ASM Metals HandBook Vol. 14 - Forming and Forging
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• Microduplex stainless steels (ferrite-austenite two-phase system)
Of these, only the hypereutectoid steels have been extensively investigated for applications involving superplastic
forming.
The development and existence of superplasticity in all of the above categories of iron-base alloys have been reviewed in
detail (Ref 53, 54). This article will describe the key characteristics of superplastic behavior in each of these iron-base
systems, with emphasis on hypereutectoid plain carbon steels, which currently have the most industrial potential.
Hypoeutectoid and Eutectoid Plain Carbon Steels
Early attempts at making superplastic plain carbon steels were not particularly successful (Ref 53). In these early efforts,
extensive heat treatments were used to produce microstructures of fine ferrite grains pinned by spheroidized cementite
particles in compositions containing between 0.2 and 1.0% C. Even though the overall grain structure was correct,
elongations of only about 130% were observed, primarily because the grain boundaries were low-angle dislocation
boundaries. The flow stress of materials with low-angle grain boundaries is usually not sensitive to changes in strain rate;
therefore, superplasticity was not observed. However, later work showed that thermally cycling such materials across the
eutectoid transformation temperature after thermomechanical processing changes the grain boundaries from low angle to
high angle; tensile elongations of the order of 1000% have been achieved in thermally cycled, hypoeutectoid plain carbon
steels (Ref 51).
Hypereutectoid Plain Carbon Steels and White Cast Irons
Hypereutectoid plain carbon steels, also known as ultrahigh-carbon (UHC) steels, will exhibit a microstructure
consisting of a continuous network of proeutectoid carbide (at prior-austenite grain boundaries) surrounding pearlite
colonies when slowly cooled from a temperature in the single-phase austenite region. This carbide network in turn
imparts poor mechanical properties to these steels. However, work conducted in the 1970s showed that, with proper
thermomechanical processing, microstructures can be achieved in UHC steels that consist of ultrafine, equiaxed grains of
ferrite and a uniform distribution of fine, spherical, discontinuous proeutectoid carbide particles (Ref 55). Such
microstructures in UHC steels have led to superplastic behavior; tensile elongations up to 1500% have been reported (Ref
56).
The requisite thermomechanical processing of UHC steels to obtain the desired fine spheroidized microstructure usually
involves two steps. In the first step, the UHC ingot or billet is solution heat treated at typically 1150 °C (2100 °F),
followed by hot and warm working during cooling down to 750 °C (1380 °F). This step refines the austenite grains,
uniformly distributes the proeutectoid carbides at prior-austenite grain and subgrain boundaries, and forms pearlite in the
matrix of the prior-austenite grains. The second step incorporates a divorced eutectoid transformation in which the UHC
steel is heated to just above the eutectoid transformation temperature and air cooled (with or without accompanying
deformation). This step transforms the fine pearlite from step 1 into a fully spheroidized fine structure. The divorced
eutectoid transformation is discussed in more detail in Ref 57.
Although not yet fully commercialized, the industrial potential of UHC steels has been thoroughly investigated. One
promising application is in the press forging of gears. Figure 44 illustrates a bevel gear that was warm forged from a fine-
grain 1.25% C UHC steel in a single operation. An additional advantage of using superplastic UHC steels is that, because
of their high carbon content, the carburization step in normal gear production is eliminated. A second example of the
superplastic press forming of fine-grain UHC steels is the aft closure for a guided missile shown in Fig. 45. In this case, a
1.6% C steel casting was liquid atomized, and the resulting fine-grain powders were warm pressed into a billet at 800 °C
(1470 °F). A third application for superplastically formed fine-grain UHC steels is in the manufacture of die components
by superplastic hobbing operations (Ref 58).
Fig. 44 Warm precision forging of a 1.25% C UHC steel billet into a bevel gear.
Forging temperature was 650
°C (1200 °F).
Fig. 45 Aft closure, for a guided missi
le, that was superplastically formed at 815 °C (1500 °F) from a 1.6% C
UHC steel. The processing procedure consisted of warm pressing (800 °C, or 1470 °F) liquid-
atomized powders
into a billet and forging the resulting billet into plate. The plate was then
superplastically formed to the final
shape (a and b). (c) Cross section of the part
The original work on UHC steels demonstrated that, given the proper microstructure, superplasticity could be developed
over the composition range of about 0.8 to 2.1% C and the temperature range of about 650 to 800 °C (1200 to 1470 °F).
