metals
Article
Hard Copper with Good Electrical Conductivity
Fabricated by Accumulative Roll-Bonding to
Ultrahigh Strains
Gongcheng Yao 1 , Qingsong Mei 1, *, Juying Li 2 , Congling Li 1 , Ye Ma 1 , Feng Chen 1 ,
Guodong Zhang 1 and Bing Yang 1
1
2
*
School of Power and Mechanical Engineering, Wuhan University, Wuhan 430072, China;
yaogch@gmail.com (G.Y.); conglingli2011@whu.edu.cn (C.L.); mayer@whu.edu.cn (Y.M.);
fengc20@whu.edu.cn (F.C.); gdzhang@whu.edu.cn (G.Z.); toyangbing@163.com (B.Y.)
School of Mechanical Engineering, Wuhan Polytechnic University, Wuhan 430023, China; jylimei@163.com
Correspondence: qsmei@whu.edu.cn; Tel.: +86-27-68772252
Academic Editor: Hugo F. Lopez
Received: 1 April 2016; Accepted: 11 May 2016; Published: 17 May 2016
Abstract: By modifying the accumulative roll-bonding (ARB) procedures, accumulative roll-bonding
(ARB) processing up to 30 cycles (N) with a 50% thickness reduction per cycle (equivalent strain = 24)
at room temperature was conducted on pure copper. The bonding condition, microhardness and
electrical conductivity of the ARBed Cu were studied. Results showed that good bonding condition
of the samples was achieved. As N increases, the microhardness of ARBed Cu increases, reaching
~2.9 times that of annealed Cu for N = 30. The electrical conductivity of ARBed Cu decreases slightly
but with periodic fluctuations for N > 10, with a minimum of 90.4% IACS for N = 30. Our study
indicated that ARB can be an effective way to produce high-hardness and high-conductivity pure
copper better than or comparable to Cu alloys and Cu based composites as reported.
Keywords: accumulative roll-bonding; pure copper; microhardness; electrical conductivity
1. Introduction
High-strength and high-conductivity materials are highly needed in industries for manufacturing
lead frames, electric resistance welding electrodes, contact wires, etc., among which copper alloys are
the most widely used [1–3]. While element alloying can increase the strength of copper effectively,
great sacrifices in electrical conductivity are also induced [4].
Accumulative roll-bonding (ARB) has been an effective method to induce significant
microstructural refinement and material strengthening by deforming metals and alloys to very large
plastic strains [5]. It was reported that the hardness of tough pitch copper (99.9%) processed by
eight cycles (N) of ARB was 2.7 times that before ARB [6]. The reduction in electrical conductivity
of ARBed Cu-0.02%P and Cu-0.15%Fe-0.02%P with an equivalent strain of 8.9 was less than 10%
IACS (International Annealed Copper Standard) [7]. However, in previous studies, the ARB process
applied to pure copper mostly had less than 10 cycles (equivalent strain < 9) because of the initiation
of large cracks in samples at large strains [8,9]. The properties of Cu subjected to ultrahigh strains
remain therefore unclear based on previous studies. Therefore, it is of great interest to investigate the
properties of pure copper deformed by ARB to ultrahigh strains.
In this study, we successfully conducted up to 30 cycles of ARB with a 50% thickness reduction
per cycle (equivalent strain = 24, far beyond those in previous studies) on pure copper by modified
ARB procedures. The bonding condition, microhardness and electrical conductivity of the ARBed
Cu were investigated. The importance of this study is to validate the capability of the modified ARB
processing of Cu to vary high N, and to check the variation of properties of ARBed Cu with N up to 30.
Metals 2016, 6, 115; doi:10.3390/met6050115
www.mdpi.com/journal/metals
Metals 2016, 6, 115 2 of 7 Cu were investigated. The importance of this study is to validate the capability of the modified ARB Metals 2016, 6, 115
2 of 7
processing of Cu to vary high N, and to check the variation of properties of ARBed Cu with N up to 30. 2. Experimental Procedure
2. Experimental Procedure Pure Cu (99.97%) sheets (Luoyang Copper Co., Ltd., Luoyang, China) with 1 mm in thickness were
Ltd., Luoyang, China) with 1 mm in thickness usedPure Cu (99.97%) sheets (Luoyang Copper Co., in this study. As-received sheets were annealed in
Ar atmosphere at 600 ˝ C for 2 h before being
were used in this study. As‐received sheets were annealed in Ar atmosphere at 600 °C for 2 h before cut into strips with dimensions of 200 mmL ˆ 10 mmW ˆ 1 mmT . The schematic of the modified ARB
T. The schematic of the modified being cut into strips with dimensions of 200 mm
× 1 mmof
process is shown in Figure 1. As shown in FigureL × 10 mm
1, the topWsurface
the Cu strip was wire-brushed
ARB process is in Unlike
Figure prior
1. As shown in Figure 1, the top the Cu strip was and degreased
byshown acetone.
