Available via license: CC BY 4.0
Content may be subject to copyright.
Citation: Rogachev, S.O.; Bazhenov,
V.E.; Komissarov, A.A.; Li, A.V.; Ten,
D.V.; Yushchuk, V.V.; Drobyshev, A.Y.;
Shin, K.S. Effect of Hot Rolling on
Structure and Mechanical Properties
of Mg–Y–Zn–Mn Alloys. Metals 2023,
13, 223. https://doi.org/10.3390/
met13020223
Academic Editor: Mohammad Jahazi
Received: 29 December 2022
Revised: 15 January 2023
Accepted: 21 January 2023
Published: 25 January 2023
Copyright: © 2023 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/).
metals
Article
Effect of Hot Rolling on Structure and Mechanical Properties of
Mg–Y–Zn–Mn Alloys
Stanislav O. Rogachev 1, 2, * , Viacheslav E. Bazhenov 3, Alexander A. Komissarov 1,4, Anna V. Li 1,
Denis V. Ten 1, Viacheslav V. Yushchuk 1, Alexey Yu. Drobyshev 4and Kwang Seon Shin 4,5
1Department of Physical Metallurgy and Physics of Strength, National University of Science and Technology
MISiS, 119049 Moscow, Russia
2Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, 119334 Moscow, Russia
3Casting Department, National University of Science and Technology MISiS, 119049 Moscow, Russia
4Laboratory of Medical Bioresorption and Bioresistance, Moscow State University of Medicine and Dentistry,
127473 Moscow, Russia
5
Magnesium Technology Innovation Center, Department of Materials Science and Engineering, Seoul National
University, 1 Gwanak-ro, Gwanak-gu, Seoul 08826, Republic of Korea
*Correspondence: csaap@mail.ru; Tel.: +7-903-967-5782
Abstract:
The effect of hot rolling on the structure and mechanical properties of three Mg–Y–Zn–Mn
alloys was studied depending on the process temperature and the reduction ratio. The original plates
of cast WZM111, WZM211, and WZM321 alloys after heat treatment were subjected to rolling from an
initial thickness of 7 mm to a final thickness of 0.2 mm at two temperatures, namely 400 and 450 ◦C.
Optical and scanning electron microscopy, the microhardness measurement, and tensile testing were
used to characterize the material. The rolling regimes that provide a good balance between the
strength and ductility of the alloys were established.
Keywords: Mg–Y–Zn–Mn magnesium alloys; rolling; mechanical properties; microstructure
1. Introduction
In past decades, the healthcare systems in many countries used permanent bone fix-
ation implants because titanium alloys are a gold standard for osteosynthesis. However,
its disadvantages, including temperature and tactile sensitivity, growth restriction, and
titanium particles in tissue together with frequently symptomatic removal, make looking
for other solutions necessary [
1
]. At present, temporary biodegradable implants that gradu-
ally dissolve as the healing process progresses and reduce healthcare costs by eliminating
secondary surgery for implant removal are gaining popularity [
2
,
3
]. Mg alloys, due to
enough mechanical properties, biocompatibility, and acceptable biodegradation rate, are at-
tractive candidates as materials for temporary fixation devices used in
osteosynthesis [4–6]
.
In comparison with permanent Ti implants, Mg ones have a density and Young’s modulus
closer to cortical bone [
7
,
8
]. The commercial NOVAMag
®
and MAGNEZIX
®
fixation de-
vices produced by Botiss biomaterials GmbH and Syntellix AG (both Germany) are used in
orthopedic practices’ in many countries and provide equal performance with Ti permanent
implants [9,10].
It is well known that the grain size significantly affects the mechanical properties of
Mg alloys because of their high Hall–Petch strengthening coefficient (~300 MPa
·µ
m
1/2
) [
11
].
Various methods, such as hot extrusion, equal channel angular pressing, and hot rolling,
are used to provide the fine-grained structure and the high mechanical properties of mag-
nesium alloys [
12
–
17
]. At the same time, in the works where biodegradable magnesium
alloys are developed, the manufacturing technique for the implant type is not considered.
For example, in many works, biodegradable plates are cut from extruded bars [
18
–
20
]. The
production yield in this case is extremely low and a high amount of scraps are produced
Metals 2023,13, 223. https://doi.org/10.3390/met13020223 https://www.mdpi.com/journal/metals
Metals 2023,13, 223 2 of 15
that lead to a large environmental impact. Rolling is typically used for the large-quantity
manufacturing of titanium alloy permanent plates, and this technique best fits biodegrad-
able magnesium alloy plate production. In this work, the influences of Zn and Y content
in Mg–Y–Zn–Mn alloys on the structure and mechanical properties during rolling are
investigated in order to choose the best alloy composition and rolling path to produce the
rolled sheet for further manufacturing of biodegradable plates.
In the last decades, a great importance has been on Mg–Zn–Y alloys, which have
a high strength after deformation processing [
21
–
23
]. According to the composition of
Mg–Zn–Y alloys, the long-period stacking-ordered (LPSO) phase (Mg
12
ZnY), W-phase
(Mg
3
Zn
3
Y
2
), I-phase (Mg
3
Zn
6
Y), and Mg
24
Y
5
phase can be found in their structure [
24
–
26
].
It was shown previously that the LPSO phase is preferable in view of the higher corrosion
resistance of the alloy [
27
]. Because of that, in this work, alloys with the LPSO phase in
their structure are under investigation.
Mn addition promotes effective grain refinement in magnesium alloys in
general [28–32]
and for Mg–Zn–Y alloys in particular [
33
–
35
], and hence it improves the mechanical
properties of the alloys. The effectiveness of Mn is from the formation of
α
-Mn precipitates
that hinder the grain growth during recrystallization [
15
]. It was shown that the addition
of Mn improves the corrosion resistance of Mg alloys due to the formation of a Mn-rich
oxide film, which prevents chloride ion penetration and also forms intermetallics with Fe,
which is the most harmful impurity in Mg alloys [5,32].
No works were found focused on the hot rolling of Mg–Y–Zn–Mn alloys. Thus, the
aim of this study is to investigate the effect of the Zn and Y content in rolled Mg–Y–Zn–
Mn alloys on the microstructure and mechanical properties to evaluate their potential for
application in biodegradable orthopedic plates.
