Question: What are the solution heat treatment and aging parameters (time in hours) for 7xxx series aluminum alloys like 7075, 7050, 7150?

ASM Handbook, Volume 4: Heat Treating
ASM Handbook Committee, p 841-879
DOI: 10.1361/asmhba0001205
Copyright © 1991 ASM InternationalⓇ
All rights reserved.
www.asminternational.org
Heat Treating of Aluminum Alloys
HEAT TREATING in its broadest sense,
refers to any of the heating and cooling
operations that are performed for the pur-
pose of changing the mechanical properties,
the metallurgical structure, or the residual
stress state of a metal product. When the
term is applied to aluminum alloys, howev-
er, its use frequently is restricted to the
specific operations employed to increase
strength and hardness of the precipitation-
hardenable wrought and cast alloys. These
usually are referred to as the "heat-treat-
able" alloys to distinguish them from those
alloys in which no significant strengthening
can be achieved by heating and cooling. The
latter, generally referred to as "non-heat-
treatable" alloys, depend primarily on cold
work to increase strength. Heating to de-
crease strength and increase ductility (an-
nealing) is used with alloys of both types;
metallurgical reactions may vary with type
of alloy and with degree of softening desired.
Except for the low-temperature stabilization
treatment sometimes given for 5xxx series
alloys (which is a mill treatment and not
discussed in this article), complete or partial
annealing treatments are the only ones used
for non-heat-treatable alloys. A general
overview of these heat treatments is covered
in the article "Principles of Heat Treating of
Nonferrous Alloys" in this Volume.
Precipitation from Solid Solution
One essential attribute of a precipitation-
hardening alloy system is a temperature-
dependent equilibrium solid solubility char-
acterized by increasing solubility with
increasing temperature (see, for example,
the phase diagrams in Fig 1 and 2). Al-
though this condition is met by most of the
binary aluminum alloy systems, many ex-
hibit very little precipitation hardening, and
these alloys ordinarily are not considered
heat treatable. Alloys of the binary alumi-
num-silicon and aluminum-manganese sys-
tems, for example, exhibit relatively insig-
nificant changes in mechanical properties as
a result of heat treatments that produce
considerable precipitation. The major alu-
minum alloy systems with precipitation
hardening include:
•
systems
Aluminum-copper-magnesium
(magnesium intensifies precipitation)
systems
• Aluminum-magnesium-silicon
with strengthening from Mg₂Si
Aluminum-zinc-magnesium systems with
strengthening from MgZn2
•
• Aluminum-zinc-magnesium-copper
tems
sys-
The general requirement for precipitation
strengthening of supersaturated solid solu-
tions involves the formation of finely dis-
persed precipitates during aging heat treat-
ments (which may include either natural aging
or artificial aging). The aging must be accom-
plished not only below the equilibrium solvus
temperature, but below a metastable miscibil-
ity gap called the Guinier-Preston (GP) zone
solvus line. The supersaturation of vacancies
allows diffusion, and thus zone formation, to
occur much faster than expected from equi-
librium diffusion coefficients. In the precipi-
tation process, the saturated solid solution
first develops solute clusters, which then be-
come involved in the formation of transitional
(nonequilibrium) precipitates.
Temperature, °C
800
600
400
200
ΑΙ
Al + CuAl2
The mechanism of strengthening from
precipitation involves the formation of co-
herent clusters of solute atoms (that is, the
solute atoms have collected into a cluster
but still have the same crystal structure as
the solvent phase). This causes a great deal
of strain because of mismatch in size be-
tween the solvent and solute atoms. Conse-
quently, the presence of the precipitate par-
ticles, and even more importantly the strain
fields in the matrix surrounding the coher-
ent particles, provide higher strength by
obstructing and retarding the movement of
dislocations. The characteristic that deter-
mines whether a precipitate phase is coher-
ent or noncoherent is the closeness of
match or degree of disregistry between
atomic spacings on the lattice of the matrix
and on that of the precipitate. These
changes in properties result from the forma-
tion of solute-rich microstructural domains,
or GP zones.
The exact size, shape, and distribution of
GP zones depend on the alloy in which they
form and on the thermal and mechanical
history of the specimen. Their shape can
(b)
L
A1 + L
Temperature range for
solution heat treating
Temperature range
for annealing
Temperature range for
precipitation heat
treating
2
4
6
8
10
12
Copper.%
1400
1000
600
200
Temperature, °F
Aluminum-copper systems with strength- Fig 1 Portion of aluminum-copper binary phase diagram. Temperature ranges for annealing, precipitation heat
•
ening from CuAl2
treating, and solution heat treating are indicated. The range for solution treating is below the eutectic
melting point of 548 °C (1018 °F) at 5.65 wt% Cu.
842 / Heat Treating of Nonferrous Alloys
Temperature, °C
(a)
Magnesium, wt%
(b)
700
600
Solidus
500
400
Solvus
300
200
100
1200
1000
595 °C
at 1.85%
Mg2Si
800
600
400
200
Mg-Si ratio of 1.73:1
1
0
0.5
1.0
1.5
2.0
Mg2Si, %
570
1.0
Temperature, °F
660 750 840 930 1020 1110
1.4
1.2
0.8
Solvus with silicon
and Mg2Si present
1.0
0.6
0.8
0.6
0.4
0.4
0.2
0
300 350 400 450 500 550 600
Temperature, °C
0.2
0
Temperature, °F
Silicon, wt%
Fig 2 Equilibrium solubility as function of tempera-
ture for (a) Mg2Si in aluminum with an Mg-Si
ratio of 1.73-to-1 and (b) magnesium and silicon in solid
aluminum when both Mg2Si and silicon are present
sometimes be deduced by refined studies of
diffuse x-ray scattering. Under favorable
conditions, GP zones can be seen in trans-
mission electron micrographs. Spherical
solute-rich zones usually form when the
sizes of the solvent and solute atoms are
nearly equal, as in the aluminum-silver and
aluminum-zinc systems. If there is a large
difference in atom sizes, as in the alumi-
num-copper system, the GP zones usually
form as disks whose planes lie parallel with
some low-index plane of the matrix lattice.
Sometimes, the solute atoms occupy pre-
ferred lattice sites within the GP zone, and
thus form a small region of lattice order.
The GP zones are of the size range of tens
of angstroms in diameter. They are essential-
ly distorted regions of the matrix lattice, rath-
er than discrete particles of a new phase
having a different lattice. As such, they are
completely coherent with the matrix, impos-
ing local but often large strains on it. These
mechanical strains, as well as the presence of
a locally solute-rich, sometimes ordered lat-
tice, can account for large changes in me-
chanical properties of the alloy before any
long-range microstructural changes occur.
The GP zones are characteristically meta-
stable and thus dissolve in the presence of a
more stable precipitate. This dissolution
causes a precipitate-free, visibly denuded
region to form around the stable precipitate
particles. The final structure consists of
equilibrium precipitates, which do not con-
tribute as significantly to hardening. More
detailed information about preprecipitation
phenomena can be found in the article
"Structures Resulting From Precipitation
From Solid Solution" in Volume 9 of the
9th Edition of Metals Handbook.
Precipitation in Aluminum-Copper Alloys.
Figure 1, which illustrates the required sol-
ubility-temperature relationship needed in
precipitation strengthening, shows the tem-
perature ranges required for solution treat-
ment and subsequent precipitate hardening
in the aluminum-copper system. The equi-
librium solid solubility of copper in alumi-
num increases as temperature increases-
from about 0.20% at 250 °C (480 °F) to a
maximum of 5.65% at the eutectic melting
temperature of 548 °C (1018 °F). (It is con-
siderably lower than 0.20% at temperatures
below 250 °C.) For aluminum-copper alloys
containing from 0.2 to 5.6% Cu, two distinct
equilibrium solid states are possible. At
temperatures above the lower curve in Fig 1
(solvus), the copper is completely soluble,
and when the alloy is held at such temper-
atures for sufficient time to permit needed
diffusion, the copper will be taken com-
pletely into solid solution. At temperatures
below the solvus, the equilibrium state con-
sists of two solid phases: solid solution, a,
plus an intermetallic-compound phase 0
(Al₂Cu). When such an alloy is converted to
all solid solution by holding above the sol-
vus temperature and then the temperature is
decreased to below the solvus, the solid
solution becomes supersaturated and the
alloy seeks the equilibrium two-phase con-
dition; the second phase tends to form by
solid-state precipitation.
The preceding description is a gross over-
simplification of the actual changes that
occur under different conditions even in
simple binary aluminum-copper alloys. A
variety of different nonequilibrium precipi-
tate structures is formed at temperatures
below solvus. In alloys of the aluminum-
copper system, a succession of precipitates
is developed from a rapidly cooled super-
saturated solid solution (SSS). These pre-
cipitates develop sequentially either with
increasing temperature or with increasing
time at temperature between room temper-
ature and the solvus. The several stages are
identified by the following notation:
SSSGP zones →0" →
0'→0 (Al₂Cu)
At temperatures in the natural aging
range (about 20 to 60 °C, or 0 to 140 °F),
the distribution of copper atoms changes
with time from random to the disklike pla-
nar aggregates (GP zones), which form on
particular crystallographic planes of the alu-
minum matrix. These aggregates create co-
herency strain fields that increase resis-
tance to deformation, and their formation is
responsible for the changes in mechanical
properties that occur during natural aging.
At higher temperatures, transition forms of
approximate composition Al₂Cu develop
and further increase strength. In the highest
strength condition, both the 0" and 0' tran-
sition precipitates may be present. When
time and temperature are increased suffi-
ciently to form high proportions of the equi-
librium 0, the alloy softens and is said to be
"overaged."
The commercial heat-treatable aluminum
alloys are, with few exceptions, based on
ternary or quaternary systems with respect
to the solutes involved in developing
strength by precipitation. Commercial al-
loys whose strength and hardness can be
significantly increased by heat treatment
include 2xxx, 6xxx, and 7xxx series wrought
alloys (except 7072) and 2xx.0, 3xx.0, and
7xx.0 series casting alloys. Some of these
contain only copper, or copper and silicon,
as the primary strengthening alloy addi-
tion(s). Most of the heat-treatable alloys,
however, contain combinations of magne-
sium with one or more of the elements
copper, silicon, and zinc. Characteristical-
ly, even small amounts of magnesium in
concert with these elements accelerate and
accentuate precipitation hardening, while
alloys in the 6xxx series contain silicon and
magnesium approximately in the propor-
tions required for formulation of magnesium
silicide (Mg2Si). Although not as strong as
most 2xxx and 7xxx alloys, 6xxx series al-
loys have good formability, weldability, ma-
chinability, and corrosion resistance, with
medium strength.
In the heat-treatable wrought alloys, with
some notable exceptions (2024, 2219, and
7178), such solute elements are present in
amounts that are within the limits of mutual
solid solubility at temperatures below the
eutectic temperature (lowest melting tem-
perature). In contrast, some of the casting
alloys of the 2xx.0 series and all of the 3xx.0
series alloys contain amounts of soluble
elements that far exceed solid-solubility
limits. In these alloys, the phase formed by
combination of the excess soluble elements
with the aluminum will never be dissolved,
although the shapes of the undissolved par-
ticles may be changed by partial solution.
Most of the heat-treatable aluminum alloy
systems exhibit multistage precipitation and
undergo accompanying strength changes
analogous to those of the aluminum-copper
system. Multiple alloying additions of both
major solute elements and supplementary
elements employed in commercial alloys are
strictly functional and serve with different
heat treatments to provide the many differ-
ent combinations of properties—physical,
mechanical, and electrochemical-that are
required for different applications. Some al-
loys, particularly those for foundry produc-
tion of castings, contain amounts of silicon
far in excess of the amount that is soluble or
needed for strengthening alone. The function
here is chiefly to improve casting soundness
and freedom from cracking, but the excess
silicon also serves to increase wear resis-
tance, as do other microstructural constitu-
ents formed by manganese, nickel, and iron.
Parts made of such alloys are commonly
used in gasoline and diesel engines (pistons,
cylinder blocks, and so forth).
Alloys containing the elements silver,
lithium, and germanium are also capable of
providing high strength with heat treatment,
and in the case of lithium, both increased
elastic modulus and lower density, which
are highly advantageous-particularly for
aerospace applications (see the following
section "Aluminum-Lithium Alloys" in this
article). Commercial use of alloys contain-
ing these elements has been restricted either
by cost or by difficulties encountered in
producing them. Such alloys are used to
some extent, however, and research is be-
ing directed toward overcoming their disad-
vantages.
In the case of alloys having copper as the
principal alloying ingredient and no magne-
sium, strengthening by precipitation can be
greatly increased by adding small fractional
percentages of tin, cadmium, or indium, or
combinations of these elements. Alloys
based on these effects have been produced
commercially but not in large volumes be-
cause of costly special practices and limita-
tions required in processing, and in the case
of cadmium, the need for special facilities to
avoid health hazards from formation and
release of cadmium vapor during alloying.
Such alloys, as well as those containing
silver, lithium, or other particle-forming el-
ements, may be used on a selective basis in
the future.
Aluminum-Lithium Alloys. Like other age-
hardened aluminum alloys, aluminum-lithi-
um alloys achieve precipitation strengthen-
ing by thermal aging after a solution heat
treatment. The precipitate structure is sen-
sitive to a number of processing variables,
including, but not limited to, the quenching
rate following the solution heat treatment,
the degree of cold deformation prior to
aging, and the aging time and temperature.
Minor alloying elements can also have a
significant effect on the aging process by
changing the interface energy of the precip-
itate, by increasing the vacancy concentra-
tion, and/or by raising the critical tempera-
ture for homogeneous precipitation. Like
some other age-hardened 2xxx aluminum
alloys, aluminum-lithium-base alloys also
gain increased strength and toughness from
deformation prior to aging. This unusual
phenomenon has given rise to a number of
thermomechanical processing steps for alu-
minum-lithium alloys aimed at optimizing
mechanical properties after artificial aging.
The age hardening of aluminum-lithium
alloys involves the continuous precipitation
of d' (AlLi) from a supersaturated solid
solution. The aluminum and lithium in the 8'
precipitates are positioned at specific loca-
tions. The eight shared corner sites are
occupied by lithium, and the six shared
faces are occupied by aluminum. This gives
rise to the aluminum-lithium composition of
8' precipitates. The geometrical similarity
between the lattice of the precipitates and
the face-centered cubic lattices of the solid
solution facilitates the observed cube/cube
orientation. The lattice parameters of the
precipitate are also closely matched to
those of the matrix. Consequently, the mi-
crostructure of an aluminum-lithium alloy
solution heat treated and aged for short
times below the d' solvus is characterized
by a homogeneous distribution of coherent,
spherical d' precipitates.
Aluminum-lithium-base alloys are micro-
structurally unique. They differ from most
of the aluminum alloys in that once the
major strengthening precipitate (8') is ho-
mogeneously precipitated, it remains coher-
ent even after extensive aging. In addition,
extensive aging at high temperatures (>190
°C, or 375 °F) can result in the precipitation
of icosahedral grain-boundary precipitates
with five-fold symmetry. Although the qua-
si-crystalline structure and the composition
of these grain-boundary precipitates are not
yet exactly known, it has been suggested
that both the precipitates and the precipi-
tate-free zones (PFZs) near the grain bound-
aries might play a major role in the fracture
process.
The low ductility and toughness of binary
aluminum-lithium alloys can be traced, at
least in part, to the inhomogeneous nature
of their slip, resulting from coherent-parti-
cle hardening of spherical d' precipitates.
The presence of equilibrium 6 (aluminum-
lithium) precipitates at grain boundaries can
also cause PFZs, which can induce further
strain localization and promote intergranu-
lar failure. Consequently, for the develop-
ment of commercial alloys, slip has been
homogenized by introducing dispersoids
(manganese, zirconium) and semicoherent/
incoherent precipitates, such as
(Al₂CuLi), 0' (Al₂Cu), or S (Al₂LiMg),
through copper or magnesium additions.
Magnesium and copper improve the
strength of aluminum-lithium alloys through
solid-solution and precipitate strengthening,
and they can minimize the formation of
PFZs near grain boundaries. Zirconium,
which forms the cubic Al3Zr coherent dis-
persoid, stabilizes the subgrain structure
and suppresses recrystallization.
T₁
Development of commercially available
aluminum-lithium-base alloys was started
by adding lithium to aluminum-copper, alu-
minum-magnesium, and aluminum-copper-
Heat Treating of Aluminum Alloys / 843
magnesium alloys. These alloys were cho-
sen to superimpose the precipitation-
hardening characteristics of aluminum-
copper-, aluminum-copper-magnesium-,
and aluminum-magnesium-base precipitates
to the hardening of lithium-containing pre-
cipitates. Proceeding in this manner, alloys
2020 (Al-Cu-Li-Cd), 01429 (Al-Mg-Li), 2090
(Al-Cu-Li), and 2091 and 8090 (Al-Cu-Mg-
Li) evolved. Besides these registered al-
loys, other commercial aluminum-lithium
alloys include Weldalite 049 and CP276.
Properties and applications of these alloys
are discussed in the article "Aluminum-
Lithium Alloys" in Volume 2 of the 10th
Edition of Metals Handbook.
In terms of d' precipitation, the only
effect of magnesium appears to be a reduc-
tion in the solubility of lithium. The micro-
structure of an aluminum-magnesium-lithi-
um alloy in the early stages of aging is
similar to that of an aluminum-lithium alloy.
Precipitation in the aluminum-copper-lithi-
um system is more complicated than that in
either the aluminum-lithium or aluminum-
magnesium-lithium systems.
Effects on Physical and Electrochemical
Properties. The above description of the
precipitation processes in commercial heat-
treatable aluminum alloys (as well as the
heat-treatable binary alloys, none of which
is used commercially in wrought form) af-
fect not only mechanical properties but also
physical properties (density and electrical
and thermal conductivities) and electro-
chemical properties (solution potential). On
the microstructural and submicroscopic
scales, the electrochemical properties de-
velop point-to-point nonuniformities that
account for changes in corrosion resistance.
Measurements of changes in physical and
electrochemical properties have played an
important role in completely describing pre-
cipitation reactions and are very useful in
analyzing or diagnosing whether heat-treat-
able products have been properly or im-
properly heat treated. Although they may
be indicative of the strength levels of prod-
ucts, they cannot be relied upon to deter-
mine whether or not the product meets
specified mechanical-property limits. Since
elements in solid solution are always more
harmful to electrical conductivity than the
same elements combined with others as
intermetallic compounds, thermal treat-
ments are applied to ingots used for fabri-
cation of electrical conductor parts. These
thermal treatments are intended to precipi-
tate as much as possible of the dissolved
impurities. Iron is the principal element
involved, and although the amount precip-
itated is only a few hundredths of a percent,
the effect on electrical conductivity of the
wire, cable, or other product made from the
ingot is of considerable practical impor-
tance. These alloys may or may not be heat
treatable with respect to mechanical prop-
erties. Electrical conductor alloys 6101 and
844 / Heat Treating of Nonferrous Alloys
6201 are heat treatable. These alloys are
used in tempers in which their strengthening
precipitate, the transition form of Mg₂Si, is
largely out of solid solution to optimize both
strength and conductivity.
Strengthening by Heat Treatment
•
•
Heat treatment to increase strength of
aluminum alloys is a three-step process:
Solution heat treatment: dissolution of
soluble phases
Quenching: development of supersatura-
tion
Age hardening: precipitation of solute at-
oms either at room temperature (natural
aging) or elevated temperature (artificial
aging or precipitation heat treatment)
Each of these steps and the use of quench-
factor analysis are described in the follow-
ing four sections. Typical solution and pre-
cipitation heat treatments for mill products
are given in Tables 1(a, b, and c) and 2, and
treatments for castings are given in Table 3.
Temper designations are defined at the end
of this article.
Solution Heat Treating
To take advantage of the precipitation-
hardening reaction, it is necessary first to
produce a solid solution. The process by
which this is accomplished is called solution
heat treating, and its objective is to take into
solid solution the maximum practical
amounts of the soluble hardening elements
in the alloy. The process consists of soaking
the alloy at a temperature sufficiently high
and for a time long enough to achieve a
nearly homogeneous solid solution.
Nominal commercial solution heat-treat-
ing temperature is determined by the com-
position limits of the alloy and an allowance
for unintentional temperature variations.
Although ranges normally listed allow vari-
ations of ±6 °C (±10 °F) from the nominal,
some highly alloyed, controlled-toughness,
high-strength alloys require that tempera-
ture be controlled within more restrictive
limits. Broader ranges may be allowable for
alloys with greater intervals of temperature
between their solvus and eutectic melting
temperatures.
Overheating. Care must be exercised to
avoid exceeding the initial eutectic melting
temperature. If appreciable eutectic melting
occurs as a result of overheating, properties
such as tensile strength, ductility, and frac-
ture toughness may be degraded. Materials
that exhibit microstructural evidence of
overheating are generally categorized as
unacceptable by specification. Evidence of
grain-boundary melting that occurs above
the eutectic melting temperature of the alloy
usually is not detectable by either visual
examination or nondestructive testing.
Although maximum temperature must be
restricted to avoid melting, the lower limit
should, when possible, be above the temper-
ature at which complete solution occurs (sol-
vus). In the alloy represented by line (a) in Fig
1, these temperatures would be about 575 and
515 °C (1065 and 960 °F), respectively. How-
ever, under production conditions, the tem-
perature interval for solution treatment
(shown in Fig 1 for typical 2xxx or 2xx.x)
alloys provides a margin to safeguard against
eutectic melting and a cushion on the low side
for increased solution and diffusion rates.
For alloys containing more than 5.65%
Cu, complete solution can never occur. For
these alloys, such as alloy 2219 (which has
5.8 to 6.8% Cu), the minimum solution heat-
treating temperature is established so that it
is as close as practical to the eutectic tem-
perature while providing a margin of safety
commensurate with the capability of the
equipment. Line (b) in Fig 1 is another
example of a composition above 5.65% Cu
that does not allow complete dissolution of
aluminum-copper precipitates.
For more complex ternary and quaterna-
ry systems, solution treatments are modi-
fied according to the effect of new elements
on the solid solubility and/or the eutectic
melting points of the basic binary system. In
aluminum-lithium alloys, for example, mag-
nesium reduces the solubility of lithium in
aluminum. In the aluminum-copper system,
magnesium also lowers the eutectic melting
point. The proximity of typical solution-
treating temperature ranges to eutectic
melting temperatures for three common alu-
minum-copper-magnesium alloys is shown
in the following table:
Solution-treating
temperature
Eutectic melting
temperature
°F
Alloy
°C
°F
°C
2014
2017
2024
496-507
496-507
488-499
925-945
925-945
910-930
510
950
513
502
955
935
Similar considerations apply to other age-
hardenable alloy systems such as aluminum-
magnesium-silicon alloys. For example, ac-
cording to Fig 2(a), a 1.08% Mg₂Si alloy
would be soaked at a temperature in excess of
500 °C (930 °F) but below the solidus of 595 °C
(1100 °F) to avoid incipient melting. Howev-
er, because some alloy constituents may form
complex eutectics that melt at temperatures
below the equilibrium eutectic temperature,
the upper limit for solution treatment of alu-
minum-magnesium-silicon alloys is in the
range of 515 to 540 °C (960 to 1000 °F). At 540
°C (1000 °F), about 0.6% Mg can be placed in
solution (Fig 2b).
Nonequilibrium Melting. When high heat-
ing rates are employed, the phenomenon of
nonequilibrium melting must be considered.
This phenomenon can also be explained with
the help of the aluminum-copper phase dia-
gram (Fig 1). The room-temperature micro-
structure of an F-temper product containing
4% Cu consists of a solid solution of copper in
aluminum and particles of Al₂Cu. When this
product is heated slowly, the Al₂Cu begins to
dissolve, and if heating is slow enough, all of
the Al,Cu is dissolved when temperatures
above the solvus (500 °C, or 932 °F) are
reached. When the heating rate is high, how-
ever, much of the Al₂Cu remains undis-
solved. If a material with this microstructure
is heated at or above the eutectic temperature
of 548 °C (1018 °F), melting will begin at the
interface between the Al₂Cu and the matrix.
With sufficient time above the eutectic tem-
perature, this metastable liquid will dissolve
to form a solid solution and will leave no trace
provided that hydrogen gas has not con-
densed at the interface to form a void. If the
product is quenched before the liquid has
time to equilibrate, however, it will solidify
and form fine eutectic rosettes. This nonequi-
librium melting should not be confused with
true equilibrium melting, which would occur
in any alloy containing more than 5.65% Cu.
In such an alloy, eutectic melting is equilibri-
um melting. No matter how long such an
alloy is held above the eutectic temperature,
the liquid will never solidify. In commercial
alloys, which usually are ternaries or quater-
naries of the major alloying elements, the
situation is more complex. Different phases
have different solvus temperatures, and non-
equilibrium melting may occur at different
temperatures depending on composition, size
of precipitates, and rate of heating. When
new solution heat-treating equipment (which
provides higher heating rates) is employed,
careful examination of alloy microstructures
should be included as part of the certification
process.
