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8/4/2019 Nano Structured Fe
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In: Severe Plastic Deformation ISBN 1-59454-508-1
Editor: Altan, Burhanettin, pp. 95-112 2005 Nova Science Publishers, Inc.
Chapter 1.6
STRUCTURE AND PROPERTIES
OF NEAR-NANOSTRUCTURED IRON
Bing Q. Hana, Farghalli A. Mohamed
band Enrique J. Lavernia
a
Department of Chemical Engineering and Materials ScienceaUniversity of California, Davis, CA
bUniversity of California, Irvine, CA
ABSTRACT
In the present study, the evolution of microstructure in pure iron during equal-
channel-angular pressing (ECAP) is investigated. The present work shows that a grain
size of approximately 200 nm was obtained after 8 passes. Because of the presence of
near-nanostructured microstructure and non-equilibrium grain boundaries after severe
plastic deformation, the material displays a distinct mechanical behavior as compared tothat of coarse-grained iron. During tensile deformation of the ECAP Fe, plastic
deformation with geometrical softening was observed, which differs from the behavior of
significant work hardening in the annealed Fe. In compression, a brief work-hardening
region followed by a long elastic-perfectly plastic deformation was observed. Asymmetry
of yield strength between tension and compression was observed, which was attributed to
the residual tensile internal stress after equal channel angular pressing, resulting in the
Bauschinger effect. The mobile dislocations in high-density dislocation regions are
believed to interact with dislocation cell blocks, triggering a local fast dynamic recovery,
which causes the material loss of strain hardening ability locally and necking starts
immediately. The elastic-perfectly plastic deformation in compressive deformation of
ultrafine-grained iron is attributed to strain instability or localization by shear banding.
Key words: equal-channel angular pressing, iron, microstructure, mechanical properties.
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia96
INTRODUCTION
Nanostructured (
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Structure and Properties of Near-nanostructured Iron 97
MATERIALS AND EXPERIMENTAL PROCEDURES
A commercial grade of 99.95% iron having a composition, in ppm, of Ni100, O86, Si75,
Co34, Al27, N11, P4.8, Ge4.6, Cr4.3, Cu3.9, B2.8, Ti1.3, C
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia98
EXPERIMENTAL RESULTS
The evolution of microhardness with pressing sequence is shown in Figure 1. The value
of the microhardness increases significantly after the first pass, modestly after the second
pass, and slightly during subsequent pressing. There is no significant difference of
microhardness in different orientations.
0
50
100
150
200
250
300
0 2 4 6 8
99.95% Fe
Longitudinal
TransverseMicrohardn
ess(Hv)
Number of passes
Figure 1. Evolution of microhardness of Fe with the number of pressing.
The intensity of peak (110) of X-ray diffraction patterns of annealed Fe and ECAP-8 Fe
is shown in Figure 2. The figure indicates that the intensity of peak (110) in ECAP-8 Fe
decreases and that the half-maximum intensity of diffraction peak is broadened as a result of
the ECAP processing. The peak broadening is attributed to both the small size of the
diffracting grains and the high internal strain introduced during ECAP. From five strong Fe
peaks (110), (200), (211), (220) and (310), the volume-averaged grain sizes (d) and the lattice
microstrain (e) can be estimated using the following equation (Klug and Alexander 1974):
2
2
)(4
)(
1
=
oo
s
sded
s (1)
wheres is the reciprocal space variable s=2sin/, and (s)o is the measured peak width. Byperforming a least-squares fit to 1/(s)o plotted against [s/(s)o]
2for all the measured peaks of
ECAP-8 Fe, dand e are determined to be about 235 nm and 0.046%, respectively.
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Structure and Properties of Near-nanostructured Iron 99
0
200
400
600
800
1000
1200
44.2 44.4 44.6 44.8 45
99.95% Fe Annealed
8 passes
Intensity
2
Peak (110)
Figure 2. X-ray diffraction pattern of Fe.