Therefore, superplastic UHC steels include both austenite-carbide as well as ferrite-carbide two-phase structures.
Ultrahigh-carbon steels are generally not susceptible to cavitation, because the ferrite, austenite, and carbide all have
about the same strength in the intermediate temperature range in which superplasticity occurs.
Furthermore, in the original work, superplasticity in UHC steels was limited to intermediate strain rates (that is, forming
rates): 10
-5
to 10
-3
s
-1
. The upper limits on strain rate and temperature are primarily related to the destruction of the fine
uniform carbide structure. Commercialization of superplastic UHC steels has in fact been hampered by these somewhat
limited strain rate and temperature ranges. However, more recent work has shown that these ranges could be substantially
extended by careful alloying additions of silicon or aluminum (Ref 59). These elements enhance superplasticity in UHC
steels because they:
• Increase the eutectoid transformation temperature (increase the stability of ferrite)
• Inhibit carbide coarsening at high temperatures (increase the activity of carbon in ferrite)
• Increase the volume fraction of proeutectoid carbides
• Do not form active sites for cavitation to occur
Ultrahigh-carbon steels containing 3% Si or 1.6% Al have exhibited superplasticity at strain rates to 10
-2
s
-1
at a
temperature of 800 °C (1470 °F). Superplasticity has been projected to occur in a UHC 12% Al alloy at strain rates to 3 ×
10
-1
s
-1
at 950 °C (1740 °F); this would make superplastic forming of UHC steels both economical and feasible for many
operations.
An additional benefit of adding aluminum or silicon to UHC steels is a low, fairly constant (14 ± 1.5 MPa, or 2.00 ± 0.22
ksi) flow stress over the temperature range of 750 to 925 °C (1380 to 1700 °F). This suggests that thin-sheet UHC steels
containing the proper amounts of aluminum or silicon can be readily blow formed over a wide temperature range.
Hypereutectoid White Cast Irons. Application of rapid solidification processing to UHC steels has also helped to
extend the compositional range for superplasticity well into the white cast irons. Such techniques can create very high
carbon content powders that, upon annealing at intermediate temperatures (595 to 700 °C, or 1100 to 1290 °F), exhibit the
requisite fine microstructure (Ref 60). These powders are then readily consolidated into fully dense compacts at
temperatures below the subcritical annealing temperature (A
1
) (Ref 61). White cast irons processed in this manner have
been shown to exhibit superplasticity at intermediate temperatures; a maximum tensile elongation of 1410% has been
observed in a Fe-3.0C-1.5Cr (Ref 62). The shaded region on the iron-carbon phase diagram shown in Fig. 46 illustrates
the range of carbon content and temperature over which superplasticity has been documented in fine-grain UHC steels
and cast irons.
Fig. 46 Iron-carbon phase diagra
m. The shaded area illustrates the temperature and composition over which
UHC steels and white cast irons have been made superplastic.
Hypereutectoid Steel Properties. Because of their carbon content, UHC steels, after superplastic forming, can be
heat treated to very high hardness levels (65 to 68 HRC). This can be important in a wide variety of applications,
including gears, tool bits, abrasion-resistant surfaces, and military vehicle armor. However, because of their very fine
grain size, plain carbon UHC steels also exhibit poor hardenability. This hardenability problem can be alleviated by dilute
alloying additions of hardenability-enhancing elements, such as chromium, manganese, and molybdenum. The dramatic
effect of alloying additions on the hardenability of UHC steels is shown in Table 2.
Table 2 Hardenability of UHC steels as a function of alloying additions
Austenitizing
temperature
(A
1
+ 50)
Critical brine
temperature to
achieve
62 HRC, T
c
Critical cylinder
diameter, D
c
Material composition
°C °F °C °F mm in.
M.F.
