Microstructure and Microhardness Evolution of Mg–8Al–1Zn Magnesium Alloy Processed by Differential Speed Rolling at Elevated Temperatures

Abstract

Mg–8Al–1Zn magnesium alloy was successfully processed using deferential speed rolling (DSR) at temperatures of 400 and 450 °C for thickness reduction of 30, 50, and 70% with no significant grain growth and dynamic recrystallization. Using optical microscopy (OM), scanning electron microscopy (SEM), and transmission electron microscopy (TEM), the rolled microstructures were examined. Although the results indicate a slight reduction in grain size from the initial condition, the DSR processing of alloy at an elevated temperature was associated with a significant number of twins and a distribution of the fine particles of the second phase. The strength in terms of microhardness measurements and strain hardening in terms of shear punch testing was significantly improved in the rolled microstructure at room temperature. The existence of twins and widely distributed second-phase fine particles at twin boundaries reflected positively on the extent of the elongations in terms of shear displacements when microstructures were tested at elevated temperatures in the shear punch testing.

Full text

Citation: Alsubaie, S.A.; Al-Zubaydi,
A.S.J.; Hussein, E.A.; Alawadhi, M.Y.
Microstructure and Microhardness
Evolution of Mg–8Al–1Zn
Magnesium Alloy Processed by
Differential Speed Rolling at Elevated
Temperatures. Materials 2024,17, 4072.
https://doi.org/10.3390/ma17164072
Academic Editor: Liyuan Sheng
Received: 13 July 2024
Revised: 5 August 2024
Accepted: 7 August 2024
Published: 16 August 2024
Copyright: © 2024 by the authors.
Licensee MDPI, Basel, Switzerland.
This article is an open access article
distributed under the terms and
conditions of the Creative Commons
Attribution (CC BY) license (https://
creativecommons.org/licenses/by/
4.0/).
materials
Article
Microstructure and Microhardness Evolution of Mg–8Al–1Zn
Magnesium Alloy Processed by Differential Speed Rolling at
Elevated Temperatures
Saad A. Alsubaie 1, *, Ahmed S. J. Al-Zubaydi 2,3, Emad A. Hussein 4and Meshal Y. Alawadhi 1
1Department of Manufacturing Engineering Technology, College of Technological Studies, The Public
Authority of Applied Education and Training (PAAET), P.O. Box 42325, Shuwaikh 70654, Kuwait;
my.alawadhi@paaet.edu.kw
2School of Applied Sciences, University of Technology-Iraq, Baghdad 10001, Iraq;
ahmed.s.alzubaydi@uotechnology.edu.iq
3School of Engineering, Faculty of Engineering and Physical Sciences, University of Southampton,
Southampton SO17 1BJ, UK
4School of Production and Metallurgy Engineering, University of Technology-Iraq, Baghdad 10001, Iraq;
emad.a.hussein@uotechnology.edu.iq
*Correspondence: sa.alsubaie@paaet.edu.kw
Abstract: Mg–8Al–1Zn magnesium alloy was successfully processed using deferential speed rolling
(DSR) at temperatures of 400 and 450
◦
C for thickness reduction of 30, 50, and 70% with no significant
grain growth and dynamic recrystallization. Using optical microscopy (OM), scanning electron
microscopy (SEM), and transmission electron microscopy (TEM), the rolled microstructures were
examined. Although the results indicate a slight reduction in grain size from the initial condition, the
DSR processing of alloy at an elevated temperature was associated with a significant number of twins
and a distribution of the fine particles of the second phase. The strength in terms of microhardness
measurements and strain hardening in terms of shear punch testing was significantly improved
in the rolled microstructure at room temperature. The existence of twins and widely distributed
second-phase fine particles at twin boundaries reflected positively on the extent of the elongations in
terms of shear displacements when microstructures were tested at elevated temperatures in the shear
punch testing.
Keywords: differential speed rolling; magnesium alloy; shear punch test
1. Introduction
In automotive applications, magnesium alloys are promising alternative lightweight
metal alloys to replace denser ones like steel and aluminum due to their low density
(
1.74 gm/cm3
) and their high specific strength (158 KN
·
m/kg). However, magnesium has
an HCP structure, which illustrates poor workability and low ductility at room temperature
due to its limited slip systems [
1
]. In HCP crystal structure materials, deformation at room
temperature can be achieved by two dominant mechanisms: slip in the basal plane and/or
mechanical twinning. To achieve homogenous deformation without cracks developing,
five independent slip systems are required, according to von Mises’ criteria. Since materials
that have HCP structure (such as magnesium alloys) have three slip systems in the basal
plane (two of which are independent), this limited slip system resulted in brittleness and
lack of ductility at room temperature, and these materials are considered to be hard-to-
work materials [
1
,
2
]. Extruded Mg alloys have attracted attention because of their better
mechanical properties than the cast Mg alloys [
3
]. The purpose of increasing the strength
of Mg alloys is to compete with lightweight metal alloys and to expand its application
in modern industry. For this to be achieved, further effort must be made. Even though
the two popular techniques of severe plastic deformation (SPD), equal-channel angular
Materials 2024,17, 4072. https://doi.org/10.3390/ma17164072 https://www.mdpi.com/journal/materials

