EFFECT OF Y ADDITION
ON MG-ZN RAPIDLY SOLIDIFIED ALLOYS
Xuefeng Guo*, and Dan Shechtman**
* Department of Materials Science and Engineering, Xi'an University of
Technology, Xi'an 710048, Shaan Xi, People Republic of China.
** Department of Materials Engineering, Technion, Haifa 32000, Israel
1. Introduction
The development of new materials is usually driven by a need
for improved properties. In recent years the need of weight reduction in
the transportation and aircraft industries brought new-life to strong Mg
alloys for use in structural applications [1-3], since magnesium alloys, the
lightest structural materials [4, 5] have met many of the requirements.
This includes high specific strengths [1-9], machinability [10, 11], good
casting qualities [12] and high damping capacity [13].
It has been established that melt spinning is a convenient way to
rapidly solidify (RS) alloys [14], and melt spun ribbons were used in
many studies to investigate both the microstructure and mechanical
properties of a wide variety of alloys [14]. RS enables to modify and
make novel metastable microstructures, which can not be made by
conventional casting. These microstructures are characterized by [14]:
(1) Substantial extensions of the solid solubility, beyond
equilibrium limits,
(2) A more uniform microstructure, with less segregation and
finer distribution of impurities and defects,
(3) Formation of new metastable phases including crystalline,
quasicrystalline or amorphous phases, and
(4) Refined grain sizes.
These microstructural changes provide opportunities to produce
a wide variety of novel materials that exhibit improved mechanical
properties and chemical stability.
Using melt spinning to develop high-strength (ultimate tensile
strength higher than 500 MPa [6]) magnesium alloys by microstructural
refinement or even glass formation has been the focus of many research
programmers [1-3, 5-9, 15-21]. Although a vast effort was made to develop
other technologies [22-25], one of the best ways to get high mechanical
properties and improved corrosion resistance is RS processing.
Interesting achievements [6] are shown in Table 1-1. Strengthening
267
mechanisms including Orowan processes [26], deposition strengthening
[26]
, and fine-grain strengthening [6, 26] all contributing to the high
strength, were used to explain the novel results obtained.
Reviewing these studies, one can conclude that a future trend for
development of Mg could be RS of alloys containing RE and Y [1-3, 5, 9,
11,15, 18-19,21,26-29]
, since additions of RE and Y increase the mechanical
properties both at room temperature and at elevated temperatures [4, 11],
and simultaneously improve their corrosion resistance [15]. New
magnesium alloys, such as Mg-Gd-Nd-Zr, Mg-Dy-Nd-Zr and Mg-Gd-YZr alloys [30], Mg-Sc[31], Mg-Sc-Mn[31] and Mg-Y-RE[11, 27-29] have been
thoroughly studied, and in particular, Mg-Y-RE (WE series). Y and RE
elements are considered the most useful additives to magnesium
systems. Such alloys can work at temperatures above 250 oC [32-33].
Table 1-1.
High strength magnesium alloys produced by of RS and extrusion
Alloy composition,
wt%
s/
MPa
Mg-18Al-10Ca
b/
MPa
/%
E/
GPa
720
Mg-20Al-15Ca
600
Mg-21Y-10Cu
724
Mg-8Al-5Ga3Zn
518
646
1.2
Mg-9Al-6.5Ca
538
567
3.2
Mg-9Al-5Ca
496
565
Mg-9Al-1Zn1.5Si
468
541
Mg-6Zn-4Ce0.4Zr
710
3.4
506
2
Mg-5Al-5Zn5Nd
EA55RS
476
513
5.0
164
2.21
330
48
1.91
168
1.78
1.78
1.82
46
124
(b/)/
MPam3Mg-1
400
50
46
455
1.81
(s/)/
MPam3Mg-1
330
5.4
Mg-11Si-4Al
H/
kgmm-2
1.81
41
554
/
Mg m-3
270
340
300
320
280
320
260
300
1.88
300
1.8
250
280
1.94
250
260
s tensile yield strength; b ultimate tensile strength;  elongation to
fracture; E Yong’s modulus;  density; H hardness.
