Technical Field
[0001] The present invention relates to forged aluminum alloys (forged aluminum alloy materials)
and methods for producing the forged aluminum alloys, where the forged aluminium alloys
are advantageously used as or for automobile suspension members.
BackgroundArt
[0002] Aluminum alloys such as 6xxx-series (Al-Mg-Si) alloys defined in Japanese Industrial
Standards (JIS) or Aluminum Association (AA) Standards have been used in structural
components of transportation equipment such as vehicles, ships, air craft, motor-bicycles,
and automobiles. The 6xxx-series aluminum alloys have relatively excellent corrosion
resistance and can be recycled satisfactorily, because scrap of these aluminum alloys
can be reused (recycled) as raw materials to be molten to give 6xxx-series aluminum
alloys.
[0003] Such aluminum alloys are often used in the form of aluminum alloy castings and forged
aluminum alloys in the structural components of the transportation equipment, from
the viewpoints of reduction in production cost, and of processability of materials
into parts having complicated shapes. Among them, forged aluminum alloys are mainly
used in strength members that require mechanical properties such as higher strength
and higher toughness. The strength members are exemplified by automobile suspension
members such as upper arms and lower arms. The forged aluminum alloys are produced
by homogenizing aluminum alloy castings, subjecting the homogenized castings to hot
forging typically via mechanical press forming or oil hydraulic press forming, and
performing tempering (heat treatments) such as solution heat treatment, quenching,
and artificial aging. The artificial aging is hereinafter also simply referred to
as "aging". The forging materials may also be selected from extrusions prepared by
homogenizing the castings and extruding the homogenized castings.
[0004] The strength members of the transportation equipment require still further reduction
in weight (reduction in thickness) with increasing requirements for lower fuel consumption
and lower carbon dioxide (CO
2) emission. However, forgings of 6061, 6151, and other 6xxx-series aluminum alloys
have been conventionally used in these applications, but inevitably offer insufficient
strength (0.2% yield strength) and/or insufficient toughness.
[0005] As possible solutions to these problems, various forged aluminum alloys have been
developed. For example, Patent Literature (PTL) 1 discloses an automobile suspension
part including a forged aluminum alloy. The forged aluminum alloy contains, in mass
percent, Mg in a content of 0.5% to 1.25%, Si in a oontent of 0.4% to 1.4%, Cu in
a content of 0.01% to 0.7%, Fe in a content of 0.05% to 0.4%, Mn in a content of 0.001%
to 1.0%, Cr in a content of 0.01% to 0.35%, and Ti in a content of 0.005% to 0.1%,
where Zr content is controlled to less than 0.15%, with the remainder consisting of
Al and inevitable impurities. In the microstructure of a transverse section at a maximum-stress-receiving
region, crystallized grains are present in a density of 1.5% or less in terms of average
area percentage, where the crystallized grains are observed in the microstructure
of a cross-sectional region receiving the maximum stress. Grain boundary precipitates
are present at an average spacing between the precipitates of 0.7 pm or more, where
the grain boundary precipitates are observed in a microstructure in a cross-sectional
region including a parting line formed upon forging.
[0006] The technique disclosed in PTL 1 is intended to allow automobile suspension parts
to have strength, toughness, and corrosion resistance all at higher levels, even when
the parts have shapes designed for weight reduction. With the technique, the forged
aluminum alloy is allowed to have an unrecrystallized structure by controlling the
chemical composition (amounts of transition elements to be added) and production conditions
. (mainly forging temperature and homogenization conditions).
[0007] In particular, various properties and microstructure (phases) in the maximum-stress-receiving
region are specified, because thickness reduction for the purpose of weight reduction
often causes recrystallization.
[0008] PTL 2 discloses a forged aluminum alloy for high-strength members. The forged aluminum
alloy contains Mg in a content of 0.6% to 1.8%, Si in a content of 0.8% to 1.8%, Cu
in a content of 0.2% to 1.0%, and at least one element selected from the group consisting
of Mn in a content of 0.1% to 0.6%, Cr in a oontent of 0.1% to 0.2%, and Zr in a content
of 0.1% to 0.2%, with the remainder consisting of Al and inevitable impurities, where
the mass ratio of Si to Mg is 1 or more. The forged aluminum alloy has a thickness
of 30 mm or less in the thinnest portion. After artificial aging, the forged aluminum
alloy has an electric conductivity of 41.0 to 42.5 percent IACS and has a 0.2% yield
strength of 350 MPa or more, where the electric conductivity is measured at the surface
of the forged aluminum alloy.
[0009] The technique disclosed in PTL 2 specifies that the electric conductivity of the
forged aluminum alloy surface after aging is from 41.0 to 42.5 percent IACS, so as
to allow the forged aluminum alloy to surely have a strength (0.2% yield strength)
of 350 MPa or more and to stably have the strength.
Citation List
Patent Literature
[0010]
PTL 1: Japanese Patent No. 5110938
PTL 2: Japanese Patent No. 3766357 The European patent application EP 2 003 219 A2 discloses an aluminium alloy forged member used for automotive underbody parts such
as suspension components. The disclosed 6XXX series alloy and its method of manufacture
has higher strength, toughness and higher corrosion resistance even when the overall
weight of the component is reduced.
Summary of Invention
Technical Problem
[0011] As described above, attempts have been made on conventional forged aluminum alloys
to have strength and toughness at higher levels. Increasing demands have been made
on such materials to have higher strength so as to reduce weights of suspension parts
from the viewpoint of providing better fuel efficiency of automobiles. However, the
materials, if designed to have still higher strength, may highly possibly cause significant
reduction in corrosion resistance and toughness. In particular, the materials, when
combined with Cu, Si, Mg, and other elements that contribute to precipitation strengthening,
may have significantly reduced corrosion resistance.
[0012] Under these circumstances, the present invention has an object to provide a forged
aluminium alloy suspension part and a method for its production, where the forged
aluminum alloy offers high strength and high toughness and still has excellent corrosion
resistance, even when having a smaller thickness.
Solution to Problem
[0013] After intensive investigations, the inventors of the present invention have made
the present invention particularly in consideration of followings. Specifically, a
forged aluminum alloy was designed to have higher strength by adding Si, Cu, and Mg
in larger amounts, where these elements contribute to age precipitation. However,
Si, Cu, and Mg in larger amounts cause the forged aluminum alloy to have lower toughness
and lower corrosion resistance To eliminate or minimize this, the amounts of Mn, Cr,
and Zr are controlled, and over-aging is performed under predetermined conditions.
In addition, the forged aluminum alloy is allowed to have still higher strength by
performing drying in a shorter time as compared with conventional techniques, where
the drying is performed after quenching.
[0014] Specifically, to achieve the object, the present invention provides, according to
a first embodiment, a forged aluminium alloy suspension part. The forged aluminum
alloy includes Mg in a content of 0.70 to 1.50 mass percent, Si in a content of 0.80
to 1.30 mass percent, Cu in a content of 0.30 to 0.90 mass percent, Fe in a content
of 0.10 to 0.40 mass percent, Ti in a content of 0.005 to 0.15 mass percent, and at
least one element selected from the group consisting ofMn in a content of 0.10 to
0.60 mass percent, Cr in a content of 0.10 to 0.45 mass percent, and Zr in a content
of 0.05 to 0.30 mass percent, with the remainder consisting of Al and inevitable impurities.
In a maximum-stress-receiving region, a Q phase has a major axis of 50 to 500 nm.
[0015] The forged aluminum alloy having the configuration has strength, toughness, and corrosion
resistance at higher levels, because the forged aluminum alloy includes Mg, Si, Cu,
Fe, and Ti in the predetermined contents and further includes at least one of Mn,
Cr, and Zr each in the predetermined content In addition, the forged aluminum alloy
has higher strength (tensile strength and 0.2% yield strength), because the Q phase
in the maximum-stress-receiving region has a major axis controlled within the specific
range.
[0016] Preferably, the forged aluminum alloy according to the present invention has an average
grain size of 50.0 pm or less in terms of minor axis in the maximum-stress-receiving
region and has an area percentage of recrystallized grains of 30.0% or less in a transverse
section including the maximum-stress-receiving region.
[0017] The forged aluminum alloy having the configuration has still higher strength, because
the forged aluminum alloy has an average grain size within the specific range in the
maximum-stress-receiving region. The forged aluminum alloy has still higher strength
and toughness, because the forged aluminum alloy has an area percentage of recrystallized
grains within the specific range in a transverse section including the maximum-stress-receiving
region.
[0018] The forged aluminium alloy suspension part according to the invention preferably
has undergone a surface treatment.
[0019] The forged aluminum alloy, when having undergone the surface treatment, has still
better corrosion resistance.
[0020] The forged aluminium alloy suspension part according to the present invention preferably
has undergone shot blasting.
[0021] The forged aluminum alloy, when having undergone shot blasting, has still better
stress corrosion cracking resistance.
[0022] The present invention also provides, according to a second embodiment, a method for
producing a forged aluminium alloy suspension part. The forged aluminum alloy includes
Mg in a content of 0.70 to 1.50 mass percent, Si in a content of 0.80 to 1.30 mass
percent, Cu in a content of 0.30 to 0.90 mass percent, Fe in a content of 0.10 to
0.40 mass percent, Ti in a content of 0.005 to 0.15 mass percent, and at least one
element selected from the group consisting ofMn in a content of 0.10 to 0.60 mass
percent, Cr in a content of 0.10 to 0.45 mass percent, and Zr in a content of 0.05
to 0.30 mass percent, with the remainder consisting of Al and inevitable impurities.
