1. No category

ARTICLES
The microstructure length scale of strain rate sensitivity in ultrafinegrained aluminum
Adam D. Kammers
Department of Mechanical Engineering, The University of Michigan, Ann Arbor, Michigan 48109, USA
Jittraporn Wongsa-Ngam
Department of Mechanical Engineering, Faculty of Engineering, King Mongkut’s Institute of Technology
Ladkrabang, Bangkok 10520, Thailand
Terence G. Langdon
Departments of Aerospace & Mechanical Engineering and Materials Science, University of Southern California,
Los Angeles, California 90089-1453, USA; and Materials Research Group, Faculty of Engineering and the
Environment, University of Southampton, Southampton SO17 1BJ, UK
Samantha Dalya)
Department of Mechanical Engineering, The University of Michigan, Ann Arbor, Michigan 48109, USA; and
Department of Materials Science & Engineering, The University of Michigan, Ann Arbor, Michigan 48109, USA
(Received 2 September 2014; accepted 24 February 2015)
The mechanical properties of ultrafine-grained aluminum produced by equal-channel angular
pressing (ECAP) are strongly influenced by strain rate. In this work, an experimental investigation
of local strain rate sensitivity as it relates to microstructure was performed using a combination of
scanning electron microscopy and digital image correlation. Uniaxial tension tests were carried
out at 200 °C and strain rates alternating between 2.5 ! 10"5 s"1 and 3.0 ! 10"3 s"1. The results
demonstrate that the heterogeneous microstructure generated by ECAP has a strong effect on the
microstructure scale strain rate sensitivity. Deformation centered at grain boundaries separating
regions of banded microstructure exhibits the greatest strain rate sensitivity. Strain rate sensitivity
is limited in deformation occurring in regions of microstructure composed of ultrafine grains
separated by low-angle grain boundaries. The tensile specimens all failed by shear bands at
200 °C and at room temperature they failed by necking with little plastic deformation apparent
outside of the neck.
I. INTRODUCTION
Ultrafine-grained (UFG) metals possess an exceptional
combination of mechanical properties including considerable ductility and higher strength as compared to their
coarse grained counterparts.1–4 While the strength increase
is a result of the high fraction of grain boundaries
impeding dislocation motion as described by the Hall–
Petch relationship, the high ductility is generally associated with an enhanced strain rate sensitivity.5–10 This
research was motivated by the need to provide a better
understanding of the relationship between this enhanced
strain rate sensitivity and the material microstructure.
UFG metals are commonly produced by the severe
plastic deformation technique of equal-channel angular
pressing (ECAP).1,11–15 ECAP-processed Al has a shear/
torsion texture16–18 and consists primarily of ultrafine
Contributing Editor: Jürgen Eckert
a)
Address all correspondence to this author.
e-mail: [email protected]
DOI: 10.1557/jmr.2015.58
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
grains separated by high-angle grain boundaries
(HAGBs).19 But it may also contain “supergrains”
(defined by Davidson20 as groups of grains in which
the slip planes in adjacent grains are oriented within 15°
of each other) and “deformation bands”18,21 (referred to
as “microstructure bands” here to avoid confusion with
measured surface deformations). These different microstructures each play a role in affecting the mechanical
behavior of the material.22 This work builds on the room
temperature experiments carried out by the authors in
Ref. 22 by relating strain rate sensitivity apparent at
elevated temperatures to the underlying microstructure.
This research appears to be the first to relate full-field
microscale elevated temperature strain rate sensitivity to
the underlying microstructure in ECAP-processed UFG
aluminum. The elevated temperature strain rate sensitivity of ECAP-processed aluminum and aluminum alloys
was investigated in earlier experimental work5–7 and the
results showed that the enhanced strain rate sensitivity
arises from deformation mechanisms active at the
grain boundaries, including varying contributions from
grain boundary (GB) sliding and thermally activated
! Materials Research Society 2015
1
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
diffusion processes. Here, the effects of the ECAPprocessed microstructure on strain rate sensitivity are
characterized through a combination of scanning electron
microscopy (SEM), digital image correlation (DIC), and
electron backscatter diffraction (EBSD).
II. EXPERIMENTAL PROCEDURE
ECAP-processed high purity (99.99%) aluminum was
investigated in this work. The coarse-grained starting
material was purchased in 10 mm diameter ! 65 mm
long extruded rods from ESPI Metals (Ashland, OR). The
rods were first annealed in air at 773 K for 1 h19 to
achieve an average grain size of approximately 500 lm
(measured as the grain diameter by EBSD). An inverse
pole figure map of the material following annealing is
shown in Fig. 1. ECAP processing was carried out using
a 90° channel angle die possessing an outer arc of
curvature of 20°. The rods were lubricated with a molybdenum disulfide (MoS2) lubricant and then processed at
room temperature at a pressing rate of ;7 mm/s through
four ECAP passes following route BC in which the rods
are rotated by 90° in the same sense between each ECAP
pass.1,23–25 Flat dogbone-shaped tensile test specimens
with a gage cross-section of 2 mm ! 1 mm and length of
8 mm were fabricated by electro-discharge machining
(EDM). The tensile directions were aligned with the
ECAP pressing axis and the specimens were cut from the
central parts of the billets to avoid the presence of any
nonuniform deformation at either end of each billet.