268
390
However, the effects of alloying additives of RE elements and Y
in Mg-Zn base alloys, especially in the initial stages of the processing,
i.e. solidification, remain unclear. It is important to understand the
alloying process because Mg-Zn base alloys have excellent mechanical
properties at room temperature, and when Y is added, also at elevated
temperatures. In addition, thermal-shock and micro-shrinkage of Mg-Zn
based alloys can be overcome by RS process. Therefore, the aim of this
study is to optimize the performance of such alloying elements by
developing a clear understanding of their effect on hardening and
strengthening mechanisms. This article describes part of a study
performed by the authors at the Technion during 2001-2002 when XG.
was on Sabbatical at the Technion from Xian University.
2. Experimental
2.1 Chill Casting
Based on the commercial alloys ZK31 and ZK60 [34] which
contain 3%Zn (in weight per cent hereafter, unless indicated otherwise)
and 6%Zn respectively, two groups of Mg-Zn based alloys containing
different levels of Y, Ce and Zr were prepared by chill casting and melt
spinning, as listed in Table 2-1. The raw materials for making the alloys
were Mg ingots of 99.95% purity, Zn bars of 99.95% purity, and Mg47%Y, Mg-90%Ce and Mg-30 %Zr master alloys.
Initial melting of billets, 700 gr each, were carried out in low
carbon steel crucibles in a resistance furnace under protection of S6F and
CO2 gas mixture, to avoid excessive oxidation. Melting temperatures
were between 760 oC and 800 oC for the different alloys. Ingots were
made by chill casting into copper molds 60 mm in diameter and 150 mm
high. Both melting and casting were processed. Each BCI was then cut
into several pieces, weighing ~70 g, each and remelted in carbon
crucible covered by pre-heated salt-mixture of 50%KCl and 50%NaCl.
The remelted alloys were made into small ingots (SCI) by chill casting
into low carbon steel molds of 22 mm in diameter and 110 mm high.
269
Table 2-1
Experimental alloys
Grou
p
Designat
ion
A1
I
A2*
A3*
B1
B2
B3
B4
B5
II
B6
Cast
state
CS
RS
CS
CS
CS
RS
CS
RS
CS
RS
CS
RS
CS
RS
CS
RS
CS
B7
RS
CS
B8*
RS
B9*
RS
EDS analyzed composition
at. %
wt. %
Mg-1Zn0.4Y
Mg-2.7Zn1.3Y
Mg-1.4Zn0.3Y
Mg-3.6Zn1Y
Mg-1.7Zn0.3Y0.14Zr
Mg-4.5Zn1Y0.5Zr
Mg-1.2Zn0.4Ce0.14Zr
Mg-3Zn2Ce0.5Zr
Mg-2.1Zn
Mg-5.4Zn
Mg-2.3Zn
Mg-6Zn
Mg-2.5Zn0.3Y
Mg-6.2Zn1.3Y
Mg-2.4Zn0.3Y
Mg-6Zn1Y
Mg-2.1Zn1Y
Mg-5.4Zn3.5Y
Mg-2.5Zn1Y
Mg-6.3Zn3.2Y
Mg-2Zn1.4Y
Mg-5Zn4.7Y
Mg-2.2Zn1.3Y
Mg-5.4Zn4.3Y
Mg-2.6Zn0.4Y0.2Ce
Mg-6.3Zn1.5Y1Ce
Mg-2.7Zn0.4Y0.2Ce
Mg-6.5Zn1.4Y0.7Ce
Mg-1.9Zn0.4Y0.2Ce
Mg-4.8Zn1.5Y1Ce
Mg-1.9Zn0.4Y0.2Ce
Mg-4.8Zn1.5Y1Ce
MgMg-5Zn1.3Y0.4Ce0.8Zr
2.0Zn0.4Y0.1Ce0.2Zr
MgMg-5.3Zn1.5Y0.4Ce0.7Zr
2.1Zn0.4Y0.1Ce0.2Zr
MgMg-5.5Zn0.7Y0.6Ce0.6Zr
2.1Zn0.2Y0.1Ce0.2Zr
MgMg-6.0Zn1.0Y0.6Ce0.6Zr
2.2Zn0.3Y0.1Ce0.2Zr
MgMg-5.5Zn1.7Y1Ce0.6Zr
2.2Zn0.5Y0.2Ce0.2Zr
*Samples for mechanical property tests were made from those alloys, CS
means chill casting,
2.2 RS
Rapidly solidified ribbons, 3 mm wide and 80 μm thick were
produced by single-roller melt spinning under low-pressure argon
environment from SCI. Melt-spinning temperature was between 700 oC
and 730 oC according to alloy composition. The estimated cooling rate
was no less than 106 oC/s.