The method includes the steps of melting, casting, homogenization, forging, and tempering.
In the melting step, an aluminum alloy having the chemical composition is melted to
give a molten metal. In the casting step, the molten metal is cast to give an ingot
In the homogenization step, the ingot is subjected to homogenization. In the forging
step, the homogenized ingot is used as a forging material, is heated, and subjected
to hot forging. The tempering step is performed after the forging step and includes,
in the following sequence, solution heat treatment, quenching at 20°C to 70°C for
30 minutes or shorter, drying for one hour or shorter, and overaging at 180°C to 220°C
for 2 to 24 hours.
[0023] According to the procedure (method), the steps are performed under the predetermined
conditions to allow the forged aluminum alloy to have strength, toughness, and corrosion
resistance at higher levels. In particular, control of the conditions for quenching,
drying, and over-aging allows the Q phase in the maximum-stress-receiving region to
have a major axis within the predetermined range and allows the forged aluminum alloy
to have higher strength. The over-aging performed under the predetermined conditions
allows grain boundary precipitates to distribute at larger spacing, and this allows
the forged aluminum alloy to have better corrosion resistance and to still offer higher
strength and higher toughness.
Advantageous Effects of Invention
[0024] The forged aluminium alloy suspension part according to the present invention offers
high strength and high toughness and still has excellent corrosion resistance, even
when designed to have a smaller thickness. Accordingly, the forged aluminum alloy
can be used in wider uses in transportation equipment and has industrially significant
value.
[0025] The method according to the present invention for producing a forged aluminum alloy
enables production of a forged aluminum alloy that offers high strength and high toughness
and still has excellent corrosion resistance, even when the forged aluminum alloy
is designed to have a smaller thickness.
Brief Description of Drawings
[0026]
Fig. 1 is a schematic view of a plane of a forged aluminum alloy according to the
present invention as observed with a transmission electron microscope (TEM) at 500000-fold
magnification;
Fig. 2 is a schematic view of a surface or cut section of a forged aluminum alloy
and illustrates how to measure the average grain size;
Fig. 3 is a plan view of a forged aluminum alloy according to an embodiment of the
present invention;
Fig. 4 is a cross-sectional view taken along the line A-A of Fig. 3;
Fig. 5 is a schematic view of a microstructure of a sample subjected to long-time
drying after quenching;
Fig. 6 is a schematic view of a microstructure of a sample subjected to short-time
drying after quenching; and
Fig. 7 is a graph illustrating how the strength of a forged aluminum alloy varies
depending on the aging time in a sample subjected to long-time drying, and in a sample
subjected to short-time drying.
Description of Embodiments
Forged Aluminum Alloy
[0027] First, the forged aluminum alloy which is used to produce the suspension part according
to the invention will be described. The forged aluminum alloy is hereinafter also
referred to as an "Al alloy forging" as appropriate.
[0028] The Al alloy forging according to the present invention includes (is made of) an
aluminum alloy containing Mg, Si, Cu, Fe, and Ti in predetermined contents and further
containing at least one of Mn, Cr, and Zr in a predetermined content, with the remainder
consisting of Al and inevitable impurities. The Q phase in a maximum-stress-receiving
region is controlled to have a major axis of 50 to 500 nm.
[0029] The Al alloy forging preferably has an average grain size in the maximum-stress-receiving
region of 50.0 pm or less in terms of minor axis and an area percentage of recrystallized
grains in a transverse section including the maximum-stress-receiving region of 30.0%
or less.
[0030] The configurations will be illustrated below.
[0031] The chemical composition of the Al alloy forging according to the present invention
will be described. The Al alloy forging according to the present invention has a chemical
composition corresponding to an Al-Mg-Si (6xxx-series) alloy. The chemical composition
is specified so as to ensure high strength, high toughness, and high durability such
as stress corrosion cracking resistance, for use as structural components or parts
of transportation equipment such as automobiles and ships. The chemical composition
of the Al alloy forging according to the present invention acts as one of significant
factors in conditions specified relating typically to grains.
[0032] Accordingly, the chemical composition of the Al alloy forging according to the present
invention is specified to include Mg in a content of 0.70 to 1.50 mass percent, Si
in a content of 0.80 to 1.30 mass percent, Cu in a content of 0.30 to 0.90 mass percent,
Fe in a content of 0.10 to 0.40 mass percent, Ti in a content of 0.005 to 0.15 mass
percent, and at least one element selected from the group consisting of Mn in a content
of 0.10 to 0.60 mass percent, Cr in a content of 0.10 to 0.45 mass percent, and Zr
in a content of 0.05 to 0.30 mass percent, with the remainder consisting ofAl and
inevitable impurities.
[0033] Next, critical significance and preferred ranges of contents of the individual elements
in the chemical composition of the Al alloy forging according to the present invention
will be described.
Mg: 0.70 to 1.50 mass percent
[0034] Magnesium (Mg) precipitates, together with Si and Cu, as Mg
2Si (β° phase) and Q phase as a result of over-aging and is necessary for imparting
a high 0.2% yield strength to the Al alloy forging. Mg, if present in a content less
than 0.70 mass percent, causes age hardening to be insufficient and causes the Al
alloy forging to have a lower 0.2% yield strength. In addition, Mg in an excessively
low content may cause the forged aluminum alloy to have a smaller major axis of the
Q phase and to have elongation, toughness, and/or corrosion resistance at lower levels.
In contrast, Mg, if present in a content greater than 1.50 mass percent, may cause
the forged aluminum alloy to have an excessively high 0.2% yield strength and adversely
affect the ingot forgeability. In addition, Mg in a content greater than 1.50 mass
percent causes Mg
2Si crystallized grains, which do not contribute to higher 0.2% yield strength, to
increase upon casting. This causes the Al alloy forging to have toughness and corrosion
resistance both at lower levels and may cause the Al alloy forging to have lower elongation.
To eliminate or minimize these, the Mg content is controlled to be 0.70 to 1.50 mass
percent, and is preferably 0.80 to 1.20 mass percent.
Si: 0.80 to 1.30 mass percent
[0035] Silicon (Si) also precipitates, together with Mg and Cu, as Mg
2Si (β° phase) and Q phase as a result of over-aging and is necessary for imparting
a high 0.2% yield strength to the Al alloy forging. Si, if present in a content less
than 0.80 mass percent, causes the Al alloy forging to have a smaller size of the
Q phase in terms of major axis, causes age hardening to be insufficient, and causes
the Al alloy forging to have a lower 0.2% yield strength. This may also cause the
Al alloy forging to be insufficient in properties such as tensile strength, elongation,
toughness, and corrosion resistance. In contrast, Si, if present in a content greater
than 1.30 mass percent, causes coarse grains of Si alone to form and to precipitate
upon casting and during quenching after the solution heat treatment. In addition,
such Si in an excessively high content causes Mg
2Si and Al-Fe-Si-(Mn, Cr, Zr) crystallized grains at grain boundaries to fail to have
a smaller average grain size and to fail to be present at larger average spacing between
them. This causes the Al alloy forging to have corrosion resistance and toughness
at lower levels, as with Mg, and still causes the Al alloy forging to offer inferior
workability such as low elongation. To eliminate or minimize these, the Si content
is controlled to be 0.80 to 1.30 mass percent, and is preferably 0.90 to 1.10 mass
percent.
Cu: 0.30 to 0.90 mass percent
[0036] Copper (Cu) contributes to higher 0.2% yield strength via solid-solution strengthening
and effectively significantly promotes age hardening of the Al alloy forging upon
over-aging. Cu, if present in a content less than 0.30 mass percent, fails to offer
these effects sufficiently and causes the Al alloy forging to have lower 0.2% yield
strength. Such Cu in an insufficient amount may also cause the Q phase to precipitate
insufficiently and may cause the Al alloy forging to have lower tensile strength.
In contrast, Cu, if present in a content greater than 0.90 mass percent, causes the
Q phase to have a larger major axis and causes the microstructure of the Al alloy
forging to have significantly higher susceptibility (sensitivity) to stress corrosion
cracking and/or grain-boundary corrosion. This causes the Al alloy forging to have
lower corrosion resistance. In addition, such excessive Cu may cause the Al alloy
forging to have lower elongation and/or lower toughness. To eliminate or minimize
these, the Cu content is controlled to be 0.30 to 0.90 mass percent, and is preferably
0.40 to 0.70 mass percent
Fe: 0.10 to 0.40 mass percent
[0037] Iron (Fe) is added to the Al alloy forging for better productivity and for less recrystallization
upon casting. However, Fe forms Al
7Cu
2Fe, Al
12(Fe, Mn)
3Cu
2, (Fe, Mn)Al
6, and/or Al-Fe-Si(Mn, Cr, Zr) crystallized grains. These coarse precipitates act as
fracture origins and cause the Al alloy forging to be inferior in properties such
as toughness and fatigue properties. In particular, Fe, if present in a content greater
than 0.40 mass percent, causes Al-Fe-Si-(Mn, Cr, Zr) crystallized grains at grain
boundaries to have a larger average grain size and to be present at a smaller average
spacing between the crystallized grains. This causes the Al alloy forging to have
toughness and corrosion resistance at lower levels and may cause the Al alloy forging
to have lower elongation. In contrast, Fe, if present in a content less than 0.10
mass percent, may cause disadvantages such as cracking upon casting and/or an abnormal
microstructure. To eliminate or minimize these, the Fe content is controlled to be
0.10 to 0.40 mass percent, and is preferably 0.20 to 0.30 mass percent.
