention is c~n~-orn~d with aluminium
foil having improved strength. In the current Al-Fe-Mn
based foil alloys, such as AP. 8006 and AA 8014, the
good balance of strength and formability of thin gauge
10 foil is obtained by achieving a combination of fine
grain size after final Anne~l ;n5 and dispersion
strengthening. This invention describes the use of an
additional strengthening ~ An; om to achieve increased
strength; namely solid solution strengthening, and
5 specif ies the range within which the solute level must
be controlled in order to avoid loss of other
beneficial properties associated with the solute-free
versions of these alloys.
British Patent Specification 1 479 429
20 described dispersion-strengthened aluminium alloys
based on the Al-Fe-Mn system, such as AA ~006 and
AA 8014. (from Registration record of ;nt~rnAtional
alloy designations and chemical composition limits for
wrought Al and wrought Al alloys, A~ Inc. May 1987) .
25 The as-cast ingot comprised unaligned intermetallic
rods. These were broken up during working to provide a
wrought aluminium alloy product cr~n~;n;n~ dispersed
;nt~ -tAlliC particles. The invention was applicable
to the production of rolled sheet, which was to some
3 extent anisotropic. It was possible to reduce the
relative proportions of the anisotropy by introducing
small proportions of Cu and/or Mg which 1~ ; n~d in
solid solution in the Al phase and had known strength
providing properties. The 1055 of anisotropy implies
35 discontinuous recrystallisation and 1055 of grain size
control, which changes would have been acceptable in
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the sheet products mainly envisaged and ~X~ r 3~
The successful production of aluminium foil having useful
properties depends on several critical pd~d~ r~. The metal to be rolled
must not be too hard, otherwise rolling down to the very low ~l ,ickl,e:.a~s
below 100 ,um required is not co"""e",i~"y viable. After rolling, the foil has
to be heated, to a temperature sufficient to remove rolling lubricant but not
so high that adjacent sheets of foil stick together. This temperature
window is quite narrow, generally 220 - 300C, and results in a final
annealing treatment of the foil. During this annealing treatment,
recry~ takes place, and it is necessary that this be continuous
recry: ' " " n, which retains a desired small grain size, rather than
discontinuous recry~ , which results in grain growth. If large
grains are present, the foil has reduced Ill~ulldlli~dl properties. While
these critical pdldlllt:L~ have long been achieved using Al-Fe-Mn alloys, it
was not apparent that they could be achieved in CUI I Ibi~ IdliUI I with solid
solution hardening. And indeed, as the inventors have discovered, the
nature and amount of solute that can be added is critically circumscribed.
In one aspect this invention provides aluminium foil
composed of an alloy of co",l,o~i~ion by weight %:
Fe 1.2 - 2.0%
Mn 0.2 - 1.0%
Mg and/or Cu 0.1 - 0.5%
Si up to 0.4%
Zn uptoO.1%
Ti up to 0.1%
balance Al of at least c~" ", It31 Uidl purity
which foil has an average grain size beiow 5 ~um after annealing.
In another aspect, the invention provides
~MENDED SHEET
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Wo 9~2582
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aluminium foil of the stated composition, wherein at
least 50g~ by volume of the as-rolled texture is
retained after final anneal.
In another aspect, the invention provides
aluminium foil of the stated composition, wherein
the crystallographic texture of the final
~nn~ 1 product is a retained rolling texture.
The aluminium foil preferably ha9 a thickness
below 100 ~Lm, particularly in the range 5 - 40 llm e.g.
10 - 20 /~m. The improved strength of foil according to
this invention should enable thinner gauge9 to be
marketed .
Fe and Mn are present to provide dispersion
strengthening properties, as described in the aforesaid
GB 1 479 42g. Preferably the Fe content is 1.4 - 1.896;
the Mn content is 0 . 3 - 0 . 6%; and the Fe + Mn content
is 1.8 - 2.15~6.
If the Fe + Mn concentration in the melt
exceeds this value of 2.15 then coarse primary
intermetallic particles (typically up to 100 ~lm length)
can form during soli~;fir~t;nn as a consec~uence of
nucleation of these phases on the cooler parts of the
molten metal distribution system. These coarse
particles will break-up somewhat during subse~uent
processing but will still persist as relatively coarse
non-deformable particles in the final product. For the
case of sheet products this will not cause 9i~nif icant
problems, but in the case of foil products will give
rise to problems with pin-hole formation in the rolled
3 strip and give rise to excessive strip breaks during
processing. It is thus preferred to ca9t a composition
where primary intermetallic particles cannot form, and
this imposes an upper limit on the Fe and Mn levels for
use of this invention for foil products.
