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Stable ductility of an electrodeposited nanocrystalline Ni-20wt%Fe alloy in
tensile plastic deformation
Junwei Mu, Xuesong Li, Lei Zhao, Zhonghao Jiang, Jianshe Lian, Qing Jiang
PII: S0925-8388(12)02132-9
DOI: http://dx.doi.org/10.1016/j.jallcom.2012.11.137
Reference: JALCOM 27372
To appear in:
Received Date: 27 September 2012
Revised Date: 21 November 2012
Accepted Date: 22 November 2012
Please cite this article as: J. Mu, X. Li, L. Zhao, Z. Jiang, J. Lian, Q. Jiang, Stable ductility of an electrodeposited
nanocrystalline Ni-20wt%Fe alloy in tensile plastic deformation, (2012), doi: http://dx.doi.org/10.1016/j.jallcom.
2012.11.137
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Stable ductility of an electrodeposited nanocrystalline Ni-20wt%Fe alloy in
tensile plastic deformation
Junwei Mua, Xuesong Li
b, Lei Zhao
a, Zhonghao Jiang
a,, Jianshe Lian
a, Qing Jiang
a
aKey laboratory of Automobile Materials, Jilin University, Nanling Campus, Changchun 130025, China
bKey Laboratory of Advanced Structural Materials, Ministry of Education,Changchun University of Technology,
Changchun 130012, China;
Abstract
Tensile behavior of an electrodeposited nanocrystalline Ni-20wt%Fe (average grain size
d ~32nm) alloy was investigated. With the variety of strain rate, high ultimate strength
(1762-1939MPa) and stable fracture ductility (8.5-9.3%) were observed during tensile tests. The
ductility of the nanocrystalline Ni-20wt%Fe alloy is more stable than that of the nanocrystalline
Ni (5.6-11.3%) with similar microstructures. The stable ductility of the nanocrystalline
Ni-20wt%Fe alloy can be attributed to its higher work hardening ability. Transmission electron
microscope analysis revealed that there are massive dislocations, deformation twins and stacking
faults in the deformed nanocrystalline Ni-20wt%Fe alloy. The decrease of stacking fault energy,
caused by alloying of Fe element, should be responsible for the crystal defect microstructures and
this increases work hardening rate, which can improve the ductility at last.
Keywords: A. Nanocrystalline; D. Ni-20wt%Fe alloy; E. Tensile; D. Stacking fault energy; D. Work
hardening rate
Corresponding author. Tel.+86 516 83591876; fax: +86 516 83591870.
E-mail address: [email protected] (Z.H. Jiang).
1. Introduction
Compared with their coarse-grained counterparts, nanocrystalline (nc) metals [1-2] have higher
strength but instable ductility. For structural materials, strength and ductility are two key material
parameters. Therefore, it is necessary to improve the ductility of nc metals without sacrificing the
strength for their structural applications. Several approaches have been proposed to improve the
ductility of nc metals, such as developing bi-modal grain size distribution [3-4], introducing
second-phase particles [5] and preexisting twins [6]. However, most studies about the ductility of
high-strength nc metals have focused on the maximum value [4, 7-8], few attention have been paid
to the stability of the ductility varied with strain rates. The instable ductility of nc metals has been
attributed to processing artifacts (e.g. impurities) and mechanical instability due to a lack of work
hardening ability [9].
