!!Comparison of superplastic behavior in two 5083 aluminum alloys

Comparison of superplastic behavior in two 5083 aluminum alloys

R.M. Cleveland a, A.K. Ghosh a,*, J.R. Bradley b

a Department of Materials Science and Engineering, The University of Michigan, 2300 Hayward Street, Ann Arbor, MI 48109-2136, USAb General Motors R&D and Planning, Warren, MI 48090-9055, USA

Received 8 May 2002; received in revised form 28 October 2002

Abstract

Superplastic elongation is generally controlled by both strain-rate sensitivity and particles causing cavitation. The superplastic

tensile behavior of 5083 Al (nominally Al�/4.6% Mg) alloys from two different sources has been examined. It is found that for nearly

the same grain size, one of the alloys (alloy A) has a slightly lower strain-rate sensitivity m . This alloy also contains a significantly

larger number of hard particles and inclusions. Because this alloy exhibits somewhat lower flow stress during constant strain-rate

tests, a reduced tendency for cavity initiation was expected. However, the larger density of particles tends to increase cavitation by

providing more nucleation sites. Then the lower m value leads to more rapid cavity growth and interlinkage. Thus, while the

presence of fine grain size and high-angle grain boundaries are important for superplastic flow, this research shows that in similarly

processed alloys small changes in m values and variations in alloy chemistry that generate cavity-causing particles can have a

medium to large effect on superplastic elongation.

# 2002 Elsevier Science B.V. All rights reserved.

Keywords: Cavity; Superplastic; Strain-rate sensitivity

1. Introduction

Currently, there is a strong interest in superplastic-

forming technology for the fabrication of automotive

sheet parts from 5000-series aluminum (i.e., Al�/Mg

alloys). It is well known that fine grain size and the

presence of high-angle grain boundaries are important

requirements for good superplastic behavior. The degree

of superplasticity in engineering alloys is influenced by

grain size and grain size distribution because of the role

of grain boundary sliding in the deformation mechan-

ism. Smaller grain size generally leads to lower flow

stresses, higher values of strain-rate sensitivity (m ), and

greater ductility [1�/3]. The distribution of grain sizes

present in production alloys can also have an important

impact on the strain-rate range in which m is high (and

peak ductility is achieved) [4]. Grain size is controlled in

these alloys by adding dispersoid-forming elements to

pin grain boundaries.

Superplastic flow is generally terminated by internal

cavitation at pre-existing particles and by the growth

and the interlinkage of such cavities [5�/9]. Cavities tend

to nucleate on second-phase particles or other high-

energy boundaries. It is well known that damage can

occur at the matrix�/particle interface during cold rolling

in alloys with hard particles. Recent evidence [10]

suggests that such damage at the particle interface may

not be entirely eliminated by the recrystallization

process and therefore could be a source of nucleation

sites for cavities during subsequent superplastic forming.

Diffusion and plastic flow during tensile straining lead

to cavity growth and, as cavities become larger or

numerous enough to interlink, failure can occur. Gen-

erally, increasing strain rate is found to increase cavita-

tion in these alloys, while higher temperature is found to

decrease cavitation. For example, cavitation was found

to be significantly reduced in a 5083 Al alloy by raising

the temperature from 525 to 550 8C [11].

In addition to affecting the grain size and cavitation,

variations in the alloy chemistry and microstructure of

the selected alloy may also influence the achievable

forming severity in parts. It is well known in industry

that the superplastic formability of 5083 Al sheet varies* Corresponding author.

E-mail address: [email protected] (A.K. Ghosh).

Materials Science and Engineering A351 (2003) 228�/236

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considerably with source and processing. While many

trends in the superplastic behavior of 5083 alloys have

been well documented, little attention in the literature

has been directed at the effect of small differencesbetween sheets of nominally the same composition on

the superplastic formability of these alloys.

