Thermal aging behavior in Cu–Al–Ni–xCo shape memory alloys
Safaa N. Saud • E. Hamzah • T. Abubakar •
H. R. Bakhsheshi-Rad
Received: 29 July 2014 / Accepted: 21 October 2014 / Published online: 7 November 2014
� Akademiai Kiado, Budapest, Hungary 2014
Abstract The effects of aging on the phase transforma-
tion temperatures, morphology, and mechanical properties
were investigated in the Cu–Al–Ni–xCo (x = 0.4, 0.7, and
1.0 mass%) SMAs via differential scanning calorimetry,
field emission scanning electron microscopy, X-ray dif-
fraction, tensile, vickers microhardness, and shape memory
effect test. The results show that the phase morphology and
the presence of c2 precipitates in different sizes, volume
fractions, and distributions have an obvious effect on the
phase transformation characteristics and properties. In
addition, a noticeable variation in the transformation tem-
peratures and thermodynamic parameters were occurred
with the aging. In tension, the highest tensile strength and
tensile strain of 750 MPa and 5.5 % were indicated in the
Cu–Al–Ni–1 mass% Co alloy after being aged at 523 K for
24 h and 373 K for 48 h, respectively. However, the results
of the strain recovery by shape memory effect were varied
in accordance with the variation of the c01 and b01 mor-
phology and volume fraction of c2 precipitates. In aged
alloy of 523 K for 48 h, the thickness ofc01 and b01 phases
and the volume fraction of c2 precipitates increase, thus,
the movement of martensitic interfaces is restricted causing
an increase in et by SME.
Keywords Shape memory alloys � Cu–Al–Ni–Co � Aging
treatment � Shape memory effect
Introduction
Shape memory alloys have a remarkable capability to act
as a sensor and an actuator, which makes them a multi-
functional material with the unique ability to precisely fit
the requirements of a particular usage [1, 2]. Among the
shape memory alloys, NiTi SMAs have been used with the
greatest amount of feedback in a numerous applications
including biomedical and industrial applications [3, 4].
However, the low transformation temperatures (-373 to
373 K) and high producing cost of NiTi alloys render their
usage for many applications impractical, especially for
high temperature usage. Therefore, Cu–Al–Ni SMAs are
the most preferred option as successful alternatives to NiTi
due to their high transformation temperatures (-573 to
573 K), low producing cost, and high thermal stability
[5–8]. However, the brittleness [9], low strength, large
elastic anisotropy, and large grain size hinder its practical
applications. As a result, many researchers have tried to
refine the grain size of Cu-based SMAs through the addi-
tion of alloying elements and/or applying different thermal
aging treatment conditions [10–18]. The unique properties
of Cu–Al–Ni SMA are mainly attributed to the thermo-
elastic martensitic transformation that takes place in a
temperature range of (173–473 K) depending on the
composition of alloy. These properties are significantly
affected by the movements of interfaces (twin boundaries,
martensitic variants, and parent/martensite phase bound-
aries) [19]. Some other parameters also exhibited a sig-
nificant influence on the properties of these alloys, for
instance, the formation of vacancies, dislocations, grain
boundaries, and precipitates, along with the way of their
distribution and volume fraction. On the other hand, the
Cu–Al–Ni alloys are liable to aging treatment [20, 21],
which it leads to form a different types, volume fractions,
S. N. Saud � E. Hamzah (&) � T. Abubakar �H. R. Bakhsheshi-Rad
Faculty of Mechanical Engineering, Universiti Teknologi
Malaysia UTM, 81310, Johor Bahru, Johor, Malaysia
e-mail: [email protected]
123
J Therm Anal Calorim (2015) 119:1273–1284
DOI 10.1007/s10973-014-4265-6
and distributions of precipitates in the microstructure,
associated with the successive of martensitic transforma-
tion from b01 to c01 phase [22]. However, it is well estab-
lished that different types of martensitic phases, c01 (2H), b01(18R), and a01 (6R) form in the Cu–Al–Ni SMAs,
depending on the chemical composition, mode of applied
load, test temperature, and crystal orientation [23–25].
Nevertheless, Balo’s research on Cu–Al–Ni alloys dis-
covered that an appropriate aging process may significantly
vary the transformation temperatures [14]. Furthermore, a
proper aging process may produce a favorable combination
of features, such as high strength and a large strain
recovery [11, 26, 27]. As these alloys are susceptible to
post-quench aging at high temperature service conditions,
their transformation temperatures, martensitic phases, and
the mechanical properties can change with the time of
treatment. As a consequence, a large number of studies
were conducted on various aspects of aging in these
alloys and their influence on the shape memory properties
[10, 11, 13, 14, 28, 29]. However, only a limited number of
studies that too confined to a small number of aging con-
ditions were conducted on the mechanical response of the
aged Cu–Al–Ni alloys [10, 14, 28, 29]. The present work
was initiated to conduct a comprehensive study on the phase
transformation temperatures, morphology, and mechanical
properties in the Cu (84.1-x)–11.9 Al–4 Ni–xCo (x = 0.4,
0.7, and 1) (in mass%) SMAs after conducting an aging
treatment at 373, 423, and 523 K for 24 and 48 h.