The microstructure of ECAP Fe was examined by TEM. After the first pass, well-defined
banded dislocation cell-blocks (CBs) are formed in the microstructure, with the length and
width of 0.5 1 m and 0.15 - 0.4 m, respectively, as shown in Figure 3 (a). In thefollowing passes, dislocation CBs in microstructure are further refined, as the evidence of
shorter length banded blocks. The length of the dislocation CBs decreases to approximately
0.5 m after 4 passes in the transverse direction although the width of the blocks has
insignificant change, and is 0.13 0.34 m, as shown in Figure 3 (b). A high density array ofdislocations in blocks is observed. Moreover, reasonably high proportions of GBs with high-
angle misorientations are observed since discontinuous circular rings in the selected area
electron diffraction (SAED) patterns. The grain sizes are approximately 0.2 m and 0.4 mon the transverse and longitudinal cross sections, respectively.
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia100
(a)
(b)
(c)
Figure 3. Microstructure (a) after 1 pass, (b) after 4 passes, and (c) after 8 passes viewing from the
direction transverse to the pressing direction; (d) microstructure and (e) selected area electron
diffraction patterns viewing from the direction parallel to the pressing direction after 8 passes.
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Structure and Properties of Near-nanostructured Iron 101
(d)
(e)
Figure 3. Microstructure (a) after 1 pass, (b) after 4 passes, and (c) after 8 passes viewing from the
direction transverse to the pressing direction; (d) microstructure and (e) selected area electron
diffraction patterns viewing from the direction parallel to the pressing direction after 8 passes
(Continued)
The microstructure of the ECAP Fe after 8 passes, as viewed from a cross section normal
to and parallel to the pressing direction, is shown in Figures 3 (c) and (d), respectively. Thereare some finer grains with dimensions less than 200 nm and some larger grains with
dimensions exceeding 500 nm in the microstructure of the ECAP-8 Fe. Although there were
not enough grains analyzed for a statistically significant average grain size determination, the
average grain size is estimated to be in the range of 200 - 300 nm. Inspection of the
microstructure reveals that some grains contain uniformly distributed dislocations with
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia102
slightly lighter diffraction contrast and lattice distortions near the grain boundaries (GBs).
The approximate integrity of circular rings in SAED patterns (Figure 3 (e)), suggests that
there are high proportions of the GBs with high-angle misorientations. Nevertheless, short
arcs and spots, which indicate the existence of preferred orientations (a fiber texture of
direction parallel to the pressing direction), were also observed from the SAED patterns. The
observation is in an agreement with the preferred orientation of an ultrafine-grained low-carbon steel after equal-channel-angular pressing (Shin et al. 2001).
The tensile behavior of the annealed Fe, ECAP-4 Fe and the ECAP-8 Fe is shown in
Figure 4 in terms of engineering stress as a function of engineering strain. The yield strength
of the annealed Fe at a strain of 0.2 pct is 79 MPa. There is an extensive region of work
hardening after yielding and a large elongation to failure for the annealed Fe. The ECAP Fe
exhibits a much higher tensile strength than that of the annealed Fe. The yield strength of the
ECAP-8 Fe, more than ten times stronger than annealed pure Fe, is almost identical to the
ultimate tensile strength with a value of 840 MPa. The plastic behavior of the ECAP Fe is
noticeably different from that of the annealed Fe. The ECAP-8 Fe exhibits a very brief low
work hardening region, for a strain of ~ 0.25%, then shows a continuous drop in the stress-
strain curve, indicating the occurrence of necking immediately after yielding. Moreover, theelongation to failure is much shorter than that of the annealed Fe.
0
200
400
600
800
1000
0 10 20 30 40 50 6
99.95% Fe
0
8 passes4 passes
Annealedat 1203 Kbefore ECAP
Engineeringstress(MPa)
Engineering strain (%)
= 1.0 x 10-3
s-1
T = 298 K
.
Figure 4. Tensile behavior of annealed Fe, ECAP-4 Fe and ECAP-8 Fe in the engineering stress-strain
curve.