In this study, two slightly different thermomechanical

processing methods were used to prepare samples from

one alloy chemistry and the behavior of this alloy was

compared with a commercially available 5083 Al sheet

material. These samples were examined to determine the

grain size and particulate size distributions in the alloys.Both these factors can influence superplastic ductility in

definite ways [4,7]. The influence of these parameters on

the superplastic mechanical behavior and fracture of the

alloys was examined using constant strain-rate tensile

tests. Such tests permit a comparison with the overall

superplastic ductility. Additional tests were used to

determine the strain-rate sensitivity of flow stress and

the progression of cavitation as a function of strain todevelop an understanding of the mechanisms leading to

failure.

2. Experimental work

5083 Al alloys were procured from two different

sources for investigation in this study. Alloy A was

produced by ARCO aluminum and alloy B was

produced by SKY aluminum. Their compositions arelisted in Table 1. Alloy A contains more Mg, Mn, and

Fe than alloy B. Alloy A was received in 4.5 mm thick

hot-rolled plates with an elongated coarse grain struc-

ture, whereas alloy B was received in cold-rolled and

annealed sheets (1.5 mm) with a fine grain structure.

More Mg in alloy A would not generally be regarded

as detrimental to the magnitude of superplastic ductility

even though more Al3Mg2 particles would be present inthe alloy at room temperature. A higher amount of Mg

can lead to higher strength in the alloy at room

temperature. At superplastic temperatures, Mg-rich

particles are in solid solution and thus higher Mg

content can produce enhanced diffusivity in the alloy

arising from the atomic size differences. Higher diffu-

sivity should in turn lead to lower flow stresses. In Al�/

Mg alloys that exhibit solute pinning of dislocations atlow temperatures, lower flow stress at superplastic

temperatures is often observed with increased levels of

Mg [12]. A higher level of Mg can also cause greater

strain hardening of the alloy during thermomechanical

processing and is likely to produce internal damage.While the intermetallic dispersoids formed by the

addition of Fe and Mn are important for grain size

stability, a high level of Fe and Mn is known to produce

larger hard particles that can act as nucleation sites for

cavitation.

To check the potential for grain refinements of alloy

A, two different thermomechanical treatments were

examined. The first method (TMT-1) involved coldrolling the plate (4.5 mm) down to a thin sheet (1.2

mm), a 73% reduction, with intermediate stress relief

anneals of 15 min at 170 8C to prevent cracking. The

sheet was then recrystallized at 500 8C for 30 min. The

second method (TMT-2) was similar to the first except

that the plate was given a solution heat treatment

followed by an overaging step to precipitate Al3Mg2

particles. The particles are expected to generate a greateramount of stored strain energy during the rolling step.

This in turn can lead to finer recrystallized grains. In

TMT-2, the plate was solution heat treated at 500 8Cfor 30 min, quenched in water, and then aged at 230 8Cfor 20 h prior to rolling and final recrystallization.

2.1. Microstructures

The microstructures of alloy A in the as-received and

TMT-2 conditions are shown in Fig. 1(a) and (b), andthe average grain diameters of alloy A in the as-received,

TMT-1, and TMT-2 conditions are listed in Table 2 (the

rolling direction of the sheet is horizontal and the

thickness direction is vertical). These samples were also

aged for 16 h at 120 8C, polished in colloidal silica, and

then etched in 10% phosphoric acid at 50 8C. The TMT-

2 condition was found to produce the smallest grain size

(5.6 mm) for alloy A, and was therefore chosen fortensile testing. The microstructure of alloy B obtained in

the sheet form is also shown in Fig. 1(c) for comparison.

Alloy B was found to have an average grain diameter of

4.5 mm, which is slightly finer than alloy A in TMT-2

condition. A histogram comparing the grain size dis-

tributions of alloy A in the TMT-2 condition and alloy

B is shown in Fig. 2. For each material, the areas of

more than 500 grains were individually measured fromdigital micrographs using NIH Image software and the

Table 1

Compositions of 5083 Al alloys by weight (remainder aluminum)

Mg (%) Cr (%) Mn (%) Si (%) Fe (%) Cu (%) Zn (%) Ti (%) Source

A 4.75 0.09 0.85 0.12 0.26 0.08 0.032 0.013 ARCO

B 4.36 0.086 0.68 0.13 0.06 0.011 0.001 0.001 SKY

Remainder is aluminum.