Following standard ECAP processing naming conventions,1,11–15 the tension specimens were cut from the
billet so that their flat faces were parallel to the flow plane
(Y plane) of the billet in the final pass.
Surface preparation began with polishing, followed by
the deposition of alignment markers, EBSD, and the
application of a DIC tracking pattern. Polishing was
performed with diamond suspensions and a final abrasive
of Buehler MasterMet 2 non-crystallizing colloidal silica
suspension. 500 nm diameter and 500 nm tall platinum
alignment markers were then applied with a focused ion
beam (FIB) to enable alignment of each EBSD field
of view (FOV) and DIC FOV. EBSD was performed,
followed by plasma cleaning and surface patterning by
self-assembly of 40 nm diameter gold nanoparticles as
described in Ref. 26.
Tests were performed using an in situ tension–
compression stage (Kammrath and Weiss) equipped
with a 1 kN load cell and mounted in a FEI Quanta
200 3D FIB/SEM chamber. Experiments were carried
out at 200 °C and at strain rates that jumped between
2.5 ! 10"5 s"1 and 3.0 ! 10"3 s"1. Note that these were
not instantaneous jump tests, as the displacement was
held constant for 7 min between load steps to capture
images for DIC. Initial loading occurred at 2.5 ! 10"5
s"1 to a strain of 2.5% to ensure that the microstructure
within the FOVs had plastically deformed. Following
this, the specimens were loaded at alternating strain
rates, starting with the high strain rate. The strain rate
was alternated after strain steps of 2.5% until a final
macroscopic strain of 17.5% was achieved. Four of
these load steps are presented here out to a macroscopic
strain of 12.5%.
DIC was performed on the test images (2048 ! 1768
pixels) with a step size of 3 pixels and a subset size
containing between 3 ! 3 speckles and 5 ! 5 speckles on
average. Single image scans were used to avoid blurring
of integrated images caused by stress relaxation. As
a result, horizontal displacement shifts (as described in
Refs. 27–29) needed to be removed from the DIC data,
leading to blank lines visible in the displacement and
strain fields. Lagrangian strain was calculated from local
quadratic fitting of 15 ! 15 datum point subsets of the
distortion corrected27–30 displacement fields.
III. REPRESENTATIVE MICROSTRUCTURES
INVESTIGATED
FIG. 1. Inverse pole figure map of the initial coarse-grained microstructure following annealing at 773 K for 1 h.
2
The FOVs investigated in these 200 °C experiments
were selected to contain distinct representative microstructures created by ECAP processing. EBSD analysis
of 135,000 lm2 of the ECAP-processed microstructure
across seventeen tension specimens revealed that the
ECAP material consisted of 40% UFG microstructure,
36% supergrains, and 24% microstructure bands.
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
Note that the bands in the banded microstructure are
considered supergrains, but are not included in the
percentage of supergrains.
The microstructure details for the four FOVs are
provided in Table I and inverse pole figure maps are
shown in Fig. 2. The FOVs denoted by “S1-1” and
“S2-1” were composed primarily of ultrafine grains
separated by a large fraction of HAGBs. The FOV
“S3-1” was composed of microstructure bands, whereas
“S3-2” consisted of a large supergrain region and
a region of ultrafine-grains with orientations similar to
the supergrain region. In the FOV labels, the first digit is
the ECAP specimen number and the second digit is the
FOV number. Thus, S3-1 and S3-2 were from the same
tensile specimen with the FOVs spaced approximately
200 lm apart.
All the FOVs had a high percentage of HAGBs. This
was particularly true in S2-1 (984 lm HAGB
length:1563 lm total GB length), S3-1 (677 lm HAGB
length:916 lm total GB length), and S3-2 (505 lm
HAGB length:830 lm total GB length), where over half
of the grain boundaries were HAGBs. Slightly less than
half (748 lm HAGB length:1543 lm total GB length) of
the grain boundaries in S1-1 were HAGBs. In addition,
all the FOVs had similar inclination of the grain
boundaries from ECAP processing. This angle preference
was centered about the predicted grain inclination
angle of 26.6° for a cubic element passed through
TABLE I. Microstructural characteristics for the four FOVs investigated.