ARB
procedures
in literature,
the surface Cu stripof was
folded
in half
by
wire‐brushed and degreased by acetone. Unlike prior ARB procedures in literature, the Cu strip was pliers instead of being cut in half, and then sandwiched in a stainless steel envelope. The envelope was
folded in half by pliers instead of being cut in half, and then sandwiched in a stainless steel envelope. made by folding a stainless steel plate with dimensions of 400 mmL ˆ 100 mmW ˆ 0.5 mmT in half and
L × 100 mmW × The envelope was made by folding a stainless steel plate with dimensions of 400 mm
then rolling it with no thickness reduction. The sample in the envelope was then roll-bonded
with a
T in half and then rolling it with no thickness reduction. The sample in the envelope was then 0.5 mm
50% reduction in thickness at room temperature without lubrication, at a rolling speed of 187 mm/min
roll‐bonded with a 50% reduction in thickness at room temperature without lubrication, at a rolling using a two-high rolling mill. The above procedures were repeated to the intended N and ARBed Cu
speed of with
187 mm/min a two‐high rolling mill. The above procedures were on
repeated to was
the samples
N = 1–30 using were fabricated
respectively.
The
equivalent
strain imposed
samples
intended N and ARBed Cu samples with N = 1–30 were fabricated respectively. The equivalent strain calculated as approximately 0.8 N [8].
imposed on samples was calculated as approximately 0.8 N [8]. Figure 1. Schematic of the modified ARB process. Figure 1. Schematic of the modified ARB process.
To investigate the bonding condition of the ARBed samples, the rolling direction‐normal To investigate the bonding condition of the ARBed samples, the rolling direction-normal direction
direction (RD‐ND) plane of each sample was observed. Microhardness of the samples was measured (RD-ND) plane of each sample was observed. Microhardness of the samples was measured on the
on the rolling direction‐transverse direction (RD‐TD) plane of each sample with a load of 100 gf and rolling direction-transverse direction (RD-TD) plane of each sample with a load of 100 gf and a duration
a duration of 10 s using a HXS‐1000A microhardness tester (Shangguang Microscope Co., Ltd., of 10 s using a HXS-1000A microhardness tester (Shangguang Microscope Co., Ltd., Shanghai, China).
Shanghai, China). Each data represents an average of ten randomly selected points. Electrical Each data represents an average of ten randomly selected points. Electrical conductivity was measured
conductivity was measured on the RD‐TD plane by an eddy‐current electrical conductivity meter on the RD-TD plane by an eddy-current electrical conductivity meter (FQR7501A, Xingsha Instrument
(FQR7501A, Xingsha Instrument Co., Ltd., Xiamen, China). Each piece of data represents an average Co., Ltd., Xiamen, China). Each piece of data represents an average of three measurements.
of three measurements. 3. Results and Discussion
3. Results and Discussion In the conventional ARB process, the roll-bonded sample is cut in half, stacked after surface
In the and
conventional process, the is cut and
in half, stacked surface treatment,
then rolledARB [5]. However,
theroll‐bonded independentsample upper layer
lower
layer ofafter the stacked
treatment, and then rolled [5]. However, the independent upper layer and lower layer of the stacked sample can move freely during rolling, which will make the two layers diverge. The overlapping
sample freely during will make the two deformations,
layers diverge. The overlapping part andcan move non-overlapping
part ofrolling, the two which layers thus
have different
which
facilitate the
part and non‐overlapping part of the two layers thus have different deformations, which facilitate initiation and propagation of cracks in the sample, leading to the failure of the sample. To overcome this
the initiation and the
propagation of or
cracks in the sample, leading to was
the used
failure of the studies
sample. To problem,
fastening
four corners
two ends
of the
sample
by wires
in some
[8,10].