Ti: 0.005 to 0.15 mass percent
[0038] Titanium (Ti) is added for refinement of ingot grains and for better workability
upon extrusion, rolling, and forging. However, Ti, if present in a content less than
0.005 mass percent, may fail to offer effective grain refinement In contrast, Ti,
if present in a content greater than 0.15 mass percent, forms coarse precipitates,
causes the Al alloy forging to be inferior in the workability, and may cause the Al
alloy forging to have lower toughness. To eliminate or minimize these, the Ti content
is controlled to be 0.005 to 0.15 mass percent, and is preferably 0.01 to 0.10 mass
percent.
[0039] At least one element selected from the group consisting of Mn: 0.10 to 0.60 mass
percent, Cr: 0.10 to 0.45 mass percent, and Zr: 0.05 to 0.30 mass percent
[0040] These elements form dispersed particles (dispersoids) upon homogenization and subsequent
hot forging. The dispersed particles are Al-Mn, Al-Cr, and/or Al-Zr intermetallic
compounds formed by bonding of elements such as Fe, Mn, Cr, Zr, Si, and Al selectively
according to the contents thereof These intermetallic compounds are generically referred
to as (Fe, Mn, Cr, Zr)
3SiAl
12 compounds.
[0041] The dispersed particles effectively eliminate or minimize grain boundary migration
after recrystallization. This configuration eliminates or minimizes coarsening of
the average grain size of a parting line microstructure in the ST direction (short
transverse direction; thickness direction) during the forging step and, in addition,
allows fine grains and fine subgrains to be present overall the Al alloy forging according
to the present invention. The elements Mn, Cr, and Zr are expected to allow the Al
alloy forging to have higher 0.2% yield strength via solid-solution.
[0042] Manganese (Mn) forms dispersed particles of a size of about 1 pm upon homogenization
and is effective for restraining of recrystallization. However, this element increases
susceptibility to grain-boundary corrosion and is controlled in the content In addition,
Mn is readily combined with Fe to form fragile, coarse precipitates. Accordingly,
control of the content allows the Al alloy forging to have better toughness.
[0043] Mn, Cr, and Zr, if present in contents respectively of less than 0.10 mass percent,
less than 0.10 mass percent, and less than 0.05 mass percent, are not expected to
offer the effects and may cause the Al alloy forging to have tensile strength, elongation,
and toughness at lower levels. In contrast, Mn, Cr, and Zr, if present in contents
respectively of greater than 0.60 mass percent, greater than 0.45 mass percent, and
greater than 0.30 mass percent, readily form crystallized grains as coarse Al-Fe-Si-(Mn,
Cr, Zr) intermetallic compounds upon melting and casting. These coarse intermetallic
compounds act as fracture origins and cause the Al alloy forging to be insufficient
in at least one of tensile strength, elongation, 0.2% yield strength, toughness, and
corrosion resistance. To eliminate or minimize these, at least one of these elements
is to be contained, in contents within the ranges of 0.10 to 0.60 mass percent for
Mn, 0.10 to 0.45 mass percent for Cr, and 0.05 to 0.30 mass percent for Zr. The Mn,
Cr, and Zr content are preferably respectively 0.30 to 0.50 mass percent, 0.15 to
0.30 mass percent, and 0.05 to 0.15 mass percent.
Remainder consisting of Al and inevitable impurities
[0044] The remainder of the chemical composition of the Al alloy forging consists of Al
and inevitable impurities. Possible inevitable impurities are exemplified by Ni, Zn,
Be, V, and other elements. Each of these elements may be contained at such a level
as not to adversely affect the advantages or features of the present invention. Specifically,
the contents of elements as the inevitable impurities are preferably 0.05 mass percent
or less per each element, and the total content of them is preferably 0.15 mass percent
or less.
[0045] Boron (B) is an inevitable impurity, but contributes to refinement of ingot grains
and effectively allows the Al alloy forging to have better workability upon extrusion,
rolling, and forging, as with Ti. Boron, however, if present in a content greater
than 500 ppm, also forms coarse precipitates and causes the Al alloy forging to be
insufficient in the workability. Accordingly, the Al alloy forging may contain boron
in a content up to 500 ppm. However, boron, if present in a content less than 1 ppm,
fails to offer the advantageous effects. Thus, boron may be contained in a content
of 1 ppm or more.
[0046] Next, the requirement on the Q phase major axis in the Al alloy forging will be described.
Major axis of Q phase in maximum-stress-receiving region: 50 to 500 nm
[0047] As used herein the term "Q phase" refers to a Q phase or a Q' phase. The Q phase
is a precipitate including Al
5Cu
2Mg
8Si
6, is precipitated via aging, and contributes to higher strength. The Q phase is precipitated
slower as compared with the B phase and β° phase (Mg
2Si). Thus, aging, even when performed as over-aging in the Al alloy forging production
method, can restrain reduction in strength. The major axis of the Q phase is controlled
to be 50 nm or more in particular for higher strength and is controlled to be 500
nm or less for higher toughness and/or better corrosion resistance. The Q phase, if
having a size in terms of major axis out of the range, may cause the Al alloy forging
to be insufficient in one or more of strength, elongation, toughness, and corrosion
resistance. To eliminate or minimize these, the major axis of the Q phase is controlled
to be from 50 to 500 nm. The "site which receives the maximum stress" is also referred
to as a "maximum-stress-receiving region", as appropriate. The maximum-stress-receiving
region is a site as indicated in Fig. 3 and will be described later.
[0048] The Q phase measurement may be performed by the following method.
[0049] First, a sample is cut out from the maximum-stress-receiving region of the Al alloy
forging. Next, the sample is subjected to electropolishing using two different solutions
to give a thin-film sample for transmission electron microscope (TEM) observation.
The two solutions are a 1:9 mixture of perchloric acid and ethanol, and a 1:3 mixture
of nitric acid and methanoL The microstructure of the thin-film sample is observed
in five view fields, while electron beams are applied in the <001> direction with
respect to the matrix, and the (100) plane is observed, at an acceleration voltage
of the transmission electron microscope of 120 kV. The observation is performed at
500000-fold magnification. The major axes of the Q phases are measured based on the
observed microstructure, and the average of the major axes in all the Q phases in
the five views fields is calculated. Specifically, the average major axis is calculated
by summing up the lengths in major axes of the Q phases in the five view fields, and
dividing the total length by the number of the Q phases in the five view fields. The
average major axis thus determined is controlled to 50 to 500 nm herein. Fig. 1 is
a schematic view of a plane upon observation with the TEM at 500000-fold magnification.
The plane in Fig. 1 includes Q phases 30 and β phases 31. The Q phases 30 each have
a relatively long, black needle-like shape. The 6 phases 31 also have a needle-like
shape, but appear as coffee beans, because the β phases are precipitated in conformance
with the matrix and thereby distort the matrix The "major axis" of each Q phase 30
refers to the longitudinal dimension of the needle. For example, the image in the
schematic view is defined as one view field, and the average of the major axes of
the Q phases 30 in five view fields is defined as the "major axis (dimension) of Q
phase".
[0050] The measurement region of the major axis of Q phase is not limited, as long as being,
for example, a transverse section (cross section in the transverse direction) including
the maximum-stress-receiving region as described below.
[0051] The major axis of Q phase may be controlled by chemical composition and conditions
for the tempering step including quenching, drying, and over-aging.
[0052] Next, the conditions on the grains in the Al alloy forging will be described.
Average grain size in maximum-stress-receiving region: 50.0 pm or less in terms of
minor axis
[0053] The average grain size affects mechanical properties. The Al alloy forging, when
having an average grain size of 50.0 µm or less in terms of minor axis in the maximum-stress-receiving
region, has higher strength. Consequently, the average grain size in the maximum-stress-receiving
region is preferably 50.0 µm or less in terms of minor axis. The average grain size
is more preferably 45.0 pm or less, and furthermore preferably 40.0 pm or less, from
the viewpoint of offering still higher strength. The lower limit is not limited. Theoretically,
the smaller the average grain size in terms of minor axis is, the better, but the
average grain size is substantially limited to 5.0 µm or more in terms of lower limit.
The maximum-stxess-receiving region is the site indicated in Fig. 3 and will be described
later.
[0054] The average grain size can be calculated by a minor-axis section method. Specifically,
the average grain size can be determined by calculation in the following manner, as
illustrated in Fig. 2. The surface or cut section of the Al alloy forging is etched
with an appropriate etchant, and an image of the etched surface or cut section is
taken with an optical microscope at 50-fold magnification. A straight line is drawn
in a direction perpendicular to the major axes of grains, the number of grains on
the straight line is counted, and the straight line length is divided by the grain
number to give the average grain size.
[0055] The measurement region of the average grain size is not limited, as long as being,
for example, a transverse section including the after-mentioned maximum-stress-receiving
region.
[0056] The average grain size may be controlled by chemical composition, forging conditions,
and tempering step conditions.
Area percentage of recrystallized grains in the transverse section including the maximum-stress-receiving
region: 30.0% or less
[0057] In the Al alloy forging, the area percentage of recrystallized grains in the transverse
section including the maximum-stress-receiving region is preferably controlled to
be 30.0% or less. The Al alloy forging, when having an area percentage of recrystallized
grains of 30.0% or less, has higher strength and higher toughness. The lower limit
is not specified, but the smaller the area percentage is, the better. The term "transverse
section" refers to such a cross section as to have a minimum area.