Experimental
Materials preparation
The alloy was produced by melting high purity metals of
Cu (99.999 %), Al (99.999 %), Ni (99.995 %), and Co
(99.95 %) using an induction furnace. These metals were
melted in a silicon carbide crucible at a temperature about
1,573 K with continuous stirring and then poured into a
cast iron mold with dimensions of 270 9 50 9 20 mm3.
Three ingots were produced with different percentages of
Co (0.4, 0.7, and 1.0 mass%) and each ingot was cut into
six samples using Electrical discharge machining wire and
then homogenized at the 1,173 K for 30 min, and then
quenched in water which led to the formation of martensite.
The chemical composition analysis for the Cu–Al–Ni–Co
SMAs was investigated using inductively coupled plasma
mass spectrometry (ICP–MS). The aging treatments were
carried out at 373, 423, and 523 K for 24 and 48 h in the
normal atmosphere. The optimization of the aging times
and temperatures has been selected in accordance with
extensive studies done by the authors.
Material characterization
Flat specimens were cut from the aged samples with
dimensions of 10 9 10 9 2 mm3 for the microstructural
and X-ray diffraction (XRD) characteristics. Filings of the
alloys removed of about 2–6 mg were taken for the differ-
ential scanning calorimetry measurements using a Mettler
Toledo DSC 822e, where the scanning rate was 10 K min-1
in the 323–673 K range. The phase identifications and
crystal structure determinations were carried out using a
D5000 Siemens X-Ray diffractometer fitted with CuKaX-ray source with a locked couple mode, 2h range between
30–80�, and 0.05� sec-1, is the scanning step. The quenched
samples were ground and polished, and then etched in a
solution containing 2.5 g ferric chloride (FeCl3.6H2O) and
48 mL methanol (CH3OH) in 10 mL HCl for 4 min.
Mechanical tests
Tensile test
The tensile test was performed using an Instron 5982-type
universal testing machine operated at a constant strain rate
of 0.1 mm min-1. The tests were carried out at room
temperature until failure occurred, and then the fracture
stress–strain levels were determined under the tensile load.
Shape memory effect test
The shape memory effect test was carried out using a spe-
cially designed machine. The specially designed contents
were analyzed using an Instron 5982-type universal testing
machine operated with special program parameters
according to the shape memory test which was connected to
a heater tape and digital thermocouple in order to control the
applied temperature, and an external extensometer to mea-
sure the shape extension and recovery. The tests were car-
ried out at a temperature below Mf, which was about 373 K,
where the alloys would be able to obtain shape recovery.
Then the deformed sample that still had an unrecoverable
shape was subsequently heated above the austenite finish
temperature (Af ? 333 K) for 10 min followed by a water
quench to recover the residual strain (er). The recovered
shape was attributed to the transformation of the detwinned
martensite to the austenite phase, which had been termed as
a transformation strain (et). After the cooling process, the
martensite again formed in a self-accommodated structure.
1274 S. N. Saud et al.
123
Results and discussion
Microstructural analysis
The typical microstructures of the Cu–Al–Ni–xCo
(x = 0.4, 0.7, and 1.0 mass%) under different aging con-
ditions are shown in Fig. 1a–r. In the aged alloys of
0.4 mass% Co at 373 K for 24 and 48 h, a fine needle-like
and plate-like of b01 and c01 are formed in different orien-
tations, volume fractions, and sizes along with some pre-
cipitates/intermetallic compounds are formed, as indicated
by the red dot arrow in Fig. 1a, j. Perhaps it seems that with
the Co addition, there is certainly an innovative new phase
formed which begin to grow up into the matrix, and
523
423
Agi
ng te
mpe
ratu
re/K
Cobalt concentration/mass%
Agi
ng te
mpe
ratu
re/K
373
0.4 0.7 1.0
Cobalt concentration/mass%0.4
(a) (b) (c)
(d) (e) (f)
(g) (h) (i)
(j) (k) (l)
(m) (n) (o)
(p) (q) (r)
0.7 1.0
γ2
γ2
γ 1
β 1
γ2
γ 1
β 1
523
48 hr
423
373
24 hr
Fig. 1 FESEM micrographs of
the microstructures of Cu–Al–
Ni–xCo SMA (x = 0.4, 0.7, and
1.0 mass%) under different
aging conditions
Thermal aging behavior in Cu–Al–Ni–xCo 1275
123
consequently above the needle-like and plate-like of b01phases, which is typically known as c2 phase. In accor-
dance with an EDS analysis of a spot scanned for the c2
phase area, it was found that these precipitates are Co-rich,
which are an amalgamation of Co, Ni, and Al in compound
of Al75Co22Ni3, as is being pointed out in Fig. 2. With
increasing aging time and temperature to 423 K for 24 and
48 h, the volume fraction and size of c2 phase will increase.