Mg and/or Cu is added to provide solution
strengthening, in a cnnt-~ntration of 0.1 - 0.5%
. - 2185216
Wo 95l25825 _ r~ . s.
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preferably 0.15 - 0.359~. At the lower end of these
ranges, little strengthening is observed. At the upper
end of these ranges, there is a risk that the solute
will encourage disrnnt; nllnus recrystallisation and will
5 result in undesired grain growth. This risk iB
particularly apparent at relatively high ~nn~Al in~
temperatures. As shown in the examples, Mg provides a
better solution strengthening effect than Cu at
equivalent cnnr~ntrations and i8 accordingly preferred.
The inventors have tried other solution
strengthening elements, but have f ound that they tend
to encourage discnnt;nllnllc recrystallisation during
final anneal or are otherwise lln~tic~Actnry. It is
therefore believed that Mg and Cu are the only two
15 usable solution strengthening additives.
Si and Zn are included in the AA
specifications of AA 8006 and AA 8014. But they are
preferably not deliberately included here. It is an
advantage of the invention that recycled scrap metal
20 can be used to make the foil.
The foil is specified as having an average
(or mean) grain size below 5 ~m, preferably below 3 ~m.
The grain size is preferably subst~nti~lly uniform, and
is achieved as a result of rnnt;nllnus recryst:~ll;R;~t;nn
25 during final anneal. Alternatively a non-uniform grain
size may be acceptable provided that gross
discnnt;nllmlc recryst~ll;qa~t;nn during final anneal is
avoided. For example, the majority of grains may have
a size of 2-3 ~Lm with a minor proportion of grains of
3 10-30 l~m. This duplex grain size structure may reduce
the ductility of the f oil, but the overall properties
may nevertheless be satisfactory.
Grain size may be det~rmi nPd by the mean
linear intercept method. On a micrograph of a section
35 of the alloy under test, a line (e.g. a straight line or
a circle) of known length is drawn, and a count is made
21 85216
Wo95l25825 1~I/~1. ~ :
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of the number of intercepts of: that line with grain
boundaries. The mean linear intercept grain size (mean
grain size) is the length of the line divided by the
number of intercepts.
The foil is generally anisotropic. Cold-
rolling develops an as-rolled texture typical of
dilute Al alloys. Texture i8 conv-nt;nn~lly measured
from an orientation distribution function in terms
of six parameters (cube, goss, copper, S, brass
and random). C~onventionally, these are measured
a8 a volume fraction of crystals orientated
over a ~15 spread about the c.~r.~ iate
Miller indices which are {001}~100~, lllO}<001~,
{112}<111~, {123}<634~, {011}<211~ respectively, the
5 random t being the r~ ;n;n~ volume fraction.
The copper, S and brass c ^ntq are generated by
cold rolling . DiSrnnt; nllnus recrystallisation would
tend to destroy the as-rolled texture and favour the
formation of cube and/or goss and/or random. In the
20 foil of this invention, at least 509,i by volume,
preferably 75~ of this as-rolled texture as represented
by the copper, S and brass components is retained after
f inal anneal . Preferably the crystallographic texture
of the final ~nn~ product is substantially the same
25 as the as-rolled product with no significant levels
of recryst~ll;c~t;nn texture r_ ^ntC.
It has surprisingly been found that the foil
of this invention may have a surface roughness greater
than that of its solute-free counterpart. This
3 increase in ro~hn~cq was confirmed by optical
profilometry (Perthometer) measurements, giving an Ra
of 0.38 for foil of this invention ~Example 2) compared
with an Ra of 0.24 for a commercial foil of
corr~qp~ nr1; ng composition without Mg . The rougher
35 surface improves the matt appearance of the foil.
In making the aluminium foil of this
: 2185216 ~
WO95l2~825 1~I,... ~c
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invention, a molten aluminium alloy of desired
composition iE cast, e.g. by direct chill (D.C. ~
casting, or alternatively by roll casting or belt
casting or other known casting techniques. The cast
5 metal is rolled by successive rolli~g steps in
conventional manner down to the re~uired foil
th; nkn~qq . These steps typically involve hot rolling
followed by cold rolling, possibly with one or more
inter~nn-~l in~ steps. Finally, the fQil is heated to a
10 temperature su-ficient to remove the rolling lubricant.