[0058] The area percentage of recrystallized grains in the transverse section including
the maximum-stress-receiving region may be controlled by homogenization temperature,
forging start temperature (forging initiation temperature) and forging finish temperature
(forging end temperature) upon forging, and solution heat treatment conditions.
[0059] Next, the maximum-stress-receiving region of the Al alloy forging will be described,
while taking an automobile suspension part illustrated in Figs. 3 and 4 as an example.
[0060] Figs. 3 and 4 illustrate a representative shape of the automobile suspension part
as an Al alloy forging according to an embodiment of the present invention. Fig. 3
is a plan view of the automobile suspension part 1 and illustrates the shape of the
entire part and the specific site of an arm portion where the maximum stress occurs.
Fig. 4 is a cross-sectional view taken along the line A-A of Fig. 3 and is a cross-sectional
view in the transverse direction of the specific site of the arm portion receiving
the maximum stress.
[0061] As illustrated in Fig. 3, the automobile suspension part 1 includes the Al alloy
forging which has been forged into this shape. The automobile suspension part 1 has
an approximately triangular shape as a whole as illustrated in Fig. 3 and includes
joint portions 5a, 5b, and 5c, such as ball joints, at the apices of the triangle;
and arm portions 2a and 2b that couple the joint portions to each other. The arm portions
2a, and 2b include ribs in peripheries (both side edges) in the transverse direction.
The ribs extend in the longitudinal direction in each arm portion. The arm portion
2a includes ribs 3a and 3b, and the arm portion 2b includes ribs 3a and 3c. The arm
portions 2a and 2b each include a web at the center part in the transverse direction,
where the web extends in the longitudinal direction of the arm portion. The arm portion
2a includes a web 4a, and the arm portion 2b includes a web 4b.
[0062] The ribs 3a, 3b, and 3c each have a relatively small width and a relatively large
thickness in common in the automobile suspension part. In contrast, the webs 4a and
4b each have a smaller thickness of about 10 mm or less and a relatively larger width
in common in the automobile suspension part, as compared with the ribs 3a, 3b, and
3c. The arm portions 2a and 2b each have an approximately H-shaped cross section as
the cross section in the transverse direction, in common in the automobile suspension
part. In the approximately H-shaped cross section, the both vertical wall portions
correspond to the ribs 3a, 3b, and 3c, and the central lateral wall portion corresponds
to the webs 4a and 4b.
[0063] A general automobile suspension part may be designed in structure of the arm portions
2a and 2b and the ball joint portions 5a, 5b, and 5c so that a specific region receiving
the maximum stress during usage (maximum-stress-receiving region) is located in the
rib near to the ball joint portion, based on the premise that the automobile suspension
part has the entire structure and shape as described above. The maximum-stress-receiving
region may naturally vary depending on the structure design conditions, but is often
located in any of the ribs.
[0064] In the automobile suspension part illustrated in Fig. 3, the specific site where
the maximum stress occurs during use (the maximum-stress-receiving region) is a shaded
region that is located in the rib near to the ball joint portion and extends in the
longitudinal direction, as indicated with slanted lines in Fig. 3. Specifically, the
specific site in the embodiment illustrated in Fig. 3 is a site that is indicated
with slanted lines, is located in one side of the arm portion 2a near to the ball
joint portion 5a, and includes part of the rib 3a and part of the web 4a. In addition,
the maximum-stress-receiving region in the area of the arm portion is not uniform
in the transverse section and corresponds to a region 6a in the top end of the rib
3a, as circled in Fig. 4. Assume that the specific maximum-stress-receiving region
in use extends not only to the rib 3a, but also to the rib 3b. In this case, a region
6b in the top end of the rib 3b also acts as a maximum-stress-receiving region in
use, where the region 6b is circled in Fig. 4.
[0065] In the automobile suspension part, other portions such as the joint portions 5a,
5b, and 5c with other members also naturally receive large stress, but the stress
is not the maximum stress. The maximum stress in the automobile suspension part occurs
in a region of a specific rib near to a ball joint portion, as illustrated in Fig.
3, where the specific arm portion is determined by the entire shape and dimensional
conditions of the arm portion. The maximum -stress-receiving region, however, may
vary depending typically on the shape and manufacturer's demand characteristics of
the automobile suspension part. However, wherever area the maximum-stress-receiving
region is located, the average grain size may be specified in the maximum-stress-receiving
region, and the area percentage of recrystallized grains may be specified in a transverse
section including the maximum-stress-receiving region.
[0066] Assume that, out of the maximum-stress-receiving region in the arm portion, the rib,
or the web including the rib becomes susceptible to grain coarsening, where the rib
or the web is to have strength at certain level In this case, it is difficult to reduce
the weight while maintaining the strength of the arm portion and the strength of the
entire automobile suspension part at high levels.
[0067] Thus, the area percentage of recrystallized grains in a transverse section including
a specific site of the arm portion is specified in the present invention, where the
specific site receives the maximum stress. The specific site is indicated with slanted
lines in Fig. 3 and is a site that is present in one side of the arm portion 2a near
to the ball joint portion 5a and includes part of the rib 3a and part of the web 4a.
When producible, the microstructure is preferably controlled as mentioned above not
only in the maximum-stress-receiving specific site of the arm portion, but also in
the entire arm portions 2a and 2b.
[0068] In a preferred embodiment of the present invention, the area ratio of recrystallized
grains is controlled in two sites including the parting line (PL region), out of the
transverse section microstructure in the arm portion 2a receiving the maximum load
The parting line is most susceptible to recrystallization, as described above. The
"area ratio of recrystallized grains" is also referred to as a "recrystallization
area ratio". With reference to Fig. 4, the two sites are a site including the entire
microstructure in the transverse section of the rib 3a; and a site including the entire
microstructure in the transverse section of the adjacent web 4a. The recrystallization
area ratio of the arm portion including both the ribs and the web is preferably controlled
in the above manner.
[0069] The web 4a is also susceptible to recrystallization, as with the PL region. The size
(recrystallization area ratio) of grains in the web also significantly affects the
strength. Since the web is forged to a working ratio different from that of the rib,
the web may highly possibly have a recrystallization area ratio different from that
of the rib. Accordingly, the recrystallization area ratio of the arm portion receiving
the maximum stress is preferably specified (controlled) both in the web and the rib.
[0070] This configuration restrains recrystallization in the arm portion receiving the maximum
stress (in particular, the rib and the web), increases subgrains, allows grains to
be refined to about 50.0 µm or less, and restrains grain boundary failure in the arm
portion. Thus, the automobile suspension part is allowed to have higher strength and
higher toughness.
[0071] The specific sites of the rib are specified (measured) as regions receiving the maximum
stress in the transverse section, out of the entire microstructure in the transverse
section of the rib 3a in Fig. 4. Specifically, the specifying (measurement) of the
specific sites of the rib is performed at two sites, i.e., a site 7 and a site 8 in
Fig. 4. The site 7 includes the region 6a in the top end of the rib 3a, where the
region 6a is circled in Fig. 4. The site 8 includes the parting line (PL region),
where the PL region is most susceptible to recrystallization, as mentioned above.
Namely, it is preferred that microstructures in the measurement sites 7 and 8 are
treated as representative microstructures of the entire microstructure in the transverse
section of the rib, and the area ratios of recrystallized grains in the two sites
are controlled to be 30.0% or less in terms of average area percentage. This configuration
is preferred to increase subgrains and to allow grains to be refined to an average
grain size of about 50.0 µm or less. This restrains grain boundary failure of the
rib and allows the automobile suspension part to have higher strength and higher toughness.
[0072] The specific site in the web is specified (measured) as a site 9 including the parting
line (PL region), where the PL region is most susceptible to recrystallization, out
of the entire microstructure in the transverse section of the web 4a in Fig. 4. Specifically,
it is preferred that a microstructure in the measurement site 9 is treated as a representative
microstructure of the entire microstructure in the transverse section of the web,
and the area ratio of recrystallized grains in this site is controlled to be 30.0%
or less in terms of average area percentage. This configuration is preferred to increase
subgrains and to allow grains to be refined to an average grain size of about 50.0
µm or less. This restrains grain boundary failure of the web and allows the automobile
suspension part to have higher strength and higher toughness.
Recrystallization Area Ratio Measurement
[0073] The area ratio of recrystallized grains can be measured in the following manner.
First, each of samples (cross-section microstructures) at measurement sites of the
rib and the web are each mechanically polished by 0.05 to 0.1 mm and etched with cupric
chloride. Images of the specific sites (measurement sites) are taken typically with
a digital camera, subjected to image processing, and the ratio of the recrystallized
grain area to the observed view-field area is calculated. The recrystallized grains
have large sizes, readily reflect light, and appear light-colored. In contrast, other
grains including subgrains have small sizes and appear dense-colored. The recrystallized
grains and the other grains can be distinguished from each other by the differences
in size and in color density as mentioned above, and this enables image processing.
[0074] The conditions specified for the microstructure as mentioned above contribute to
higher strength and higher toughness particularly of the rib and web of the arm portion,
where the rib and web include the maximum-stress-receiving region. In short, the microstructural
conditions contribute to higher strength and higher toughness of the region receiving
the maximum stress in the arm portion. Thus, the microstructural conditions allow
an automobile suspension part to have higher strength, higher toughness, and better
corrosion resistance, even when the automobile suspension part includes an arm portion
having an approximately H-shaped cross section. The H-shaped cross section includes
a web and ribs, where the web has a small thickness of about 10 mm or less and a relatively
large width and is present in the central part of the arm portion, and the ribs have
a smaller width and a larger thickness as compared with the web and are present in
peripheries of the arm portion. Namely, the microstructural conditions allow an automobile
suspension part to have higher strength, higher toughness, and better corrosion resistance,
even when the suspension part has a shape designed for weight reduction.
