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Journal of Alloys and Compounds
journal homepage: www.elsevier.com/locate/jalcom
A comparative study of iron, cobalt or cerium micro-alloying on
microstructure and apparent viscosity of Al-5Ni alloy
Kang Wang
a,b
, Mingguang Wei
b,c
, Zhongmiao Liao
a
, Shuoxun Jin
a
, Bingbing Wan
a
,
Zhiqin Lei
a
, Peng Tang
c
, Jun Tian
a
, Lijuan Zhang
b,⁎
, Wenfang Li
a,⁎
a
School of Materials Science and Engineering, Dongguan University of Technology, Dongguan, Guangdong 523808, China
b
Institute of Science and Technology Innovation, Dongguan University of Technology, Dongguan, Guangdong 523808, China
c
State Key Laboratory of Featured Metal Materials and Life-cycle Safety for Composite Structures, Nanning, Guangxi 530004, China
article info
Article history:
Received 11 October 2022
Received in revised form 26 March 2023
Accepted 8 April 2023
Available online 10 April 2023
Keywords:
Al-Ni alloy
Micro-alloying
Solidification course
Apparent viscosity
Microstructure evolution
abstract
The micro-alloying is promising in improving the mechanical properties, thermal conductivity, and cast-
ability of the near-eutectic Al-Ni alloys, enabling them to be formed via the special casting or additive
manufacturing technique. But the rheological behavior of the molten near-eutectic Al-Ni alloys is not clear.
The present work investigates the effect of three kinds of micro-alloying agents (iron, cobalt, and cerium) at
0.3 wt% on the solidification and rheological behaviors of the Al-5Ni alloy. Results indicate that the addition
of the micro-alloying element can significantly refine the α-Al grains in the Al-5Ni alloy, leading to the
improvement of the thermal conductivity and mechanical properties. The cooling curves thermal analysis
suggests that the incorporated micro-alloying element can notably change the cooling rate of the melts
during the solidification course. Phase-field simulations illustrate that the different cooling rates cause the
change of the shape factor and the fraction of α-Al grains during the solidification course of the alloy,
leading to the remarkable difference of the viscosities at different temperatures and shear rates of the Al-
5Ni melts. This work provides an in-depth understanding of the correlation between the solidification route
and rheological behavior of the Al-5Ni melts with micro-alloying treatment.
© 2023 Elsevier B.V. All rights reserved.
1. Introduction
Nowadays, the demand for structural materials with superior
thermal conductivity, high strength, and lightweight has increased
due to the rapid expansion in the electronic communications and
electric vehicles (EVs) [1,2]. The traditional casting alloys, such as Al-
Si, Al-Cu and Al-Mg systems, are extensively used in engineering
applications due to their advantages including low density and
sound castability [3,4]. But the thermal conductivities of them are
significantly smaller than 200 W/(m·K), which is getting more
challenging to satisfy the heat dissipation requirements of power
intensive electronic devices [5].
The fins in the aluminum heat dissipation parts must be ex-
tremely thin in order to meet the criteria of thermal management
components. Conventional cast alloys have great fluidity, making
them appropriate for additive manufacturing or custom casting.
However, the high solid solubility of alloying elements in the
aluminum lattice frequently imposes a restriction on their thermal
conductivity [6]. Although the solid solubility of the alloying ele-
ments in the aluminum lattice can be homogenized by annealing or
aging heat treatment, it causes uncontrollable deformation of the
cast parts [7,8]. As a result, developing a non-heat-treated cast alu-
minum alloy with high thermal conductivity is required.
The solid solubility of several industrial transition elements (such
as iron [9], cobalt [10], and nickel [11]) and rare earth elements
[12–16] in the aluminum lattice are extremely low, which are po-
tential to be used as the main alloying elements. The micro-alloying
treatment is deemed as the good strategy to effectively enhance the
mechanical properties aluminum alloys. By utilizing this relatively
simple method, the grain size or secondary phases in the alloy can be
significantly refined [15,16]. Luo et al. [10] found that the excellent
thermal conductivity of Al-2Fe-0.3Co hypereutectic alloy reached
235 W/(m K), resulting from the low solid solubility of iron and
cobalt in the aluminum lattice. Nevertheless, there might be chal-
lenging while this kind of alloy is used for the casting formation.
Firstly, the eutectic point of the Al-Fe system is only about 1.8 wt%.
Additionally, the solidification interval of eutectic Al-Fe alloy is re-
latively short. These characters suggest that the latent heat of Al-Fe
Journal of Alloys and Compounds 952 (2023) 170052
https://doi.org/10.1016/j.jallcom.2023.170052
0925-8388/© 2023 Elsevier B.V. All rights reserved.
]]]]
]]]]]]
⁎
Corresponding authors.
E-mail addresses: ljz201709@126.com (L. Zhang), mewfli@163.com (W. Li).
eutectic alloy released from the solidification course might be in-
sufficient to maintain the low fluidity of the melts during the mold-
filling process.
One of the most promising alloy groups being developed is based
on the Al-Ni eutectic system. Owing to the fibrous Al-Ni eutectic
structure coupling the low solid solubility of Ni atom in aluminum
and relatively high eutectic composition (at about 6.0 wt%), the Al-
Ni-based alloys exhibited excellent electrical conductivity and
strength [11, 17–18]. The Al-Ni eutectic alloy possesses great feeding
ability and very low susceptibility to hot-tearing, which it is suitable
for the additive manufacturing [19]. Although eutectic alloys are
regarded to possess excellent castability [20], the molten alloy
generally solidifies to the ingot under the non-equilibrium condition
in the practical manufacturing process. This kind of condition
usually leads to the segregation of coarse hypereutectic phases in a
eutectic alloy [21,22]. Hence, compared with the eutectic alloy, the
hypoeutectic or near-eutectic composition system is more suitable
for producing structural parts with good mechanical or thermo-
physical performances.
It is noteworthy that the semi-solid interval of the near-eutectic
Al-Ni alloys is not as wide as the traditional cast aluminum alloys.