R.M. Cleveland et al. / Materials Science and Engineering A351 (2003) 228�/236 229

diameters calculated by assuming that the grains are

spherical. The diameter is given by D�ffiffiffiffiffiffiffiffiffiffiffi4A=p

p; where

A is the measured area in a two-dimensional section.

Relative frequency is the percent of the measured grains

falling in each bin. The distributions are generally

similar except that alloy A is shifted about 1 mm toward

the larger size compared with alloy B.

As-polished micrographs of alloys A and B are shown

in Fig. 3(a) and (b). A significant difference in the

particulate size and volume fraction in the two alloys is

evident. Clearly, alloy A has larger particles than alloy

B. Such digital micrographs were analyzed using NIH

Image to count and measure the particles. The resolu-

tion of the micrographs allowed particles about 0.5 mm

in diameter and larger to be counted. A minimum

sample area of 0.04 mm2 was analyzed for each material.

Measurements of particles from micrographs (Table 3)

indicate that the volume fraction of particles in alloy B is

much lower and that the average particle size is also

smaller. Only particles of about 0.5 mm in diameter and

larger could be detected by the optical method used. A

histogram of the particle size distribution is shown in

Fig. 4, where the relative frequency is determined by the

number of particles in each diameter group divided by

the total number of particles counted for each alloy. The

plot clearly shows that alloy A has larger particles than

Fig. 1. Micrographs of the grain structures of alloys A and B. Alloy A is shown in the as-received condition (a) and after TMT-2 (b). The sheet

structure of alloy B is shown in (c).

Table 2

Grain diameters of 5083 Al alloys (mm)

A: as-received A: TMT-1 A: TMT-2 B

9.4 6.4 5.6 4.5

Fig. 2. The grain size distribution in Al 5083 alloys A (in TMT-2

condition) and B.

Table 3

Measurements of particles in Al 5083 alloys

Alloy A Alloy B

Area fraction of particles (%) 4.2 0.5

Average particle diameter (mm) 2.9 1.3

Standard deviation (mm) 2.3 0.9

R.M. Cleveland et al. / Materials Science and Engineering A351 (2003) 228�/236230

alloy B. The compositions of the particles were exam-

ined using X-ray energy-dispersive spectroscopy on a

Philips XL30 SEM. Fig. 5 is an SEM micrograph of

alloy A with labels indicating the two types of particles

Fig. 3. Micrographs showing the particulates in aluminum 5083 alloy

A (a) and alloy B (b).

Fig. 4. The particle size distribution in Al 5083 alloys A and B. The

relative frequency is the number of particles in each diameter group

normalized by the total number of particles counted for each alloy.

Fig. 5. SEM micrograph of particles in alloy A. The particles labeled

‘‘a’’ are found to be rich in iron and manganese. The particles labeled

‘‘b’’ are found to be rich in silicon and oxygen.

Fig. 7. Stress�/strain curves for both Al 5083 alloys tested at constant

strain rate at 500 and 550 8C.

Fig. 6. Photographs of samples after various test conditions.


Fig. 8. An example of a stress�/strain curve used for the determination of m values.

Fig. 9. Flow stress as a function of strain rate for alloy A in the TMT-2 condition (a, b) and alloy B (c, d).Plots (a, c) are for 500 8C and (b, d) are for

550 8C. The flow stresses increase for higher strain levels.


found. The particles labeled ‘‘a’’ were found to be rich in

iron and manganese, while those labeled ‘‘b’’ were found

to be rich in silicon and oxygen.

2.2. Tensile tests

Superplastic tensile samples were machined from

alloy A sheet in the TMT-2 condition and from the

alloy B sheet that was already in fine grain condition.