(a)
1.25C-0.5Mn 773 1425 66 151 6.9 0.27
1.0
1.25C-0.5Mn-0.2P 795 1465 72 162 10.9 0.43
1.6
1.25C-1Cr-0.5Mn-0.25Mo
790 1455 73 163 11.4 0.45
1.7
1.25C-3Si-1.4Cr-0.5Mn 825 1520 78 172 15.5 0.61
2.2
1.25C-1.6Al-1.5Cr-0.5Mn
860 1580 93 199 22.4 0.88
3.2
1.25C-2Mn-1Cr 795 1465 >100 >212 >23.1 >0.91 >3.3
(a)
M.F., multiplying factor = D /D
The room-temperature tensile properties of fine-grain UHC steels have been extensively studied (Ref 63). As expected,
the properties are very sensitive to heat treatment. Figure 47 compares the tensile properties of UHC steels in two
different heat treatment conditions with the properties of other plain carbon and low-alloy structural steels. In addition, if
properly heat treated to achieve an extremely fine (optically unresolvable) martensite, fully hardened UHC steels will
exhibit a surprisingly high strain to failure in compression. As shown in Fig. 48, a 1.3% C steel that was water quenched
from 770 °C (1420 °F) (Steel A) will exhibit a compression fracture strength of 4.5 GPa (650 ksi) and a strain to failure of
10%. However, austenitizing the same steel at 1100 °C (2010 °F) prior to water quenching from 770 °C (1420 °F) (Steel
B) will coarsen the resulting martensite and reduce the compression strain to failure to less than 2%.
Fig. 47 Tensile strength versus elongation to failure of heat-treated fine-grain UHC steels compared to low-
carbon steel, high-strength low-alloy (HSLA) steels, and dual-phase steels
Fig. 48 Influence of prior heat treatment on the microstructure (a and b) and compression stress-
strain
behavior (c) of a 1.3% C UHC steel quenched from 770 °C (1420 °F)
Because they contain a high volume fraction of carbide, UHC steels exhibit only moderate impact resistance (<25 J, or 18
ft · lbf, Charpy V-notch at 25 °C, or 80 °F). The impact resistance of fine-grain UHC steels can be enhanced by
laminating them to tougher (but weaker) materials. Fine-grain UHC steels can be readily solid state bonded to themselves
or other ferrous materials, partly because of their superplastic nature (Ref 64). The lamination procedure can consist of
conventional roll bonding, press bonding, or explosive techniques as long as a good metallurgical bond results. The effect
of such lamination on the impact properties of UHC steels is illustrated in Fig. 49. The impact properties of a UHC
steel/1020 steel laminated composite are clearly superior to those of either of the monolithic UHC steel or the monolithic
1020 low-carbon steel. Most UHC steel-base laminates will also exhibit superplasticity over limited temperature and
strain rate ranges and therefore can be superplastically formed (Ref 65, 66). The room-temperature tensile strength of a
UHC steel laminate will of course be reduced relative to that of a monolithic UHC steel by an amount governed by the
rule of mixtures; however, this loss in strength can be partially compensated for by heat treating the laminate to
selectively harden the UHC steel component to very high strength levels.
Fig. 49 UHC steel/1020 steel laminated composite to improve impact resistance of fine-
grain UHC steels. (a)
Orientation of mechanical test samples taken fr
om a laminated composite of UHC steel and 1020 steel. (b)
Optical photomicrograph of interface in laminated composite of UHC steel and 1020 showing an excellent
metallurgical bond. (c) Bend test sample of the laminated composite. (d) Impact properties of U
HC steel, 1020
steel, and UHC steel/1020 steel laminated composite including photographs of the tested samples
Low-to-Medium Alloy Steels
Superplastic studies have been conducted on several alloy steels (Ref 53). Compositions and thermomechanical
processing procedures are generally chosen to generate fine-grain (1 to 2 m, or 40 to 80 in.) microstructures consisting
of roughly equal amounts of ferrite and austenite. These phases will coexist only in a narrow temperature range, but
within that range each phase will inhibit the other from growing, and superplasticity will occur. Elongations of 300 to
500% have been achieved in steels containing 1 to 2% Mn and 0.1 to 0.4% C, and elongations of up to 600% have been
achieved in Fe-4Ni-3Mo-1-2Ti alloys when tested at temperatures in the dual-phase region. However, the potential of
superplastic forming these alloys has not been exploited, because of the narrow temperature range over which superplastic
flow occurs and because of the rapid growth that can occur even in the two-phase region at high temperatures.