On the other hand, this increment can lead to form and
accumulates the c2 phase at the grain boundaries, thus, the
thickness of the grain boundaries has increased as shown in
Fig. 1d, m. Generally speaking, with an increase in aging
time and temperature, the driving force of the precipitation
hardening is increasing over the subsequent aging [11, 28].
With further increases in the aging time and temperature to
523 K for 24 and 48 h, the c2 precipitates are starting to
form and diffuse into the matrix associated with the
decreases of their sizes as shown in Fig. 1g, p. From
another point of view, it was noticed that with the aging
treatment, the volume fraction of b01 phase increased, which
is consequently associated with the increases of c01 thick-
ness, even though both phases are formed in the matrix and
orientated randomly in accordance to the aging treatment
condition.
At 0.7 mass% Co-aged alloys, the microstructures have
exhibited the same variations in the volume fraction, size,
and morphology of the c01 phase, b01 phase, and c2 precip-
itate as shown in Fig. 1. These variations occur according
to the principles of the aging treatment condition. It can be
seen that within 373 K aging temperature for 24 h, the
volume fraction and size of c01 phase are increased, while
the volume fraction of c2 precipitates is shown much less
and look like they have been penetrated into the micro-
structure as shown in Fig. 1b. With increasing the aging
time to 48 h at 373 K, the volume fraction of b01 phase and
c2 precipitates will be increased, and they have formed
randomly in the matrix as observed in Fig. 1k. However, at
the aging temperature of 423 K and time of 24 h, the c2
precipitates are starting to create a coarse boundary, and
therefore, clear and thick boundaries have occurred along
with a thick and sharp plate-like of c01 phase are formed,
while the remaining area is full of the needle-like of b01phase as shown in Fig. 1e. The c2 precipitates are formed
randomly in the microstructure of the 423 K for 48 h aged
alloy and the volume fraction of b01 phase has increased that
is associated with the decrement of plate thickness of c01phase as shown in Fig. 1n. With a further increase in the
aging time and temperature to 523 K for 24 and 48 h, the
plate thickness of c01 phase has increased and started to
form in a crisscross with each other obtaining a V shape
crossed with the needle-like of b01 phase, whereas, the c2
precipitates have shown the same decrement of the previ-
ous alloy as shown in Fig. 1h, q.
From the microstructures of 1.0 mass% Co alloys aged
at 373 K for 24 and 48 h, it was seen that the plate sizes of
b01 and c01 phases became more visible, thus corresponding
with the increase in the thickness of the coarse variation of
the c01 phase. In addition, these variants started to obtain a
discontinuous growth with the increase of the distance
between them, as shown in Fig. 1c, l. However, after
increasing the aging temperature to 423 K, the coarse
variants showed a complete discontinuous growth along
with obtaining random orientations. Moreover, the size of
the c2 precipitates was reduced due to the dissolution of
these particles into the microstructure, as shown in Fig. 1f
and o. When the aging temperature increased to 523 K for
24 and 48 h, it was found that a completed growth of the c01phase with a lamella structure was occurred, and the plate-
10 μm Electron image 1
1
AlCo
Ni
Cu
Co
Elements Wt.% At.%
Al 9.58 19.89Ni 0.96 0.91Co 4.71 4.49Cu 84.75 74.71
Spectrum 2
Ni
Full Scale 6308 cts Cursor: 3.202 (56 cts) keV2 3 4 5 6 7 8 9 10
1Full Scale 2783 cts Cursor: 5.338 (22 cts) keV
2 3 4 5 6 7 8 9 10
CoNi Al Elements Wt.% At.%
Al 56.34 73.8Co 38.52 23.1Ni 5.14 3.1
Spectrum 1
(a)
(b)
(c)
Fig. 2 A spot scanned of the EDS analysis of the aged Cu–Al–Ni–0.4 mass% Co a Micrograph of scanned area; b Spectrum 1; c Spectrum 2
1276 S. N. Saud et al.
123
like groups of the b01 phase become thicker, as shown in
Fig. 1i, r, along with an increase in the size of the c2
precipitates.
Figure 3(a–c) shows the XRD diffraction patterns of the
aged Cu–Al–Ni–xCo (x is 0.4, 0.7, and 1.0 mass%).