overcome this problem, fastening the four corners or two ends of the sample by wires was used in The maximum cycle that can be performed was still limited as the large force during rolling could
some studies [8,10]. The maximum cycle that can be performed was still limited as the large force easily overcome the fastening. In this study, the Cu strip was folded as a whole for rolling. The folding
during rolling could easily overcome the fastening. In this study, the Cu strip was folded as a whole end of the sample greatly limited the relative movement of the upper and lower layers of the sample
during rolling. The stainless steel envelope further fixed the sample and homogenized the force during
Metals 2016, 6, 115
3 of 7
Metals 2016, 6, 115 3 of 7 rolling. As a result, the crack of samples was inhibited and up to 30 cycles of ARB with a 50% thickness
for rolling. The folding end of the sample greatly limited the relative movement of the upper and reduction
per cycle (equivalent strain = 24) were successfully conducted on a copper sample at room
lower layers of the sample during rolling. The stainless steel envelope further fixed the sample and temperature
without intermediate annealing.
homogenized the force during rolling. As a result, the crack of samples was inhibited and up to 30 Microstructures
of the
RD-NDreduction planes of
samples
with various
N= (N
1, 5,successfully 10, 20 and 30) and
cycles of ARB with a 50% thickness per cycle (equivalent strain 24) =were conducted on a copper sample at room temperature without intermediate annealing. annealed
Cu before ARB (N = 0) are shown in Figure 2. In contrast to the clear grain boundaries
Microstructures of the RD‐ND planes of samples with various N (N = 1, 5, 10, 20 and 30) and in Figure
2a, it is difficult to recognize the individual grains in Figure 2b–f as the grains have been
annealed Cu before ARB (N = 0) are shown in Figure 2. In contrast to the clear grain boundaries in severely deformed and significantly refined by ARB [5]. According to previous studies, the grains
2a, Cu
it is have
difficult to significantly
recognize the individual grains Figure 2b–f the grains have been nm for
ofFigure ARBed
been
refined, with
anin average
grainas size
of ~260
nm–300
severely deformed and significantly refined by ARB [5]. According to previous studies, the grains of N = 6–8 [5,11]. In this study, the optical micrographs are used to illustrate the bonding condition of
ARBed Cu have been significantly refined, with an average grain size of ~260 nm–300 nm for N = 6–
ARBed
samples on a relatively large scale, which is crucial for bulk structural materials. For ARBed
8 [5,11]. In this study, the optical micrographs are used to illustrate the bonding condition of ARBed samples,
the theoretical number of interfaces between Cu layers is 2N ´ 1 (for N = 30, the theoretical
samples on a relatively large scale, which is crucial for bulk structural materials. For ARBed samples, N − 1 (for N = 30, the theoretical number of number
of interfaces is as large as 109 ). In Figure 2, no interfaces
between the un-bonded layers can be
the theoretical number of interfaces between Cu layers is 2
9
seen,
indicating
the
excellent
bonding
condition
of
the
ARBed
Cu
samples.
interfaces is as large as 10 ). In Figure 2, no interfaces between the un‐bonded layers can be seen, indicating the excellent bonding condition of the ARBed Cu samples. (a)
(b)
(d)
(e)
(c)
(f)
Figure 2. Micrographs of (a) annealed Cu before ARB, RD‐ND planes of ARBed Cu samples with N = Figure 2. Micrographs of (a) annealed Cu before ARB, RD-ND planes of ARBed Cu samples with
(b) 1, (c) 5, (d) 10, (e) 20 and (f) 30.
N = (b) 1, (c) 5, (d) 10, (e) 20 and (f) 30.
The microhardness of the ARBed Cu as functions of both N and corresponding equivalent strain is shown in Figure 3. As shown Figure 3, the microhardness of ARBed Cu increases with The microhardness of the
ARBedin Cu
as functions
of both N and corresponding equivalent strain
but 3.