Underheating. When the temperatures at-
tained by the parts or pieces being heat treat-
ed are appreciably below the normal range,
solution is incomplete, and strength some-
what lower than normal is expected. In the
aluminum-copper system (Fig 1), the shallow
slope of the solvus at its intersection with the
composition line indicates that a slight de-
crease in temperature will result in a large
reduction in the concentration of the solid
solution and a correspondingly significant de-
crease in final strength. The effect of solution-
treating temperature on the strength of two
aluminum alloys is illustrated by the following
data:
°C
Solution-
treating
temperature
°F
MPa
Tensile strength
ksi
Yield strength
MPa
ksi
6061-T6 sheet 1.6 mm (0.064 in.) thick
493
920
301
43.7
272
39.4
504
940
316
45.8
288
41.7
516
960
333
48.3
305
44.3
527
980
348
50.5
315
45.7
2024-T4 sheet 0.8 mm (0.032 in.) thick
488
910
419
60.8
255
37.0
491
915
422
61.2
259
37.5
493
920
433
62.8
269
39.0
496
925
441
63.9
271
39.3
Table 1(a)
alloying
Heat Treating of Aluminum Alloys / 845
Typical solution and precipitation heat treatments for commercial heat-treatable aluminum alloy mill products with copper
Alloy
Product form
Al-Cu alloys without magnesium alloying
Solution heat treatment(a)
Precipitation heat treatment
Metal temperature(b)
Metal temperature(b)
Temper
Time(c),
°C
°F
designation
°C
°F
h
Temper
designation
2011
Rolled or cold finished rod and bar
525
975
T3(d)
160
320
14
T8(d)
T4
T451(e)
2025
Die forgings
515
960
T4
170
340
2219(f)
Flat sheet
535
995
T31(d)
175
350
T37(d)
165
325
T42
190
375
Plate
535
995
T31(d)
175
350
T37(d)
175
350
T351(e)
175
350
T42
190
375
2219(f)
Rolled or cold finished wire, rod, and bar
Extruded rod, bar, shapes, and tube
535
995
T351(e)
190
375
535
995
T31(d)
190
375
T3510(e)
190
375
T3511(e)
190
375
T42
190
375
Die forgings and rolled rings
535
995
T4
190
375
Hand forgings
Al-Cu-Mg alloys
535
995
T4
190
375
T352(f)
175
350
00000000000000
10
T6
18
T81(d)
24
T87(d)
36
T62
18
T81(d)
18
T87(d)
18
T851(e)
36
T62
18
T851(e)
18
T81(d)
18
T8510(e)
18
36
26
26
18
2018
2024(h)
Die forgings
Flat sheet
510(g)
950(g)
T4
170
340
495
920
T3(d)
190
375
T361(d)
190
375
T42
190
375
190
375
2024(h)
Coiled sheet
495
920
T4
T42
190
375
190
375
Plate
495
920
Rolled or cold finished wire, rod, and bar
495
920
22
T351(e)
190
375
T361(d)
190
375
T42
190
375
T4
190
375
T351(e)
190
375
T36(d)
190
375
T42
190
375
Extruded rod, bar, shapes, and tube
495
920
T3
190
375
T3510(e)
190
375
T3511(e)
190
375
T42
190
375
Drawn tube
2036
Sheet
2038
Sheet
2218
Die forgings
495
920
T3(d)
T42
500
930
T4
540
1000
T4
205
400
510(g)
950(g)
T4
170
340
10
510(i)
950(i)
T41
240
460
02826:26282228622261 206
T8511(e)
T62
T6
T6
T852f)
T61
T81(d)
T861(d)
9
T62
T72
9
T62
T72
T851(e)
T861(d)
9
T62
T6
T851(e)
T86(d)
T62
T81
T8510(e)
T8511(e)
T62
T6
T61
T72
Al-Cu-Mg-Si alloys
2008
Sheet
510
950
T4(d)(j)
205
400
1
T62(e)
2014(h)
Flat sheet
500
935
T3(d)
160
320
18
T62
T42
160
320
18
T6
Coiled sheet
500
935
T4
160
320
18
T6
T42
160
320
18
T62
Plate
500
935
T42
160
320
18
T62
T451(e)
160
320
18
T651(e)
Rolled or cold finished wire, rod, and bar
500
935
T4
160(k)
320(k)
18
T6
T42
160(k)
320(k)
18
T451(e)
160(k)
320(k)
18
T62
T651(e)
Extruded rod, bar, shapes, and tube
500
935
T4
160(k)
320(k)
18
T6
T42
160(k)
320(k)
18
T4510(e)
160(k)
320(k)
18
T62
T6510(e)
Drawn tube
500
935
T4
160(k)
320(k)
18
T6
T42
160(k)
320(k)
18
T62
Die forgings
500(1)
935(1)
T4
170
340
10
T6
2017
Rolled or cold finished wire, rod, and bar
500
935
T4
T42
2117
Rolled or cold finished wire and rod
500
935
T4
T42
2618
Forgings and rolled rings
530
985
T4
200
390
20
T61
4032
Die forgings
510(h)
950(h)
T4
170
340
10
T6
Al-Cu-Li alloys
2090
2091
8090
CP276
Sheet
Sheet
Extruded bar
Extruded bar
540
1000
T3(d)
165
325
24
T83(d)
530
990
T3(d)
120
250
24
T84(d)
530
990
T3(d)
190
375
12
Peak aged(d)
530
990
T3(d)
190
Extruded bar
540
1000
T3(d)
190
375
375
12
Peak aged(d)
12-15
Peak aged(d)
(a) Material should be quenched from the solution-treating temperature as rapidly as possible and with minimum delay after removal from the furnace. When material is quenched by total immersion in water.
unless otherwise indicated, the water should be at room temperature, and should be suitably cooled so that it remains below 38 °C (100 °F) during the quenching cycle. Use of high-velocity, high-volume
jets of cold water also is effective for some materials. (b) The nominal temperatures listed should be attained as rapidly as possible and maintained within ±6 °C (±10 °F) of nominal during the time at
temperature. (c) Approximate time at temperature. The specific time will depend on the time required for the load to reach temperature. The times shown are based on rapid heating, with soak time measured
from the time the load reaches a temperature within 6 °C (10 °F) of the applicable temperature. (d) Cold working subsequent to solution heat treatment and prior to any precipitation heat treatment is necessary
to attain the specified properties for this temper. (e) Stress relieved by stretching to produce a specified amount of permanent set subsequent to solution heat treatment and prior to any precipitation heat
treatment. (f) Stress relieved by 1 to 5% cold reduction after solution treatment and prior to precipitation heat treatment. (g) Quenched in water at 100 °C (212 °F). (h) These heat treatments also apply to
alclad sheet and plate of these alloys. (i) Quenched with room-temperature air blast. (j) See U.S. Patent 4,840,852. (k) An alternative heat treatment of 8 h at 177 °C (350 °F) may also be used. (1) Quenched
in water at 60 to 80 °C (140 to 180 °F).
846 / Heat Treating of Nonferrous Alloys
Table 1(b) Typical solution and precipitation heat treatments for Mg-Si aluminum alloys (6xxx series alloys)
Solution heat treatment(a)
Precipitation heat treatment
Metal temperature(b)
Metal temperature(b)
Alloy
6005
Product form
°C
°F
Temper
designation
°C
°F
Time(c), h
Temper
designation
Extruded rod, bar, shapes,
530(d)
985(d)
TI
175
350
8
T5
and tube
6009(e)
Sheet
555
1030
T4
205
400
1
T6(e)
6010
Sheet
565
1050
T4
205
400
1
T6(e)
6053
Die forgings
520
970
T4
170
340
10
T6
6061(f)
Sheet
530
985
T4
160
320
18
T6
T42
160
320
18
T62
Plate
530
985
T4(g)
160
320
18
T6(g)
T42
160
320
18
T62
T451(h)
160
320
18
T651(h)
Rolled or cold finished
530
985
T4
160(i)
320(i)
18
T6
wire, rod, and bar
160(i)
320(i)
18
T89(j)
160(i)
320(i)
18
T93(k)
160(i)
320(i)
18
T913(k)
160(i)
320(i)
18
T94(k)
T42
160(i)
320(i)
18
T62
T451(h)
160(i)
320(i)
18
T651(h)
Extruded rod, bar, shapes,
and tube
530(d)
985(d)
T4
175
350
8
T6
T4510(h)
175
350
8
T4511(h)
175
350
8
T6510(h)
T6511(h)
530
985
T42
175
350
8
T62
6061(f)
Drawn tube
530
985
T4
160(i)
320(i)
18
T6
T42
160(i)
320(i)
18
T62
Die and hand forgings
530
985
T4
175
350
8
T6
Rolled rings
530
985
T4
175
350
8
T452(1)
175
350
8
T6
T652(1)
6063
Extruded rod, bar, shapes,
(d)
(d)
TI
205(m)
400(m)
1
T5
and tube
520(d)
970(d)
T4
175(n)
350(n)
520
970
T42
175(n)
350(n)
Drawn tube
520
970
T4
175
350
175
350
175
350
175
350
T42
175
350
6013(0)
Sheet
570
1055
W(p)
190
375
Plate
570
1055
W(p)
190
375
∞ ∞ ∞ ∞ ∞ ∞ ∞ ∞
8
T6
8
T62
8
T6
8
T83(j)(d)
8
T831(j)(d)
8
T832(j)(d)
8
T62
4
T6
4
T651
6066
Extruded rod, bar, shapes,
530
990
T4
175
350
8
T6
and tube
T42
Drawn tube
530
990
T4
T42
Die forgings
530
990
T4
6070
Extruded rod, bar, shapes,
545(d)
1015(d)
T4
and tube
T42
6111
Sheet
560
1040
T4
6151
Die forgings
515
960
T4
Rolled rings
515
960
T4
6262
Rolled or cold finished
540
1000
T4
wire, rod, and bar
T42
6262
Extruded rod, bar, shapes,
and tube
540(d)
1000(d)
T4
540
Drawn tube
540
1000
1000
T42
T4
T42
6463
Extruded rod, bar, shapes,
and tube
(d)
(d)
TI
520(d)
6951
Sheet
520
530
970(d)
970
985
T4
T42
E = FEE FEE FERRE PREFECE
175
350
T4510(h)
175
350
T4511(h)
175
350
175
350
175
350
175
350
ao ao ao ao ∞ ∞ ∞
8
T62
8
T6510(h)
8
T6511(h)
8
T6
8
T62
8
T6
160
320
18
T6
160
320
18
T62
175
350
8
T6(q)
170
340
10
T6
170
340
10
T6
T452(1)
170
340
10
T652(1)
170
340
8
T6
170
340
12
T451
170
340
170
340
175
350
12
2882
T9(k)
T651(h)
T62
T6
T4510(h)
175
350
12
175
350
12
T6510(h)
T62
170
340
8
T6
170
340
8
T9(k)
170
340
8
T62
205(m)
400(m)
1
T5
175(n)
350(n)
8
T6
175(n)
350(n)
8
T62
T4
160
320
18
T6
T42
160
320
18
T62
(a) Material should be quenched from the solution-treating temperature as rapidly as possible and with minimum delay after removal from the furnace. When material is quenched by total immersion in water,
unless otherwise indicated, the water should be at room temperature, and should be suitably cooled so that it remains below 38 °C (100 °F) during the quenching cycle. Use of high-velocity, high-volume
jets of cold water also is effective for some materials. (b) The nominal temperatures listed should be attained as rapidly as possible and maintained within ±6 °C (±10 °F) of nominal during the time at
temperature. (c) Approximate time at temperature. The specific time will depend on the time required for the load to reach temperature. The times shown are based on rapid heating, with soak time measured
from the time the load reaches a temperature within 6 °C (10 °F) of the applicable temperature. (d) By suitable control of extrusion temperature, product may be quenched directly from extrusion press to
provide specified properties for this temper. Some products may be adequately quenched in room-temperature air blast. (e) Alternate heat treatments of 4 h at 190 °C (375 °F) or 8 h at 175 °C (350 °F) may
also be used. See U.S. Patent 4,082,578. (f) These heat treatments also apply to alclad sheet and plate in these alloys. (g) Applicable to tread plate only. (h) Stress relieved by stretching to produce a specified
amount of permanent set prior to precipitation heat treatment. (i) An alternative heat treatment of 8 h at 170 °C (340 °F) also may be used. (j) Cold working after solution treatment is necessary to attain
specified properties during precipitation heat treatments. (k) Cold working after precipitation heat treatment is necessary to attain specified properties. (1) Stress relieved by 1 to 5% cold reduction subsequent
to solution heat treatment and prior to precipitation heat treatment. (m) An alternative treatment of 3 h at 182 °C (360 °F) also may be used. (n) An alternative treatment of 6 h at 182 °C (360 °F) also may
be used. (o) See U.S. Patent 4,589,932. (p) Two weeks of natural aging to a T4 condition. (q) Artificially aged in laboratory from T4 to T6.
Table 1(c)
Heat Treating of Aluminum Alloys / 847
Typical solution and precipitation heat treatments for heat-treatable Zn-Mg aluminum alloys from the 7xxx series
Solution heat treatment(a)
Precipitation heat treatment
Metal temperature(b)
Metal temperature(b)
Alloy
7001
Product form
°C
°F
Temper
designation
°C
°F
Time(c),
h
Temper
designation
Extruded rod, bar, shapes,
and tube
465
870
W
120
250
24
T6
120
250
24
T62
W510(d)
120
250
24
T6510(d)
W511(d)
120
250
24
T6511(d)
7005
Extruded rod, bar, and
:
T53(e)
shapes
7050
Plate
475
890
W51(d)
(f)
(f)
(f)
T7651(g)
(h)
(h)
(h)
T7451(g)
Extrusions
475
890
W510(d)
(f)
(f)
(f)
T76510(g)
W511(d)
(f)
(f)
(f)
T76511(g)
Die and hand forgings
475
890
W
(h)
(h)
(h)
W52(d)
(h)
(h)
(h)
T74(g)
T7452(g)
7075(i)
Sheet
480
900
W
120(j)
250(j)
24
T6
120(j)
250(j)
24
T62
(f)
(f)
(f)
T76(g)
(h)(k)
(h)(k)
(h)(k)
T73(g)
Plate
480
900
W
ޑޑ
120(j)
250(j)
24
T62
W51(d)
(h)(k)
(h)(k)
(h)(k)
120(j)
250(j)
24
(f)
(f)
(f)
T7351(d)(g)
T651(d)
T7651(g)
7075(i)
Rolled or cold finished
wire, rod, and bar
490
915
W
120
250
24
T6
120
250
24
(h)(k)
(h)(k)
(h)(k)
W51(d)
120
250
(h)(k)
(h)(k)
24
(h)(k)
T62
T73(g)
T651(d)
T7351(d)(g)
Extruded rod, bar, shapes,
and tube
465
870
W
120(1)
250(1)
24
T6
120(1)
250(1)
24
T62
(h)(k)
(h)(k)
(h)(k)
T73(g)
(f)
(f)
(f)
W510(d)
120(1)
250(1)
24
(h)(k)
(h)(k)
(h)(k)
T76(g)
T6510(d)
T73510(d)(g)
(f)
(f)
(f)
W511(d)
120(1)
250(1)
24
(h)(k)
(h)(k)
(h)(k)
(f)
(f)
(f)
T76510(g)
T6511(d)
T73511(d)(g)
T76511(g)
Drawn tube
465
870
=
W
120
250
24
T6
120
250
24
(h)(k)
(h)(k)
(h)(k)
T62
T73(g)
Die forgings
470(m)
880(h)
3
120
250
24
(h)
(h)
(h)
W52(n)
(h)
(h)
(h)
Hand forgings
470(m)
880(h)
W
120
250
24
(h)
(h)
(h)
W52(n)
120
250
24
(h)
(h)
(h)
Rolled rings
470
880
7175
Die forgings
(o)
(o)
(o)
(o)
(0)
Hand forgings
(o)
(0)
(0)
(o)
BBBBBB
W
120
250
24
T6
W
(o)
(0)
(o)
W
(o)
(o)
(o)
W52(n)
(o)
(o)
W
(0)
W52(n)
(0)
(0)
(o)
T6
T73(g)
T7352(n)(g)
T6
T73(g)
T652(n)
T7352(n)(g)
T66(0)
T74(g)(0)
T7452(n)(g)(0)
T74(g)(0)
T7452(n)(g)(0)
7475
Sheet
515(p)
960(p)
W
120
250
3
plus 155
315
3
(f)
(f)
(f)
T61(p)
T761(g)(p)
Plate
510(p)
950(p)
W51(d)
120
250
24
(f)
(f)
T651(p)
T7651(g)(p)
(h)
(h)
(h)
T7351(g)(p)
Alclad
7475
Sheet
495
920
W
120
250
3
plus 155
(f)
315
3
(f)
(f)
T61(p)
T761(g)(p)
(a) Material should be quenched from the solution-treating temperature as rapidly as possible and with minimum delay after removal from the furnace. When material is quenched by total immersion in water,
unless otherwise indicated, the water should be at room temperature, and should be suitably cooled so that it remains below 38 °C (100 °F) during the quenching cycle. Use of high-velocity, high-volume
jets of cold water also is effective for some materials. (b) The nominal temperatures listed should be attained as rapidly as possible and maintained within ±6 °C (±10 °F) of nominal during the time at
temperature. (c) Approximate time at temperature. The specific time will depend on the time required for the load to reach temperature. The times shown are based on rapid heating, with soak time measured
from the time the load reaches a temperature within 6 °C (10 °F) of the applicable temperature. (d) Stress relieved by stretching to produce a specified amount of permanent set after solution treatment and
prior to precipitation heat treatment. (e) No solution heat treatment; 72 h at room temperature following press quench, followed by two-stage precipitation heat treatment comprised of 8 h at 107 °C (225 °F)
plus 16 h at 149 °C (300 °F). (f) Aging practice varies with product, size, nature of equipment, loading procedures, and furnace-control capabilities. The optimum practice for a specific item can be ascertained
only by actual trial treatment of the item under specific conditions. Typical procedures involve a two-stage treatment comprised of 3 to 30 h at 121 °C (250 °F) followed by 15 to 18 h at 163 °C (325 °F) for
extrusions. An alternative two-stage treatment of 8 h at 99 °C (210 °F) followed by 24 to 28 h at 163 °C (325 °F) also may be used. (g) Aging of aluminum alloys 7050, 7075, 7175, and 7475 from any temper
to the T73 or T76 temper series requires closer-than-normal controls on aging variables such as time, temperature, heatup rate, and so forth, for any given item. In addition, when material in a T6-type temper
is reaged to a T73- or T76-type temper, the specific condition of the T6 material (such as property levels and other effects of processing variables) is extremely important and will affect the capability of the
reaged material to conform to the requirements specified for the applicable T73- or T76-type temper. (h) Two-stage treatment comprised of 6 to 8 h at 107 °C (225 °F) followed by: 24 to 30 h at 163 °C (325 °F)
for sheet and plate: 8 to 10 h at 177 °C (350 °F) for rolled or cold finished rod and bar; 6 to 8 h at 177 °C (350 °F) for extrusions and tube; 8 to 10 h at 177 °C (350 °F) for forgings in the T73 temper; and 6
to 8 h at 177 °C (350 °F) for forgings in the T7352 temper. (i) These heat treatments also apply to alclad sheet and plate of these alloys. (j) An alternative two-stage treatment comprised of 4 h at 96 °C (205 °F)
followed by 8 h at 157 °C (315 °F) also may be used. (k) For sheet, plate, tube, and extrusions, an alternative two-stage treatment comprised for 6 to 8 h at 107 °C (225 °F) followed by 14 to 18 h at 168 °C
(335 °F) may be used, provided that a heatup rate of approximately 14 °C/h (25 °F/h) is employed. For rolled or cold finished rod and bar, the alternative treatment is 10 h at 177 °C (350 °F). (I) An alternative
three-stage treatment comprised of 5 h at 99 °C (210 °F), 4 h at 121 °C (250 °F), and then 4 h at 149 °C (300 °F) may also be used. (m) Quenched in water at 60 to 80 °C (140 to 180 °F). (n) Stress relieved
by 1 to 5% cold reduction after solution treatment and prior to precipitation heat treatments. (o) 7175-T74 and -T7452 heat treatments are directed to specific results, may vary from supplier to supplier and
are either proprietary or patented. (p) Must be preceded by soak at 466 to 477 °C (870 to 890 °F). See U.S. Patent 3,791,880.
848 / Heat Treating of Nonferrous Alloys
Table 2 Soak times and maximum quench delays for solution treatment of wrought
aluminum alloys
See Table 1 for solution-treating temperatures.
Maximum
quench
delay, s
7
Soak time, minutes
Air furnace(b)
Salt bath(c)
Thickness(a), mm (in.)
min
max(d)
min
max(d)
≤0.41 (0.016)
20
25
10
15
5
0.51 (0.020)
20
30
10
20
7
0.64 (0.025)
25
35
15
7
0.81 (0.032)
25
35
15
1.02 (0.040)
30
40
20
30
10
1.27 (0.050)
30
40
20
30
10
1.35 (0.053)
30
40
20
30
10
1.80 (0.071)
35
45
25
10
2.03 (0.080)
35
45
25
10
35
45
25
10
40
55
30
15
40
55
30
15
50
60
35
15
50
60
35
15
55
65
35
15
65
75
45
55
15
+30
+20
30
+20
(e)
2.29 (0.090)
2.54 (0.100)
3.18 (0.125)
4.06 (0.160)
4.57 (0.180)
6.35 (0.250)
For each additional 12.7 (12) or fraction
Rivets (all)
>6.35 (0.250)-12.7 (0.500)
222222222222222328
+30
60
28339999972288678
22222222222222292
2822288333444445
5
(a) Minimum dimension of thickest section. (b) Soak time begins when all pyrometer instruments recover to original operating
temperature. (c) Soak time begins at time of immersion except when a heavy charge causes bath temperature to drop below specified
minimum, in which case soak time begins when bath regains minimum temperature. (d) Applicable to alclad materials only. (e) Increases
in thickness above 12.7 mm (12 in.) do not affect maximum quench delay, which remains constant at 15 s.
In the tabulation above, note especially the
effects of small increments of temperature,
within the normal range, on the properties
of 0.8 mm (0.032 in.) 2024-T4 sheet.
Solution-Treating Time. The time at the
nominal solution heat-treating temperature
(soak time) required to effect a satisfactory
degree of solution of the undissolved or
precipitated soluble phase constituents and
to achieve good homogeneity of the solid
solution is a function of microstructure be-
fore heat treatment. This time requirement
can vary from less than a minute for thin
sheet to as much as 20 h for large sand or
plaster-mold castings. Guideline informa-
tion for soak times required for wrought
products of various section thicknesses is
given in Table 2. Similar guidelines for
castings are presented in Table 3. The time
required to heat a load to the treatment
temperature in furnace heat treatment also
increases with section thickness and fur-
nace loading, and thus total cycle time
increases with these factors.
Soak time for alclad sheet and for parts
made from alclad sheet must be held to a
minimum, because excessive diffusion of
alloying elements from the core into the
cladding reduces corrosion protection. For
the same reason, reheat treatment of alclad
sheet less than 0.75 mm (0.030 in.) thick
generally is prohibited, and the number of
reheat treatments permitted for thicker al-
clad sheet is limited.
The soak times for wrought alloys take
into account the normal thermal lag be-
tween furnace and part and the difference
between surface and center temperatures
for commercial equipment qualified to the
standards of MIL-H-6088. The rapid heat-
ing rates of salt baths permit all immersion
•
time to be counted as soak time unless the
bath temperature drops below the minimum
of the range. Even then, soak time begins as
soon as the bath temperature returns to the
minimum. In air furnaces, soak time does
not begin until all furnace instruments re-
turn to their original set temperature-that
is, the temperature reading before insertion
of the load.
In air furnaces, thermocouples may also
be attracted to, or buried in, parts located in
the load in such a manner as to represent
the hottest and coldest temperatures in each
zone. In this way, it is possible to ensure
that adequate soaking is obtained.
Special consideration is given also to es-
tablishing soak times for hand and die forg-
ings; soak time in some specifications is
extended to complete solution and homog-
enization in areas that received marginal
reduction during forging. Considerable vari-
ation exists in the amount of soak time
added; some specifications call for an arbi-
trary addition, such as one hour, and others
require one hour per inch of thickness of the
original forging.
In air furnaces, careful attention should
be given to arrangement of the load. Air
flow and natural temperature distribution
within the furnace should be arranged to:
•
Offer minimum resistance to air flow
Produce the least disturbance in the natu-
ral temperature distribution
• Afford constant replenishment of the en-
velope of air around each part
It is common practice to specify a minimum
spacing of 50 mm (2 in.) between parts, but
large complex shapes may require consider-
ably greater spacing. Many operators have
found conservative loading practices to be
more economical in the long run than heavi-
er loading, because with lighter loads heat-
ing rates are higher and fewer rejections and
service failures are encountered.
High-Temperature Oxidation. There is a
condition, commonly but erroneously
known as HTO or high-temperature oxida-
tion, which can lead to deterioration of
properties in aluminum alloys. High-tem-
perature oxidation is a misnamed condition
of hydrogen diffusion that affects surface
layers during elevated-temperature treat-
ment. This condition can result from mois-
ture contamination in the furnace atmo-
sphere and is sometimes aggravated by
sulfur (as in heat-treatment furnaces also
used for magnesium alloy castings) or other
furnace refractory contamination.
Moisture in contact with aluminum at
high temperatures serves as a source of
nascent hydrogen, which diffuses into the
metal. Foreign materials, such as sulfur
compounds, function as decomposers of the
natural oxide surface film, eliminating it as a
barrier either between the moisture and the
aluminum or between the nascent hydrogen
and the aluminum. The most common man-
ifestation of high-temperature oxidation is
surface blistering, but occasionally the only
manifestations are internal discontinuities
or voids, which can be detected only by
careful ultrasonic inspection or by metallo-
graphic techniques.
It is important to recognize that the symp-
toms of high-temperature oxidation are
identical to those of unsoundness or high
gas content in the original ingot or of other
improper mill practice. Blisters resulting
from ingot defects, improper extrusion or
improper rolling may be lined up in the
direction of working. However, it usually is
impossible to distinguish among defect
sources, and therefore the possibility that a
contaminated atmosphere is the cause of
the defects must be checked.
Not all alloys and product forms are
equally vulnerable to this type of attack.
The 7xxx series alloys are most susceptible,
followed by the 2xxx alloys. Extrusions
undoubtedly are the most susceptible form;
forgings are probably second. Low-strength
alloys and alclad sheet and plate are rela-
tively immune to high-temperature oxida-
tion. (Blistering of alclad material as a result
of inadequate bonding is not the same as the
blistering caused by high-temperature oxi-
dation.)
If the protective oxide film formed during
mill operations is removed from the mill
product by a subsequent mechanical condi-
tioning operation such as sanding, the con-
ditioned surface will be more susceptible to
high-temperature oxidation than those from
which the film was not removed.
Moisture can be minimized by thoroughly
drying parts and racks before they are
charged. Drain holes often are needed in
racks of tubular construction to avoid en-
Table 3 Typical heat treatments for aluminum alloy sand and permanent mold castings
Solution heat treatment(b)
Heat Treating of Aluminum Alloys / 849
Aging treatment
Alloy
Temper
Type of
casting(a)
Temperature(c)
Temperature(c)
°C
°F
Time, h
°C
°F
Time, h
201.0(d)
T4
S or P
490-500(e)
+525-530
910-930(e)
+980-990
2
14-20
Minimum of 5 days at room temperature
T6
S
T7
S
510-515(e)
+525-530
510-515(e)
950-960(e)
+980-990
2
14-20
155
+525-530
950-960(e)
+980-990
2
14-20
190
310
370
20
5
T43(f)
525
980
20
T71
490-500(e)
910-930(e)
2
+525-530
+980-990
14-20
200
204.0(d)
T4
S or P
530
985
12
24 h at room temperature + ½ to 1 h at 160 °C
390
Minimum of 5 days at room temperature
4
T4
S or P
520
970
10
T6(g)
S or P
530
985
12
(g)
(g)
206.0(d)
T4
S or P
490-500(e)
+525-530
910-930(e)
+980-990
2
14-20
Minimum of 5 days at room temperature
T6
S or P
490-500(e)
+525-530
910-930(e)
+980-990
2
14-20
155
310
12-24
T7
S or P
490-500(e)
910-930(e)
2
+525-530
+980-990
14-20
200
390
4
T72
S or P
490-500(e)
910-930(e)
2
+525-530
+980-990
14-20
243-248
470-480
208.0
T55
222.0
O(h)
T61
T551
SSSP
155
310
16
315
600
3
510
950
12
155
310
11
170
340
16-22
T65
510
950
4-12
170
340
7-9
242.0
O(i)
345
650
3
T571
205
400
8
165-170
330-340
22-26
T77
S
515
960
5(j)
330-355
625-675
2 (minimum)
T61
S or P
515
960
4-12(j)
205-230
400-450
3-5
295.0
T4
515
960
T6
515
960
T62
515
960
T7
515
960
296.0
T4
510
950
T6
510
950
T7
510
950
319.0
T5
T6
505
940
505
940
328.0
T6
515
960
332.0
T5
333.0
T5
P
T6
P
505
950
T7
P
505
940
336.0
T551
P
T65
P
515
960
354.0
(k)
525-535
980-995
355.0
T51
S or P
T6
S
525
980
P
525
980
T62
525
980
T7
525
980
525
980
T71
525
980
525
980
C355.0
T6
525
980
T61
525
980
£sts£stts:E;g;sEs:ཨEsES
12
12
155
310
3-6
12
155
310
12-24
12
260
500
46
8
...
8
155
310
1-8
8
260
500
46
205
400
8
12
155
310
2-5
4-12
155
310
2-5
12
155
310
2-5
205
400
7-9
205
400
7-9
6-12
155
310
2-5
6-12
260
500
46
205
400
7-9
8
205
400
7-9
10-12
(h)
(h)
(1)
225
440
7-9
12
155
310
3-5
4-12
155
310
2-5
4-12
170
340
14-18
12
225
440
3-5
4-12
225
440
3-9
12
245
475
46
4-12
245
475
3-6
12
155
310
3-5
6-12
Room temperature
8 (minimum)
155
310
10-12
356.0
T51
T6
T7
T71
A356.0
T6
T61
SSPSPSesa
S or P
225
440
7-9
540
1000
12
155
310
3-5
540
1000
4-12
155
310
2-5
540
1000
12
205
400
3-5
540
1000
4-12
225
440
7-9
540
1000
10-12
245
475
3
540
1000
4-12
245
475
3-6
540
1000
12
155
310
3-5
540
1000
6-12
Room temperature
155
310
8 (minimum)
6-12
(continued)
(a) S, sand; P, permanent mold. (b) Unless otherwise indicated, solution treating is followed by quenching in water at 65-100 °C (150-212 °F). (c) Except where ranges are given, listed temperatures are ±6 °C
or ±10 °F. (d) Casting wall thickness, solidification rate, and grain refinement affect the solution heat-treatment cycle in alloys 201.0, 204.0, and 206.0, and care must be taken in approaching the final solution
temperature. Too rapid an approach can result in the occurrence of incipient melting. (e) For castings with thick or other slowly solidified sections, a pre-solution heat treatment ranging from about 490 to
515 °C (910 to 960 °F) may be needed to avoid too rapid a temperature rise to the solution temperature and the melting of CuAl2. (f) Temper T43 for 201.0 was developed for improved impact resistance with
some decrease in other mechanical properties. Typical Charpy value is 20 J (15 ft lb). (g) The French precipitation treatment technology for the heat treatment of 204.0 alloy requires 12 h at temperature.