According to the influence of the aging treatment, the
obtained peaks vary in terms of presence, 2h value, and
intensity, which are completely reflected on the structure
parameters, such as miller induced or crystallite size. These
variations are attributed to the morphology and orientation
of the martensitic phase, and to the volume fraction and size
of the precipitates. Generally, the pattern peaks of the XRD
diffraction represent the existing phases into the micro-
structure, in which the pattern peaks of the peaks of (200)
and (202) represent the c01 phase, and (122), (0018), (128),
(1210), (2010), (1123), (208), (320), (040), and (311) rep-
resent the b01 phase. The pattern peaks of (-421), (712), and
(314) represent the c2 precipitates. According to the
matching standard of this alloy, it was found that these
peaks are related to the (Al75Co22Ni3) phase that we refer to
as c2. Furthermore, it was observed that the c2 precipitate
pattern peaks vary in intensity values and shift according to
the volume fraction of these precipitates in the micro-
structure, which is in complete agreement with the micro-
structural variations. At 0.4 mass% Co addition-aged
alloys, the results show that peak patterns varied in terms of
presence and intensities due to changes in aging time and
temperature. It was also found that the highest intensities
30
Inte
nsity
/a.u
.
(200
)
(122
)(2
02) (0
018)
(128
)
Inte
nsity
/a.u
.
Inte
nsity
/a.u
.
Inte
nsity
/a.u
.
35 40 45 50 552θ/°
60 65 70 75 80 30 35 40 45 50 552θ/°
2θ/°
60 65 70 75 80
30 35 40 45 50 552θ/°
60 65 70 75 80
(121
0)
(040
)
(311
)
(421
)
(112
3)(2
08)
(712
)
(314
)
(200
)
(122
)
(200
)
(122
)(2
02)
(128
)(1
210)
(201
0)
52 54 56 58 60 62 64 66 68 70
(001
8)
(112
3)
(208
)
(712
)
(314
)
(202
) (001
8)(1
28)
(121
0)(2
010)
(320
)(0
40)
(311
)
(112
3)
(208
)
(712
)
(314
)
(421
)
(421
)
(320
) (040
)
(311
)
373 K-24 hr
373 K-48 hr423 K-24 hr
423 K-48 hr523 K-24 hr
523 K-48 hr
JCPDS XRDAl75Co22Ni3(49-1278)
Inte
nsity
/a.u
.
Inte
nsity
/a.u
.
2θ/°
2θ/°
5550 60 65 70
48 50 52 54 56 58 60 62 64 66 68 70
(a) (b)
(c)
Fig. 3 X–ray diffraction patterns of a Cu–Al–Ni–0.4 Co SMA, b Cu–Al–Ni–0.7 Co SMA, c Cu–Al–Ni–1.0 Co SMA
Thermal aging behavior in Cu–Al–Ni–xCo 1277
123
were observed with the 423 K aging temperatures for 24 h.
This may be attributed to higher volume fraction of c2
precipitates into the matrix and at the grain boundaries of
this alloy. Almost the same behavior has been repeated for
the 0.7 and 1.0 mass% aged alloys as shown in Fig. 3b, c.
The lattice parameters and crystallite size of aged alloys
of Cu–Al–Ni–Co SMA were determined from the XRD
patterns and are recorded in Table 1. The lattice parameters
were evaluated in accordance to the orthorhombic 18R
structure, which was proven by the XRD indexing patterns.
Thus, the lattice parameters were determined using the
following relation [30]:
1
d2¼ 1
a2
h2
sin2b
� �þ k2
b2þ 1
c2þ l2
sin2b
� �� 2hl cosb
ac sin2b: ð1Þ
The crystallite size was determined by a Scherrer
equation [31, 32] for the highest intensity of the two peaks
of (0018) and (128), as follows:
Crystallite size dð Þ ¼ 0:9 � ðkÞB � cos h
; ð2Þ
where k is the XRD wavelength, b is the full width at half
maximum, and h is the Bragg’s angle. It was found that the
lattice parameters and crystallite size are varied with the
variation of aging time and temperature. However, the
highest crystallite sizes were observed 19.4, 17.4, and
24.8 nm with the 373 K–24 h for 0.4, 0.7, and 1 mass%
Co-aged alloys, respectively. The maximum value was
obtained with the aged 1.0 mass% Co alloys.
Transformation temperatures
Figure 4 shows the endothermic and exothermic curves of
the Cu–Al–Ni–1.0 mass% Co SMA aged at 523 K for
24 h. After applying the aging treatment at different times
and temperatures, the transformation temperatures are
slightly increased. The determined data from the DSC
curves for the aged and unaged alloys are presented in
Table 2. The transformation temperature increased gradu-
ally with an increase in the percentage of Co addition from
0.4 to 1 mass%, on the other hand, as a comparison with
the aged samples, it was also found that the transformation
temperatures decreased with increase in the aging times
and temperatures. This may attributed to the fact that the
aging treatment led to increase in the dislocation density,
and thus decrease the transformation temperatures [33].
