A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
663–674
INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION
AND MECHANICAL PROPERTIES OF LM6-ALLOY-BASED
FUNCTIONALLY GRADED COMPOSITE TUBES
VPLIV VSEBNOSTI NIKLJA NA MIKROSTRUKTURO IN RAZVOJ
MEHANSKIH LASTNOSTI KOMPOZITNIH FUNKCIONALNO
GRADUIRANIH CEVI NA OSNOVI ZLITINE VRSTE LM6
A. Saiyathibrahim
1
, S. Santhosh
2
, G. Raja Kumar
3
, S. Bharani Kumar
4
1
Department of Mechanical Engineering, Karpagam Institute of Technology, Coimbatore, Tamil Nadu 641 105, India
2
Department of Mechanical Engineering, Sri Krishna College of Technology, Coimbatore, Tamil Nadu 641 042, India
3
Department of Mechanical Engineering, Swarna Bharati Institute of Science and Technology, Khammam, Telangana 507 002, India
4
Department of Mechanical Engineering, Rajalakshmi Institute of Technology, Kuthambakkam, Tamil Nadu 600 124, India
Prejem rokopisa – received: 2023-02-20; sprejem za objavo – accepted for publication: 2023-10-22
doi:10.17222/mit.2023.799
Despite the growing demand for new materials, we used the horizontal centrifugal casting technique to synthesize functionally
graded composite (FGC) tubes using an LM6 alloy containing (3, 6 and 9) w/% nickel. All the fabricated tubes were evaluated
for variations in microstructure, hardness and tensile properties along the radial cross-section in three distinct zones (inner, tran-
sition and outer). X-Ray diffraction (XRD) results indicated the formation of in-situ Al3Ni in all FGC tubes, and these in-situ
tri-aluminides increased further with the addition of Ni. A comprehensive microstructural analysis across the tubes utilising
scanning electron microscopy (SEM) images showed that gathering in-situ Al3Ni particles keeps growing nearer the outer zone,
and primary Si cuboids increase in the inner zone. This accumulation of particles improved the mechanical properties at all three
zones of the FGC tubes compared to the LM6 tube having no nickel. The results of the hardness investigation showed that pre-
cipitated Al3Ni in FGCs has a beneficial impact on the enhancement of hardness. Furthermore, the observed UTS improvement
in the FGC tubes was clearly associated with the precipitation and strengthening action of Al3Ni intermetallic phases, while a
significant reduction in elongation has been noted. Due to the influence of centrifugal force and density variation, the tube con-
taining 9 w/% of Ni demonstrated good gradation among composite alloy-fabricated FGC tubes, with in-situ Al3Ni particles set-
tling largely in the outer zone and primary Si particles settling primarily in the inner zone across the radial thickness.
Keywords: functionally graded composite, microstructure, hardness, tensile strength
Zaradi povpra{evanja po novih inventivnih materialih so se avtorji posvetili razvoju proizvodnje funkcionalno graduiranih
kompozitnih cevi (FGC; angl.: Functionally Graded Composite) s tehniko horizontalnega centrifugalnega litja. Za sintezo
trislojne kompozitne cevi z razli~no strukturo in lastnostmi so kot osnovo uporabili Al zlitino vrste LM6 s (3, 6 in 9) w/% niklja
(Ni). Vse izdelane kompozitne cevi so okarakterizirali glede na variranje mikrostrukture, trdote in mehanskih lastnosti vzdol`
radialnega preseka posameznih treh slojev (notranje prehodne in zunanje cone oziroma plasti). Rezultati analiz z rentgensko
difrakcijo (XRD; angl.: X-Ray Diffraction) ka`ejo na in-situ tvorbo Al3Ni v vseh FGC ceveh in vsebnost teh in-situ nastalih
tri-aluminidov nara{~a z nara{~ajo~o vsebnostjo Ni. Obse`ne mikrostrukturne preiskave z uporabo vrsti~ne elektronske
mikroskopije (SEM; angl.: Scanning Electron Microscopy) preko cevi so pokazale, da se in-situ nastali Al3Ni delci zbirajo in
rastejo v bli`ini zunanje cone cevi, medtem ko se primarne Si kockice nahajajo znotraj notranje plasti cevi. To zbiranje oziroma
kopi~enje delcev izbolj{a mehanske lastnosti vseh treh slojev FGC v primerjavi s cevjo iz osnovne Al zlitine brez dodatka Ni.
Meritve trdote so pokazale, da izlo~eni delci Al3Ni v FGCs ugodno vplivajo na povi{anje trdote. Nadalje so natezni preizkusi
izvedeni na preizku{ancih iz posameznih slojev cevi jasno pokazali izbolj{anje kon~ne natezne trdnosti (UTS; angl.: ultimate
tensile strength) zaradi izlo~anja in utrjevalnega u~inka intermetalne faze Al3Ni. Isto~asno so opazili pomembno zmanj{anje
raztezka zlitin. Zaradi vpliva centrifugalne sile in variranja gostote ima cev z vsebnostjo 9 w/% Ni najbolj{o lastno gradacijo
med vsemi izdelanimi kompozitnimi zlitinami FGC cevi z in-situ nastalimi Al3Ni delci, ki se ve~insko nahajajo v zunanji plasti
kompozitne cevi in primarni Si delci, ki se ve~insko nahajajo v notranjem sloju notranje plasti vzdol` radialne debeline cevi.