Microduplex Stainless Steels
Microduplex stainless steels are so called because their stable microstructure consists of both ferrite and austenite in fine
grain size form--about 2 to 3 m (80 to 120 in.). They typically contain 18 to 26% Cr and 5 to 8% Ni and can also
contain molybdenum, titanium, copper, silicon, manganese, carbon, and diatomic nitrogen. Elongations to failure in these
alloys are commonly in excess of 500% at strain rates between 10
-3
and 10
-2
s
-1
and at temperatures between 900 and 1000
°C (1650 to 1830 °F) (Ref 66). Although microduplex stainless steels have found a number of commercial applications in
the chemical industry because of their high strength and corrosion resistance, their superplastic forming potential has not
been exploited. This is because microduplex stainless steels are susceptible to cavitation during superplastic flow due to
the strength differential between the ferrite and the austenite at these high temperatures (Ref 67).
Outlook for Superplastic Ferrous Alloys
The number of iron-base alloys exhibiting superplasticity is relatively limited. This is primarily because of the difficulty
in satisfying all of the microstructural requirements generally required for superplasticity in ferrous materials. Of the
alloys identified to date for possible superplastic forming, the hypereutectoid (or UHC) steels show the most promise.
These steels can be stretched without cavitation over 1000% at reasonable forming rates and intermediate temperatures.
Ultrahigh-carbon steels are heat treatable to high hardnesses, exhibit good room-temperature mechanical properties, and
can be readily laminated to other ferrous materials to enhance toughness and impact resistance. Future efforts in the
superplastic forming of iron-base alloys will undoubtedly focus on this class of steels.
References cited in this section
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W. Rosenhain and D. Ewen, J. Inst. Met., Vol 8, 1912, p 149
49.
E.E. Underwood, J. Met., Vol 914, 1962, p 919
50.
A.K. Mukherjee, in Annual Reviews of Materials Science,
Vol 9, R.A. Huggins, Ed., Annual Reviews, Inc.,
1979
51.
O.D. Sherby and J. Wadsworth, in Deformation, Processing, and Structure,
G. Krauss, Ed., American
Society for Metals, 1984, p 355
52.
O.D. Sherby and R.D. Caligiuri, in Superplasticity:AGARD(NATO) Lecture Series No. 154,
Advisory
Group for Aerospace Research and Development, 1987, p 3-1
53.
N. Ridley, in Superplastic Forming of Structural Alloys,
N.E. Paton and C.H. Hamilton, Ed., The
Metallurgical Society, 1982, p 191
54.
B. Walser and U. Ritter, in International Conference on Superplasticity,
B. Baudelet and M. Suery, Ed.,
Editions du Centre Nationale de la Recherche Scientifique, 1985, p 15.1
55.
O.D. Sherby, B. Walser, C.M. Young, and E.M. Cady, Scr. Metall., Vol 9, 1975, p 569
56.
T. Oyama, J. Wadsworth, M. Korchynsky, and O.D. Sherby, in Proceedings of the Fifth Internat
ional
Conference on the Strength of Metals and Alloys,
International Series on the Strength and Fracture of
Materials and Structures, Pergamon Press, 1980, p 381
57.
T. Oyama, O.D. Sherby, J. Wadsworth, and B. Walser, Scr. Metall., Vol 18, 1984, p 799
58.
R. Pearce and E.W.J. Miller, in Superplastic Forming of Structural Alloys,
N.E. Paton and C.H. Hamilton,
Ed., The Metallurgical Society, 1982, p 191
59.
O.D. Sherby, T. Oyama, D.W. Kum, B. Walser, and J. Wadsworth, J. Met., Vol 37 (No. 6), 1985, p 50
60.
L.E. Eiselstein, O.A. Ruano, and O.D. Sherby, J. Mater. Sci., Vol 18, 1983, p 483
61.
R.D. Caligiuri, R.T. Whalen, and O.D. Sherby, Int. J. Powder Metall. Powder Technol., Vol 3, 1976, p 154
62.
O.A. Ruano, L.E. Eiselstein, and O.D. Sherby, Metall. Trans. A, Vol 13A, 1982, p 1785
63.
