nano structured fe

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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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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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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(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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(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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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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(a)

(b)

Figure 5. (a) The morphology of necking area after failure and (b) fracture surface of ECAP-8 Fe.


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(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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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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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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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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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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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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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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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.

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