with a reduced rate accompanied by fluctuations from = 5 to N = 30. isincreasing shown in N, Figure
As shown
in Figure
3, the microhardness
of ARBed
CuN increases
with
increasing
Microhardness of the ARBed Cu sample with N = 30 is 140.1 HV, ~2.9 times that of initial annealed N, but with a reduced rate accompanied by fluctuations from N = 5 to N = 30. Microhardness of the
Cu. The evident increase of the microhardness of ARBed Cu mainly results from the strengthening ARBed
Cu sample with N = 30 is 140.1 HV, ~2.9 times that of initial annealed Cu. The evident increase
effect due to accumulation of lattice defects and significant grain refinement at ultrahigh strains of the microhardness of ARBed Cu mainly results from the strengthening effect due to accumulation
[6,12,13]. The reduced rate of the increase of microhardness of ARBed Cu with increasing strain can ofbe lattice
defects and significant grain refinement at ultrahigh strains [6,12,13]. The reduced rate of the
attributed to the balance of the strengthening effect of strain hardening, fine grain boundary increase
of
microhardness of ARBed Cu with increasing strain can be attributed to the balance of the
hardening and the softening effect of dynamic recovery [7,14,15]. strengthening effect of strain hardening, fine grain boundary hardening and the softening effect of
dynamic recovery [7,14,15].
Variation in microhardness of tough pitch copper (99.9%) processed by eight-cycle ARB reported
by Shaarbaf et al. [6] is also shown in Figure 3 for comparison. The normalized hardness (ratio of the
final microhardness to initial microhardness) of ARBed tough pitch copper with N = 8 is 2.7. The same
ratio is obtained for ARBed Cu with N = 5 in the present study, and the ratio for ARBed Cu with N = 30
is 2.9. We found that the most evident increase in the microhardness of ARBed Cu occurred for N ď 5
in present study. Further deformation by ARB to N = 30 only brought slight increase in microhardness,
which clearly indicates that the strengthening effect on pure copper induced by ARB has an upper
Metals 2016, 6, 115
4 of 7
limit. Meanwhile, the continuous increase of microhardness up to N = 30, although with much less
evidently for N = 5–30, indicates the good bonding condition and intactness of the ARBed Cu samples
with very high N since possible debonding and cracking in samples can induce significant decrease of
hardness
and electrical conductivity when N is very large.
Metals 2016, 6, 115 4 of 7 Equivalent Strain
0
4
8
12
16
20
24
4.0
140
3.5
120
3.0
100
2.5
Measured data of pure copper
Measured data of tought pitch copper [6]
Normalized data of pure copper
Normalized data of touch pitch copper [6]
80
60
40
0
5
10
15
N
20
25
30
2.0
1.5
Ratio after Normalization
Microhardness (HV)
160
1.0
Figure 3. Microhardness of ARBed Cu in the present study and ARBed tough pitch copper [6] as Figure 3. Microhardness of ARBed Cu in the present study and ARBed tough pitch copper [6] as
functions of N and corresponding equivalent strain, before and after normalization. functions
of N and corresponding equivalent strain, before and after normalization.
Variation in microhardness of tough pitch copper (99.9%) processed by eight‐cycle ARB Ductility is critical, especially for metals prepared by high-strain deformation. As for the ductility
reported by Shaarbaf et al. [6] is also shown in Figure 3 for comparison. The normalized hardness of ARBed Cu, Shaarbaf et al. [6] reported that the elongation of ARBed tough pitch copper (99.9%)
(ratio of the final microhardness to initial microhardness) of ARBed tough pitch copper with N = 8 is decreases from ~50% to ~2% in the 1st cycle and then increases up to ~6% with N = 8. Jang et al. [11]
2.7. The same ratio is obtained for ARBed Cu with N = 5 in the present study, and the ratio for ARBed reported
a similar phenomenon that the elongation of ARBed Cu decreases greatly in the 1st cycle and
Cu with N = 30 is 2.9. We found that the most evident increase in the microhardness of ARBed Cu then
shows
a gradual
following
cycles,
suggestingby that
the to strain
alone can
not
occurred for N ≤ 5 in increase
present instudy. Further deformation ARB N = hardening
30 only brought slight explain
this
phenomenon.
The
effect
of
ultrahigh
strains
on
the
ductility
of
ARBed
Cu
is
two-fold,
i.e.,
increase in microhardness, which clearly indicates that the strengthening effect on pure copper it
can induce loss in ductility due to strain hardening and induce increase in ductility due to dynamic
induced by ARB has an upper limit. Meanwhile, the continuous increase of microhardness up to N = recovery.
However,
further
study
is needed
point. On
other
hand, condition for applicants
30, although with much less evidently for to
N clarify
= 5–30, this
indicates the the
good bonding and that
require
a
combination
of
high
microhardness
and
electrical
conductivity,
e.g.,
lead
frames
and
intactness of the ARBed Cu samples with very high N since possible debonding and cracking in welding
electrodes, the ductility may not be a main requirement. For example, the elongations of
samples can induce significant decrease of hardness and electrical conductivity when N is very large. C19400,
NK-120
C70250
used for
lead
frames
are 4%–5%,
5% and 2%–6%
respectively
[1,16].