Microstructure identifier
FOV size (lm ! lm)
S1-1
S2-1
S3-1
S3-2
45
46.5
42.7
42.7
!
!
!
!
Average grain size (lm)
% HAGBs
1.97
1.92
3.11
2.82
48
60
74
61
38.8
40.1
36.9
36.9
Mode of GB trace angle (°)
27.2°
17.0°
22.4°
28.8°
6
6
6
6
25.1°
23.3°
14.4°
21.4°
FIG. 2. Inverse pole figure maps of the FOVs investigated in the 200 °C strain rate jump tension tests. These FOVs are composed of UFG regions,
microstructure bands, and supergrains common in the ECAP-processed microstructure. Areas (a) S1-1 and (b) S2-1 are composed primarily of ultrafine grains,
whereas (c) S3-1 is composed of microstructure bands and (d) S3-2 contains a large supergrain region and a large area of ultrafine grains.
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
3
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
a single ECAP pass.15,31 There were also large areas of
grains with orientations near f111gÆ!211æ, denoted with
purple in the inverse pole figure maps. This orientation,
while present in the initial coarse-grained microstructure
in Fig. 1, was prevalent in the ECAP-processed material.
IV. EXPERIMENTAL RESULTS AND DISCUSSION
A. Macroscopic observations
The macroscopic response of UFG Al at 200 °C was
significantly different from the room temperature
response. At 200 °C, the material possessed a lower
yield strength (59.1 6 6.9 MPa versus 93.0 6 11.2 MPa)
and a 3x greater steady-state strain rate sensitivity
(m 5 0.057 6 0.012 versus m 5 0.018 6 0.007). Note
that the strain rate sensitivity was calculated with the
equation m ¼ d ln r=d ln e_ , where r is the engineering
stress calculated from the stage mounted load cell and e_ is
the strain rate input into the stage control software. Thus,
this value represents the macroscopic strain rate sensitivity of the material. Figure 3 shows the stress–strain
curves for strain rate jump tests at room temperature
and 200 °C. The room temperature stress–strain curve,
which followed the same loading procedure as the 200 °C
tests, did not show strain hardening for any load steps
except for the initial loading. This can be explained by
the failure of the room temperature specimen by a narrow
neck that formed shortly after the macroscopic yield
strength was reached, as is evident in Fig. 4(a). The neck
formed during the first high strain rate load step. The
second high strain rate load step from 7.5 to 10% strain
resulted in a significant narrowing of the gage section in
the necked region. This explains the small increase
FIG. 3. Macroscopic stress–strain curve calculated from load cell and
grip displacement data, displaying differences in the strain rate
sensitivity at room temperature and 200 °C. Work hardening is not
apparent in the room temperature experiment with the exception of the
initial load step. Work hardening is present in all low strain rate load
steps of the 200 °C experiments except the last. Segments labeled
“"5” and “"3” were loaded at 2.5 ! 10"5 s"1 and 3 ! 10"3 s"1,
respectively.
4
in flow stress that occurred upon the high strain rate
loading, as shown in Fig. 3. An angled crack, which may
have formed during the high strain rate step, grew during
the subsequent low strain rate step from 10 to 12.5%
strain. This is observed as a steep drop in the stress–strain
curve during this load step at a macroscopic strain of
approximately 11%.
Shear bands, marked in Figs. 4(b)–4(d), were more
prevalent than necking at 200 °C and propagated more
quickly at high applied strain rate. As shown in Fig. 3 for
specimen S3-1/2, strain hardening (defined here as
a continuous increase in flow stress throughout the load
step) was apparent throughout the first three low strain
rate load steps of the 200 °C test. However, a shear band
which formed early in the loading grew rapidly during
the third high strain rate load step, after which work
hardening was no longer apparent. S1-1 showed a narrow
shear band with no observable necking, whereas S2-1
and S3-1/2 show a combination of necking and wider
shear banding. The cause of this is not known for S2-1,
but for S3-1/2, the wider shear band probably resulted
from the additional localized deformation at the right of
the gage section.
At 200 °C, it is likely that diffusion at low strain rates
improved ductility through the recovery of dislocations
and the initiation of diffusion deformation mechanisms
which allowed plastic deformation to spread over a wider
area of the gage section. This can be seen in Fig. 4, where
FIG. 4. Post-test gage sections for (a) room temperature specimen and
(b–d) 200 °C specimens: (b) S1-1, (c) S2-1, and (d) S3-1/2. The 200 °C
gage sections (b), (c), and (d) show similar shear band formation and
plastic deformation across the entire gage section, whereas the room
temperature gage section has a narrow necked region and minimal
surface deformation visible away from the neck. All the specimens
were subjected to the same load profile. The black rectangles mark the
locations of the FOVs. Note that the rectangles are approximately
500% larger than the actual FOV size.