The uniform plastic strain of nc metals, related to the work hardening ability, play a significant
role in ductility. In order to improve the ductility of nc metals, the uniform plastic strain should be a
main factor considered. One of the reasons to limit the ductility, particularly uniform plastic strain,
is the propensity for plastic instability (inhomogeneous deformation such as necking in tension and
shear banding) in the early stage of plastic strain. In these cases, the tensile engineering stress-strain
curve peaks at small plastic strain and then plunges down as the localized deformation leads to
fracture failure at an accelerated pace. This propensity of plastic instability in nc metals during ten-
sile deformation is related to the diminishing work hardening ability [10]. As we known, the sam-
ples of nc metals prepared by electrodeposition have few impurities [11] and the work hardening
ability of nc metals can be enhanced by the method of alloying [12-15]. Alloying can decrease the
stacking fault (SF) energy of nc metals, so that partial dislocations can emit easily from grain
boundaries (GBs) and the emitted partial dislocations can further induce twinning. Previous inves-
tigations of nc Ni-Fe alloy revealed that deformation processing could lead to complicated structur-
al evolutions, producing a high density of deformation twins accompanied by dislocations in nc
grains [16-17]. This indicated that twins are effective in increasing the dislocation storage capacity,
which could increase work hardening ability and thus improve ductility [18-20] and strength [21].
Recent reports [12, 22-24] showed that some face-centered cubic (fcc) nc alloys prepared by elec-
trodeposition exhibit good ductility with a satisfactory strength.
In this work, the effect of strain rate on the ductility of nc Ni-20wt%Fe (Ni-20Fe) alloy prepared
by direct current electrodeposition was investigated by tensile tests under a board strain rate ranges.
A primary aim of this work is to present the results elucidating SF energy effect on the deformation
and fracture behavior of the nc Ni-20Fe alloy with the average grain size of ~32 nm, compared with
the nc Ni with similar microstructures reported by Shen et al. [4].
2. Experimental procedures
Bulk nc Ni-20Fe alloy was prepared by a direct current electrodeposition technique. The plating
bath was composed of 210 g/L nickel sulfate, 15 g/L ferrous sulfate, 42 g/L boric acid, 26 g/L so-
dium chloride, 20 g/L sodium citric acid and a few additives. The pH value was adjusted to 3.5 at a
temperature of 62±1. The current density of the cathode is 5 A/dm2. In order to obtain high pure
nc alloy, the electrolyte was strictly purified for a week under a low current density range of 0.1-0.5
A/dm2 before electrodeposition. The as-deposited sheets have a dimension of ~120 mm 100 mm
0.5 mm, which were deposited on low carbon steel sheets of a thinness of ~2 mm. The dog-bone
shaped samples with a gauge cross-section of ~2.5 mm × 0.5 mm and a gauge length 8.0 mm were
cut from the as-deposited sheets using an electrodischarging machine and then were polished to
mirror-like finish surface using 0.5 m diamond suspension. Tensile tests (at least 3 tests for each
strain rate) were carried out on MTS-810 system at a strain rate range from 1.35×10−5
to 1.35 s−1
and room temperature (RT). Morphologies of deformation and fracture surface of the samples were
examined by scanning electron microscope (SEM, JSM-5600). Foil samples for transmission elec-
tron microscope observation were prepared by ion beam thinning. Microstructures and grain sizes
of the samples before and after deformation were observed using transmission electron microscope
(TEM, JEM-2100F) under accelerating voltage of 200 kV. TEM foil samples after tensile tests were
prepared as follows. The fractured tensile samples were mechanically thinned down to ~24 µm
thick. At the location of the fracture surface, 3 mm×2 mm wide slices were then cut parallel to the
tensile axis, followed by Ar ion milling using a RES 101 Rapid Etching System.
3. Results and discussion
3.1. Microstructures of as-deposited samples
Fig. 1a and b show the TEM bright field image of the as-deposited nc Ni-20Fe alloy and nc Ni
with the corresponding selected area diffraction (SAD) pattern, respectively. Fig. 1c and d show the
grain size statistical distribution of the nc Ni-20Fe alloy and nc Ni obtained using the information
presented in Fig. 1a and b. Statistical analysis of 200 grains in each of the two images indicated that
the average grain size of the nc Ni-20Fe alloy is ~32 nm calculated based on the number frequency,
while the nc Ni is ~37 nm. Most of grains in the nc Ni-Fe alloy and nc Ni have a size ranging from
25 nm to 45 nm.