Ductility is and
critical, especially for metals prepared by high‐strain deformation. As for the Thus,
the
ARBed
Cu
in
the
present
study
can
be
a
potential
candidate
for
these
applications
despite
ductility of ARBed Cu, Shaarbaf et al. [6] reported that the elongation of ARBed tough pitch copper the
reduced ductility.
(99.9%) decreases from ~50% to ~2% in the 1st cycle and then increases up to ~6% with N = 8. Jang et Figure 4 illustrates the electrical conductivity (σ) of the ARBed Cu versus N as well as
al. [11] reported a similar phenomenon that the elongation of ARBed Cu decreases greatly in the 1st corresponding
strain. increase As shown
in Figure cycles, suggesting 4 σ decreases onlythat slightly
with increasing
cycle and then equivalent
shows a gradual in following the strain hardening N.
For
N
=
30,
σ
of
the
ARBed
Cu
is
still
90.4%
IACS,
only
8.7%
IACS
smaller
than
initial
annealed Cu
alone can not explain this phenomenon. The effect of ultrahigh strains on the ductility of ARBed Cu (N
= 0). Most i.e., interestingly,
we observed
periodic
in the σ and of ARBed
Cuincrease for N > 10.
is two‐fold, it can induce loss in evident
ductility due to fluctuations
strain hardening induce in Such
fluctuations
are believed
to be associated
with
alternating
accumulation
and annihilation
of lattice
ductility due to dynamic recovery. However, further study is needed to clarify this point. On the defects.
Whenfor Cu is
deformed by
ARB
to higha strains,
accumulation
of lattice
defects (e.g.,
dislocations)
other hand, applicants that require combination of high microhardness and electrical can
induce
a
decrease
in
the
electrical
conductivity.
Meanwhile,
dynamic
recovery
can
induce
an
conductivity, e.g., lead frames and welding electrodes, the ductility may not be a main requirement. increase
in the electrical conductivity due to annihilation of defects. Usually, dynamic recovery is
For example, the elongations of C19400, NK‐120 and C70250 used for lead frames are 4%–5%, 5% driven
by
the respectively increase of strain
energy
to the a critical
level.
dynamic
recovery,
the strain
in
and 2%–6% [1,16]. Thus, ARBed Cu After
in the present study can be a energy
potential the
sample is released and then further accumulation of lattice defects is possible. This may lead to the
candidate for these applications despite the reduced ductility. alternating
and increase
of lattice
defects during
sample,
which
will
Figure decrease
4 illustrates the electrical conductivity (σ) the
of deformation
the ARBed of
Cu the
versus N as well as corresponding equivalent strain. As shown in Figure 4 σ decreases only slightly with increasing N. For N = 30, σ of the ARBed Cu is still 90.4% IACS, only 8.7% IACS smaller than initial annealed Cu (N = 0). Most interestingly, we observed evident periodic fluctuations in the σ of ARBed Cu for N > 10. Such fluctuations are believed to be associated with alternating accumulation and annihilation of lattice defects. When Cu is deformed by ARB to high strains, accumulation of lattice defects (e.g., Metals 2016, 6, 115 5 of 7 Metalsrecovery 2016, 6, 115
is driven by the increase of strain energy to a critical level. After dynamic recovery, the 5 of 7
conductivity
Electrical Electrical
conductivity
(% IACS)(% IACS)
strain energy in the sample is released and then further accumulation of lattice defects is possible. Metals 2016, 6, 115 5 of 7 This may lead to the alternating decrease and increase of lattice defects during the deformation of induce
the
alternating
and
increase
ofenergy the
electrical
conductivity.
Nevertheless,
further
study,
recovery is driven decrease
by the increase of strain to a critical level. of After dynamic conductivity. recovery, the the sample, which will induce the alternating decrease and increase the electrical e.g., aNevertheless, further study, e.g., a direct characterization of the lattice defects is needed to clarify direct
characterization
of
the
lattice
defects
is
needed
to
clarify
this
point.
strain energy in the sample is released and then further accumulation of lattice defects is possible. This may lead to the alternating decrease and increase of lattice defects during the deformation of this point. Equivalent
strain
the sample, which will induce the alternating decrease and increase of the electrical conductivity. Nevertheless, further study, e.g., a direct characterization of the lattice defects is needed to clarify 0
4
8
12
16
20
24
100
this point. Equivalent strain
98
0
4
8
12
16
20
24
100
96
98
94
96
92
94
90
92
88
90
0
5
10
15
N
20
25
30
Figure 4. Electrical conductivity of ARBed Cu versus N and equivalent strain. 88
Figure 4. Electrical conductivity
of10
ARBed15Cu versus
strain.