The aging temperatures of 140, 160, or 180 °C (285, 320, or 355 °F) are selected to meet the required combination of properties. (h) Stress relieve for dimensional stability as follows: hold 5 h at 413 ± 14 °C
(775 ± 25 °F): furnace cool to 345 °C (650 °F) over a period of 2 h or more: furnace cool to 230 °C (450 °F) over a period of not more than 2 h; furnace cool to 120 °C (250 °F) over a period of approximately
2 h; cool to room temperature in still air outside the furnace. (i) No quench required: cool in still air outside the furnace. (j) Air-blast quench from solution-treating temperature. (k) Casting process varies
(sand, permanent mold, or composite) depending on desired mechanical properties. (1) Solution heat treat as indicated, then artificially age by heating uniformly at the temperature and for the time necessary
to develop the desired mechanical properties. (m) Quench in water at 65-100 °C (150-212 °F) for 10-20 s only. (n) Cool to room temperature in still air outside the furnace.
850 / Heat Treating of Nonferrous Alloys
Table 3 (continued)
Solution heat treatment(b)
Aging treatment
Temperature(c)
Temperature(c)
Alloy
Temper
Type of
casting(a)
°C
°F
Time, h
°C
°F
Time, h
357.0
T6
P
T61
S
A357.0
359.0
A444.0
T4
P
520.0
T4
S
535.0
T5(h)
705.0
T5
S
andzannn
540
1000
8
175
350
6
540
1000
10-12
155
310
10-12
(k)
540
1000
8-12
(h)
(h)
(h)
(k)
540
1000
10-14
(h)
(h)
(h)
540
1000
8-12
430
810
18(m)
400
750
5
Room temperature
21 days
100
210
8
P
:
Room temperature
21 days
100
210
10
707.0
T5
SP
155
310
3-5
Р
Room temperature, or
21 days
100
210
8
T7
710.0
711.0
RE
T5
T1
712.0
T5
SPSPS
530
530
990
990
8-16
175
350
4-10
4-8
175
350
4-10
Room temperature
21 days
Room temperature
21 days
713.0
T5
S or P
Room temperature, or
155
Room temperature, or
120
21 days
315
6-8
21 days
250
16
771.0
T53(h)
T5
T51
T52
T6
T71
850.0
T5
851.0
T5
SSSSSSSS
415(n)
775(n)
5(n)
180(n)
360(n)
4(n)
180(n)
355(n)
3-5(n)
205
405
6
(h)
(h)
(h)
590(n)
1090(n)
6(n)
130
265
3
590(i)
1090(i)
6(i)
140
285
15
S or P
220
430
7-9
T6
S or P
P
220
430
7-9
480
900
6
220
430
4
852.0
T5
S or P
220
430
7-9
(a) S, sand; P, permanent mold. (b) Unless otherwise indicated, solution treating is followed by quenching in water at 65-100 °C (150-212 °F). (c) Except where ranges are given, listed temperatures are ±6 °C
or ±10 °F. (d) Casting wall thickness, solidification rate, and grain refinement affect the solution heat-treatment cycle in alloys 201.0, 204.0, and 206.0, and care must be taken in approaching the final solution
temperature. Too rapid an approach can result in the occurrence of incipient melting. (e) For castings with thick or other slowly solidified sections, a pre-solution heat treatment ranging from about 490 to
515 °C (910 to 960 °F) may be needed to avoid too rapid a temperature rise to the solution temperature and the melting of CuAl2. (f) Temper T43 for 201.0 was developed for improved impact resistance with
some decrease in other mechanical properties. Typical Charpy value is 20 J (15 ft lb). (g) The French precipitation treatment technology for the heat treatment of 204.0 alloy requires 12 h at temperature.
The aging temperatures of 140, 160, or 180 °C (285, 320, or 355 °F) are selected to meet the required combination of properties. (h) Stress relieve for dimensional stability as follows: hold 5 h at 413 ± 14 °C
(775 +25 °F); furnace cool to 345 °C (650 °F) over a period of 2 h or more; furnace cool to 230 °C (450 °F) over a period of not more than 2 h; furnace cool to 120 °C (250 °F) over a period of approximately
2 h; cool to room temperature in still air outside the furnace. (i) No quench required; cool in still air outside the furnace. (j) Air-blast quench from solution-treating temperature. (k) Casting process varies
(sand, permanent mold, or composite) depending on desired mechanical properties. (j) Solution heat treat as indicated, then artificially age by heating uniformly at the temperature and for the time necessary
to develop the desired mechanical properties. (m) Quench in water at 65-100 °C (150-212 °F) for 10-20 s only. (n) Cool to room temperature in still air outside the furnace.
trapment of water. Another common re-
quirement is adjustment of the position of
the quench tank with respect to furnace
doors and air intake. Because it is unlikely
that all moisture can be eliminated from the
atmosphere in a production heat-treating
furnace, it is extremely important to elimi-
nate all traces of other contaminants from
both the parts and the furnace atmosphere.
The most virulent contaminants in attack-
ing aluminum are sulfur compounds. Resi-
dues from forming or machining lubricants, or
from a sulfur dioxide protective atmosphere
used in prior heat treatment of magnesium,
are potential sources of sulfur contamination.
In one plant, surface contamination resulted
from sulfur-containing materials in tote boxes
used to transport parts. In another, an epi-
demic of blistering was cured by rectifying a
"sour" degreaser. In a third instance, it was
found that a vapor-degreasing operation was
not completely removing a thin, hard waxy
residue, and an alkaline cleaning operation
was added.
Very often, the source of contamination
is obscure and difficult to detect, and the
problem must be combated in another way.
The most common of the alternative meth-
ods is use of a protective fluoborate com-
pound in the furnace. Such a compound
usually is effective in minimizing the harm-
ful effects of moisture and other undesirable
contaminants because it forms a barrier
layer or film on the aluminum surface. The
additive is not a universal solution; in some
applications, high-temperature oxidation
has occurred even though a fluoborate com-
pound was employed. Also, the use of such
compounds, particularly ammonium fluo-
borate, may present a hazard to personnel if
used in poorly sealed furnaces or in furnac-
es that discharge their atmospheres into
enclosed areas.
Protective fluoborate compounds accen-
tuate staining or darkening of the parts
being treated. (At times, this attack, partic-
ularly on parts located near the protective-
compound container during heat treatment,
has been severe enough to be termed "cor-
rosion.") Although this minor nuisance
might be considered a small price to pay for
solution of a problem of high-temperature
oxidation, the residual compound in the
furnace dissipates slowly. Therefore, subse-
quent loads of alloys and product forms
whose end uses require bright surfaces, and
that are not susceptible to high-temperature
oxidation, may be detrimentally affected.
Successful use of fluoborate protective
compounds appears to depend on specify-
ing the right amount for each furnace; this
must be established on a trial-and-error
basis. One aircraft manufacturer adds 4
g/m³ (0.004 oz/ft³) of furnace chamber to
each load. Another adds 0.45 kg (1 lb) per
shift to a metal container hung on the fur-
nace chamber wall, thus avoiding loss of the
compound during quenching.
A second method of combating high-tem-
perature oxidation is to anodize the work
before it is heat treated. The resultant alu-
minum oxide film prevents attack by con-
taminants in the furnace atmosphere. The
only deterrents to the use of anodizing are
its cost (in money and time) and the slight
surface frostiness which results from the
subsequent stripping operation.
The usual objection to the blistered sur-
face produced by high-temperature oxida-
tion is its unsightly appearance. This often
can be improved (for salvage purposes) by
applying local pressure to flatten each blis-
ter and then finishing by a mechanical pro-
cess such as polishing, buffing, sanding, or

abrasive blasting. In general, the effect of
HTO on static properties and fatigue
strength is slight. However, if a void result-
ing from HTO is located close to another
stress concentration, such as a hole, much
greater degradation of fatigue strength is
likely. In critical aluminum alloy forgings,
any blistering must be evaluated carefully
for its effect on the integrity of the part. Any
"cosmetic" salvage should be performed
only after it has been established that the
blisters are superficial and will not remain in
the finished product.
Precipitation Heat Treating without Prior
Solution Heat Treatment. Certain alloys that
are relatively insensitive to cooling rate
during quenching can be either air cooled or
water quenched directly from a final hot-
working operation. In either condition,
these alloys respond strongly to precipita-
tion heat treatment. This practice is widely
used in producing thin extruded shapes of
alloys 6061, 6063, 6463, and 7005. Upon
precipitation heat treating after quenching
at the extrusion press, these alloys develop
strengths nearly equal to those obtained by
adding a separate solution heat treating op-
eration. Changes in properties occurring
during the precipitation treatment follow
the principles outlined in the discussion of
solution heat-treated alloys.
Quenching
Quenching is in many ways the most
critical step in the sequence of heat-treating
operations. The objective of quenching is to
preserve the solid solution formed at the
solution heat-treating temperature, by rap-
idly cooling to some lower temperature,
usually near room temperature. From the
preceding general discussion, this statement
applies not only to retaining solute atoms in
solution, but also to maintaining a certain
minimum number of vacant lattice sites to
assist in promoting the low-temperature dif-
fusion required for zone formation. The
solute atoms that precipitate either on grain
boundaries, dispersoids, or other particles,
as well as the vacancies that migrate (with
extreme rapidity) to disordered regions, are
irretrievably lost for practical purposes and
fail to contribute to the subsequent
strengthening.
In most instances, to avoid those types of
precipitation that are detrimental to me-
chanical properties or to corrosion resis-
tance, the solid solution formed during so-
lution heat treatment must be quenched
rapidly enough (and without interruption) to
produce a supersaturated solution at room
temperature the optimum condition for
precipitation hardening. The resistance to
stress-corrosion cracking of certain copper-
free aluminum-zinc-magnesium alloys,
however, is improved by slow quenching.
Most frequently, parts are quenched by
immersion in cold water or, in continuous
heat treating of sheet, plate, or extrusions in
primary fabricating mills, by progressive
flooding or high-velocity spraying with cold
water. However, parts of complex shape,
often with both thin and thick sections (such
as die forgings, most castings, impact extru-
sions, and components formed from sheet)
are commonly quenched in a medium that
provides somewhat slower cooling. This
medium may be water at 65 to 80 °C (150 to
180 °F), boiling water, an aqueous solution
of polyalkylene glycol, or some other fluid
medium such as forced air or mist.
If appreciable precipitation during cool-
ing is to be avoided, two requirements must
be satisfied. First, the time required for
transfer of the load from the furnace to the
quenching medium must be short enough to
preclude slow precooling into the tempera-
ture range where very rapid precipitation
takes place. For alloy 7075, this range was
determined to be 400 to 290 °C (750 to 550
°F), and some sources quote this range (or a
slightly different range) as the most critical
range for quenching of any aluminum alloy.
Later work has shown that the most critical
range is alloy-dependent, and as will be
discussed in detail under "Quench-Factor
Analysis,” significant errors can result from
the assumption that precipitation is negligi-
ble outside of a so-called "critical range.'
The second requirement for avoidance of
appreciable precipitation during quenching
is that the volume, heat-absorption capaci-
ty, and rate of flow of the quenching medi-
um be such that little or no precipitation
occurs during cooling. Any interruption of
the quench that might allow reheating into a
temperature range where rapid precipitation
can occur must be prohibited.
For maximum dimensional stability,
some forgings and castings are fan cooled or
still-air cooled. In such instances, precipita-
tion-hardening response is limited, but sat-
isfactory values of strength and hardness
are obtained. Extrusions produced without
separate solution heat treatment can be air
or mist quenched, but thicker sections may
require water quenching by immersion or
spraying. Alloys that are relatively dilute,
such as 6063 and 7005, are particularly well
suited to air quenching, and their mechani-
cal properties are not greatly affected by its
low cooling rate. Lower quenching rates are
also employed for forgings, castings, and
complex shapes to minimize warpage or
other distortion and the magnitude of resid-
ual stresses developed as a consequence of
temperature nonuniformity from surface to
interior.
Effect of Quench Rate on Properties. As a
broad generalization, the highest strengths
attainable and the best combinations of
strength and toughness are those associated
with the most rapid quenching rates. Resis-
tance to corrosion and stress-corrosion
cracking are other characteristics that are
generally improved by maximum rapidity of
quenching. Some of the alloys used in arti-
Heat Treating of Aluminum Alloys / 851
ficially aged tempers, and in particular the
copper-free 7xxx alloys, are exceptions to
this rule. The effect of quench rate on
mechanical properties may also depend on
the desired temper. In the underaged con-
dition, for example, a slow quench rate is
more detrimental on ductility and fracture
toughness. Strength would be more affected
after near-to-peak aging.
Because of these effects, much work has
been done over the years to understand and
predict how quenching conditions and prod-
uct form influence properties. The relative
effects of quench methods can be compared
in terms of average quench rates. In Fig 3,
for example, the effects of quenching on the
yield strength of four alloys are compared in
terms of average quenching rates through
the range from 400 to 290 °C (750 to 550 °F).
For alloys relatively high in sensitivity to
quenching rate, such as 7075, rates of about
300 °C/s (540 °F/s) or higher are required in
order to obtain near-maximum strength af-
ter precipitation heat treatment. The other
alloys in Fig 3 maintain their strengths at
cooling rates as low as about 100 °C/s (180
°F/s). Similar comparisons in terms of aver-
age quench rates are shown in Tables 4 and
5.
Average quench rates are useful in com-
paring experimental results from various
quench methods. In Table 4, for example, a
severe reduction in strength occurred at the
average quench rate of 36 °C/s (65 °F/s).
However, average quench rates only com-
pare results in a "critical" temperature
range, where precipitation is most likely to
occur. This method is not entirely accurate,
because significant precipitation can also
occur outside the specified critical temper-
ature range of average quench rates. More-
over, for high-strength alloys, toughness
and corrosion resistance may be impaired
without significant loss of tensile strength.
Therefore, a more sophisticated compar-
ison, known as quench-factor analysis, is
needed for quantitative property prediction
or property optimization. Quench-factor
analysis, as discussed in a later section, is
useful when cooling rates are nonuniform.
Delay in Quenching. Whether the transfer
of parts from the furnace to the quench is
performed manually or mechanically, it
must be completed in less than the specified
maximum time. The maximum allowable
transfer time or “quench delay" varies with
the temperature and velocity of the ambient
air and the mass and emissivity of the parts.
From cooling curves such as those illustrat-
ed in Fig 4, maximum quench delays (see
table accompanying Fig 4) can be deter-
mined that will ensure complete immersion
before the parts cool below 400 °C (750 °F).
MIL-H-6088 specifies maximum quench de-
lays for high-strength alloys of 5, 7, 10, and
15 s for thickness ranges of up to 0.016 in.
(0.41 mm), 0.017 to 0.031 in. (0.43 to 0.79
mm), 0.032 to 0.090 in. (0.81 to 2.29 mm),
852 / Heat Treating of Nonferrous Alloys
Yield strength, MPa
(a)
Fraction of WQ yield strength, %
(c)
Average quenching rate
from 750-550 °F, °F/s
10
600
500 -7050-T736
400
300
200
100
102
103
104
7050-T6
80
60
7050-T73
キーキ
40
6061-T6
-2024-T4
20
1
10
102
103
104
Average quenching rate
from 400-290 °C, °C/s
100
Furnace cooling
80
60
50
40
20
20
20
Yield strength, ksi
Tensile strength, MPa
(b)
Average cooling rate
from 750-550 °F, °F/s
10
102 103 104 105
700
100
600
500
7178-T6
7075-T6
7050-T73
80
7075-T73
2014-T6
2024-T4
6070-T6
1
6061-T6
60
40
400
300
200
1
10
Average cooling rate
from 400-290 °C, °C/s
102 103 104 105
Air cooled
Forced air cooling
Alloy and condition (Source: Ref 1)
O 8090, peak aged
⚫ 2090, peak aged
▲ 7150, aged 24 h at 120 °C
▲ 7475, aged 24 h at 120 °C
нн
0
0.01
0.1
1
10
Average cooling rate, °C/s
WQ
Tensile strength, ksi
Fig 3 Quench sensitivity of various aluminum alloys as a function of average quench rates. (a) Yield strength
after aging of four wrought alloys. (b) Tensile strength after aging of eight wrought alloys. (c) Relative
quench sensitivity of two aluminum-lithium alloys (2090 and 8090, both solution treated for 1 h at 520 °C, or 970
°F) and two Zn-Mg-Cu aluminum alloys (7150 and 7475, both solution treated for 40 min at 480 °C, or 895 °F)
and over 0.090 in., respectively. Quench de-
lay is conservatively defined as commencing
"when the furnace door begins to open or the
first corner of a load emerges from a salt
bath" and ending "when the last corner of the
load is immersed in the water quench tank."
Recommended maximum quench-delay times
are listed in Table 2. However, exceeding the
maximum delay time is permitted if tempera-
ture measurements of the load prove that all
parts are above 415 °C (775 °F) when
quenched. The C-curves used in quench-fac-
tor analysis can also assist in determining a
maximum allowable delay.
It is relatively easy to control quench
delay in day-to-day operations by using a
Table 4 Effect of average quench rate on tensile properties of aluminum-lithium alloy
2090
Average quench rate at center of plate
0.5 °C/s (13 mm plate, air
cooled)
36 °C/s (38 mm plate, quenched
in room-temperature water)
46 °C/s (13 mm plate, quenched
in boiling water)
48 °C/s (13 mm plate, quenched
in room-temperature water)
85 °C/s (13 mm plate, quenched
in ice brine)
(a) Data are averages from 4 specimens.
Elongation(a), %
Yield
strength(a)
Tensile
strength(a)
Condition
MPa
ksi
MPa
ksi
As-quenched
162
334
2
6% stretch + aged
448
513
5
8 h at 190 °C
As-quenched
128
312
12
6% stretch + aged
338
476
6
8 h at 190 °C
As-quenched
138
331
16
6% stretch aged
530
570
9
8 h at 190 °C
As-quenched
139
331
17
6% stretch aged
526
570
7
8 h at 190 °C
As-quenched
135
349
19
6% stretch aged
8 h at 190 °C
535
575
7
stopwatch or, if necessary, by attaching
thermocouples to parts. However, although
the cooling rate between 400 and 260 °C (750
and 500 °F) is most critical and must be
extremely high for many high-strength al-
loys, it cannot be directly measured in pro-
duction operations. It is usual to rely on
standardized practices, augmented by re-
sults of tension tests and tests of suscepti-
bility to intergranular corrosion.
Water-immersion quenching normally is
controlled in practice by stipulating maxi-
mum quench-delay time and maximum wa-
ter temperature. The first requirement con-
trols the cooling rate during transfer and,
for high-strength alloys, often is based on
the criterion of complete immersion before
the metal cools below 415 °C (775 °F). This
specification of 415 °C (775 °F) is based on a
critical temperature for alloy 7075, which
has one of the more severe C-curves (Fig 5).
Therefore, the criterion for complete im-
mersion of other alloys might be based on a
temperature lower than the 415 °C (775 °F)
specification, depending on the characteris-
tics of the particular C-curve.
The second requirement controls the
cooling rate during immersion. MIL-H-6088
specifies that for water-immersion quench-
ing, except quenching of forgings and cast-
ings, the temperature of the water shall not
exceed 38 °C (100 °F) upon completion of
quenching. This requirement controls both
the temperature of the quench water prior
to immersion and the ratio of the combined
mass of load and rack to the volume of
water. However, to ensure adequate
quenching effectiveness, it is necessary also
that the cooling fluid flow past all surfaces
of each part during the first few seconds
after immersion. Before parts enter the fur-
nace, their placement in racks or baskets
should be compatible with this requirement.
During the first few seconds of quenching,
agitation of the parts or the water should be
sufficient to prevent local increases in tem-
perature due to the formation of steam
pockets.
In one application, it was found that
2024-T4 plates 13 by 760 by 760 mm (1/2 by
30 by 30 in.), quenched singly into a large
volume of still water, were quite susceptible
to intergranular corrosion. This susceptibil-
ity disappeared completely when the
quenching practice was modified by adding
sufficient agitation to break up the insulat-
ing blanket of steam that formed on the
surface of the hot metal. Quenching prac-
tices for small parts such as fasteners and
hydraulic fittings have been modified for the
same reason. Dumping in bulk from baskets
has been replaced by methods, such as the
use of shaker hearth furnaces or special
racking, which permit parts to be quenched
singly.
Spray Quenching. For spray quenching,
the quench rate is controlled by the velocity
of the water and by volume of water per unit
Heat Treating of Aluminum Alloys / 853
Table 5 The effect of quench rate on the mechanical properties of age-hardened aluminum-lithium alloy 8090
Alloy composition
Al-2.28Li-0.86Cu-
0.90Mg-0.13Zr-
0.13Fe-0.06Si
Cooling from
solution treatment(a)
Yield strength(b)
Ultimate tensile
strength(b)
Stretch, %
Aging treatment
MPa
ksi
MPa
ksi
Elongation in 50 mm
(2 in.)(b), %
Air cool
2
190 °C for 16 h
380
55
446
64.5
7.7
(-0.25 °C/s)
4
170 °C for 24 h
401
Polymer quench
(-18 °C/s)
2
190 °C for 16 h
415
4
170 °C for 24 h
415
Water quench
2
190 °C for 16 h
428
(~120 °C)
4
170 °C for 24 h
417
Al-2.58Li-1.36Cu-
0.89Mg-0.13Zr-
0.17Fe-0.04Si
Air cool
2
190 °C for 16 h
417
88 28 80
58
465
67.5
6.0
60
481
70
8.0
60
481
70
7.2
62
492
71.4
8.1
60
483
70
7.5
60
485
70.3
6.5
(-0.25 °C/s)
4
170 °C for 24 h
442
Polymer quench
(-18 °C/s)
2
190 °C for 16 h
448
170 °C for 24 h
448
Water quench
(~120 °C/s)
2
190 °C for 16 h
464
170 °C for 24 h
448
2 22 28
64
503
73
4.5
65
524
76
65
519
75
5.0
67
535
77.5
23 24
6.8
8.2
65
517
75
6.3
(a) Solution treatment of 550 °C (1020 °F) for 1 h. (b) Data are averages from two specimens.
area per unit time of impingement of the
water on the workpiece. Rate of travel of
the workpiece through the sprays is an
important variable.
Local increases in temperature that occur
within the first few seconds of quenching,
caused by a phenomenon such as plugged
spray nozzles, are particularly deleterious.
The remaining “internal heat” may be suf-
ficient to reheat the surface region. When
this happens, a large loss in strength occurs
at the previously quenched surface. The
loss of strength in the affected area of a
heavy part is much more severe than that
caused by an inadequate quenching rate
alone. This is illustrated for 75 mm (3 in.)
thick 7075-T62 plate in Fig 6, which com-
pares, at various depths, the properties of a
plate for which quenching was interrupted
on one side after 3 s with those of a plate
that was quenched from one side only.
Quench Severity and Quenchant Selec-
tion. Quench severity is commonly ex-
pressed in terms of an H-value (or Gross-
mann number), where the H-value is related
to the thermal conductivity (k) of the part(s)
and the coefficient of heat transfer (C) be-
tween the quenchant and the part. These
quantities are related by the equation H =
C/2k, where the coefficient of heat transfer
(C) is affected by the quenchant velocity at
the surface of the part and several inherent
characteristics of the quenchant (such as
quenchant boiling point, viscosity, density,
thermal conductivity, and specific heat).
Water, which is the most widely used and
effective quenching medium, can obtain
cooling rates up to about 200 °C/s (400 °F/s)
at the midplane of 25 mm (1 in.) thick
aluminum alloy plate (see the dashed line in
Fig 7). No rates higher than those defined
by this line have been observed, although
rates approaching them were measured with
impinging spray quenches. Lower cooling
rates are achieved by immersion in heated
water (Fig 7) or by reducing the velocity of
the quenchant around the part (Table 6).
Cooling rates can also be reduced by low-
ering surface tension or by increasing the
stability of the vapor film around the part.
Polymer quenchants, which retard cool-
ing rates by the formation of films around
the part, are compared with water in Table
6. The effective film coefficient is essential-
ly the heat transfer coefficient (C), which is
related to the Grossmann number (H). The
application of polymer quenchants is cov-
ered in AMS specifications 3025 and 2770,
although many aluminum and aerospace
companies have developed internal specifi-
cations that differ from AMS-2770. Typical
parameters for quenching wrought products
(other than forgings) in glycol-water solu-
tions are presented in Table 7.
Other Factors Affecting Quench Rate.
Quenching rates are very sensitive to the
surface condition of the parts. Lowest rates
are observed with products having freshly
machined or bright-etched, clean surfaces,
or products that have been coated with
materials that decrease heat transfer. The
presence of oxide films or stains increases
cooling rates. Further marked changes can
be effected through the application of non-
reflective coatings, which also accelerate
heating (Fig 8). Surface roughness exerts a
similar effect; this appears related to vapor
film stability. The manner in which complex
products, such as engineered castings and
die forgings, enter the quenching medium
can significantly alter the relative cooling
Temperature, °C
Time per unit thickness, s/in
0
600
1000
2000
3000
1000
800
400
Alclad
600
200
Nonclad
400
Temperature, °F
mm
0.41
0.51
0.64
0.81
200
1.02
0
0
50
100
150
Thickness
Maximum quench delay, s
in.
Alclad
Nonclad
0.016
6.4
4.4
0.020
8.0
5.5
0.025
10.0
6.8
0.032
12.8
8.8
0.040
20.0
11.0
Time per unit thickness, s/mm
Fig 4 Cooling curves for alclad and nonclad aluminum products cooled from 495 °C (920 °F) in forced air. Air temperature, 25 °C (80 °F); air velocity, 2.3 m/s (450 ft/min).
shown.
Tabulated values of quench delay (maximum delay before the material being quenched has cooled below 400 °C, or 750 °F) were determined from cooling curves
854 / Heat Treating of Nonferrous Alloys
Temperature, °C
Fig 5
600
500
400
A
300
200
A: 7075
B: 2017
100
C: 6061
D: 6063
0
1
10
100
Time, s
103
1110
930
750
570
390
212
32
Temperature, °F
Time-temperature-property curves at 95% of maximum tensile stress for various alloys. See the section
"Quench-Factor Analysis" for discussion. Source: Ref 2
rates at various points, thereby affecting
mechanical properties and residual stresses
established during quenching. Similarly,
quenching complex extruded shapes whose
wall thicknesses differ widely poses special
problems if distortion and stresses are to be
minimized. In batch heat-treating opera-
tions, placement and spacing of parts on the
racks can be a major factor in determining
the quenching rates. In immersion quench-
ing, adequate volumes of the quenching
medium must be provided to prevent an
excessive temperature rise in the medium.
