high strength mg-zn-y alloys containing quasicrystalline particles
TRANSCRIPT
High strength Mg-Zn-Y alloys containing quasicrystalline particles
D. H. Bae and D. H. Kim Center for Non-crystalline Materials, Department of Metallurgical Engineering, Yonsei University,
134 Shinchon-dong Seodaemun-gu, Seoul, 120-749, Korea
Abstract
High strength Mg-rich Mg-Zn-Y(-Zr) alloys, strengthened by
quasicrystalline particles, have been developed by
thermomechanical processes. The deformation behaviors of these
alloys at room and high temperatures have been investigated and
compared to those of commercial alloys AZ31, AZ61 and AZ91.
Yield strengths of the Mg-Zn-Y alloys, which increase with an
increase in the volume fraction of quasicrystalline phase, are
relatively high due to the strengthening effect of quasicrystalline
particles. At high temperatures (300 – 425oC), the flow stress
levels of the Mg-Zn-Y alloys are lower than those of AZ31, AZ61
and AZ91, while their elongations to failure are larger.
Quasicrystalline particles in the Mg-Zn-Y alloys resist coarsening
during deformation at temperatures up to the eutectic temperature,
leading to a stable distribution of quasicrystalline particles.
Stability of both quasicrystalline particles and microstructure of
the α-Mg matrix in the Mg-Zn-Y alloys provides large elongation
with no void opening at the interface between the quasicrystalline
phase and the α-Mg matrix.
I. Introduction
Significant interest currently exists in the development of
wrought magnesium alloys with high strength, high corrosion
resistance, and excellent formability at elevated temperatures for
applications as structural parts. The typical wrought alloys, for
example, AZ31, AZ61 and ZK60, normally exhibit the moderate
strength at ambient temperature and poor creep resistance [1]. For
the purpose of the practical usage of a wrought Mg alloy as
structural parts, however, its mechanical properties of strength,
fracture toughness, and creep resistance should be further
improved in a low temperature regime, typically up to 200oC, and
final products can be easily fabricated with low-cost using
conventional forming technologies at elevated temperatures.
Recently, it was reported that as-cast Mg-rich Mg-Zn-Y
alloys (produced by permanent mould casting), which were
consisted of a thermally stable icosahedral quasicrystalline phase
(I-phase) in-situ formed as a second phase of the eutectic in the α-
Mg matrix during solidification, exhibited yield stress from 180 to
480MPa at room temperature, depending on the volume fraction of
I-phase [2]. Quasicrystals are isotropic and posses specially
ordered lattice structure called the quasiperiodic lattice structure
[3]. When an alloy possesses quasicrystals as a second phase, they
are stable against coarsening at high temperatures due to the low
interfacial energy of quasicrystals [4], providing the improved
bonding properties in the I-phase/matrix interface. The
disadvantage of the low eutectic temperature of a Mg-Zn binary
alloy (~ 340oC) was also surmounted in the Mg-Zn-Y alloy since
small amount of Yttrium could increase the eutectic temperature of
the Mg-Zn-Y alloy significantly [5].
In this study, several Mg-rich Mg-Zn-Y alloys, each
containing different amount of I-phase, were selected in the
Magnesium Technology 2002Edited by H.I. Kaplan
TMS (The Minerals, Metals & Materials Society), 2002
compositional range up to the eutectic Mg73.2Zn23Y3.8 (where the
ratio of Zn to Y is around 6) since these alloys can be rolled at
high temperatures. We utilized the conventional thermomechanical
processes, i.e. hot-rolling and annealing, for the as-cast alloys.
During the hot-rolling process, as-cast I-phase structure formed in
the interdendritic region was broken to be changed to small
particles distributed in the α-Mg matrix. Uniaxal tension tests
were performed on alloy sheets at room and high temperatures,
and the effect of I-phase particles on flow behavior and correlated
microstructural evolution during deformation was investigated.
2. Materials and Experimental Procedure
For this study, test materials were Mg96Zn3.4Y0.6 and
Mg95Zn4.3Y0.7 alloys in atomic percentage. A Mg94.8Zn4.3Y0.7Zr0.2
alloy was also prepared since an addition of small amount of Zr in
the Mg-Zn-Y alloy may provide finer solidified microstructure [6].