2. Materials and Methods
2.1. Materials
For alloy preparation, the following raw materials were used: Mg (99.95 wt.% purity;
SOMZ, Solikamsk, Russia), Zn (99.995 wt.%; UGMK, Verkhnaya Pyshma, Russia),
Mg–3Mn
(wt.%) master alloy (self-made using Mn 99.8 wt.%), and Mg–20Y (wt.%) master alloy
(Uralredmet, Verkhnaya Pyshma, Russia). The melts were prepared using a graphite
crucible in a resistance furnace under Ar + 2 vol.% SF
6
atmosphere. Before pouring, the
melt was purged with Ar. The rectangular ingots 12
×
60
×
200 mm
3
were cast into graphite
permanent molds preheated to 150
◦
C. Three Mg–Y–Zn–Mn alloys with different Y and
Zn content were prepared as listed in Table 1. The alloy compositions were determined
using energy-dispersive X-ray spectroscopy (EDS, Oxford Instruments, Oxford, UK) on
the metallographic sections with 0.1 wt.% accuracy. Three areas of size 1
×
1 mm
2
were
analyzed for each specimen.
Table 1. Chemical composition of the studied magnesium alloys.
Alloy Designation Content of Element (wt.%)
Mg Y Zn Mn
WZM111 Bal. 1.2 0.6 0.8
WZM211 Bal. 2.5 1.1 0.8
WZM321 Bal. 3.4 1.7 0.8
Before rolling, the castings were subjected to solution treatment at a temperature of
520
◦
C for 10 h, followed by quenching in water. After that, the castings were processed
using a milling machine to the required size, namely plates 54 mm wide and 70 mm long.
The solution treatment temperature (520
◦
C) was taken as 20
◦
C lower than the solidus
temperature determined from the plotted nonequilibrium solidification pathway of alloys
in accordance with the Sheil–Gulliver model (Figure 1).
Metals 2023,13, 223 3 of 15
Metals 2023, 13, x FOR PEER REVIEW 3 of 16
The solution treatment temperature (520 °C) was taken as 20 °C lower than the solidus
temperature determined from the plotted nonequilibrium solidification pathway of al-
loys in accordance with the Sheil–Gulliver model (Figure 1).
Figure 1. (a) The nonequilibrium solidification pathway of the studied alloys according to the
Sheil–Gulliver model and (b) the enlarged fragment for the mass fraction of the solid 0.8–1
The heat-treated plates with an initial thickness of 7.1 ± 0.1 mm were subjected to
rolling to a final thickness of 0.23 ± 0.2 mm through 18 passes. Rolling was carried out at
two temperatures, namely 400 and 450 °C. The reduction ratio per pass was calculated by
the equation %100
0
0⋅
−
=
h
hh k
ε
where 0
h and k
h are the plate’s thickness before
and after each specific pass, respectively. The average reduction ratio per pass was 17%.
The sample designations and characteristics (total reduction ratio and final thickness)
selected for further research and testing are listed in Table 2.
Table 2. Characteristics of the studied samples.
Sample Designation Total Reduction
Ratio (%)
Final Thickness
(mm)
E3 35–37 4.42–4.65
E6 60–65 2.47–2.89
E12 88–89 0.80–0.83
E15 94 0.39–0.41
E18 96–97 0.21–0.25
2.2. Materials
For structural studies and microhardness measurements, ND-RD longitudinal sec-
tions of size 15 × 5 mm2 were cut from rolled plates with different reductions (thickness)
by the electro-spark method. The sections were embedded in epoxy resin using a Sim-
pliMet 1000 machine (Buehler, Leinfelden-Echterdingen, Germany); after that, the sec-
tions were prepared. The sections’ surfaces were carefully ground and polished to a
mirror finish. Grinding was carried out sequentially on abrasive paper with a roughness
of P400, P1000, and P2500. Polishing was carried out on fabric using diamond paste. The
structural studies were carried out on pre-etched sections in a solution of the following
composition: 20 mL of acetic glycol, 1 mL of nitric acid, 60 mL of ethylene glycol, and 60
mL of distilled water; the microhardness was tested on unetched sections.
Figure 1.
(
a
) The nonequilibrium solidification pathway of the studied alloys according to the
Sheil–Gulliver model and (b) the enlarged fragment for the mass fraction of the solid 0.8–1.
The heat-treated plates with an initial thickness of 7.1
±
0.1 mm were subjected to
rolling to a final thickness of 0.23
±
0.2 mm through 18 passes. Rolling was carried out
at two temperatures, namely 400 and 450
◦
C. The reduction ratio per pass was calculated
by the equation
ε=h0−hk
h0·
100% where
h0
and
hk
are the plate’s thickness before and after
each specific pass, respectively. The average reduction ratio per pass was 17%. The sample
designations and characteristics (total reduction ratio and final thickness) selected for
further research and testing are listed in Table 2.
Table 2. Characteristics of the studied samples.
Sample Designation Total Reduction Ratio (%) Final Thickness (mm)
E3 35–37 4.42–4.65
E6 60–65 2.47–2.89
E12 88–89 0.80–0.83
E15 94 0.39–0.41
E18 96–97 0.21–0.25
2.2. Materials
For structural studies and microhardness measurements, ND-RD longitudinal sections
of size 15
×
5 mm
2
were cut from rolled plates with different reductions (thickness) by
the electro-spark method. The sections were embedded in epoxy resin using a SimpliMet
1000 machine (Buehler, Leinfelden-Echterdingen, Germany); after that, the sections were
prepared. The sections’ surfaces were carefully ground and polished to a mirror finish.
Grinding was carried out sequentially on abrasive paper with a roughness of P400, P1000,
and P2500. Polishing was carried out on fabric using diamond paste. The structural studies
were carried out on pre-etched sections in a solution of the following composition: 20 mL
of acetic glycol, 1 mL of nitric acid, 60 mL of ethylene glycol, and 60 mL of distilled water;
the microhardness was tested on unetched sections.