Klju~ne besede: funkcionalno graduirani kompoziti, mikrostrukture, trdota, natezna trdnost
1 INTRODUCTION
Functionally Graded Composites (FGCs) are sophis-
ticated multifunctional structures that have recently
sparked widespread interest. The distribution of rein-
forcements in all such FGCs varies linearly from one end
to the other in the constructed component.
1
This gradient
composition makes these materials preferable in many
applications, such as cylinder liners, brake rotor discs,
pulleys, optical devices, armors, etc. Moreover, tailoring
properties at specific locations in a matrix phase by in-
troducing one or more reinforcements is more attractive
for FGCs, which is impossible in conventional compos-
ites.
2,3
Among the various FGCs investigated recently,
aluminium-based FGCs have been reported as having
better mechanical and tribological properties due to the
inclusion of reinforcements in appropriate volume pro-
portions. Nowadays, aluminides such as Al
3
Ti, Al
2
Cu,
AlB
2
,Al
3
Ni, and Al
3
Zr are used instead of ceramics to
impart advantages like high thermal stability, elastic
Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674 663
UDK 620.186:669.715 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 57(6)663(2023)
*Corresponding author's e-mail:
imsaiyath@gmail.com (Saiyathibrahim A.)

modulus, hardness and low thermal expansion coeffi-
cient. These aluminides also provide a strong bond be-
tween reinforcements and matrix, lowering the residual
stresses. Hence, the degree of failure is reduced during
the thermal loading cycle of components.
4
Aluminium-silicon (Al-Si) alloys are widely utilised
as engineering materials to manufacture castings, partic-
ularly in the automotive, military and aerospace sectors,
attributed to their superior strength-to-weight ratio,
strong thermal conductivity, great castability, simplicity
of fabrication and lightweight. The LM6 alloy is one of
several Al-Si alloys that are extensively used to make
components such as cylinder blocks and pistons. How-
ever, the mechanical properties of this alloy have been
observed to be reduced by microstructural flaws like
dendrites and silicon morphology.
5–7
In general, alloying
metals such as Cu, Ce, B, Ni and Mg encourage the for-
mation of intermetallic compounds, altering the proper-
ties of Al-Si alloys. These alloying elements can influ-
ence the properties by way of precipitation hardening
and solid-solution strengthening to create new low-den-
sity, high-durability components.
8
Nickel has a very low
solid solubility in aluminium and therefore forms Al
3
Ni
intermetallic after the solubility limit is reached. Further-
more, Ni has been identified to be the most effective con-
stituent for improving the high-temperature characteris-
tics of Al-Si alloys. Golmohammadi et al.
9
proposed that
adding Ni to LM6 alloy stimulated the development of
Al
3
Ni particles, which seemed to be related to expanded
wear resistance and hardness. Yang et al.
10
reported that
incorporating Ni into an Al-Si alloy boosted its ultimate
tensile strength while decreasing its elongation. Research
related to Al
3
Ni tri-aluminide reported that this
intermetallic has high hardness (841HV), which can be
used for the production of in-situ composites.
11
These
hard intermetallic phases are usually dispersed in a duc-
tile matrix to overcome low ductility and fracture tough-
ness, resulting in a good structural composite material.
Centrifugal casting is a simple process for producing
cylindrical FGCs with great microstructural control and
good mould filling, providing better properties. The con-
tinuous gradient in FGC using centrifugal casting is
achieved by centrifugal force created during mould rota-
tion. The change in density between the molten matrix
and reinforcement particles encourages the formation of
a gradient microstructure over the thickness of a cylindri-
cal tube constructed with a matrix material that includes
reinforcing particles (in-situ or ex-situ). During the cen-
trifugal casting of an Al-Si alloy, it was revealed that due
to the application of centrifugal force, primary Si parti-
cles migrated toward the inner surface across the radial
thickness of the produced tube. The lower density of pri-
mary Si particles in comparison to the molten matrix has
been identified as the origin of this reinforcement dis-
placement and the production of non-uniform micro-
structures known as gradation.
12,13
Lin et al.
14
fabricated
an Al-Si-Mg alloy tube in a centrifugal casting route,
which possessed a particle-free outer zone and an inner
zone with a lot of both in-situ Mg
2
Si and primary Si rein-
forcements. The high proportion of these in-situ Mg
2
Si
particles was observed to increase the mechanical prop-
erties in Al-Mg
2
Si FGCs, such as hardness and tensile
strength in the inner zone.
15,16
Duque et al.
17
claimed that
adding boron to aluminium increased the hardness in the
outer zone by precipitating and segregating in-situ AlB
2
reinforcement in a centrifugally cast tube. Rajan et al.
18
achieved an outstanding gradient structure in an Al-Ni
alloy-based FGC tube by constituting a high-density
in-situ Al
3
Ni-rich outer zone and an intermetallic free in-
ner zone. Further, it was highlighted that in FGCs, cen-
trifugal force is vital in creating gradation. The sorting of
reinforcements over the radial thickness has been ob-
served to be good at high centrifugal force, and particle
settling at various zones is determined by their density.