270
2.3 Extrusion
Two extrusion procedures were made, of the RS, comminuted
and canned ribbons and of the regular cast billets.
In preparation for the extrusion of the RS alloys, ribbons were
comminuted and canned in cylinders 30 mm in diameter. The canned
powder was degassed in a vacuum chamber for 1 h, under ~ 10-5 Pa,
heated to 150oC and pressed to form pre-extrusion bulk. The preextrusion bulk was then extruded at 300 oC via a die to form a 12 mm
diameter round rod. The extrusion processed was performed at a
constant strain rate of 0.5 min-1.
Prior to extrusion the as-cast ingot it was machined to a cylinder
30 mm in diameter and 82 mm in length, and then extruded at 300 oC. at
strain rate of 0.5 min-1.
2.4 Hear Treatment
Heat treatments were performed on coupons cut from both
extruded and as-cast bars. The coupons were encapsulated in quartz
tubes that had been backfilled with high purity argon after evacuation.
As-cast samples were heat-treated for 100 h at 300 3 oC, extruded
samples were heat-treated for 1 h at 100 oC. The temperature was
monitored by high precision equipment.
2.5 Microstructural Analysis
The alloys at different stages were analyzed by XRD with CuKα radiation operating at 40 KV and 40 mA on a Philips X-ray
diffractometer.
For microstructural study, ribbon samples were mounted edgeon (along longitudinal cross section) in cold-setting resin. All the
samples at different experiment stage were abraded to 1000 grit,
polished to 1µm, etched with 5% of HNO3 in ethanol, observed with an
optical microscope, then coated with carbon and analyzed by a SEM
equipped with EDS.
Transmission electron microscopy (TEM) specimens were made
from cross-section of the ribbons. These were mounted in cold-setting
resin within copper tubes. The tubes were then cut, polished to a
thickness less than 0.08 mm, and ion beam thinned to final thickness on
a Gatan-691 with beam energy 5.0 KeV, Characterization was
performed in a Philips CM20 TEM operating at 200 kV.
271
In order to analyze critical phase transforming temperatures,
specimens were tested on a Perkin-Elmer DTA 1700 high temperature
differential thermal analyzer (DTA) system with heating rate of 15 oC
min-1 under flowing argon.
3. Microstructure of RS Alloy Ribbons with Different Y
Content
Y addition to magnesium alloys is the most effective way to
improve creep resistance and high strength at elevated temperatures
thanks to fine and thermally stable Y containing particles [35] that can
impede dislocation motion. In this study we prepared Mg-Zn alloys with
different Y concentrations by melt spinning and evaluated their
microstructures. Following is the microstructural study.
3.1 Solidification Microstructures
Fig. 3-1 shows X-ray diffraction patterns of the ribbons. Except
for two or three peaks at low angles which cannot be identified, A1
consists of supersaturated magnesium solid solution (Mg); B2 and B3
consist of (Mg) and W phase (Mg3Y2Zn3) and B4 consists of (Mg), W
and X phase (Mg12YZn).
Fig. 3-2 shows SEM images of the ribbons. The microstructure
of A1 consists of characteristic columnar grains that originate at the
wheel side and grow through the cross section of the ribbon. In addition,
a discontinuous white layer was found at cell boundaries, as shown in
Fig. 3-2 (a) and (b). EDS analysis reveals that the white layer contains
higher Zn and Y than the average composition.