When jet agitation is used to induce water
Table 6 Grossmann numbers and heat transfer coefficients (C) of quenchant-to-part films
Quenchant
Effective film heat transfer
coefficient (C)
Btu/ft2 h °F
2460
3105
Temperature
Velocity
Туре
°C
°F
m/s
ft/min
Grossmann
Number
(H = C/2k)
W/cm². K
Water
27
80
0.00
0
1.07
3.55
0.25
50
1.35
4.78
0.50
100
1.55
5.14
3565
Water
38
100
0.00
0
0.99
3.28
2275
0.25
50
1.21
4.01
2785
0.50
100
1.48
4.91
3400
Water
49
120
0.00
0
1.10
3.65
2530
0.25
50
1.29
4.29
2970
0.50
100
1.60
5.31
3680
Water
60
140
0.00
0
0.86
2.85
1980
0.25
50
1.09
3.62
2510
0.50
100
1.33
4.41
3060
Water
71
160
0.00
0
0.21
0.70
485
0.25
50
0.57
1.89
1310
0.50
100
0.79
2.62
1815
0.00
0
0.11
0.36
255
0.25
50
0.21
0.69
485
0.50
100
0.27
0.89
620
Water
93
200
0.00
0
0.06
0.20
138
0.25
50
0.08
0.27
0.50
100
0.09
0.30
207
100
212
0.00
0.04
0.13
0.25
50
0.04
0.13
92
0.50
100
0.04
0.13
92
Polyalkylene glycol
(UCON A)(a)
30
85
0.00
0.19
0.63
429
0.25
50
0.21
0.70
475
0.50
100
0.23
0.77
529
Polyvinyl
pyrrolidone
(PVP90)(a)
30
85
0.00
0
0.44
1.49
0.25
50
0.40
1.34
0.50
100
0.42
1.41
Water
82
180
Water
80
88088
088
184
92
1012
912
966
(a) Polymer quenchants with concentrations of 25%. K is equal to the thermal conductivity of the aluminum alloy (7075). Source: Ref 4
flow between parts, jets should not impinge
directly and cause rapid localized cooling.
Quenching to Minimize Residual Stress
and Warpage. Although cold-water immer-
sion or flushing is most common, because it
produces the most effective quench (and
has been required by MIL-H-6088 for 2014,
2017, 2024, 2117, 7075, and 7178 alloys
except forgings), it presents problems in-
volving residual stress and warpage.
Residual stresses in heavy sections of
aluminum alloys originate from differential
thermal expansion during quenching ―that
is, the still-warm central material contracts,
pulling in the already cooled outer shell.
The magnitude of stresses increases with
section size, as shown in Fig 9.
The distribution pattern of residual
stresses in as-quenched parts (compression
in the outer layers and tension in the central
portion) is usually desirable in service.
Compressive stresses inhibit failure by fa-
tigue and stress corrosion-two mecha-
nisms that initiate in the outer fibers. Un-
fortunately, metal-removal operations
required after heat treating often expose
material that is stressed in tension. Also,
metal-removal operations that are asym-
metrical (with respect to residual stresses)
cause distortion by redistributing residual
stresses. When close-tolerance parts are
being fabricated, the resulting warpage can
be costly and difficult to correct.
Although service performance is some-
times a factor, the major incentive for re-
ducing residual stress differentials has been
a reduction in warpage during machining or
an improvement in shape before machining.
One approach to reducing the cooling-
rate differential between surface and center
is the use of a milder quenching medium-
water that is hotter than that normally used
or water-glycol solutions. Boiling water,
which is the slowest quenching medium
used for thick sections, is sometimes em-
ployed for quenching wrought products
even though it lowers mechanical properties
and corrosion resistance. Quenching of
castings in boiling water, however, is stan-
dard practice, and is reflected in design
allowables.
Another approach to the minimization of
residual stresses that is generally successful
consists of rough machining to within 3.2
mm (0.125 in.) or less of finish dimensions,
heat treating, and then finish machining.
This procedure is intended to reduce the
cooling-rate differential between surface
and center by reducing thickness; other
benefits that accrue if this technique is used
to reduce or reverse surface tension stress-
es in finished parts are improvements in
strength, fatigue life, corrosion resistance,
and reduced probability of stress-corrosion
cracking.
Several factors (especially quenching
warpage) sometimes preclude general use of
this procedure. The thinner and less sym-
Fig 6
Longitudinal tensile strength, MPa
Longitudinal yield strength, MPa
Hardness, HRB
90
70
Depth, in.
0
0.5
1.0
1.5
2.0
2.5
3.0
60 O Control specimen
Quenched from side A only
A Quenched from side B, interrupted after 3 s
50
0
10
20
30
Side A
0
0.5
600
500
400
300
0
Side A
b
40
Depth, mm
Depth, in.
1.0
1.5
10
20
30
40
Depth, mm
Depth, in.
0
0.5
1.0
1.5
550
T
450
350
250
о
50
60
70
80
Side B
2.0
2.5
3.0
18
50
60
70
B
80
70
60
-50
80
788
Side B
2.0
2.5
3.0
T
70
60
50
50
Longitudinal tensile strength, ksi
30
Longitudinal yield strength, ksi
40
40
30
150
0
10
20
30
40
50
60
70
80
Side A
Depth, mm
Side B
Through-thickness property variations due to quench rate and temperature-rise effects in 7075-T62 plate
75 mm (3 in.) thick
metrical a section, the more it will warp
during quenching, and the residual stresses
resulting from straightening of warped parts
(plus straightening costs) often are less de-
sirable than the quenching stresses. Holding
fixtures and die quenching may be helpful,
but precautions must be taken to ensure
that they do not retard quenching rates
excessively. Other factors that must be con-
sidered are the availability of heat-treating
facilities and whether or not the advantages
of such a manufacturing sequence offset the
delay and cost entailed in a double-machin-
ing setup.
Warpage of thin sections during quench-
ing is also a problem. Even in the same
load, symmetry of cooling usually varies
significantly among identical parts and the
resultant inconsistent warpage usually re-
quires costly hand straightening. Conse-
quently, a significant amount of effort has
been devoted to reducing or eliminating
warpage by changing racking positions to
achieve symmetry of cooling.
For sheet-metal parts, one manufacturer
uses a double screen floor in the quenching
rack to reduce the force of initial contact
between water and parts. Others allow parts
to "free fall" from rack to quench tank.
Spacing and positioning on the rack are
carefully controlled so that parts will enter
the water with minimum impact. With this
technique, water turbulences must be
avoided, because it will often cause parts to
float for a few seconds, greatly reducing
their cooling rate.
Heat Treating of Aluminum Alloys / 855
Because of the difficulties encountered
with quenching in cold water, milder quen-
chants have been employed. Indiscriminate
use of milder quenchants can have cata-
strophic effects; however, when their use is
based on sound engineering judgment and a
metallurgical knowledge of the effects on
the specific alloy, significant cost savings or
performance improvements can be realized.
The most frequent advantage is the re-
duction in costly straightening operations
and in resultant uncontrolled residual
stresses. For example, one aircraft manu-
facturer utilizes water-spray and air-blast
quenching for weldments and complex
formed parts made from 6061, an alloy
whose corrosion resistance is insensitive to
quenching rate. Straightening requirements
are negligible and, through careful control
of racking and coolant flow, the decrease in
mechanical properties is minimized, as
shown by the data in Fig 10.
Another development for reducing straight-
ening costs is quenching in water-polymer
solutions. Quenching of formed sheet-metal
parts in aqueous solutions of polyalkylene
glycol or in similar inversely soluble media
has significantly reduced the cost of straight-
ening these parts after quenching. The SAE
heat-treatment specification AMS-2770 rec-
ommends, for several alloys, maximum thick-
nesses that can be quenched in solutions of
specific concentrations while maintaining ac-
ceptable property levels. Typical parameters
for quenching wrought products (other than
forgings) in glycol-water solutions are pre-
sented in Table 7. Additional information on
polymer quenchants for aluminum alloys can
be found in Ref 5.
Forming and Straightening after Quench-
ing. Immediately after being quenched,
most aluminum alloys are nearly as ductile
as they are in the annealed condition. Con-
sequently, it is often advantageous to form
or straighten parts in this temper. More-
over, at the mill level, controlled mechani-
cal deformation is the most common meth-
od of reducing residual quenching stresses.
Because precipitation hardening will occur
at room temperature, forming or straighten-
ing usually follows as soon after quenching
as possible. In addition, maximum effec-
tiveness in stress relief is obtained by work-
ing the metal immediately after quenching.
Forming and straightening operations
vary in degree from minor corrections of
warpage to complete forming of complex
parts from solution-treated flat blanks. Par-
ticular value is gained when enough forming
can be done at this stage of processing to
eliminate the distortion caused by quench-
ing. However, production operations must
be adjusted so that most of the plastic
deformation is accomplished before an ap-
preciable amount of precipitation hardening
takes place.
Although the most severe forming opera-
tions may have to be arranged to avoid
856 / Heat Treating of Nonferrous Alloys
Average cooling rate at 400-290 °C, °C/s
Thickness, mm
0.1
1
250
75
25
25
7.5
Air cool
2.5
0.75
0.25
0.1
1
Fig 7
10
10
100
103
10
Computed maximum
(assumes instantaneous
cooling of surface from
875-210 °F)
212 °F
180 °F
75 °F
150 °F
200 °F
Immersion in water at
indicated temperature
100
Average cooling rate at 750-550 °F, °F/s
0.1
0.01
103
104
Thickness, in.
Effects of thickness and quenching medium on average cooling rates at midplane of aluminum alloy
sheet and plate quenched from solution temperatures. The dashed line delineates the maximum cooling
rates theoretically obtainable at the midplane of plate, assuming an infinite heat transfer coefficient (C) and a
diffusivity factor of 1400 cm²/s. Source: Ref 3
natural aging, it often is desirable to allow
some natural aging to occur and thus avoid
formation of Lüders lines. This condition of
nonuniform deformation is most likely to
occur shortly after quenching and diminish-
es significantly after a few hours of natural
aging. Complete freedom from Lüders
lines, however, may require one or two
days of natural aging prior to forming. Thus,
the forming operation may have to be timed
so as to obtain the most appropriate trade-
off of these characteristics for the specific
parts involved. Lüders lines also can be
reduced by employing low strain rates or by
forming at temperatures of 150 to 175 °C
(300 to 350 °F).
Residual stresses in sheet-metal parts
formed in the quenched condition are higher
than those in parts formed in the annealed
condition. Consequently, forming in the
quenched condition should be selected judi-
precipitation rate is low despite the high
degree of supersaturation. At intermediate
temperatures, precipitation rate is highest.
Consequently, times to produce equal
amounts of precipitation follow a C-shape
pattern.
Using isothermal quenching techniques,
Fink and Willey pioneered the attempts to
describe the effects of quench rates with the
use of C-curves (Ref 6). The C-curves plot
the time required at different temperatures
to precipitate a sufficient amount of solute
to: reduce strength by a certain amount (Fig
5); cause a change in the corrosion behavior
from pitting to intergranular (Fig 12); pro-
duce a given electrical conductivity (Fig
13); or relate other properties, such as frac-
ture toughness, to isothermal quench con-
ditions. The nose of the C-curves identifies
the critical temperature range (the region of
highest precipitation rates). Investigators
use critical temperature ranges in conjunc-
tion with properties of samples quenched
continuously from the solution temperature
to compare relative sensitivities of alloys to
quenching condition.
Although average quench rates through a
critical temperature range can provide rea-
sonable property predictions if cooling rates
are fairly uniform, average quench rates can-
not provide quantitative predictions when
cooling rates vary considerably during the
ciously for parts that are critical in fatigue quench. For such instances, a procedure
(Fig 11) or stress corrosion.
of parts
Re-solution heat treatment
formed after quenching often causes exces-
sive grain growth in critically strained re-
gions and thus is not recommended.
Quench-Factor Analysis
During the quenching of alloys from a
solid-solution temperature condition, the
rate of precipitation during quenching is
maximized in a so-called "critical" temper-
ature range, because the diffusion of dis-
solved species and the subsequent nucle-
ation of precipitates exhibit opposite
behavior as a function of temperature. At
high temperatures, nucleation rates are
small because of the low degree of super-
saturation, and so precipitation rates are
low despite the high diffusion rates. At low
temperatures, diffusion rate is low, and thus
Table 7 Limits for quenching in glycol-water solutions
Data are for wrought aluminum alloy products other than forgings.
Glycol
Maximum thickness
concentration,
vol%
Alloys
mm
12-16
2014, 2017, 2117, 2024, 2219
7075, 7175
2.03
25.4
17-22
2014, 2017, 2117, 2024, 2219
1.80
in.
0.080
1.000
0.071
7075, 7079, 7175, 7178, 6061
12.7
0.500
23-28
2014, 2017, 2117, 2024, 2219
1.60
0.063
7075, 7079, 7175, 7178, 6061
9.53
0.375
29-34
2014, 2017, 2117, 2024, 2219
1.02
7075, 7079, 7175, 7178, 6061
6.35
35-40
7075, 7079, 7175, 7178, 6061
2.03
0.040
0.250
0.080
known as "quench-factor analysis" uses in-
formation from the entire C-curve to predict
how any quench curve affects properties.
Quench-factor analysis is useful in designing
suitable limits for quench delays, or when it is
not sufficient just to ensure that the cooling
curve misses the nose of the C-curve.
The method of quench-factor analysis, as
outlined by Evancho and Staley (Ref 7), is
based on the determination of a quench
factor (T), which is the major variable in the
following equation for precipitation kinetics
during continuous cooling:
(=1- exp (KT)
(Eq 1)
where is the fraction transformed and k is
a constant related to the transformation
fraction of a given C-curve. The quench
factor (T) is defined as:
T=
- ၂
dt
(Eq 2)
where is time and C, is critical time as a
function of temperature to transform a spec-
ified fraction (x). The locus of critical times
for a given transformation fraction x (or a
percentage of mechanical properties from
precipitation) is the C-curve, and the value
of k is related to x as follows: k = ln(1 - x),
or e⭑
= 1 x. Therefore, when T = 1, the
fraction transformed, 5, equals the fraction
value designated by the C-curve. Equation
2 is based on the assumption that the reac-
tion rate is a function only of the amount
transformed and temperature.
Temperature, °C
(a)
Temperature, °C
(b)
Fig 8
500
400
300
200
100
500
400
300
200
Black oxide
etched coating
1
2
3
Time, s
4
100
Powdered oxide
sprayed coating
0
20 °C water quench
As-rolled
surface
Sanded
surface
5
Boiling water quench
Black oxide
etched coating
As-rolled
surface
6
800
600
400
200
800
600
Surface
400
Sanded
surface
10
20
30
40
50
60
70
80
Time, s
200
Temperature, °F
Temperature. °F
Effect of surface conditions on the midplane cooling of a 13 mm (0.5 in.) thick plate of 7075 from
quenching in (a) 20 °C (70 °F) water and (b) boiling water. Source: Ref 5
Tension
Longitudinal stress, MPa
Compression
Cross section of solid cylindrical
2
specimen, in.
a
10
20 30
40
100
50
50
100
0
10
1
Quenched
in boiling
water
10
20
30
Cross section of solid cylindrical
specimen, 103 mm²
5
5
Tension →
10
✓ Compression
Tension→→
Longitudinal stress, ksi
Longitudinal stress, MPa
- Compression
Cross section of solid cylindrical
specimen, in.2
0
10
100
20 30
T
40
0
50
100
0
10
Quenched
10
in cold
water
20
T
30
Cross section of solid cylindrical
specimen, 103 mm²
10
- Compression
Tension
Longitudinal stress, ksi
Fig 9 Effect of quenching from 540 °C (1000 °F) on residual stresses in solid cylinders of alloy 6151
Heat Treating of Aluminum Alloys / 857
The numerical evaluation of the quench
factor involves the integration of Eq 2. This
integral can be graphically integrated using
the method illustrated in Fig 14. Examples
of the way to use the quench factor (T) in the
analysis of quench methods are described
below. Neither the average quenching rate
through a critical temperature range nor
quench-factor analysis can predict strength
when the temperature increases during
quenching after it is cooled below some
critical temperature. Under this condition,
strength in the affected areas can be signif-
icantly lower than in other areas of the
material. The most likely way for this phe-
nomenon to occur is during spray quench-
ing, when the surface cools rapidly by the
impinging spray, but reheats by heat flow
from the hotter interior when the spray is
interrupted.
Predicting Strengths of Thick Products.
Effects of the quenching rate on alloy
strengths can be represented on a general-
ized graph of the type shown in Fig 3, and
the expected quenching rates of products
having various dimensions can be deter-
mined from Fig 7. Nevertheless, combining
these two kinds of information to predict
mechanical properties must be done with
caution. Inconsistencies were encountered,
for example, in correlating properties of
thick sections quenched in high-cooling-rate
media with properties of thinner sections
quenched in media affording milder quench-
ing action. One of the reasons for the incon-
sistencies is believed to be the different
shapes of the cooling curves. This difficulty
can be overcome by using quench-factor
analysis. The other reason is that the degree
of recrystallization and texture of the thick
and thin sections may be different.
Predicting Corrosion Behavior. Alloy
2024-T4, for example, is susceptible to in-
tergranular corrosion when a critical
amount of solute is precipitated during
quenching, but will corrode in the less se-
vere pitting mode when lesser amounts are
precipitated. For predicting the effects of
proposed quenching conditions on the cor-
rosion characteristics of 2024-T4, the postu-
lated quench curve is drawn and the quench
factor is calculated using the C-curve in Fig
12. Corrosion characteristics are predicted
from the plot in Fig 15. When the quench
factor (T) is less than 1.0, continuously
quenched 2024-T4 will corrode by pitting.
These relationships are applied to studies
of effects of proposed changes in quench
practice on design of new quenching sys-
tems. For example, consider that the goal of
a proposed quenching system for 2024-T4
sheet products is to minimize warpage while
preventing susceptibility to intergranular
corrosion. Warpage occurs when the stress-
es imposed by temperature differences
across the parts exceed the flow stress. As
quenching rate decreases, the tendency for
large differences in temperature to occur
% of maximum yield strength
858 / Heat Treating of Nonferrous Alloys
Thickness, 0.001 in.
100
150
100
70
50
60
50
90
6061-T6 sheet
1250
2500
3750
200
Water
spray
250
Air blast
Stress, MPa
400
3
Curve Bend radius
123
Not bent
3.2 mm
Condition during flattening
Not applicable
50
Annealed
3.2 mm
4
3.2 mm
300
5
1.6 mm
As-quenched + 3 days storage
As-quenched + 14 days storage
As-quenched + 3 days storage
40
200
5000
6250
7500
100
0.01
Thickness, um
Thickness, 0.001 in.
Stress ratio, 0.1
0.1
4
5
1
30
1
2
20
3
10
Stress, ksi
% of maximum tensile strength
100
150
100
90
80
200
250
Water spray
Air blast
70
6061-T6 sheet
601250
2500
3750
5000
6250
Thickness, um
Fig 10
7500
Effect of quenching medium on strength of
6061-T6 sheet. Water-immersion quench
equals 100%. Control of coolant flow will minimize
decrease in mechanical properties.
decreases but the tendency for intergranular
corrosion to occur increases.
The C-curve in Fig 12 indicates that
quenching rate can be decreased near the
solution heat-treating temperature and near
room temperature without greatly sacrific-
ing corrosion characteristics, but this infor-
mation does not provide a quantitative an-
swer. Simple calculations, however, can
reveal a multitude of hypothetical cooling
curves that provide slow quenching during a
large portion of the quench cycle but suffi-
ciently rapid quenching where critical times
are short so that desirable corrosion char-
acteristics are obtained.
As an example, one-, two-, and three-step
quench curves that would ensure accept-
Millions of cycles to failure
Fig 11 Fatigue characteristics of 1 mm (0.04 in.) alclad 2024-T4 sheet after 90° bending in the annealed
condition and subsequent flattening as indicated. Flattening (unbending) was done either in the
annealed condition (curve 2), or in the solution-treated and quenched condition (curves 3, 4, 5) with indicated
storage times at -18 to -12 °C (0 to 10 °F).
able corrosion behavior in 2024-T4 sheet
(quench factor, 0.99) were calculated. Some
of these curves are plotted in Fig 16. This
illustration shows that 2024 can be
quenched at a rate of 470 °C/s (850 °F/s) or
higher and still develop acceptable corro-
sion characteristics if the quenching rate is
linear from the solution temperature to 150
°C (300 °F). If sheet 3.2 mm (0.125 in.) thick
is air-blast quenched (rate of heat removal,
5.68 W/m² °C) to 395 °C (740 °F), however,
the quenching rate from 395 to 150 °C must
be at least 945 °C/s (1700 °F/s) to maintain
the acceptable corrosion behavior. It may
also be air-blast quenched to 395 °C (740
°F), spray quenched at 3300 °C/s (6000 °F/s)
to 250 °C (480 °F), then air-blast quenched
to 150 °C (300 °F).
Holding temperature, K
700
650
600
550
Other curves could be drawn, of course,
but the important points are that air-blast
quenching cannot be continued to more
than a few degrees below 395 °C (740 °F)
and cannot be initiated at more than a few
degrees above 270 °C (520 °F) even if infi-
nite quenching rates are attained from 395
to 270 °C (740 to 520 °F).
Predicting yield strength is more complex
than predicting corrosion behavior and re-
quires some knowledge of the relationship
between extent of precipitation and loss in
ability to develop property. Because attain-
able strength of precipitation-hardening alu-
minum alloys is a function of the amount of
solute remaining in solid solution after
quenching, relationships between strength
(σ) attainable after continuous cooling and
IACS
17%
18%
425
375
19%
20%
325
21%
275
Holding temperature, °C
Temperature, °C
500
900
400
300
200
100
0.1
Predominantly
pitting
1
Predominantly
intergranular
corrosion
700
500
300
Temperature, °F
10
100
103
Critical time, s
Fig 12
C-curve indicating type of corrosion attack
on 2024-T4 sheet
500
17%
450
10
100
225
19%
20%
21%
18%
22%
175
103
104
105
106
Holding time, s
Fig 13 Change in electrical conductivity of an Al-2.5% Li binary alloy after the following: solution treated at 540
°C (1000 °F) for 12 h, immersed into an adjacent salt or oil bath for the appropriate isothermal holding
temperature and time, then quenched into water. Source: Ref 1
Temperature
ドドド
Quench curve
Temperature
(T₁ + T₂)/2
け
+ T3)/2
C-curve
Temperature, °C
Heat Treating of Aluminum Alloys / 859
500
900
Air-blast quench
3.2-mm (0.125-in.) sheet
on both sides
800
400
470 °C/s (850 °F/s)
700
2780 °C/s (5000 °F/s)
Air-blast quench
600
300
3.2-mm (0.125-in.)
sheet
500
Temperature, 'F
TEITI
1++AG
|
1+AtF-1
TF-1 tF
Elapsed time
(TF + TF)/2
С2
C₁₁ CF-1
Critical time
x =
+
Δια Δια
C₂
+
AtF-1
CF-1
Fig 14 Method of determining quench factor, &, using a cooling curve and a C-curve
Average depth, mm
0.20
Quench
o Air-water-air
• Water-air
0.15
A Air-water
▲ Water
0.10
□ Air
0.05
Type of corrosion
T
QA
0
0.1
1
0.2
է
1
L
1
0.4
0.6 0.8 1.0
2.0
4.0
6.0 8.0 10.0
20.0
Quench factor, T
P + SI
P +
0.1
Quench
Air-water-air
• Water-air
▲ Air-water
▲ Water
□ Air
P-Pitting
0.006
0.004
0.002
2 P+SI - Pitting and slight intergranular
P+1 Pitting and intergranular
Intergranular
0800
0.2
0.4
0.6 0.8 1.0
2.0
Quench factor, 1
4.0 6.0 8.0 10.0
20.0
Fig 15 Type and depth of attack on 2024-T4 sheet versus quench factor
quench factor (T) can be expressed as fol-
lows:
(Eq 3)
σx=σmax exp (k₁T)
where σ max
is the strength attainable with
an infinite quenching rate and:
T =
Sa
dt
(Eq 4)
where t is time and Cx is the C-curve for
σ that is, critical time as a function of
temperature to reduce attainable strength to
x of σmax. The constant k₁ is related to the
natural logarithm of x. For example, if 7 is
based on the C-curve for 99.5% of maxi-
mum yield strength, then k₁ = -0.005013 =
In (0.995).
The advantage of predicting yield
strength from quench factor instead of from
average quenching rate is illustrated by the
Average depth, in.
following comparison. Four specimens of
alloy 7075-T6 quenched by various means
(see Fig 17) were selected. Yield strengths
were predicted both from average quench-
ing rate between 400 and 290 °C (750 and
550 °F) and from quench factor. Quench
factor was calculated using the C-curve for
99.5% maximum yield strength for 7075-T6
(Fig 18), and yield strength was estimated
from the above equation defining the
quench factor (7) (see Fig 19).
A comparison of predicted yield strength
with actual yield strength is given in Table
8. Yield strengths predicted from quench
factor agree very well with measured yield
strengths for all specimens, the maximum
error being 19.3 MPa (2.8 ksi). Yield
strengths predicted from average quenching
rates, however, differ from measured val-
ues by as much as 226 MPa (32.8 ksi).
200
400
3330 °C/s
(6000 F/s)
945 °C/s
300
(1700 F/s)
100
1
0
2
4
6
Time, s
Fig 16 Quench curves for 2024-T4 sheet, to elimi-
nate susceptibility to intergranular corrosion
The advantage of using the quench factor
for predicting yield strength from cooling
curves is apparent. Cooling curves that
have long holding times either above or
below the critical temperature range from
400 to 290 °C (750 to 550 °F) cannot be used
to predict yield strength from average
quenching rate. In such instances, predic-
tion of yield strength on the basis of quench
factor is particularly advantageous.
With the use of finite-element analysis,
quench factors can also be plotted as a
function of Grossmann quench severity val-
ues (H) or the heat transfer coefficients (C)
between the quenchant and a particular part
(Fig 20). However, an underlying assump-
tion of both quench-factor analysis and av-
erage-cooling-rate estimation is that the
only effect of temperature is on the kinetics
of precipitation. This assumption is not val-
id, however, when portions of the metal are
quenched locally but reheated significantly
before quenching is complete.
Age Hardening
After solution treatment and quenching,
hardening is achieved either at room tem-
perature (natural aging) or with a precipita-
tion heat treatment (artificial aging). In
some alloys, sufficient precipitation occurs
in a few days at room temperature to yield
stable products with properties that are
adequate for many applications. These al-
loys sometimes are precipitation heat treat-
ed to provide increased strength and hard-
ness in wrought or cast products. Other
alloys with slow precipitation reactions at
room temperature are always precipitation
heat treated before being used.
In some alloys, notably those of the 2xxx
series, cold working of freshly quenched
material greatly increases its response to
later precipitation heat treatment. Mills take
advantage of this phenomenon by applying
a controlled amount of rolling (sheet and
plate) or stretching (extrusion, bar, and
plate) to produce higher mechanical proper-
ties. However, if the higher properties are
used in design, reheat treatment must be
avoided.
Temperature, °C
860 / Heat Treating of Nonferrous Alloys
500
Quenched in still air to 370 °C (700 °F) + cold-water quench
900
500
800
400
300
Quenched in boiling water to
315 °C (600 °F) + cold-water quench
700
-600
00
Temperature, °F
- 400
Temperature, "C
300
" 500
200
F6 Quenched in denatured alcohol to 290 °C
†(550 °F) cold-water quench
Cold-water quench
200
100
0
5
10
15
20
25
30
35
Time, s
400
300
300
100
0.1
1
10
100
Critical time, s
1000
10 000
900
700
500
Fig 17 Cooling curves for 7075-T6 sheet
Natural Aging. The more highly alloyed
members of the 6xxx wrought series, the
copper-containing alloys of the 7xxx group,
and all of the 2xxx alloys are almost always
solution heat treated and quenched. For
some of these alloys-particularly the 2xxx
alloys the precipitation hardening that re-
sults from natural aging alone produces
useful tempers (T3 and T4 types) that are
characterized by high ratios of tensile to
yield strength and high fracture toughness
and resistance to fatigue. For the alloys that
are used in these tempers, the relatively
high supersaturation of atoms and vacan-
cies retained by rapid quenching causes
rapid formation of GP zones, and strength
increases rapidly, attaining nearly maxi-
mum stable values in four or five days.
Tensile-property specifications for products
in T3- and T4-type tempers are based on a
nominal natural aging time of four days. In
alloys for which T3- or T4-type tempers are
standard, the changes that occur on further
natural aging are of relatively minor magni-
tude, and products of these combinations of
alloy and temper are regarded as essentially
stable after about one week.