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
plastic deformation is visible in the gage sections of the
200 °C specimens [Figs. 4(b)–4(d)], and not the gage
section of the room temperature specimen [Fig. 4(a)].
To provide the reader with context on what is
occurring on the macroscopic scale when viewing the
DIC results, an example of shear band formation with
increasing load is shown in Fig. 5 for specimen S2-1.
Please note the “I” and “II” areas of the gage section as
shown in the top image of Fig. 5. Table II provides
measurements of the incremental strain in area I that
eventually developed a shear band and area II where the
DIC FOV was located. Even during the early load step
from 2.5 to 5% macroscopic strain, strain began to
localize more strongly in area I which would eventually
develop a shear band. A large jump in the incremental
strain of area I, corresponding to rapid growth of the
developed shear band, occurred during the load step
from 7.5 to 10% macroscopic strain. This shear band
growth during the high strain rate load step from 7.5 to
10% was observed in the gage sections outside of the
DIC FOVs in all the 200 °C strain rate sensitivity tests.
The growth of the shear band corresponded to a drop in
the mean axial strain during this load step in the DIC
FOVs, located away from the shear band, as shown in
line graphs at the bottom of Figs. 6–9 which will be
discussed in more detail in the following paragraphs.
However, even after this drop at high strain rate, strain
localization returns to the DIC FOVs at low strain rate.
Strain localization nearly disappeared in the DIC FOV at
the high strain rate load step from 12.5 to 15% (not
presented here) yet returned in the following low strain
rate load step (strain values were just above the noise
threshold and are not presented here).
B. General DIC observations
While the macroscopic material responses were quite
different at 200 °C and room temperature, the microscale
strain localization showed many similarities. In both
cases, the following behavior was observed.
(i) Dislocation slip was active in larger grains, supergrains, and microstructure bands on the slip system best
aligned with the final pass ECAP theoretical shear plane
and possessing the highest Schmid factor.
(ii) 80% of strain localization by length occurred on
HAGBs (mean misorientation of 29.5° 6 14.1°) angled at
near 26.6° from horizontal (mean trace angle of 22.5° 6
18.1°), which is the predicted grain inclination angle for
a single ECAP pass. Details on these calculations are
provided in the following paragraph.
(iii) In the banded microstructure of S3-1, shear strain
(exy) localized nearly exclusively at the HAGBs separating microstructure bands. Sharp shear offsets resulting
from GB sliding were also apparent at these boundaries
during low strain rate load steps.
FIG. 5. Surface deformation throughout loading of S2-1. Strain
increments (2.5–5.0%), (7.5–10.0%), and (12.5–15.0%) were at a strain
rate of 3 ! 10"3 s"1. Strain increments (5.0–7.5%), (10.0–12.5%), and
(15.0–17.5%) were at a strain rate of 2.5 ! 10"5 s"1. Data presented in
the DIC strain fields are from macroscopic incremental strains ranging
from 2.5 to 12.5%.
C. Dependence of strain localization on strain rate
Incremental strain localization was measured to determine the impact of microstructure on strain rate sensitivity.
Incremental strain fields are shown in Figs. 6–9 for each
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
5
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
TABLE II. Tabular data corresponding to the images in Fig. 5. Even
in the incremental strain step from 2.5–5%, strain more favorably
localized in area I suggesting that a shear band was already formed
although it was not visible on the surface. After the shear band was
visible to the eye, strain nearly exclusively localized there.
Incremental strain (%)
Incremental
strain step
Strain rate (s"1)
2.5–5.0%
5.0–7.5%
7.5–10.0%
10.0–12.5%
12.5–15.0%
15.0–17.5%
3
2.5
3
2.5
3
2.5
!
!
!
!
!
!
10"3
10"5
10"3
10"5
10"3
10"5
I
II
4.0
2.0
6.8
4.0
9.1
8.1
3.3
2.4
2.1
2.3
0.4
0.6
FOV at macroscopic strain increments of (2.5–5%),
(5–7.5%), (7.5–10%), and (10–12.5%) as determined
by linear variable differential transformer (LVDT).
These coincide to two high strain rate load steps and
two low strain rate load steps, as shown in Fig. 3.
In Figs. 6–9, the bottom center plot shows the DICcalculated mean incremental strain for the entire FOV
incurred during each load step. The mean strain
response of S2-1 (Fig. 7) dropped with each load
increment, except for the last load increment during
which it remained nominally constant. In S1-1, S3-1,
and S3-2, the mean strain increased with each low strain
rate load step. In each test, the mean strain in the first
two load steps was near the applied strain of 2.5%. After
the second load step, the mean DIC FOV strain dropped
to approximately 1% as a result of strain localization
preferentially occurring at shear bands.