The heating rate is preferably 1C - 100C per hour.
As noted above, this temperature is typically in the
range 220 - 3D0C, preferably 230 - 280C, more
preferably 230 - 250C, and also effects continuous
5 recrystallisation of the foil. The aluminium foil of
this invention is preferably subst;~nt;Ally free of
surface ~nnt~m;n;ltion by rolling lubricar,t.
The technical basis of the invention, as
presently understood by the i~ventors, is ~ ; n~d in
20 the following paragraphs.
A range of aluminium alloys are known to
achieve a fine grain size after final ~nn~l inS by a
gradual coarsening of the cold-rolled substructure,
sometimes called ~nnt;nllml~ recryst~ll;q~t;nn, which
25 allows a good ~ ' ;n~tion of strength and formability
to be achieved . During the f inal Ann~A l; n~ of
~ lmin;llm foil products it is important to avoid the
oc~uLLt:~lce of large recrystallised grains which
severely ~;m;n;sh formability, often as a result of
3 strain loc~ t; on leading to premature f ailure during
loading. These grains are formed in the ~1Aqq;~
di8-nnt;nllnus manner whereby individual grainE nucleate
and grow to a large size. It is known that in these
type of alloyE this transition from dis~-nnt;nllnus to
35 cnnt~ nllnus recrystallisation occurs when the level of
:`
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WO ~snss2~ J/~ '
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cold work is increased above a critical level typical
of foil rolling.
If there is a sufficient high concentration
of non-deformable intermetallic particles, such as the
5 FeA16 and/or (FeMn)A16 eutectic rods formed during
solidification of Al-Fe-Mn alloys such as AA 8006 and
AA 8014, then after deformation to high strains these
particles must have increased dislocation activity
associated with them in order to Tn-int~;n r ^nt;nll;ty
10 across the aluminium/particle interface. Under
conYentional solute-free conditions, the8e dislocations
are capable of rearranging themselves into dislocation
walls, or sub-grain boundaries. As deformation
proceeds the geometrically necessary dislocations
5 generated during the rolling process crntinll~ to
migrate to, and recover into, the sub-grain boundaries,
increasing their misoriont~ti~n. Eventually these
boundaries will attain high misorientations with their
neighbours, i . e . high angle grain boundaries . When
20 these boundaries are then annealed, they can all
migrate at similar rates, thus encouraging r~ntinll~uS
recryst~l 1 i q~tion In addition this is helped by the
ability of the now broken up rod eutectic to pin grain
boundaries and prevent excessive rates of grain growth.
25 Dispersoids formed during hot processing of the ingot
will also assist this pinning process.
Thus, the conventional (solute-free) AA 8006
achieves a fine grain size after anneal, which imparts
the good balance of strength and ductility associated
3 with these alloys. The strength is inversely
proportional to the grain (or sub-grain) size, and
follows a d~l r.ol~t;r~nch;r
This invention still m~;nt~;nq this
strengthening m~h~n; r- whilst using the additional
35 strengthening ~ h~n; ~-- of solid solution
strengthening. If the amount of solute added is too
; ` 21 852~ 6
w09s/2s82s r~"~
-- 8
high then the ability to control the grain size during
the final anneal is lost, giving rise to a decrease in
grain size strengthening and formability. This
presumably is because dynamic recovery is prevented
5 during rolling, and 80 the driving force for
di8cnnt;nllnus recrystallisation is increased. This
al80 makes it increasingly difficult to roll the foil
to the reguired thin gauge because of the increased
rolled strength, giving a loss of the roll softening
10 normally found in solute-free alloys of this type.
Another aspect of the reaLLa.,ly t of
dislocations into high angle grain boundaries during the
rolling process is that the strength of the foil
decreases as the rolling s~rain is increased (roll
5 sof tening), instead of the usual roll hardening
associated with most aluminium alloys. Adding solute
to the alloy will hinder the ability of the
dislocations to rearrange themselves into low energy
configuration in the sub-grain boundaries, and will
20 prevent roll softening from occurring. Thus, if too
much solute is added the cold rolled strength of the
foil will be significantly increased, losing the
ability to roll the material to the thin gauges needed
for household foil and packaging applications (within
25 the range 40 ~Lm to 5 /~m).