To be simpler, greater density reinforcements settle in
the outside zone, and their concentration decreases as
they approach the inner zone, whereas lighter particles
do the opposite.
19
Considering the existing literature, it can be asserted
that certain aluminium alloy-based FGCs created by cen-
trifugal casting and reinforced with several in-situ parti-
cles (Al
3
Ti, Mg
2
Si, Al
2
Cu, AlB
2
,Al
3
Zr, Al
3
Ni, etc) were
methodically explored. However, there are many alu-
minium alloys to be investigated in this in-situ reinforce-
ment context. The magnesium (Mg) effect on Al-Si al-
loys as well as the generation and dispersion of Mg
2
Si
particles along the radial cross-section of a centrifugally
cast tube, has been extensively researched in the previous
decade. However, there has been no reporting on the
characteristics of graded composites created using an
Al-Si alloy when introducing Ni, as well as the conse-
quent morphological alterations and reinforcement set-
tling processes. To bridge this gap, Ni was added to the
LM6 alloy in different concentrations to form FGC tubes
using a horizontal centrifugal casting in this study. Fur-
thermore, the microstructural characteristics and me-
chanical properties have been studied across the thick-
ness of the tubes. Also, the fracture mechanisms
associated with the composites were discussed.
2 MATERIALS AND METHODS
The intended nickel concentrations of 3, 6 and 9wt%
in the LM6 aluminium alloy were achieved by resistance
melting of a commercially available LM6 alloy with
Al-26%Ni master alloy in a graphite crucible. Figure 1
shows the microstructural features examined using an
optical microscope (OM) for an LM6 alloy with
Al-26%Ni master alloy used to synthesize tubular speci-
mens.
Table 1 shows the chemical composition of the alloys
developed for this investigation as well as the designa-
tion of the manufactured tubes. To fabricate tubes, the
prepared cast alloy bars were melted at a processing tem-
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
664 Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674

perature of 820°C under cover flux. A spinning graphite
stirrer was used to mix the composite melt for 10 min af-
ter degassing (500 min
–1
). The liquid metal was then
poured into a cylindrical grey cast iron (GCI) mould
(120 mm outer diameter & 110 mm length) that had been
heated to 300 °C and was spinning at 1900 min
–1
in a
horizontal centrifugal casting machine.
A quantity of 950g melt was used to fabricate the tu-
bular specimens (Figure 2) with an inner diameter of 50
mm, thickness of 15 mm, and length of 110 mm. Sam-
ples at three different zones (inner, middle, and outer) for
quantitative and microstructural observations were pre-
pared from the cast tubes, as shown in Figure 2.
Microstructural examinations were carried out along the
radial direction to observe zone-wise variations using
scanning electron microscopy with energy-dispersive
spectroscopy (SEM/EDS - JEOL JSM6360) following
standard metallographic routines. As a further step, an
image analyzer was used to evaluate the proportional
volume of the reinforcement phase segregation.
Vickers hardness measurements were made at five
spots (Figure 2) across the radial thickness in all tubes.
The Vickers hardness (HV) is expressed as the indenta-
tion force/indentation area ratio, as seen in Equation (1).
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674 665
Figure 2: Microstructure of a) LM6 alloy b) Al-26%Ni master alloy
Figure 1: Microstructure of a) LM6 alloy b) Al-26%Ni master alloy
Table 1: Alloys used in the production of cast tubes and their chemical composition
Designation Alloy used
Chemical composition in w/%
Si Fe Ni Cu Zn Mn Mg Al
LM6 Tube LM6 alloy 11.83 0.52 0.02 0.27 0.09 0.16 0.06 Balance
FGC
A
Tube LM6+3 w/%Ni 11.61 0.52 3.02 0.27 0.11 0.02 0.03 Balance
FGC
B
Tube LM6+6 w/%Ni 10.96 0.41 6.1 0.26 0.11 0.02 0.01 Balance
FGC
C
Tube LM6+9 w/%Ni 10.48 0.44 9.07 0.26 0.11 0.02 0.01 Balance

HV
P/
d
=
22
2
sin( )  (1)
where P is the indentation load (in N), d is the indenta-
tion diagonal length (in mm) and is the indenter angle
(in degrees).
20
The evaluations were conducted with a
50 g load and a 15 s dwell, and the values presented
were the average of five repetitions. Tensile specimens
were cut from three different zones by following
ASTM: B-557M standard as shown in Figure 3.Inor -
der to find the crack path and fracture mode, fractured
surfaces were examined by SEM.
3 RESULTS AND DISCUSSION
3.1 Microstructural Characterization
Cross-sections of the cylindrical tubes were taken in
the radial direction and microstructures of both LM6 al-
loy and FGCs were accessed in three distinct zones:
outer zone (within 5 mm), transition zone (5-10 mm),
and inner zone (10-15 mm) as shown in Figure 4. The
LM6 alloy tube microstructures (Figure 4a to 4c)) con-
sisted of a large quantity of flake-like eutectic Si phase
embedded in the aluminium matrix with some cuboids of
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
666 Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674
Figure 4: SEM images of LM6 tube and FGCs (FGC
A
,FGC
B
&FGC
C
) at three different zones
Figure 3: Dimensions of tensile test specimen

primary Si. In general, the eutectic LM6 alloy contains
eutectic Si phase and  -Al matrix.