Materials 2024,17, 4072 2 of 17
pressing (ECAP) [
4
] and high-pressure torsion (HPT) [
5
], have become important in recent
years because of their potential to produce bulk nanostructured materials [
6
], the previous
processes are still on the laboratory scale and did not succeed in producing a large bulk
material. Other techniques, such as rolling, have been developed to achieve grain size
reduction by inducing intense plastic straining in a larger dimension material.
The differential speed rolling (DSR) process, which is one of the rolling subdivision
types, can introduce a plastic strain at high values to metallic materials in comparison to
traditional rolling within the cross-section of the deformed workpiece. This can be achieved
by using two identical rotated rolls but at dissimilar speeds. The DSR process imposes
shear and compression deformations simultaneously, leading to significant microstructural
and properties alterations in the processed materials. The shear deformation is introduced
by virtue of the differential speeds of rollers, whereas the compressive deformation is
introduced by a reduction in the thickness of the processed material [
7
,
8
]. Advantages of
the DSR processing of metallic materials are enhancement of both strength and ductility via
bimodal grain refinement, applicability of large-scale parts, utilization of a lower number
of rolling passes, and efficient deformity in comparison with conventional rolling by virtue
of differential speeds and subsequent imposed shear deformation, leading to lower power
consumption by the DSR facility in comparison with conventional rollers [
9
,
10
]. Some
limitations come with the DSR process, such as increased friction and wear in the roller
equipment due to the effect of differential speeds and the non-uniformity of deformation
in the processed materials across thickness cross-section, leading to heterogeneity in the
microstructure and properties [
11
,
12
]. However, this process can sometimes be considered
continuous SPD processing due to its ability to improve strength and ductility at large-
dimension manufacturing of worked pieces in comparison with HPT and ECAP [13,14].
The DSR process was used for the development of various metallic alloys based on
Mg [
15
], Al [
16
], and Cu [
17
]. The outcomes in these aforementioned studies revealed an
improvement in the mechanical properties of the deformed alloys due to intense imposed
shear deformation and, consequently, grain refinement. Also, it was found that the rolling
temperature and thickness reduction have an impact on the resultant microstructure and
mechanical properties [
18
]. Wrought alloys are considered to display the lowest workability
of magnesium alloys at room temperature yet have the greatest strength. In this category,
the magnesium–aluminum system is the most important in alloys such as AZ31, AZ61,
Mg–8Al–1Zn
, and AZ91 [
19
–
22
]. It is well known that increasing the Al content in Mg
alloys will increase the strength of the alloys. Although Mg–8Al–1Zn is a widely known
magnesium alloy, there is still a lack of information on it as a material processed by
different SPD processes [
23
]. In the current research, the DSR process will be employed to
improve the properties of Mg–8Al–1Zn alloy sheets by examining the effect of processing
temperatures on microstructural evolution and mechanical properties of the alloy, such as
microhardness and shear punch properties of Mg–8Al–1Zn alloy samples.
2. Experimental Material and Procedure
The material used in this study was a commercial as-cast Mg–8Al–1Zn magnesium
alloy (Mg–8.11% Al–0.45% Zn). This alloy came in the form of an extruded rod of 10 mm in
diameter and 200 mm in length. The extrusion of the as-cast alloy was achieved at
300 ◦C
using an extrusion ratio of 12; then, the resultant rod was solution heat treated at 400
◦
C
for 12 h under an argon atmosphere followed by water quenching. This extruded alloy
bar was cut into sheets for DSR processing with dimensions of 60 mm in length, 8 mm in
width, and 2 mm in thickness. The average grain size and Vickers microhardness in the
as-received extruded alloy bar were about 25
µ
m and 65 Hv, respectively. The sheet samples
of Mg–8Al–1Zn alloy were rolled in DSR at 400 and 450
±
10
◦
C, at reduction ratios of 30,
50, and 70%, corresponding to 1 pass, 2 passes, and 4 passes for each reduction ratio. Before
the DSR process, the sheet samples were kept at the required DSR temperature for
5 min
in the furnace and then rolled at the same temperature in order to keep the processing
temperature during the DSR at the same level. The diameter ratio between the upper and

Materials 2024,17, 4072 3 of 17
lower rollers was 1.15, with rotation speeds of 2.5 m/min and 3 m/min for the upper and
lower rollers. The DSR process was conducted along the length dimension of the samples
(rolling direction, RD), which is aligned to the extrusion direction (ED) of the alloy bar. The
final thickness was 1.4 mm by one pass (thickness reduction 10%), where each sample was
rotated by 180
◦
around the rolling direction (RD) between each two successive passes, as
represented in Figure 1[
7
]. The samples were reheated between each of the two successive
passes, and the temperature of the rollers was maintained at the required DSR temperature
using an embedded heating element inside the rollers, where the rolling was achieved
under no lubrication condition.
Materials 2024, 17, x. hps://doi.org/10.3390/xxxxx www.mdpi.com/journal/materials
Normal Direction
,
ND
Transverse Direction, TD
Rollin
g
Direction
,
RD
180°
Figure 1. Schematic representation of the differential speed rolling process associated with sample
rotation procedure between two successive passes [7].
The Mg–8Al–1Zn sheet samples prior to and post-DSR process were mechanically
ground and polished to a final mirror-like surface, then etched using acetic-glycol solu-
tion, rinsed with ethanol, and then air-dried. The microstructure was examined using
optical microscopy (OM), scanning electron microscopy (SEM), and Transmission electron
microscopy (TEM). X-ray diffraction analysis was also performed on alloy sheet samples
to determine the presence of phases and calculate crystallite size and defect dislocations.
Vicker microhardness measurements were achieved for sheet samples on normal direction
surfaces of the sheets, and data were then plotted. The shear punch test (SPT) was con-
ducted on alloy sheets at testing temperatures of 400 and 450
±
10
◦
C using initial strain
rates of 10
−1
, 10
−2
, and 10
−3
s
−1
. The SPT samples were cut from the longitudinal direction
of the DSR sheets with dimensions of 0.8 mm in thickness and 10 mm in diameter. A
punch fixture with a cylindrical flat punch was used for shear application with a
3.175 mm
diameter and 3.225 mm diameter of the receiving hole, as reported in [
24
]. The applied
load in SPT was measured against the recorded punch displacement. All the data were
acquired by a computer to determine the relation between the shear stress in the tested
sheet samples and to draw load–displacement curves.
3. Results and Discussion
3.1. Microstructural Evolution
The initial microstructures of Mg–8Al–1Zn magnesium alloy prior to DSR processing
are shown in Figure 2, where the average grain size of the microstructure was 25
µ
m, as seen
by OM and SEM observations. These observations showed that the microstructure, prior to
the process, consists of two main phases: the
α
-Mg phase, which forms the grain structure
of the alloy (this phase appears as light and dark grey in OM and SEM micrographs,
respectively) and the
β
-phase, which consists mainly of Mg
17
Al
12
, which decorates the
grain boundaries of the
α
-Mg phase grains as seen in OM and SEM. This phase appears
to have a black-and-white appearance in the aforementioned observations. A closer look
at the microstructure shows that the
β
-phase appears as lamella and coarse particles at