Reference is directed to the ~~ ~nying
drawings in w~lich :-
Figure 1 shows the effect of ::nnP~l ;n~temperature on tensile strength of laboratory processed
3 alloys rolled to 140 l~m;
Figure 2 shows the effect of ~nn~ l in~
temperature oll tensile yield stress of laboratory
proceæsed alloys rolled to 140 ~m;
Figure 3 shows the effect of ~nn~l ;n~
35 temperature o~ tensile elongation of laboratory
processed alloys rolled to 140 llm; and
2185216
Wo95l25825 5 ~
Figures 4a and 4b are pole diagrams of a foil
sample before and after ;InnP~l ;n~,
Exa~Ple
5 Laboratorv Processinq
The effect of different levels of copper and
magnesium additions have been investigated using
laboratory processing of 200 mm x 75 mm cross-section
D.C. ingots of 1.6% Fe, 0.40% Mn, 0.15~ Si (denoted by
0 8006 in the figures) and modified alloys r~nt~;n;nr~
0 . 2% and 0 . 4~ of Cu or Mg . At the cooling rates
associated with this ingot cross-section, the addition
of the solutes does not prevent the formation of the
preferred rod eutectic, with only slight coarsening
15 being observed.
The above ingots have been heated to 525C,
hot rolled to 20 mm, and annealed at 330C for 3 hours
to simulate commercial hot processing. The materials
have then been cold rolled to 4.5 mm, inter~nnP~1Pd at
360C, and cold rolled to 145 ~m. This reproduces the
strain levels achieved during rolling of 14 ~m
household f oil . Table 1 shows the ef f ect of the
rolling rP~lllrt;r~n on the tensile strength of the
materials. Adding solid solution strengtheners
25 prevents the usual roll softening associated with
AA 8006 from occurring, thus; ~o'::;n~ an upper limit on
how much solute could be added and still enable thin
gauge products to be rolled commercially.
The 140 ~m foil has been annealed for 2 hours
3 at a range of temperatures using a simulation of batch
~nn~l; nrj, involving heating to temperature at
25C/hour and longitudinal tensile properties measured.
The variation of UTS, O . 2~6 proof stress, and
elongation-to-failure are shown in Figures 1, 2 and 3,
35 respectively, for the five alloys. All alloys
r- nt~;n;nrj the solute additions show an i, L~V. ~ in
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W0 9~2s82s F~~ a
-- 10 -
UTS over the solute-free AA 8006, with the ; , uv - t
being of the order of 20 to 40 MPa after the
commercially u~able anneal at temperatures in the
region of 220 - 260C ~owever, after ~nnP~l ;n~ at the
highest temperature investigated (300C)_ the more
c~n~-Pntr~tPd allûys have lower strengths than the 0 . 2~
additions. This is more prnn~unrP~ in the proof stress
data, and indicates that in the more r~n-Pntrated
alloys there is a loss in strength as a consequence of
loss of grain size control caused by rl;~cnntiml~lus
recrystallisation. This is confirmed by the optical
metallography of the grain structures after ;3nnP:~l in~
where coarse grained regions are apparent in the solute
~;n;n~ alloyg annealed at the highest temperatures.
The loss of grain size control is often associated with
loss of formability and ductility, although the
ductilities do not shûw any reduction here, pûssibly as
a consequence of the much thicker gauges P~m; nP~l here
(140 ~um vs 14 ~m) preventing strain localisation.
This 1088 of grain size control at the higher
temperatures in the 0.49~ rr~nt~;n;n~ alloys shows that
there will be ~n upper limit on the amount of solute
which can be added for solid solution strengthening
without running intû problems with loss of strength (in
particular yield strength) caused by coarser
recrystallised grains.
r 1~- 2
p~ ;~nt Trials
3 Based on the wish to achieve a significant
strength increase over the standard AA 8006 composition, a
full scale processing trial has been performed with
A~ 8006 plus 0.2 wt.~ Mg. This was DC cast as an ingot of
1600 mm x 600 mm cross section. The ingot was then
processed, the processing route consisting of hot
rolling to 3 mm, cold rolling to 450 ~m, and
21 85216
Wo 95/25825 ~ .~ 5,~c ~
interannealing at 360C. It was then cold rolled to
the f inal gauge of 14 ~lm .
Tensile testing of the as-rolled foil showed
the yield stress to be significantly higher than the
standard Mg-free version (Table 2).