21
Rapid cooling and
directed solidification during the centrifugal casting pro-
cess resulted in the nucleation of primary Si particles in
the microstructure of the eutectic Al-Si alloy in its
as-cast form.
22,23
Since solidification occurs in the above
manner during centrifugal casting, it is conceivable for
the LM6 alloy to include both eutectic Si phases and a
substantial number of primary Si particles. Furthermore,
it is frequently reported that increasing the silicon per-
centage in a cast Al-Si alloy increases the hardness by
reducing the -Al matrix.
24
Thus, it is reasonable to con-
clude that the Si phases are the key factor in increasing
the hardness of the Al-Si alloy. Primary Si particle segre-
gation in the LM6 tube was found to commence in the
outer zone and reached its maximum in the inner zone.
According to the microstructures, primary Si became
less accessible in the direction of the centrifugal force.
Besides, all three zones of the LM6 tube included
eutectic Si flakes. Although there were some differences
between the microstructures of the three zones consid-
ered for examination, they were not substantial enough
to claim that a fine gradation is achieved in the LM6
tube. This minor gradation is because of the eutectic na-
ture of the LM6 alloy and the presence of so few primary
Si particles.
On the other hand, the microstructure of functionally
graded composites (FGC
A
, FGC
B
& FGC
C
) mainly con-
sisted of the  -Al dendrite, primary silicon, eutectic sili-
con, primary Al
3
Ni intermetallic and eutectic Al
3
Ni
intermetallic surrounded by their boundaries as illus-
trated in Figure 4d–4l). When compared to the LM6
tube sample, it was observed that the addition of Ni
changes the eutectic silicon morphology to form primary
Si cuboids.
23
X-Ray diffraction (XRD) analysis of the LM6 tube
showed the presence of two phases such as  -Al and Si;
however, the FGC resulted in three phases such as Al
3
Ni,
Si, and  -Al, as seen in Figure 5. Considering that Ni
and Si are only marginally soluble in the  -Al matrix
(0.05 % and 0.1 % respectively), the generation of Al
3
Ni
tri-aluminide seems logical. It is important to note that
intermetallic Al
3
Ni has a very low solid solubility of Si,
at only 0.5 %.
25
As a result, the interaction between the
aluminium matrix and nickel was clearly the only way
for an intermetallic compound to occur in the melt. It can
also be said that the precipitation of the Al
3
Ni phase de-
pends on the quantity of nickel present in the composite
melt. The stable Al
3
Ni intermetallic phase was estab-
lished in the melt through a ternary eutectic reaction be-
tween the nickel and the Al-Si alloy as mentioned in the
following Equation (2):
26
LLAl+Al
3
Ni + Si (2)
Centrifugal casting can be divided into two distinct
subfields: centrifugal solid-particle method (CSPM) and
centrifugal in-situ method (CISM) based on the target
operating temperature.
27
The solid reinforcements, such
as ceramic particles, are added externally into the melt in
a CSPM. Also in these methods, the dispersed phase re-
mains unchanged both chemically and physically
throughout the process. In a centrifugal in-situ method,
the reinforcements (intermetallics) are precipitating
through chemical reactions inside the melt. The crystall-
isation phenomenon used here is very comparable to that
used in in-situ composite manufacturing. The density
difference between the primary crystals in the melt is re-
sponsible for the creation of compositional gradients
during the CISM of fabricating FGCs.
12
These primary
Al
3
Ni particles and Si cuboids in melt under centrifugal
force are amenable to the aforementioned formation
mechanism. The primary Si particle has a density of
2.33 g/cm
3
, whereas the Al
3
Ni precipitate is reported to
have a density of 4.0 g/cm
3
.
18,28
These values are differ-
ent from the molten aluminium (2.37 g/cm
3
). Gradient
distribution of the primary Al
3
Ni and Si particles inside
the FGC tubes was therefore anticipated.
The centripetal acceleration (a
p
) and radial moving
velocity (V
c
) are responsible for the reinforcing particle
segregation in different zones of the tube in centrifugal
casting. They are expressed in Equations (3) and (4) as
follows:
aN r
p
p
=
− 
 1 22
4π (3)
V
dr
c
p
c
=
−
2
1
2
  
 )
(4)
where N is the mould rotational velocity (s
–1
),  p
is the
reinforcement particle density,  l
is the density of com-
posite melt (kg/m
3
), r is the particle distance from the
mould center (m),  is the angular velocity of mould
(rad/s), d is the diameter of the reinforcement particle
(m) and  c
is the composite melt viscosity (Pa·s). Parti-
cles will migrate to the outer zone of the tube if  p
>  l
and they will flow into the inner zone of the cast tube if
 p
< l
. The primary Si particles in the LM6 tube sample
have a density (2.33 g/cm
3
) that is less than the liquid
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674 667
Figure 5: X-ray diffractogram of LM6 and FGC
C
tube samples

metal (2.37 g/cm
3
) and this density difference pushed
the particles into the inner zone. Primary Al
3
Ni particles
in FGCs have a density of 4.0 g/cm
3
that is greater than
that of the liquid metal, forcing the particles to the outer
zone by centrifugal force.