Materials 2024,17, 4072 4 of 17
different areas around the grain boundaries. These phases were identified chemically using
energy-dispersive spectroscopy (EDS) attached to the SEM unit, with the following ratios of
(91.44 wt.% Mg–8.11 wt.% Al–0.45 wt.% Zn) for
α
-Mg phase microstructure and (
55.2 wt.%
Mg–44.3 wt.% Al–0.5 wt.% Zn) for β-phase microstructure.
Materials 2024, 17, x FOR PEER REVIEW 4 of 20
3. Results and Discussion
3.1. Microstructural Evolution
The initial microstructures of Mg–8Al–1Zn magnesium alloy prior to DSR processing
are shown in Figure 2, where the average grain size of the microstructure was 25 µm, as
seen by OM and SEM observations. These observations showed that the microstructure,
prior to the process, consists of two main phases: the α-Mg phase, which forms the grain
structure of the alloy (this phase appears as light and dark grey in OM and SEM micro-
graphs, respectively) and the β-phase, which consists mainly of Mg
17
Al
12
, which decorates
the grain boundaries of the α-Mg phase grains as seen in OM and SEM. This phase ap-
pears to have a black-and-white appearance in the aforementioned observations. A closer
look at the microstructure shows that the β-phase appears as lamella and coarse particles
at different areas around the grain boundaries. These phases were identified chemically
using energy-dispersive spectroscopy (EDS) aached to the SEM unit, with the following
ratios of (91.44 wt.% Mg–8.11 wt.% Al–0.45 wt.% Zn) for α-Mg phase microstructure and
(55.2 wt.% Mg–44.3 wt.% Al–0.5 wt.% Zn) for β-phase microstructure.
Figure 2. The microstructures of the as-received alloy as seen by OM in (a) and SEM in (b,c). The α-
Mg matrix phase and β-phase are in white and dark appearances in (a), whereas these phases appear
in SEM at reverse appearances to OM, as shown in (b). A magnified micrograph shows the mor-
phology of β-phase in (c).
The processing of the alloy by DSR, starting from a thickness reduction ratio of 30%,
50%, and 70%, has been achieved successfully at elevated temperatures without surface
cracking. The DSRolled microstructure showed initial grain refinement down to 15 and
Figure 2. The microstructures of the as-received alloy as seen by OM in (a) and SEM in (b,c). The
α
-Mg matrix phase and
β
-phase are in white and dark appearances in (a), whereas these phases
appear in SEM at reverse appearances to OM, as shown in (b). A magnified micrograph shows the
morphology of β-phase in (c).
The processing of the alloy by DSR, starting from a thickness reduction ratio of
30%, 50%, and 70%, has been achieved successfully at elevated temperatures without
surface cracking. The DSRolled microstructure showed initial grain refinement down to
15 and
20 µm
for the samples rolled at 400 and 450
◦
C, respectively. The alloy at a rolling
temperature of 400
◦
C showed extensive twinning and deformation bands as seen by OM
observations in Figure 3a,b and SEM observations in Figure 4a,b, whereas less twinning
and deformation bands were seen in DSRolled microstructures at a rolling temperature of
450
◦
C as seen by OM observations in Figure 3c,d and SEM observations in Figure 4c,d.
The TEM observations of the DSRolled alloy at a rolling temperature of 400
◦
C are shown
in Figure 5a–d, where the deformation bands of the nanoscale appeared clearly. Moreover,
nanotwins associated with dislocations were also seen in the DSRolled alloys. These
nanotwins were found to be intersected at some areas in the microstructure, resulting in
the fragmentation of the grains down to nanoscales. This fragmentation has resulted in
the formation of sub-grains of nanoscales surrounded by dislocations. A closer look at an
area of dislocations shows that nanoparticles of
β
-phase were found due to fragmentation
by virtue of imposed rolling strain in DSR processing. The
β
-phase in the rolled alloys at

Materials 2024,17, 4072 5 of 17
rolling temperatures of 400 and 450
◦
C has significantly fragmented into nanoparticles,
especially at the lower reduction ratio of 30%. However, at higher reduction ratios up to
70%, this phase disappeared and seems to be dissolved within the
α
-Mg phase, as seen
in SEM observations in Figure 4b,d. The disappearance of the
β
-phase was proportional
to the increase in rolling temperature, as the effective dissolution of this phase happened
at a rolling temperature of 450
◦
C rather than 400
◦
C, where it normally starts to dissolve
at 400
◦
C as reported earlier due to the lower melting point of this phase in this alloy
(460
◦
C). At a rolling temperature of 450
◦
C and at all ratios of reduction, the dissolution
process happened at a faster rate, as seen in SEM observation in Figure 4c,d, which was
associated with relatively grain growth and a small number of
β
-phase particles in the
areas of dissolution [
25
,
26
]. However, the
β
-phase could not dissolve completely as each
rolling pass took no more than one minute, whereas the complete dissolution process
requires a time from 2 to 24 h [
25
–
28
]. In the initial stage of DSR deformation, i.e., at a
reduction ratio of 30%, specifically at a rolling temperature of 400
◦
C, diffusional paths
were provided for the solute atoms in the
β
-phase to diffuse within the
α
-Mg grains via
introducing microstructural defects such as twins, dislocations and fragmentation of coarse
particles which belonged to the
β
-phase. At an advanced stage of deformation, i.e., at a
reduction ratio of 70% and a rolling temperature of 400
◦
C, a balance would appear between
the dissolution and dynamic precipitation of the β-phase [25,29].
Materials 2024, 17, x FOR PEER REVIEW 6 of 20
Figure 3. The DSRolled microstructures as seen by OM at a (a) rolling temperature of 400 °C for 30%
of thickness reduction, (b) rolling temperature of 400 °C for 70% of thickness reduction, (c) rolling
temperature of 450 °C for 30% of thickness reduction, and (d) rolling temperature of 450 °C for 70%
of thickness reduction.
Figure 3. The DSRolled microstructures as seen by OM at a (a) rolling temperature of 400
◦
C for 30%
of thickness reduction, (b) rolling temperature of 400
◦
C for 70% of thickness reduction, (c) rolling
temperature of 450 ◦C for 30% of thickness reduction, and (d) rolling temperature of 450 ◦C for 70%
of thickness reduction.

Materials 2024,17, 4072 6 of 17
Materials 2024, 17, x FOR PEER REVIEW 7 of 20
Figure 4. The DSRolled microstructures as seen by SEM at a (a) rolling temperature of 400 °C for
30% of thickness reduction, (b) rolling temperature of 400 °C for 70% of thickness reduction, (c)
rolling temperature of 450 °C for 30% of thickness reduction, and (d) rolling temperature of 450 °C
for 70% of thickness reduction.
Figure 4. The DSRolled microstructures as seen by SEM at a (a) rolling temperature of 400
◦
C for 30%
of thickness reduction, (b) rolling temperature of 400
◦
C for 70% of thickness reduction, (c) rolling
temperature of 450 ◦C for 30% of thickness reduction, and (d) rolling temperature of 450 ◦C for 70%
of thickness reduction.
The XRD patterns for the initial and rolled alloys are represented in Figure 6a, where
the rolled alloys showed a clear tendency for twin generation as observed by the twin
planes of (10
1
1), (10
1
2), and (10
1
3) for
α
-Mg phase. This tendency for twinning generation
in the
α
-Mg phase was higher in the rolled alloy at a rolling temperature of 400
◦
C rather
than 450
◦
C, with the disappearance of the
β
-phase at these rolling temperatures as shown
by XRD patterns. A reduction in the crystallite size was observed in the rolled alloys
at both rolling temperatures in comparison with the as-received alloy, as represented in
Figure 6b. However, the value of the crystallite size was slightly higher in the alloy rolled
at a rolling temperature of 450
◦
C rather than 400
◦
C. The reduction in the crystallite size
was associated with a slight increase in the defects in the rolled alloys at both rolling
temperatures in comparison with the as-received alloy. The measurements of the grain and
crystallite sizes were consistent for both as-received and rolled alloys.