[0075] The maximum-stress-receiving region has been illustrated herein while taking the
automobile suspension part (Al alloy forging) having the shape illustrated in Figs.
3 and 4. The conditions for the maximum-stress-receiving region, however, will be
applied also to automobile suspension parts having other shapes.
Cross-Section Shapes Other than H-shape
[0076] The microstructural conditions specified in the present invention for automobile
suspension parts may be applied also to any other cross-section shapes than the H-shaped
cross-section shape including ribs and a web. For example, the microstructural conditions
specified in the present invention may be applied to a microstructure in a maximum-stress-receiving
region, out of the microstructure of a transverse (crosswise) section. Specifically,
control of the area percentage of recrystallized grains observed in the microstructure
in a transverse section including a maximum-stress-receiving region to 30.0% or less
allows the maximum-stress-receiving region to have higher strength and higher toughness
in the cross section.
[0077] The aluminium alloy forging is used to produce an automobile suspension member.
[0078] Examples of the automobile suspension members include, but are not limited to, upper
arms and lower arms. A non-limiting example in shape of final products is an automobile
suspension part having the shape as illustrated in Figs. 3 and 4. The automobile suspension
part includes an arm portion that includes ribs and a web and has an approximately
H-shaped cross section. The ribs are located in peripheries and have a relatively
small width and a large thickness. The web are located in a central portion and has
a small thickness of about 10 mm or less and a relatively large width.
[0079] The aluminium alloy suspension member according to the present invention may have
undergone a
surface treatment. The Al alloy forging, when having undergone the surface treatment,
has still better corrosion resistance. The surface treatment will be described in
the after-mentioned surface treatment step.
[0080] The aluminium alloy forging preferably has a hydrogen gas concentration controlled
within the range as follows.
Hydrogen: 0.25 ml or less per 100 g of Al
[0081] Hydrogen (H
2) often causes forging defects such as bubbles, thereby acts as fracture origin, and
causes the Al alloy forging to readily have lower toughness and lower fatigue properties,
particularly when the Al alloy forging is processed at a low working ratio.
[0082] Hydrogen has significant effects particularly on high-strength structural components
of transportation equipment and other parts or members. To eliminate or minimize this,
the hydrogen content is preferably minimized and controlled to 0.25 ml or less per
100 g of Al.
Forged Aluminum Alloy Production Method
[0083] Next, the method according to the present invention for producing a forged aluminium
alloy suspension part will be illustrated. The production method according to the
present invention is a method for producing the above-mentioned forged aluminum alloy
and includes the steps of melting, casting, homogenization, forging, and tempering.
The production method may further include a surface treatment step and/or a degassing
step, as needed.
Melting Step
[0084] The melting step is the step of melting an Al alloy having the chemical composition
into a molten metal
Casting Step
[0085] The casting step is the step of casting the molten metal to give an ingot, where
the molten metal is prepared by melting (in the melting step) so as to have the chemical
composition. The casting is performed by a common melting casting technique selected
as appropriate. The melting-casting technique is exemplified by, but is not limited
to, continuous casting-directed rolling, semicontinuous casting (direct chill casting
(DC casting)), and hot top casting (hot top direct chill casting). The ingot is not
limited in shape and may be in any form such as round bars and other ingots, and slabs.
[0086] In the casting step, the cooling of the molten metal to give the ingot is preferably
performed at a cooling rate of 10°C/sec or more. The cooling, when performed at a
cooling rate in this range, allows Al-Fe-Si-(Mn, Cr, Zr) crystallized grains present
at grain boundaries to have a smaller average grain size and to be present at larger
average spacing between them. This allows the Al alloy forging to have strength, toughness,
and corrosion resistance at still higher levels. The "cooling rate" of the molten
metal in this step is defined as an average cooling rate from the liquidus temperature
to the solidus temperature.
Homogenization Step
[0087] The homogenization step is the step of subjecting the ingot to homogenization. The
homogenization of the ingot in the homogenization step is preferably performed at
a holding temperature of 400°C to 560°C.
[0088] The homogenization, when performed at a holding temperature of 560°C or lower, causes
the (Fe, Mn, Cr, Zr)
3SiAl
12 dispersed particles themselves to resist coarsening and to less increase in number.
This allows a relatively large number of fine dispersed particles dispersed in grains,
readily gives an unrecrystallized structure, and consequently allows the Al alloy
forging to have strength, toughness, and corrosion resistance at higher levels.
[0089] In contrast, the homogenization, when performed at a holding temperature of 400°C
or higher, often allows the (Fe, Mn, Cr, Zr)
3SiAl
12 dispersed particles to maintain their sizes at such a level as to restrain the recrystallization
and to be present in a larger number, where the dispersed particles themselves contribute
to restrained recrystallization. In addition, the homogenization allows the Al-Fe-Si-(Mn,
Cr, Zr) crystallized grains to be dissolved as solutes sufficiently and allows Mg
2Si and Al-Fe-Si-(Mn, Cr, Zr) crystallized grains to readily have a smaller average
grain size and to readily have a larger average spacing between them, where the crystallized
grains are present at grain boundaries in the microstructure of the Al alloy forging
after the tempering step mentioned below. This allows the Al alloy forging to have
strength, toughness, and corrosion resistance at higher levels. For these reasons,
the holding temperature is preferably controlled within the range of 400°C to 560°C.
[0090] For stable precipitation of dispersed particles, the holding time at the holding
temperature is preferably 3 hours or longer. The homogenization may employ a furnace
as appropriately selected typically from air furnaces, induction heating furnaces,
and salt-bath furnaces. The "rate of temperature rise" of the ingot in this process
refers to an average rate of temperature rise from room temperature up to the holding
temperature.
Forging Step
[0091] The forging step is the step of heating the ingot after homogenization as a forging
material and subjecting the heated ingot to hot forging.
[0092] In the forging step, the forging material in the form typically of an ingot or extruded
bar is hot-forged typically via mechanical press forming or oil hydraulic press forming.
The hot forging of the forging material is preferably started at a temperature of
500°C or higher. The hot forging, when performed at a start temperature of 500°C or
higher, allows subgrain phases to occupy a larger proportion of the as-forged structure,
thereby increases grain boundaries in the as-forged structure, and accelerates precipitation
of Mg
2Si. This allows the Al alloy forging to have strength, toughness, and corrosion resistance
at higher levels. For these reasons, the start temperature is preferably controlled
to be 500°C or higher. The start temperature is more preferably 520°C or higher from
the viewpoint of restrainment of recrystallization.
[0093] The hot forging of the forging material is preferably finished at a temperature of
400°C or higher. Such high-temperature plastic working in the present invention accelerates
dynamic recovery and lowers the dislocation density after working. This restrains
grain coarsening caused by recrystallization, allows the Al alloy forging to have
an unrecrystallized structure as the microstructure, and allows the Al alloy forging
to have strength, toughness, and corrosion resistance at higher levels. The forging,
when performed at an end temperature of 400°C or higher, accelerates dynamic recovery
and allows the Al alloy forging to have strength, toughness, and corrosion resistance
at higher levels. For these reasons, the forging finish temperature is preferably
controlled to be 400°C or higher, and is more preferably 420°C or higher.
[0094] The forging material may be one prepared by subjecting the homogenized ingot to working
such as extrusion and/or rolling, so as to eliminate or minimize an as-cast structure
remained in the Al alloy forging and to allow the Al alloy forging to have strength
and toughness at still higher levels.
[0095] The control of the forging finish temperature in the hot forging of the forging material
to preferably 400°C or higher, and more preferably 420°C or higher, may require a
technique such as reheating before hot forging or use of a tool with which the work
can be held at high temperatures.
[0096] The hot forging is preferably performed via mechanical forging, and the number of
forging procedures is preferably 3 or less. The Al alloy forging may have any shape
not limited, such as a near net shape near to the final product shape. A non-limiting
example of the final product shape is the shape of the automobile suspension part
as illustrated in Fig. 3. The Al alloy forging after the forging may be subjected
to trimming so as to remove unnecessary portions.
Tempering Step
[0097] The tempering step is performed after the forging step and is the step of performing
solution heat treatment, quenching, drying, and over-aging. The tempering step is
performed so as to offer strength, toughness, and corrosion resistance necessary as
an Al alloy forging. In general, non-limiting examples of the tempering step include
T6, T7, and T8 tempers. The T6 temper sequentially includes solution heat treatment,
quenching, and artificial aging to give maximum strength. The T7 temper sequentially
includes solution heat treatment, quenching, and over-aging under conditions exceeding
the conditions of artificial aging to give maximum strength. The T8 temper sequentially
includes solution heat treatment, quenching, cold working, and artificial aging to
give maximum strength.
[0098] The tempering step in the present invention is performed after the forging step and
includes solution heat treatment; quenching at 20°C to 70°C for 30 minutes or shorter;
drying for one hour or shorter; and over-aging at 180°C to 220°C for 2 to 24 hours
performed in this sequence.