There have been several attempts to improve the fluidity of the Al-Ni
hypoeutectic alloys by incorporating the third alloying element,
manganese [22,23], thereby can be cast by special casting methods,
e.g., high-pressure die-casting, squeeze casting, and semi-solid die-
casting. But the solid solubility of manganese in aluminum is rela-
tively high, which is harmful to the thermal conductivity. Our pre-
vious work [13] reported that the trace addition of ytterbium in the
Al-5Ni alloy was effective in improving the strength and maintaining
excellent thermal conductivity. The rare earth element with con-
siderably low solid solubility in aluminum, and thereby can sig-
nificantly refine the α-Al grains in the alloy. Regarding the selection
strategy of the micro-alloying elements, rare earth, especially
cerium, has been widely used to modify the coarse secondary phase
in cast aluminum alloys [2,24]. The atom radiuses of transit elements
and rare earth elements are significantly larger than the aluminum,
which are inclined to interact with the nickel atoms rather than
dissolving in the aluminum lattice. For the sake of manufacturing
the thin-wall parts with special casting methods, there is an urgent
need for understanding the effect of micro-alloying treatment on the
rheology behavior of the near-eutectic Al-Ni alloys.
Several researchers in the past two decades tried to find the
correlation between micro-alloying treatments and the fluidity of
cast aluminum alloys. Taghaddos et al. [25] pointed out that the
fluidity of the molten Al-Si cast alloys with high iron content was
relatively poor, owing to the formation of coarse iron phases with
irregular shapes in the melts. But the fluidity was remarkably im-
proved by micro-alloying the manganese into the melts. Prukkanon
et al. [26] found that the addition of Sc or Zr elements at 0.2–0.4 wt%
were favorable to the castability improvement of the commercial
A356 alloy. Li et al. [27] indicated that the fluidity of commercial
ADC12 alloys was greatly promoted while adding lanthanum at
0.3–0.9 wt%. Lü et al. [28] stated that the fluidity and electrical
conductivity of commercial-purity aluminum was able to be sig-
nificantly upgraded by adding cerium at 0.3 wt%. These studies il-
lustrated that the fluidity of aluminum alloys was notably increased
while limiting the addition content of the micro-alloying elements
below 1 wt%. However, most of the fluidity evaluations of the alu-
minum alloys in these studies are based on the spiral fluidity mold,
and the rheology behavior, relating to the solidification course, has
been rarely explored. Using the differential scanning calorimetry
(DSC) technique, Mukherjee et al. [29] unveiled the effect of trace
scandium addition on the solidification route of the Al-Zn alloys.
Nevertheless, because the DSC experiment requiring a sample with a
small volume, the non-equilibrium solidification behavior of a
molten ingot might not be equivalent to the solidification route as
displayed in the DSC result. The cooling curve thermal analysis
(CCTA) technique [30] has been found to be justified in describing
the solidification characters of the large ingot.
In the present study, the near-eutectic Al-5Ni ingots with the
trace addition of iron, cobalt, or cerium were prepared for the mi-
crostructures and physical properties investigation. The rheological
properties and solidification characters of the molten samples were
studied in depth by simultaneously using a high-temperature type
Couette viscometer and the CCTA apparatus. The microstructure
observations at room temperature and theoretical investigations of
the solidification were provided to analyze the rheology behavior of
the alloys.
2. Experiments and methodologies
2.1. Alloys preparation
The high-purity aluminum (purity > 99.9 wt%) and the Al-10Ni,
Al-10Fe, Al-20Ce, and Al-10Co master alloys (the numbers represent
the mass percentage of solute metals, the same below) were selected
as the raw materials. To fabricate the Al-5Ni alloys, firstly, the pure
Al ingots were molten in a mid-frequency induction furnace with 2
lpm argon gas blowing. After that, the Al-10Ni ingots with pre-
determined amount were added into the molten aluminum at
750 °C. To further prepare the Al-5Ni alloys with 0.3 wt% addition
amount of iron, cobalt, or cerium, the different master alloy ingots
were individually added into the Al-5Ni alloys in separate furnaces.
A graphite rod was utilized to stir the molten alloy for homo-
genization. After stirring for 10 s, the liquid with high temperature
was held in the furnace for 5 min at 700 °C. Then, it was poured into
several steel molds to fabricate the rod-like sample (Φ15 × 30 mm)
for metallographic inspection, the cylinder sample (Φ29 × 50 mm)
for rheology test and ingots (10 mm × 90 mm × 45 mm) for physical
properties samples preparations.
2.2. Microstructures and phases identifications
The prepared rod-like samples were cut into small ingots
(Φ15 × 5 mm) for microscope sample preparation. These ingots were
polished and etched with 0.5 % aqueous hydrofluoric acid for 20 s.
Microstructures of the as-cast alloys were observed using an optical
microscope (OM, Leica DMI-3000) and scanning electron microscope
(SEM, FEI Verios G4). The energy dispersive spectrum (EDS, Oxford
X-Max-N) was utilized to determine the composition of secondary
phases in these alloys. Phase constitutions of the fabricated alloys
were analyzed by an X-ray diffraction system (XRD, Bruker D8
Advance) with Cu-Kα radiation. The scanning rate of XRD tests was
set as 2°/s at the 2θ range from 10° to 90°.
2.3. Physical properties characterizations
The specimens for tensile tests were machined out from the in-
gots with a gauge length of 50 mm. Tensile tests were measured
using an Instron 5982 testing machine with a strain rate of 1 mm/
min at room temperature. The NETZSCH LFA 457 laser con-
ductometer was utilized to measure thermal conductivity. The
samples for this measurement were also cut from the ingot to the
dimension of Φ12.6 × 3 mm. The tensile test and the thermal con-
ductivity measurement were repeated three times for each alloy.
The rheology properties of the specimens were characterized via
the high-temperature type Couette rheometer (Anton Paar FRS-
1600) according to the DIN 53019 standard [31]. The sketch map of
the apparent viscosity test is shown in Fig. 1, and the shearing rate
range and temperature range of this apparatus are 0–73 s
−1
and
25–1600 °C. The precision of temperature controlling of the furnace
in this apparatus is ± 2 °C. The as-cast cylinder samples were
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
2
machined out to the dimension of Φ29 × 40 mm. Initially, the cy-
linder bar was deposited into the graphite container and heated to
750 °C in the furnace of the rheometer. After the ingot was molten,
the rotor of the Couette viscometer was immersed into the melts and
stayed at 1.5 mm above the bottom of the graphite crucible. In this
test, the rotation process of each testing point lasted for 180 s to
obtain 60 apparent viscosity values at a specific shearing rate.