Materials 2024,17, 4072 7 of 17
Materials 2024, 17, x FOR PEER REVIEW 8 of 20
Figure 5. (a) Deformation bands, (b) nanotwins with dislocations, (c) intersected twins with dislo-
cations, and (d) sub-grain formation surrounded with dislocations in the DSRolled alloy at a rolling
temperature of 400 °C.
Figure 5. (a) Deformation bands, (b) nanotwins with dislocations, (c) intersected twins with disloca-
tions, and (d) sub-grain formation surrounded with dislocations in the DSRolled alloy at a rolling
temperature of 400 ◦C.
The materials with hexagonal-close packed crystal structures, such as magnesium and
titanium alloys, have a limited number of independent slip systems, which makes twinning
a vital deformation mechanism at high imposed strains and elevated temperatures [
20
].
Normally, the magnesium would be crystallographically orientated towards the basal
system before the deformation processing by rolling because the system requires the lowest
critical resolved shear stress in comparison to the counter values for non-basal systems,
such as the twin pyramidal system. However, at high imposed strains and/or elevated
temperatures, the critical resolved shear stress of the twin pyramidal system decreases.
This will lead to the activation of additional systems to accommodate the deformation at
the aforementioned conditions. The activation of pyramidal deformation systems can be
noticed with the increase in processing temperature above 225
◦
C [
30
,
31
]. In the current
research, the rolling temperatures were 400 and 450
◦
C, which induced effectively under
DSR equivalent straining (see Figure 6e) to activate the considerable pyramidal systems
as seen by twinning that appeared in OM, SEM and via XRD peaks and represented by

Materials 2024,17, 4072 8 of 17
twin distribution in Figure 7. The DSR equivalent strain (
ε
) was calculated according to the
following equation [12]:
ε=rε2
r+γ2
3=s[2
√3·ln1
1−r]2+1
3[1
(h0+h)·R·cos−1(1−h0−h
2R)(1−v2
v1
)]2,
where
εr
and
γ
represent the compression and shear deformation components, respectively;
r
represents the reduction in thickness and calculated based on initial thickness
(h0)
and
final thickness
(h)
as
r=
1
−h
h0
,
R
represents the ratio of roll radii, and
v2
and
v1
are the
high and low speeds of rollers, respectively. It is noted that this equation does not directly
imply the effect of rolling temperature and generated temperature due to friction between
the sample and rolls in the calculation of the imposed strain in DSR processing. However,
the DSR at a relatively low-speed ratio (<1) would not impose any significant increase in
the temperature, whereas at a high-speed ratio (>2), the increase in the temperature would
be considerable [32].
The existence of compressive load components during straining processing via DSR
processing also contributed to the activation of these deformation planes. As the thickness
reduction per each pass was maintained at about 10%, preceding the deformation, the
activation of contraction twinning will be dominated rather than extension twinning to
accommodate the deformation homogeneity at higher strain levels and temperatures. This
can be attributed to the decrease in the sample thickness, which leads to a decrease in the
angle between the normal direction (ND) and the radial force at the edge of the rollers. The
radial force has two components: the radial force in the normal direction (ND) and the
other in the rolling direction (RD). With further thinning in the sample thickness, radial
force in the normal direction will be higher than in the rolling direction, leading to the
reorientation of grain structure in the Mg-based alloys towards the activation of additional
deformation twinning systems at elevated temperatures [
33
,
34
]. It seems that the grain
structure in the magnesium alloy that rolled at 400 and 450
◦
C was refined extensively
by twinning fragmentation, where the twins were generated at these temperatures and
then intersected under rolling strain deformation. Due to the limited active slip systems
in magnesium and its alloys at ambient temperatures, thus the expected deformation to
occur in these materials would be via the activation of additional deformation systems.
The activation of such systems, such as pyramidal twining systems, requires elevated
temperatures, as in the case of deformation conditions in the current investigation. As a
result of this extensive activity of twin deformation, deformation bands and shear bands
have formed here to dominate the deformation mechanisms under such rolling conditions
where no dynamic recrystallization has appeared but grain growth instead that was resisted
by the twin formation and intersection here [
30
,
35
]. It can be seen from the microstructural
observation that the deformation bands have extended to form shear bands to generate
a homogenous plastic deformation due to the imposed strain via DSR processing. The
generated shear bands were seen over many grains in the rolled microstructures despite
the occurrence of grain growth. These bands normally come with intense shear regions,
leading to the deformation of localized regions associated with localized dislocation areas
and second-phase fragmentation, as seen by TEM observations, where the twinning cannot
accommodate the intense imposed straining with further deformation. The occurrence of
shear bands appears to be in the orientation of propagation of twinning or intersecting
them [
36
,
37
]. The shear bands were also found to generate via double twin areas in the
rolled alloy, where these areas were considered as nucleation sources for the generation
of the sheared material to accommodate the localized deformation at these regions and
keep the deformation continuity under DSR straining rather than grain growth at such
elevated temperatures of rolling processing [
33
,
34
]. Normally, the existence of shear bands
during rolling processing at temperatures below 300
◦
C is considered as nucleation sites
for dynamic recrystallized misstructure [
38
,
39
]. However, in the current investigation,
the rolling temperatures were 400 and 450
◦
C, where no dynamic recrystallization was

Materials 2024,17, 4072 9 of 17
noticed, but the development of subgrains, as seen by TEM observation, were formed
instead at such rolling temperatures. Thus, the domination deformation mechanism was
not only the slip but mainly the twinning over the orientations of planes of (10
1
2) and
(10
1
3) twinning as indicated by XRD data [
30
,
40
]. It should be noted that the relatively
initial grain structure in the current alloy with a high amount of aluminum content would
also support the activation of twinning in magnesium alloys. Thus, the grain refinement
via twinning segmentation of initial grains in the rolled alloy has facilitated the plastic flow
of the alloy under rolling conditions despite the localized regions of deformation due to
shear banding [30,41].
Materials 2024, 17, x FOR PEER REVIEW 9 of 20
(a)
(b)
(c)
30 40 50 60 70
Intensity a.u.
2θ º
α-Mg (1010)
α-Mg (0002)
α-Mg (1011)
α-Mg (1012)
α-Mg (1020)
α-Mg (1013)
α-Mg (2020)
β-phase (411)
β-phase (322)
As−received
Hot rolled
450 °C
Hot rolled
400 °C
As–received Hot rolled – 400 ºC Hot rolled – 450 ºC
10
20
30
40
50
60
Grain size (μm)
Crystallite size (nm)
Dislocation density (×10
10
m
−2
)
Average microhardrness (Hv)
Sample
50
100
150
200
250
300
0.0
0.5
1.0
1.5
2.0
2.5
3.0
55
60
65
70
75
80
85
−30 −20 −10 0 10 20 30
0
20
40
60
80
100
Vickers microhardness (Hv)
Distance (cm)
Hot rolled – 400 ºC – 4 passes (RE:70%)
Hot rolled – 450 ºC – 4 passes (RE:70%)
As–received
Figure 6. Cont.