In the second alloy group, the images show that the
microstructures consist of Mg cells, submicron particles dispersed
within them and cell boundary network. The (Mg) cells appear in two
different layers along the cross section. One layer that originates at the
wheel side consists of directionally solidified columnar cells reaching 10
to 20 m in size. The cells’ axes are not perpendicular to the quenching
surface but are tilted some degrees away from it and are perpendicular to
the solidification front. The other layer that consists of fine equiaxed
grains about 5 m in diameter appears above the columnar grains. The
size, shape and distribution of the submicron particles are similar in the
second alloy group as shown in Fig. 3-2 (c) to (h). In B4, although X-ray
272
analysis (Fig. 3-1 (d)) indicate formation of X phase is difficult to be
distinguished it from W phase by SEM and EDS perhaps because of its
limited quantity. The network at grain boundaries is not the general
eutectic mixture, but rather an intermetallic compound.
2500
*
* (Mg)
? Unknown phase
Intensity
2000
1500
(a)
*
1000
*
*
500
?
**
*
*
?
?
*
*
*
*
*
*
*
*
*
*
0
20
40
60
80
*
100
*
120
2[deg.]
10000
*
Intensity
8000
* (Mg)
W Mg3Y2Zn3
? Unknown
phase
(b)
6000
4000
2000
*
?W
?
? *?
W
0
20
40
*
*
*
** *
60
*
80
*
*
* *
*
100
2 [deg.]
Fig. 3-1. X-ray diffraction patterns of alloys (a) A1; (b) B2
273
*
120
2000
1800
*
1600
1400
Intensity
(c)
* (Mg)
W Mg3Y2Zn3
? Unknown phase
*
1200
1000
800
600
200
*
*
400
*
?W
?
?
W
*
W
** *
*
*
*
* *
*
*
*
*
0
20
40
60
80
100
120
2[deg.]
3000
(d)
*
2500
* (Mg)
W Mg3Y2Zn3
X Mg12YZn
? Unknown phase
*
Intensity
X
2000
1500
*
1000
500
X
W
? ?
X
W
*
X
W
*
*
W
W
W
*
**
*
*W *
*
* *
*
* *
*
0
20
40
60
80
100
2[ded.]
Fig. 3-1. X-ray diffraction patterns of alloys (c) B3; (d) B4
274
120
(a)
(b)
Wheel cooling side
Wheel cooling side
(c)
(d)
Wheel cooling side
(e)
Free surface
(f)
Free surface
(g)
(h)
Fig. 3-2. Microstructures of RS alloys (a) and (b) A1; (c) and (d) B2; (e) and (f)
B3; and (g) and (h) B4
275
Fig. 3-3 shows DTA results. Each of the A1, B2 and B3 shows
only one endothermic reaction, the melting of (Mg) matrix. B4 shows
two endothermic peaks, the lower temperature peak is the melting of
(Mg) + W + X eutectic mixture, and the higher temperature peak is the
melting of (Mg) matrix.
Fig. 3-3 DTA results of the RS ribbons (a) A1, (b) B2, (c) B3, and (d) B4
3.2 Solidification Process
It has been demonstrated that the RS process consists of two
steps [36-38]. First, rapid solidification takes place and coarse cells,
composed of fine super-saturate dendrite branches from highly
undercooled melt. This stage results in rapid latent heat release and
temperature recalescence. During this stage only about 20% of the total
melt solidified [38]. In the second stage, the residual melt undergoes
relatively slower solidification, and the fine dendrite branches reheat and
possibly remelt due to the solute redistribution and heat flux. The
remelting temperature is determined by the local solidification time [38],
solute redistribution and heat flux rate [38]. In the ribbons, the local
solidification time is limited due to their high ratio of surface to volume.
In the lower solute content A1, solute accumulation around dendrite
276
branch necks is not high enough to remelt the dendrite, therefore, coarse
grains are formed during first step of solidification. In higher solute
content B2, B3 and B4, on the other hand, the dendrite branches of the
coarse grains are partly remelted and separated into small dendritic
groups, these groups grow into the undercooled melt and eject solute to
the solid/liquid front to form the final fine grains, mainly away from the
wheel side (as shown in Fig. 3-2 (c) and (g)) or along the cross section
of the ribbons (as shown in Fig. 3-2 (e)).