The transformation temperatures of the aged alloys are
varied in according to the variations of the structures and
morphologies as well as the variations of volume fraction
of precipitates/intermetallic compounds that are associated
during the aging treatment. In other words, the phenomena
behind the variations in transformation temperatures are
related to the nucleation and growth of c2 precipitates that
may induce a stress field by the coherent boundaries
between the existing phases of c01 and b01 and precipitates of
c2 [34]. Remarkably, during the aging process, this stress
field can affect the transformation temperatures. Further
growth of the c2 precipitates can lead to loss of coherent
interface, and therefore, its influence on the transformation
Table 1 Lattice parameters and crystallite size of Cu–Al–Ni–XCo SMA under different aging conditions
Alloys Conditions a/A´
b/A´
c/A´ b Crystallite size/A
´
Cu–Al–Ni–0.4 mass% Co 373 K–24 h 3.711 5.79 43.985 96.364 194
373 K–48 h 3.877 5.125 39.285 74.467 168
423 K–24 h 4.132 5.268 38.406 91.924 167
423 K–48 h 3.892 5.145 39.573 73.219 180
523 K–24 h 3.914 5.211 39.605 73.419 157
523 K–48 h 4.0913 5.289 39.449 94.687 185
Cu–Al–Ni–0.7 mass% Co 373 K–24 h 4.138 5.234 38.186 87.570 174
373 K–48 h 4.138 5.238 38.189 87.569 172
423 K–24 h 4.108 5.254 38.625 94.091 143
423 K–48 h 4.144 5.234 38.206 87.757 169
523 K–24 h 3.747 5.287 38.523 98.532 172
523 K–48 h 3.821 5.183 38.569 79.557 161
Cu–Al–Ni–1 mass% Co 373 K–24 h 4.139 5.250 38.086 87.912 248
373 K–48 h 4.002 5.155 38.869 75.639 204
423 K–24 h 4.131 5.237 38.051 88.771 198
423 K–48 h 3.988 5.153 38.938 75.351 226
523 K–24 h 3.753 5.267 38.592 94.777 187
523 K–48 h 3.753 5.268 38.580 95.016 204
1278 S. N. Saud et al.
123
temperatures. A number of the aged alloys have shown
multiple endothermic/exothermic peaks in their transfor-
mation curves; their multiple peaks are attributed to the
interface transformations. These intermartensitic transfor-
mations are a first-order phase transformation between
martensites with different structures at temperatures below
the Ms. So far, several intermartensitic phases have been
found in the shape memory alloys, which have modulated a
lattice with different periodicity of stacking sequences [35].
However, the structures and transformation temperatures of
these intermartensitic phases depend on the levels of
applied stress and the chemical composition. On the other
hand, these interphases are very sensitive to the internal
stress of the alloy [36].
Meanwhile, the differences of the austenitic transfor-
mation temperatures (Af–As) are mainly larger than the
46018.5
19
19.5
20
20.5
21
21.5
470
Heating
480 490 500 510
Temperature/K
Hea
t flo
w/m
W
520 530 540 550 560
503 503.5 504 504.5 505Temperature/K
Hea
t flo
w/m
WH
eat f
low
/mW
505.5 506 506.5 507
512
20.94
20.96
20.98
21
21.02
21.04
21.06
18.7
18.75
18.8
18.85
18.9
18.95
513 514 515 516 517 518 519Temperature/K
End
o
Cooling
Fig. 4 DSC diagrams on the
heating and cooling cycle of the
Cu–Al–Ni–1.0 mass% Co SMA
aged at 523 K for 24 h
Table 2 Transformation temperatures of aged Cu–Al–Ni–xCo SMA
Alloys Conditions Transformation temperatures/K
As Af Ms Mf Af–As Ms–Mf To
Cu–Al–Ni–0.4 mass% Co 0 513.4 521.5 510 505.3 8.1 4.7 515.75
373 K–24 h 508.2 515.6 504.5 490.9 7.4 13.6 510.05
373 K–48 h 499.8 514.5 503 490 14.7 13 508.75
423 K–24 h 511.86 516.15 508.1 498.3 4.29 9.8 512.125
423 K–48 h 510 517.56 506 503.2 7.56 2.8 511.78
523 K–24 h 510.1 517.8 506.1 503.7 7.7 2.4 511.95
523 K–48 h 510 517.86 506 503.1 7.86 2.9 511.93
Cu–Al–Ni–0.7 mass% Co 0 515.68 525.2 514.2 507.7 9.52 6.5 519.7
373 K–24 h 508.91 515.66 504.36 490.1 6.75 14.26 510.01
373 K–48 h 500 514.41 503.1 489.8 14.41 13.3 508.755
423 K–24 h 511.8 516.3 508.1 498.6 4.5 9.5 512.2
423 K–48 h 510 517.5 506 503.4 7.5 2.6 511.75
523 K–24 h 510.3 518 506.7 503.4 7.7 3.3 512.35
523 K–48 h 509.9 515 506 503 5.1 3 510.5
Cu–Al–Ni–1 mass% Co 0 522.88 532 521.2 514.4 9.12 6.8 526.6
373 K–24 h 508.4 516 504.5 491.3 7.6 13.2 510.25
373 K–48 h 500.4 514 503 490.6 13.6 12.4 508.5
423 K–24 h 511.4 516.3 510.1 498 4.9 12.1 513.2
423 K–48 h 510.2 518 506.1 503.4 7.8 2.7 512.05
523 K–24 h 513.4 518.2 506.3 503.6 4.8 2.7 512.25
523 K–48 h 510.2 518.1 506.1 503.5 7.9 2.6 512.1
Thermal aging behavior in Cu–Al–Ni–xCo 1279
123
martensitic transformation temperatures (Ms–Mf), except
the aged alloys of 373 K for 24 h of the 0.4 mass% Co
addition, and with 373 and 423 K for 24 h of the
0.7 mass% Co and 1.0 mass% Co addition, which means
the (Ms–Mf) [ (Af–As). On the other hand, a small elastic
strain occurred associated with the occurrence of the plastic
relation and/or lattice softening [37]. However, within the
aging treatment, there will be some aging-induced lattice
softening or maybe some variations in the long-range of the
martensitic structure order.