Grain boundaries (particularly HAGBs) exhibited the
greatest strain rate sensitivity, and the incremental mean
strain decreased (increased) when moving from a low to
high strain rate (high to low strain rate) as demonstrated
in Table III. Relationships between strain localization and
the underlying microstructure were determined by first
removing all strain data below the noise threshold and
manually selecting visible strain localization in each
strain component of each test. This yielded a binary
image where strain localization was black and the remainder of the FOV was white. The binary image was
then imported into Matlab where the black highlighted
strain was thinned to lines 200 nm wide and used to select
EBSD data that was aligned with the DIC FOV. Strain
localization which was aligned within 15° of a GB trace
and contained a GB was defined to be at that GB.
Otherwise the strain localization was assigned to the
grain where it was contained. Once the localization was
assigned to the GB or grain, relationships between the
strain localization and microstructural characteristics
from the EBSD data could be studied. An example
demonstrating this procedure can be found in the Supplemental Figure S1.
6
The greatest change in mean strain magnitude generally occurred at HAGBs, followed by low-angle grain
boundaries (LAGBs) and grain interiors. This is true even
in S2-1 in which strain rate sensitivity is not easily visible
in the strain fields of Fig. 7. This is demonstrated in
Table III where strain rate sensitivity is presented as the
mean of the absolute values of the change in incremental
strain when moving between strain rates. Boundaries
between microstructure bands in S3-1 showed even greater
strain rate sensitivity than HAGBs. The microstructure
band GB incremental strain always exhibited strain rate
sensitivity, with changes in mean incremental strain being
of a greater magnitude than at HAGBs, LAGBs, or within
grains. Other microstructural characteristics such as the
difference in adjacent grain size, angle between the highest
Schmid factor slip directions and plane normals, and
crystalline orientation were investigated, but there were
no apparent relationships to the strain rate sensitivity.
At high strain rates (strain increments of 0.025–0.050%
and 0.075–0.100%), lower strain localization occurred at
grain boundaries with magnitudes near the same level as
dislocation slip. This is best observed in Figs. 6 and 8
where strain localization was nearly absent from the
second high strain rate load step. Dislocation slip was
identified as strain localization aligned with a f111g plane
trace. To further confirm that it was dislocation slip, the
DIC in-plane displacement measurements were compared
to the calculated in-plane displacement components for the
Æ!110æ directions on the aligned f111g plane. At low strain
rates, diffusion-based deformation mechanisms allowed
for significant strain localization. We propose that
diffusion-based deformation mechanisms are active as in
other works5,7 due to the fact that high strain localization
occurred at HAGBs at low strain rates in these elevated
temperature tests. The same behavior was not witnessed in
room temperature experiments where diffusion is less
likely. Again, Figs. 6 and 8 demonstrate the return of
strain localization in the final low strain rate load step.
Additional strain fields and plots of the strain fields’
cumulative distribution can be found in Supplemental
Figures S2–S5.
Intragranular dislocation slip, which was active in all
observed supergrains and often apparent in the transverse
(eyy) strain fields, as labeled “IDS” in the first eyy strain
fields in Figs. 8 and 9, exhibited low strain rate
sensitivity. In turn, the small strain rate sensitivity in
supergrains caused by dislocation slip limited the strain
rate sensitivity at neighboring grain boundaries. The
reader should note that dislocation slip can appear in
different strain fields or be nearly unobservable in the
DIC strain fields depending on the orientation of the
active slip system. Slip occurring directly into the sample
surface, which is highly unlikely in uniaxial tension, will
probably not be measured by two-dimensional DIC. Slip
occurring on a slip plane perpendicular to the surface and
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
FIG. 6. Strain fields for S1-1 (200 °C). Rows (a) and (c) show the incremental strain fields for a strain rate of 3 ! 10"3 s"1. Rows (b) and (d) show
the incremental strain fields for a strain rate of 2.5 ! 10"5 s"1. Image size: 2048 ! 1768 pixels. Image resolution: 22.0 nm/pixel. DIC parameters
25 ! 25 pixel subset size, 3 pixel step size.
in a direction which lies in the sample face plane will
appear as primarily shear strain. For slip to be most
apparent in eyy, slip must be occurring primarily in the
y-direction. In S3-2, which contained a supergrain, active
dislocation slip caused the mean eyy to continually
decrease in magnitude. This is due to the fact that,
as observed in tests on coarse-grained aluminum, dislocation slip shows limited strain rate sensitivity.
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
7
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
FIG. 7. Strain fields for S2-1 (200 °C). Rows (a) and (c) show the incremental strain fields for a strain rate of 3 ! 10"3 s"1. Rows (b) and (d) show
the incremental strain fields for a strain rate of 2.5 ! 10"5 s"1. Image size: 2048 ! 1768 pixels. Image resolution: 22.7 nm/pixel. DIC parameters
25 ! 25 pixel subset size, 3 pixel step size.