For comparison, commercial alloys, AZ31 and AZ61, were
provided by Korea Institute of Industrial Technology in the as-cast
condition. Also an AZ91 alloy was provided in the hot-extruded
condition. Each alloy was separately prepared by melting from
high purity 99.9%Mg, 99.9%Zn, 99.95%Y and 99.9%Zr under a
dynamic air/carbon dioxide/sulfur hexafluoride atmosphere. The
as-cast Mg-Zn-Y(-Zr) alloys were consisted of two phases of α-
Mg and I-phase of Mg3Zn6Y1 [7] as shown in Fig. 1 (a) of a SEM
micrograph of a Mg94.8Zn4.3Y0.7Zr0.2 alloy, in which eutectic
pockets were seen. All alloys except for an AZ91 alloy were hot-
rolled from 12 mm to 1 mm final thickness (reduction ~ 92%). An
AZ91 alloy was hot-rolled 5 mm to 1 mm final thickness
(reduction ~ 80%). The rolls were heated to around 100oC prior to
rolling. The block samples were heated at 400oC for 10min and
then rolled with a reduction of ~ 15% per one pass. The sheets
were annealed at 400OC for 0.5 hour in an air circulating furnace.
For Mg-Zn-Y(-Zr) alloys, the as-cast eutectic pocket structure was
destroyed during rolling, providing the distribution of particles
(0.5 –2.0 µm in size) in the α-Mg matrix as shown in Fig. 1(b) of a
SEM image of a hot-rolled Mg94.8Zn4.3Y0.7Zr0.2 alloy, in which the
structure of I-phase remains [8]. The test materials were cut in
three orthogonal sections, mechanically polished, and then etched
with an etchant (1 ml HNO3, 24 ml water, and 75 ml ethylene
glycol). The α-Mg grains developed during hot-rolling via
dynamic recrystallization (DRX) process were found to be
equiaxied in all alloys, and the linear intercepted grain sizes of the
alloys were listed in Table 1. Fine grained structure of an initial
grain size (DO) of 7.7µm was developed in Mg-Zn-Y and Mg-Zn-
Y-Zr alloys due to the effect of large amount of particles in the
DRX process. Grain size of an AZ91 alloy was found to be larger
than that of an AZ61 alloy. This may be due to the lower reduction
amount of the alloy in the hot- rolling process.
Fig. 1. Secondary electron micrographs of a Mg94.8Zn4.3Y0.7Zr0.2
alloy: (a) as-cast; (b) hot-rolled on the L-T plane, where L and T
refer to the longitudinal and transverse directions, respectively.
The volume fraction of I-phase, measured by an image analysis
method, was around 8% for an Mg96Zn3.4Y0.6 alloy (see Table 1).
During continuous heating at a heating rate of 0.67 K/s in
differential thermal analysis (DTA), first endothermic peak
appeared with onset temperature of around 440oC for Mg-Zn-Y
and Mg-Zn-Y-Zr alloys as shown in Fig. 2. This peak was
recognized as the melting of the eutectic pockets of the alloys. The
finishing temperature of the second endotherm was around 640oC.
Uniaxial tensile tests were carried out on dog-bone
specimens of the hot-rolled sheets (specimen gauge length = 7mm)
under a constant cross-head speed condition of an initial strain rate
of 10-3s-1 at room temperature and under a constant strain rate
condition at high temperatures up to 425 OC in the strain rate range
of 10-4s-1 and 10-1s-1. Cross-head speed was controlled by
computer through a digital interface board on an Instron-type
machine. Before the test, the load train was preheated to the test
temperature within a clamshell furnace having three heating zones
independently controlled, and then the test was performed by
(b)
T
L
(a)
300 400 500 600 700-1.0
-0.5
0.0
0.5
1.0
1.5
2.0As-cast
Mg94.8
Zn4.3
Y0.7
Zr0.2
Mg96Zn3.4Y0.6
∆T (
OC
)
Temperature, OC
Fig. 2. DTA traces showing melting endotherms obtained during
heating at a heating rate of 0.67 K/s for Mg96Zn3.4Y0.6 and
Mg94.8Zn4.3Y0.7Zr0.2 alloys.
placing the specimen into the load train. Typically the heating-
plus-holding time prior the test was around 10 min. The
fluctuation of the temperature during the test was ±1 oC.
Microstructures of the tested specimen were observed by optical
microscope (OM) and scanning electron microscope (SEM). The
α-Mg grain size was determined by the linear intercept method. To
understand the role of the I-phase during deformation, the
interface between I-phase and the matrix was carefully examined
using TEM. Thin foil specimens for TEM observation were
prepared by an ion beam milling method and were observed in a
JEOL 200CX microscope.