Materials 2024,17, 4072 10 of 17
Materials 2024, 17, x FOR PEER REVIEW 10 of 20
(d)
(e)
Figure 6. (a) XRD paerns for α-Mg phase and β-phase in the as-received and DSRolled alloys, (b)
variation in grain size, crystallite size, dislocation density, and average microhardness in the as-
received and DSRolled alloys, (c) variation in microhardness over the normal direction to the rolling
direction for the as-received and DSRolled alloys for 4 passes (at 70% of reduction in thickness), (d)
variation in microhardness with the reduction in thickness in as-received and DSRolled alloys, and
(e) the total DSR equivalent strain imposed in the hot DSRolled alloys at each number of passes and
ratio in thickness reduction.
The materials with hexagonal-close packed crystal structures, such as magnesium
and titanium alloys, have a limited number of independent slip systems, which makes
twinning a vital deformation mechanism at high imposed strains and elevated tempera-
tures [20]. Normally, the magnesium would be crystallographically orientated towards
the basal system before the deformation processing by rolling because the system requires
the lowest critical resolved shear stress in comparison to the counter values for non-basal
systems, such as the twin pyramidal system. However, at high imposed strains and/or
elevated temperatures, the critical resolved shear stress of the twin pyramidal system de-
creases. This will lead to the activation of additional systems to accommodate the defor-
mation at the aforementioned conditions. The activation of pyramidal deformation sys-
tems can be noticed with the increase in processing temperature above 225 °C [30,31]. In
the current research, the rolling temperatures were 400 and 450 °C, which induced
30% 50% 70%
0
20
40
60
80
100
Vickers microhardness (Hv)
Reduction in Thickness
Hot rolled – 400 ºC
Hot rolled – 450 ºC
1234
30
40
50
60
70 Reduction in Thickness %
DSR equivalent strain
Number of passes
Reduction in Thickness %
0.6
0.9
1.2
1.5
1.8
DSR equivalent strain
Figure 6. (a) XRD patterns for
α
-Mg phase and
β
-phase in the as-received and DSRolled alloys,
(b) variation in grain size, crystallite size, dislocation density, and average microhardness in the
as-received and DSRolled alloys, (c) variation in microhardness over the normal direction to the
rolling direction for the as-received and DSRolled alloys for 4 passes (at 70% of reduction in thickness),
(d) variation in microhardness with the reduction in thickness in as-received and DSRolled alloys,
and (e) the total DSR equivalent strain imposed in the hot DSRolled alloys at each number of passes
and ratio in thickness reduction.
Materials 2024, 17, x FOR PEER REVIEW 11 of 20
effectively under DSR equivalent straining (see Figure 6e) to activate the considerable py-
ramidal systems as seen by twinning that appeared in OM, SEM and via XRD peaks and
represented by twin distribution in Figure 7. The DSR equivalent strain (𝜀) was calculated
according to the following equation [12]:
𝜀=

𝜀
+
=

[
√·ln󰇡
󰇢]+
[
()·𝑅·cos(1−
 )(1−
)],
where 𝜀 and 𝛾 represent the compression and shear deformation components, respec-
tively; 𝑟 represents the reduction in thickness and calculated based on initial thickness
(ℎ) and final thickness (ℎ) as 𝑟=1−
 , 𝑅 represents the ratio of roll radii, and 𝑣
and 𝑣 are the high and low speeds of rollers, respectively. It is noted that this equation
does not directly imply the effect of rolling temperature and generated temperature due
to friction between the sample and rolls in the calculation of the imposed strain in DSR
processing. However, the DSR at a relatively low-speed ratio (<1) would not impose any
significant increase in the temperature, whereas at a high-speed ratio (>2), the increase in
the temperature would be considerable [32].
(a) (b)
Figure 7. Twin distribution in the DSRolled alloy at rolling temperatures of (a) 400 °C and (b) 450
°C.
The existence of compressive load components during straining processing via DSR
processing also contributed to the activation of these deformation planes. As the thickness
reduction per each pass was maintained at about 10%, preceding the deformation, the
activation of contraction twinning will be dominated rather than extension twinning to
accommodate the deformation homogeneity at higher strain levels and temperatures. This
can be aributed to the decrease in the sample thickness, which leads to a decrease in the
angle between the normal direction (ND) and the radial force at the edge of the rollers.
The radial force has two components: the radial force in the normal direction (ND) and
the other in the rolling direction (RD). With further thinning in the sample thickness, ra-
dial force in the normal direction will be higher than in the rolling direction, leading to
the reorientation of grain structure in the Mg-based alloys towards the activation of addi-
tional deformation twinning systems at elevated temperatures [33,34]. It seems that the
grain structure in the magnesium alloy that rolled at 400 and 450 °C was refined exten-
sively by twinning fragmentation, where the twins were generated at these temperatures
and then intersected under rolling strain deformation. Due to the limited active slip sys-
tems in magnesium and its alloys at ambient temperatures, thus the expected deformation
to occur in these materials would be via the activation of additional deformation systems.
The activation of such systems, such as pyramidal twining systems, requires elevated tem-
peratures, as in the case of deformation conditions in the current investigation. As a result
of this extensive activity of twin deformation, deformation bands and shear bands have
123456
0
20
40
60
80
100
Frequency %
Twins per grain
Hot rolled at 400ºC – 30%
Hot rolled at 400ºC – 50%
Hot rolled at 400ºC – 70%
123456
0
20
40
60
80
100
Frequency %
Twins per grain
Hot rolled at 450ºC – 30%
Hot rolled at 450ºC – 50%
Hot rolled at 450ºC – 70%
Figure 7. Twin distribution in the DSRolled alloy at rolling temperatures of (a) 400
◦
C and (b) 450
◦
C.
3.2. Mechanical Behaviour Evolution
The rolled alloys showed an increase in strength in terms of Vickers microhardness,
as shown in Figure 6c, in comparison with the as-received alloy. This increase in the
microhardness was proportional to the increase in the number of passes or reduction in
thickness in the rolled alloys, as represented in Figure 6d. However, the alloy rolled at