Mechanical properties
Stress–Strain curve behavior and microhardness test
The typical stress–strain curves of the Cu–Al–Ni–xCo-
aged alloys are presented in Fig. 5a–c. It can be observed
that the tensile properties of the aged alloys are varied in
terms of fracture stress and fracture strain in accordance to
the condition of aging treatment as shown in Table 3.
In the 0.4 mass% Co-aged alloys, the strength is
increased with increasing the aging temperatures from 373
to 423 K, and then dropped at 523 K as shown in Fig. 5a. It
may be attributed to the higher or the more heterogeneous
dispersion of precipitate, which has shown the highest
volume fraction associated with a large size with 423 K for
24 h. Furthermore, it was found that the presence of pre-
cipitate can pin the mobility of dislocations, and therefore
the dislocations are required for a higher stress, either to
shear the precipitate particles, if the particles are large,
which lead to increase the fracture stress or bow around
them resulting in high fracture strain. Figure 5b shows the
stress–strain curve of the 0.7 mass% Co-aged alloy. It
found that the strength of the alloys is increased with
increasing the aging temperatures, where it has observed
the highest strength with the 523 K for 24 h. This may be
00 0.5 1 1.5 2 2.5 3 3.5
100 °C - 24 hr
250 °C-48 hr
100 °C - 48 hr
150 °C - 24 hr150 °C - 48 hr250 °C - 24 hr250 °C - 48 hr
4
50100150200250300350400450
Tens
ile s
tres
s/M
Pa
Tensile strain/% 0 0.5 1 1.5 2 2.5 3 3.5 4 4.5 5 5.5Tensile strain/%
Tens
ile s
tres
s/M
Pa
Tens
ile s
tres
s/M
Pa
500550600650700750
050
100150200250300350400450500550600650700750
050
100150200250300350400450500550600650700750800
0 0.5 1 1.5 2 2.5 3 3.5 4 4.5 5 5.5 6Tensile strain/%
100 °C - 24 hr
100 °C - 48 hr
150 °C - 24 hr
150 °C - 48 hr
250 °C - 24 hr
250 °C - 48 hr
100 °C - 24 hr
100 °C - 48 hr
150 °C - 24 hr
150 °C - 48 hr
250 °C - 24 hr
250 °C - 48 hr
100
°C-4
8 hr
100
°C-2
4 hr
150 °C-24 hr
150 °C-48 hr
250
°C-2
4 hr
250 °C-48 hr
100
°C-4
8 hr
100
°C-2
4 hr
150 °C-24 hr
150 °C-48 hr
250
°C-2
4 hr
250 °C-48 hr
100
°C-4
8 hr
100
°C-2
4 hr
150 °C-24 hr
150 °C-48 hr
250
°C-2
4 hr
(a) (b)
(c)
Fig. 5 Stress–strain curves obtained from the tensile test performed at room temperature of a Cu–Al–Ni–0.4 mass% Co SMA, b Cu–Al–Ni–
0.7 mass% Co SMA, and c Cu–Al–Ni–1.0 mass% Co SMA
1280 S. N. Saud et al.
123
attributed to the high volume fraction of c01 plate-like that
has affected the dislocation movements, and thus,
increased the strength of the alloy. On the other hand, the
volume fraction, distribution, and size of the c2 precipitates
may also influence the dislocation movements, therefore
influencing the strength and ductility of the alloy [38–40].
However, the highest strain was observed with the 423 K
for 24 h alloy. This may be attributed to the large size of
the c2 precipitates, which force the dislocation to bow
around them, and thus increase the strain of the alloy. The
aged samples of 1.0 mass% Co alloy showed an increase
in the values of fracture strain and fracture stress as shown
in Fig. 5c. The determined data are given in Table 3.
From this point of view, it has been proven that the
brittleness of Cu-based SMA can be slightly reduced
using an aging treatment. These enhancements were
mainly attributed to the microstructure changes, which can
be considered as variations of martensitic phase mor-
phologies and orientations, as well as the behavior and
morphology of the c2 precipitates. It was also seen that the
elongation had reached the maximum value at the aging
condition of 373 K for 48 h, which improved the frac-
tured strain to 5.5 %.