3. Results and Discussion
3.1 Mechanical behavior at room temperature
The mechanical properties of 0.2 percent yield stress (σ0.2),
ultimate tensile strength (UTS) and elongation to failure are listed
in Table 1. The Mg96Zn3.4Y0.6 alloy shows the high level of yield
stress around 210MPa and UTS 355MPa, similar to those of an
AZ91 alloy, and large elongation around 23 percent. The large
elongation observed in the large number of I-phase particles
reinforced Mg-Zn-Y alloy is unusual. Generally, elongation is low
for the alloy containing large amount of intermetallic particles
since geometrically necessary dislocations are formed in the
region surrounding the hard particle, eventuating in decohesion
from the matrix [9]. The Mg-Zn-Y alloys can be strengthened by
large number of I-phase particles and solute atoms of Zn and Y in
the α-Mg matrix. However, 18% drop in yield stress for the
Mg94.8Zn4.3Y0.7Zr0.2 alloy, compared to that of the Mg95Zn4.3Y0.7
alloy having similar microstructure, indirectly infers that the
strengthening effect of I-phase is only moderate. A Zr addition in
the Mg-Zn-Y alloy may decrease the solubility of alloying
elements in the α-Mg matrix, softening the alloy. Further studies
are necessary to clarify the compositional difference in the α-Mg
matrix between two different alloys. Thus, we believe that the
improvement of mechanical properties by I-phase particles is
somewhat different from those by intermetallic particles (i.e.
Mg17Al12) in the Mg-Al-Zn alloys. The quasi-periodic lattice
structure of I-phase provides the stable I-phase particle/matrix
interface [4], which may be achieved by accommodating the
lattice spacing between them, providing the lower lattice
mismatching strain in the α-Mg matrix.
3.2 Flow behavior and microstructural evolution during
deformation at high temperatures
Typically for magnesium alloys the extensive
microstructural evolution by DRX or grain growth is occurred at
high temperatures, depending on the deformation mechanism [10].
These processes influence the stress level during deformation and
formability of the alloy. Stress vs strain curves at a strain rate of
10-3s-1 and a temperature of 300oC are shown in Fig. 3 for AZ31,
AZ91, and Mg96Zn3.4Y0.6 alloys. The I-phase reinforced alloy
exhibits the lower level of yield stress but higher elongation. The
stress level is dependent on the strain. Typically the alloys
containing second phase particles reach a peak at low strain and
then decrease continuously with increasing strain. However, the
AZ31 alloy exhibits strain hardening in the low strain range and
Table 1. Properties of tested Mg alloys at room temperature a.
Alloy Grain Size (DO), µm Vol. Fraction of
Second Phase
Yield Stress
(σ0.2), MPa
UTS,
MPa
Elongation to
failure, %
AZ31 17.6 152 275 22.0
AZ61 9.9 175 320 19.8
AZ91 13.4 225 395 18.2
Mg96Zn3.4Y0.6 7.8 0.08 210 355 23.4
Mg95Zn4.3Y0.7 7.7 0.09 220 370 19.7
Mg94.8Zn4.3Y0.7Zr0.2 7.8 0.09 180 325 23.5 a Specimen dimension for tensile testing: gauge thickness, 1.0mm; gauge length, 7mm. Uniaxial tension test: initial strain rate, 10-3s-1.
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 1.1 1.20
10
20
30
40
50
60
70
80
90
100
Temp. = 300OC
Strain Rate = 1x10-3s
-1
Mg96
Zn3.4
Y0.6
AZ31
AZ91
Stre
ss (
MPa
)
Strain
Fig. 3. Stress vs strain curves for four different Mg alloys of AZ31,
AZ91, and Mg96Zn3.4Y0.6 obtained from the uniaxial tension test
conducted at a strain rate of 1x10-3s-1 and a temperature of 300OC.
then the flow stress decreases with increasing strain. This
incubation strain may be necessary for dynamic recrystallization in
the AZ31 alloy since the nucleation site of recrystallization is
limited only to the grain boundaries. The AZ91 alloy shows the
significant stress drop mainly due to the extensive DRX process
(i.e. grain refinement from 13.4 to 7µm). The intermetallic
particles of Mg17Al12 in the AZ91 alloy can effectively provide the
site for an initiation of recrystallization. However, the variation of
flow stress for large amount of an I-phase particle reinforced alloy
is much low. Since the particles are closely spaced in the Mg-Zn-Y
alloy, particles can exert a significant pinning effect on both low
and high angle grain boundaries, hindering recrystallization [11].