Materials 2024,17, 4072 11 of 17
a rolling temperature of 450
◦
C showed less proportionality with the number of passes
of DSR in comparison with the alloy rolled at a rolling temperature of 400
◦
C. It is worth
noting that the reduction ratios in thickness of 30, 50, and 70% correspond to the number of
rolling passes of one, two, and four passes, which correspond to DSR equivalent strains of
0.71, 1.16, and 1.77, respectively, as represented in Figure 6e.
The development in strength with regard to the microhardness measurements in
the rolled alloy depends on the evolution in the grain structure regarding the grain size,
crystallite size, and the distribution of
β
-phase. The grain refinement, twin generation, and
fragmentation in the rolled alloy resulted in a hardness increase by virtue of obstruction
of the generated dislocation motion during plastic deformation generated by the indenter.
Normally, the deformation of magnesium alloys at room temperature would imply the
activity of basal slip systems at low strains and non-basal slip systems at high strains. The
existence of twins in the hexagonal-closed packed magnesium alloys leads to an increase
in strain hardening under the deformation process due to obstruction of the generated
dislocation motion at the twin boundaries. Additionally, the accompanying reduction in
the grain size due to the intersection of twins and their boundaries leads to the occurrence
of the dynamic effect of Hall–Petch at these conditions [
26
,
42
]. This can be explained
according to dislocation pile-ups at the twin boundaries that result in regions of high-stress
concentration, leading to shearing of the generated twins via shearing bands that extend
over many grains. Thus, the dislocation pile-ups in these sheared regions, as well as at
twin boundaries and around the formed subgrains, will impose significant obstruction
to slip deformation, leading to obvious strain hardening in the rolled alloy. However,
the hardening in the rolled alloy at a rolling temperature of 450
◦
C was slightly lower in
comparison with the rolled alloy at 400
◦
C due to the effect of softening that lowers the
density of dislocations [
25
,
30
]. This softening arises from the intersection of generated
twins that leads to the formation of dislocation-free new grains that nucleated within
the twin bands and, at grain boundaries, that slide over each other, leading to a notable
material flow [
20
,
36
]. As a matter of fact, the dynamically recrystallized new grains were
hardly visible in this work since their formations need time and area, which were strongly
restricted by very short processing times for each sample in the DSR, where the grain
growth appears. However, this does not negate their potential contribution to the softening
of the alloy [36,43].
Despite the fact that the alloy was rolled in DSR processing at elevated temperatures,
the strength in terms of microhardness measurements was higher than the initial unrolled
alloy, which reflects the effect of the DSR imposed strain on the microstructural develop-
ment and property evolution. It was anticipated that by processing the alloy at elevated
temperatures, the grain growth would introduce a detrimental effect on the performance of
the hot rolled alloy, but the current results showed contrary to this view. The distribution of
the
β
-phase would also introduce an effect of strain hardening in the hot rolled alloy. From
SEM observation, this phase was in the form of brittle lamella and/or agglomerates in the
unrolled initial alloy. In this case, the effect of this phase in obstructing the dislocation mo-
tions and accumulating these dislocations at interfaces of coarse grain/
β
-phase would have
lower significance. However, the hot rolled alloys showed a finer
β
-phase with wide-spread
distribution due to the nature of rolling processing. Thus, the fine particles of
β
-phase act
as excellent obstructions to the dislocation activities at subgrain/fine grain/
β
-phase and
twin boundaries, leading to additional contribution for the strain hardening at ambient
temperatures [38,39].
The mechanical behavior of the as-received and rolled alloys, as tested by shear
punch testing, is represented by shear stress–strain plots in Figures 8and 9. The rolled
as-received alloy at room temperature was tested using shear punch testing at testing
temperatures of 400 and 450
◦
C, as shown in Figure 8. The samples of rolled as-received
alloy showed high values of strain hardening at all rolling passes. A noticeable increase in
the measured displacement is expressed by shear strain in the plots of shear stress–strain
with increasing the number of rolling passes and testing temperature. However, the rolled

Materials 2024,17, 4072 12 of 17
as-received samples at 450
◦
C at all passes showed higher values of strain hardening and
shear displacements with increasing the number of rolling passes, in comparison with the
rolled as-received samples at 400
◦
C that approximately the same shear displacements at
the same strain rates despite the number of rolling passes as in Figure 8. On the other hand,
the alloy rolled at 400 and 450
◦
C and then tested in the shear punch testing at 400 and
450 ◦C
showed considerable shear displacements, as in Figure 9, in comparison with the
rolled as-received alloy at room temperature. These hot-rolled alloys showed higher values
of strain hardening and lower shear displacements in an inverse manner with the increase
in the strain rate and decrease in the number of rolling passes. However, the hot-rolled alloy
at 400
◦
C showed a higher strain hardening than the hot-rolled alloy at 450
◦
C at all strain
rates and a number of rolling passes. The values of strain rate sensitivity for the hot-rolled
alloys after shear punch testing at 400 and 450
◦
C were calculated and represented in
Figure 10. The aforementioned values were used to represent the deformation mechanism
of the hot-rolled alloys at the testing temperatures. These values were slightly higher
at a testing temperature of 400
◦
C rather than 450
◦
C. The positive value of strain rate
sensitivity here reflects the increase in the shear strength with increasing strain rate at both
testing temperatures.
Materials 2024, 17, x FOR PEER REVIEW 14 of 20
(a)
(b)
Figure 8. Shear stress–strain curves for the DSRolled alloy at room temperatures for different ratios
of thickness reduction and then tested in shear punch testing at testing temperatures of (a) 400 °C
and (b) 450 °C at different strain rates.
0 102030405060
0
50
100
150
200
Shear stress (MPa)
Shear strain (%)
Rolled at RT−30% − SPT at 400
o
C−10
−1
s
−1
Rolled at RT−30% − SPT at 400
o
C−10
−2
s
−1
Rolled at RT−30% − SPT at 400
o
C−10
−3
s
−1
Rolled at RT−50% − SPT at 400
o
C−10
−1
s
−1
Rolled at RT−50% − SPT at 400
o
C−10
−2
s
−1
Rolled at RT−50% − SPT at 400
o
C−10
−3
s
−1
Rolled at RT−70% − SPT at 400
o
C−10
−1
s
−1
Rolled at RT−70% − SPT at 400
o
C−10
−2
s
−1
Rolled at RT−70% − SPT at 400
o
C−10
−3
s
−1
0 20406080100120140
0
50
100
150
200
Shear stress (MPa)
Shear strain (%)
Rolled at RT−30% − SPT at 450
o
C−10
−1
s
−1
Rolled at RT−30% − SPT at 450
o
C−10
−2
s
−1
Rolled at RT−30% − SPT at 450
o
C−10
−3
s
−1
Rolled at RT−50% − SPT at 450
o
C−10
−1
s
−1
Rolled at RT−50% − SPT at 450
o
C−10
−2
s
−1
Rolled at RT−50% − SPT at 450
o
C−10
−3
s
−1
Rolled at RT−70% − SPT at 450
o
C−10
−1
s
−1
Rolled at RT−70% − SPT at 450
o
C−10
−2
s
−1
Rolled at RT−70% − SPT at 450
o
C−10
−3
s
−1
Figure 8. Shear stress–strain curves for the DSRolled alloy at room temperatures for different ratios
of thickness reduction and then tested in shear punch testing at testing temperatures of (a) 400
◦
C
and (b) 450 ◦C at different strain rates.