18
At the same time, primary
Si particles were being drawn inward (inner zone) by
centripetal force in the case of the FGCs. This transpor-
tation occurred due to convection, which was increased
by the considerable temperature difference formed dur-
ing solidification between the outer and inner zones of
the tube.
22,29
From the microstructures of the FGC tubes (Fig-
ure 4d–4l)) it can be concluded that the gradation has
been established in all three tubes. The formation of a
gradient in the thickness of the FGCs was obtained by
four phases. The primary Si cuboids in the melt were
only partially retained in the outer zone of the tube at the
beginning stage due to the cooling behaviour of the
mould. As can be noticed in Figure 4, in the subsequent
phase, numerous primary Si particles were pushed to-
ward the inner zone by centripetal force. The third stage
involved the segregation of the primary Al
3
Ni particles
that have precipitated within the melt. Intermetallic nu-
cleation growth, development and transportation are all
controlled by the temperature disparity established be-
tween the mould and poured liquid metal during solidifi-
cation.
28
Due to their high density, primary Al
3
Ni parti-
cles travelled to the outer zone, where they settled with
the help of centrifugal force. The primary Si particles
were likewise now heading towards a gradual inward mi-
gration at this stage. As the melt viscosity increases, pri-
mary Al
3
Ni particles slow down and settle to a lower ve-
locity in the final phase. Hence, both particles were not
reaching their destinations as expected. They settled
down quickly during the traveling time itself because of
the rapid cooling nature of the centrifugal casting tech-
nique.
Figure 6 shows that the LM6 tube has primary Si cu-
boids traceable throughout the whole structure. The per-
centage of primary Si particles within the overall volume
was calculated to be between 4.6 % and 6.0 %. Due to its
eutectic composition, LM6 alloy had very few primary
Si particles. The least volume fraction of primary Si
(4.6  /%) was present in the outer zone; the inner and
transition zones have far better segregation. These find-
ings are in accordance with the respective microstruc-
tures of the LM6 tube (Figure 4a–4c).
As illustrated in Figure 4, it became evident that the
primary Al
3
Ni and Si particles had been segregated and
accumulated in the inner and outer zones of all the FGCs
with Ni levels of (3, 6 and 9) w/%. In FGC
A
, the addition
of 3 w/% of Ni promotes the formation of primary Al
3
Ni
particles to a value of 28  /% in the outer zone and pri-
mary Si particles to a value of 9.9 /% in the inner zone.
The same trend has been observed in the other two FGC
tubes. The largest volume fractions of the two kinds of
reinforcements were noted in the FGC
C
tube sample as
13  /% of primary Si at the inner zone and 36.5  /% of
primary Al
3
Ni at the outer zone. As shown in Figure 6b,
the volume fraction of primary Al
3
Ni particles in the
transition zones of FGCs varied from (17 to 28) /%, as-
sociated with primary Si particles. When FGCs were
compared to the LM6 tube, the segregation of primary Si
particles increased by 39.4 % in FGC
A
, 51.6 % in FGC
B
,
and 53.9 % in FGC
C
in their inner zone.
It can be summarised that two kinds of primary parti-
cles, such as primary Si and Al
3
Ni, were plentiful in the
inner and outer zones of the gradient composite tubes,
respectively. An amount of eutectic Al
3
Ni and Si existed
in all the zones. Furthermore, eutectic Si flakes were
found in greater quantity in the LM6 tube and decreased
as the nickel concentration climbed. Because of the cool-
ing action of the mould, very limited primary Si cuboids
remained within the outer zone, whereas the fewest pri-
mary Al
3
Ni particles settled in the inner zone owing to
the rise in the viscosity of the melt. Primary Si and Al
3
Ni
particulates exhibited an evident gradient pattern along
the radial direction of the FGC tubes from the inner to
the outer zone. In centrifugal casting, a mould wall is a
good place for beginning the solidification at a higher
cooling rate, so directional solidification occurred from
the wall to the mould center during the casting process.
It is also believed that this situation improved the feeding
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
668 Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674
Figure 6: a) Volume fraction of primary Si particles, b) volume fraction of Al
3
Ni particles at different zones of the cast tubes

potential of interdendritic regions and decreased the for-
mation of porosity in the LM6 and FGC tubes.
3.2 Hardness
The findings from the hardness assessments of both
LM6 and FGC tubes at five spots in the thickness are
presented in Figure 7. The hardness of the FGC tubes in-
creased compared to LM6 alloy tubes, owing mostly to
precipitating hard in-situ Al
3
Ni reinforcements (841HV).
Figure 7 shows a declining trend in hardness measured
from the outside edge of the FGC tube samples. In con-
trast, the outer-to-inner edge hardness assessment in an
LM6 tube sample revealed a distinct pattern. The peak
hardness of this sample was 83 HV at its inner edge
(considered as inner zone) and 76 HV at its outer edge
(considered as outer zone). This minor variation between
these two zones was associated with a deficit of primary
Si cuboids, according to microstructural investigations of
the LM6 tube. It was evident that the method in which
the reinforcement phases segregate during centrifugal
casting of composite melt has a direct impact on the
hardness values.
29
The hardness readings for FGC tube
samples followed the same trend, with hardness decreas-
ing towards the inner edge across the radial thickness.