Thus, a region of microstructure such as supergrains
that deformed by dislocation slip showed little strain
rate sensitivity. Once shear bands formed outside of the
DIC FOV, strain more favorably localized there,
8
causing the strain localization from slip to decrease in
magnitude.
The small strain rate sensitivity in supergrains from
dislocation slip in S3-2 limited the strain rate sensitivity
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
FIG. 8. Strain fields for S3-1 (200 °C). Rows (a) and (c) show the incremental strain fields for a strain rate of 3 ! 10"3 s"1. Rows (b) and (d) show
the incremental strain fields for a strain rate of 2.5 ! 10"5 s"1. The black areas cover failed correlation resulting from the Pt markers. Image size:
2048 ! 1768 pixels. Image resolution: 20.8 nm/pixel. DIC parameters 25 ! 25 pixel subset size, 3 pixel step size.
at the neighboring grain boundaries as specified by
“IG/GB” for intragranular/grain boundary in Fig. 9(b).
The highlighted strain localization appears similarly in
each strain component, meaning that the deformation is
occurring in the same direction and possesses a similar
magnitude at grain boundaries and slip bands. This
resulted in the strain rate sensitivity at HAGBs in eyy,
where strain localization from slip was most apparent, to
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
9
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
FIG. 9. Strain fields for S3-2 (200 °C). Rows (a) and (c) show the incremental strain fields for a strain rate of 3 ! 10"3 s"1. Rows (b) and (d) show
the incremental strain fields for a strain rate of 2.5 ! 10"5 s"1. The black areas cover failed correlation resulting from the Pt markers. Image size:
2048 ! 1768 pixels. Image resolution: 20.8 nm/pixel. DIC parameters 25 ! 25 pixel subset size, 3 pixel step size.
be the lowest of any test. This is probably the result of
several grain boundaries being elongated and parallel
to the traces of the slip planes active in the neighboring
supergrain. Since plastic deformation occurred on
10
these grain boundaries to accommodate slip, which
showed little strain rate sensitivity, the strain rate
sensitivity of these grain boundaries was similarly
small.
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
TABLE III. Mean of the absolute value of the change in strain that occurs when switching between strain rates. A “–” is marked in entries where
strain rate sensitivity, defined as a decrease (increase) when moving from a low to high strain rate (high to low strain rate), was not observed in all
strain rate changes. Strain rate sensitivity is most prominent at microstructure band boundaries, followed by HAGBs, LAGBs, and grain interiors.
S1-1
S1-1
S1-1
S2-1
S2-1
S2-1
S3-1
S3-1
S3-1
S3-1
S3-1
S3-2
S3-2
S3-2
HAGBs
LAGBs
Grain interiors
HAGBs
LAGBs
Grain interiors
HAGBs
LAGBs
Grain interiors
HAGBs (not microstructure band boundaries)
Microstructure band boundaries
HAGBs
LAGBs
Grain interiors
The fact that dislocation slip showed little strain rate
sensitivity and remained active at high strain rate can
explain shear band formation in tests at 200 °C. Recovery
of dislocations during the low strain rate load steps
permitted easier dislocation slip at high strain rate. This
phenomenon, in combination with the favorable alignment of active slip systems with the plane of maximum
resolved shear stress, encouraged shear banding.
Cooperative diffusion along grain boundaries, many of
which had orientations similar to that of the shear bands,
was probably also a contributing factor.
V. CONCLUSIONS
The strain rate sensitivity of UFG high purity aluminum was examined at room temperature and 200 °C
for strain rates alternating between 2.5 ! 10"5 s"1 and
3.0 ! 10"3 s"1. The results presented here are the first
full-field, quantitative investigations of the effect of
microstructure on strain rate sensitivity. This work
yielded the following conclusions:
(1) High strain localization was visible in the gage
section of the 200 °C specimens as shear bands. These
shear bands grew at a faster rate at high strain rate,
whereas at low strain rate, strain localization was distributed across the gage section.
(2) At room temperature, the specimens failed by
necking. Little plastic deformation was apparent outside
of the neck. At 200 °C, plastic deformation was apparent
across the entire length of the gage section.
(3) Even after shear bands were visible in the gage
section, strong strain localization occurred at the grain
boundaries at low strain rate. This was apparent in the
DIC strain fields, where little localization was apparent at
the second high strain rate load step, but strain localization returned at the following low strain rate load step.