The tensile specimens had a gage width of 6.2 mm and

gage length of 12.7 mm. Samples of both Al 5083 alloys

were tested at temperatures of 500 and 550 8C. The tests

were conducted on a computer-controlled Instron 4505

with a three-zone split furnace that wrapped around thesample. Load was measured using a 1000 kN load cell

and strain was calculated from cross-head displacement

and checked by final measurements of the samples. A

constant strain rate of 0.001 s�1 was maintained by the

method developed by Friedman and Ghosh [13]. Photo-

graphs of the initial sample geometry and samples after

testing are shown in Fig. 6. At this strain rate, the

average elongations for alloy A were 190% at 500 8Cand 340% at 550 8C. For alloy B, the average elonga-

tions were 270% at 550 8C and 310% at 550 8C. These

elongations represent a reasonable degree of super-

plasticity.

Stress�/strain plots for these tensile tests are shown in

Fig. 7. Alloy A exhibited lower flow stresses than alloy B

at both temperatures*/a result unexpected on the basis

of its slightly coarser grain size, but reasonable from the

standpoint of higher diffusivity resulting from the higher

Mg content. Alloy B failed at a higher elongation at

500 8C (�/270%), but both alloys failed at nearly the

same final elongation at 550 8C (310�/340%). Micro-

graphs of samples deformed at 500 8C showed that the

average grain size increased to between 7 and 8 mm in

alloy A (for strains of �/0.5) and to between 9 and 10

mm in alloy B (for strains of �/0.7), indicating similar

grain growth characteristics.

To obtain a clear understanding of the above features,

the strain-rate sensitivity of the alloy must be examined.

The variation of flow stress as a function of strain rate

and the corresponding strain-rate sensitivity of flow

stress, m , were characterized using step strain-rate tests.

In these tests, a succession of strain-rate decrements and

increments were performed to determine steady-state

Fig. 10. The effect of strain rate on strain-rate sensitivity index, m , for alloy A in TMT-2 condition (a, b) and alloy B (c, d). Plots (a, c) are for 500 8Cand (b, d) are for 550 8C.


flow stress under different strain rates (and also over

different strain ranges). The scheme of these tests was

optimized previously [2]. The stress�/strain plot from

one such test is shown in Fig. 8. Similar tests were

performed at 500 and 550 8C. The resulting values of

flow stress for each imposed strain rate are plotted in

Fig. 9(a)�/(d) on a log�/log scale for the two alloys. The

general shapes of these curves are sigmoidal as those

found in the previous studies [1�/3]. The flow stress is

not simply a function of strain rate but is seen to

increase somewhat at higher strain levels. This strain-

induced hardening in superplastic materials has been

found to be a result of concurrent grain growth in the

alloy. Although the flow stress values for alloys A and B

do not vary a great deal from each other, upon closer

examination, alloy A is found to have somewhat lower

flow stress at higher strain rates, but higher flow stresses

at lower strain rates compared with alloy B.

The above observation of the trend of flow stress

variation with strain rate is consistent with the results

shown in Fig. 10, which shows strain-rate sensitivity

(m ), given by the slope of the log stress vs. log strain rate

curves from Fig. 9. The peak m values of the two alloys

in Fig. 10 are in the range 0.4�/0.65 for different

temperatures and generally increase with increasing

test temperature. However, for alloy A, the peak values

of m were about 0.47 at 500 8C and 0.54 at 550 8C and

the peak values of m for alloy B were consistently

higher, ranging from 0.52 at 500 8C to 0.64 at 550 8C.

This difference in m values was repeatable and is

consistent with the higher tensile elongation of alloy B

at 500 8C. It should be noted that the grain size

distribution plot in Fig. 2 indicates that in the finer

grain size range (2�/4 mm), alloy B has considerably

more grains present. As the model presented in [4]

indicates, this is exactly the reason responsible for the

slightly higher m value in alloy B. We believe that this

grain size distribution is a result of local variations in

stored cold work in alloy A with more Mg and more

particles that provide local areas of grain growth. The

uniformity of composition in alloy B may result in a

more homogeneously distributed strain energy and

Fig. 11. Micrographs showing the progression of cavitation in the two alloys at 500 8C: alloy A (a, b) and alloy B (c, d). True strains at the position

on the sample, where each micrograph was taken, are indicated.


hence a finer grain sizes. Since m values are generally

higher for smaller grain sizes [1,2], the decrease in m

value with increasing strain observed in the plots of Fig.