[0099] The solution heat treatment is preferably performed at a holding temperature of from
500°C to 580°C. The solution heat treatment, when performed at a holding temperature
of 500°C or higher, accelerates solutionization, increases solute Mg
2Si, and allows the Al alloy forging to have higher 0.2% yield strength. In contrast,
the solution heat treatment, when performed at a holding temperature of 580°C or lower,
less causes local melting and grain coarsening, and allows the Al alloy forging to
have higher 0.2% yield strength. For these reasons, the holding temperature is preferably
controlled to be from 500°C to 580°C.
[0100] The solution heat treatment is preferably performed for a holding time of 20 minutes
to 20 hours at a rate of temperature rise of 100°C/ hour or more, so as to ensure
0.2% yield strength. The "rate of temperature rise" of the Al alloy forging herein
refers to the average rate of temperature rise from the solution heat treatment start
temperature up to the holding temperature.
[0101] The quenching is performed at a temperature of 20°C to 70°C. The quenching, if performed
at a temperature lower than 20°C, excessively rapidly cools the work to cause a larger
difference in temperature between the inside and the outside of the forging to thereby
cause strain. In contrast, the quenching, if performed at a temperature higher than
70°C, cools the work at an excessively low cooling rate and causes coarse precipitates
to form during cooling, where the coarse precipitates do not contribute to higher
strength. This causes the Al alloy forging to fail to have sufficient 0.2% yield strength.
In addition, the quenching temperature affects the Q phase. For these reasons, the
quenching temperature is controlled to be 20°C to 70°C, and is preferably 30°C to
60°C.
[0102] The quenching is performed for a time of 30 minutes or shorter. The quenching, if
performed for a time longer than 30 minutes, may cause precipitation to initiate during
quenching and causes the Al alloy forging to fail to have sufficient 0.2% yield strength.
In addition, the quenching time affects the Q phase. For these reasons, the quenching
time is controlled to be 30 minutes or shorter. The lower limit of the quenching time
may vary depending on the size and weight of the Al alloy forging and may be set so
as to give a time necessary for advantageous effects of the quenching.
[0103] The quenching may be performed via cooling by immersing the work in water or warm
water, or by showering the work with water or warm water. The cooling (quenching)
is preferably performed at a cooling rate of 40°C/sec or more so as to eliminate or
minimize reduction in toughness and fatigue properties. The solution heat treatment
may employ equipment selected typically from air furnaces, induction heating furnaces,
salt-bath furnaces as appropriate.
[0104] If moisture remains as attached to the surface of the Al alloy forging, hydrogen
atoms migrate from the moisture into the Al alloy forging. The migrated hydrogen atoms
form hydrogen molecules upon aging and expand to cause surface defects called "blister".
To eliminate or minimize this, the Al alloy forging after quenching is sufficiently
dried before aging. However, the drying is performed for a time of one hour or shorter.
The drying, if performed for a time longer than one hour, causes formation of precipitates
(cluster I) due to natural aging, as mentioned below, where the precipitates (cluster
I) do not contribute to higher strength.
[0105] Next, drying will be specifically described. Fig. 5 is a schematic view of a microstructure
when drying after quenching is performed for a long time. Fig. 6 is a schematic view
of a microstructure when drying after quenching is performed for a short time. Fig.
7 is a graph illustrating how the strength of the Al alloy forging varies depending
on the aging time in a sample subjected to long-time drying and a sample subjected
to short-time drying.
[0106] Drying causes formation of a cluster I and a cluster II in grains in the microstructure
after quenching. The cluster I is a cluster of precipitates that do not contribute
to higher strength. The cluster II is a cluster of precipitates that contribute to
higher strength. However, drying performed for a short time does not cause the formation
of the cluster II in some cases, as mentioned below. The cluster I and the cluster
II are formed by aggregation of Si, Mg, and Cu.
[0107] As illustrated in Fig. 5, drying after quenching performed for a long time (longer
than one hour) causes formation of a large number of cluster I (21) and a small number
of cluster II (22) in grains 10 after drying.
[0108] In early stages of aging, the cluster I (21) gradually disappears due to reversion.
In contrast, the cluster II (22) acts as precipitates, is converted via a Guinier-Preston
zone (GP-zone) 25 into a Q phase 30 (or into a β phase 31 in some cases), where the
GP-zone 25 is a precursor of the Q phase 30 and the β phase 31. During this process,
new GP-zones 25 are also formed. The new GP-zones 25 are derived from clusters II
(22) newly formed during the aging.
[0109] At the completion of the aging, the GP-zones 25 are converted into the Q phases 30
and/or the 6 phases 31; and the clusters I (21) disappear due to reversion.
[0110] As illustrated in Fig. 7, the strength in the sample subjected to long-time drying
(sample L) reaches its peak slower as compared with the sample subjected to short-time
drying (sample S).
[0111] In the long-time drying, a large number of the clusters I (21) form, and the strength
reaches its peak slower by the time for the clusters I (21) to be reversed. Accordingly,
the Q phases 30 and the β phases 31 formed from the clusters II (22) undergo excessive
over-aging, and the Al alloy forging subjected to long-time drying (sample L) has
a lower peak strength as compared with the Al alloy forging subjected to short-time
drying (sample S), as illustrated in Fig. 7.
[0112] In contrast, short-time drying (for one hour or shorter) after quenching causes formation
of a small number of the cluster I (21) in the grains 10 after drying, as illustrated
in Fig. 6. In this process, few or approximately no cluster II (22) is formed. In
early stages of the aging, the cluster I (21) gradually disappears due to reversion,
but GP-zones 25 are formed The GP-zones 25 are derived from clusters II (22) newly
formed during the aging.
[0113] The sample subjected to short-time drying includes a smaller amount of the cluster
I(21) formed during the drying and includes a larger amount of clusters II (22) newly
formed during the aging, as compared with the sample subjected to long-time drying.
Accordingly, the sample subjected to short-time drying includes a larger amount of
GP-zones 25 formed in the early stages of the aging, as compared with the sample subjected
to long-time drying.
[0114] At the completion of the aging, the GP-zones 25 are converted into the Q phases 30
and/or the B phases 31; and the cluster I(21) disappears due to reversion.
[0115] In the sample subjected to short-time drying, precipitates that contribute to higher
strength are not formed even when the aging further proceeds. Thus, the strength reaches
its peak earlier in the sample subjected to short-time drying (sample S) as compared
with the sample subjected to long-time drying (sample L), as illustrated in Fig. 7.
[0116] In addition, in the sample subjected to short-time drying, a very small amount of
the cluster I (21), which does not contribute to higher strength, is formed, but larger
amounts of Si, Mg, and Cu aggregate to form the clusters II (22) which contribute
to higher strength. Accordingly, the sample Al alloy forging (sample S) subjected
to short-time drying has higher strength as compared with the sample Al alloy forging
(sample L) subjected to long-time drying.
[0117] As described above, drying, when performed for a time of one hour or shorter, allows
the A1 alloy forging to have higher strength. The short-time drying also allows the
strength to reach its peak earlier, and this allows the aging to be performed for
a shorter time and allows the Al alloy forging to offer better productivity. For these
reasons, the drying time is controlled to be one hour or shorter, and is preferably
0.5 hour or shorter. The drying may be performed according to a known common procedure,
as long as the surface can be sufficiently dried. The surface of the Al alloy forging
has only to be dried such that no moisture remains. For example, the surface may be
dried typically via forced-drying for 10 seconds using a fan.
[0118] Next, aging will be described.
[0119] The corrosion resistance and stress corrosion cracking resistance of the Al alloy
forging have a significant relationship with grain boundary precipitates. The age
hardening, if performed as subaging or peak aging, causes grain boundary precipitates
to precipitate as fine precipitates in a high density, and the grain boundary precipitates
act as origins of corrosion and often cause continuous corrosion. In contrast, the
age hardening, when performed as over-aging, causes the grain boundary precipitates
to coarsen. Such age hardening performed as over-aging allows the grain boundary precipitates
to be present at wider spacing between each other, and this retards the proceeding
of corrosion after corrosion proceeds to a certain extent. Thus, the over-aging allows
the Al alloy forging to resist corrosion.
[0120] The temperature and time of the averaging significantly affect the Q phase after
over-aging. This requires selection of such conditions as to give a required 0.2%
yield strength and to give other required properties such as toughness, elongation,
and corrosion resistance, where the selection is done in consideration of the production
history prior to this process. In this regard, the over-aging conditions may be selected
within the range of 180°C to 220°C for 2 to 24 hours for strength, toughness, and
corrosion resistance at higher levels, where the conditions may vary also depending
on the alloy element contents and the production history (conditions) prior to the
over-aging, and actual conditions should be checked on each production process and
production equipment.
[0121] The over-aging, if performed at a temperature lower than 180°C and/or for a time
shorter than 2 hours, is performed insufficiently and causes the Al alloy forging
to have strength and corrosion resistance at lower levels. In contrast, the over-aging,
if performed at a temperature higher than 220°C and/or for a time longer than 24 hours,
causes the Al alloy forging to have strength, elongation, and toughness at lower levels.
To eliminate or minimize these, the overaging is performed at a temperature of 180°C
to 220°C for 2 to 24 hours. The over-aging is preferably performed at a temperature
of 180°C to 200°C for a time of 4 to 12 hours.
[0122] The over-aging may be performed using equipment selected typically from air furnaces,
induction heating furnaces, and oil baths as appropriate.
Surface Treatment Step
[0123] The surface treatment step is the step of performing a surface treatment on the Al
alloy forging after the tempering step.
[0124] Exemplary techniques of the surface treatment include surface treatments typically
via cathodic electrodeposition, or surface coating (e.g., GEOMET® treatment and powder
coating). The cathodic electrodeposition and surface coating may be performed according
to common, known procedures without limitation. The surface treatment allows the Al
alloy forging to have still better corrosion resistance.