Similar to the microstructural observations, the reported
hardness values demonstrated proper gradient as a result
of the segregating manner of both primary Al
3
Ni blocky
particles and primary Si cuboids over the radial
cross-section. LM6 tube had an outside edge hardness
value of 76 HV, which elevated by 15.55 %, 23.23 %,
and 35.75 % in the FGC
A
, FGC
B,
and FGC
C
, respectively.
The hardness improved when intermetallic Al
3
Ni precip-
itated in a eutectic Al-Si melt.
21
The observed increase in
hardness near the outside edge of FGC tubes was mostly
attributable to tri-aluminide segregation and nucleation.
The FGC
C
tube sample had the highest hardness val-
ues tested from the outside edge to the inner edge, with
the outer edge having the maximum hardness value of
118.30HV. This enhancement was connected to its
microstructure and reinforcements segregation in such a
manner that the presence of a substantial amount of
in-situ Al
3
Ni precipitates owing to the action of centrifu-
gal force and accelerated precipitate growth rate in the
composite melt. It was recognized that the rate of devel-
opment of in-situ Al
3
Ni nuclei and the temperature gra-
dient induced were predominantly accountable for the
increase in hardness in a eutectic Al-Si alloy.
30
Due to a
temperature disparity (about 520 °C) between the com-
posite melt and the GCI mould, a long time started to
emerge during solidification, which promoted inter-
metallic Al
3
Ni nucleation. Furthermore, when adding Ni
in the range of (3 to 9) w/%, the inner edges of fabricated
FGC samples having improved primary Si settlements
showed hardness improvements in the order of 1.2 %,
3.5 %, and 16 % when compared to LM6 tube (inner
edge = 83 HV). In comparison, it was obvious that the
presented hardness values were closely related to the vol-
ume proportions of both in-situ Al
3
Ni and primary Si
precipitates (Figures 6).
3.3 Tensile Characteristics
The obtained tensile properties of both the LM6 tube
and FGCs at three different zones are presented in Fig-
ure 8. All of the FGC tube specimens demonstrated
higher tensile strength than the LM6 tube in all three
zones. The microstructural observations suggested that
the UTS improvement was due to the precipitation and
strengthening influence of the Al
3
Ni intermetallic phase
in the FGC tubes.
10
They were wetted with the molten
Al-Si alloy, resulting in good and persistent bonding be-
tween particles and matrix. This reliable interface was
presumed to effectively transmit and disperse load from
the matrix to the reinforcement. Generally, the tensile
characteristics of the in-situ composites based on Al-Si
alloy are dictated by microstructural features such as
 -Al matrix, precipitated intermetallics, Si phase, and
microporosities.
31
Owing to the existence of more pri-
mary Al
3
Ni dispersoids that had greater bonding with the
matrix, the outer zones of FGCs exhibited elevated ten-
sile strength. The young’s modulus and UTS of Al
3
Ni
tri-aluminide have been reported to be 116 GPa to
152 GPa and 2160 MPa, respectively, which are much
higher than Al-Si matrix (E = 70 GPa &
U T S=3 7M P a ) .
11
Inspite of these high properties, the
load acting was transferred to the intermetallic particles,
which performed as load-bearing structures to postpone
the fracture during the tensile test.
As indicated in Figure 8a, the FGC
C
tube outer zone
specimen with the highest proportion of Al
3
Ni (31  /%)
had a maximum tensile strength of 203 MPa. This could
possibly be viewed as a result of the action of hard Al
3
Ni
reinforcing particles, which impeded the plastic flow of
the aluminium matrix. In contrast, the LM6 tube pre-
sented only 137 MPa in the same zone, which is 32.5 %
less than FGC
C
. However, the nickel-added FGCs
showed a reduction in the percentage of elongation to
rupture compared with the LM6 tube, as shown in Fig-
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674 669
Figure 7: Variations in hardness from the outer region along the radial
thickness of centrifugally cast tubes

ure 8b. It was noticed that when the nickel content of the
composite melt increases, the elongation decreases sig-
nificantly.
10
This might be owing to the segregation of
hard intermetallic Al
3
Ni in the Al-Si grain boundaries.
The UTS at the inner zone of the LM6 tube specimen
was 130 MPa, which was greatly improved in the case of
all the FGC tube specimens due to the addition of nickel.
It has also been found that the fracture propensity of an
Al-Si alloy increases if the microstructure consists of
coarser eutectic Si flakes than cuboidal primary Si.
32
It
was evident from the microstructures of the FGCs that
all three zones contain eutectic phases of Si and Al
3
Ni.
These hard eutectic phases were believed to form an in-
terconnected network that significantly enhanced the
UTS of the FGCs by improving the load-transfer capac-
ity.
33
As seen in microstructures (Figure 4), all the FGC
tubes had more primary Si particles and fewer primary
Al
3
Ni particles in their inner zones. Furthermore, these
less in-situ Al
3
Ni intermetallics were identified as inca-
pable to withstand much-pulling force. Figure 8a shows
that the tensile strength increased at the inner zone from
130 MPa for the LM6 tube to 189 MPa for FGC
C
. The
elongation, on the other hand, reduced from 5.7 % to
2.6 %. The substantial increase in UTS at the inner zone
of FGCs should be attributable to the existence of nu-
merous primary Si particles accompanied by a few
intermetallic particles, which is not the case in the LM6
tube. The LM6 tube specimen only comprised more
eutectic silicon flakes with very few primary silicon par-
ticles, as shown in Figure 4.