To further investigate the temperature and strain rate effect
on the flow behavior and microstructural evolution, stress vs strain
curves at a test temperature of 425oC and strain rates of (a) 10-1s-1
and (b) 10-3s-1 are plotted in Fig. 4 for AZ31, AZ61, Mg96Zn3.4Y0.6
and Mg94.8Zn4.3Y0.7Zr0.2 alloys. The concurrent variation of
average grain size (D) as a function of strain for largely strained
samples can be related to the following form [12-13].
D = DO + β ε (1)
where β is a constant and ε is a strain. Based on Eq. (1), Fig. 5
shows the β values normalized initial grain size, (1/DO)(dD/dε),
are plotted as a function of strain rate for different alloys.
At a high strain rate of 10-1s-1, I-phase reinforced alloys also
exhibit the lower level of yield stress but higher elongation. The
Mg-Zn-Y alloy is found to be no grain elongation to loading
direction but only a grain refinement as shown in Fig. 5. However,
the flow stress slightly increases with increasing strain. This may
be due to the particle pinning effect, decreasing the contribution of
grain boundary sliding to total elongation. For the AZ31 and AZ61
alloys, however, strain hardening is observed in the low strain
range and then the flow stress decreases with increasing strain.
This stress drop in the high stain range is found to be due to not
recrystallization but necking. Both alloys exhibit an abnormal
grain growth in this test condition as shown in
0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.40
10
20
30
40
50
60
70
80
90
100(a)
Mg94.8
Zn4.3
Y0.7
Zr0.2
Mg96
Zn3.4
Y0.6
AZ31AZ61
Temp. = 425OC
Strain Rate = 1x10-1s
-1
Stre
ss (
MPa
)Strain
0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.40
5
10
15
20
25
30
35
40
45
50
Mg94.8
Zn4.3
Y0.7
Zr0.2
Mg96
Zn3.4
Y0.6
AZ61AZ31
Temp. = 425OC
Strain Rate = 1x10-3s
-1
(b)
Str
ess
(MP
a)
Strain
Fig. 4. Stress vs strain curves at two different strain rates of (a) 10-
1s-1 and (b) 10-3s-1 at a test temperature of 425OC for four different
Mg alloys of AZ31, AZ61, Mg96Zn3.4Y0.6 and Mg94.8Zn4.3Y0.7Zr0.2.
Fig. 5, providing only small uniform elongation. Fine grained Mg-
Zn-Y(-Zr) alloys showing large elongation under the high strain
rate deformation condition is very attractive in industry for the
fabrication of complex parts by the practically used forming
technologies.
At a strain rate of 10-3s-1, typically grains grow dynamically
with increasing strain for all alloys as shown in Fig. 5. However,
10-4
10-3
10-2
10-1
100
-1.0
-0.5
0.0
0.5
1.0
1.5
Abnormal grain growth
Grain Refinement
Grain Growth
Mg96Zn3.4Y0.6
AZ31
AZ61
(1/D
O)
dD /
dε
Strain Rate, s-1
Fig. 5. Grain size variation, (1/DO) dD/dε, plotted as a function of
strain rate at a test temperature of 425OC for three different Mg
alloys of AZ31, AZ61 and Mg96Zn3.4Y0.6.
the variation of the stress level is relatively low for the AZ31 alloy
having larger grains, compared to that for the AZ61 alloy
exhibiting superplastic deformation (which is evinced by no
diffuse necking in the deformed specimen and void formation at
the grain boundaries), since the contribution of grain boundary
sliding to total deformation is relatively low for the large grained
alloy [14]. The Mg-Zn-Y alloy shows slight grain growth with no
grain elongation to the load direction, providing a slight increase
in the flow stress during deformation. The most interesting thing
observed in the Mg-Zn-Y alloys is that elongation is quite large
with no void formation, comparable to that in the superplastic
AZ61 alloy containing large number of voids at the grain
boundaries. Due to the pinning effect of large number of I-phase
particles in the fine grained Mg-Zn-Y alloys, the contribution of
grain boundary sliding to total elongation may be weak, providing
no void formation, but the large elongation with the maintenance
of the equiaxied grain shape can be achieved by grain boundary
diffusion with the relatively lower grain growth processes. In
addition, void formation at the particle/matrix interface may be
difficult since the eutectic can be softened at this high temperature.