For tensile testing, small-sized tensile specimens with a total length of 12 mm and a
length and width of the gauge part of 5 and 1.45 mm, respectively, were cut from the plates
along the rolling direction by the electro-spark method. The thickness of the specimens
corresponded to the thickness of the plate after rolling to a specific reduction ratio. Before
the tensile test, the surfaces of the tensile specimens were not subjected to additional
grinding or polishing.
Metals 2023,13, 223 4 of 15
2.3. Microhardness Measurement
The Vickers microhardness values were measured on samples’ longitudinal sections
using a Micromet 5101 tester (load 0.5 N, load exposure time 10 s, Buehler, Leinfelden-
Echterdingen, Germany). Six measurements were made randomly with the calculation of
the arithmetic average and standard deviation.
2.4. Tensile Test
Tensile testing of small-sized specimens was carried out using an Instron 5966 uni-
versal testing machine (Instron, Norwood, MA, USA) with special adapters installed. The
deformation rate was 0.002 s
−1
. During the test, the stress–strain curves were plotted
using the built-in software; after that, the mechanical characteristics of the specimens were
determined: the yield strength (YS), the ultimate tensile strength (UTS), and the relative
elongation (RE). Three specimens per point were tested.
2.5. Structural Studies
The study of the longitudinal sections’ microstructures was carried out using an optical
microscope (OM) Axio Observer D1m Carl Zeiss (Carl Zeiss, Oberkochen, Germany) with
a built-in digital camera at magnifications from 50
×
to 1000
×
as well as a scanning electron
microscope (SEM) Tescan Vega 3 SBH (Tescan, Brno, Czech Republic) with an Oxford
energy-dispersive microanalysis attachment in the back-scattered electron mode (sensitive
to the atomic mass of a chemical element) at magnifications from 200×to 5000×.
3. Results and Discussion
3.1. Microstructure Analysis
3.1.1. Microstructure of the Alloys in As-Cast and Heat Treated States
Figure 2shows the SEM images representing the typical microstructures of the studied
samples in the as-cast state and after heat treatment. As-cast alloys have a dendritic
structure of a magnesium solid solution (
α
-Mg) and an LPSO phase with specific broken
boundaries located along the boundaries of dendritic cells. According to the EDS data,
almost all manganese is in
α
-Mg. In addition to the LPSO phase, a small amount of the
Y-rich phase is observed in the structure, the composition of which is close to Mg
24
Y
5
. It
corresponds to the calculations of the solidification pathway using the Sheil–Gulliver model
(see Figure 1). After heat treatment, the structures of the WZM111 and WZM211 alloys
are mainly homogenized, but the Mg
24
Y
5
phase can be observed, while the open contours
of the dendritic structure with fragments of the LPSO phase remain in the WZM321 alloy.
These results contradict the calculated equilibrium phase composition of the WZM111 and
WZM211 alloys for a temperature of 520
◦
C (see Table 3), according to which the WZM111
and WZM211 alloys should contain a small amount of the LPSO phase, about 3% and 8%,
respectively, but are consistent with the calculation for the WZM321 alloy for the fraction
of the LPSO phase which should reach 12%.
Table 3. Equilibrium phase composition (wt%) of the studied alloys at different temperatures.
Alloy
Phase Fraction at Temperature
520 ◦C 450 ◦C 400 ◦C
α-Mg LPSO α-Mn α-Mg LPSO α-Mn α-Mg LPSO α-Mn
WZM111 96.91 3.09 0 95.27 4.38 0.35 94.80 4.67 0.53
WZM211 92.18 7.78 0.04 91.09 8.51 0.40 90.80 8.64 0.56
WZM321 87.64 12.30 0.06 86.68 12.92 0.40 86.42 13.02 0.56
Metals 2023,13, 223 5 of 15
Metals 2023, 13, x FOR PEER REVIEW 5 of 16
Figure 2. Microstructure (SEM) of magnesium alloys (a,b) WZM111, (c,d) WZM211, and (e,f)
WZM321: (a,c,e) as-cast state and (b,d,f) after heat treatment (HT).
Table 3. Equilibrium phase composition (wt%) of the studied alloys at different temperatures.
Alloy
Phase Fraction at Temperature
520 °С
450 °С
400 °С
α-Mg
LPSO
α-Mn
α-Mg
LPSO
α-Mn
α-Mg
LPSO
α-Mn
WZM111
96.91
3.09
0
95.27
4.38
0.35
94.80
4.67
0.53
WZM211
92.18
7.78
0.04
91.09
8.51
0.40
90.80
8.64
0.56
WZM321
87.64
12.30
0.06
86.68
12.92
0.40
86.42
13.02
0.56
3.1.2. Microstructure of the Alloys after Rolling
The plates made of the WZM211 and WZM321 alloys failed during rolling at 400 °C
(the WZM211 alloy after the 6th pass, and the WZM321 alloy after the 1st pass). The rest
Figure 2.
Microstructure (SEM) of magnesium alloys (
a
,
b
) WZM111, (
c
,
d
) WZM211, and (
e
,
f
) WZM321:
(a,c,e) as-cast state and (b,d,f) after heat treatment (HT).
3.1.2. Microstructure of the Alloys after Rolling
The plates made of the WZM211 and WZM321 alloys failed during rolling at 400
◦
C
(the WZM211 alloy after the 6th pass, and the WZM321 alloy after the 1st pass). The
rest of the plates for all rolling regimes were successfully rolled to a final thickness of
0.23 ±0.2 mm through 18 passes.
Figures 3–6show the OM images representing the typical microstructures of the
studied samples after rolling with different reductions. Since both variants of rolling were
carried out at high homologous temperatures (0.6–0.7 T
melt
), dynamic recrystallization
played the main role in the microstructure formation and grain refinement [36,37].
Metals 2023,13, 223 6 of 15
Metals 2023, 13, x FOR PEER REVIEW 6 of 16
of the plates for all rolling regimes were successfully rolled to a final thickness of 0.23 ±
0.2 mm through 18 passes.
Figures 3–6 show the OM images representing the typical microstructures of the
studied samples after rolling with different reductions. Since both variants of rolling
were carried out at high homologous temperatures (0.6–0.7 Tmelt), dynamic recrystalliza-
tion played the main role in the microstructure formation and grain refinement [36,37].