The fracture surface areas of the Cu–Al–Ni–xCo SMAs
are displayed in Fig. 5a–c. It was determined that the
popular fracture features of the aged alloys are varied based
on the changing the time and temperature of aging treat-
ment. Figure 5a reveals the fracture area of the aged Cu–
Al–Ni–0.4 mass% Co SMA. The fracture feature of 373 K
for 24 h alloy shows a mix mode fracture of the inter-
granular and transgranular as shown in Fig. 5a. By
increasing the aging time to 48 h, the fracture feature
transferred from the mix mode to a dimple rupture along
with a little intergranular area exhibited on the left side of
the fracture area as shown in Fig. 5a, whereby the
appearance of these dimples represents a ductile fracture.
The 423 K for 24 h alloy displayed an aggregate mode
fracture between a little area of a quasi-cleavage and a
large area of intergranular feature as shown in Fig. 5a.
With further aging time and temperatures, the alloys show
an intergranular fracture feature, except the alloy 523 K for
24 and 48 h exhibited a little area of quasi-cleavage along
with the intergranular fracture as shown in Fig. 5a. How-
ever, Fig. 5b, c show the fracture features of the aged Cu–
Al–Ni–0.7 mass% Co, and Cu–Al–Ni–1.0 mass% Co
SMAs, respectively. It was indicated that the fracture fea-
tures of both alloys are behaving similar to the 0.4 mass%
Co-aged alloys, whereas the types of fracture are varied
between the transgranular, intergranular, dimple rupture
with a little intergranular, quasi-cleavage, and mix mode.
However, the variations of the fracture modes are attributed
to the amount of the elastic anisotropy at the grain
boundaries and the grain size of the aged alloys. Therefore,
as much as the amount of the elastic anisotropy at the grain
boundaries is small, the dimple rupture is occurring, which
represents the ductile fracture. When the elastic anisotropy
at the grain boundaries is large, a quasi-cleavage fracture is
observed, which represents a brittle fracture mode. This
Table 3 Results obtained from the tensile, microhardness, and shape memory tests on the aged Cu–Al–Ni–xCo SMA at different conditions
Alloys Aging conditions Stress/MPa Strain/% Microhardness/MPa Strain recovery/%
Cu–Al–Ni–0.4 mass% Co 373 K–24 h 620 2.2 331 73
373 K–48 h 600 2.45 317 76
423 K–24 h 680 2.8 345 82
423 K–48 h 710 3.4 358 77
523 K–24 h 670 2.6 340 80
523 K–48 h 700 2.7 351 81
Cu–Al–Ni–0.7 mass% Co 373 K–24 h 600 4.1 340 85
373 K–48 h 600 4.8 342 84
423 K–24 h 700 5 360 87
423 K–48 h 610 3.6 347 80
523 K–24 h 712 3.8 376 82
523 K–48 h 700 3.2 370 80
Cu–Al–Ni–1 mass% Co 373 K–24 h 625 4.5 360 76
373 K–48 h 700 5.5 372 73
423 K–24 h 680 4 367 71
423 K–48 h 650 3.5 356 80
523 K–24 h 750 3.75 387 82
523 K–48 h 740 5 381 100
Thermal aging behavior in Cu–Al–Ni–xCo 1281
123
study can conclude that aging treatment is one of the most
important factors that can influence the amount of an
elastic anisotropy associated with controlling the grain size
of the alloys, and thus, the variations of the aging times and
temperatures are able to vary the fracture features of the
alloys. On the other hand, the fracture surface of the alloys
can also be affected by the structural compounds, for
instance, parent phases, precipitates, and intermetallic
compounds along with their size and distribution into the
matrix [41, 42]. Microhardness of the aged Cu–Al–Ni–xCo
SMA with different aging times and temperatures are
shown in Table 3. The microhardness value varied sub-
stantially through increasing the aging time and tempera-
ture. This study found that the highest microhardness
values were 358, 370, and 387 MPa with the 423 K for
48 h alloy for 0.4 mass% Co, 523 K for 48 h alloy for
0.7 mass% Co, and 523 K for 24 h alloys for 1.0 mass%
Co addition-aged alloys, respectively. The increment of
microhardness after the aging treatment in the shape
memory alloys is essentially linked to dislocation move-
ments [43, 44]. The reason for the microhardness variations
is believed to be related to the c2 precipitates. When the
alloy is aged at a lower temperature such as 373 K, the
nucleation and growth rates of the c2 precipitates are
slower in which this phenomena is greatly depended on the
temperature. As the aging temperature increases to 423 K
and 523, the c2 precipitates will come out gradually, which
led to harden the alloys.