0
5
20 N and
25 equivalent
30
N
The above results show that the ARBed Cu has a good combination of hardness and electrical conductivity
Electrical Electrical
conductivity
(% IACS)(% IACS)
conductivity. Figure 5 presents the electrical conductivity‐strength relationships of the ARBed Cu The
above results
show
that the
ARBed Cu
has a good combination
of hardness
and electrical
Figure 4. Electrical conductivity of ARBed Cu versus N and equivalent strain. samples in the present study (N = 5, 10, 12, 20, 27, 30), and high‐strength, high‐conductivity Cu alloys conductivity. Figure 5 presents the electrical conductivity-strength relationships of the ARBed Cu
The above results show that the ARBed Cu has a good combination of hardness and electrical [1,16,17] for lead (including Cu/GNPs (graphene Cu
samples
in
theused present
studyframes (N = 5,and 10, Cu 12, based 20, 27, composites 30), and high-strength,
high-conductivity
conductivity. Figure 5 presents the electrical conductivity‐strength relationships of the ARBed Cu nanoplatelets) [18], Cu/CNTs (carbon nanotubes) [19], Cu/Al
2O3 [20], Cu/SiC [21], Cu/TiB
2 [22]) are alloys [1,16,17] used for lead frames and Cu based composites (including Cu/GNPs (graphene
samples in the present study (N = 5, 10, 12, 20, 27, 30), and high‐strength, high‐conductivity Cu alloys 2O3 and Cu/SiC was also included for comparison. The strength of the ARBed Cu, Cu/GNPs, Cu/Al
nanoplatelets) [18], Cu/CNTs (carbon nanotubes) [19], Cu/Al2 O3 [20], Cu/SiC [21], Cu/TiB2 [22])
[1,16,17] used for lead frames and Cu based composites (including Cu/GNPs (graphene calculated as H/3 [23,24], where H is the microhardness of the material. As shown in Figure 5, the are also
included
for comparison.
Thecombination strength ofof the
ARBedand Cu,electrical Cu/GNPs,
Cu/Al2 Othan and
Cu/SiC
32 [22]) are nanoplatelets) [18], Cu/CNTs (carbon nanotubes) [19], Cu/Al
2O3 [20], Cu/SiC [21], Cu/TiB
ARBed Cu samples have a better strength conductivity most was calculated
as H/3
[23,24], where
is the Cu/Al
microhardness
of the material.2O
As
shownto inCu/C Figure 5,
3 and Cu/SiC was also included for comparison. The strength of the ARBed Cu, Cu/GNPs, Cu/Al
high‐strength, high‐conductivity Cu Halloys, 2O3, Cu/SiC, Cu/TiB2, and are similar the ARBed
Cu samples have a better combination of strength and electrical conductivity than most
calculated as H/3 [23,24], where H is the microhardness of the material. As shown in Figure 5, the composites (including Cu/GNPs and Cu/CNTs). Obviously, our results indicate the effectiveness of ARBed Cu high-conductivity
samples have a better of 2strength and electrical than most high-strength,
Cu combination alloys, Cu/Al
O3 , Cu/SiC,
Cu/TiB2conductivity , and are similar
to Cu/C
ARB for producing high‐strength and high‐electrical‐conductivity Cu. high‐strength, high‐conductivity Cu alloys, Cu/Al
2
O
3
, Cu/SiC, Cu/TiB
2
, and are similar to Cu/C composites (including Cu/GNPs and Cu/CNTs). Obviously, our results indicate the effectiveness of
100
Present study *
composites (including Cu/GNPs and Cu/CNTs). Obviously, our results indicate the effectiveness of ARB for
producing high-strength and high-electrical-conductivity
Cu.