Figure 3. Microstructure (OM) of the WZM111 alloy after rolling at 400 °C to different reductions:
(a)–E3, (b)–E6, (c)–E18.
Figure 4. Microstructure (OM) of the WZM111 alloy after rolling at 450 °C to different reductions:
(a,b)–E3, (c,d)–E6, (e)–E12, (f)–E18.
Figure 3.
Microstructure (OM) of the WZM111 alloy after rolling at 400
◦
C to different reductions:
(a)–E3, (b)–E6, (c)–E18.
Metals 2023, 13, x FOR PEER REVIEW 6 of 16
of the plates for all rolling regimes were successfully rolled to a final thickness of 0.23 ±
0.2 mm through 18 passes.
Figures 3–6 show the OM images representing the typical microstructures of the
studied samples after rolling with different reductions. Since both variants of rolling
were carried out at high homologous temperatures (0.6–0.7 Tmelt), dynamic recrystalliza-
tion played the main role in the microstructure formation and grain refinement [36,37].
Figure 3. Microstructure (OM) of the WZM111 alloy after rolling at 400 °C to different reductions:
(a)–E3, (b)–E6, (c)–E18.
Figure 4. Microstructure (OM) of the WZM111 alloy after rolling at 450 °C to different reductions:
(a,b)–E3, (c,d)–E6, (e)–E12, (f)–E18.
Figure 4.
Microstructure (OM) of the WZM111 alloy after rolling at 450
◦
C to different reductions:
(a,b)–E3, (c,d)–E6, (e)–E12, (f)–E18.
Metals 2023, 13, x FOR PEER REVIEW 6 of 16
of the plates for all rolling regimes were successfully rolled to a final thickness of 0.23 ±
0.2 mm through 18 passes.
Figures 3–6 show the OM images representing the typical microstructures of the
studied samples after rolling with different reductions. Since both variants of rolling
were carried out at high homologous temperatures (0.6–0.7 Tmelt), dynamic recrystalliza-
tion played the main role in the microstructure formation and grain refinement [36,37].
Figure 3. Microstructure (OM) of the WZM111 alloy after rolling at 400 °C to different reductions:
(a)–E3, (b)–E6, (c)–E18.
Figure 4. Microstructure (OM) of the WZM111 alloy after rolling at 450 °C to different reductions:
(a,b)–E3, (c,d)–E6, (e)–E12, (f)–E18.
Figure 5.
Microstructure (OM) of the WZM211 alloy after rolling at 450
◦
C to different reductions (the
grain boundaries in the insert of Figure 5c were drawn in a graphics editor): (
a
)–E3, (
b
)–E6, (
c
)–E18.
Metals 2023,13, 223 7 of 15
Metals 2023, 13, x FOR PEER REVIEW 7 of 16
Figure 5. Microstructure (OM) of the WZM211 alloy after rolling at 450 °C to different reductions
(the grain boundaries in the insert of Figure 5c were drawn in a graphics editor): (a)–E3, (b)–E6,
(c)–E18.
Figure 6. Microstructure (OM) of the WZM321 alloy after rolling at 450 °C to different reductions
(the grain boundaries in the insert of Figure 6c were drawn in a graphics editor): (a)–E3, (b)–E6,
(c)–E18.
The evolution of the microstructure with an increase in the reduction ratio differed
in the three alloys and, in addition, depended on the rolling temperature. Thus, in the
WZM111 alloy, after rolling at a temperature of 400 °C to a low reduction ratio (E3), a
deformed structure with numerous intersecting shear bands was observed (Figure 3a).
With an increase in the reduction ratio to E6, equiaxed grains with a predominant size of
4–8 μm were formed in the alloy structure as a result of dynamic recrystallization (Figure
3b). The grains were elongated in the rolling direction. With a further increase in the re-
duction ratio to E18, the predominant grain size decreased to 3–5 μm, and the grain
shape tended to be equiaxed (Figure 3c).
With an increase in the rolling temperature to 450 °C, the equiaxed grains in the
structure of the WZM111 alloy were already formed at a reduction ratio of E3 (Figure 4a),
although the deformed areas remained (Figure 4b). The predominant size of such grains
was 10–20 µm. With an increase in the reduction ratio to E6–E18, an alternation of the
deformed areas and the recrystallized grains was observed in the structure, and the grain
size both increased and decreased with the increasing reduction ratio (Figure 4c–f). The
body of many grains contained parallel slip bands. This indicates the incompleteness of
the dynamic recrystallization process. In addition, the precipitating process of the α-Mn
particles, which can contribute both to the recrystallized grain refinement and additional
strengthening, was superimposed on the microstructure formation process. According to
the calculated equilibrium phase composition of the alloys for the rolling temperatures,
the fractions of the α-Mn phase in all three alloys were about 0.4% at 450 °C and about
0.6% at 400 °C (see Table 3).
In the WZM211 alloy, the dynamic recrystallization process was difficult, which is
associated with the precipitation of numerous LPSO phase particles during deformation.
At a low reduction ratio (E3), the structure was close to that of the WZM111 alloy after
rolling at 400 °С (Figure 5a). With an increase in the reduction ratio, the deformed areas
and the recrystallized grains alternated in the structure, and the developed grain struc-
ture was formed only after high reductions (E18) (Figure 5c). At the reduction value of
E18, the grain size was 4–12 µm.
The dynamic recrystallization process in the WZM321 alloy was the most difficult,
which is also associated with the precipitation of numerous LPSO phase particles. Thus,
at a low reduction (E3), the original large grains with strongly distorted boundaries,
elongated in the direction of rolling, were retained in the alloy structure (Figure 6a). The
recrystallized grains began to form only at high reductions (E12). However, even after
maximum reduction (E18), a slight elongation of the grains along the direction of rolling
was observed (Figure 6c). At the reduction value of E18, the grain size was 3–8 µm.
Figure 6.
Microstructure (OM) of the WZM321 alloy after rolling at 450
◦
C to different reductions (the
grain boundaries in the insert of Figure 6c were drawn in a graphics editor): (
a
)–E3, (
b
)–E6, (
c
)–E18.