MeanjDexxj
MeanjDeyyj
MeanjDexyj
0.0151
0.0145
0.0098
–
–
–
0.0177
0.0153
0.0086
0.0168
0.0181
0.0116
0.0114
–
0.0115
0.0099
0.0089
0.0043
0.0038
0.0033
0.0059
0.0048
–
0.0049
0.0063
0.0034
0.0029
–
0.0049
0.0065
0.0082
0.0033
–
–
0.0144
0.0114
0.0060
0.0128
0.0150
0.0062
0.0047
0.0031
(4) The greatest strain rate sensitivity was observed at
HAGBs, particularly those separating microstructure
bands, due to diffusion deformation mechanisms that
are more active at 200 °C during the low strain rate load
steps. Intragranular strain localization displayed the
lowest strain rate sensitivity.
(5) Strain rate sensitivity was reduced by the presence
of supergrains because of the relatively high activity of
dislocation slip in these structures.
(6) Dislocation slip in supergrains can limit the strain rate
sensitivity in neighboring regions of ultrafine grains separated
by HAGBs. This was observed in S3-2 where the grain
boundaries were elongated and parallel to the traces of slip
planes active in neighboring supergrains, causing significant
plastic deformation in these boundaries to accommodate slip.
ACKNOWLEDGMENTS
This work was supported by the National Science
Foundation under Grant No. 092753 and by the Rackham
Graduate School Non-Traditional Student Fellowship from
the University of Michigan. The authors would like to
acknowledge Sara Nitz for her sample preparation work and
experimental assistance. Portions of this work were performed at the Electron Microbeam Analysis Laboratory
(EMAL) at the University of Michigan and at the Lurie
Nanofabrication Facility (LNF), a member of the National
Nanotechnology Infrastructure Network, which is supported
in part by the National Science Foundation. This work was
also supported in part by the National Science Foundation
of the United States under Grant No. DMR-1160966.
REFERENCES
1. R.Z. Valiev and T.G. Langdon: Principles of equal-channel
angular pressing as a processing tool for grain refinement. Prog.
Mater. Sci. 51, 881 (2006).
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
11
IP address: 169.231.136.53
A.D. Kammers et al.: The microstructure length scale of strain rate sensitivity in ultrafine-grained aluminum
2. R.Z. Valiev, I.V. Alexandrov, Y.T. Zhu, and T.C. Lowe: Paradox
of strength and ductility in metals processed by severe plastic
deformation. J. Mater. Res. 17, 5 (2002).
3. J. Dvorak, V. Sklenicka, and Z. Horita: Microstructural evolution
and mechanical properties of high purity aluminium processed by
equal-channel angular pressing. Mater. Trans. 49, 15 (2008).
4. Y. Estrin and A. Vinogradov: Extreme grain refinement by severe
plastic deformation: A wealth of challenging science. Acta Mater.
61, 782 (2013).
5. J.M. Wheeler, V. Maier, K. Durst, M. Göken, and J. Michler:
Activation parameters for deformation of ultrafine-grained aluminium as determined by indentation strain rate jumps at elevated
temperature. Mater. Sci. Eng., A 585, 108 (2013).
6. J. May, H.W. Höppel, and M. Göken: Strain rate sensitivity of
ultrafine-grained aluminium processed by severe plastic deformation. Scr. Mater. 53, 189 (2005).
7. N.Q. Chinh, P. Szommer, T. Csanádi, and T.G. Langdon: Flow
processes at low temperatures in ultrafine-grained aluminum.
Mater. Sci. Eng., A 434, 326 (2006).
8. A. Böhner, V. Maier, K. Durst, H.W. Höppel, and M. Göken:
Macro- and nanomechanical properties and strain rate sensitivity
of accumulative roll bonded and equal channel angular pressed
ultrafine-grained materials. Adv. Eng. Mater. 13, 251 (2011).
9. H.W. Höppel, J. May, P. Eisenlohr, and M. Göken: Strain rate
sensitivity of ultrafine-grained materials. Z. Metallkd. 96, 566
(2005).
10. I. Sabirov, M.R. Barnett, Y. Estrin, and P.D. Hodgson: The effect
of strain rate on the deformation mechanisms and the strain rate
sensitivity of an ultra-fine-grained Al alloy. Scr. Mater. 61, 181
(2009).
11. Y. Iwahashi, J. Wang, Z. Horita, M. Nemoto, and T.G. Langdon:
Principle of equal-channel angular pressing for the processing of
ultra-fine grained materials. Scr. Mater. 35, 143 (1996).
12. T.G. Langdon: The principles of grain refinement in equal-channel
angular pressing. Mater. Sci. Eng., A 462, 3 (2007).
13. T.G. Langdon: Twenty-five years of ultrafine-grained materials:
Achieving exceptional properties through grain refinement. Acta
Mater. 61, 7035 (2013).
14. M. Furukawa, Z. Horita, and T.G. Langdon: Factors influencing
the shearing patterns in equal-channel angular pressing. Mater.
Sci. Eng., A 332, 97 (2002).