In contrast to the relatively stable condi-
tion reached in a few days by 2xxx alloys
that are used in T3- or T4-type tempers, the
6xxx alloys and to an even greater degree
the 7xxx alloys are considerably less stable
at room temperature and continue to exhibit
significant changes in mechanical properties
for many years. The differences in rate and
Fig 18 C-curve for 99.5% maximum yield strength of 7075-T6 sheet
duration of changes in tensile yield strength
of representative alloys of the three types
are illustrated in Fig 21. Because of the
relative instability of the 7xxx alloys, the
naturally aged temper (after solution heat
treatment and quenching) is designated by
the suffix letter W. For a specific descrip-
tion of this condition, the time of natural
aging should be included (example: 7075-W,
1 month).
Aging characteristics vary from alloy to
alloy with respect to both time to initial
change in mechanical properties and rate of
change, but aging effects always are less-
ened by reductions in aging temperature
(see Fig 21). With some alloys, aging can be
suppressed or delayed for several days by
holding at a temperature of -18 °C (0 °F) or
lower. It is usual practice to complete form-
ing and straightening before aging changes
mechanical properties appreciably. When
scheduling makes this impractical, aging
may be avoided in some alloys by refriger-
ating prior to forming. It is conventional
practice to refrigerate alloy 2024-T4 rivets
to maintain good driving characteristics.
Full-size wing plates for current-generation
jet aircraft have been solution heat treated
and quenched at the primary fabricating
mill, packed in dry ice in specially designed
insulated shipping containers and transport-
ed by rail about 2000 miles to the aircraft
manufacturer's plant for forming.
Unanticipated difficulties may arise as a
result of failure to control refrigerator or
part temperature closely enough. If opening
of the cold box to insert or remove parts is
done too frequently, the cooling capacity of
the refrigerator may be exceeded. At times,
the rate at which heavy-gage parts can be
cooled in a still-air cold box has been found
to be insufficient. This problem has been
solved in one plant by immersing parts in a
solvent at -40 °C (-40 °F) before placing
them in the refrigerator.
The T3-type tempers are distinguished
from T4-type tempers by significant me-
chanical-property differences resulting from
cold work strain hardening associated with
certain mechanical operations performed
after quenching. Roller or stretcher leveling
to achieve flatness or straightness introduc-
es modest strains (on the order of 1 to 4%)
that cause changes in mechanical properties
(primarily, increases in strength). Further
increases in strength can be obtained by
cold rolling, additional stretching, combina-
tions of these operations, or for products
such as hand forgings, compressive defor-
mation. The tempers produced by these
operations followed by natural aging alone
(no precipitation heat treatment) are classi-
fied as T3-type tempers, and an additional
digit is used to indicate a variation in strain
hardening that results in significant changes
in properties. In the most recently intro-
Table 8 Yield-strength values for 7075-T6 sheet predicted from cooling curves using
average quench rate and quench factor
% of maximum yield strength
Average quench
rate from 400 to
290 °C (750 to
100
550 °F)
Quench
factor,
Measured yield
strength
Yield strength
predicted from
average quench
rate
Yield strength
predicted from
quench factor
80
Quench
°C/s
°F/s
T
MPa
ksi
MPa
ksi
MPa
ksi
60
100 o/max exp(-0.005013 t)
Cold water
935
1680
0.464
506
73.4
499
72.4
498
72.3
Denatured alcohol to
40
290 °C (550 °F),
then cold water
50
90
8.539
476
69.1
463
67.2
478
69.4
20
Boiling water to
0
315 °C (600 °F),
0.01
0.1
1
10
Quench factor, t
100
103
then cold water
30
55
15.327
458
66.4
443
64.2
463
67.1
Still air to 370 °C
(700 °F), then cold
water
5
9
21.334
468
67.9
242
35.1
449
65.1
Fig 19 Yield strength versus quench factor
Temperature, "F

Quench factor for 99.5% of
attainable yield strength for alloy 7075-T73
10
5
1
40
45
Sheet thickness, in.
Plate thickness, in.
0
0.05
0.1
0.15
0.2
0.25
0.3
0
0.5
1
1.5
2
2.5
3
50
550
50
0
1.25
2.5
(a)
C=0.036 W/cm2.K
3.75
Sheet thickness, mm
C = 0.072
C=0.108
C = 0.144
C = 0.216
C = 0.288
C = 0.432
5
C = 0.720
C = 1.44
6.25
7.5
Quench factor for 99.5% of
attainable yield strength for alloy 7075-T73
10
5
(b)
1
45
C = 0.216
Heat Treating of Aluminum Alloys / 861
C = 2.16
C = 2.88
C = 4.32
C=7.20
Plot of quench factors derived from finite element analysis with given product sizes and film (heat transfer) coefficients (C). Heat transfer coefficients between
Fig 20
the quenchant and part are expressed in W/cm² K. Source: Ref 4
12.5
25
37.5
50
62.5
75
Plate thickness, mm
duced 2xxx aircraft alloy, 2324, high
strength is achieved by cold rolling plate to
a T39 temper.
Precipitation heat treatments generally
are low-temperature, long-term processes.
Temperatures range from 115 to 190 °C (240
to 375 °F); times vary from 5 to 48 h.
Choice of time-temperature cycles for
precipitation heat treatment should receive
careful consideration. Larger particles of
precipitate result from longer times and
higher temperatures; however, the larger
particles must, of necessity, be fewer in
number with greater distances between
them. The objective is to select the cycle
that produces optimum precipitate size and
distribution pattern. Unfortunately, the cy-
cle required to maximize one property, such
as tensile strength, is usually different from
that required to maximize others, such as
yield strength and corrosion resistance.
Consequently, the cycles used represent
compromises that provide the best combi-
nations of properties.
Production of material in T5- through
T10-type tempers (see the section on tem-
per designations near the end of this article)
necessitates precipitation heat treating at
elevated temperatures (artificial aging). Al-
though the hardening precipitate developed
by this operation is submicroscopic, struc-
tures before and after precipitation heat
treatment often can be distinguished by
etching metallographic specimens. In alumi-
num alloys in the solution heat treated and
quenched condition, coloration contrast be-
tween grains of differing orientation is rela-
tively high, particularly in 2xxx series
wrought alloys and 2xx.0 series casting al-
loys. This contrast is noticeably decreased
by precipitation heat treatment.
Differences in type, volume fraction,
size, and distribution of the precipitated
particles govern properties as well as the
changes observed with time and tempera-
ture, and these are all affected by the initial
state of the structure. The initial structure
may vary in wrought products from unre-
crystallized to recrystallized and may ex-
hibit only modest strain from quenching or
additional strain from cold working after
solution heat treatment. These conditions,
as well as the time and temperature of
precipitation heat treatment, affect the final
structure and the resulting mechanical prop-
erties.
Because mechanical properties and other
characteristics change continuously with
time and with temperature, as shown in Fig
22(a), (b), and (c) by typical curves for three
wrought alloys, treatment to produce a
combination of properties corresponding to
a specific alloy-temper combination re-
quires one or more rather specific and co-
ordinated combinations of time and temper-
ature, with both parameters being subject to
practical limitations. Recommended com-
mercial treatments often are compromises
between time and cost factors and the prob-
ability of obtaining the intended properties,
with consideration of allowances for varia-
bles such as composition within specified
range and temperature variations within the
furnace and load. Use of higher tempera-
tures may reduce treatment time; but if the
temperature is too high, characteristic fea-
tures of the precipitation-hardening process
reduce the probability of obtaining the re-
quired properties.
T6 and T7 Tempers. Precipitation heat
treatment following solution heat treatment
and quenching produces T6- and T7-type
tempers. Alloys in T6-type tempers gener-
ally have the highest strengths practical
without sacrifice of the minimum levels of
other properties and characteristics found
by experience to be satisfactory and useful
for engineering applications. Alloys in T7
tempers are overaged, which means that
some degree of strength has been sacrificed
or "traded off" to improve one or more
other characteristics. Strength may be sac-
rificed to improve dimensional stability,
particularly in products intended for service
at elevated temperatures, or to lower resid-
ual stresses in order to reduce warpage or
distortion in machining. T7-type tempers
frequently are specified for cast or forged
engine parts. Precipitation heat-treating
temperatures used to produce these tem-
pers generally are higher than those used to
produce T6-type tempers in the same al-
loys.
Two important groups of T7-type
tempers-the T73 and T76 types-have
been developed for the wrought alloys of
the 7xxx series, which contain more than
about 1.25% copper. These tempers are
intended to improve resistance to exfolia-
tion corrosion and stress-corrosion crack-
ing, but as a result of overaging, they also
increase fracture toughness and, under
some conditions, reduce rates of fatigue-
crack propagation. The T73-type temper
862 / Heat Treating of Nonferrous Alloys
Tensile strength, MPa
Yield strength, MPa
Elongation, % in 50 mm or 2 in.
600
2014
500
RT
400
0 °C (32 °F).
60
300
200
80
2024
600
80
T
500
70
70
RT
Tensile strength, ksi
Tensile strength, MPa
0 °C (32°F)
60
400
-18 °C (0 °F)
50
--18°C (0°F)
40
300
T
T
200
1 year
20
30 min
1 day
1 week 2 months.
100
0.1
1
102
103
104
30
1 year
30 min
1 day
1 week 2 months
100
0.1
1
102
103
104
500
400
10
Elapsed time after quenching, h
1
70
T
20
I
1
60
50
50
300
RT
40
0°C (32°F)
30
200
100
Yield strength, ksi
Yield strength, MPa
500
400
300
200
10
Elapsed time after quenching, h
RT
T
50
40
30
20
10
70
60
50
Tensile strength, ksi
30
Yield strength, ksi
0°C (32 F)
40
30
-18 C (OF)
20
20
20
-18 °C (0 °F)
100
1 year
10
1 year
10
30 min
1 day
1 week 2 months
30 min
1 day
1 week 2 months
11
0
0
0
0.1
1
102
103
104
0
0.1
1
102
103
104
40
30
20
20
10
10
Elapsed time after quenching, h
-18 °C (0 ˚F)
0 °C (32°F)
Elongation, % in 50 mm or 2 in.
40
30
10
Elapsed time after quenching, h
-18 °C (0°F)
0°C (32°F)
RT
20
10
10
1 year
1 week 2 months
1 year
30 min
1 day
1 week 2 months
30 min
1 day
I
I
0
0
0.1
1
10
102
103
104
0.1
1
10
1
102
Elapsed time after quenching, h
11
103
Elapsed time after quenching, h
104
Fig 21 Aging characteristics of aluminum sheet alloys at room temperature, at 0 °C (32 °F), and at -18 °C (0 °F)
has greatly minimized stress-corrosion
cracking of large and complex machined
parts made of these alloys, which occa-
sionally occurred with T6-type tempers.
The precipitation heat treatment used to
produce the T73- and T76-type tempers
consist either of a two-stage isothermal
precipitation heat treatment or of heating
at a controlled rate to a single treatment
temperature. The microstructural/
electrochemical relationships that are re-
quired in order to achieve the desired
corrosion-resisting characteristics can be
developed by using only a single-stage
precipitation heat treatment above about
150 °C (300 °F), but higher strength is
obtained by preceding this with a lower-
temperature stage or with a slow-con-
trolled heatup. Extended natural aging can
provide the same results, but the times
required at room temperature are imprac-
tical. Either during the preliminary stage
or during slow heatup, a fine, high-density
dispersion of GP zones is nucleated. Either
the time and temperature of the first step
or the rate of heating must be controlled to
produce GP zones that will not dissolve
but will transform to the n' precipitate
when heated to the aging temperature
above 150 °C (300 °F). The aging practice
that produces the results in the shortest
time depends on the GP-zone solvus tem-
perature. This temperature, in turn, de-
pends on vacancy concentration, a factor
influenced by solution heat-treating tem-
perature and quench rate, and on compo-
sition. If first-step aging time is too short,
if first-step aging temperature is too far
below the GP-zone solvus, or if heating
rates are too high, the GP zones will
dissolve above 150 °C (300 °F), and the
resultant coarse and widely distributed
precipitate will provide lower strength.
The T76-type treatments have the same
operational sequence but employ second-
stage heating only long enough to develop
a resistance to exfoliation corrosion higher
than that provided by the T6-type tempers.
Materials in the T73-type temper also have
high resistance to exfoliation corrosion.
Recommended treatments to produce T5-
and T6-type tempers, and those of the T7-
type employed for dimensional and proper-
ty stabilization, provide adequate tolerance
for normal variations encountered with
good operating practices. On the other
hand, the T73, T74 (formerly T736), and
T76 tempers for alloys 7049, 7050, 7075,
7175, and 7475 involve changes in strength
that occur significantly more rapidly at the
temperatures employed in the second stage
of the T7x precipitation heat-treatment cy-
cle compared to the changes occurring at
the temperatures employed to produce the
T6 temper.
As illustrated in Fig 23, variations in soak
time of several hours, and variations in soak
temperature of up to 11 °C (20 °F) from the
nominal aging practice of 24 h at 120 °C (250
°F) affect the strength of 7075-T6 by as
much as 28 MPa (4 ksi). In contrast, similar
variations in second-step soak time and
temperature for 7075-T73-that is, varia-
tions for 24 h at 165 °C (325 °F)-affect
strength by up to 150 MPa (22 ksi).
Consequently, control of both tempera-
ture and time to achieve the mechanical
properties and corrosion resistance speci-
fied for these tempers is more critical than
the control required in producing the T6
temper. Moreover, rate of heating from the
first to the second aging step must be con-
sidered, because precipitation occurs dur-
ing this period.
Heat treaters attempt to adjust these new
problems by empirically modifying soak
times to compensate for precipitation dur-
ing heating and for effects of soaking at
temperatures above or below the nominal.
A method has been developed (Ref 8) that
permits quantitative compensation for the
effects of precipitation during heating and of
soaking either above or below the recom-
mended temperature. For overaging, these
effects can be described by the following
equation:
YS = Y exp
-
(
+ 0
FYS
(Eq 5)
where YS is yield strength; Y is a term
having units of strength that is dependent on
alloy, fabrication, and test direction; to is
time at soak temperature; Fys is a temper-
ature-dependent term; and
YS
Tensite strength, MPa
Yield strength, MPa
Elongation, " in 50 mm or 2 in.
6061
600
500
400
300
200
30 min
100
0.1
1
500
400
300
200
10
RT
0 °C (32°F)
T
T
T
--
80
70
60
50
40
10
Tensile strength, ksi
Tensile strength, MPa
600
7050
00
80
500
70
RT
60
400
300
T
50
40
Heat Treating of Aluminum Alloys / 863
Tensile strength, ksi
Tensile strength, MPa
600
500
7075
80
70
0°C (32°F)
RT
60
400
300
-18 °C (0°F)
50
40
40
T
T
30
200
30
200
-18 °C (0 °F)
30
30
1 year
1 year
-
1 day 1 week 2 months
20
20
1 year
30 min
1 day
1 week
2 months
20
| L
102
11
100
1
☐ 1
30 min
☐
100
1 day
1
1 week 2 months
11
103
104
0.1
1
10
102
103
104
0.1
1
10
102
103
104
24
20
Elapsed time after quenching, h
T
T
T
500
70
60
50
40
30
Yield strength, ksi
Yield strength, MPa
400
300
200
Elapsed time quenching, h
RT
500
70
0
60
T
T
T
60
400
50
50
40
40
30
Yield strength, ksi
Yield strength, MPa
Elapsed time after quenching, h
300
RT
0 °C (32°F)
200
T
70
660
650
40
30
RT
20
20
T
20
20
100
0 °C (32 F:
100
-18 °C (0 °F)
100
-18 °C (0 °F) 1 year 10.
1 year
10
10
30 min
1 day 1 week
2 months
30 min
1 day
1 week 2 months
1 year
0
0.1
1
☐☐
10
11
102
0
0
11
☐
☐
0
30 min
1 day
1 week 2 months
103
104
0.1
1
10
102
103
104
0
1
11
0
0.1
1
10
102
103
104
Elapsed time after quenching, h
Elapsed time after quenching, h
40
30
20
10
-18 °C (0°F)
0 °C (32°F)
RT
1 year
30 min
1 day
T
0
0.1
1
10
1 week 2 months
☐
103
102
Elapsed time after quenching, h
Fig 21 (continued)
Elongation, % in 50 mm or 2 in.
40
30
RT
1 year
30 min
1 day
1 week 2 months.
104
0
0.1
☐
1
☐ 1
1
10
102
103
104
Elapsed time after quenching, h
Elongation, % in 50 mm or 2 in.
40
30
Elapsed time after quenching, h
20
0 °C (32°F)
10
-18 °C (0 °F)
RT
1 year
30 min
0
1 day
1
0.1
1
10
1 week 2 months
☐
102
103
104
Elapsed time after quenching, h
Tensile strength, ksi
Yield strength, ksi
0=
S
dt
Fys
(Eq 6)
Fys 1.45×10-16
exp (
18 090
)
TK
(Eq 7b)
where t is time during heating.
Equation 5 provides the basis for selec-
tion of a nominal aging time that will result
in the desired yield strength and gives the
furnace operator a method of compensating
for heating rate and for differences between
desired and attained soak temperatures.
Specifics will be illustrated using data for
alloy 7050. The value of Fys (in units of
hours) for 7050 can be calculated by the
following equation:
Fys 1.45×10-16 exp (
32 562
TF+460
)
(Eq 7a)
where T is temperature in °F, or
where Tk is temperature in K.
In one experiment, lengths of 7050-W (4
days) extrusions were aged at 24 h at 120
°C (250 °F) plus the equivalent of 3 to 42 h
at 165 °C (325 °F). For the second step, a
logarithmic heatup was used in which 10 h
were required for the load to reach 155 °C
(315 °F), and nominal soak temperature
was 165 °C (325 °F). Figure 24 indicates
that yield strength generally agreed with
values predicted using Eq 5. The deviation
of the curve for short-transverse strength
at the short aging times indicates that the
method is inadequate for predicting
strength on the underaging side of the
aging curve.
The effects of neglecting to compensate
for soaking at temperatures other than the
nominal can be large (Fig 25). For exam-
ple, the calculated difference in strength
between alloy 7050 extrusions soaked 29 h
at 160 °C (320 °F) and at 165 °C (325 °F) is
about 50 MPa (7 ksi), and the calculated
difference in strength between 7050 extru-
sions soaked 29 h at 155 °C (315 °F) and at
170 °C (335 °F) is about 100 MPa (14 ksi).
Neglecting to compensate for time spent
heating the work to the soak temperature
will increase the variability. Strength loss
attributed to heatup was 14 MPa (2 ksi).
These kinetic relationships also can as-
sist in selection of equivalent aging times
for alternate second-step aging tempera-
tures. Equations 5 and 7 can be rearranged
Elongation, % in 50 mm or 2 in.
Yield strength, MPa
Tensile strength, MPa
864 / Heat Treating of Nonferrous Alloys
50
500
450
400
350
300
260 "C
(500 °F)
250
200
0
0.01
0.1
500
400
550
500
105 C
RT
(225 F) -70
450
60
135 "C
(275 F)
150 °C
(300 "F)
175 "C
(350 (F)
190 "C
(375 °F)
205 "C
230 "C
(450 "F)
(400 "F)
T
Tensile strength, ksi
Tensile strength, MPa
400
350
300
40
250
☐
100 C
(212 F)
150 C
NW(320F) (300)
190 C
(375 °F)
IT
70
50
60
Tensile strength, ksi
205 C
260 °C
(500 ˚F)
40
(400 °F)
30 min
30 min
1 day
1
10
1 day 1 week 2 months 1 year
1 week 2 months 1 year.
30
00
200
1
100
103
104
105
0
0.01
0.1
1
10
100
103
1|
104
30
105
Duration of precipitation heat treatment, h
Aging time, h
500
70
105 °C
(225°F)
60
+
135°C
400
RT
(275 °F)
50
175 C
(350 °F)
130 °C (265°F) 150 °C
(300 ˚F)
70
10
60
60
300
150 °C
40
(300 °F)
200
175 °C (350 °F) - 30
190 °C (375 °F)
Yield strength, ksi
Yield strength, MPa
100°C
(212 "F)
50
300
190 °C
(375 "F)
160 °C
(320 °F)
200
40
40
-205 °C (400 °F)
30
205 C
(400 F)
20
260 °C (500 °F).
100
20
230 °C (450 °F).
100
00
260 ˚C
10
10
(500 "F)
30 min
1 day 1 week 2 months 1 year
30 min
1 day 1 week 2 months 1 year
0
0
0.01
0.1
1
10
100
11
103
1
0
1
☐
☐
0
0
104
105
0
0.01
0.1
1
10
100
103
104
105
Duration of precipitation heat treatment, h
Aging time, h
40
30 205 °C
190 °C
(375 °F)
(400 °F)
175 °C
(350 °F)
150 "C 135 C
(300 ˚F) (275 °F)
20
Elongation, % in 50 mm or 2 in.
20
RT
15
10
260 °C
(500 °F)
100 C
(212 °F)
130 C
(265°F)
205 C
(400 F)
150 C
(300 F)
10
105 C
(225 FI
5
190 °C 1
(375 "F)
160 °C
260 °C
(500 °F)
230 C
(450 °F)
30 min
1 day
0
1 week 2 months 1 year
☐ 1
30 min
†175 C
(350 F) (320 ˚F)
1 day 1 week 2 months 1 year
0
0.01
0.1
1
10
100
103
☐
104
0
1
1
105
0
0.01
0.1
1
10
100
11
103
||
104
105
Aging time, h
Duration of precipitation heat treatment, h
Fig 22(a) Aging characteristics of alloy 2014 sheet
to yield the following equation:
t2=t₁ exp
32 562 32 562
T₁+460 T₂+460
(Eq 8)
where 11 is aging time at temperature T₁, t₂ is
aging time at temperature T₂ that will provide
equivalent yield strength, and T, and T₂ are in
°F. For example, the time at 175 °C (350 °F)
equivalent to aging alloy 7050 for 29 h at 165
°C (325 °F) is calculated as follows:
Fig 22(b) Aging characteristics of alloy 2024 sheet (see also Fig 22d)
1350 29/exp (1.28)=29/3.6=8 h
Thermomechanical effects on aging occur
from deformation after solution heat treat-
ment. The deformation step may be warm
or cold and before, after, or during aging.
The simplest thermomechanical practices
are those of the conventional T3, T8, or T9
tempers. The rate and extent of precipita-
tion strengthening are distinctly increased
in some alloys by cold working after
quenching, whereas other alloys show little
or no added strengthening when treated by
this sequence of operations.
Alloys of the 2xxx series such as 2014,
2124, and 2219 are particularly responsive
to cold work between quenching and aging,
and this characteristic is the basis for the
higher-strength T8 tempers. The strength
improvement accruing from the combina-
tion of cold working and precipitation heat
treating is a result of nucleation of addition-
Yield strength, ksi
Tensile strength, MPa
Yield strength, MPa
Elongation, % in 50 mm or 2 in.
350
325
(250 °F)
300
275
250
225
200
150 °C
(300 °F)
170 °C
(340 °F)
205 °C
(400 °F)
230 °C
(450 °F)
260 °C
(500 °F)
30 min
1 day
175
1 week 2 months 1 year
☐ 1
J
0
0.01
0.1
1
10
100
103
104
105
300
250
200
150
100
Duration of precipitation heat treatment, h
50
0
0.01
0.1
40
445
40
35
35
30
40
120 °C
(250 °F)
150 °C
(300 °F)
35
170 °C
(340 °C)
30
205 °C
(400 °F)
1
230 °C
(450 °F)
25
20
20
15
260 °C
(500 °F)
30 min
1 day
1 week 2 months 1 year
10
103
104
105
1
10
100
Duration of precipitation heat treatment, h
30
230 °C 205 °C
(450 °F) (400 °F)
170 °C
(340 °F)
_ 150 °C
(300 °F)
260 °C
(500 °F)
120°C
(250 °F)
201
10-260 °C
(500 °F)
0
0
0.01
0.1
30 min
10
1 day 1 week 2 months 1 year
100
1
103
Duration of precipitation heat treatment, h
Fig 22(c) Aging characteristics of alloy 6061 sheet
al precipitate particles by the increased
strain. In some alloys of the 2xxx series,
strain introduced by cold working after so-
lution heat treatment and quenching also
induces nucleation of a finer precipitate
dispersion that increases strength. Depend-
ing on the aging temper, however, tough-
ness may be adversely affected, as illustrat-
ed in Fig 26 for 2024 sheet.
Strengthening from thermomechanical
processing is the basis for the higher-
11
104
105
Tensile strength, ksi
Yield strength, ksi
strength T8-type tempers of alloys 2011,
2024, 2124, 2219, and 2419, which are
produced by applying controlled amounts
of cold rolling, stretching, or combinations
of these operations. Normally, cold work
is introduced by stretching; however, oth-
er methods such as cold rolling can be
used. Recently, 2324-T39 was developed.
The T39 temper is obtained by cold rolling
approximately 10% after quenching fol-
lowed by stretching to stress relieve. This
Heat Treating of Aluminum Alloys / 865
type of approach results in strengths sim-
ilar to those obtained with T8 processing
but with the better toughness and fatigue
characteristics of T3 products. Alloys
2024, 2124, and 2219 in T8-type tempers
are particularly well suited for supersonic
and military aircraft; alloy 2219 in such
tempers, and alloy 2014-T65, were the
principal materials for the fuel and oxidiz-
er tanks (which also served as the primary
structure) of the Saturn V space vehicles.
Re-solution heat treatment of mill prod-
ucts supplied in these tempers can result in
grain growth and in substantially lower
strength than is normal for the original
temper. Such reheat treatment is not rec-
ommended.
Alloys of the 7xxx series do not respond
favorably to the sequence of operations
used to produce T8-type tempers, and no
such tempers are standard for these alloys.
The strains associated with stretching or
compressing of 7xxx alloys have relatively
little effect on the mechanical properties of
material precipitation heat treated to T6-
type tempers. On the other hand, these
operations have measurable detrimental
effects on final strength when T73-, T736,
or T76-type tempers are produced, partic-
ularly in the direction opposite the direc-
tion of cold work. Accordingly, specifica-
tion properties are somewhat lower for
the stress-relieved versions of these tem-
pers. Decreasing the overaging time to
compensate for the loss in strength is not
advisable, because this would impair de-
velopment of the desired corrosion char-
acteristics.
Temperature control and uniformity pre-
sent essentially the same problems in pre-
cipitation heat treating as they do in solu-
tion heat treating.
Good temperature control and uniformity
throughout the furnace and load are re-
quired for all precipitation heat treating.
Recommended temperatures are generally
those that are least critical and that can be
used with practical time cycles. Except for
7xxx alloys in T7x tempers, these tempera-
tures generally allow some latitude and
should have a high probability of meeting
property specification requirements. Fur-
nace radiation effects seldom are trouble-
some except in those few furnaces that are
used for both solution and precipitation heat
treating. Generally, such situations should
be avoided, because the high heat capacity
needed for the higher temperatures may be
difficult to control at normal aging temper-
atures.
Soak time in precipitation heat treating is
not difficult to control; the specified times
carry rather broad tolerances. Heavier loads
with parts racked closer together, and even
nested, are not abnormal. The principal haz-
ard is undersoaking due to gross excesses in
loading practices. Some regions of the load
may reach soak temperature long after soak
866 / Heat Treating of Nonferrous Alloys
Elongation, % in
50 mm or 2 in.
Yield strength, MPa
Tensile strength, MPa
600
2024-T4
(not cold worked)
80
500
190 °C (375 °F).
70
205 °C (400 °F)
60
60
400
300
600
500
220 °C (425 °F) - 50
Tensile strength, ksi
Tensile strength, MPa
600
2024-T3
(cold worked 1 to 2%)
500
400
220 °C (425 °F)
80
190 °C (375 °F)
205 °C (400 °F)
70
60
50
50
Tensile strength, ksi
Tensile strength, MPa
600
2024-T36
(cold worked 5 to 6%)
190 °C (375 °F) 175 °C (350 °F)-80
000
500
400
F205 °C (400 °F)
70
220 °C (425 °F)
300
300
0
4
8
12
16
0
4
8
12
16
0
4
Aging time, h
Aging time, h
80
70
60
400
190 °C (375 °F)-
300
205 °C (400 °F)
40
220 °C (425 °F)
200
0
4
30
8
Aging time, hi
Yield strength, ksi
Yield strength, MPa
600
500
600
80
80
400
190 °C (375 °F)
205 °C (400 °F)
220 °C (425 °F)
T
70
60
50
300
30
200
12
16
0
4
190 °C (375 °F)
205 °C (400 °F)
220 °C (425 °F)
0
4
8
Elongation, % in
50 mm or 2 in.