3.2. Tensile results of nc Ni-20Fe alloy and nc Ni
3.2.1. Tensile properties
Fig. 2a shows the engineering stress-strain curves of the nc Ni-20Fe alloy deformed under a
broad strain rate range of 1.35×10−5
to 1.35 s−1
. The nc Ni-20Fe alloy shows the ultimate tensile
strength ( UTS ) of 1762-1939 MPa and the yield strength ( 2.0 ) of 1134-1368 MPa with the fracture
ductility ( f ) of 8.5-9.3% and uniform plastic strain ( u ) of 6.3-7%. While the nc Ni shown in Fig.
2b has UTS of 1605-1916 MPa, 2.0 of 876-1307 MPa, f of 5.8-10.7% and u of 4.6-6.9%.
To compare the performance of the nc Ni-20Fe alloy with the nc Ni visually, the variations of UTS ,
2.0 , f and u of the nc Ni-20Fe alloy and nc Ni with strain rate were shown in Fig. 2c and d,
respectively. A noticeable characteristic of the nc Ni-20Fe alloy shown in Fig. 2c and d is that the
strength and ductility are insensitivity to strain rate as compared with nc Ni.
3.2.2. Strain rate sensitivity and activation volume
The strain rate sensitivity and the activation volume have been widely used to interpret the un-
derlying deformation mechanism in nc metals. The strain rate sensitivity (m) of flow stress can be
defined as:
log
log
σm (1)
where and are the flow stress and strain rate, respectively. The m value was estimated from
the logarithm plot of flow stress with , as shown in Fig. 3a and b, for the nc Ni-20Fe alloy
and nc Ni, respectively. It can be seen that the m value (measured from the slopes) of the nc
Ni-20Fe alloy decreases from 0.0166 to 0.0080 with increasing strain from 0.2% to 2.5%, while the
m value of the nc Ni decreases from 0.0293 to 0.0158. The variations of m values of the nc Ni-Fe
alloy and nc Ni during deformation are shown in Fig. 3c. The activation volume of flow stress (V )
can be given by:
ln3KTV (2)
where K is the Boltzmann constant and T is the absolute temperature. The variations of the V values
of the nc Ni-Fe alloy and nc Ni during deformation are shown in Fig. 3d, where b is the Burgers
vector. The V value of the nc Ni-20Fe alloy and nc Ni increase from 22.2 b3 to 24.9 b
3 and 12.7 b
3
to 15.2 b3, respectively, with increasing strain from 0.2% to 2.5%. It has been accepted widely that
the plastic deformation of nc metals can be explained by the dislocation-based and GB-based me-
chanisms due to their high volume fractions of GBs and triple junctions [24]. As shown in Fig. 3d,
the V value of the nc Ni-20Fe alloy increases from 22.2 b3
to 25.1 b3 with increasing strain from
0.2% to 1.5%. However, after the strain of 1.5%, the V value is almost constant with increasing
strain, which indicates that GB activities should participate in the initial deformation for the small V
value. With increasing strain, an increase in the dislocation emission will occur during further de-
formation, which leads to the increasing of the V value. While at large strains (>1.5%), the V value
will become more stable because the dislocation density no longer increases significantly. Com-
pared with the nc Ni-20Fe alloy, the V value of the nc Ni increases monotonously from 12.7 to 15.2
b3. This indicates both dislocation activities and GB activities participate in the whole deformation
process and the dislocation density increases with increasing strain. The above results are in the
reasonable agreements with the corresponding variations of the m values shown in Fig. 3c and indi-
cate a significant role of the Fe element in affecting the deformation mechanism of the Ni-20Fe al-
loy.