The morphology of necking area and the fracture surface of the ECAP-8 Fe is shown in
Figures 5 (a) and (b), respectively. They are remarkably different from those of the annealedFe. The plastic deformation of in the ECAP Fe was concentrated in the necking area. The
vein-like patterns, resembling fracture via cleavage but different from the dimpled ductile
fracture, were observed on the fracture surface of the ECAP-8 Fe. Inspection of the
cleavage surface (see insert with a higher magnification) reveals that there is a subtle
banding structure with a width of ~ 0.3 m.
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Structure and Properties of Near-nanostructured Iron 103
(a)
(b)
Figure 5. (a) The morphology of necking area after failure and (b) fracture surface of ECAP-8 Fe.
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia104
(a)
(b)
Figure 6. (a) Microstructure and (b) selected area electron diffraction pattern taken within the shear
bands after tensile failure of ECAP-8 Fe.
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Structure and Properties of Near-nanostructured Iron 105
The microstructure at the position of necking after tensile failure is shown in Figure 6 (a).
A typical picture taken within the shear band displays elongated grains with a width of ~ 200
nm, which contain high dislocation densities. The elongated structure in shear bands is similar
to the columnar structure in rolling. The presence of circular rings in SAED patterns (Figure 6
(b)) suggests that the high proportions of the GBs with high-angle misorientations still exist.
The variation of microhardness measured from the failed tensile specimen of the ECAP-8Fe was plotted in Figure 7 as a function of distance along tensile direction starting from the
fracture surface. It is observed that the value of microhardness in the necking area is slightly
higher than that in the other areas. This observation suggests that there is slight work
hardening, instead of work softening, in the neck region. Therefore, the stress drop in the
tensile stress-strain curve is attributed to geometrical softening, i.e., the rapid decrease of the
cross-sectional area in the neck region.
0
50
100
150
200
250
300
0 2 4 6 8
99.95% Fe
10
Microhardness(HV)
Distance from fracture surface (mm)
8 passes
Gage section Shoulder section
Figure 7. Variation of microhardness along tensile direction.
Compressive testing results of the ECAP-8 Fe were plotted in Figure 8 in the form of the
true stress-strain curve, which was constructed by using the concept of volume constancy. An
elastic-perfectly plastic deformation was observed in compression of the ECAP Fe, whereas
the work-hardening behavior was observed in compression of the annealed Fe. Neither
bucking nor barreling was observed on specimens after compression. Nevertheless, shear
banding after compression, which is inclined at an angle of approximately 57.5 deg to the
compression axis was observed in the ECAP Fe.
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia106
0
200
400
600
800
1000
0 0.1 0.2 0.3 0.4 0.5
99.95% Fe
Tension
Compression
Truestress(MPa)
True strain
= 1.0 x 10-3
s-1
.
8 passes
IF
M
Figure 8. Tensile and compressive behavior of ECAP-8 Fe in the true stress-strain curve. M: point of
maximum tensile load; I: point of interrupted tensile test; F: point of tensile failure.
For comparison, tensile results in the form of the true stress-strain curve of the ECAP-8
Fe were also plotted in Figure 8. A close examination of compression and tension curves
indicates that they are similar with respect to the following aspect: a continuous increase of
the true stress in plastic deformation occurs. This finding again indicates that the engineering
stress drop in Figure 4 reflects geometrical softening due to a neck formation.The true tensile stress-true strain curve was constructed by using: (a) the concept of
volume constancy up to the point of maximum load (M on the tensile curve), and (b) the
actual cross sectional area beyond the point of the maximum load (onset of necking). There is
only one datum for true stress and true strain at the failure point (F on the tensile curve)
after the point of maximum load from a failed specimen. In order to reveal whether the
necking deformation is attributed to the formation of Lders bands, an additional tensile
testing was performed, which was interrupted at an engineering strain of 2.5%, a Lders band
was formed in the gage section, as shown in Figure 9. The additional value of true stress and
true strain at the position of Lders band was also plotted in Figure 9 (I on the tensile
curve). An inspection of the neck formed during deformation shows the presence of two
primary characteristics: (a) the neck assumes the shape of a narrow band with width nearly
equal to the thickness of the sample, and (b) the neck is inclined at angle of approximately58 deg to the testing axis. These characteristics are consistent with those reported for the local
necking in a sheet specimen. Also, there is a very small, narrow diffusive neck extending to
the two sides of the local neck.
