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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