30
8
Aging time, h
205 °C (400 °F)
50
Yield strength, ksi
Yield strength, MPa
- 40
8
Aging time, h
T
12
16
190 °C (375 °F)
175 °C (350 °F) - 80
500
70
205 °C (400 °F)
60
400
220 °C (425 °F)
300
T
50
40
30
30
200
12
16
0
4
8
12
16
Aging time, h
190 °C (375 °F)
220 °C (425 °F)
12
16
0
4
8
Elongation, % in
50 mm or 2 in.
30
190 °C (375 °F)
175 °C (350 °F)
205 °C (400 °F)
220 °C (425 °F)
12
16
0
4
8
12
16
Aging time, h
60
60
50
50
Yield strength, ksi
Aging time, h
Aging time, h
Fig 22(d) Effects of cold work after quenching and before aging on tensile properties of alloy 2024 sheet
time has been called. Placement of load ther-
mocouples is critical, and limiting the size and
spacing of a load may be necessary for aging
Aging temperature, °C
225
200
175
60
150
to the T73 and T76 tempers. As discussed
above, soak time is not as critical for peak-
aged (T6 and T8) tempers.
360
385
400
410
MPa
440
470
350
500
T73 aging
525
550
300
64
68
125
ksi
72-
76
80
100
0.01
0.1
1
10
Aging time, h
250
T6 aging
100
1000
Hardening of Cast Alloys
In general, the principles and procedures
for heat treating wrought and cast alloys are
similar. The major differences between so-
lution-treating conditions for castings and
those for wrought products are found in
soak times and quenching media. Solution
of the relatively large microconstituents
present in castings requires longer soaking
Aging temperature, °F
Yield strength, MPa
700
600
Longitudinal
500
Short transverse
Predicted from Eq 3
о
• Actual
400
2
4 6 8 10
Equivalent aging time
100
90
70
88
80
]
60
20 30 40 60
at 165 °C (325 °F), h
Yield strength, ksi
Fig 23 Iso-yield-strength curves for alloy 7075
Fig 24
Actual versus predicted yield strengths for
alloy 7050 extrusions
Tensile strength, ksi
100
Aging temperature, °F
310
320
330
50
0
Effect on yield strength, MPa
50
-50
Compared for standard
-100
conditions of 29 h
at 165 °C
340
10
10
5
C
G
Effect on yield strength, ksi
Tear strength/Yield strength
-15
-150
-20
150
155
160 165 170 175
Aging temperature, °C
Effect of aging temperature on yield strength
Fig 25 of alloy 7050-1736
periods than those used for wrought prod-
ucts (Table 3). When heat treatment of
castings must be repeated, solution times
become similar to those for wrought prod-
ucts, because the gross solution and homog-
enization has been accomplished and is
irreversible under normal conditions. Re-
duction of stresses and distortion from
quenching are also important, because cast-
ings generally are complex shapes with vari-
ations in section thickness.
Different casting processes and foundry
practices also result in microstructural dif-
ferences with relevance to heat-treatment
practice, because the coarser microstruc-
tures associated with slow solidification
rates require a longer solution heat treat-
ment exposure. Therefore, the time re-
quired at temperature to achieve solution is
progressively shorter for investment, sand,
and permanent mold castings. Foundry
practice (chills, gating, type of mold) also
plays an important role in the response of a
casting, or a portion of a casting, to heat
treatment. For example, thin-wall sand
castings produced with extensive use of
chills can often display finer microstruc-
tures than heavy-section permanent mold
parts produced in such a way that process
advantages are not exploited.
For these reasons, solution heat-treat-
ment practices can be optimized for any
specific part to achieve solution with the
shortest reasonable cycle once production
practice is finalized, even though most
foundries and heat treaters will standardize
a practice with a large margin of safety.
There also exists a fundamental difference
between unmodified and modified alloys in
which the size and shape of silicon crystals
are modified with additions of elements
such as calcium, sodium, strontium, or an-
timony. Modified alloys undergo rapid
spheroidization while complete sphe-
roidization is not achieved in unmodified
alloys even after very long times. The prac-
tical implication is that shorter solution heat
treatment could be employed in fully mod-
Transverse yield strength, ksi
40 45 50 55 60 65 70
1.8
5000
T4
Fe Si
Cu Mg Mn
4.2 1.4 0.6 0.34 0.14
4500
1.6
T3
Stretched 5%
T7
4000
1.4
1.2
T6
Not stretched
1.0
0.8
250
Fig 26
3500
3000
2500
T7
T8)
2000
300 350 400 450 500
Transverse yield strength, MPa
Estimated K, MPav mm
Effect of stretching and aging on the tough-
ness and yield strength of 2024 sheet
ified castings. The microsegregation of sili-
con and magnesium is not severe in the
aluminum-silicon-magnesium casting al-
loys, and hence it takes only a short time to
homogenize the alloy and to place the
Mg2Si into solution.
Quenchants. Quenching of aluminum
castings is often done in boiling water or a
milder medium to reduce quenching stress-
es in complex shapes. A commercially im-
portant variety is a mixture of polyalkylene
glycol and water, which has no detrimental
effect on properties for thicknesses under
approximately 3.2 mm (0.125 in.). Quen-
chant additions can be made for the follow-
ing purposes:
• To promote stable vapor film boiling by
the deposition of compounds on the sur-
face of parts as they are submerged in the
quench solution
• To suppress variations in heat flux by
increasing vapor film boiling stability
through chemically decreased quench so-
lution surface tension
• To moderate quench rate for a given water
temperature
The key to the compromise between
goals involving property development and
the physical consequences of quenching is
uniformity of heat extraction, which is in
turn a complex function of the operable heat
extraction mechanism. Nucleate, vapor
film, and convective boiling occur with dra-
matically different heat extraction rates at
different intervals. Differences in section
thickness, load density, positioning, racking
methods, surface condition, and casting ge-
ometry also influence the results.
Property Development, Yield strength is
largely controlled by the limiting hardening-
element level, and tensile strength (in a
general sense) is related to the ductility at a
given yield strength. Ductility, however, is
controlled for a given yield strength by
soundness and microstructural fineness,
and is thus determined in the foundry and
Heat Treating of Aluminum Alloys / 867
not by the heat treater. This effect of casting
methods on property development is shown
in Fig 27. Because of the finer cast structure
and higher supersaturation of the more rap-
idly solidified permanent mold castings,
their tensile properties are superior to those
of sand castings of the same composition
similarly heat treated.
Tempers. Cast products of heat-treatable
aluminum alloys have the highest combina-
tions of strength, ductility, and toughness
when produced in T6-type tempers. Devel-
oping T6-type tempers in cast products re-
quires the same sequence of operations
employed in developing tempers of the
same type in wrought products-solution
heat treating, quenching, and precipitation
heat treating. Premium-quality casting spec-
ifications such as MIL-A-21180 can require
different strengths and ductility levels in the
same casting.
Among precipitation treatments unique to
castings are those resulting in the T5 and T7
tempers. The T5 temper is produced merely
by applying a precipitation treatment to the
as-cast casting, without previous solution
treatment. A moderate increase in strength
is achieved without warpage and subse-
quent straightening. High hardness and di-
mensional and strength stability at elevated
temperatures account for the almost univer-
sal use of materials in T5 tempers for pis-
tons and other engine parts. Some applica-
tions demand combinations of strength,
toughness, and dimensional stability that
cannot be met by heat treating to T5-, T6-,
or T8-type tempers. For these applications,
T7-type tempers are developed by solution
heat treating, quenching in a medium that
provides a moderate cooling rate, and then
precipitation heat treating at a temperature
higher than those used to develop T5-, T6-,
and T8-type tempers. Heat treating to T7-
type tempers results in lower strength than
that of material in T6- or T8-type tempers,
develops high ductility and toughness, and
carries precipitation far enough to minimize
further precipitation during service.
Stress Relief
Immediately after being quenched, most
aluminum alloys are nearly as ductile as
they are in the annealed condition. Conse-
quently, it is often advantageous to stress
relieve parts by working the metal immedi-
ately after quenching. Numerous attempts
also have been made to develop a thermal
treatment that will remove, or appreciably
reduce, quenching stresses. Normal precip-
itation heat-treating temperatures are gener-
ally too low to provide appreciable stress
relief. Exposure to higher temperatures (at
which stresses are relieved more effective-
ly) results in lower properties. However,
such treatments are sometimes utilized
when even moderate reduction of residual
stress levels is important enough so that
Tensile strength, MPa
868 / Heat Treating of Nonferrous Alloys
350
300
Permanent mold castings
Sand castings
350
50
150 °C (300°F)
40
250
205 °C (400)
°F)-
200
150
260 °C (500 °F)
20
Tensile strength, ksi
Tensile strength, MPa
300
250
200
150 °C (300 °F)
40
205 °C (400 °F)
30
150
260 °C (500 °F)
20
7079
50
Spar forging
Tensile strength, ksi
B
A
Planes of
saw cuts
-160 mm
B
A
Direction of
Yield strength, MPa
Elongation, %
in 50 mm or 2 in.
100
100
0
10
20
30
0
10
Aging time, h
250
200
30
-150 °C (300 °F)-
150
205 °C (400 °F)
260 °C (500 °F)
100
20
10
50
50
0
10
20
30
Aging time, h
2
1
260 °C (500 °F)
205 °C (400 °F)
10
Aging time, h
150 °C (300 °F).
20
30
Yield strength, ksi
Yield strength, MPa
Elongation, %
in 50 mm or 2 in.
Aging time, h
250
200
150 °C (300 °F)
150
150
100
50
1
20
30
-205 °C (400 °F):
260 °C (500 °F)
T
20
10
0
10
20
30
Aging time, h
Yield strength, ksi
Deflection, μm
Deflection, um
compressive deformation
7500
Saw cut A-A
T6 parallel
250
200
4500
T6 normal
150
100
1500
50
0
-1500
T652 parallel
T652 normal
-50
Length, 1.8 m
2500
Saw cut B-B
T6 parallel
80
T6 normal
1500
60
40
T652 normal
500
T652 parallel
20
0
-20
205 °C (400 °F)
260 °C (500 °F)
-500
-150 °C (300 °FH
10
20
30
Aging time, h
Fig 27 Comparison of the precipitation-hardening characteristics of 356.0-T4 sand and permanent mold
castings
some sacrifice in mechanical properties can
be accepted. The T7 temper for castings is a
typical example of this kind of treatment.
Mechanical Stress Relief. Deformation
consists of stretching (bar, extrusions, and
plate) or compressing (forgings) the product
sufficiently to achieve a small but controlled
amount (1 to 3%) of plastic deformation. If
the benefits of mechanical stress relieving
are needed, the user should refrain from
reheat treating.
Figure 28 illustrates the beneficial effect
of 3% permanent deformation in compres-
sion on a large forging.
These methods are most readily adapt-
able to mill and forge shop products and
require equipment of greater capacity than
that found in most manufacturing plants.
Application of these methods to die forgings
and extrusions usually requires construc-
tion of special dies and jaws. Stretching
generally is limited to material of uniform
cross section; however, it has been applied
successfully to stepped extrusions and to a
3 by 14 m (10 by 47 ft) aircraft wing skin
roll-tapered to a thickness range of 7.1 to
3.2 mm (0.280 to 0.125 in.).
Specific combinations of the supplemen-
tal digits are used to denote the tempers
produced when mechanical deformation is
used primarily to relieve residual stresses
induced during the quenching operation.
For products stress relieved by stretching,
the digits 51 follow the basic Tx designation
(T451, for example). For products stress
relieved by compressive deformation, the
supplementary digits are 52.
An additional digit is added to designa-
tions for extrusions: an added zero specifies
that the product has not been straightened
after final stretching; an added one indicates
that straightening may have been performed
after final stretching.
Effect of Precipitation Heat Treating on
Residual Stress. The stresses developed dur-
ing quenching from solution heat treatment
are reduced during subsequent precipitation
heat treatment. The degree of relaxation of
stresses is highly dependent upon the time
and temperature of the precipitation treat-
ment and the alloy composition. In general,
the precipitation treatments used to obtain
the T6 tempers provide only modest reduc-
tion in stresses, ranging from about 10 to
35%. To achieve a substantial lowering of
quenching stresses by thermal stress relax-
ation, higher-temperature treatments of the
T7 type are required. These treatments are
used when the lower strengths resulting
from overaging are acceptable.
Other thermal stress-relief treatments,
known as subzero treatment and cold stabi-
Length, -1.8 m
Deflection, 0.001 in.
Deflection, 0.001 in.
Effect of 3% permanent deformation in com-
Fig 28
pression (T652 treatment) on distribution of
stress in a large forging. Parallel and normal refer to
warpage directions with respect to the plane of the
saw cut.
lization, involve cycling of parts above and
below room temperature. The temperatures
chosen are those that can be readily ob-
tained with boiling water and mixtures of
dry ice and alcohol-namely, 100 and -73
°C (212 and -100 °F)-and the number of
cycles ranges from one to five. The maxi-
mum reduction in residual stress that can be
effected by these techniques is about 25%.
The maximum effect can be obtained only if
the subzero step is performed first, and
immediately after quenching from the solu-
tion-treating temperature while yield
strength is low. No benefit is gained from
more than one cycle.
A 25% reduction in residual stress is
sometimes sufficient to permit fabrication
of a part that could not be made without this
reduction. However, if a general reduction
is needed, as much as 83% relief of residual
stress is possible by increasing the severity
of the uphill quench—that is, more closely
approximating the reverse of the cooling-
rate differential during the original quench.
This may be accomplished by a patented
process that involves extending the subzero
step to 195 °C (-320 °F) and then very
rapidly uphill quenching in a blast of live
steam (Fig 29). The rate of reheating is
extremely critical, and therefore, to ensure
proper application of the steam blast, a
Plate
thickness
50 mm
Heat Treating of Aluminum Alloys / 869
Table 9 Reheating schedules for wrought aluminum alloys
The schedules given in this table normally will not decrease strength more than 5%.
Reheating time at a temperature of:
150 °C
165 °C
Alloy and temper
(300 °F)
(325 °F)
175 °C
(350 °F)
190 °C
(375 °F)
205 °C
(400 °F)
220 °C
(425 °F)
230 °C
(450 °F)
A B C D
Range of
residual
2014-T4
(a)
(a)
(a)
(a)
(a)
(a)
(a)
stress
MPa ksi
2014-T6
2-50 h
8-10 h
2-4 h
1/2-1 h
5-15 min
(b)
(b)
2024-T3, 2024-T4
(a)
(a)
A 28
B 86
4.0
12.4
2024-T81, 2024-T86
20-40 h
(a)
2-4 h
(a)
(a)
(a)
(a)
1 h
12 h
15 min
5 min
6061-T6, 6062-T6,
C 130
6063-T6
19.0
D 165
24.0
7075-T6, 7178-T6
100-200 h
10-12 h
50-100 h
1-2 h
8-10 h
1-2 h
1-2 h
1/2-1 h
1/2 h
5-10 min
15 min
(b)
5 min
(a)
(a) Reheating not recommended. (b) Bring to temperature
Fig 29
Treatment
A: Cooled to -195 °C, then uphill
quenched in a steam blast
B: Cooled to -75 ° C, then uphill
quenched in a steam blast
C: Cooled to -75 or -195 ° C, then
uphill quenched in boiling water
D: Standard specimen, quenched and
aged to T6 temper in conventional
manner with no further treatment
Effectiveness of various uphill quenching
treatments in reducing residual quenching
stresses in 2014 plate. Note: uphill quenching treat-
ments (single-cycle only) were applied from 1/2 to 11/2 h
after quenching from the recommended solution-
treating temperature. All specimens were aged to the
T6 temper after uphill quenching.
special fixture usually is required for each
part.
This process will not solve all problems of
warpage in machining. It may reduce
warpage internally but increase warpage of
the extreme outer layers, although in the
opposite direction (Fig 30). Also, the effect
of the altered residual-stress pattern on per-
formance must be evaluated carefully for
each part. This is particularly important for
parts subjected to cyclic loading or exposed
to corrosive environments such as marine
atmospheres, especially if the process is
introduced after the start of production and
original performance tests are not repeated.
Further disadvantages are the cost and haz-
ard involved in handling liquid nitrogen and
live steam.
Effects of Reheating
The precipitation characteristics of alumi-
num alloys must be considered frequently
during evaluation of the effects of reheating
on mechanical properties and corrosion re-
sistance. Such evaluations are necessary for
determining standard practices for manu-
facturing operations, such as hot forming
and straightening, adhesive bonding, and
paint and dry-film lubricant curing, and for
evaluating the effects of both short-term
and long-term exposure in elevated temper-
atures in service.
The stage of precipitation that exists in an
alloy at the time of reheating plays a signif-
icant role in the effects of reheating. Con-
sequently, it is extremely dangerous to re-
heat material in a solution heat-treated
temper without first carefully testing the
effects of such reheating. In one such test,
2024-T4 sheet was found to be very suscep-
tible to intergranular corrosion when sub-
jected to a 15-min drying operation at 150 °C
(300 °F) during the first 8 h after quenching;
no susceptibility was evident when the
same drying operation was performed more
than 16 h after quenching. In another test,
7075-W (0.2 to 600 h) bar and plate were
reheated for hot forming at 175 °C (350 °F)
for 20 min. Strengths after aging to the T6
temper were 10 to 15% lower than those for
standard 7075-T6. In contrast, similar re-
heating of T6 material for up to 1 h at 175 °C
(350 °F) produced no detrimental effect.
If reheating is performed on material in
the W or T4 condition, its effect can be
estimated from families of precipitation
heat-treating curves such as those present-
ed in Fig 22. Such curves can also be used
for reheating of precipitation heat-treated
material at the precipitation heat-treating
temperature. For reheating at other temper-
atures, other data may be needed (Fig 31).
The heat-treating and reheating curves may
be used as the bases for limitations on
reheating (Table 9).
Annealing
Annealing treatments employed for alu-
minum alloys are of several types that differ
in objective. Annealing times and tempera-
tures depend on alloy type as well as on
initial structure and temper.
Full Annealing. The softest, most ductile,
and most workable condition of both non-
heat-treatable and heat-treatable wrought
alloys is produced by full annealing to the
temper designated “O.” Strain-hardened
products in this temper normally become
recrystallized, but hot-worked products
may remain unrecrystallized. In the case of
heat-treatable alloys, the solutes are suffi-
ciently thoroughly precipitated to prevent
natural age hardening. A higher maximum
temperature than that used for stress-relief
annealing, controlled cooling to a lower
temperature, and additional holding time at
the lower temperature generally are em-
ployed.
For both heat-treatable and non-heat-treat-
able aluminum alloys, reduction or elimina-
tion of the strengthening effects of cold work-
ing is accomplished by heating at a
temperature from about 260 to about 440 °C
(500 to 825 °F). The rate of softening is
strongly temperature-dependent; the time re-
quired to soften a given material by a given
amount can vary from hours at low tempera-
tures to seconds at high temperatures.
If the purpose of annealing is merely to
remove the effects of strain hardening, heat-
ing to about 345 °C (650 °F) will usually
suffice. If it is necessary to remove the
hardening effects of a heat treatment or of
cooling from hot-working temperatures, a
treatment designed to produce a coarse,
widely spaced precipitate is employed. This
usually consists of soaking at 415 to 440 °C
(775 to 825 °F) followed by slow cooling (28
°C/h, or 50 °F/h, max) to about 260 °C (500
°F). The high diffusion rates that exist dur-
ing soaking and slow cooling permit maxi-
mum coalescence of precipitate particles
and result in minimum hardness.
As a result of this treatment, only partial
precipitation occurs in 7xxx alloys, and a
second treatment (soaking at 230 ± 6 °C, or
450 ± 10 °F, for 2 h) is required. When the
need arises for small additional improve-
ments in formability, cooling at 28 °C/h (50
°F/h) should be extended to 230 °C (450 °F),
and the material should be soaked at 230 °C
for 6 h. The effects of eliminating or pro-
longing the 230 °C second step on the duc-
tility of 7075-0 sheet are compared with the
standard treatment in Table 10.
In annealing, it is important to ensure that
the proper temperature is reached in all
portions of the load; therefore, it is common
to specify a soaking period of at least 1 h.
The maximum annealing temperature is
moderately critical; it is advisable not to
exceed 415 °C (775 °F), because of oxidation
and grain growth. The heating rate can be
critical, especially for alloy 3003, which
usually requires rapid heating for preven-
tion of grain growth. Relatively slow cool-
ing, in still air or in the furnace, is recom-
mended for all alloys to minimize distortion.
Typical annealing conditions used for some
alloys in common use are listed in Table 11.
870 / Heat Treating of Nonferrous Alloys
Deflection, m
-1300
-2600
2600
+1300
+100
25-by-25-mm (1-by-1-in.) bar
0
0
-100
-3900
-5200
7075
o Control specimen
Quenched from 75 °C -200
(-100 °F) to steam
A Quenched from liquid
nitrogen to steam
-6500
1
2
3
4
Deflection, 0.001 in.
Tensile strength, MPa
475
450
65
425
O 120 °C
60
400
• 135 °C
Reheating
A 150 °C
375
temperature
55
99
Tensile strength, ksi
425
60
400
375
55
350
50
O 120 °C
325
135 °C
Yield strength, ksi
A 150 °C
1
300
Elongation in
50 mm (2 in), %
Yield strength, MPa
10
0
10
100
103
Deflection, um
Deflection, um
+1300
50-by-50-mm
Tine number
50
+50
0
0
(2-by-2-in.) bar
-50
-1300
-2600
50
-100
7075
-150
-3900
。 Control specimen
• Quenched from -75 °C (-100 °F) to steam
-5200
2600
+1300
0
-1300
-2600
A Quenched from liquid nitrogen to steam
2
3
4
Tine number
-3900
1
2
3
-200
5
6
Deflection, 0.001 in.
1 2 3 4 5 6
100
75-by-75-mm (3-by-3-in.) bar
7075
o Control specimen
Quenched from -75 °C (-100 °F) to steam
A Quenched from liquid nitrogen to steam
4
5
Tine number
6
7
500
+50
-50
-100
A-150
8
Deflection, 0.001 in.
Effect of uphill quenching on deflection of tines. Six-tine specimen was machined from 50 by 50 mm
Fig 30
(2 by 2 in.) bar. Similar specimens machined from 25 by 25 mm (1 by 1 in.) and 75 by 75 mm (3 by 3 in.)
bars had four and eight tines, respectively.
Products that can be heated and cooled
very rapidly, such as wire, are annealed by
continuous processes that require a total
heating and cooling time of only a few
seconds. Continuous annealing of coiled
sheet is accomplished in a total time of a
few minutes. For these extremely rapid
operations, maximum temperature may ex-
ceed 440 °C (825 °F).
Although material annealed from the pre-
cipitation-hardened condition usually has
sufficient ductility for most forming opera-
Fig 31
Duration of reheating, days
Effects of reheating on tensile properties of
alclad 2024-T81 sheet
tions, this ductility often is slightly lower
than that of material that has not been
subjected to prior heat treatment—that is,
material annealed at the producing source.
Therefore, when maximum ductility is re-
quired, annealing of a previously heat-treat-
ed product is sometimes unsuccessful.
Partial Annealing. Annealing of cold-
worked non-heat-treatable wrought alloys to
obtain intermediate mechanical properties
(H2-type tempers) is referred to as partial
annealing or recovery annealing. Tempera-
tures used are below those that produce ex-
tensive recrystallization, and incomplete soft-
ening is accomplished by substructural
changes in dislocation density and rearrange-
ment into cellular patterns (polygonization).
Bendability and formability of an alloy an-
nealed to an H2-type temper generally are
significantly higher than those of the same
alloy in which an equal strength level is de-
veloped by a final cold-working operation
(H1-type temper). Treatments to produce H2-
type tempers require close control of temper-
ature to achieve uniform and consistent me-
chanical properties.
Figure 32 shows changes in yield strength
as functions of temperature and time for
sheet of two non-heat-treatable alloys (1100
and 5052) initially in the highly cold-worked
condition (H18 temper). From these curves,
it is apparent that, by selection of appropri-
ate combinations of time and temperature,
mechanical properties intermediate to those
of cold-worked and fully annealed material
can be obtained. It is also evident that yield
strength depends much more strongly on
temperature than on time of heating.
Stress-Relief Annealing. For cold-worked
wrought alloys, annealing merely to remove
the effects of strain hardening is referred to
as stress-relief annealing. Such treatments

Table 10 Effects of annealing treatments on ductility of 7075-0 sheet
Annealing
treatment
for thickness of:
Elongation in bending(c), %
in 50 mm (2 in.) for
thickness of:
Elongation in tension(a), % in 50 mm (2 in.)
Bend angle(b), degrees, for
thickness of:
0.5 mm
(0.020 in.)
1.6 mm
(0.064 in.)
2.6 mm
(0.102 in.)
1.6 mm
(0.064 in.)
2.6 mm
(0.102 in.)
1.6 mm
(0.064 in.)
Treatment 1(d)
12
12
Treatment 2(e)
14
12
14
82
73
48
Treatment 3(f)
16
91
92.5
58
56
14
16
76
84
2.6 mm
(0.102 in.)
50
57
60
(a) Uniform elongation of gridded tension specimens. (b) Bend angle at first fracture. (c) Elongation in bend test for 1.3 mm (0.05 in.)
gage spanning fracture. (d) Soak 2 h at 415 14 °C (775 ± 25 °F); furnace cool to 260 °C (500 °F) at 30 °C/h (50 °F/h); air cool. (c) Soak
2 h at 425 °C (800 °F), air cool; soak 2 h at 230 °C (450 °F), air cool. (f) Soak 1 h at 425 °C (800 °F); furnace cool to 230 °C (450 °F) at
30 °C/h (50 °F/h); soak 6 h at 230 °C (450 °F), air cool
employ temperatures up to about 345 °C
(650 °F), or up to 400 ± 8 °C (750 ± 15 °F)
for 3003 alloy, and cooling to room temper-
ature. No appreciable holding time is re-
quired. Such treatment may result in simple
recovery, partial recrystallization, or full
recrystallization. Age hardening may follow
stress-relief annealing of heat-treatable al-
loys, however, because a concentration of
soluble alloying elements sufficient to cause
natural aging remains in solid solution after
such treatments.
A special form of stress-relief temper is
used for heat-treatable alloy products that
subsequently will be inspected ultrasonical-
ly. The product is heated to its normal
solution heat-treating temperature, then
cooled in still air to room temperature. This
temper is referred to as the O1 temper.
Controlled-Atmosphere Annealing and
Stabilizing. Aluminum alloys that contain
Table 11
even very small amounts of magnesium will
form a surface magnesium oxide unless the
atmosphere in the annealing furnace is free
of moisture and oxygen. Examples include
alloy 3004, which is used for cooking uten-
sils, and alloys of the 5xxx series.
Another problem that control of the an-
nealing atmosphere helps to overcome or
avoid is oil staining by oil-base roll lubri-
cants that do not burn off at lower annealing
temperatures. If the oxygen content of the
furnace atmosphere is kept very low during
such annealing, the oil will not oxidize and
stain the work.
Temperature control for full and partial
annealing is somewhat more critical than for
stress-relief annealing; the temperatures
and times specified are selected to produce
recrystallization and, in the case of heat-
treatable alloys, a precipitate of maximum
size; for this the cooling rate must be close-
Typical full annealing treatments for some common wrought aluminum alloys
These treatments, which anneal the material to the O temper, are typical for various sizes and methods of
manufacture and may not exactly describe optimum treatments for specific items.