According to the model proposed by Asaro and Suresh [25], the critical athermal radius of partial
dislocation cr can be given by:
0
2 )]ˆ5[2
1ˆ51( ereerc (3)
where 1/ˆ GbSF is a reduced stacking fault energy, SF is the stacking fault energy (0.184
J/m2 for Ni-20Fe alloy and 0.214 J/m
2 for Ni [26]), for convenience of notation we have represented
the magnitude of the nucleating, leading, partial dislocation as b1 and 0r is an inner cutoff radius
on the order of b1. The cr could be calculated to be 1.92 b for the nc Ni-20Fe alloy and 2.01 b for
the nc Ni. Using this result to the estimate an activation volume 8.5V b3
for nc Ni-20Fe alloy
and 3.6V b3
for nc pure Ni, respectively. It is noticed that the equation predicted values are
smaller than the V values obtained from the experiments. As we known, if a partial dislocation is
emitted from a stress concentrator site such as a grain boundary facet, both of leading and trailing
partial dislocations will be emitted with a splitting width that depend on the stacking fault energy.
The partial dislocation emission is illustrated schematically in Fig. 4. During thermal activation, as
researched by Wei [27], the splitting width will in turn determine the activation distance and accor-
dingly the activation volume. Base on the model proposed by Asaro and Suresh [25] and the analy-
sis of Wei [27], one could expect the activation volume of plastic deformation of fcc metals should
be a function of stacking fault energy, which can be expected to be:
1
2
2
1* brV c (4)
where )12/(2
SFGb is the equilibrium splitting width between two partials, G is shear mod-
ulus (8.5×1010
Pa for Ni-20Fe alloy and 8.2×1010
Pa for Ni [28]). Based on Eq. (4), *V for the
nc Ni-20Fe alloy can be calculated to be 22.6 b3, which is almost consistent with the experimental V
values (22.2 b3
-24.9 b3). Thus, the partial dislocation activities should be the main deformation
mechanism. However, *V for the nc Ni is calculated to be 18.7 b3, which is slightly larger than
the experimental V values (12.7 b3 - 15.2 b
3). If deformation is controlled by GB activities, *V is
on the order of b3- b
3 [25]. That is, for the nc Ni, both dislocation activities and GB activities
should participate in the deformation and the GB activities would play a minor role for the experi-
mental activation volume of the nc Ni. Eq. (4) shows that *V is a monotonically increasing
function of stacking fault energy, when the average grain size of the nc metals is fixed. However,
the dislocations emitted from triple junctions were not considered, so that Eq. (4) could only be
used to estimate approximately the *V values, but nonetheless, such activation volume values
are in fact to rationalize for partial dislocation-based deformation. In particular, for fcc nc metals in
this grain size range, it was forecasted to be in the range 10-30 b3 consistent with available experi-
mental evidence, for the nc Ni with d ~37 nm is 11 b3-26 b
3 and the nc Ni-Co alloy with d ~
15nm is 12 b3 [4, 12].
3.2.3. Work hardening rate
At low strain rate, the dislocation activity from the GB source is partly restrained due to the lack
of the effective high stress concentration and in turn the work hardening rate will become small.
However, the GB activities will result in the enhanced strain rate sensitivity and the evidently de-
creased flow stress [29]. The GB activities at low strain rate will release the local high stress con-
centration, delay the immature onset of necking and sustain a large u . Therefore, both of the nc
Ni-20Fe alloy and nc Ni exhibits a large f of 9-11% at low strain rate.
At high strain rate, the deformation mechanism of the nc Ni-Fe alloy and the nc Ni is mediated
by the dislocation activities, but the ductility of the nc Ni-Fe alloy is obviously higher than that of
the nc Ni. A possible reason for this difference is high work hardening rate of the nc Ni-Fe alloy.
According to the theoretical model proposed by Gutkin, it is energetically favorable for the
Ni–20Fe alloy with low SF energy successively emitting partial dislocations from GBs [30]. Due to
the partial dislocation emission, twinning should be considered to have a significant contribution to
plasticity during tensile deformation. Further details of the microstructure are provided in the next
section. As twinning usually occurs as dislocation glide obstacles, therefore the dislocation storage
capacity is maintained in nc Ni-Fe alloy with low SF energy [30-31], which could enhance work
hardening ability and improve the mechanical properties of nc Ni-Fe alloy. Fig. 5 show the work
hardening rate (WHR) plotted of nc Ni-20Fe alloy and nc pure Ni versus true strain at = 1.35 s-1
.