Generally, the fluidity of the molten alloy would be affected by the
temperature, shear rate applied to the melts and the resting time of
the melts [32]. Hence, the present study was primarily concerned
with the apparent viscosity and rheology behavior of the molten
alloy under the steady state with constant temperatures and shear
rates. The apparent viscosities at five shear rate points (0.1, 1.0, 10.0,
33.0 and 73.0 s
−1
) were measured for each sample at a specific
temperature. Five constant temperatures, 660, 650, 645, 640, 635,
and 630 °C, were set for the viscosity tests.
2.4. Solidification analysis
The CCTA detections were conducted to investigate the correla-
tion between solidification and the apparent viscosity of the alloys.
During the solidification course, the K-type thermocouples (with a
measurement accuracy of ± 0.01 °C) and the NI 9212 series high-
speed data acquisition system were carried out for recording the
temperature variation of the melts. The solidification behavior of the
molten alloy was characterized by analyzing the temperature-time
(T-t) curves on the basis of Newtonian thermal analysis theory [30].
For the sake of understanding the evolution rule of the micro-
structure, the formation route of the phases in the alloy was de-
termined by referring equilibrium phase diagrams of the alloys [33].
Due to the dominent phase of the Al-5Ni alloys is α-Al, its evolution
behavior can notably influence the rheology behavior of the melts at
a specific temperature. After understanding the solidification beha-
vior of the molten alloy, the evolution route of these α-Al grains will
be discussed in depth by using the quantitative phase-field simula-
tion [34]. Meanwhile, morphologies of the α-Al grains in the mi-
crostructures of the samples were quantified and compared with the
phase-field simulation results. The simulations and image quantifi-
cations were conducted using the Fortran and MATLAB software,
respectively.
3. Results
3.1. Microstructures and phases identifications
Fig. 2(a) shows the optical microstructures of Al-5Ni alloys,
which reveals the α-Al grains in this alloy are coarse dendrite shape.
Meanwhile, the Al-Ni secondary phases are clustered to the polygons
and distributed around α-Al grains. The addition of iron or cobalt at
0.3 wt% greatly changes the morphologies of Al-Ni secondary phases.
As seen, the α-Al grains are significantly refined in these two sam-
ples (see Fig. 2(b) and (c)). It is also evident that the short-rod-like
secondary phase is precipitated at the boundaries of α-Al grains.
From Fig. 2(d), it can be seen that the grain refinement effect of
cerium addition in Al-5Ni alloy is not as significant as the iron or
cobalt addition, revealing a rounded α-Al grains in the micro-
structure.
The SEM images present the morphology details of secondary
phases in these Al-5Ni samples with different micro-alloying ele-
ments. Fig. 3(a) reveals that two kinds of secondary phases, with
particulate shape (labeled as A, eutectic boundary phase) and fi-
brous-like shape (labeled as B, eutectic internal phase), exist in Al-
5Ni alloy. It can be seen that the fibrous-like phase forms intensive
clusters, and the particulate phase distributes around these clusters.
Fig. 3(b) and (c) present the microstructures of Al-5Ni alloys with
0.3 wt% iron or cobalt added, respectively. Compared with Fig. 3(a), it
is found that secondary phases in these alloys are short-rod shape
(labeled as ‘C′ in Fig. 3(b) and ‘E′ in Fig. 3(c)) and particulate shape
(labeled as ‘D′ in Fig. 3(b) and ‘F′ in Fig. 3(c)). These secondary phases
are uniformly distributed at the grain boundaries of α-Al grains in-
stead of forming clusters. Fig. 3(d) displays the morphologies of
secondary phases in Al-5Ni-0.3Ce alloy. Different from the other
samples, the morphologies of secondary phases are elliptic shape
(labeled as ‘G′) and particulate shape (labeled as ‘H′). The phases also
form clusters which are similar to the Al-5Ni sample.
The EDS data corresponding to the labeled secondary phases in
Fig. 3 are given in Table 1. Neither nickel nor the micro-alloying
elements can be detected in the aluminum matrix. It is due to the
low solid solubilities of these elements in aluminum. According to
the relative studies [22], the secondary phase in the hypoeutectic Al-
Ni alloy is Al
3
Ni. By comparing the EDS results, it is found that the
content of nickel or the other elements in the eutectic external phase
with short-rod shapes is a few higher than the eutectic internal
phase with particulate or fibrous shape. This illustrates the nickel
atoms will agglomerate at the Al-Ni eutectic external phases. The
modification effects of iron, cobalt, and cerium on the Al
3
Ni eutectic
phase are quite different. The XRD patterns, as exhibited in Fig. 4,
unveil the phase compositions in these fabricated alloys. The trace
additions of Fe, Co, and Ce elements in Al-5Ni alloys lead to the
formation of Al
9
FeNi, Al
9
Co
2
, and Al
23
Ce
4
Ni
6
, respectively.
3.2. Solidification curves
Fig. 5 presents the cooling curves of the Al-5Ni alloys with dif-
ferent micro-alloying elements incorporated. To know the solidifi-
cation characters, the first- and second-derivatives of cooling curves
of the temperature-time curve are also provided in these diagrams.
The first-derivative curve is able to reflect the exothermic behavior
of the alloy, and the second-derivative curve is used for determining
the onset, termination, and phase transformation characters during
Fig. 1. Sketch map of the Couette-type viscometer for the viscosity measurements of the Al-5Ni alloys.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
3
the solidification course [30]. From the results, several peaks can be
clearly observed in the fluctuating second-derivative curves. A phase
transformation point can be determined when a peak with high
intensity intersects with the zero line for the first time. The symbols
of ‘A’, ‘B’, ‘C’, and ‘D’ are marked adjacent to the characteristic points
in the diagrams, representing the onset of solidification, the onset of
the eutectic structure formation, the end of the solidification, and
the formation of Al-Ni-Ce phase, respectively.