Materials 2024,17, 4072 13 of 17
Materials 2024, 17, x FOR PEER REVIEW 15 of 20
(a)
(b)
Figure 9. Shear stress–strain curves for the DSRolled alloy at rolling temperatures of 400 and 450 °C
for different ratios of thickness reduction and then tested in shear punch testing at testing tempera-
tures of (a) 400 °C and (b) 450 °C at different strain rates.
0 40 80 120 160 200
0
10
20
30
40
Shear stress (MPa)
Shear strain (%)
Rolled at 400
o
C−30% − SPT at 400
o
C−10
−1
s
−1
Rolled at 400
o
C−30% − SPT at 400
o
C−10
−2
s
−1
Rolled at 400
o
C−30% − SPT at 400
o
C−10
−3
s
−1
Rolled at 400
o
C−50% − SPT at 400
o
C−10
−1
s
−1
Rolled at 400
o
C−50% − SPT at 400
o
C−10
−2
s
−1
Rolled at 400
o
C−50% − SPT at 400
o
C−10
−3
s
−1
Rolled at 400
o
C−70% − SPT at 400
o
C−10
−1
s
−1
Rolled at 400
o
C−70% − SPT at 400
o
C−10
−2
s
−1
Rolled at 400
o
C−70% − SPT at 400
o
C−10
−3
s
−1
0 40 80 120 160 200
0
10
20
30
40
Rolled at 450
o
C−30% − SPT at 450
o
C−10
−1
s
−1
Rolled at 450
o
C−30% − SPT at 450
o
C−10
−2
s
−1
Rolled at 450
o
C−30% − SPT at 450
o
C−10
−3
s
−1
Rolled at 450
o
C−50% − SPT at 450
o
C−10
−1
s
−1
Rolled at 450
o
C−50% − SPT at 450
o
C−10
−2
s
−1
Rolled at 450
o
C−50% − SPT at 450
o
C−10
−3
s
−1
Rolled at 450
o
C−70% − SPT at 450
o
C−10
−1
s
−1
Rolled at 450
o
C−70% − SPT at 450
o
C−10
−2
s
−1
Rolled at 450
o
C−70% − SPT at 450
o
C−10
−3
s
−1
Shear stress (MPa)
Shear strain (%)
Figure 9. Shear stress–strain curves for the DSRolled alloy at rolling temperatures of 400 and
450
◦
C for different ratios of thickness reduction and then tested in shear punch testing at testing
temperatures of (a) 400 ◦C and (b) 450 ◦C at different strain rates.
In shear punch testing, the as-received samples that rolled at room temperature have
presented a significant strain hardening in comparison to the hot rolled samples. This can
be attributed to the strain hardening in the rolled alloy that increased significantly at room
temperature as the number of passes increased or with the higher reduction in the thickness
of sheet samples of the rolled alloy. In the rolled samples at room temperature, the level of
defects during DSR processing is anticipated to reach higher levels in comparison with the
hot-rolled samples, leading to a considerable obstruction to the tension deformation even at
elevated testing temperatures that leads to high values of tensile strengths in the as-received
samples post cold rolling. However, these samples showed moderate elongations in terms
of shear displacements with slower strain rates and a higher number of passes. This can
be explained in terms of grain refinement in the as-received samples that rolled at room
temperature, which is expected to be much finer than for the hot-rolled samples. This would
provide paths for grain-boundary sliding during the hot deformation in tensile testing at