Commercial batch anneal is carried out at a
temperature of 220 - 260C during a heating cycle of at
least 8 hours f rom room temperature to the ~nn~ 1; ng
temperature. Preferably the metal is held in the
temperature range f or at least 3 0 minutes . The total
cycle time depends on the coil width. Tensile
properties of the plant annealed foil are shown in
Table 2, showing that a significant strength
imp, ,-v~ t is achieved over AA 8006 . The results of
the plant trial clearly demonstrate that there will be
an upper limit on the amount of magnesium which can be
added to large cross-section D.C. cast ingot and still
give the re~uired microstructure for rnnt;nllmlR
recrySt;,11; R~t inn .
The process of rolling is highly anisotropic
as a result of crystal plasticity and inevitably leads
to a product with preferred orientations or
crystallographic texture. In order to describe
crystallographic texture a system has been devised that
25 enables reference directions on the sample to be
related to the crystallographic directions of a large
number of grains on a simple diagram called a pole
figure. The t~rhn;-r~ s for measuring crystallographic
texture in metals are well established and an excellent
3 reference is Hatherley and ~ trh;n~on "An Introduction
to Textures in Metals ~ The Institute of Metallurgists,
r/lnnnrr~rh No 5, 1979.
The crystallographic texture of the foil
samples before and after ~nnl~l ;nr, have bee~ rlf~t~rl~;nf~d
35 using x-ray diffraction from a laminate made from the
14 llm f oil . Figures 4a and 4b show the pole f igures
`~ '. ` 21 8521 6
Wo 95/25825 1
-- 12 --
generated from the ~111} aluminium planes orientated
relative to the rolling direction (vertical),
transverse direction (horizontal) and the foil plane
normal ~into the page). Figure 4a ls the as-rolled
5 f oil . Fi~ure 4b is the annealed f oil . The contour
levels are 1.00 1.60 2.20 2.80 3.40 4.00 4.60. This
shows that the crystallographic texture i8 PC5Pnt i ~ y
unaltered by the anneal, i . e . the texture is a retained
rolling texture. Pole figures corresponding to other
0 aluminium reflections have also been obtained from
which the Or;Pnt~t;nn Dlstribution Function (ODF) in
the rolled and annealed conditions have been generated.
The volume fractions of specific texture rn~monpnts
have been extracted from the ODF ' s and these are shown
5 in Table 3.
The grain size of the 14 llm foil has been
determined after commercial i~nnPAl ;n~ using the mean
linear intercept technique . This has been perf ormed on
micrographs ~ht~; nPr~ in the Transmission Electron
0 Microscope (TEM). A total line length of lmm has been
m;nPcl and the mean linear intercept grain size
~lPtPrm;nP-l to be 3.1 ~m.
.` 21~5216
Wo 95l25825 1 ~l, ,.,~.~C ~ -~
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Table 1 - Effect o~ solute additions on the as-rolled
strength of laboratory procesaed alloys rolled to give
the eguivalent strain as commercially rolled housefoil.
Alloy 0 . 29~ Proof Stressl UTS ~longation
(MPa ) (MPa ) ( % )
8006 159 210 6 . 1
8006 + 0.2~ Cu 217 268 5.9
8006 + 0.4~ Cu 249 309 3.3
8006 + 0.2~6 Mg 259 311 2.2
8006 + 0.4~6 Cu 274 331 2.4
~ ~ 2185216
WO 95/2~825 - 14 - r~
Table 2 - Te~8i.le properties O~ comnLercially pr.,.luc~d
14 ~n foil.
Alloy Cond:Ltion 0.2% Proo~ Stress ~TS 17lnr1~t;nn
(MPa) (MPa) (~)
AA8006 P.~-rolled 165 190 l.0
AA8006 Plant annealed 98 115 2.7
8006 + 0.2 Mg ~s-rolled 231 255 0.6
8006 + 0.2 Mg Plant annealed 102 123 l.9
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WO 9S/2S825 - 15 - r ~ l l~. ,,, _ - '~ E
Table 3 - 14 ~m Foil
r~ _ .1 Volume ~6 i 15~
As-rolled Annealed
I
Cube {001)<100> 2.2 2.8
Goss { 110 } ~001~ 3 . 2 2 . 4
Copper {112}<111~ 20.2) 25.3)
S {123}c634~ 35.2) 72.5 39.5) 79.2
Brass { 011 } ~211~ 17 .1) 14 . 4 )
Random 22 .1 15 . 6