From Fig. 5(a) to (c), two stairs can be found in the first-derivative
curves. The first stair reflects the exothermic processes of the
Fig. 2. Optical microstructures of (a) Al-5Ni alloy, (b) Al-5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d) Al-5Ni-0.3Ce alloy.
Fig. 3. Morphologies of secondary phases in (a) Al-5Ni alloy, (b) Al-5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d) Al-5Ni-0.3Ce alloy.
Table 1
The EDS results of the secondary phases correspond to the labeled points in Fig. 3.
Compositions Labeled points inFig. 3
A B C D E F G H
Al/at % 91.4 92.6 84.8 96.9 89.5 97.3 86.4 91.1
Ni/at % 8.6 7.4 13.7 2.7 9.5 0.3 12.3 8.62
Fe/at % – – 1.5 0.4 – – – –
Co/at % – – – – 1.0 2.3 – –
Ce/at % – – – – – – 1.3 0.27
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
4
preliminary phase, and the second one is the exothermic course of
(α-Al+Al
3
Ni) eutectic structures and the secondary phases with the
micro-alloying elements. In terms of the Al-5Ni-0.3Ce alloy, one
more sharp peak occurrs after the second stair of the first-derivative
of the cooling curve (see Fig. 5(d)), illustrating that the solidification
route of the molten Al-5Ni-0.3Ce sample is notably different from
the other samples.
In order to understand the solidification behavior in depth, the
feature points of temperature and time of the solidification pro-
cesses are concluded and given in Table 2. Additionally, the time
durations of the solidification processes, t
Total
, are also provided. It
suggests that the addition of micro-alloying elements can prolong
the solidification durations. From the time characteristic points, the
solidification onset of the Al-5Ni alloy is earlier than the other alloys,
corresponding to a relatively higher nucleation temperature of the
preliminary phase. Owing to the thermodynamic colligative prop-
erty, the initial nucleation temperatures of the Al-5Ni alloys with
micro-alloying elements are lower. It is noted that the T
A
point of the
Al-5Ni-0.3Co alloy is slightly higher than that of the Al-5Ni alloy,
suggesting that the undissolved secondary phases might be gener-
ated before the nucleation of the α phase. In terms of the Al-5Ni
alloys and the alloys with 0.3 wt% iron or cobalt, it is found that the
solidification courses are terminated at 627.2–628.1 °C. But the so-
lidification duration of Al-5Ni-0.3Ce alloy is terminated at 610.8 °C.
This phenomenon illustrates that the micro-alloying elements are
able to significantly affect the solidification mode of the Al-5Ni alloy.
To understand the solidification modes of these molten alloys in-
depth, exothermic baselines of these specimens are determined by
utilizing the method as reported in Ref. [35]. According to this
theory, the correlation between f
s
and temperature is established
and given in Fig. 6(a). This diagram reveals that the solidification
processes of the alloys start from about 650 °C, and the fraction solid
of the alloys is sharply increased at the temperature range of
642–638 °C. From the top right corner of Fig. 6(a), the enlarged f
s
-T
diagram to displays the preliminary stage of solidification. In this
stage, it is obvious that the growth rates of f
s
of the Al-5Ni, Al-5Ni-
0.3Fe, and Al-5Ni-0.3Co alloys are higher than the Al-5Ni-0.3Ce alloy.
Although the T
A
of the Al-5Ni-0.3Co alloy is larger than that of the
Al-5Ni and Al-5Ni-0.3Fe alloy, the sharp increase of f
s
of the Al-5Ni-
0.3Co alloy is later than the other two alloys. Moreover, it is noted
Fig. 4. X-ray diffraction patterns of Al-5Ni alloys with different micro-alloying ele-
ments.
Fig. 5. Cooling curve thermal analysis of (a) Al-5Ni alloy, (b) Al-5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d) Al-5Ni-0.3Ce alloy.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
5
that an inflection point can be seen in this process, dividing the f
s
-T
curve as the nucleation stage and the growth stage of the pre-
liminary phase.
Fig. 6(b) shows the end stages of f
s
-T curves. As seen, a notable
inflection point is revealed in each of the f
S
-T curves, implying the
decrease in the growth rate of the solid phase. During the middle
stage of the solidification course, the f
S
values of the molten samples
are continuously elevating except for the Al-5Ni-0.3Co alloy. As the
fraction solid increases, the temperature is slightly decreasing and
then gradually rises. This phenomenon relates to the significant re-
calescence behavior of the Al-5Ni-0.3Co alloy. The recalescence point
of the Al-5Ni-0.3Co alloy reveals at f
S
= 0.28.
By comparison, the inflection points in the f
s
-T curves correspond
to the cooling feature of the T-t curve (see Fig. 6(c)–(f)). It can be
found that there are two courses with different growth rates of f
S
during the initial solidification stage of the samples. In this period,
for most of the samples, the T-t curves rapidly decline. But the re-
duction rates of T are continuously decreased. However, in the case
of Al-5Ni-0.3Ce alloy, the temperature is sharply decreased to a
minimum point, then it climbs up again to a relatively smooth stage.
This trend can also be observed in the other samples, but the
minimum point is not as notable as the Al-5Ni-0.3Ce sample. In
terms of the samples except for the Al-5Ni-0.3Co alloy, the f
S
curves
suddenly increase in the middle stage of solidification. As seen, in
the cases of the other samples, the sharp increase of the f
S
curve is
related to the increasing trend of the cooling rate As seen in the first
derivative curves (marked by an arrow in Fig. 6(c)–(e)).
It is noteworthy that the fraction solid curve of the Al-5Ni-0.3Ce
sample is relatively smooth (see Fig. 6(f)). It might be corresponding
to the shape of the T-t curve of this sample is quite different from the
others. The other samples possess a lower and more steady cooling
rates, whereas the cooling rate of the Al-5Ni-0.3Ce alloy is suddenly
promoted to a maximum point that is larger than 0. In this process,
the melts show a strong undercooling point, facilitating the growth
of the α-Al grains. Right after the maximum point, the cooling rate of
this sample is decreased to a steady stage, providing a considerably
stable latent heat to maintain the steady growth of the phases.