15. Y. Iwahashi, M. Furukawa, Z. Horita, M. Nemoto, and
T.G. Langdon: Microstructural characteristics of ultrafine-grained
aluminum produced using equal-channel angular pressing. Metall.
Mater. Trans. A 29, 2245 (1998).
16. I.J. Beyerlein and L.S. Tóth: Texture evolution in equal-channel
angular extrusion. Prog. Mater. Sci. 54, 427 (2009).
17. S.D. Terhune, D.L. Swisher, K. Oh-ishi, Z. Horita, T.G. Langdon,
and T.R. McNelley: An investigation of microstructure and grainboundary evolution during ECA pressing of pure aluminum.
Metall. Mater. Trans. A 33, 2173 (2002).
18. K. Oh-ishi, A.P. Zhilyaev, and T.R. McNelley: Effect of strain
path on evolution of deformation bands during ECAP of pure
aluminum. Mater. Sci. Eng., A 410–411, 183 (2005).
19. M. Kawasaki, Z. Horita, and T.G. Langdon: Microstructural
evolution in high purity aluminum processed by ECAP. Mater.
Sci. Eng., A 524, 143 (2009).
20. D.L. Davidson: The effect of a cluster of similarly oriented grains
(A supergrain) on fatigue crack initiation characteristics of clean
materials. In Fourth International Conference on Very High Cycle
Fatigue (VHCF-4), J.E. Allison, J.W. Jones, J.M. Larsen, and
R.O. Ritchie eds.; TMS: Warrendale, PA, 2007; pp. 23–28.
21. A.P. Zhilyaev, D.L. Swisher, K. Oh-ishi, T.G. Langdon, and
T.R. McNelley: Microtexture and microstructure evolution during
processing of pure aluminum by repetitive ECAP. Mater. Sci.
Eng., A 429, 137 (2006).
22. A.D. Kammers, J. Wongsa-Ngam, T.G. Langdon, and S. Daly:
The effect of microstructure heterogeneity on the microscale
deformation of ultrafine-grained aluminum. J. Mater. Res. 29,
1664 (2014).
23. K. Nakashima, Z. Horita, M. Nemoto, and T.G. Langdon: Development of a multi-pass facility for equal-channel angular
pressing to high total strains. Mater. Sci. Eng., A 281, 82 (2000).
24. M. Furukawa, Y. Iwahashi, Z. Horita, M. Nemoto, and
T.G. Langdon: The shearing characteristics associated with equalchannel angular pressing. Mater. Sci. Eng., A 257, 328 (1998).
25. K. Oh-ishi, Z. Horita, M. Furukawa, M. Nemoto, and
T.G. Langdon: Optimizing the rotation conditions for grain
refinement in equal-channel angular pressing. Metall. Mater.
Trans. A 29, 2011 (1998).
26. A.D. Kammers and S. Daly: Self-assembled nanoparticle surface
patterning for improved digital image correlation in a scanning
electron microscope. Exp. Mech. 53, 1333 (2013).
27. M.A. Sutton, N. Li, D.C. Joy, A.P. Reynolds, and X. Li: Scanning
electron microscopy for quantitative small and large deformation
measurements. Part I: SEM imaging at magnifications from 200 to
10,000. Exp. Mech. 47, 775 (2007).
28. M.A. Sutton, N. Li, D. Garcia, N. Cornille, J-J. Orteu,
S.R. McNeill, H.W. Schreier, X. Li, and A.P. Reynolds: Scanning
electron microscopy for quantitative small and large deformation
measurements. Part II: Experimental validation for magnifications
from 200 to 10,000. Exp. Mech. 47, 789 (2007).
29. A.D. Kammers and S. Daly: Digital image correlation under
scanning electron microscopy: Methodology and validation. Exp.
Mech. 53, 1743 (2013).
30. M.A. Sutton, N. Li, D. Garcia, N. Cornille, J-J. Orteu,
S.R. McNeill, H.W. Schreier, and X. Li: Metrology in a scanning
electron microscope: Theoretical developments and experimental
validation. Meas. Sci. Technol. 17, 2613 (2006).
31. Y.T. Zhu and T.C. Lowe: Observations and issues on mechanisms
of grain refinement during ECAP process. Mater. Sci. Eng., A 291,
46 (2000).
Supplementary Material
To view supplementary material for this article, please visit http://dx.doi.org/jmr.2015.58.
Reprinted from J. Mater. Res., A. Kammers, J. Wongsa-Ngam, T. Langdon, S. Daly.
Microstructure Length Scale Strain Rate Sensitivity in Ultrafine-Grained Aluminum. Copyright
(2015) Cambridge Journals.
12
J. Mater. Res., 2015
http://journals.cambridge.org
Downloaded: 01 Apr 2015
IP address: 169.231.136.53