The quality index (Q) has been used to evaluate the
overall tensile properties of any alloy, and it is thought to
be a better representation of the actual tensile property
compared to either tensile strength or elongation alone.
34
It is commonly computed using Equation 6.
Q = UTS (MPa) + 150 × log(%E) (6)
As can be stated from Figure 8c that the incorpora-
tion of nickel into eutectic Al-Si alloy improves the Q
greatly and hence Q values of all FGC tube specimens
were greater than the LM6 tube in all three zones. When
FGC tubes were chilled from the process temperature,
variation in the Coefficient of thermal expansion (CTE)
between the Al-Si matrix (24 × 10
–6
/K) and in-situ Al
3
Ni
(14.3 × 10
–6
/K) causes dislocations to occur. In addition,
such dislocations are produced as a result of the exis-
tence of Al
3
Ni intermetallics within the matrix, obstruct-
ing plastic deformation and contributing to the better ten-
sile characteristics of the FGCs. As expected, the highest
quality index was achieved in the FGC
C
outer zone, be-
cause of the abundant presence of primary Al
3
Ni precipi-
tates and the notable presence of primary Si. The lowest
Q computed for the transition zone of the LM6 tube was
242 MPa, which is 10.29 % lower than the transition
zone of FGC
B
. However, the presence of more in-situ
Al
3
Ni and Si phases greatly reduced the -Al phase, low-
ering the ductility of the FGCs.
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
670 Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674
Figure 8: Tensile properties of centrifugally cast tubes in different zones: a) UTS, b) Elongation (%) and c) Quality Index (Q)

3.4 Tensile Fractography
The fractured surfaces of the LM6 and FGC tubes
were analyzed perpendicular to the tensile axis. SEM im-
ages of The LM6 tube fractured specimens were display-
ing a clear brittle fracture at all three zones, as shown in
Figure 9. However, because of the existence of primary
Si particles, some ductile dimples were also observed.
During the examination of the fractured surfaces, all
three zones of the LM6 tube showed many cleavage fac-
ets and tear ridges. These cleavage facets were formed
by the fracture of Si particles, whereas the tear ridges
were created by plastic deformation and fracturing of the
aluminium matrix. The glowing traces surrounding the
tear ridges clearly revealed the brittle cracking of the Si
particles. The inner zone of this LM6 tube is composed
of comparatively more primary silicon particles, which
are responsible for the decrement in tensile strength than
the other two zones. Because of the stress concentration,
primary Si particles were observed to crack and debond.
The typical ductile dimples were also shown in the inner
zone. The formation of void nuclei was reported to occur
predominantly at the interfaces between the primary Si
particles and the matrix, wherein stress concentrations
were higher.
35
Two kinds of explanations were attained through
microstructural and fractographic evidence of LM6 tube:
a) small dendrites and silicon particles showed less dam-
age rate since they required a huge strain to attain the
critical degree of disruption for causing damage and b)
elongated Si phase with a long length contained coarser
microstructures were more likely to crack at low loads,
reducing the ductility. The presence of eutectic Si phase
in all three zones across the LM6 tube caused fractures
to develop quickly in the aluminium matrix, attributable
to eutectic Si fragmentation. These cracks became the
sources of stress, forming a network between cracks and
causing an uncontrollable fracture in the specimen, as
shown in Figure 9b.
16
The fracture mechanism associ-
ated with the LM6 tube specimens can be understood in
the following ways. The imposed tensile strain caused
significant plastic deformation around the Si particles in
the aluminium matrix, resulting in Al-Si debonding, tear
ridges, and the creation of microvoids at the interface.
The microvoids linked up with each other at the Al-Si in-
terface during tensile stress and constitute a microscopic
crack. Subsequently, these microcracks became unstable,
causing the ultimate failure. Also, in the eutectic, cracks
at the interfaces among the aluminium matrix and Si
were regularly found to develop; however, some frac-
tured Si particles were also identified.
Figure 10 depicts the SEM fractography of the FGC
C
tube at its three zones, which have the highest addition of
nickel. The fractured surfaces of this specimen showed a
mixed mode of ductile dimples and cleavages of brittle
fracture. In particular, brittle fracture was observed in the
outer zones of all three FGCs with very few dimples.
The other two zones presented a mixed mode of brit-
tle-ductile fracture. It was identified that the increase of
nickel directly improved the nucleation of in-situ Al
3
Ni
phase from microstructural observation, which promoted
more brittle cleavages across the radial cross-section of
tubes. As can be seen in the outer zone of the FGC
C
tube
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674 671
Figure 9: SEM fractographs of LM6 tube zones (a, b & c), d) EDS pattern of  -Al phase and e) Si phase

specimen (Figure 10a), more brittle fracture arose,
which presented low ductility. Fractography of FGC
C
tube specimen has the following failure characteristics: i)
interfacial debonding of both primary Al
3
Ni and Si parti-
cles from aluminium matrix (intergranular – brittle), ii)
crack propagation through Si particles (transgranular –
brittle) and iii) torn out or decohered Si particles (duc-
tile).