3.3. Physical properties
3.3.1. Rheological properties and behaviors
The effects of temperature and shear rate (dγ/dt) on the apparent
viscosity (η) of the samples are concluded in Fig. 7, and the corre-
sponding shear stresses (τ) are given in Fig. 8. At the temperature
ranged from 660 to 645 °C, the η value of the molten Al-5Ni alloy is
about 1 × 10
1
Pa·s at the shear rate of 0.1 s
−1
(see Fig. 7(a)). Then, it
shows a downward trend by enhancing the shear rate, revealing a
general shear-thinning character. However, the rates of decrease of
the η numbers slow down when the shear rate is larger than 10 s
−1
. A
similar phenomenon can also be found in the Al-5Ni alloys with the
Table 2
Characteristic temperature (T) and time (t) of the solidification route.
Alloys Time character (t/s) Temperature character (T/°C)
t
A
t
B
t
C
t
D
t
Total
T
A
T
B
T
C
T
D
Al-5.0Ni 464 900 1537 – 1073 653.2 638.9 627.2 –
Al-5Ni-0.3Fe 565 897 1706 – 1141 650.9 639.7 625.8 –
Al-5Ni-0.3Co 498 963 1677 – 1179 654.3 639.8 629.1 –
Al-5Ni-0.3Ce 573 865 1734 1650 1161 648.7 639.7 610.8 620.1
Fig. 6. The comparison of the Al-5Ni fraction solid curves at different temperatures (a) the initial stage of solidification, (b) the end stag of solidification and the comparison of
solidification evolution of the (c) Al-5Ni sample, (d) Al-5Ni-0.3Fe sample, (e)Al-5Ni-0.3 Co sample and (f) Al-5Ni-0.3Ce sample.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
6
micro-alloying elements incorporated (see Fig. 7(b)–(d)). Regarding
the Al-5Ni-0.3Fe alloy, it is evident that the viscosities at 645 °C are
apparently higher than the other samples.
The apparent viscosities and the rheology behavior of these
samples display significant differences when the temperature
reaches 640 °C. Herein, the η numbers are 1–3 orders of magnifica-
tion higher than the numbers of the corresponding sample mea-
sured at 645 °C, relating to the dramatical rise of the fraction solid
values adjacent to this temperature. The η values of the Al-5Ni-0.3Fe
and Al-5Ni-0.3Co specimens are linearly decreased by promoting the
shear rate. But in terms of the other two samples, it can be seen that
the reduction rates of η are slightly slowed down.
When the temperature is decreased to 635 °C, the η values no
longer exhibit the descending trend except for the Al-5Ni-0.3Ce
alloy. In this case, it is found that the viscosities are continuously
elevated by increasing the shear rates (From Fig. 7(a)–(c)). As far as
the Al-5Ni, Al-5Ni-0.3Fe, and Al-5Ni-0.3Co samples, the η numbers
at the shear rate of 73 s
−1
are about 1, 2.5, and 4.5 orders higher than
the numbers at 0.1 s
−1
. However, the viscosity of the Al-5Ni-0.3Ce
sample still reveals a relatively prominent decrease trend.
Fig. 8 displays the shear stresses of the molten alloys at the
corresponding temperatures and shear rates. According to the theory
of Newton-Couette flow [36], viscosity is the ratio of shear stress to
shear rate. Hence, the variation trends τ values at different shear
rates are similar to the η-dγ/dt patterns. When it comes to 635 °C,
most of the shear stresses applied to the melts reach 10
4
Pa at 635 °C,
approaching to the upper limit of the measuring range of the
viscometer. In this case, the viscosities are higher than 10
7
–10
8
Pa·s,
which is not acceptable in the casting formation.
3.3.2. Effect of micro-alloying on tensile properties and thermal
conductivities
The thermal conductivity and tensile properties of these samples
are summarized in Fig. 9. As seen, the yield strength (YS) and ulti-
mate tensile strength (UTS) of the 4 groups of alloys vary greatly. The
Al-5Ni alloy has a YS value of 89 ± 3 MPa, the lowest among these
samples. The micro-alloying treatments by adding iron, cobalt, or
cerium are able to improve the YS property of the alloy to some
extent, with optimized YS values of 141 ± 4, 138 ± 1 and
130 ± 2 MPa, respectively. The UTS property of the Al-5Ni alloy is
150 ± 5 MPa, which is close to the UTS value of the Al-5Ni-0.3Fe
sample. The UTS values of Al-5Ni-0.3Co and Al-5Ni-0.3Ce are dra-
matically increased to 180–190 MPa. The elongation values of these
Al-5Ni-0.3Ce alloy exceed 19 %, which are significantly higher than
the other alloys.
The micro-alloying treatment is demonstrated to be meaningful
in promoting the thermal conductivity of the Al-5Ni alloy. The
thermal conductivity (TC) of the Al-5Ni alloy is 208 ± 2 W/(m·K),
which is close to the Al-5Ni-0.3Co alloy. The Al-5Ni-0.3Ce alloy
possesses the maximum TC value among these samples, which is
about 5 % higher than the Al-5Ni alloy. By comparison, it is found
that the UTS and TC values of Al-5Ni-0.3Ce alloy exceed 190 MPa and
210 W/(m·K), respectively.
Fig. 7. Apparent viscosity measurements at different temperatures and shear rates of: (a) Al-5Ni alloy, (b) Al-5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d) Al-5Ni-0.3Ce alloy.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
7
4. Discussion
4.1. Improvement mechanism of strength and thermal conductivity
In the present work, the thermal conductivity of the Al-5Ni alloy
is mostly larger than 200 W/(m·K), which is significantly higher than
the traditional Al-Si cast alloys. The excellent thermal conductivities
are due to the relatively low solid dissolubility degree of Ni in the
aluminum lattice. From the phase diagrams [33], the solid solubility
of iron, cobalt, or cerium is also extremely low in the aluminum
lattice.