The evolution of the microstructure with an increase in the reduction ratio differed
in the three alloys and, in addition, depended on the rolling temperature. Thus, in the
WZM111 alloy, after rolling at a temperature of 400
◦
C to a low reduction ratio (E3), a
deformed structure with numerous intersecting shear bands was observed (Figure 3a). With
an increase in the reduction ratio to E6, equiaxed grains with a predominant size of 4–8
µ
m
were formed in the alloy structure as a result of dynamic recrystallization (Figure 3b). The
grains were elongated in the rolling direction. With a further increase in the reduction ratio
to E18, the predominant grain size decreased to 3–5 µm, and the grain shape tended to be
equiaxed (Figure 3c).
With an increase in the rolling temperature to 450
◦
C, the equiaxed grains in the
structure of the WZM111 alloy were already formed at a reduction ratio of E3 (Figure 4a),
although the deformed areas remained (Figure 4b). The predominant size of such grains
was 10–20
µ
m. With an increase in the reduction ratio to E6–E18, an alternation of the
deformed areas and the recrystallized grains was observed in the structure, and the grain
size both increased and decreased with the increasing reduction ratio (Figure 4c–f). The
body of many grains contained parallel slip bands. This indicates the incompleteness of
the dynamic recrystallization process. In addition, the precipitating process of the
α
-Mn
particles, which can contribute both to the recrystallized grain refinement and additional
strengthening, was superimposed on the microstructure formation process. According to
the calculated equilibrium phase composition of the alloys for the rolling temperatures, the
fractions of the
α
-Mn phase in all three alloys were about 0.4% at 450
◦
C and about 0.6% at
400 ◦C (see Table 3).
In the WZM211 alloy, the dynamic recrystallization process was difficult, which is
associated with the precipitation of numerous LPSO phase particles during deformation.
At a low reduction ratio (E3), the structure was close to that of the WZM111 alloy after
rolling at 400
◦
C (Figure 5a). With an increase in the reduction ratio, the deformed areas
and the recrystallized grains alternated in the structure, and the developed grain structure
was formed only after high reductions (E18) (Figure 5c). At the reduction value of E18, the
grain size was 4–12 µm.
The dynamic recrystallization process in the WZM321 alloy was the most difficult,
which is also associated with the precipitation of numerous LPSO phase particles. Thus, at
a low reduction (E3), the original large grains with strongly distorted boundaries, elongated
in the direction of rolling, were retained in the alloy structure (Figure 6a). The recrystal-
lized grains began to form only at high reductions (E12). However, even after maximum
reduction (E18), a slight elongation of the grains along the direction of rolling was observed
(Figure 6c). At the reduction value of E18, the grain size was 3–8 µm.
Before rolling, as noted above, the LPSO phase particles were present only in the
WZM321 alloy. During rolling, the LPSO phase particles precipitated in the WZM211 and
WZM321 alloys, which hinders dynamic recrystallization. As an example, Figure 7shows
SEM images of the WZM321 alloy structure after rolling with different reductions. At low
Metals 2023,13, 223 8 of 15
reductions (E3-E6), the structure contains large particles of the LPSO phase, up to 50
µ
m
in length (particles of Type 1 in Figure 7). These particles were present in the structure of
the alloy before rolling. With an increase in the reductions (E12–E18), large particles of the
LPSO phase were elongated in the direction of rolling, and fine particles of the LPSO phase,
0.5–2 µm in size, precipitated in the structure (particles of Type 2 in Figure 7).
Metals 2023, 13, x FOR PEER REVIEW 8 of 16
Before rolling, as noted above, the LPSO phase particles were present only in the
WZM321 alloy. During rolling, the LPSO phase particles precipitated in the WZM211 and
WZM321 alloys, which hinders dynamic recrystallization. As an example, Figure 7 shows
SEM images of the WZM321 alloy structure after rolling with different reductions. At low
reductions (E3-E6), the structure contains large particles of the LPSO phase, up to 50 µ m
in length (particles of Type 1 in Figure 7). These particles were present in the structure of
the alloy before rolling. With an increase in the reductions (E12–E18), large particles of
the LPSO phase were elongated in the direction of rolling, and fine particles of the LPSO
phase, 0.5–2 µm in size, precipitated in the structure (particles of Type 2 in Figure 7).
Figure 7. SEM images of the microstructure of the WZM321 alloy after rolling at 450 °C to different
reductions: (a)–E6, (b)–E12, (c)–E18.
3.2. Mechanical Properties
3.2.1. Microhardness Measurement
Table 4 shows the average microhardness values of the studied alloys depending on
the rolling temperature and the reduction ratio. It can be seen that an increase in the
rolling temperature does not have a strong effect on the change in microhardness. At the
same rolling temperature (450 °C), the average (for all reductions) microhardness in-
creases in the following order: the WZM111 alloy, the WZM211 alloy, and the WZM321
alloy, 75, 85, and 91 HV, respectively.
Table 4. Microhardness (HV) of the studied alloys after rolling.
Reduction ratio
Rolling Temperature of the Alloy
400 °С
450 °С
450 °С
450 °С
WZM111
WZM211
WZM321
Е3
77 ± 5
72 ± 5
92 ± 3
93 ± 4
Е6
76 ± 4
75 ± 3
81 ± 3
93 ± 3
Е12
80 ± 4
76 ± 2
86 ± 3
91 ± 3
Е18
73 ± 3
77 ± 2
82 ± 3
91 ± 5
Changes in the microhardness values ambiguously correlate with changes in the
microstructure. Thus, only for the WZM211 alloy, a noticeable decrease in the micro-
hardness was observed with an increase in the reduction ratio from E3 to E6–E18, which
can be associated with the development of the dynamic recrystallization process [38]. For
other alloys, with an increase in the reduction ratio from E3 to E18, the microhardness
either did not change or fluctuated slightly, which can be associated with the competition
of the hardening process caused by the accumulation of dislocations and the precipita-
tion of the α-Mn/LPSO-phase particles and the softening process caused by a decrease in
the dislocation density due to dynamic recrystallization. At the same time, it is difficult to
estimate the contribution of the dynamically recrystallized grain size to the hardening of
Figure 7.