[0125] The surface treatment may also be shot blasting. The shot blasting may be performed
according to a common, known procedure without limitation.
[0126] The shot blasting applies compressive residual stress to the Al alloy forging surface
and reduces tensile stress, where the tensile stress causes stress corrosion cracking.
[0127] An Al alloy forging having a relatively high Cu content within the chemical composition
range specified in the present invention did not undergo stress corrosion cracking
in a stress corrosion cracking test for 90 days. However, this Al alloy forging, when
subjected to a stress corrosion cracking test for a longer period of 180 days to obtain
higher reliability, suffered from stress corrosion cracking. In contrast, the Al alloy
forging having the chemical composition, when subjected to shot blasting using glass
beads at 0.3 MPa for 90 seconds, did not undergo stress corrosion cracking even after
180-days testing.
[0128] The production method according to the present invention preferably further includes
a degassing step between the melting step and the casting step.
Degassing Step
[0129] The degassing step is the step of removing hydrogen gas from the molten metal melted
in the melting step so as to control the hydrogen gas concentration to 0.25 ml or
less per 100 g of the aluminum alloy. Namely, the degassing step is the step of performing
a degassing treatment. The hydrogen gas removal may be performed in a holding furnace
by fluxing, chlorine refining, or, in-line refining of the molten metal, where the
holding furnace is used for control of the molten metal chemical composition and for
removal of inclusions. The hydrogen gas is preferably removed by blowing an inert
gas such as argon gas into the molten metal using spinning nozzle inert flotation
(SNIF) or a porous plug (see Japanese Unexamined Patent Application Publication No.
2002-146447) in an apparatus for removing the hydrogen gas.
[0130] The determination of the hydrogen gas concentration is performed by measuring the
hydrogen gas concentration of the ingot produced in the casting step, or of the forging
produced in the forging step. The hydrogen gas concentration of the ingot can be determined
typically by cutting a sample out of the ingot prior to the homogenization, ultrasonically
cleaning the sample with an alcohol and acetone, and measuring the hydrogen gas concentration
typically by the inert gas fusion thermal conductivity detection (LIS A06-1993). The
hydrogen gas concentration of the forging can be determined typically by cutting a
sample out of the forging, immersing the sample in a NaOH solution, removing an oxide
coating from the surface of the sample with nitric acid, ultrasonically cleaning the
sample with an alcohol and acetone, and measuring the hydrogen gas concentration typically
via the volumetric method using hot vacuum extraction (LIS A06-1993).
[0131] The production method according to the present invention may further include a preforming
step typically using a forging roll, prior to the forging step. Examples
[0132] Next, the present invention will be illustrated in further detail with reference
to examples (experimental examples).
First Experimental Example
[0133] Al alloys having chemical compositions given in Table 1 were cast via hot-top casting
at a cooling rate of 20°C/sec into ingots in the form of round bars having a diameter
of 68 mm and a length of 580 mm. The ingots were subjected to homogenization at a
rate of temperature rise of 5°C/min and a holding temperature of 550°C for a holding
time of 4 hours.
[0134] Further, hot forging was performed three times so that the total forging working
ratio reached 75%. The hot forging was performed via mechanical forging using upper
and lower dies (tools) at a forging start temperature of 520°C and a forging finish
temperature of 420°C. The hot forging yielded Al alloy forgings having the shape of
an automobile suspension member as illustrated in Figs. 3 and 4. The forgings each
had a thickness in the thinnest portion of 6 mm.
[0135] Next, the Al alloy forgings were subjected to solution heat treatment at 550°C for
4 hours in an air furnace, followed by water cooling (water quenching) at 40°C for
15 minutes, and drying for 10 minutes to dryness. Subsequently, the Al alloy forgings
were subjected to overaging at 190°C for 5 hours in an air furnace.
[0136] Three test specimens were sampled from a maximum-stress-receiving region of each
Al alloy forging and subjected to measurements on average grain size in terms of minor
axis, area percentage, and major axis of Q phase as microstructural forms, as presented
in Table 2. The test specimens were also subjected to determinations of Al alloy forging
properties including tensile properties such as tensile strength, 0.2% yield strength,
and elongation, and Charpy impact value (mechanical property). The tensile properties
are indices of strength, and the Charpy impact value is an index of toughness. Each
of data in Table 2 is an average of values measured on the three test specimens sampled
from each Al alloy forging. The test specimens were sampled from the site 7 in Fig.
4 and had a parallel portion diameter of 7 mm, a gauge length of 25 mm, and a total
length of 90 mm, which size is a subsize of the No. 4 test specimen prescribed in
JIS Z 2241:2011.
[0137] The average grain size (in micrometer) was measured in the site 7 illustrated in
Fig. 4. The average grain size (in micrometer) was determined by etching the cut section
of each Al alloy forging with Barker's solution, taking an image of the etched cut
section at 50-fold magnification via polariscopy using an optical microscope, drawing
a straight line in a direction perpendicular to the major axes of grains, counting
the number of grains on the straight line, and dividing the length of the straight
line by the measured number of grains (see Fig. 2). Each measurement site was defined
herein typically as a region of the Al alloy forging excluding a region of 2 mm deep
from the surface, so as to exclude the recrystallized layer in the surface of the
Al alloy forging.
[0138] The area ratio of recrystallized grains was measured in the following manner. For
the rib, the measurement was performed at two sites, ie., the sites 7 and 8 illustrated
in Fig. 4. For the web, the measurement was performed at the site 9 illustrated in
Fig. 4.
[0139] The samples from the (cross-section microstructure) observation sites in the rib
and the web were mechanically polished by 0.05 to 0.1 mm and etched with cupric chloride.
Images of the specified sites were taken with a digital camera, subjected to image
processing, and the ratio of the area of recrystallized grains to the area of observation
view field was calculated and defined as the area ratio of recrystallized grains.
[0140] The major axis of Q phase was measured at the site 7 illustrated in Fig. 4. The major
axis of Q phase was measured in the following manner. First, a thin-film sample (thickness:
700 to 1200 nm) for transmission electron microscope (TEM) observation was prepared
by electropolishing using two solutions, i.e., a 1:9 mixture of perchloric acid and
ethanol, and a 1:3 mixture of nitric acid and methanol. The microstructure of the
thin-film sample was observed in five view fields at an acceleration voltage of the
transmission electron microscope of 120 kV while applying beams at <001> to the matrix
and performing the observation in the (100) plane with respect to the matrix. The
observation was performed at 500000-fold magnification. The major axes of Q phase
were measured based on the observed microstructures, and the average of major axes
of Q phases in all the five view fields was calculated and defined as the major axis
of Q phase.
[0141] The tensile strength, 0.2% yield strength, and elongation were measured under conditions
in conformance with the conditions prescribed in JIS Z 2241:2011. The Charpy impact
value was measured under conditions in conformance with the conditions prescribed
in JIS Z 2242:2005. Criteria for good properties are 380 MPa or more for the tensile
strength; 360 MPa or more for the 0.2% yield strength; 10% or more for the elongation;
and 10 J/cm
2 or more for the Charpy impact value.
[0142] Independently, a C-ring shaped specimen was sampled from each Al alloy forging and
subjected to a stress corrosion cracking test. The stress corrosion cracking test
was performed using the C-ring shaped specimen under conditions in conformance with
the conditions of the alternate immersion test prescribed in ASTM G47 (2011). In the
test, the C-ring shaped specimen was immersed in a sodium chloride solution for 10
minutes, and retrieved from the sodium chloride solution and air-dried for 50 minutes;
and this procedure (immersion and drying) was repeated for 90 days, while applying
stress of 75% of the yield strength in the LT direction (long transverse direction)
to the specimen. The specimen was then observed to determine whether stress corrosion
cracking occurred. A sample suffering from stress corrosion cracking was evaluated
as having poor stress corrosion cracking resistance (×); a sample suffering from not
stress corrosion cracking, but grain-boundary corrosion was evaluated as having somewhat
poor stress corrosion cracking resistance (Δ), where the grain-boundary corrosion
may highly possibly lead to stress corrosion cracking; a sample suffering from neither
stress corrosion cracking nor grain-boundary corrosion was evaluated as having good
stress corrosion cracking resistance (○); and a sample not suffering from corrosion
at all in the aluminum alloy portion was evaluated as having very good stress corrosion
cracking resistance (⊚). The results are presented in Table 2.