From this point of view, the thickness of ribbon is also
important. The thicker it is, the slower is the solidification rate, and the
more serious is the remelt of the dendrite. As seen in Fig. 3-2 (e), where
the microstructure consists of small grains along the whole section of the
ribbon which is thicker than those in Fig. 3-2 (c) and (g).
The white submicron particles and network are isolated
compound rather than eutectic mixtures as shown in Fig. 3-4. According
to the X-ray analysis results and equilibrium phase diagram [39], the
compositions of the last solidification melts are located at (Mg) + W
two-phase region in B2 and B3, and at (Mg) + W +X three-phase region
in B4. However, during non-equilibrium solidification the last growing
(Mg) grows on the primary (Mg) directly, resulting in the growing of
isolated compounds at the last solidifying regions.
(a)
(b)
Fig. 3-4. The compound morphologies of rapidly solidified (a) B3 and, (b) B4
3.3 Potential Effect of Dispersiods and Grain Boundary
Network
It is well known that refined grain microstructure of Mg
alloys [40] can effectively improve their mechanical properties at room
temperature. Refined grains, obtained by rapid solidification, is therefore
a benefit. However, it is also known that elevated temperature strength
277
and creep resistance cannot be improved by refined grain structure
because of the rapid diffusion [41] through interior of the grains and along
grain boundaries. Therefore, the smaller the grain size, the greater the
total grain-boundary area available for grain boundary diffusion [42]
resulting in lower elevated temperature strength and creep resistance.
High melting point compound dispersed uniformly within the
magnesium solid solution and at grain boundaries is the best way to
increase elevated temperature strength and creep resistance [35]. The
increase in the volume fraction of particles within the magnesium
benefits mechanical properties by impeding dislocation mobility.
Moreover, these particles are thermally stable (as shown in Fig.3-3) due
to their higher melting point (> 800 K) [39], and limited solubility in
magnesium matrix [35].
High melting point compound network at grain boundary is
broken up during the extrusion process and have no bad effect on the
extruded material.
3.4 Y Contents
Even with 1.3%Y content the B2 ribbon can get reasonable
amount of Y containing compounds, as shown in Fig. 3-2 (c), due to two
reasons. First, in the Mg-Zn-Y system the solubility of Y in matrix is
limited as [42] at ambient temperature. During solidification, Y will be
ejected onto the solid/liquid interface and, then accumulate around the
necks of dendrite branches to form submicron compound and at cell
boundaries to form a particle network. Second, the resulted compounds
involve a large number of magnesium, or magnesium and zinc. On the
other hand, higher Y content will result in low melting point eutectic
mixture as indicated by the DTA results (Fig. 3-3 (d)) and X-ray
diffraction patterns (Fig. 3-1 (d)) of B4. Based on our experimental
results, we can conclude that when the alloy contains 5.5%Zn to 6%Zn,
Y contents of less than 4.3% can be beneficial.
Conclusion
When the magnesium alloy with 5.5%Zn to 6%Zn contains also
an addition of 1.3%Y-3.2%Y a reasonable amount of dispersed micro W
particles are formed and improve properties.
The microstructures of rapidly solidified Mg-Zn based alloys
consist of cellular structure. The cell wall contains higher concentration
278
of alloying elements than the center of the cell, as expected. The ribbon
solidifies under thermal diffusion controlled condition. For alloy B2, the
evaluated dendrite growth rate is excess 2  10-3 m/s.
Yttrium tends to segregate at grain boundaries and between
dendrite arms in Mg-Zn-Y system during conventional solidification.
Small amounts of Y additives to Mg-Zn base alloy is in solid solution if
rapidly solidified.
Using 1.3%Y-3.2%Y to alloy Mg-Zn base alloys can get
reasonable amount of dispersed micro W particles when the alloys
contain 5.5%Zn to 6%Zn.
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