SEM images of the microstructure of the WZM321 alloy after rolling at 450
◦
C to different
reductions: (a)–E6, (b)–E12, (c)–E18.
3.2. Mechanical Properties
3.2.1. Microhardness Measurement
Table 4shows the average microhardness values of the studied alloys depending on
the rolling temperature and the reduction ratio. It can be seen that an increase in the rolling
temperature does not have a strong effect on the change in microhardness. At the same
rolling temperature (450
◦
C), the average (for all reductions) microhardness increases in
the following order: the WZM111 alloy, the WZM211 alloy, and the WZM321 alloy, 75, 85,
and 91 HV, respectively.
Table 4. Microhardness (HV) of the studied alloys after rolling.
Reduction Ratio
Rolling Temperature of the Alloy
400 ◦C 450 ◦C 450 ◦C 450 ◦C
WZM111 WZM211 WZM321
E3 77 ±5 72 ±5 92 ±3 93 ±4
E6 76 ±4 75 ±3 81 ±3 93 ±3
E12 80 ±4 76 ±2 86 ±3 91 ±3
E18 73 ±3 77 ±2 82 ±3 91 ±5
Changes in the microhardness values ambiguously correlate with changes in the mi-
crostructure. Thus, only for the WZM211 alloy, a noticeable decrease in the microhardness
was observed with an increase in the reduction ratio from E3 to E6–E18, which can be
associated with the development of the dynamic recrystallization process [
38
]. For other
alloys, with an increase in the reduction ratio from E3 to E18, the microhardness either
did not change or fluctuated slightly, which can be associated with the competition of
the hardening process caused by the accumulation of dislocations and the precipitation
of the
α
-Mn/LPSO-phase particles and the softening process caused by a decrease in the
dislocation density due to dynamic recrystallization. At the same time, it is difficult to
estimate the contribution of the dynamically recrystallized grain size to the hardening of the
alloy through the Hall–Petch relation due to the highly inhomogeneous microstructure and
dislocation structure: the presence of the second phase particles, the presence of deformed
regions with a high dislocation density, and the presence of a dynamically recrystallized
structure with a reduced dislocation density.
Metals 2023,13, 223 9 of 15
The weak sensitivity of the microhardness of different alloys to the reduction ratio
in the E6–E12 range can be associated either with the rapid activation of the dynamic
recrystallization process already at low reductions (for the WZM111 alloy) or vice versa
with the difficulty of the dynamic recrystallization process (for the WZM211 and WZM321
alloys).
3.2.2. Tensile Tests
Tables 5–8present the averaged (over three specimens) values of the mechanical prop-
erties of the studied alloys obtained in tensile testing, depending on the rolling temperature
and the reduction ratio. The typical stress–strain curves for the samples are shown in
Figures 8–11. For clarity, the histograms of the distribution of mechanical properties for
each alloy were plotted depending on the rolling temperature and the reduction ratio
(Figures 12–15).
The change in the mechanical properties correlates well with the change in the mi-
crostructure. Thus, for the WZM111 alloy, after rolling at 400
◦
C with an increase in the
reduction ratio from E6 to E12, an increase in strength was observed (especially in the
yield strength, by 30%), and with a further increase in the reduction ratio to E15–E18, the
strength decreased monotonically (by 7–10%), which is associated with more complete
dynamic recrystallization. For the WZM111 alloy, after rolling at 450
◦
C with an increase in
the reduction ratio from E6 to E12, the strength, on the contrary, first decreased (by 12–21%)
and then increased (by 6%) with a further increase in the reduction ratio to E18. At the
same time, an increase in the rolling temperature from 400 to 450
◦
C at low reductions
caused an increase in strength, and at high reductions, this caused a decrease in ductility.
Table 5. Mechanical properties of the WZM111 alloy after rolling at a temperature of 400 ◦C.
Reduction Ratio YS (MPa) UTS (MPa) RE (%) UTS/YS
E6 202.6 ±14.4 265.2 ±2.1 12.3 ±4.9 1.31
E12 262.9 ±2.6 282.6 ±1.1 12.3 ±0.6 1.07
E15 241.6 ±10.3 265.4 ±3.3 17.0 ±1.0 1.10
E18 237.1 ±16.7 263.7 ±5.3 11.7 ±1.5 1.11
Table 6. Mechanical properties of the WZM111 alloy after rolling at a temperature of 450 ◦C.
Reduction Ratio YS (MPa) UTS (MPa) RE (%) UTS/YS
E6 253.7 ±6.5 306.9 ±5.2 13.0 ±1.0 1.21
E12 222.1 ±10.0 243.2 ±7.3 13.0 ±2.6 1.09
E18 236.0 ±10.4 258.7 ±5.7 7.0 ±1.7 1.10
Table 7. Mechanical properties of the WZM211 alloy after rolling at a temperature of 450 ◦C.
Reduction Ratio YS (MPa) UTS (MPa) RE (%) UTS/YS
E6 103.6 ±20.1 120.6 ±15.4 1.5 ±1.0 1.16
E12 282.9 ±10.4 311.3 ±6.1 10.3 ±1.5 1.10
E18 261.8 ±22.2 283.0 ±15.1 7.0 ±1.0 1.08
Table 8. Mechanical properties of the WZM321 alloy after rolling at a temperature of 450 ◦C.
Reduction Ratio YS (MPa) UTS (MPa) RE (%) UTS/YS
E6 186.5 ±28.3 189.5 ±31.7 0 1.02
E12 367.5 ±11.5 386.7 ±19.3 7.0 ±1.4 1.05
E18 300.4 ±4.4 322.2 ±5.8 5.7 ±1.5 1.07
Metals 2023,13, 223 10 of 15
Metals 2023, 13, x FOR PEER REVIEW 10 of 16
Figure 8. Typical stress–strain curves of the WZM111 alloy samples after rolling at 400 °C.
Figure 9. Typical stress–strain curves of the WZM111 alloy samples after rolling at 450 °C.
Figure 10. Typical stress–strain curves of the WZM211 alloy samples after rolling at 450 °C.
Figure 8. Typical stress–strain curves of the WZM111 alloy samples after rolling at 400 ◦C.