The structure of the as-cast hypoeutectic Al-5Ni alloys consists of
an α-Al dendritic matrix with a eutectic Al-Ni mixture. In the inter-
dendritic region, the eutectic mixture is formed by the α-Al and the
intermetallic compound Al
3
Ni [37,38]. Such a eutectic mixture nu-
cleates in a cooperative and alternating way during growth and re-
mains located between the dendritic arms. As indicated by several
studies [17,19,21,38], the fibrous Al-Al
3
Ni structure can significantly
enhance the strength of Al-5Ni and Al-5Ni-0.5Ce alloys, which may
allow the design of Al-Ni components with sound mechanical
properties. From the microstructures of the Al-5Ni-0.3Fe sample, it
can be observed that a more flake-like secondary phase is generated,
leading to the poor UTS property. Although several flake-like phases
also exist in the Al-5Ni-0.3Co alloy, the relatively high UTS and YS
properties are attributed to the refined α-Al grains and uniformly
distributed secondary phases.
Hence, utilizing iron, cobalt or cerium as the micro-alloying agent
is favorable to synergistically enhance the thermal conductivity and
the strength of the Al-5Ni alloys [39]. The experiment results illus-
trate that the addition of the micro-alloying elements is favorable to
the improvements of thermal conductivity and strength. These en-
hancements are closely related to the microstructures of the Al-5Ni
alloys modified by different micro-alloying elements. As a result,
before understanding the effect of the alloys’ composition on the
rheology behavior of the melts, the correlation between the solidi-
fication routes and the microstructures should be established.
Fig. 8. Averaged shear stress applied to the melts at different temperatures and shear rates of: (a) Al-5Ni alloy, (b) Al-5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d) Al-5Ni-0.3Ce
alloy.
Fig. 9. Tensile properties and thermal conductivities of the (a) Al-5Ni alloy, (b) Al-
5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d) Al-5Ni-0.3Ce alloy.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
8
4.2. Correlation between solidification and the microstructure
4.2.1. Evolution of the solid fraction in the molten samples
The incorporation of micro-alloying elements affects the solidi-
fication route of the melts. Fig. 10 shows the f
S
values at different
temperatures of the Al-5Ni melts with different micro-alloying ele-
ments incorporated at different temperatures. The f
S
values related
to the viscosity measurements are marked in these graphs. From
Fig. 10(a), it is evident that the f
S
values of the Al-5Ni samples are
sharply increasing when the temperatures are lower than 640 °C. At
640 °C, the fraction solid values of Al-5Ni and Al-5Ni-0.3Fe samples
reach 0.188 and 0.195, respectively (see Fig. 10(a) and (b)). Such a
high quantity of solid phase as revealed in the Al-5Ni-0.3Fe alloy is
resulted from the simultaneous precipitation of Al
9
FeNi and α-Al
during the initial stage of solidification [33, 40–41]. From Fig. 10(c),
the fraction solid of Al-5Ni-0.3Co alloy is significantly lower than the
other three samples at 640 °C, but a sudden elevation of f
S
number
occurrs right after 640 °C. This solidification pattern might corre-
spond to the α-Al, Al
9
Co
2
, and Al
3
Ni phases are simultaneously
formed at about 644 °C (see Fig. 6(c)) [33,42]. The f
S
value of the Al-
5Ni-0.3Ce sample at 640 °C, as revealed in Fig. 10(d), is 3 times
higher than that of the Al-5Ni-0.3Co sample owing to the significant
undercooling effect (see Fig. 6(d)). It is noteworthy that all of the f
S
numbers of the samples are sharply rose and approached to 1.0 at
635 °C, leading to the extremely high apparent viscosities of the
alloys. However, in terms of the Al-5Ni-0.3Ce sample, it can be seen
that an exothermic peak reveals right after the second stage of the
first-derivation curve (see Fig. 5(d)), resulting in the lower fraction
solid at 635 °C.
By comparison, the shapes of the f
S
-T curves of the four alloys are
relatively similar. Due to the dosage (0.3 wt%) of the micro-alloying
elements is small, the precipitation of the secondary phase might
not notably influence the viscosity of the alloy. But it will influence
the formation and morphology of the α-Al grains during the solidi-
fication process.
4.2.2. Morphological evolution of the α-Al grains during solidification
Several researchers [43,44] indicated that the viscosity of the
alloy is mainly influenced by the morphology of the α-Al grains.
Although the addition amount of the micro-alloying elements in the
Al-5Ni alloy is relatively tiny, it significantly impacts the T-t curve of
the alloy. As seen in Fig. 6, the cooling rates (the first-derivative
curve) of the molten alloys are notably different, which will re-
markably affect the morphology of the α-Al grains. In the present
work, the dimensionless phase-field modeling was used to under-
stand the effect of the solidification behavior on the evolution of the
α-Al grains. The governing equations of the phase field Φ and the
rescaled solute concentration U are respectively provided as [45]:
+
= + +
=
R t
m C Nt
W n W n W n
UR t
k m C
1| | ( )
[ ( ) ] | | ( ) ( )
( )
[(1 )| | ] (1 )
C
i x y
i
i
c
0
2
,
2 2
0
2 2
(1)
+
= ·
+
+
+
·
+
+
+
k k U
t
D
k D
DU
W K U
t
k U
t
(1 ) (1 )
2
1
2
(1 )
2
[1 (1 ) ]
2 2 | |
1 (1 )
2
L
S
L
0
(2)
with
=
+
+
U
k k
k k k
[(1 ) (1 ) ]
(1 )[ (1 ) (1 ) ]
C
C
2
0
.
where C is the local solute concentration, C
0
is the initial con-
centration, k is the solute partition coefficient, R
c
is the cooling rate,
and m is the liquidus slope. The W(n) is related to the surface energy
anisotropy factor (ε) [45]. The solid and liquid phases are re-
presented by Φ = 1 and Φ = −1, respectively.
The spatial length and time are rescaled by the interface width
W
0
and relaxation time τ
0
, respectively. It is reasonable to assume
that the micro-alloying elements will not significantly change the
parameters of C, C
0
, m, and k. Hence, the phase field value is dom-
inantly determined by the R
c
value. As seen in Fig. 6(c–f), two stairs
with different cooling rate values are revealed in each first-deriva-
tive curve, representing the cooling and recalescence process during
the solidification. Due to the phase-field simulation being based on
the dimensionless time and scale, the relative periods of the cooling
and recalescence process of these Al-5Ni samples are non-
dimensionalized and given in Fig. 11.