3.4 Role of quasicrystals in the improvement of mechanical
properties
To understand the role of I-phase particles in the alloy at
high temperature, microstructures of a deformed specimen were
examined. Figs. 6(a) and (b) show, respectively, a SEM image and
a bright-field TEM image of the Mg95Zn4.3Y0.7 alloy deformed to a
strain of around 1.0 with a strain rate of 10-4s-1 at a test
temperature of 400oC. The flow stress of the alloy was constant (~
10 MPa) over the strain range. The particles initially distributed
near the eutectic pocket before testing (see Fig. 1(b)) move away
each other during deformation, being distributed more randomly in
the test specimen. Surprisingly, the particles do not coarsen, but
only the shape of the particle, initially faceted, evolves to be
somewhat rounded as shown in Fig 6(b). The inset in Fig. 6(b)
shows the diffraction pattern taken from the particle, in which a 5-
fold symmetry can be identified as an I-phase. The existence of I-
phase in the deformed specimen clearly indicates that I-phase
thermally equilibrates with α-Mg phase. Furthermore, any
debonding or nanoscale defect at the particle/matrix interface
cannot be seen in the test specimen. In general, particle/matrix
debonding has been considered as an initial stage of the failure
mechanism at high temperature deformation in the alloy consisting
of intermetallic compound particles at grain boundaries, and the
coalescence of many cavities at large strains can induce the
Fig. 6. (a) SEM on the L-T plane and (b) bright-field TEM images
of a I-phase particle in the Mg95Zn4.3Y0.7 alloy deformed to a strain
of 1.0 at the test condition of a strain rate of 10-4s-1 and a test
temperature of 400 OC. An electron diffraction pattern of a particle
is inserted in (b).
(a)
L
T
(b)
failure of such an alloy. For example, an intermetallic particle of
Mg17Al12 in an AZ91 alloy [15]. Many defects have been reported
to develop in the particle/matrix interface during the
thermomechanical processing possibly due to the mismatched
lattice structure of the particle to the matrix. Sometimes these
defects can grow rapidly due to the constrained material flow near
the particle at high temperature [16]. However, careful
examination of the I-phase/matrix interface in the Mg95Zn4.3Y0.7
alloy does not reveal any defects. The structure of I-phase in Mg-
Zn-Y alloys is the face-centered I-phase [17], but the details on
atomic structure are clearly unknown yet. However, the I-phase
exhibits more isotropic characteristic than crystalline particles due
to its high symmetric structure. This may provide the reasonably
stable bonding with low strain energy at the matrix adjacent to the
I-phase particle. In general, the matrix adjacent to the microscale
intermetallic particle is highly stressed due to the mismatched
lattice constants between the matrix and the particle. However, due
to the quasiperiodic lattice structure of I-phase, the mismatched
strain may be compromised by the I-phase particle, decreasing the
stress concentration in the matrix near the I-phase particle. Within
our knowledge, the diffusivity of Y in Mg is not available.
However, low diffusivity is expected by considering large
difference in atomic size between Mg and Y. Furthermore, the low
interfacial energy of quasicrystals reduces the driving force for
coarsening the particle, leading to the stable size of the particles
[4]. These may be responsible for the negligible coarsening of I-
phase during high temperature deformation as shown in Fig.6(a).
4. Conclusions
In summary, fine grained magnesium alloys reinforced by
quasicrystalline particles were easily developed by
thermomechanical processes for as-cast Mg-rich Mg-Zn-Y and
Mg-Zn-Y-Zr alloys. An addition of Zr in the Mg-Zn-Y alloy does
not influence the hot-rolled microstructure but decreases the
solubility of the alloying elements. The mechanical behaviors of
the alloys at room and high temperatures was investigated and
compared to those of commercial alloys AZ31, AZ61 and AZ91.
Yield strengths of the Mg-Zn-Y alloys are relatively high due to
the strengthening effect of quasicrystalline particles. But their
strengthening effect is moderate. At high temperatures, the levels
of flow stress of the Mg-Zn-Y alloys are lower than those of
commercial magnesium alloys due to the softness of the eutectic
region and finer grained structure. But the alloys exhibit much
higher elongation at high strain rates and/or at low temperatures
since large number of quasicrystalline particles in the Mg-Zn-Y
alloys can effectively prohibit against microstructural evolution of
the α-Mg matrix during deformation. Furthermore, I-phase
particles in the Mg-Zn-Y alloys are stable against coarsening
during deformation near the melting temperature of the eutectic,
forming the stable quasicrystalline particle/matrix interface. Those
stabilities provide large elongation with no void opening at the
interface between the quasicrystalline particle and the α-Mg
matrix. The observed mechanical properties of high strength at
room temperature and large elongation at high temperatures can
open interesting perspectives for the applications of such
quasicrystalline paticle reinforced composite materials as sheet
components.
Acknowledgements
This work was funded by Creative Research Initiatives of the
Korea Ministry of Science and technology.
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