Cu-Fe
[16,17]
ARB for producing high‐strength and high‐electrical‐conductivity Cu. Cu-Ni-Si [1,16]
90
Cu/C
100
80
90
70
Cu alloys
Cu/C
80
60
70
50
Cu alloys
Cu/SiC
60
40
50
30
Cu/TiB2
Cu/SiC
Cu-Sn [1,16]
Cu/GNPs
[18]**
Present study
Cu/CNTs
[19]
Cu-Fe [16,17]
Cu/Al
O [20] *
Cu-Ni-Si
2 3 [1,16]
Cu-Sn [1,16]
Cu/SiC
[21] *
Cu/GNPs
[18] *
Cu/TiB
[22]
2
Cu/CNTs [19]
Cu/Al2O3 [20] *
Cu/Al2O3
Cu/SiC [21] *
Cu/TiB2 [22]
Cu/Al2O3
40
20
250 300 350 400 450 500 550 600 650 700 750 800
30
Tensile
strength (MPa)
Cu/TiB
2
20
250 300 350 400 450 500 550 600 650 700 750 800
Tensile strength (MPa)
Figure 5. Electrical conductivity versus tensile strength of ARBed Cu (N = 5, 10, 12, 20, 27, 30) in this
study, high-strength, high-conductivity Cu alloys [1,16,17] and common Cu based composites [18–22].
* tensile strength is calculated as H/3 [23,24].
Metals 2016, 6, 115
6 of 7
4. Conclusions
By modifying the ARB procedures, we successfully conducted ARB processing up to 30 cycles
with a 50% thickness reduction per cycle (equivalent strain = 24) at room temperature on pure copper.
We found that good bonding condition was achieved for the ARBed samples. The microhardness
of ARBed Cu samples increases rapidly at first and then is almost saturated from the 5th cycle to
the 30th cycle, ~2.9 times that of the annealed pure copper for N = 30. The electrical conductivity
decreases with increasing N, all larger than 90% IACS for various N. The variations of microhardness
and electrical conductivity with ARB strains can be understood considering the strengthening effects
due to accumulation of lattice defects and significant grain refinement and the softening effect due
to dynamic recovery. The ARBed Cu has a better combination of strength and electrical conductivity
than common Cu alloys, and is even comparable to Cu matrix composites reinforced by GNPs or
CNTs. Our study indicates the effectiveness of ARB processing for fabrication of high-strength and
high-conductivity pure copper.
Acknowledgments: This work was financially supported by the National Natural Science Foundation of China
(grant 51371128) and the research foundation of Wuhan University.
Author Contributions: Q.S.M. and J.L. conceived and designed the experiments; G.Y., C.L., Y.M. and F.C.
performed the experiments; G.Y. and Q.S.M. analyzed the data; G.Z. and B.Y. contributed analysis tools and
discussion of results; G.Y. and Q.S.M. wrote the paper.
Conflicts of Interest: The authors declare no conflicts of interest.
References
1.
2.
3.
4.
5.
6.
7.
8.
9.
10.
11.
12.
13.
14.
Tomioka, Y.; Miyake, J. A copper alloy development for leadframe. In Proceedings of 1995 Electronic
Manufacturing Technology Symposium, Omiya, Japan, 4–6 December 1995; pp. 433–436.
Suzuki, S.; Shibutani, N.; Mimura, K.; Isshiki, M.; Waseda, Y. Improvement in strength and electrical
conductivity of Cu–Ni–Si alloys by aging and cold rolling. J. Alloy. Compd. 2006, 417, 116–120. [CrossRef]
Lu, D.P.; Wang, J.; Zeng, W.J.; Liu, Y.; Lu, L.; Sun, B.D. Study on high-strength and high-conductivity Cu-Fe-P
alloys. Mater. Sci. Eng. A 2006, 421, 254–259. [CrossRef]
Ghosh, G.; Miyake, J.; Fine, M.E. The systems-based design of high-strength, high-conductivity alloys. JOM
1997, 49, 56–60. [CrossRef]
Tsuji, N.; Saito, Y.; Lee, S.H.; Minamino, Y. ARB (Accumulative Roll-Bonding) and other new Techniques to
Produce Bulk Ultrafine Grained Materials. Adv. Eng. Mater. 2003, 5, 338–344. [CrossRef]
Shaarbaf, M.; Toroghinejad, M.R. Nano-grained copper strip produced by accumulative roll bonding process.