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Structure and Properties of Near-nanostructured Iron 107
Figure 9. The morphology of Lders bands right after yielding of ECAP-8 Fe.
DISCUSSION
Microstructural Evolution
From the significant increases in microhardness with increasing number of passes, it is
indicated that equal-channel angular pressing is an effective approach to strengthen materials.
After the initial several passes, significant shear deformation occurs in coarse grains along the
pressing direction, resulting in the significant increase in microhardness and strength. The
increment of strength in the first and second passes is due to the development of intensive
dislocation cell-blocks. The density of dislocations introduced by shear deformation increases
dramatically in the initial several pressings and rapidly to a high level after 4 passes. With
increasing strain, some dislocations around block walls may have been rearranged to form the
dislocation boundaries with high-angle misorientations, leading dislocation CBs to a granular-
type structure. It is plausible that the deformation structures are in thermodynamic
equilibrium (i.e., in low-energy dislocation structures (LEDS)) (Kuhlmann-Wilsdorf 2002).
During the subsequent pressing deformation from 4 to 8 passes, the dislocation density
gradually approaches saturation in the deformation structures. After 8 passes, the dislocation
structures may be far from thermodynamic equilibrium, and are generally referred to as the
high-energy dislocation structures (HEDS), since dislocations in excess of those required to
accommodate the misorientations between walls of dislocation CBs may be accumulated in
the vicinity of GBs. The excess dislocations at boundaries are not arranged in LEDS, which
renders the grain boundary unstable.
It is well accepted that the strength increase due to work hardening is expressed by =
MGb1/2, where M = 2.75 is Taylor orientation factor for bcc structure, = 0.4 for bccmetals (Courtney 2000). In the present study, between annealed Fe and ECAP-8 Fe isabout 761 MPa. If G = 64000 MPa and b = 2.4810-10 s-1 for-Fe (Frost and Ashby 1982), theincrement of actual dislocation densities could be estimated to be about 1.91015 m-2 forECAP-8 Fe, which is slight lower than that of pure Fe processed by torsion at 293 K which
was measured to be about 31015 m-2 at a shear strain of 8 (Schafleret al. 1997).
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia108
Plastic Deformation
The tensile localized deformation of ECAP-Fe is strikingly different from typical plastic
deformation of the annealed pure Fe. While the annealed Fe shows a large work hardening
region where the load increases with increasing strain, the ECAP-Fe is plastically unstable, as
indicated by a continuous drop of load and a lack of any work hardening, i.e., the yield stressexceeds the rate of work hardening (y > d/d). According to Considres criterion(Courtney 2000), necking starts at the maximum stress when the increase in strength of the
materials due to work hardening is less than the decrease in the load-bearing ability due to the
decrease in cross-sectional area. For the annealed Fe, after an extended work hardening
region, the ultimate tensile strength is obtained when the necking deformation starts. The
tensile plastic deformation of ECAP-Fe is very localized and is restricted to a narrow area
where the first shear band formed (Figure 9). Out of the localized deformed zone, the
measurable uniform deformation is very low. Even though the rate of geometrical softening in
the annealed Fe after necking is slightly faster than that in the ECAP Fe, geometrical
softening in the ECAP Fe seems to be attributed to the necking deformation. Inspection of
TEM results (Figure 6) reveals that grains have been substantially elongated inside the shearbands. The grain morphology was changed from spherical to an elongated shape which was
parallel to the shear band. The phenomenon indicates the tremendous dislocation activity is
involved in the formation of shear banding. The mobile dislocations in dense dislocation
regions will interact with dislocation cell blocks, triggering a local fast dynamic recovery,
which causes the material loss of strain hardening ability and necking starts right away.