[Table 1]
Alloy number |
Al alloy chemical composition |
(in mass percent, with the remainder being Al) |
Mg |
Si |
Cu |
Fe |
Ti |
Mn |
Cr |
Zr |
Zn |
1 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
2 |
0.70 |
0.80 |
0.30 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
3 |
1.50 |
0.80 |
0.30 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
4 |
1.50 |
1.30 |
0.90 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
5 |
0.70 |
1.30 |
0.90 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
6 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.60 |
0.15 |
- |
<0.01 |
7 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
- |
0.45 |
0.20 |
<0.01 |
8 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.60 |
- |
0.05 |
<0.01 |
9 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.10 |
0.45 |
- |
<0.01 |
10 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
- |
0.10 |
0.30 |
<0.01 |
11 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
- |
0.45 |
- |
<0.01 |
12 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.60 |
- |
- |
<0.01 |
13 |
1.00 |
0.80 |
0.50 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
14 |
1.00 |
0.90 |
0.50 |
0.40 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
15 |
1.00 |
0.80 |
0.50 |
0.15 |
0.005 |
0.50 |
0.15 |
0.10 |
<0.01 |
16 |
1.00 |
1.10 |
0.50 |
0.20 |
0.15 |
0.50 |
0.15 |
0.10 |
<0.01 |
17 |
1.00 |
1.00 |
0.50 |
0.20 |
0.15 |
- |
- |
0.30 |
<0.01 |
18 |
1.00 |
1.00 |
0.50 |
0.20 |
0.15 |
0.10 |
0.10 |
0.05 |
<0.01 |
19 |
0.50 |
1.30 |
0.90 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
20 |
1.80 |
1.30 |
0.90 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
21 |
1.00 |
0.60 |
0.90 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
22 |
1.00 |
1.80 |
0.90 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
23 |
1.00 |
1.30 |
0.10 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
24 |
1.00 |
1.30 |
1.20 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
25 |
1.90 |
1.70 |
1.10 |
0.20 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
25 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
- |
- |
- |
<0.01 |
27 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.70 |
0.15 |
0.10 |
<0.01 |
28 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.50 |
0.50 |
0.10 |
<0.01 |
29 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.50 |
0.15 |
0.35 |
<0.01 |
30 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.70 |
0.50 |
0.35 |
<0.01 |
31 |
1.00 |
1.10 |
0.50 |
0.45 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
32 |
1.00 |
1.10 |
0.50 |
0.03 |
0.03 |
0.50 |
0.15 |
0.10 |
<0.01 |
33 |
1.00 |
1.10 |
0.50 |
0.20 |
0.17 |
0.50 |
0.15 |
0.10 |
<0.01 |
34 |
1.00 |
1.10 |
0.50 |
0.20 |
0.03 |
0.05 |
0.05 |
0.03 |
<0.01 |
35 |
1.00 |
1.10 |
0.50 |
0.20 |
- |
0.50 |
0.15 |
0.10 |
<0.01 |
Notes: Underfined data are out of the scope of the first embodiment of the present
invention; and the symbol "-" refers to that the element in question is not contained. |
[0143] As demonstrated in Tables 1 and 2, Al alloy forgings (Sample Nos. 1 to 18; Examples)
meeting the conditions specified within the scope of the present invention had a tensile
strength, a 0.2% yield strength, a Charpy impact value, and stress corrosion cracking
resistance at excellent levels. In contrast, Al alloy forgings (Sample Nos. 19 to
35; Comparative Examples) not meeting the conditions specified within the scope of
the present invention had disadvantages as follows.
[0144] Sample No. 19 had a Mg content less than the lower limit and a major axis of Q phase
less than the lower limit and was inferior in 0.2% yield strength, elongation, Charpy
impact value, and stress corrosion cracking resistance. Sample No. 20 had a Mg content
greater than the upper limit and was inferior in elongation, Charpy impact value,
and stress corrosion cracking resistance. Sample No. 21 had a Si content less than
the lower limit and a major axis of Q phase less than the lower limit and was inferior
in all the evaluated properties. Sample No. 22 had a Si content greater than the upper
limit and was inferior in elongation, Charpy impact value, and stress corrosion cracking
resistance. Sample No. 23 had a Cu content less than the lower limit, was approximately
devoid of Q phases, and was inferior in tensile strength and 0.2% yield strength.
Sample No. 24 had a Cu content greater than the upper limit and a major axis of Q
phase greater than the upper limit and was inferior in elongation, Charpy impact value,
and stress corrosion cracking resistance. Sample No. 25 had Mg, Si, and Cu contents
each greater than the upper limit, had a major axis of Q phase greater than the upper
limit, and was inferior in elongation, Charpy impact value, and stress corrosion cracking
resistance.
[0145] Sample No. 26 did not contain Mn, Cr, and Zr and was inferior in tensile strength,
0.2% yield strength, elongation, and Charpy impact value. Sample No. 27 had a Mn content
greater than the upper limit and was inferior in tensile strength, 0.2% yield strength,
and Charpy impact value. Sample No. 28 had a Cr content greater than the upper limit
and was inferior in tensile strength, 0.2% yield strength, Charpy impact value, and
stress corrosion cracking resistance. Sample No. 29 had a Zr content greater than
the upper limit and was inferior in tensile strength, 0.2% yield strength, Charpy
impact value, and stress corrosion cracking resistance. Sample No. 30 had Mn, Cr,
and Zr contents each greater than the upper limit and was inferior in tensile strength,
0.2% yield strength, elongation, and Charpy impact value.
[0146] Sample No. 31 had an Fe content greater than the upper limit and was inferior in
elongation, Charpy impact value, and stress corrosion cracking resistance. Sample
No. 32 had an Fe content less than the lower limit, suffered from cracking upon casting,
and could not be forged. Sample No. 33 had a Ti content greater than the upper limit
and was inferior in Charpy impact value. Sample No. 34 had Mn, Cr, and Zr contents
each less than the lower limit and was inferior in tensile strength and 0.2% yield
strength. Sample No. 35 did not contain Ti, thereby had a coarse as-cast structure,
and suffered from cracking upon forging.
Second Experimental Example
[0147] Al alloy forgings were prepared under conditions given in Table 3. The Al alloy forgings
each had the shape of an automobile suspension member illustrated in Figs. 3 and 4.
The forgings had a thickness in the thinnest portion of 6 mm. In each sample, homogenization
was performed at a holding temperature of from 550°C to 570°C for a holding time of
4 hours; and forging was performed at a forging start temperature of from 500°C to
520°C and a forging finish temperature of from 380°C to 420°C. Solution heat treatment
was performed at a temperature of from 550°C to 580°C. Other conditions in the tempering
step are as given in Table 3. Other production conditions are as with First Experimental
Example.
[0148] Three test specimens were sampled from a maximum-stress-receiving region in each
Al alloy forging and subjected to measurements of average grain size in terms of minor
axis, area percentage, and major axis of Q phase as microstructure form, as presented
in Table 3. The test specimens were also subjected to determination of properties
of the Al alloy forging, including tensile properties such as tensile strength, 0.2%
yield strength, and elongation, and Charpy impact value (mechanical property), where
the tensile properties are indices of strength, and the Charpy impact value is an
index of toughness. Each of data in Table 3 is an average of values determined on
the three test specimens sampled from eachAl alloy forging. The measurement of the
microstructure form and the evaluation of the properties of the Al alloy forging were
performed by methods as in First Experimental Example. The results are presented in
Table 3.
[0149] As presented in Table 3, Al alloy forgings (Sample Nos. 36 to 40; Examples) meeting
the conditions specified within the scope of the present invention were excellent
in tensile strength, 0.2% yield strength, Charpy impact value, and stress corrosion
cracking resistance. In contrast Al alloy forgings (Sample Nos. 41 to 47; Comparative
Examples) not meeting the conditions specified within the scope of the present invention
have disadvantages as follows.
[0150] Sample No. 41 underwent quenching in the tempering step performed at a quenching
temperature lower than the lower limit, had a major axis of Q phase greater than the
upper limit, and was inferior in elongation and stress corrosion cracking resistance.
Sample No. 42 underwent quenching in the tempering step performed at a quenching temperature
higher than the upper limit for a quenching time longer than the upper limit, had
a major axis of Q phase less than the lower limit, and was inferior in tensile strength
and 0.2% yield strength. Sample No. 43 underwent drying in the tempering step performed
for a drying time longer than the upper limit, had a major axis of Q phase less than
the lower limit, and was inferior in tensile strength, 0.2% yield strength, and elongation.
[0151] Sample No. 44 underwent over-aging in the tempering step performed at an over-aging
temperature higher than the upper limit, had a major axis of Q phase greater than
the upper limit, and was inferior in tensile strength, 0.2% yield strength, elongation,
and Charpy impact value. Sample No. 45 underwent over-aging in the tempering step
performed at an over-aging temperature lower than the lower limit, had a major axis
of Q phase less than the lower limit, and was inferior in tensile strength, 0.2% yield
strength, and stress corrosion cracking resistance. Sample No. 46 underwent over-aging
in the tempering step performed for an over-aging time longer than the upper limit,
had a major axis of Q phase greater than the upper limit, and was inferior in tensile
strength, 0.2% yield strength, elongation, and Charpy impact value. Sample No. 47
underwent over-aging in the tempering step performed for an over-aging time shorter
than the lower limit, had a major axis of Q phase less than the lower limit, and was
inferior in tensile strength, 0.2% yield strength, and stress corrosion cracking resistance.
Third Experimental Example
[0152] Al alloy forgings were subjected to surface treatments under conditions given in
Table 4 and subjected to evaluation of stress corrosion cracking resistance. The evaluation
was performed by the method as in First Experimental Example.
[0153] As demonstrated in Table 4, the Al alloy forgings of Sample Nos. 48 to 50, which
underwent the surface treatments, had excellent stress corrosion cracking resistance
as compared with the Al alloy forging of Sample No. 51, which did not undergo a surface
treatment.
List of Reference Signs
[0155]
- 1
- automobile suspension part (Al alloy forging)
- 2a and 2b
- arm portion
- 3a, 3b, and 3c
- rib
- 4a and 4b
- web
- 5a, 5b, and 5c
- joint portion
- 6a and 6b
- maximum-stress-receiving region (cross sectional direction)
- 7, 8, and 9
- sampling site
- 10
- grain
- 21
- cluster I
- 22
- cluster II
- 25
- GP-zone
- 30
- Q phase
- 31
- ß phase