Shape memory effect
The shape memory effects of Cu–Al–Ni–xCo (x is 0.4, 0.7,
and 1.0 mass%) are shown Fig. 6a–c. It is found that the
SMEs are variations in the terms of strain recovery,
according to the condition of aging treatment as shown in
Table 3. For the 0.4 mass% Co, the highest results are
observed with the 423 K for 24 h alloy, which has obtained
an 82 % recovery out of the original shape after being
heated above the Af as shown in Fig. 6a. Generally
speaking, the strain recovery is associated with the mar-
tensite ? austenite transformation, and as well that the
magnitude of the plastic deformation is dependent on the
size and volume fraction of needle-like, plate-like, and
precipitates that occurred around them, and therefore,
00
50
100
150
200
250
300
350
400
450
500
550
600100 °C-24 hr100 °C-48 hr150 °C-24 hr150 °C-48 hr250 °C-24 hr250 °C-48 hr
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2 1.3
Tensile strain/%
Tens
ile s
tres
s/M
Pa
Tens
ile s
tres
s/M
Pa
0
50
100
150
200
250
300
350
400
450
500
550
600
Tens
ile s
tres
s/M
Pa
1.4 1.5 1.6 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2 1.3
Tensile strain/%1.4 1.5 1.6
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 1.2 1.3
Tensile strain/%1.4 1.5 1.6
0
50
100
150
200
250
300
350
400
450
500
100 °C-24 hr100 °C-48 hr150 °C-24 hr150 °C-48 hr250 °C-24 hr250 °C-48 hr
100 °C-24 hr100 °C-48 hr150 °C-24 hr150 °C-48 hr250 °C-24 hr250 °C-48 hr
Loading at 100 °C
Unloading at 100 °C
Recovery at 300 °C
Loading at 100 °C
Unloading at 100 °C
Recovery at 300 °C
Loading at 100 °C
Unloading at 100 °C
Recovery at 300 °C
(a) (b)
(c)
Fig. 6 Shape memory effect curves of a Cu–Al–Ni–0.4 mass% Co, b Cu–Al–Ni–0.7 mass% Co, and c Cu–Al–Ni–1.0 mass% Co
1282 S. N. Saud et al.
123
423 K for 24 h alloys has exhibited the highest strain
recovery. In addition, the flow strength of the austenite
phase may also significantly contribute to the enhancement
of SME that can be obtained as a result of the reordering of
martensite phase. On the other hand, the stiffened parent
phase is capable of accommodating the transformation
strain elastically, which can lead to maintaining the
coherent martensite/austenite interfaces during the direct
and/or reverse transformation with heat treaditions are
shown inted to a certain temperature or by applying a small
stress. However, the precipitation effects may also play an
important role in the enhancement of SME that have been
varied during the aging treatment. This effect can be
explained in terms of: (i) martensite $ transformation
temperature; (ii) the amount of induced stress during the
forward or backward transformation; (iii) the interaction
between the precipitates and dislocations. For the
0.7 mass% Co-aged alloys, the highest shape recovery
occurred with 423 K for 24 h alloy as shown in Fig. 6b.
The reason behind this improvement also relates to the
same phenomena previously mentioned with 0.4 mass%
Co-aged alloys. The 1.0 mass% Co-aged alloys show the
highest performance for the shape recovery with 523 K for
48 h alloy as shown in Fig. 6c, which obtained a completed
shape recovery (100 % of the original shape) after being
heated above the Af without obtaining any residual strain.
This may be attributed to the high volume fraction and size
of c01 phase. As mentioned previously, the increment in the
size and volume fraction of the plate-like can lead to
improving the SME by increasing the strain around the c01plate-like as proven by the micrographs in Fig. 1a–c. On
the other hand, the higher volume fraction and a large size
of c2 precipitates can hinder the movements of dislocations
and martensite variant interfaces associated with a decre-
ment of permanent strain, and thus, increase the shape
recovery.
Conclusions
Aging of the Cu–Al–Ni–Co SMAs at different tempera-
tures and times shows considerable variations in transfor-
mation temperatures and transformation hysteresis due to
successive martensitic transition and precipitate formation.
These variations are suitable factors to select this alloy to
be used for high temperature applications and fluctuating
environmental conditions. Moreover, aging also has a
noticeable influence on the stress–strain properties of the
alloys within the aging temperatures and time considered.
A significant variation in the fracture strain and fracture
stress value with the aging of the martensite alloys was
observed in the phase transition regime. With aging, the
shape memory effects were varied in terms of strain
recovery and this could be as a result of the variation of the
phase morphology and volume fraction of the c2 precipi-
tates. However, the highest strain recovery was observed
with the Cu–Al–Ni–1.0 mass% Co SMA after being aged
at 523 K for 48 h, in which it exhibited a 100 % recovery
of the original shape after being heated above the tem-
perature of Af.
Acknowledgements The author(s) would like to thank the Malay-
sian Ministry of Higher Education (MOHE) and Universiti Teknologi
Malaysia for providing the financial support and facilities for this
research, under Grant No. R.J130000.7824.4F150.
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