It is worth noting that accurate acquisition of displacement of gage section is very
important for plotting the correct stress-strain curves, since recently there are two reports on
the occurrence of unusual stress-strain curves in ECAP Fe (Fukuda et al. 2002, Sus-
Ryszkowska et al. 2004), in which there is a very long work-hardening region before the
ultimate tensile strength in the ECAP Fe. In the present study, the displacement of the gage
section was accurately acquired by a non-contact video extensometer. The present authors
also observed the similar unusual stress-strain curves (not reported in the present study), if thedisplacement of the gage section was not measured exactly by the video extensometer, but
replaced by the displacement of cross sections of the tensile machine. Therefore, the unusual
stress-strain curves of ECAP Fe as reported in above two references may be attributed to the
following reasons: the elastic deformation of the machine cross-sections and fixtures
(machine compliance) and/or the slippery distance of clips at the sample-shoulder sections
when applied under higher stresses were included in plotting the stress-strain curves.
The compressive deformation with a flow stress plateau after yield strength was observed
in the ECAP Fe. Shear bands were prevalently observed to exist in the compressive
deformation of high-strength nanocrystalline and submicron-grained Fe (Jia et al. 2003, Wei
et al. 2002). The intense, localized inhomogeneous plastic flow resulted from the deformation
of grains in these narrow bands.In related studies, the low work-hardening region in the tensile deformation of a
cryomilled ultrafine-gained Al-Mg alloy was attributed to dynamic recovery (Han et al.
2003c), since the high stacking fault energy of the Al alloy may facilitate dislocation slip and
thus dynamic recovery. In addition, the low activation energy for dynamic recovery may exist
in cryomilled Al alloys because of the presence of residual stress and an abundance of
structural defects in the cryomilled microstructure (Zhou et al. 2003).
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Structure and Properties of Near-nanostructured Iron 109
Although the occurrence of dislocation slip as well as dynamic recovery might be
difficult in overall compressive deformation in Fe because of the low stacking fault energy in
Fe, severe localized inhomogeneous plastic deformation in shear banding may accelerate
dynamic recovery. In addition, because of the saturation of dislocations and the existence of
non-equilibrium grain boundaries in the ECAP Fe (Han et al. 2004), the significant
contribution of dislocation accumulation to work hardening might be impeded. Therefore, theelastic-perfectly plastic deformation in the compression deformation of UFG Fe might be
attributed to strain instability or strain localization by shear banding.
TENSION-COMPRESSION ASYMMETRY
It is noteworthy that the yield strength of ECAP Fe in compression is lower than that in
tension, while, in several other nanostructured or ultrafine-grained materials, the compressive
yield strength is observed to be higher than that in tension (Carsley et al. 1998, Carsley et al.
1997, Han et al. 2003b, Hayes et al. 2001, Jain and Christman 1994, Jia et al. 2000a) or
equivalent to that in tension (Han et al. 2003c). In fact, the phenomenon of higher yieldstrength in tension than compression was also revealed in an earlier report on pure Fe
deformed under severe plastic deformation, e.g., via wire drawing (Langford and Cohen
1969).
In related studies, the presence of residual processing defects was considered to be
responsible for the asymmetry of low tensile strength than compressive strength in other
nanostructured materials (Berbon et al. 2001, Carsley et al. 1998, Han et al. 2003b, Jain and
Christman 1994, Rittner et al. 1997). It is well established that plastic anisotropy, the
dependence of properties on orientation, is primarily attributed to texture (Dieter 1986). The
effect of a mechanical fibering, the alignment of a second phase parallel to the direction of
extrusion, on the plastic anisotropy and the asymmetry of yield strength was analyzed in an
as-extruded two-phase Al-10Ti-2Cu alloy (Han et al. 2003b). The existence of the mechanical
fibering in the as-extruded two-phase Al-10Ti-2Cu alloy results in not only the plastic
anisotropy, stronger strength and better ductility in the longitudinal direction than in the
direction perpendicular to the extrusion direction, but also the asymmetry of yield strength,
stronger yield strength in compression than in tension.