WHR is defined as:
)/( tt (5)
where t is true stress and t is true strain. At t = 3.1 %, WHR equals to ~17630 for both of
nc Ni-20Fe alloy and nc pure Ni. With further deformation, WHR for nc pure Ni sharply decreases
to zero at t = 4.5 %, while WHR for nc Ni-20Fe alloy remains up zero to t = 6.8 %. As a result,
the onset of necking is delayed, and a higher uniform plastic strain is obtained for the nc Ni-20Fe
alloy.
3.3. Microstructures of deformed samples
For further illuminating the different ductility trend at high strain rate, SEM fracture morpholo-
gies of the nc Ni-Fe alloy and nc Ni deformation at strain rate of 1.35 s-1
are shown in Fig. 6a and b.
From these figures, it can be clearly seen that many deformation bands were observed on the profile
of fracture samples. These deformation bandings mean that local plastic deformation has occurred.
However, deformation bandings of nc Ni-Fe alloy is more uniform than those of nc Ni, which
demonstrate location deformation of the nc Ni-Fe alloy was greatly weakened, which delay onset of
plastic deformation and improve the ductility. Furthermore, there are many microcracks within
these bands as marked by white cirques in Fig. 6a and b. It should be noted that the edges of micro-
cracks of nc Ni-Fe alloy are much smoother than those of nc Ni, which could restrain microcracks
propagate effectively. Recent analysis pointed to the fact that the postuniform elongation is believed
to be caused by the propagation of the microcracks [24]. Therefore, postuniform deformation of nc
Ni-Fe alloy was delayed and this improve the uniform plastic strain and the ductility.
The above dislocation activities and GB activities during the tensile deformation may result in
some local motion of GBs and hence trigger possible grain growth, which was observed in nc Ni
[32], nc Cu [33] and nc Ni-Fe [16, 19]. However, statistical analysis of grains in TEM image of the
nc Ni-Fe alloy and nc Ni after deformation at =1.35 s-1
, as shown in Fig.7a and b, which exhi-
bited the average grain size is ~34 nm and ~38 nm, respectively. Therefore, no obvious grain growth
phenomenon was found in present samples. The reason for this discrepancy may be the limit plastic
deformation of tensile samples.
Recent studies by Liao and Gubicza et al. support that deformation twins and SFs occur upon
plastic deformation in nc fcc metals [34-35]. Fig. 8a shows a typical TEM image in <011> orienta-
tion of the nc Ni-Fe alloy after deformation at = 1.35s-1
with a Fourier transformation of the
white frame in the inset. To see the detailed structure of this area, a Fourier-filtered image of the
white frame is shown in Fig. 8b, in which microtwins and SFs could be clearly seen (marked with
white plotlines). The formation of microtwins and SFs, considered as a contributing deformation
mechanism for nc materials, indicates the operation of partial dislocation mediated processes [21,
36-37]. In order to detect lattice dislocations (marked with white T) as well as to determine their
Burgers vectors, local Burgers circuits[34] were employed to determine the Burgers vector of the
dislocation marked with a white ellipse in Fig. 8b. The Burgers vector of this lattice dislocation was
determined to be ]112[4/
ab on a )111(
plane, where a is the lattice constant of nc Ni-Fe alloy,
i.e., this lattice dislocation is a 60º full dislocation of the face-centered cubic lattice. Many other
dislocations can also be seen marked with white T. Statistical analysis of about 20 HRTEM images
of nc Ni-Fe alloy, dislocation density of nc Ni-Fe alloy is ~7.1×1016
m2. Such observations align
well with the previously reported results which frequently demonstrated the existence of trapped
dislocations in deformed nc Ni-Fe alloy [38-39]. However, there are fewer twins and dislocations in
nc Ni, the HRTEM was not shown. The frequent observation of dislocations in nc Ni-Fe alloy indi-
cates that dislocations motion is effectively blocked by the microtwins and SFs generated by partial
dislocation emission from the grain boundaries, which is similar with the previously research [21].