Metal temperature
Alloy
°C
°F
Approximate
time at
temperature, h
Metal temperature
Alloy
°C
°F
Approximate
time at
temperature, h
1060
345
650
(a)
5457
345
650
(a)
1100
345
650
(a)
5652
345
650
(a)
1350
345
650
(a)
6005
415(b)
775(b)
2-3
2014
415(b)
775(b)
2-3
6009
415(b) 775(b)
2-3
2017
415(b)
775(b)
2-3
6010
415(b) 775(b)
2-3
2024
415(b)
775(b)
2-3
6053
415(b) 775(b)
2-3
2036
385(b)
725(b)
2-3
6061
415(b) 775(b)
2-3
2117
415(b)
775(b)
2-3
6063
415(b) 775(b)
2-3
2124
415(b)
775(b)
2-3
6066
415(b) 775(b)
2-3
2219
415(b)
775(b)
2-3
7001
415(c)
775(c)
2-3
3003
415
775
(a)
7005
345(d)
650(d)
2-3
3004
345
650
(a)
7049
415(c)
775(c)
2-3
3105
345
650
(a)
7050
415(c)
775(c)
2-3
5005
345
650
(a)
7075
415(c)
775(c)
2-3
5050
345
650
(a)
7079
415(c)
775(c)
5052
345
650
(a)
7178
415(c) 775(c)
2-3
2-3
5056
345
650
(a)
7475
415(c)
775(c)
2-3
5083
345
650
(a)
5086
345
650
Brazing sheet
5154
345
650
5182
345
650
(a)
5254
345
650
(a)
No. 11 and 12
345
650
5454
345
650
(a)
No. 21 and 22
345
650
5456
345
650
(a)
No. 23 and 24
345
650
(a)
(a)
(a)
(a) Time in the furnace need not be longer than necessary to bring all parts of the load to appealing temperature. Cooling rate is
unimportant. (b) These treatments are intended to remove the effects of solution treatment and include cooling at a rate of about 30 °C/h
(50 °F/h) from the annealing temperature to 260 °C (500 °F). Rate of subsequent cooling is unimportant. Treatment at 345 °C (650 °F),
followed by uncontrolled cooling, may be used to remove the effects of cold work or to partly remove the effects of heat treatment. (c)
These treatments are intended to remove the effects of solution treatment and include cooling at an uncontrolled rate to 205 °C (400 °F)
or less, followed by reheating to 230 °C (450 °F) for 4 h. Treatment at 345 °C (650 °F), followed by uncontrolled cooling, may be used
to remove the effects of cold work or to partly remove the effects of heat treatment. (d) Cooling rate to 205 °C (400 °F) or below is less
than or equal to 30 °C/h (50 °F/h).
Heat Treating of Aluminum Alloys / 871
ly controlled. Even allowing the load to cool
in the furnace may result in an excessively
high rate. Similarly, lowering the furnace-
control instrument by 28 °C (50 °F) each
hour may produce stepped cooling, which is
not satisfactory for severe forming opera-
tions. For maximum softening, a continu-
ous cooling rate of not more than 28 °C/h (50
°F/h) is recommended.
Annealing of castings for 2 to 4 h at
temperatures from 315 to 345 °C (600 to 650
°F) provides the most complete relief of
residual stresses and precipitation of the
phases formed by the excess solute retained
in solid solution in the as-cast condition.
Such annealing treatments provide maxi-
mum dimensional stability for service at
elevated temperatures. The annealed tem-
per is designated “O." (This temper was
designated "T2" prior to 1975.)
Grain Growth
Many of the aluminum alloys in common
use are subject to grain growth during solu-
tion treatment or annealing. This phenome-
non can occur during or after recrystallization
of material that has been subjected to a small
critical amount of prior cold work. It is usu-
ally manifested by surface roughening during
subsequent fabrication operations and fre-
quently results in rejections for appearance or
functional reasons. Less frequently, some de-
terioration of mechanical properties is en-
countered, and this is undesirable regardless
of surface-roughening effects.
Degree of susceptibility to grain growth
varies with alloy, structure, and chemical-
composition variation, and from one prod-
uct form to another. The critical range of
cold work is ordinarily about 5 to 15%.
Usually, temperatures of 400 °C (750 °F)
and above must be reached before grain
growth occurs, but some growth has been
encountered at temperatures as low as 345
°C (650 °F). Grain growth that occurs during
initial recrystallization is more a function of
composition, structure, and degree of cold
work than of temperature per se; tempera-
tures in excess of 455 °C (850 °F) in common
alloys can lead to secondary-recrystalliza-
tion grain-growth problems. The common
symptom indicating moderately large-grain
material is roughening or "orange peel" on
the external surfaces of bends. Severe
growth of grains to fingernail size and larger
sometimes is evident in parts made from
annealed (O temper) material by stretch
forming and then thermal treating or similar
operations. This type of grain growth often
is detected during subsequent anodizing,
etching, and chemical milling operations.
Cracking during welding or brazing is
another characteristic which may indicate
that severe grain growth has occurred. In
such instances, cracks propagate along
grain boundaries that provide little obstruc-
tion to their progress.
872 / Heat Treating of Nonferrous Alloys
200
“100 B
150
175 °C (350 °F)-
Yield strength, MPa
Yield strength, MPa
100
50
350
300
250
205 °C (400 °F)
230 °C (450 °F)
260 °C (500 °F)
288 and 316 °C (550 and 600 °F)
1
1
0
0.5
1
1.5
2
2.5
Time, h
3
3.5
4
4.5
5
20
10
0
▬▬▬▬▬ 5052-+18 7 50
175 °C (350 °F)
230 °C (450 °F)
200
260 °C (500 °F)
150
100
50
290 °C (550 °F)
☐ 315 °C (600 °F)
205 °C (400 °F)
0
0
0.5
1
1.5
2
2.5
3
3.5
4
4.5
Time, h
Fig 32 Representative isothermal annealing curves for alloys 1100-H18 and 5052-H18
If the surface roughening is objectionable
from either an appearance or a functional
aspect, the desirability of surface-smooth-
ing operations, such as sanding or buffing,
must be evaluated. If reductions in mechan-
ical properties are suspected, these must be
established by test and evaluated in relation
to the anticipated service.
In one application, a part that had been
made by stretch forming O-temper 2 mm
(0.080 in.) sheet and heat treating exhibited
significantly lower tensile and yield
strengths in portions where severe grain
growth had occurred than in portions hav-
ing normal grain size:
Yield strength
ksi
Tensile strength
Test
Grain
structure
MPa
ksi
MPa
Transverse
1
Coarse
265
38.5
247
35.8
2
Coarse
263
38.2
241
35.0
3
Fine
311
45.1
261
37.8
Longitudinal
1
Coarse
259
37.6
243
35.3
2
Coarse
269
39.0
245
35.6
3
Fine
305
44.2
270
39.1
5
40
40
30
20
10
Yield strength, ksi
Yield strength, ksi
In other similar investigations, no detri-
mental effects have been discovered, and in
many cases such parts have served satisfac-
torily in critical applications.
When a grain-growth problem is discov-
ered, it is too late to change the condition of
the parts in question, but several possible
methods are available for preventing recur-
rence of the difficulty. The simplest of these
is relieving the causative stress by interject
ing a stress-relief anneal into the manufac-
turing sequence immediately prior to the
solution-treating or full-annealing cycle in
which the grain growth occurred. This ap-
proach is usually successful and practical.
Another possibility is to adjust the amount
of stress present in the part immediately
prior to the critical heat treatment so that
the stress level is outside the critical range.
This may be done by adding a cold-working
operation before forming, such as pre-
stretching of blanks, or by forming in mul-
tiple stages with a stress-relief anneal before
each stage.
A third method that is sometimes suc-
cessful consists of increasing the heating
rate during the critical heat treatment by
reducing the size of furnace loads or by
changing from an air furnace to a salt bath.
paired
In one application, severe grain growth was
found during bending of alloy 1100 rectan-
gular tubing. The roughening of the inside
surfaces of the parts, which occurred during
forming of the large-grain material, im-
their functioning as radar
waveguides. Investigation disclosed that, to
minimize handling marks, the material was
procured in the strain-hardened (H14) tem-
per and was stress-relief annealed at 345 °C
(650 °F) immediately prior to forming. Grain
growth occurred during annealing as a re-
sult of the moderate amount of cold work
introduced at the mill. The problem was
eliminated by changing the stress-relieving
operation to a 5-min heating period in an air
furnace operating at 540 °C (1000 °F). The
explanation advanced for the success of this
treatment was that, due to the rapid heating
rate, the temperature of the material was
raised through the recrystallization range
for the less severely cold-worked grains
before the critically cold-worked grains had
time to grow appreciably.
Heating Equipment and
Accessories
The general methods for heat treating
aluminum alloys include the use of molten
salt baths, air-chamber furnaces, and induc-
tion heaters. The choice of heating equip-
ment depends largely on the alloy and the
configuration of the parts to be processed.
The type of heat treatment can also influ-
ence the choice of heating equipment. For
example, both molten salt baths and air-
chamber furnaces are suitable for solution
treating of aluminum alloys, while induction
heating requires additional analysis to de-
fine the proper temperature range for solu-
tion treatment. Induction methods can pro-
vide high heating rates, which affect
transformation behavior (see, for example,
the section "Nonequilibrium Melting" in
this article).
Molten salt baths and air-chamber furnac-
es both have advantages and disadvantages
in solution heat treatments, as discussed
below. Oil and gas-fired furnaces, in de-
signs that allow the products of combustion
to come in contact with the work, are
usually unsatisfactory because they pro-
mote high-temperature oxidation.
Salt baths heat the work faster (see Table
2) than air furnaces, provided that the
amount of work introduced at any one time
is controlled to prevent the temperature
from falling below the desired range. If the
temperature is permitted to fall below the
minimum limit, much of the advantage of
the salt bath is lost, because of the necessity
for reheating the large mass of salt.
Salt baths are also more readily adapted
to the introduction, at any time, of small
amounts of work requiring different soaking
Table 12 Frequency selection for induction heating with a longitudinal-flux coil and a
transverse-flux coil
Minimum part thickness, mm (in.), for a frequency of:
200 Hz
1 kHz
3 kHz
10 kHz
Heat Treating of Aluminum Alloys / 873
Production rate, tons/h
0
1
2
1800
3
4
5
6
7
Solution heat treat at 590 °C
1500
Full anneal at 425 °C
Partial anneal at 315 °C
1200
Material
60 Hz
Solenoid (longitudinal-flux) coil
Steel below Curie temperature
>38 (1.5)
13 (0.5)
Steel above Curie temperature
Brass
>175 (7.0)
100 (4.0)
5 (0.2)
43 (1.7)
2.3 (0.09)
25 (1.0)
>50 (2.0)
28 (1.1)
13 (0.5)
Aluminum
>38 (1.5)
22 (0.85)
9 (0.375)
7 (0.28)
5 (0.2)
1 (0.04)
13 (0.5)
4 (0.16)
3 (0.12)
Transverse-flux coil
Aluminum
>5 (0.2)
1.3 (0.05)
Brass
Steel
>10 (0.4)
>50 (2.0)
2.5 (0.1)
13 (0.5)
0.25 (0.01)
0.5 (0.02)
2.5 (0.1)
0.08 (0.003)
0.15 (0.006)
0.9 (0.035)
0.04 (0.0016)
0.08 (0.0032)
0.5 (0.020)
Power source, kW
900
600
300
0
0
5
10
15 20
Line speed, m/min
25
30
periods. (Economical utilization of air fur-
naces usually dictates accumulation of a
large load of parts of similar thickness be-
fore charging.) Also, the buoyant effect of
the salt reduces distortion during heating,
and the large reservoir of heat facilitates
temperature control and uniformity.
Salt bath operation entails special house-
keeping requirements. Dragout is costly and
unsightly. Because residual salt on parts
may result in corrosion, all salt must be
completely removed, including that from
crevices and blind holes. In addition, salt
residue from the quench water must be kept
to a minimum by a constant water overflow
or by providing a fresh-water rinse for all
parts after quenching. When these provi-
sions are impractical, corrosion can be in-
hibited by adding 14 g (½ oz) of sodium or
potassium dichromate to each 45 kg (100 lb)
of the molten salt.
Precautions. Molten salt baths are poten-
tially hazardous and require special precau-
tions. Operators must be protected from
splashing and dripping of the hot salt. Be-
cause heated nitrates are powerful oxidizing
agents, they must never by allowed to come
in contact with combustibles and reducing
agents, such as magnesium and cyanides.
Most authorities advise against inserting
aluminum alloys containing more than a few
percent of magnesium into molten nitrate.
To avoid exposure of personnel to nitrous
fumes produced during decomposition of
nitrates, good ventilation is essential.
When molten nitrates are being used, the
possibilities of explosions resulting from
both physical and chemical reactions must
be avoided. The former result from rapid
expansion of gases entrapped beneath the
surface of the bath. Hence, parts entering
the bath must be clean and dry; they must
also be free of pockets or cavities that
contain air or other gases. Chemical-reac-
tion explosions result from rapid break-
down of the nitrates due to overheating or
reaction with the pot material. Stainless
steel pots (preferably of type 321 or 347) are
more resistant to scaling than those made of
carbon steel or cast iron and therefore pre-
sent a lower probability of local overheat-
ing. Sludge or sediment accumulations in
bottom-heated pots can also lead to local
overheating. Overheat controls are essen-
tial to ensure against temperatures exceed-
ing 595 °C (1100 °F).
It is vitally important that water be kept
away from a nitrate tank. In controlling a
nitrate fire, do not use water or any fire
extinguisher containing water. The best ex-
tinguisher is dry sand, a supply of which
should be kept near the tank.
Extra sacks of salt should be stored in a
dry place, distant from the tank. If the fresh
salt being added to the bath is even slightly
damp, it should be added very slowly or
when the bath is frozen.
Air furnaces are used more widely than salt
baths because they permit greater flexibility
in operating temperature. When production
schedules and the variety of alloys requiring
heat treatment necessitate frequent changes
in temperature, the time and cost of adjusting
the temperature of a large mass of salt makes
the use of an air furnace almost mandatory.
However, waiting periods are often required
to allow the walls of air furnaces to stabilize at
the new temperature before parts are intro-
duced. Otherwise, parts may radiate heat to
colder walls or absorb radiant heat from hot-
ter walls, and the temperature indicated by
the control instrument will not reflect actual
metal temperature in the usual manner. Air
furnaces are also more economical when the
product mix includes a few rather large parts;
holding the temperature of a large volume of
salt in readiness for an occasional large part is
far more expensive than heating an equal
volume of air.
Induction heating with either solenoid
(longitudinal-flux) coils or transverse-flux
coils provides an efficient method for in-line
heating of flat-rolled products (particularly
strip). Solenoid coils create a longitudinal
flux, which can produce efficient heating for
thicker and/or lower resistivity materials.
Solenoid coils can also be used efficiently in
the heating of thinner magnetic material
(see, for example, steel below the Curie
temperature in Table 12).
For several nonferrous materials (alumi-
num, copper, brass), however, efficiency
and power factors with solenoid coils are
significantly lower than for ferrous materi-
Power requirement for transverse-flux induc-
Fig 33
tion heating of aluminum strip 1 mm (0.04 in.)
thick and 1270 mm (50 in.) wide. Source: Ref 9
als. Therefore, transverse-flux coils are ide-
ally suited for heating nonferrous materials,
because transverse-flux lines do not exhibit
the degree of current cancellation associat-
ed with longitudinal flux lines. This aspect
of transverse-flux coils improves efficiency
and also permits the use of lower frequen-
cies (Table 12). This reduces the capital
equipment costs, and where it shifts from
requiring RF frequencies, the power source
conversion efficiency is also significantly
improved. Aluminum, brass, copper, and
austenitic stainless steel strip lines are ide-
ally suited for transverse-flux heating. Each
of these materials often requires in-line pro-
cesses like partial or full annealing or solu-
tion heat treating to provide necessary me-
chanical properties for subsequent finishing
operations.
Transverse-flux induction heating offers
several benefits for in-line strip heating and
has been known for many years. However,
it requires specially designed iron-cored
laminated inductor coils and tighter control
of the power, strip handling, and process
parameters. Frequency selection is dictated
by the resistivity and thickness of the ma-
terial, while power requirements depend on
the production rates, the specific heat, and
the processing temperatures for a given
material. Figure 33 shows typical power
source requirements for transverse-flux
heating of aluminum strip.
Furnace Temperature Control
The importance of close temperature con-
trol in solution treating has been noted in
the previous section on solution treating.
Each control zone of each furnace should
contain at least two thermocouples. One
thermocouple, with its instrument, should
act as a controller, regulating the heat input;
the other should act independently as a
safety cutoff, requiring manual reset if its
set temperature (usually the maximum of
874 / Heat Treating of Nonferrous Alloys
the specified range) is exceeded during the
solution-treating cycle.
Safety cutoffs are mandatory for salt
baths to guard against explosions and often
have paid for themselves in air furnaces by
saving a load of parts or even the furnace
itself. It is important, however, that they be
tested periodically (by deliberately over-
shooting the empty furnace) to guard
against "frozen" corroded contacts result-
ing from prolonged periods of idleness.
At least one of the instruments for each
zone should be of the recording type, and
both instruments should have restricted
scales for instance, 400 to 600 °C (750 to
1110 °F), rather than 0 to 600 °C (32 to 1110
°F). This is required for maximum accuracy
because manufacturers' guarantees are
specified in terms of percent of scale.
In the placement of instruments, expo-
sure to extremes in ambient temperature,
humidity, vibration, dust, and corrosive
fumes should be avoided. Ambient temper-
atures between 5 and 50 °C (40 and 120 °F)
are satisfactory, but temperature changes of
6 °C/h (10 °F/h) or more should be avoided.
It is also essential that instruments and
thermocouple circuits be shielded from
electromagnetic fields commonly associat-
ed with the leads of high-amperage furnace
heating elements.
Temperature-sensing elements must be
capable of responding more rapidly to tem-
perature changes than the materials being
processed. Therefore, thermocouple wire
diameter should not exceed 1½ times the
thickness of the minimum-gage material to
be heat treated, and should in no case
exceed 14 gage. Thermocouples for salt
baths should be enclosed in suitable protec-
tion tubes. Air-furnace thermocouples
should be installed in open-end protection
tubes, with the thermocouple junction ex-
tending sufficiently beyond the tube to pre-
vent any loss in sensitivity.
Temperature-sensing elements should be
located in the furnace work chamber, not in
ducts and plenums, and should be as close
as possible to the working zone. Specifica-
tion MIL-H-6088C restricts distance be-
tween the sensing element and the working
zone to a maximum of 100 mm (4 in.). The
safety-cutoff thermocouple should be locat-
ed to reflect the highest temperature in the
working zone. The control thermocouple
should be located in a position where it will
read a temperature approximately halfway
between the hottest and coldest tempera-
tures.
Probe Checks. After the temperature-
measurement equipment is properly in-
stalled, it must be checked frequently for
accuracy. This is accomplished by inserting
a calibrated probe thermocouple into the
furnace adjacent to each furnace thermo-
couple and comparing its reading on a cali-
brated test potentiometer with that indicat-
ed by the furnace instrument. Correction
factors should be applied after each probe
check, but if the correction required ex-
ceeds ±3 °C (±5 °F), the source of the
deviation should be corrected. MIL-H-6088
recommends that this check be made week-
ly, but many operators make the check as
frequently as once each shift.
Temperature-Uniformity Surveys. In con-
trolling the temperature of parts that are
being heat treated it must first be deter-
mined that the temperature indicated by the
furnace instruments truly represents the
temperature of the nearby air or salt. Sec-
ond, the uniformity of temperature within
the working zone must be shown to be
within a range of 11 °C, or 20 °F (6 °C, or 10
°F, for precipitation heat treatment of alloy
2024). This is accomplished by measuring
the temperature at several test locations,
using calibrated test thermocouples and a
calibrated test potentiometer, and reading
furnace instruments nearly simultaneously.
MIL-H-6088 recommends monthly surveys
with one test location per 1.1 m³, or 40 ft³
(0.7 m³, or 25 ft³, for air furnaces on initial
survey), but with a minimum of nine test
locations distributed as shown in Fig 34.
Despite the large size of some furnaces,
rather surprising temperature uniformities
have been reported. In one instance the
initial survey of an air furnace measuring
12.5 by 1.2 by 3.0 m (41 by 4 by 10 ft)
showed maximum temperature variations of
+1.7, −1.1 °C (+3, −2 °F). When a parti-
tion 0.3 m (1 ft) thick was lowered, convert-
ing the furnace to two chambers 6.1 by 1.2
by 3.0 m (20 by 4 by 10 ft) each, the spread
was 1.1, 0.6 °C (+2, -1 °F) in one
section and +0.6, −1.1 °C (+1, −2 °F) in
the other.
For each furnace load, one thermocouple
(the "cold" couple) should be placed in the
coldest area of the furnace and another (the
"hot" couple) in the hottest area. In addi-
tion to these two thermocouples, a load
thermocouple should be installed. The load
couple should be of approximately the same
gage as the sheet or other product being
heat treated. If heavy plate, forgings, or
castings are being heat treated, a similar
discarded item should be used at the con-
trolling load couple. The thermocouple
should be placed in a drilled hole and
packed to hold it firmly in place during the
heat-treating cycle. In some instances, the
items being heat treated can be used as the
load couples. The thermocouples can be
placed in holes drilled in areas that will be
removed in making the finished article.
It is important that items of different
thicknesses-1 mm (0.040 in.) sheet and 25
mm (1 in.) plate, for example-not be heat
treated in the same furnace load.
In salt baths, uniformity surveys usually
are made by holding a probe thermocouple
in each location until thermal equilibrium is
reached; in air furnaces, a mock heat-treat-
ing cycle is required. First, the air furnace is
6
5
5-194
8
=
Rectangular furnace
福
Cylindrical salt bath
WINE
6 5 4
8
Cylindrical air furnace
Location of thermocouples for surveying
Fig 34
temperature uniformity in the working
zones of air furnaces and salt baths
stabilized at the test temperature. Then a
rack containing the test thermocouples is
inserted into the furnace. By using multiple
switches or a multipoint recording instru-
ment, all test thermocouples and furnace
instruments are read every 5 min. As the
temperature approaches the test range, it is
advisable to increase the frequency of read-
ings to detect possible overshooting. After
thermal equilibrium is reached, readings
should be continued until the recurrent tem-
perature pattern is established.
Surveys of salt baths generally are con-
sidered acceptable whether they are made
while the bath is empty or filled with work.
It is controversial whether surveys of air
furnaces should be made with or without a
load. Undoubtedly, recovery overshoots
are most likely to occur with a very light
load and would not be detected if a heavier
load were used. Certainly, if all loads are
essentially alike, surveys should be made
with typical loads. With widely varying
loads, the optimum approach is to make
several surveys initially, including one with
an empty furnace, and then to make suc-
ceeding surveys with an empty furnace to
ensure against changes in furnace charac-
teristics. If any changes are made in the
furnace that might affect temperature distri-
bution, such as repair of vanes or louvers,
several surveys should be repeated.
Another aspect of the problem of temper-
ature control in air furnaces is the necessity
of ensuring that the temperature of the parts
is the same as that of the surrounding air.
Furnace components whose temperature
differs from the air temperature must be
suitably shielded to prevent radiation to or
from the parts being heat treated. In a
furnace used for solution heat treating of
rivets, unshielded heating elements have
been known to produce part temperatures
as much as 20 °C (35 °F) higher than the
control temperature, resulting in eutectic
melting and cracking. In two other instanc-
es, reradiation through inadequate shielding
produced a radiation effect of as much as 11
°C (20 °F). One of these problems was
solved by painting the shield with reflective
aluminum paint and the other by adding a 13
mm (1/2 in.) thick layer of asbestos to the 1.6
mm (1/16 in.) stainless steel shield.
Furnace-wall temperatures that differ ap-
preciably from the temperature of the parts
also must be avoided. Consequently, when
the operating temperature of an air furnace
is changed, waiting periods are required
after the furnace instrument indicates sta-
bility, to allow the furnace walls to stabilize
at the new temperature. The magnitude of
this limitation is directly proportional to the
efficiency of the furnace as an insulated
chamber, but possibilities of such radiation
should be recognized even in thin-wall fur-
naces.
Radiation effects are potentially danger-
ous because they often cannot be detected
by ordinary thermocouples. Specially pre-
pared radiation panels with thermocouples
attached are used, and their readings are
compared with adjacent free thermocou-
ples. These panels normally are made of
material of the same gage as the thinnest
parts to be heat treated and should have a
single surface area of about 650 cm² (100
in.2). A thermocouple is attached to the
center of the panel by welding or peening.
In order to detect the maximum effect,
panel surfaces should be darkened so that
their emissivity is at least as high as that of
any material to be processed. During the
test, the panel surfaces should be parallel to
the suspected source or recipient of radia-
tion. As an example of the number of panels
required, several aerospace companies
specify one panel for every 1.5 linear meters
(5 linear feet) of furnace wall.
Instrument Calibration. All instruments
and thermocouples must be accurately cal-
ibrated, and it is essential that the calibra-
tions be traceable directly to the National
Bureau of Standards. The chain of trace-
ability should consist of not more than four
links for sensing elements and three links
for measuring elements. To illustrate, if the
article calibrated by the National Bureau of
Standards is called a primary standard, then
the chain of traceability of measuring ele-
ments should consist of primary standard,
test potentiometer, and furnace instrument.
Similarly, the chain for sensing elements
should consist of primary standard, second-
ary standard, test thermocouple, and fur-
nace thermocouple. Every effort should be
made to ensure that the temperature indi-
cated by the furnace instruments is as close
as possible to the actual temperature. To
achieve this, it is necessary to apply correc-
tion factors obtained during calibration to
the next lower echelon of accuracy. Even
then, if all errors inherent in the chain are in
the same direction, a considerable differ-
ence will exist between the measured and
actual temperatures. Therefore, it is advis-
able to operate as close to the mean of the
desired range as possible.
Dimensional Changes during
Heat Treatment
In addition to the completely reversible
changes in dimensions that are simple func-
tions of temperature change and are caused
by thermal expansion and contraction, di-
mensional changes of more permanent char-
acter are encountered during heat treat-
ment. These changes are of several types,
some of mechanical origin and others
caused by changes in metallurgical struc-
ture. Changes of mechanical origin include
those arising from stresses developed by
gravitational or other applied forces, from
thermally induced stresses or from relax-
ation of residual stresses. Dimensional
changes also accompany recrystallization,
solution, and precipitation of alloying ele-
ments.
Solution Heat Treatment. Distortion as a
result of creep during solution heat treat-
ment should be avoided by proper loading
of parts in baskets, racks, or fixtures, or by
provision of adequate support for long piec-
es of plate, rod, bar, and extrusions heat
treated in horizontal roller hearth furnaces.
Sheet is provided with air-pressure support
in continuous heat-treating furnaces to
avoid scratching, gouging, and distortion. If
parts are to be solution heat treated in
fixtures or racks made of materials (such as
steel) with coefficients of thermal expansion
lower than that of the aluminum being treat-
ed, allowance should be made for this dif-
ferential expansion to ensure that expansion
of the aluminum is not restricted. Straight-
ening immediately after solution heat treat-
ing may be preferable to fixturing.
Solution of phases formed by major alloy-
ing elements causes volumetric expansion
or contraction, depending on the alloy sys-
tem, and this may have to be taken into
account in heat treatment of long pieces.
For example, solution heat treatment and
quenching of alloy 2219 causes lengthwise
contraction of about 2 mm/m (0.002 in./in.).
Solution heat treatment and quenching of
alloys of the 7xxx series is accompanied by
lengthwise expansion-about 0.6 mm/m
(0.0006 in./in.) for alloy 7075 rod or plate.
Quenching. The most troublesome
changes in dimensions and shape are those
that occur during quenching or that result
from stresses induced by quenching. Due to
its nonuniform cooling, quenching may pro-
duce warpage or distortion, particularly in
Heat Treating of Aluminum Alloys / 875
Longitudinal stress range, MPa
500
400
300
Average cooling rate at center
of cylinders, °F/s
200
°
100
0
Poo
10
10
100
0
100
Average cooling rate at center
of cylinders, °C/s
1660
40
20
20
Longitudinal stress range, ksi
Effect of quenching rate on longitudinal
Fig 35
stress ranges in alloy 2014-T4 cylinders
quenched in various media. Cylinders were 75 mm (3
in.) in diameter by 230 mm (9 in.) long. Cooling rate
was measured from 400 to 290 °C (750 to 555 °F). Stress
range is maximum tensile stress plus maximum com-
pressive stress.
thin material and in thin sections of parts
that contain variations in thickness. For
thick-section products or parts, changes in
external shape may be small because of
rigidity, but the interior-to-surface temper-
ature gradients that form with rapid cooling
create residual stresses; these stresses nor-
mally are compressive at the surfaces and
tensile in the interior.