10 is consistent with concurrent grain growth. At boththese temperatures, the m vs. strain rate curves are more

sharply peaked for alloy B than for alloy A. The broader

peaks in m for alloy A suggest that its formability is less

sensitive to strain rate, which may make it a more

practical alloy for production if the cavitation can be

minimized.

2.3. Cavitation study

In order to further examine the differences between

the behavior of the two Al 5083 alloys, the tendency for

the formation of cavities was studied as a function of

strain in the samples that were pulled to failure at

500 8C and an overall strain rate of 0.001 s�1. Grid

lines on the surface of the samples as well as width and

thickness measurements were used to determine the

strain levels at several points along the sample length sothat cavitation could be measured at known strain

levels. Since alloy B failed at a larger strain, the selected

strain levels were larger than for alloy A. Micrographs

of the samples in the as-polished condition from the

selected locations were examined. Several examples of

the micrographs are shown in Fig. 11(a)�/(d). In Fig.

11(a) and (b), the cavities in alloy A are shown at true

strains of 0.51 and 1.02, respectively. Alloy B is shown attrue strains of 0.85 in (c) and 1.53 in (d). Alloy A is seen

in the micrographs to have a larger number of cavities at

the depicted strain levels, but many of the cavities in the

alloy A micrographs are relatively small. It is not

surprising that alloy A form more cavities since it has

many more second-phase particles than alloy B on

which cavities can nucleate. It is believed that the high

frequency of large particles in alloy A may havecontributed to the creation of damage during cold

rolling, which ultimately influenced the levels of cavita-

tion during superplastic testing.

Micrographs such as those shown in Fig. 11 were

quantitatively analyzed using NIH Image software to

determine the volume fraction and frequency of cavities.

Fig. 12(a) depicts the measured number of cavities per

unit area as a function of strain for the two alloys. Thevariations in the measurements lie within the bounding

lines that describe the data trend. It is generally expected

that the number of cavities would approach zero at zero

strain. (Micrographs of the grip region of the sample

confirm that this is true.) Fig. 12(b) shows the average

cavity diameter as a function of strain. The average

cavity diameter in alloy A is somewhat smaller, i.e., it

has a greater number of smaller cavities at low strainuntil rapid growth of the cavities occurs near failure.

This smaller average cavity size has previously been seen

in alloys containing more particles and fits with existing

models of cavitation [14]. The volume fraction of cavities

is nearly the same for the two alloys at low strains, as

shown in Fig. 12(c). However, in alloy A, which has a

larger number of more closely spaced cavities, the rapidgrowth in the cavities leading to cavity interlinkage and

failure begins at a lower strain.

3. Conclusions

The effect of small differences in microstructure and

composition between sheets of nominally the same

Fig. 12. Measurements of cavitation for alloys A and B at 500 8C.


chemistry has been investigated to understand the

differences in their superplastic response. The two sheets

had similar grain sizes but different grain size distribu-

tions, different strain-rate sensitivity values, and sig-

nificantly different particle content resulting from small

variations in Mg, Mn, and Fe content. This research

shows that in similarly processed alloys with changes in

m values, cavity-causing particles can have a measurable

effect on superplastic elongation.

Alloy A with a slightly higher Mg level compared with

alloy B exhibited lower flow stress in constant strain-

rate tests, possibly because of enhancement in the

chemical diffusivity of the alloy.

In spite of the lower flow stress for alloy A, the

tendency for cavity initiation is not reduced because of

its larger density of particles and lower strain-rate

sensitivity value.

The peak values of strain-rate sensitivity were lower

for alloy A than for alloy B, which appears to result

from a lower fraction of fine grains. The m values

decreased with increasing strain, as expected from

concurrent grain growth. Lower m values led to more

rapid cavity growth and interlinkage in alloy A. The

combination of a large number of particles and lower m

values in alloy A resulted in reduced ductility.

Acknowledgements

The authors thank General Motors Research and

Development Center for the support of this research.

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!!Comparison of superplastic behavior in two 5083 aluminum alloys

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