As a result, the dislocation storage capacity is maintained in nc Ni-Fe alloy with low SF energy,
which could enhance work hardening ability and improve the mechanical properties of nc Ni-Fe
alloy, especially the ductility.
4 .Summary
The direct current electrodeposited nc Ni-20wt%Fe alloy with average grain size of ~32 nm were
systematically investigated by tensile tests compared with nc Ni. With increasing strain rate from
1.35×10-5
s-1
to 1.35 s-1
, both of the nc Ni-Fe alloy and the nc Ni have a high strength, however, the
ductility of nc Ni-Fe alloy is more stable than that of nc Ni. The values of strain rate sensitivity and
activation volume reveal that the deformation mechanism operated by the dislocation activities in
the nc Ni-Fe alloy, while the deformation mechanism of nc Ni may relate to the common role of
dislocation activates and GB activates. Adding of Fe atoms decrease stacking fault energy of Ni and
improve partial dislocation emission. Twinning via partial dislocation emission from grain bounda-
ries becomes a major deformation mechanism and this enhance work hardening ability. So that nc
Ni-Fe alloy exhibit stable ductility without sacrificing strength.
Acknowledgements
The authors wish to thank Dr. X.X. Shen for providing the data and samples of nc pure Ni. This
work was financially supported by the National Nature Science Foundation of China (No.
50771049).
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Figure Captions
Fig. 1 TEM bright-field images and electron diffraction patterns of (a) nc Ni-Fe alloy and (b) nc
Ni; The grain size distribution of (c) nc Ni-Fe alloy and (d) nc Ni obtained using
information presented in (a) and (c).
Fig. 2 The engineer stress-strain curves of the electrodeposited (a) nc Ni-Fe alloy and (b) nc Ni
under uniaxial test at a broad strain rate range of 1.35×10-5
to 1.35 s-1
and RT; (c) Variation
of the ultimate tensile strength ( UTS ) and the yield stress ( 2.0 ) with strain rate ( ); (d)
Variation of the tensile fracture ductility ( f ) and the uniform plastic strain ( u ) with
strain rate ( ).
Fig. 3 Logarithm plots of flow stress and different plastic strain as a function of strain rate of (a)
nc Ni-Fe alloy and (b) nc Ni; Variation of (c) m values and (d) V values of nc Ni–Fe alloy
and nc Ni with plastic strain ( p ).
Fig. 4 Schematic illustration of the emission of a leading partial dislocation (Pl) from GB
followed by its trailing partial dislocation (Pt), cr is the critical athermal radius of the
partial dislocation. The spacing between the two partial dislocations ( ) depends on the
stacking fault energy.
Fig. 5 Work hardening rate ( ) of nc Ni-20Fe alloy and nc Ni plotted versus true strain
deformed at 1.35 s−1
and RT. The inset is corresponding true stress-strain curves.
Fig. 6 Fracture morphologies of (a) nc Ni-20Fe alloy and (b) nc Ni after deformed at 1.35 s−1
and
RT.
Fig. 7 Typical TEM images of (a) nc Ni-Fe alloy and (b) nc Ni after deformed at 1.35 s−1
and RT.
Fig. 8 (a) A typical high-resolution TEM image of nc Ni-20Fe alloy after deformed at = 1.35s-1
with a Fourier transformation of the white frame at the down left corner; (b) The inverse
Fourier-filtered image from inside the white frame in (a) shows some microtwins, stacking
faults (the white plotlines) and dislocations (the white T).