Metals 2023, 13, x FOR PEER REVIEW 10 of 16
Figure 8. Typical stress–strain curves of the WZM111 alloy samples after rolling at 400 °C.
Figure 9. Typical stress–strain curves of the WZM111 alloy samples after rolling at 450 °C.
Figure 10. Typical stress–strain curves of the WZM211 alloy samples after rolling at 450 °C.
Figure 9. Typical stress–strain curves of the WZM111 alloy samples after rolling at 450 ◦C.
Metals 2023, 13, x FOR PEER REVIEW 10 of 16
Figure 8. Typical stress–strain curves of the WZM111 alloy samples after rolling at 400 °C.
Figure 9. Typical stress–strain curves of the WZM111 alloy samples after rolling at 450 °C.
Figure 10. Typical stress–strain curves of the WZM211 alloy samples after rolling at 450 °C.
Figure 10. Typical stress–strain curves of the WZM211 alloy samples after rolling at 450 ◦C.
Metals 2023,13, 223 11 of 15
Metals 2023, 13, x FOR PEER REVIEW 11 of 16
Figure 11. Typical stress–strain curves of the WZM321 alloy samples after rolling at 450 °C.
The change in the mechanical properties correlates well with the change in the mi-
crostructure. Thus, for the WZM111 alloy, after rolling at 400 °C with an increase in the
reduction ratio from E6 to E12, an increase in strength was observed (especially in the
yield strength, by 30%), and with a further increase in the reduction ratio to E15–E18, the
strength decreased monotonically (by 7–10%), which is associated with more complete
dynamic recrystallization. For the WZM111 alloy, after rolling at 450 °C with an increase
in the reduction ratio from E6 to E12, the strength, on the contrary, first decreased (by
12–21%) and then increased (by 6%) with a further increase in the reduction ratio to E18.
At the same time, an increase in the rolling temperature from 400 to 450 °C at low reduc-
tions caused an increase in strength, and at high reductions, this caused a decrease in
ductility.
During the tensile test of the WZM211 and WZM321 alloy specimens after rolling to
a low reduction ratio, their failure could occur even in the elastic region at low stresses
due to surface cracks formed during rolling. However, with an increase in the reduction
ratio, such cracks heal, which leads to an increase in strength, and the specimens de-
formed with the development of the neck, which indicates a high ductility of the materi-
al. With an increase in the reduction ratio from E12 to E18, the strength of the WZM211
and WZM321 alloys decreased (by 7–9% and 17–18%, respectively). In addition, at the
same rolling temperature (450 °C), the transition from the E12 to E18 reduction for all
three alloys caused a decrease in the relative elongation, i.e., ductility. A decrease in
strength can be explained by a decrease in the dislocation density as a result of a more
complete dynamic recrystallization process. The decrease in ductility can be explained by
the presence of a surface defective layer, the negative effect of which will be stronger
with a decrease in the sample’s thickness (or with an increase in the reduction ratio).
Figure 11. Typical stress–strain curves of the WZM321 alloy samples after rolling at 450 ◦C.
Metals 2023, 13, x FOR PEER REVIEW 12 of 16
Figure 12. Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 400 °C.
Figure 13. Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
Figure 14. Dependence of the YS, UTS, and RE of the WZM211 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
Figure 12.
Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 400 ◦C.
Metals 2023, 13, x FOR PEER REVIEW 12 of 16
Figure 12. Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 400 °C.
Figure 13. Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
Figure 14. Dependence of the YS, UTS, and RE of the WZM211 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
Figure 13.
Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 450 ◦C.
Metals 2023,13, 223 12 of 15
Metals 2023, 13, x FOR PEER REVIEW 12 of 16
Figure 12. Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 400 °C.
Figure 13. Dependence of the YS, UTS, and RE of the WZM111 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
Figure 14. Dependence of the YS, UTS, and RE of the WZM211 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
Figure 14.
Dependence of the YS, UTS, and RE of the WZM211 alloy on the reduction ratio after
rolling at a temperature of 450 ◦C.
Metals 2023, 13, x FOR PEER REVIEW 13 of 16
Figure 15. Dependence of the YS, UTS, and RE of the WZM321 alloy on the reduction ratio after
rolling at a temperature of 450 °C.
The most attractive balance of strength and ductility for all three alloys was achieved
after rolling with the E12–E18 reductions.
At the same rolling temperature (450 °C) and high reductions (E12-E18), the strength
increased in the following order: the WZM111 alloy, the WZM211 alloy, and the
WZM321 alloy, which correlates with their microhardness as well as with the amount of
the LPSO phase that precipitated in their structure after rolling. In this case, the change in
their relative elongation had an inverse relationship.
It is of interest to compare the achieved mechanical properties with those of other
Mg–Zn–Y alloys. For example, for a comparable reduction ratio (~90%) of hot rolling, the
yield strength of the WZM111 alloy significantly (by a factor of 1.3–2) exceeds that of
Mg–(3–9)Zn–(0.6–2)Y alloys, but it is inferior to them in ductility [21]. Accordingly, for
the WZM211 and WZM321 alloys, the strength advantage is even greater. Thus, the
achieved strength of the studied alloys exceeds the strength of other hot-rolled alloys of
the Mg–Zn–Y system. The mechanical properties of the hot-rolled WZM321 alloy with
the greatest strength can be compared with the properties of Mg–Zn–Y alloys after hot
extrusion [22,39].
The requirements for biomaterials purposed for bone fixtures are a strength higher
than 200 MPa and a relative elongation greater than 10% [3,11]. Through this study, we
were able to ascertain that the WZM111 and WZM211 alloys processed by hot rolling at
450 °C possess the provided requirements, which make them suitable for applications in
bone implants. The WZM321 alloy has a great strength performance but a low relative
elongation. Possibly annealing applied to the alloy can increase its relative elongation,
and this should be established in future research.
The thickness of the samples achieved after rolling with the E12–E18 reductions was
0.8–0.2 mm. In accordance with the literature, the lower value is close to the thickness of
barrier membranes in dental surgery [40], and the higher value is close to the thickness of
biodegradable plates [19,20]. Further research will be focused on corrosion properties
and biocompatibility investigation of Mg–Y–Zn–Mn alloys in order to choose the best
alloy composition with a low corrosion rate (<0.5 mm/year) [3,11]. In addition, it is im-
portant to study the texture, whose contribution has a strong influence on the formation
of the mechanical properties of magnesium alloys [13,15,38].