To reasonably simplify the calculation model, the simulation is
focused on the effect of cooling and recalescence effect on the α-Al
grains. The cooling rate in the phase-field is determined by im-
porting the first-derivative curve of the T-t curves, and the para-
meters for the phase-field analysis are listed in Table 3 [46]. The
maximum iteration of the numerical simulation is 2 × 10
4
steps are
long enough to reflect the effect of cooling and recalescence on the
morphology of the α-Al grains. From Fig. 2, it can be seen that the
dendrite-like α-Al grain was about 50–60 % area fraction of the
Fig. 10. Fraction solid of the molten alloys at the specific temperatures related to the viscosity measurements: (a) Al-5Ni alloy, (b) Al-5Ni-0.3Fe alloy, (c) Al-5Ni-0.3Co alloy and (d)
Al-5Ni-0.3Ce alloy.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
9
microstructure image. Hence, the simulations are terminated at
about 60 % of the area fraction (at 20,000 steps), which is approxi-
mated to the microstructure of the solidified samples.
Fig. 12 shows the numerical simulation results of the evolution
courses of the α-Al grains in the Al-5Ni alloys. Before the 2500 steps,
all of the molten samples stay in the cooling periods, so the α-Al
grains present the dendrite shape. Until the 5000 steps, the dendrite
tips of the α-Al grains of the samples, except for the Al-5Ni alloy, are
remarkably spherized. At the 10,000 steps, the dendrite tips of the α-
Al grains in all of the phase-field simulations are spherized. The
fractions, averaged shape factors, and the correlated temperatures of
the α-Al grains in these phase-field results are further quantified to
understand the morphological evolution of the α-Al grains.
Fig. 13(a) displays the variations in the fraction of the α-Al grains
in these samples. It reveals that the fractions of the α-Al grains are
sharply increased before the 5000 steps. However, due to the re-
calescence effect of the samples, the fractions of these semi-solid
melts are slightly decreased. Fig. 13(b) suggests the high nucleation
rate of α-Al grains in the initial solidification stage of the Al-5Ni-
0.3Ce alloy. However, the averaged shape factor (ASF) values of the
samples are continuously decreasing before the 5000 steps. Then,
the ASF of the Al-5Ni-0.3Ce alloy is sharply elevating at the steps
ranging from 10,000 to 20,000, but most of the samples are de-
clining at this time step range. Fig. 13(c) exhibits the temperature
evolution route during the growth course of the α-Al grains in the
phase-field model. The variation trends of the α-Al fractions are
generally consistent with those of the f
S
-T curves, as shown in
Fig. 10.
The shape factors of the α-Al grains in the fabricated samples are
also quantified and provided in Fig. 14 based on the microstructure
images. By comparison, it is found that the shape factor of the Al-
5Ni-0.3Ce is the highest, illustrating that the shape of α-Al grains in
this sample is spherized. The shape factor value of the α-Al grains in
the Al-5Ni-0.3Fe is the lowest. The microstructure of this sample
also reveals the dendrite-like α-Al grains, corresponding to the phase
field simulation result in Fig. 13(b). Chen et al. [43] thought that the
viscosity variation of the alloy was correlated with the shape of the
grains. They suggested that the clustering of the solid phases and the
ripening of the α-Mg grains would notably influence the viscosity of
the melts. Hence, the shape factor level of the α-Al grains should be
considered while analyzing the rheology behavior of the Al-5Ni
melts.
Fig. 11. Sketch map of the stages with different solidification behaviors of the Al-5Ni
samples with different types of the micro-alloying elements.
Table 3
Phase-field parameters of the Al-5Ni alloy [46].
Parameters Values
C
0
(initial concentration)/wt% 5
m (liquidus slope) 3.6
k (solute partition coefficient) 0.007
D
L
/m
2
/s 1.0 × 10
−9
D
S
/m
2
/s 1.0 × 10
−13
W
0
(interface width)/m 1 × 10
−7
τ
0
(relaxation time)/s 5.5 × 10
−7
ε (surface energy anisotropy factor) 0.01
Fig. 12. The phase-field simulation: growth of the α-Al grains in the Al-5Ni alloys affected by the micro-alloying elements of: (a) without micro-alloying elements, (b) with 0.3Fe,
(c) with 0.3Co, (d) with 0.3Ce.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
10
4.3. Mechanism of the rheological behavior
The phase-field simulation results indicated that the shape fac-
tors of the α-Al grains in the Al-5Ni alloys strongly relied on the
types of the added micro-alloying elements. Different micro-alloying
elements lead to varying extents of the constitutional supercooling
and recalescence of the melts, further leading to the difference in the
rheological behaviors of the molten alloys.
Fig. 15(a) shows the apparent viscosities of the Al-5Ni alloys with
different micro-alloying elements at 650 °C. At this temperature, the
fraction solid values of these samples are meager. As a result, their
viscosities, as well as the rheological behavior, of them are ap-
proximated.
At 645 °C (about 2500 steps of the phase-field), it is found that
the η value of the Al-5Ni-0.3Fe sample at 0.1 s
−1
is relatively high (see
Fig. 15(b)). It is due to the Al-5Ni-0.3Fe alloy possessing a relatively
high fraction of α-Al grains and a notably low ASF value (see Fig. 13(a)
and (b)). In this case, the fraction of α-Al grains in the melts is re-
latively low. The shear-thinning effect of the viscosities at the shear
rate ranged from 0.1 to 33 s
−1
is caused by the deagglomeration of
the α-Al nuclei and small dendrite grains. In addition, the increase of
the shear rate at this range might break the coarse dendrite grain.
Fig. 13. Evolution of (a) the area fraction of α-Al, (b) the averaged shape factors (ASF) of α-Al and (c) the phase-field f
S
-T relationship in the phase-field simulated microstructures
of the Al-5Ni samples.