Debonding of the primary Si particles occurred dur-
ing the first stage of tensile testing owing to plastic de-
formation in the  -Al matrix generated by a minor ten-
sile force applied. The crack formed through this
debonding tried to develop a severe fracture in the FGCs.
However, this crack flow was restricted by the existence
of in-situ Al
3
Ni in interdendritic areas of the FGCs. This
resistance to crack flow offered by the intermetallic was
figured to be the key factor in the enhancement of tensile
strength. With regard to the inner zones of the FGCs,
more primary Si debonding failure was observed (Figure
10c). The scarcity of Al
3
Ni tri-aluminide and availability
of more -Al matrix caused more severe crack formation
and growth in the inner zones of FGCs than in the other
two zones. Furthermore, because of very limited stability
of the  -Al matrix, the constant tensile stress induced
dislocations. These dislocations stacked up at the inter-
face ( -Al/Si) and triggered high stress concentrations,
which crushed the silicon particles and generated a
crack- nucleation site.
36,37
The presence of dimples on the fracture surfaces of
FGCs can clearly identify it (Figure 10b). As mentioned
above, these cracks were not moving through the Al
3
Ni
intermetallic particles, because there was no evidence of
broken intermetallics on the fracture surfaces. Instead,
these cracks moved along the boundaries, resulting in in-
terfacial debonding between the Al
3
Ni and the matrix.
20
According to Griffith’s theory, the particle is estimated to
be crushed only when the stress surpasses the Griffith
threshold. The stress required to fracture a particle is ex-
pressed in Equation (7) as follows,
 c
p c
p
=
k
d
(7)
Where k
c
p
is the fracture toughness and d is the diam-
eter of the reinforcement particle.
38
Because no
transgranular crack was observed in the Al
3
Ni phase, it
can be inferred that there was no superficial interface be-
tween the matrix and intermetallic. Generally, if the ma-
trix remains smooth, the interface becomes weak and
there is decoherence. EDS analysis corroborated the in-
clusion of in-situ Al
3
Ni in the interdendritic vicinity, as
shown in Figure 10d. A reinforcement particle fracture
is distinctly related to its fracture toughness; hence it was
evident that the FGCs included Al
3
Ni intermetallics with
excellent fracture toughness. The fracture initiation at the
inner zone was caused by the primary Si particles and in
the case of the remaining two zones, it was almost
eutectic Si. In all of the FGCs, the brittle Si particles
A. SAIYATHIBRAHIM et al.: INFLUENCE OF NICKEL ON THE MICROSTRUCTURAL EVOLUTION AND ...
672 Materiali in tehnologije / Materials and technology 57 (2023) 6, 663–674
Figure 10: SEM fractographs of FGC
C
tube zones (a,b&c)andEDSpattern of Al
3
Ni phase (d)

were torn out (decohered) from the  -Al matrix in these
flat areas (Figure 10a), having left a smooth facet ter-
race. Despite the fracture of the brittle Si phase, these
cleavage facets were more likely to form.
39
It should be
recorded that the inner zone had more elongation than
the outer zone. The rising concentrations of brittle Al
3
Ni
phase represented the decrease in the percentage of elon-
gation to rupture from FGC
A
to FGC
C
, which is mainly
accountable for the 3.1 % drop in elongation at all zones
while attempting to compare LM6 with FGC
C
tubes.
4 CONCLUSIONS
Four kinds of defect-free functionally graded com-
posite (FGC) tubes (LM6 tube, FGC
A
, FGC
B
, and FGC
C
)
were manufactured by horizontal centrifugal casting,
with nickel inclusion varying from 0 to 9 w/%. All three
FGC tubes demonstrated more substantial gradation than
the LM6 tube, which had very low gradation owing to
the absence of nickel addition and the eutectic nature.
The microstructures, hardness, and tensile properties of
the fabricated tubes were comprehensively investigated
zone by zone (inner, transition, and outer zone). This
study discovered that adding nickel to the LM6 alloy
generated in-situ Al
3
Ni, which inhibited the growth of
eutectic Si flakes and nucleated cuboidal primary Si.
Due to the notable segregation of primary Si cuboids
in the inner zone, zone-wise characterization of the LM6
tube revealed only a marginal gradient structure across
the radial thickness. Tensile fractography revealed the
brittle fracture, which was induced by the abundant
availability of the elongated eutectic Si phase. FGC tubes
exhibited three distinct microstructural features in their
radial cross-section due to the segregation of primary Si
and in-situ Al
3
Ni during solidification related to their
density difference. A richly in-situ Al
3
Ni accumulated
outer zone, a more cuboidal Si-deposited inner zone and
a mixed transition zone were observed and this gradient
structure influenced the properties.
It was also deduced that increasing the amount of
nickel in the composite melt increases the gradation, re-
sulting in distinct property differences. In comparison to
the LM6 tube, the hardness of the FGC tubes improved
gradually across all three zones. However, nickel-added
FGCs showed a reduction in the percentage of elonga-
tion to rupture, and fracture analysis indicated a mixed
mode of ductile dimples and brittle fracture cleavages.
It appears that different qualities may be tailored in
LM6 alloy-based FGC in this approach by promoting
various in-situ reinforcements, and greater attention is to
be provided in the future to examine other properties of
the manufactured FGC tubes.
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