4. Conclusions
Based on the results of this study on the effects of the rolling temperature and the
reduction ratio on the structure and mechanical properties of the WZM111, WZM211,
and WZM321 alloys, it was established that:
Figure 15.
Dependence of the YS, UTS, and RE of the WZM321 alloy on the reduction ratio after
rolling at a temperature of 450 ◦C.
During the tensile test of the WZM211 and WZM321 alloy specimens after rolling to a
low reduction ratio, their failure could occur even in the elastic region at low stresses due to
surface cracks formed during rolling. However, with an increase in the reduction ratio, such
cracks heal, which leads to an increase in strength, and the specimens deformed with the
development of the neck, which indicates a high ductility of the material. With an increase
in the reduction ratio from E12 to E18, the strength of the WZM211 and WZM321 alloys
decreased (by 7–9% and 17–18%, respectively). In addition, at the same rolling temperature
(450
◦
C), the transition from the E12 to E18 reduction for all three alloys caused a decrease in
the relative elongation, i.e., ductility. A decrease in strength can be explained by a decrease
in the dislocation density as a result of a more complete dynamic recrystallization process.
The decrease in ductility can be explained by the presence of a surface defective layer, the
negative effect of which will be stronger with a decrease in the sample’s thickness (or with
an increase in the reduction ratio).
The most attractive balance of strength and ductility for all three alloys was achieved
after rolling with the E12–E18 reductions.
At the same rolling temperature (450
◦
C) and high reductions (E12–E18), the strength
increased in the following order: the WZM111 alloy, the WZM211 alloy, and the WZM321
alloy, which correlates with their microhardness as well as with the amount of the LPSO
phase that precipitated in their structure after rolling. In this case, the change in their
relative elongation had an inverse relationship.
Metals 2023,13, 223 13 of 15
It is of interest to compare the achieved mechanical properties with those of other
Mg–Zn–Y alloys. For example, for a comparable reduction ratio (~90%) of hot rolling,
the yield strength of the WZM111 alloy significantly (by a factor of 1.3–2) exceeds that
of Mg–(3–9)Zn–(0.6–2)Y alloys, but it is inferior to them in ductility [
21
]. Accordingly,
for the WZM211 and WZM321 alloys, the strength advantage is even greater. Thus, the
achieved strength of the studied alloys exceeds the strength of other hot-rolled alloys of
the Mg–Zn–Y system. The mechanical properties of the hot-rolled WZM321 alloy with
the greatest strength can be compared with the properties of Mg–Zn–Y alloys after hot
extrusion [22,39].
The requirements for biomaterials purposed for bone fixtures are a strength higher
than 200 MPa and a relative elongation greater than 10% [
3
,
11
]. Through this study, we
were able to ascertain that the WZM111 and WZM211 alloys processed by hot rolling at
450
◦
C possess the provided requirements, which make them suitable for applications in
bone implants. The WZM321 alloy has a great strength performance but a low relative
elongation. Possibly annealing applied to the alloy can increase its relative elongation, and
this should be established in future research.
The thickness of the samples achieved after rolling with the E12–E18 reductions was
0.8–0.2 mm. In accordance with the literature, the lower value is close to the thickness of
barrier membranes in dental surgery [
40
], and the higher value is close to the thickness
of biodegradable plates [
19
,
20
]. Further research will be focused on corrosion properties
and biocompatibility investigation of Mg–Y–Zn–Mn alloys in order to choose the best alloy
composition with a low corrosion rate (<0.5 mm/year) [
3
,
11
]. In addition, it is important
to study the texture, whose contribution has a strong influence on the formation of the
mechanical properties of magnesium alloys [13,15,38].
4. Conclusions
Based on the results of this study on the effects of the rolling temperature and the
reduction ratio on the structure and mechanical properties of the WZM111, WZM211, and
WZM321 alloys, it was established that:
(1) All alloys were successfully rolled at 450
◦
C to reductions up to 97%, while only
the WZM111 alloy was successfully rolled at 400 ◦C;
(2) During hot rolling, the dynamic recrystallization process occurs, and this pro-
cess proceeds most easily and with most difficulty in the WZM111 and WZM321 alloys,
respectively;
(3) During hot rolling, the LPSO phase particles precipitate in the structure of the
WZM211 and WZM321 alloys;
(4) At high reductions (96–97%), an increase in the rolling temperature from 400 to
450
◦
C has little effect on the strength of the WZM111 alloy but reduces its ductility. At
the same rolling temperature (450
◦
C) and at high reductions, the strength increases in
the following order: the WZM111 alloy, the WZM211 alloy, and the WZM321 alloy, while
ductility has an inverse relationship;
(5) The most attractive balance of strength and ductility for all three alloys is achieved
after rolling to total reduction ratios of 88–97%. The achieved strength of all three alloys
exceeds the strength of other hot-rolled alloys of the Mg–Zn–Y system, but they are inferior
to them in ductility, while the mechanical properties of the hot-rolled WZM321 alloy with
the greatest strength can be compared with the properties of Mg–Zn–Y alloys after hot
extrusion.
Author Contributions:
Conceptualization, V.E.B., A.A.K. and A.Y.D.; methodology, V.E.B. and S.O.R.;
investigation, S.O.R., A.V.L., D.V.T. and V.V.Y.; writing—original draft preparation, S.O.R.; writing—
review and editing, V.E.B.; visualization, S.O.R.; funding acquisition, K.S.S. All authors have read
and agreed to the published version of the manuscript.
Funding:
The authors are grateful to the Ministry of Science and Higher Education of the Russian
Federation for financial support under the Megagrant (No. 075-15-2022-1133).
Metals 2023,13, 223 14 of 15
Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.
Data Availability Statement:
The raw/processed data required to reproduce these findings cannot
be shared at this time as the data also forms part of an ongoing study.
Acknowledgments: We thank N. Munzaferova for help with research.
Conflicts of Interest: The authors declare no conflict of interest.
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