Mater. Sci. Eng. A 2008, 473, 28–33. [CrossRef]
Jang, Y.; Kim, S.; Han, S.; Lim, C.; Goto, M. Tensile behavior of commercially pure copper sheet fabricated by
2- and 3-layered accumulative roll bonding (ARB) process. Met. Mater. Int. 2008, 14, 171–175. [CrossRef]
Saito, Y.; Utsunomiya, H.; Tsuji, N.; Sakai, T. Novel ultra-high straining process for bulk
materials—Development of the accumulative roll-bonding (ARB) process. Acta Mater. 1999, 47, 579–583.
[CrossRef]
Kunimine, T.; Fujii, T.; Onaka, S.; Tsuji, N.; Kato, M. Effects of Si addition on mechanical properties of copper
severely deformed by accumulative roll-bonding. J. Mater. Sci. 2011, 46, 4290–4295. [CrossRef]
Suresh, K.S.; Sinha, S.; Chaudhary, A.; Suwas, S. Development of microstructure and texture in Copper
during warm accumulative roll bonding. Mater. Charact. 2012, 70, 74–82. [CrossRef]
Jang, Y.H.; Kim, S.S.; Han, S.Z.; Lim, C.Y.; Kim, C.J.; Goto, M. Effect of trace phosphorous on tensile behavior
of accumulative roll bonded oxygen-free copper. Scr. Mater. 2005, 52, 21–24. [CrossRef]
Jamaati, R.; Toroghinejad, M.R. Application of ARB process for manufacturing high-strength, finely dispersed
and highly uniform Cu/Al2 O3 composite. Mater. Sci. Eng. A 2010, 527, 7430–7435. [CrossRef]
Xing, Z.P.; Kang, S.B.; Kim, H.W. Structure and properties of AA3003 alloy produced by accumulative roll
bonding process. J. Mater. Sci. 2002, 37, 717–722. [CrossRef]
Eizadjou, M.; Manesh, H.D.; Janghorban, K. Microstructure and mechanical properties of ultra-fine grains
(UFGs) aluminum strips produced by ARB process. J. Alloy. Compd. 2009, 474, 406–415. [CrossRef]
Metals 2016, 6, 115
15.
16.
17.
18.
19.
20.
21.
22.
23.
24.
7 of 7
Shih, M.H.; Yu, C.Y.; Kao, P.W.; Chang, C.P. Microstructure and flow stress of copper deformed to large
plastic strains. Scr. Mater. 2001, 45, 793–799. [CrossRef]
Ma, J.S.; Huang, F.X.; Huang, L.; Geng, Z.T.; Ning, H.L.; Han, Z.Y. Trends and development of copper alloys
for lead frame. J. Funct. Mater. 2002, 33, 1–4.
ASM handbook committee. Metals Handbook Volume 2, 9th ed.; ASM: Metals Park, OH, USA, 1979;
pp. 310–313.
Chen, F.; Ying, J.; Wang, Y.; Du, S.; Liu, Z.; Huang, Q. Effects of graphene content on the microstructure and
properties of copper matrix composites. Carbon 2016, 96, 836–842. [CrossRef]
Arnaud, C.; Lecouturier, F.; Mesguich, D.; Ferreira, N.; Chevallier, G.; Estournès, C.; Weibel, A.; Laurent, C.
High strength–High conductivity double-walled carbon nanotube–Copper composite wires. Carbon 2016, 96,
212–215. [CrossRef]
Rajkovic, V.; Bozic, D.; Stasic, J.; Wang, H.; Jovanovic, M.T. Processing, characterization and properties
of copper-based composites strengthened by low amount of alumina particles. Powder Technol. 2014, 268,
392–400. [CrossRef]
Efe, G.C.; Yener, T.; Altinsoy, I.; Ipek, M.; Zeytin, S.; Bindal, C. The effect of sintering temperature on some
properties of Cu-SiC composite. J. Alloy. Compd. 2011, 509, 6036–6042.
López, M.; Corredor, D.; Camurri, C.; Vergara, V.; Jiménez, J. Performance and characterization of dispersion
strengthened Cu-TiB2 composite for electrical use. Mater. Charact. 2005, 55, 252–262. [CrossRef]
Zhang, P.; Li, S.X.; Zhang, Z.F. General relationship between strength and hardness. Mater. Sci. Eng. A 2011,
529, 62–73. [CrossRef]
Cahoon, J.R.; Broughton, W.H.; Kutzak, A.R. The determination of yield strength from hardness
measurements. Metall. Trans. 1971, 2, 1979–1983.
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