H. Sunada, J. Wadsworth, J. Lin, and O.D. Sherby, Mater. Sci. Eng., Vol 38, 1979, p 35
64.
O.D. Sherby, J. Wadsworth, R.D. Caligiuri, L.E. Eiselstein, B.C. Snyder, and R.T. Whalen, Scr. Metall.,
Vol 13, 1979, p 941
65.
B.C. Snyder, J. Wadsworth, and O.D. Sherby, Acta Metall., Vol 32, 1984, p 919
66.
G.S. Daehn, D.W. Kum, and O.D. Sherby, Metall. Trans. A, Vol 17A, 1986, p 2295
67.
G.W. Hayden, R.C. Gibson, H.P. Merrick, and J.H. Brophy, Trans. ASM, Vol 60, 1967, p 3
Superplastic Sheet Forming
C.H. Hamilton, Washington State University; A.K. Ghosh, Rockwell International
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Editions du Centre Nationale de la Recherche Scientifique, 1985, p 15.1
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Conference on the Strength of Metals and Alloys,
International Series on the Strength and Fracture of
Materials and Structures, Pergamon Press, 1980, p 381
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Formability Testing of Sheet Metals
Brian Taylor, General Motors Corporation
Introduction
SHEET METAL FORMING is the process of converting a flat sheet of metal into a part of desired shape without fracture
or excessive localized thinning. The process may be simple, such as a bending operation, or a sequence of very complex
operations such as those performed in high-volume stamping plants. In the manufacture of most large stampings, a sheet
metal blank is held on its edges by a blankholder ring and is deformed by means of a punch and die. The movement of the
blank into the die cavity is controlled by pressure between the upper and lower parts of the blankholder ring.
This control is usually increased by means of one or more sets of drawbeads. These consist of an almost semicylindrical
ridge on the upper part of the blankholder and a corresponding groove in the lower part (the positions are sometimes
reversed). The drawbeads force the periphery of the blank to bend and unbend as it is pulled into the die; this increases the
restraining force considerably. Presses with capacities to 17.8 MN (2000 tonf) are commonly used for the manufacture of
large stampings, and presses to 26.7 MN (3000 tonf) are used for heavy-gage parts.
Sheet metal forming operations are so diverse in type, extent, and rate that no single test provides an accurate indication
of the formability of a material in all situations. However, knowledge of material properties and careful analysis of the
various types of forming involved in making a particular part are indispensable in determining the probability of
successful part production and in developing the most efficient process.
Formability Testing of Sheet Metals
Brian Taylor, General Motors Corporation
Types of Forming
Many forming operations are complex, but all consist of combinations or sequences of the basic forming operations--
bending, stretching, drawing, and coining (see the Section "Forming Processes for Sheet, Strip, and Plate" in this
Volume).
Bending is the most common type of deformation, and it occurs in almost all forming operations. Bending around small
radii can lead to splitting in the early stages of a forming process because it localizes strain and prevents its distribution
throughout the part. Ideally, strain should be distributed as uniformly as possible to maximize the amount of deformation
that can be obtained. Even a slight increase in the radius in a given location can sometimes significantly improve strain
distribution. Frequently, designs specify smaller radii than necessary, which results in manufacturing problems and
increased costs.
Lubrication is not recommended when bending over a sharp radius, because die friction reduces strain localization by
restricting metal movement away from the radius. The orientation of the sheet in relation to the rolling direction can also
be important in a bending operation. During rolling, inclusions and other defects become elongated in the rolling
direction, producing lines of weakness. When the axis of bending is in this direction, there is a tendency toward splitting
along the lines of weakness. This lowers resistance to fracture compared to when the axis is inclined to the rolling
direction. Detailed information on sheet metal lubrication is available in the article "Selection and Use of Lubricants in
Forming of Sheet Metal" in this Volume.
The outer and inner panels of a part are frequently assembled by bending (hemming) the edges of the outer panel around
the inner panel. This requires material that can be easily bent over very small radii. In the absence of other types of
deformation, bending produces tensile stresses on the outside surface. These decrease to zero at an interior level known as
the neutral axis. These stresses then become compressive on the inside of the bend. They can cause springback (shape