Materials 2024,17, 4072 14 of 17
slower strain rates [
24
,
44
]. At DSR processing at room temperature, it is anticipated that the
grain size reaches the ultrafine scale due to the high value of the defects that are imposed
in the microstructure where no effects of dynamic recovery, recrystallization, or grain
growth would appear. Thus, the achieved elongations in the as-received samples that were
processed in DSR at room temperature were expected due to the formation of equiaxed and
elongated microstructures towards the rolling direction, leading to an enhancement in the
ductility. However, these elongations were even better in the hot-rolled samples rather than
the cold-rolled counterparts, where the hot-rolled samples showed significant improvement
in the achieved elongations in terms of shear displacements with increasing the number
of passes. This can be attributed to the significant activity of twinning that was generated
during hot rolling and then persisted during the hot deformation in the shear punch
testing. The deformation of magnesium alloys at room temperature is limited mainly to the
basal slip activity that leads to constraints in the material flow and elongations at ambient
temperatures. However, at elevated temperature deformation, these alloys tend to initiate
twins to accommodate the deformation following the twin-induced mechanism, associated
with the increase in the slip deformation mechanism leading to the end in the formation of
new dynamic recrystallized grains [
45
,
46
]. In the current investigation, the hot-rolled alloys
showed no dynamic recrystallization at the range of processing temperatures but instead
significant activity of the generated twins. Thus, the edge dislocations that accumulated
and moved over the pyramidal twin system would be bound on the slip basal system, and
their mobility would increase with increasing temperature, leading to a significant flow
of the alloy under elevated temperature deformation. The values of strain rate sensitivity,
as represented in Figure 10, and the distribution of twins data, as represented in Figure 7,
confirm the domination of pyramidal cross-slip as a main mechanism of the alloy flow
during shear punch testing at temperatures 400–450
◦
C [
45
,
47
]. The distribution of the
fine particles for the
β
-phase after the hot rolling is expected to have an impact on the
alloy flow and elongation under the condition of hot deformation in the current study. The
β
-phase has fragmented and refined down to the nanoscale and distributed along the twin
boundaries and inside the grains, leading to strain hardening at the initial stages of hot
deformation followed by the softening effect of the alloy under this condition. This can
be attributed to the pinning effect of these fine particles to the motion of dislocation and
material flow at the early stage of deformation despite the elevated temperature of the test.
These particles then act as a lubricant, which facilitates the glide of grains over grains and
twin boundaries at testing temperatures of 400–450
◦
C (673–723 K); this testing temperature
represents about (0.91–0.98) of the melting point of this phase of 460 ◦C (733 K) [48,49].
Materials 2024, 17, x FOR PEER REVIEW 16 of 20
(a)
(b)
Figure 10. Values of strain rate sensitivity for the DSRolled alloy at rolling temperatures of 400 and
450 °C for different ratios of thickness reduction and then tested in shear punch testing at testing
temperatures of (a) 400 °C and (b) 450 °C at different strain rates.
In shear punch testing, the as-received samples that rolled at room temperature have
presented a significant strain hardening in comparison to the hot rolled samples. This can
be aributed to the strain hardening in the rolled alloy that increased significantly at room
temperature as the number of passes increased or with the higher reduction in the thick-
ness of sheet samples of the rolled alloy. In the rolled samples at room temperature, the
level of defects during DSR processing is anticipated to reach higher levels in comparison
with the hot-rolled samples, leading to a considerable obstruction to the tension defor-
mation even at elevated testing temperatures that leads to high values of tensile strengths
in the as-received samples post cold rolling. However, these samples showed moderate
elongations in terms of shear displacements with slower strain rates and a higher number
of passes. This can be explained in terms of grain refinement in the as-received samples
that rolled at room temperature, which is expected to be much finer than for the hot-rolled
samples. This would provide paths for grain-boundary sliding during the hot defor-
mation in tensile testing at slower strain rates [24,44]. At DSR processing at room
10
−3
10
−2
10
−1
20
30
40
50
Flow stress (MPa)
Strain rate (s
−1
)
Rolled at 400
o
C−30% − SPT at 400
o
C
Rolled at 400
o
C−50% − SPT at 400
o
C
Rolled at 400
o
C−70% − SPT at 400
o
C
m-values
0.23
0.22
0.20
10
−3
10
−2
10
−1
10
20
30
40 m-values
Flow stress (MPa)
Strain rate (s
-1
)
Rolled at 450
o
C−30% − SPT at 450
o
C
Rolled at 450
o
C−50% − SPT at 450
o
C
Rolled at 450
o
C−70% − SPT at 450
o
C
0.15
0.20
0.15
Figure 10. Cont.

Materials 2024,17, 4072 15 of 17
Materials 2024, 17, x FOR PEER REVIEW 16 of 20
(a)
(b)
Figure 10. Values of strain rate sensitivity for the DSRolled alloy at rolling temperatures of 400 and
450 °C for different ratios of thickness reduction and then tested in shear punch testing at testing
temperatures of (a) 400 °C and (b) 450 °C at different strain rates.
In shear punch testing, the as-received samples that rolled at room temperature have
presented a significant strain hardening in comparison to the hot rolled samples. This can
be aributed to the strain hardening in the rolled alloy that increased significantly at room
temperature as the number of passes increased or with the higher reduction in the thick-
ness of sheet samples of the rolled alloy. In the rolled samples at room temperature, the
level of defects during DSR processing is anticipated to reach higher levels in comparison
with the hot-rolled samples, leading to a considerable obstruction to the tension defor-
mation even at elevated testing temperatures that leads to high values of tensile strengths
in the as-received samples post cold rolling. However, these samples showed moderate
elongations in terms of shear displacements with slower strain rates and a higher number
of passes. This can be explained in terms of grain refinement in the as-received samples
that rolled at room temperature, which is expected to be much finer than for the hot-rolled
samples. This would provide paths for grain-boundary sliding during the hot defor-
mation in tensile testing at slower strain rates [24,44]. At DSR processing at room
10
−3
10
−2
10
−1
20
30
40
50
Flow stress (MPa)
Strain rate (s
−1
)
Rolled at 400
o
C−30% − SPT at 400
o
C
Rolled at 400
o
C−50% − SPT at 400
o
C
Rolled at 400
o
C−70% − SPT at 400
o
C
m-values
0.23
0.22
0.20
10
−3
10
−2
10
−1
10
20
30
40 m-values
Flow stress (MPa)
Strain rate (s
-1
)
Rolled at 450
o
C−30% − SPT at 450
o
C
Rolled at 450
o
C−50% − SPT at 450
o
C
Rolled at 450
o
C−70% − SPT at 450
o
C
0.15
0.20
0.15
Figure 10. Values of strain rate sensitivity for the DSRolled alloy at rolling temperatures of 400 and
450
◦
C for different ratios of thickness reduction and then tested in shear punch testing at testing
temperatures of (a) 400 ◦C and (b) 450 ◦C at different strain rates.
4. Conclusions
1.
A commercial Mg–8Al–1Zn magnesium alloy was successfully processed by differen-
tial speed rolling at elevated temperatures of 400 and 450 ◦C without any significant
grain growth at these rolling temperatures for a reduction percentage of 30% and 70%.
2.
Considerable twinning was observed in DSRolled alloy with no grain growth nor dy-
namic recrystallization indicating that twin deformation is the dominant deformation
mechanism in addition to slip mechanism.
3.
The DSRolled alloy showed significant elongation in terms of shear displacements at
testing temperatures of 400 and 450 ◦C.
4.
The existence of twins and the distribution of fine particles of
β
-phase (Mg
17
Al
12
)
played important roles in the behavior of DSRolled alloy under the hot shear defor-
mation. These two factors inhibit the significant grain growth, improve the strain
hardening and improve the alloy flow under the hot deformation conditions.
Author Contributions: Conceptualization, E.A.H.; Methodology, S.A.A., A.S.J.A.-Z. and E.A.H.;
Validation, M.Y.A.; Formal analysis, S.A.A. and A.S.J.A.-Z.; Investigation, E.A.H.; Resources, M.Y.A.;
Writing—original draft, S.A.A.; Writing—review & editing, A.S.J.A.-Z. and M.Y.A. All authors have
read and agreed to the published version of the manuscript.
Funding: This research received no external funding.
Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.
Data Availability Statement: The original contributions presented in the study are included in the
article, further inquiries can be directed to the corresponding author.
Conflicts of Interest: The authors declare no conflict of interest.
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Xu, S.W.; Zheng, M.Y.; Kamado, S.; Wu, K.; Wang, G.J.; Lv, X.Y. Dynamic microstructural changes during hot extrusion and
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