As previously discussed, warpage or dis-
tortion of thin-section material can be re-
duced by using a quenching medium that
provides slower cooling; however, cooling
must be sufficient to produce the required
properties. Slower quenching can also re-
duce the magnitude of residual stresses in
thicker parts or pieces, as shown in Fig 9 for
cylindrical specimens of alloy 6151
quenched in cold or boiling water. Stress
range (maximum tensile stress plus maxi-
mum compressive stress) for a cylinder with
a radius of 89 mm (3.5 in.) is about 205 MPa
(30 ksi) when the cylinder is quenched in
cold water but less than 70 MPa (10 ksi)
when it is quenched in boiling water. The
effects of average cooling rate through the
temperature range from 400 to 290 °C (750
to 550 °F) on longitudinal stress ranges
developed in alloy 2014 cylinders 75 mm (3
in.) in diameter are shown in Fig 35.
High stresses induced by rapid quenching
generally are reduced only modestly by the
precipitation heat treatments used to pro-
duce T6- or T8-type tempers. Consequent-
ly, for the alloys that require rapid cooling
to develop the properties of these tempers,
those incorporating mechanical stress relief
(Tx51, Tx52) usually are specified when
substantial metal must be removed to pro-
duce final shapes. Other T8-type tempers,
such as T86 and T87, also have low residual
stress as a result of the stretching required
to produce them.
Heat Treatments for Precipitation and Sta-
bilization. The most significant dimensional
changes associated with precipitation heat
876 / Heat Treating of Nonferrous Alloys
treatments and stabilizing heat treatments
arise from concurrent dilution of the solid
solution (which changes lattice parameter)
and formation of precipitate. Changes in
density and specific volume resulting from
these changes in metallurgical structure are
the reverse of those caused by solution of
the alloy phases. However, because the
strongest tempers are those in which the
precipitate is present in nonequilibrium
transition forms, the amount of change dur-
ing precipitation heat treatment does not
totally compensate for the previous (and
opposite) change that occurred during solu-
tion heat treatment. Most of the heat-treat-
able alloys expand (grow) during precipita-
tion heat treatment. Exceptions are alloys
of the 7xxx wrought series and the 7xx.0
casting series, which exhibit contraction.
In alloys of the 2xxx series, the amount of
growth decreases with increasing magne-
sium content. Thus, growth of about 1.5
mm/m (0.0015 in./in.) can be expected dur-
ing precipitation heat treatment of alloy
2219-T87, about 0.5 mm/m (0.005 in./in.) for
treatment of alloy 2014-T6 and less than 0.1
mm/m (0.0001 in./in.) for treatment of alloy
2024-T851. Alloys 7050 and 7075, on the
other hand, contract about 0.3 mm/m
(0.0003 in./in.) on precipitation heat treating
from the W temper to the T6 temper and
about 0.7 mm/m (0.0007 in./in.) on treating
from the W temper to the T73 temper.
Stabilizing T7-type treatments cause great-
er amounts of growth than the T5-, T6-, or
T8-type treatments for the same alloys. This
increased growth is associated either with
formation of increased amounts of transition
precipitates or with transformation of transi-
tion precipitates to equilibrium phases.
Dimensional Stability in Service
Dimensional stability of heat-treated
parts in service depends on alloy, temper,
and service conditions. Of the latter, ex-
cluding mechanical conditions such as ap-
plied loads, the most important is service
temperature range relative to the range in
which precipitation occurs. Residual stress-
es constitute another source of dimensional
changes. Stress relief minimizes changes
due to residual stresses, and most mill prod-
ucts usually are supplied in tempers that
include stress relief. Potential dimensional
change as a result of further precipitation in
parts that operate at elevated temperatures
is minimized for wrought products by use of
T7-type stabilizing treatments and for cast-
ings by use of T5-type treatments. Howev-
er, components of high-precision equip-
ment, such as instruments for aerospace
guidance systems and optical and telescopic
devices, may require special supplementary
treatments during manufacture to further
reduce stresses or subsequent precipitation.
(These treatments are discussed below, un-
der "Stability of Precision Equipment.")
The T3- and T4-type tempers are the least
stable dimensionally because of possible
precipitation in service. Alloys 2024 and its
variants have the smallest dimensional
change in aging; the total change from the
quenched to the average state is of the order
of 0.06 mm/m (0.00006 in./in.), less than the
change due to a temperature variation of 3
°C (5 °F). These alloys therefore can be used
in the T3- and T4-type tempers, except for
precision equipment. For all other alloys,
T6- or T8-type tempers should be used,
because in these tempers all the alloys have
good dimensional stability.
% of tests
Yield strength, ksi
60
65
50
T
1
40
180 specimens
from a single sheet
30
de 20
10
70
T
7075-T6
Alclad sheet
4290 routine
mill tests
75
400
425
450
Yield strength, MPa
475
500
525
in heat-treated 7075-T6 clad sheet product
with distribution in a single sheet. A is 95% probability
that not more than 1% of all material will fall below this
value; B is 95% probability that not more than 10% of
all material will fall below this value. (A and B refer
only to curve representing 4290 routine mill tests.)
Stability of Precision Equipment. Proper
maintenance of high-precision devices,
such as gyros, accelerometers, and optical
systems, requires use of materials in which
dimensional changes from metallurgical in-
stability are limited from 10 µm/m (10 μin./ Fig 36 Comparison of distribution of yield strength
in.). Several laboratory investigations and
considerable practical experience have
shown that wrought alloys 2024 and 6061
and casting alloy 356.0 are well suited and
generally preferred for such applications.
Dimensional changes were no greater than
10 μm/m when alloys 2024-T851 and -T62,
6061-T651 and -T62, and 356.0-T51, -T6,
and -T7 were tested for more than a year at
room temperature and for several months at
70 °C (160 °F), and then the same alloys
were tested with repeated thermal cycling
between 20 and -70 °C (68 and -94 °F).
Because stresses applied or induced by
acceleration in such devices generally are
not high, strength levels lower than those of
the highest-strength tempers frequently are
satisfactory. To increase precision of ma-
chining to intended dimensions, as well as
to promote maximum stability, it is com-
mon practice to apply additional thermal
treatments for stress relief and precipitation
of 1 to 2 h at temperatures of 175 to 205 °C
(350 to 400 °F) after rough machining. These
additional treatments sometimes are repeat-
ed at successive stages of processing, and
even after final machining. In addition, it
has been claimed that one or two cyclic
treatments consisting of cooling to -100 °C
(-150 °F), holding for 2 h, heating to 232 to
240 °C (450 to 465 °F) and again holding for
2 h can improve dimensional stability of
356-T6 castings.
Quality Assurance
Quality-assurance criteria that heat-treat-
ed materials must meet always include min-
imum tensile properties and, for certain
alloys and tempers, adequate fracture
toughness and resistance to detrimental
forms of corrosion (such as intergranular or
exfoliation attack) or to stress-corrosion
cracking. All processing steps through heat
treatment must be carefully controlled to
ensure high and reliable performance.
Tensile Tests. In general, the relatively
constant relationships among various prop-
erties allow the use of tensile properties
alone as acceptance criteria. The minimum
guaranteed strength is ordinarily that value
above which it has been statistically pre-
dicted with 95% probability that 99% or
more of the material will pass. The inherent
variability within lots and among specimens
from a given piece is shown in Fig 36.
Testing provides a check for evidence of
conformance; process capability and pro-
cess control are the foundations for guaran-
teed values.
Published minimum guaranteed values
are applicable only to specimens cut from a
specific location in the product, with their
axes oriented at a specific angle to the
direction of working as defined in the appli-
cable procurement specification. In thick
plate, for example, the guaranteed values
apply to specimens taken from a plane
midway between the center and the surface,
and their axes parallel to the width dimen-
sion (long transverse). Different properties
should be expected in specimens taken from
other locations, or in specimens whose axes
were parallel to thickness dimension (short
transverse). However, the specified "refer-
ee" locations and orientations do provide a
useful basis for lot-to-lot comparisons, and
constitute a valuable adjunct to other pro-
cess-control measures.
Tensile tests can be used to evaluate the
effects of changes in the process, provided
specimens are carefully selected. A varia-
tion in process that produces above-mini-
mum properties on test specimens, howev-
er, is not necessarily satisfactory. Its
acceptability can be judged only by compar-
ing the resulting properties with those de-
veloped by the standard process on similar-
ly located specimens. Finally, variations in
heat-treating procedure are likely to affect
Heat Treating of Aluminum Alloys / 877
600
Table 13 Typical acceptable hardness values for wrought aluminum alloys
Acceptable hardness does not guarantee acceptable properties; acceptance should be based on acceptable
hardness plus written evidence of compliance with specified heat-treating procedures. Hardness values
higher than the listed maximums are acceptable provided that the material is positively identified as the
correct alloy.
550
85
80
80
75
T
T
70
65
60
55
50
45
80
90
100
70
Hardness, HRB
Tensile strength versus hardness for various
aluminum alloys and tempers
• Use a specimen that has at least 19 cm² (3
in.2) of surface area
Hardness
Alloy and temper
Product form(a)
HRB
HRE
HRH
HRIST
2014-T3, T4, -T42
2014-T6, -T62, -T65
All
65-70
87-95
Sheet(b)
80-90
103-110
All others
81-90
104-110
2014-T61
2024-T3
All
100-109
Not clad(c)
69-83
97-106
111-118
82.5-87.5
Tensile strength, MPa
500
450
400
Clad, ≤1.60 mm (0.063
in.)
52-71
91-100
109-116
80-84.5
Clad, 1.60 mm (0.063
in.)
52-71
93-102
109-116
350
2024-T36
All
76-90
2024-T4, -T42(d)
Not clad
69-83
100-110
97-106
111-118
85-90
82.5-87.5
Clad, 1.60 mm (0.063
in.)
300
52-71
91-100
109-116
80-84.5
50
60
Clad, >1.60 mm (0.063
in.)
52-71
93-102
109-116
2024-T6, -T62
All
74.5-83.5
99-106
2024-T81
Not clad
74.5-83.5
99-106
84-88
84-88
Fig 37
Clad
99-106
2024-T86
All
83-90
105-110
87.5-90
6053-T6
All
79-87
74.5-78.5
6061-T4(d)
Sheet
60-75
88-100
64-75
Extrusions; bar
70-81
82-103
67-78
•
Not clad, 0.41 mm
6061-T6
(0.016 in.)
75-84
Not clad, ≥0.51 mm
(0.020 in.)
47-72
85-97
Clad
84-96
6063-T5
All
55-70
89-97
62.5-70
6063-T6
All
70-85
6151-T6
All
91-102
7075-T6, -T65
Not clad(e)
85-94
106-114
Clad:
≤0.91 mm (0.036 in.)
102-110
>0.91 1.27 mm
(>0.036 0.050 in.)
78-90
104-110
>1.27 1.57 mm
(>0.050 ≤ 0.062 in.)
76-90
104-110
>1.57 1.78 mm
(>0.062≤ 0.070 in.)
76-90
102-110
>1.78 mm (0.070 in.)
73-90
102-110
7079-T6, -T65
7178-T6
All(e)
81-93
104-114
Not clad(f)
85 min
105 min
Clad:
≤0.91 mm (0.036 in.)
102 min
>0.91 1.57 mm
(>0.036≤ 0.062 in.)
>1.57 mm (0.062 in.)
85 min
88 min
78-84
87.5-92
86-90
87.5-92
88 min
86 min
(a) Minimum hardness values shown for clad products are valid for thicknesses up to and including 2.31 mm (0.091 in.); for heavier-gage
material, cladding should be locally removed for hardness testing or test should be performed on edge of sheet. (b) 126 to 158 HB (10
mm ball, 500 kg load). (c) 100 to 130 HB (10 mm ball, 500 kg load). (d) Alloys 2024-T4, 2024-T42 and 6061-T4 should not be rejected for
low hardness until they have remained at room temperature for at least three days following solution treatment. (e) 136 to 164 HB (10
mm ball, 500 kg load). (f) 136 HB min (10 mm ball, 500 kg load)
the relationships among tensile properties
and other mechanical properties. In appli-
cations where other properties are more
important than tensile properties, the other
properties should be checked also.
Hardness tests are less valuable for ac-
ceptance and rejection of heat-treated alu-
minum alloys than they are for steel. Nev-
ertheless, hardness tests have some utility
for process control. Typical hardness val-
ues for various alloys and tempers are given
in Table 13. Figure 37 shows the general
relationship between longitudinal tensile
strength and hardness for aluminum alloys.
Intergranular-Corrosion Test. The extent
of precipitation during elevated-temperature
aging of alloys 2014, 2219, and 2024 markedly
influences the type of corrosion attack and
the corrosion resistance. With thin-section
products quenched at rates sufficiently rapid
to prevent precipitation in the grain bound-
aries during the quench, short periods of
precipitation heat treating produce localized
grain boundary precipitates adjacent to the
depleted areas, producing susceptibility to
intergranular corrosion. Additional heating,
however, induces extensive general precipi-
tation within the grains, lowering the corro-
sion potential differences between the grains
and the boundary areas, thus removing the
cause of the selective corrosion.
The most common test for susceptibility
to intergranular corrosion is carried out as
follows:
•
Remove any cladding by filing or etching
Clean the specimen by immersing it for 1
min in a solution containing 5% concen-
trated nitric acid and 0.5% hydrofluoric
acid at a temperature of 95 °C (200 °F);
rinse in distilled water. Immerse for 1 min
in concentrated nitric acid at room tem-
perature; rinse in distilled water
• Immerse the specimen for 6 h in a freshly
prepared solution containing 57 g of sodi-
um chloride and 10 mL of 30% hydrogen
peroxide per liter of water at a tempera-
ture of 30 ± 5 °C (86 ± 9 °F). More than
one specimen may be corroded in the
same container provided that at least 4.6
mL of solution is used for each square
centimeter (30 mL/in.²) of specimen sur-
face and that the specimens are electrical-
ly insulated from each other
• After the immersion period, wash the
specimen with a soft-bristle brush to re-
move any loose corrosion product. Cut a
cross-sectional specimen at least 19 mm
(3/4 in.) long through the most severely
corroded area; mount and metallographi-
cally polish this specimen
Examine the cross-sectional specimen mi-
croscopically at magnifications of 100×
and 500× both before and after etching
with Keller's reagent
•
• Describe the results of the microscopic
examination in terms of the five degrees
of severity of intergranular attack illus-
trated in Fig 38
Electrical Conductivity. For control of the
corrosion and stress-corrosion characteris-
tics of certain tempers, notably the T73 and
T76 types, the materials must meet combi-
nation criteria of yield strength plus electri-
cal conductivity. Although these criteria are
based on indirect measurements of proper-
ties, their validity for ensuring the intended
corrosion and stress-corrosion resistance
Tensile strength, ksi
878 / Heat Treating of Nonferrous Alloys
Sheet and plate
Sheet, plate, extrusions, and
forgings
Alloy
Condition
Product form
2048
2124
2419
(a)
(b)
7049
7050
** EE
T8
Sheet and plate
T3, T8
T8
T7
T7
7150
T6
7175
T6, T7
(c)
(d)
7475
T6, T7
(e)
Five degrees of severity of intergranular at-
Fig 38
tack. Severity of intergranular attack (sche-
matic), as observed microscopically in transverse sec-
tions after test for susceptibility to intergranular
corrosion. Top of each area shown in surface exposed
to corrosive solution
has been firmly established by extensive
correlation and testing.
Low tensile strengths may be accompa-
nied by high levels of electrical conductiv-
ity, so electrical conductivity is sometimes
used as a quality-assurance diagnostic
tool. However, because the correlation
between strength and electrical conductiv-
ity is strongly a function of chemical com-
position and fabricating practice, use of
electrical conductivity is not recommend-
ed except for rough screening. This
screening must be followed by hardness
testing, and then by tensile testing if the
hardness tests indicate that the heat treat-
ment was suspect.
Fracture Toughness Indices. Fracture
toughness is rarely, if ever, a design con-
sideration in the 1000, 3000, 4000, 5000,
and 6000 series alloys. The fracture tough-
ness of these alloys is sufficiently high that
thicknesses beyond those commonly pro-
duced would be required to obtain a valid
test.
Fracture toughness is a meaningful de-
sign-related parameter for some conven-
tional high-strength alloys and all the con-
trolled-toughness, high-strength alloys.
Conventional aerospace alloys for which
fracture toughness minimums may be use-
ful in design include 2014, 2024, 2219,
7075, and 7079. These alloys have tough-
ness levels that are inferior to those of
their controlled-toughness counterparts.
Consequently, these products are not used
in fracture-critical applications, although
fracture toughness can be a meaningful
design parameter. Fracture toughness is
not guaranteed in conventional high
strength alloys.
Fracture toughness quality control and
material procurement minimums are appro-
priate for controlled-toughness, high-
strength alloys. The alloys and tempers
currently identified as controlled-tough-
ness, high-strength products include:
Plate, forgings, and extrusions
Sheet, plate, forgings, and
extrusions
Sheet and plate
Sheet, plate, forgings, and
extrusions
Sheet and plate
The fracture toughness of these alloys and
tempers range in measured K₁c values from
about 20 MPa√√√m (18 ksi√in.) upward. Con-
trolled-toughness alloys are often derivatives
of conventional alloys. For example, 7475
alloy is a derivative of 7075 with maximum
compositional limits on some elements that
were found to decrease toughness.
In products of the newer controlled-
toughness high-strength alloys 2090, 2091,
2124, 2224, 2324, 7050, 7149, 7150, 7175,
7475, and 8090, which provide guaranteed
levels of fracture toughness, minimum val-
ues of the applicable indices, K₁c or Kc,
are established by accumulation of statis-
tical data from production lots as a basis
for guaranteed minimum values. If the
minimum specified fracture toughness val-
ue is not attained, the material is not
acceptable. Some specifications allow use
of less-expensive screening tests (such as
the notch tensile or chevron-notched short
bar) as a basis for release of high-tough-
ness alloy products. In these instances,
correlations between K₁ and the screening
test result is used to establish the appro-
priate notch-yield ratio as a lot-release
criterion.
Temper Designations for
Heat-Treatable Aluminum Alloys
The temper designations used in the Unit-
ed States for heat-treatable aluminum alloys
are part of the system that has been adopted
as an American National Standard (ANSI
H35.1). Used for all wrought and cast prod-
uct forms except ingot, the system is based
on the sequences of mechanical or thermal
treatments, or both, used to produce the
various tempers. The temper designation
follows the alloy designation and is separat-
ed from it by a hyphen. Basic temper des-
ignations consist of individual capital let-
ters. Major subdivisions of basic tempers,
where required, are indicated by one or
more digits following the letter. These digits
designate specific sequences of treatments
that produce specific combinations of char-
acteristics in the product. Variations in
treatment conditions within major subdivi-
sions are identified by additional digits. The
conditions during heat treatment (such as
time, temperature, and quenching rate)
used to produce a given temper in one alloy
may differ from those employed to produce
the same temper in another alloy.
Designations for the common heat-treat-
ed tempers, and descriptions of the se-
quences of operations used to produce
those tempers, are given in the following
paragraphs. (For the entire aluminum alloy
temper designation system, including desig-
nations for non-heat-treatable alloys, see
Volume 2 of this Metals Handbook series.)
Basic temper designations for heat-treated
conditions include the codes O, W, and T.
Other basic temper designations are F (as
fabricated) and H (strain hardened).
O, annealed. Applies to wrought prod-
ucts that are annealed to obtain lowest
strength temper and to cast products that
are annealed to improve ductility and di-
mensional stability. The O may be followed
by a digit other than zero.
W, solution heat treated. An unstable
temper applicable to any alloy that naturally
ages (spontaneously ages at room tempera-
ture) after solution heat treatment. This
designation is specific only when the period
of natural aging is indicated-for example,
Wh. (See also the discussion of the Tx51,
Tx52, and Tx54 tempers, in the section
below on subdivision of the T temper.)
T, heat treated to produce stable tempers
other than O. Applies to products that are
thermally treated, with or without supple-
mentary strain hardening, to produce stable
tempers. The T is always followed by one or
more digits, as discussed below.
Major Subdivisions of T Temper. In T-type
designations, the T is followed by a number
from 1 to 10; each number denotes a specif-
ic sequence of basic treatments, as de-
scribed below.
TI, cooled from an elevated-temperature
shaping process and naturally aged to a
substantially stable condition. Applies to
products that are not cold worked after an
elevated-temperature shaping process such
as casting or extrusion, and for which me-
chanical properties have been stabilized by
room-temperature aging. If the products are
flattened or straightened after cooling from
the shaping process, the effects of the cold
work imparted by flattening or straightening
are not recognized in specified property
limits.
T2, cooled from an elevated-temperature
shaping process, cold worked, and natural-
ly aged to a substantially stable condition.
Applies to products that are cold worked
specifically to improve strength after cool-
ing from a hot-working process such as
rolling or extrusion, and for which mechan-
ical properties have been stabilized by
room-temperature aging. The effects of cold
work, including any cold work imparted by
flattening or straightening, are recognized in
specified property limits.
T3, solution heat treated, cold worked,
and naturally aged to a substantially stable
condition. Applies to products that are cold
worked specifically to improve strength af-
ter solution heat treatment, and for which
mechanical properties have been stabilized
by room-temperature aging. The effects of
cold work, including any cold work impart-
ed by flattening or straightening, are recog-
nized in specified property limits.
T4, solution heat treated and naturally
aged to a substantially stable condition.
Applies to products that are not cold
worked after solution heat treatment, and
for which mechanical properties have been
stabilized by room-temperature aging. If the
products are flattened or straightened, the
effects of the cold work imparted by flatten-
ing or straightening are not recognized in
specified property limits.
T5, cooled from an elevated-temperature
shaping process and artificially aged. Ap-
plies to products that are not cold worked
after an elevated-temperature shaping pro-
cess such as casting or extrusion, and for
which mechanical properties or dimensional
stability, or both, have been substantially
improved by precipitation heat treatment. If
the products are flattened or straightened
after cooling from the shaping process, the
effects of the cold work imparted by flatten-
ing or straightening are not recognized in
specified property limits.
T6, solution heat treated and artificially
aged. Applies to products that are not cold
worked after solution heat treatment, and
for which mechanical properties or dimen-
sional stability, or both, have been substan-
tially improved by precipitation heat treat-
ment. If the products are flattened or
straightened, the effects of the cold work
imparted by flattening or straightening are
not recognized in specified property limits.
T7, solution heat treated and stabilized.
Applies to products that have been precip-
itation heat treated to the extent that they
are overaged. Stabilization heat treatment
carries the mechanical properties beyond
the point of maximum strength to provide
some special characteristic, such as en-
hanced resistance to stress-corrosion crack-
ing or to exfoliation corrosion.
T8, solution heat treated, cold worked,
and artificially aged. Applies to products
that are cold worked specifically to improve
strength after solution heat treatment, and
for which mechanical properties or dimen-
sional stability, or both, have been substan-
tially improved by precipitation heat treat-
ment. The effects of cold work, including
any cold work imparted by flattening or
straightening, are recognized in specified
property limits.
T9, solution heat treated, artificially
aged, and cold worked. Applies to products
that are cold worked specifically to improve
strength after they have been precipitation
heat treated.
T10, cooled from an elevated-tempera-
ture shaping process, cold worked, and
artificially aged. Applies to products that
are cold worked specifically to improve
strength after cooling from a hot-working
process such as rolling or extrusion, and for
which mechanical properties or dimensional
stability, or both, have been substantially
improved by precipitation heat treatment.
The effects of cold work, including any cold
work imparted by flattening or straighten-
ing, are recognized in specified property
limits.
Other Subdivisions of T Temper Codes for
Stress-Relieved Products. When it is desir-
able to identify a variation of one of the ten
major T tempers described above, addition-
al digits, the first (x) of which cannot be
zero, may be added to the designation.
The following specific sets of additional
digits have been assigned to stress-relieved
wrought products.
Tx51, stress relieved by stretching. Ap-
plies to the following products when
stretched to the indicated amounts after
solution heat treatment or after cooling
from an elevated-temperature shaping pro-
cess:
Product form
Plate
Rod, bar, shapes, extruded
tube
Drawn tube
Permanent set, %
112-3
1-3
1/2-3
Tx51 applies directly to plate and to rolled
or cold finished rod and bar. These products
receive no further straightening after
stretching. Tx51 also applies to extruded
rod, bar, shapes, and tubing, and to drawn
tubing, when designated as follows:
•
Tx510. Products that receive no further
straightening after stretching
⚫ Tx511. Products that may receive minor
straightening after stretching to comply
with standard tolerances
⚫ Tx52. Stress relieved by compressing. Ap-
plies to products that are stress relieved
by compressing after solution heat treat-
ment, or after cooling from a hot-working
process to produce a permanent set of 1
to 5%
• Tx54. Stress relieved by combining
stretching and compressing. Applies to
die forgings that are stress relieved by
restriking cold in the finish die. (These
same digits and 51, 52, and 54-may be
added to the designation W to indicate
unstable solution heat-treated and stress-
relieved tempers)
Heat Treating of Aluminum Alloys / 879
Temper designations T42 and T62 have
been assigned to wrought products heat
treated from the O or the F temper to
demonstrate response from the heat treat-
ment described below. Temper designations
T42 and T62 also may be applied to wrought
products heat treated from any temper by
the user when such heat treatment results in
the mechanical properties applicable to
these tempers.
• T42. Solution heat treated from the O or
the F temper to demonstrate response to
heat treatment and naturally aged to a
substantially stable condition
• T62. Solution heat treated from the O or
the F temper to demonstrate response to
heat treatment and artificially aged
Subdivision of the O Temper. In temper
designations for annealed products, a digit
following the O indicates special character-
istics. For example, 01 denotes that a prod-
uct has been heat treated according to a
time/temperature schedule approximately
the same as that used for solution heat
treatment, and then air cooled to room
temperature, to accentuate ultrasonic re-
sponse and provide dimensional stability;
this designation applies to products that are
to be machined prior to solution heat treat-
ment by the user.
REFERENCES
1. S. Hirano et al., Quench Sensitivity in
Al-Li Based Alloys, Proceedings of Con-
ference on Aluminum-Lithium Alloys
(Vol 1), Materials and Component Engi-
neering Publications, 1989, p 335-344
2. T. Sheppard, Mater. Sci. Technol., Vol
4, July 1988, p 636
3. J.E. Hatch, in Aluminum Properties and
Physical Metallurgy, American Society
for Metals, 1984, p 165-166
4. C.E. Bates, Selecting Quenchants to
Maximize Tensile Properties and Mini-
mize Distortion in Aluminum Parts, J.
Heat Treat., Vol 5 (No. 1), 1987, p 27-40
5. T. Croucher, Critical Parameters for
Evaluating Polymer Quenching of Alu-
minum, Heat Treat., Vol 19 (No. 12),
Dec 1987, p 21-25
6. W.L. Fink and L.A. Willey, Quenching
of 75S Aluminum Alloy, Trans. AIME,
Vol 175, 1948, p 414-427
7. J.W. Evancho and J.T. Staley, Kinetics
of Precipitation in Aluminum Alloys dur-
ing Continuous Cooling, Metall. Trans.
A, Vol 5A, Jan 1974, p 43-47
8. J.T. Staley, Industrial Heating XLIV,
Oct 1977, p 6-9
9. G.F. Bobart, J. Heat Treat., Vol 6 (No.
1), 1988, p 47-52

nominal aging practice of 24 h at 120 °C (250 °F) affect the strength of 7075-T6

7050-W (4 days) extrusions were aged at 24 h at 120 °C (250 °F)

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