Inspection of the SAD patterns in the microstructure parallel to the pressing direction
(Figure 3 (e)) reveals that a preferred orientation (texture) of grains was produced after severe
plastic deformation. Although it is difficult to understand the inverse asymmetry of yield
strength in ECAP Fe (stronger yield strength in tension than in compression) on the basis of
the existence of texture, the possible role of texture strengthening on the asymmetry of yield
strength in ECAP Fe cannot be completed ruled out.
Another aspect that should be considered for the occurrence of the asymmetry of yield
strength is the effect associated with the presence of internal microstrain in the ECAP Fe. Onthe basis of analysis of the X-ray diffraction patterns, it is found that there is a microstrain
with the magnitude of 0.046 pct in the ECAP Fe. The presence of high internal strain was
already reported in several other ultrafine-grained materials processed via severe plastic
deformation (Nazarov et al. 1994, Schafleret al. 1997), which can be described well in terms
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Bing Q. Han, Farghalli A. Mohamed and Enrique J. Lavernia110
of nonequilibrium grain boundaries containing disordered extrinsic grain boundary
dislocations of high density.
Under equal-channel angular pressing, a large shear strain of 1.15 per pass is introduced
into the materials through two channels with 90 deg via dislocation slip (Segal 1995). The
morphology of a unit cell before and after pure shear deformation during ECAP is illustrated
in Figure 10. A square unit cell (abcd) in the vertical channel is sheared into a rhombohedralshape (abcd) within the exit channel after pure shear deformation. The unit cell is
elongated along the longitudinal direction having an angle of approximately 26.6 deg with the
exit direction (Iwahashi et al. 1998, Segal 1995, Zhu and Lowe 2000). The deformation of the
unit cell is in an excellent agreement with the experimental observation on the grain
deformation after one pass (Han et al. 2003a). Therefore, the pure shear deformation during
equal channel angular pressing results in the elongation of grains, analogous to the
circumstances of a tensile deformation employed on grains after each pass. Although most of
energy loss results from dislocation annihilation and rearrangement to form granular ultrafine
structures from dislocation cell blocks after severe plastic deformation for 8 passes (Han et al.
2004), the ECAP Fe should store a small fraction of the energy of deformation, which in turn
have a significant effect on diffusion and plastic deformation (Nazarov et al. 1993, Valiev etal. 2000). In summary, the residual tensile internal stress after equal channel angular pressing
results in the so-called Bauschinger effect (Dieter 1986), which leads to the lower strength in
compression than in tension during subsequent deformation.
a b
cdb
da
c
45o
26.6o
Exit
Entrance
Figure 10. Schematic of the shape change of a unit cell before and after one pass.
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Structure and Properties of Near-nanostructured Iron 111
CONCLUSIONS
Pure Fe was processed by means of equal-channel-angular pressing. The value of
microhardness increases with increasing number of pressing, with a saturation of the eighth
pass. Dislocation cell-blocks were obtained after pressing and gradually evolved into grains
with high-angle misorientations. A grain size of approximately 200 nm was obtained after 8
passes. In tension, plastic deformation with geometrical softening was observed in the ECAP
Fe, which is different from strain hardening in the annealed Fe. In compression of the ECAP
Fe, a strain-hardening region followed by an elastic-perfectly plastic deformation was
observed. The residual tensile internal stress after equal channel angular pressing might result
in the Bauschinger effect, which leads to the lower strength in compression than in tension
during the subsequent deformation. The mobile dislocations in dense dislocation regions
might interact with dislocation cell blocks, triggering a local fast dynamic recovery, which
causes the material loss of strain hardening ability and necking starts right away. The elastic-
perfectly plastic deformation in the compression deformation of UFG Fe might be attributed
to strain instability or strain localization by shear banding.
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ACKNOWLEDGMENTS
Support from the Army Research Office under Grant No. DAAD19-03-1-0020 is
gratefully acknowledged.