Fig. 14. Shape factor evaluation of the Al-5Ni alloys (a) without any micro-alloying elements, (b) with 0.3 wt% Fe addition, (c) with 0.3 wt% Co addition, and (d) with 0.3 wt% Ce
addition.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
11
The viscosities of the melts are slightly increased at the shear rate of
33 and 73 s
−1
. The shear-thickening effect of the melts is due to the
solid phase might be agglomerated again under the centrifugal force
of the melts, providing resistance to the fluid.
In the case of 640 °C, the viscosities of the Al-5Ni-0.3Fe sample
are higher than the other alloys at different shear rates (see
Fig. 15(c)). This phenomenon mainly results from the high fraction
and low ASF value of the α-Al grains in this sample. The dendrite
grains would be interlocked with each other in the semi-solid flow of
the melts, further clogging the flow of the molten alloy. The fractions
of α-Al grains are also relatively high in the other samples. But the
spherized dendrite grains impede the interlocking of grains and
provide paths for the flowing melts. In addition, as seen in Fig. 3, the
secondary phase, Al
9
FeNi, of the Al-5Ni-0.3Fe is coarser than the Al-
5Ni and the Al-5Ni-0.3Co samples. This coarse phase will be formed
earlier than the α-Al grains in the melts [33], facilitating the inter-
locking of the solid phases.
Fig. 15(d) displays the viscosities of the samples at 635 °C, illus-
trating the high viscosities of the samples. The remarkably low
viscosities of Al-5Ni-0.3Ce alloy partly is attributed to a sharp exo-
thermic peak is revealed at the end of the solidification period of the
Al-5Ni-0.3Ce alloy (the ‘CD’ section in Fig. 5(d)). From Fig. 13(b), the
roundness values of the α-Al grains are mostly higher than the other
samples, facilitating the melts to flow through the gaps between the
grains.
From the above analysis, different rheological behavior of the
melts is correlated with the morphologies of the α-Al grains in the
samples. By referring to the phase-field simulation results, it can be
known that the change in the shearing rates might lead to different
interaction behaviors of the dendrite α-Al grains. The additions of
trace iron and cobalt influence the cooling rate of the Al-5Ni melts,
but their solidification intervals are nearly identical to the Al-5Ni
alloy. The addition of trace cerium not only significantly changes the
cooling rate of the melts, but also extends the solidification interval
due to the exothermic effect of the Al-Ni-Ce phase at the end period
of the solidification.
5. Conclusions
The present study found that the additions of different micro-
alloying elements in the Al-5Ni alloys caused a significant mod-
ification in the morphology of α-Al grains, leading to considerable
improvements in mechanical performance and thermal con-
ductivity. The castability of these alloys was in-depth investigated by
analyzing their rheological and solidification behaviors. The main
findings are concluded as follows:
(1) The α-Al grains of the Al-5Ni alloy were notably refined while
incorporating a trace amount of iron, cobalt or cerium at 0.3 wt%.
The trace addition of the transition elements iron, or cobalt in
the Al-5Ni alloy led to an over 50 % increase in the yield strength
while maintaining their high thermal conductivities
(> 208 W/(m K)).
(2) The fraction solid of the samples was less than 20 % when the
temperature ranged from 640 to 650 °C. The shear-thinning be-
havior of the melts, as revealed at the 0–33 s
−1
shear rates, was
caused by the deagglomeration of the nuclei of α-Al grains, the
secondary phases, and the dendrite grains. But these solid
phases might be agglomerated again while the shear rates were
increased, leading to the shear-thickening phenomenon.
(3) The temperature-time curves of the melts and their derivative
curves were used for describing the solidification behavior of the
melts, also providing the cooling parameters to the solidification
phase-field simulation. The averaged shape factors of the α-Al
grains, as derived from the phase-field simulations, were in ac-
cordance with the microstructures in the practical samples.
(4) The Al-5Ni semi-solid melts with the trace addition of iron re-
vealed high viscosities, owing to the relatively high solid fraction
and significant dendrite-like α-Al grains. The trace addition of
cerium in the Al-5Ni alloy not only significantly spherized the
dendrite tips of the α-Al grains, but also notably extended the
solidification interval. It led to the good fluidity of this sample at
the relatively low temperature of 635 °C.
CRediT authorship contribution statement
K. Wang: Investigation, Conceptualization, Simulations, Writing
– original draft. M. Wei: Programming and Simulations. Z. Liao:
Writing – original draft. S. Jin: Thermal analysis investigation. B.
Wan: Data curation, Microstructure investigation. Z. Lei:
Programming and Simulations. P. Tang: Thermal analysis investiga-
tion, Resources. J. Tian: Microstructure investigation. L. Zhang:
Methodology, Resources, Funding acquisition. W. Li: Supervision,
Writing – review & editing.
Fig. 15. Comparison of the apparent viscosities of the Al-5Ni samples with different micro-alloying elements at the specific viscosity measurement temperatures: (a) at 650 °C, (a)
at 645 °C, (a) at 640 °C, (a) at 635 °C.
K. Wang, M. Wei, Z. Liao et al. Journal of Alloys and Compounds 952 (2023) 170052
12
Data Availability
Data will be made available on request.
Declaration of Competing Interest
The authors declare that they have no known competing fi-
nancial interests or personal relationships that could have appeared
to influence the work reported in this paper.
Acknowledgements
This research is financially supported by Guangdong Major
Project of Basic and Applied Basic Research (2020B0301030001),
Guangdong Basic and Applied Basic Research (2021A1515010587,
2019A1515110135 and 2021B1515130010), Science and Technology
Service Network Initiative of Chinese Academy of Sciences
(Dongguan Special Project: 20211600200082); Dongguan Major
Project of Science (20211800904892); The Innovative Team Project
of Guangdong University (2021KCXTD022); National Natural Science
Foundation of China (52261024).
The authors would like to acknowledge the supplement of the
alloys preparation apparatus from Guangdong Fushengda Intelligent
Technology Co. Ltd (Dongguang, Guangdong). Meanwhile, we ac-
knowledge the help of alloys preparation from Guangdong Research
Center of High-performance Light Alloys Forming Technology, and
the Novel Light Alloy and its Process Technology Key Laboratory of
Dongguan City.
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