Effect of grain refinements on the microstructure and thermal behaviour of Mg–Li–Al alloy
Journal of Thermal Analysis and Calorimetry
Effect of grain refinements on the microstructure and thermal behaviour of Mg-Li-Al alloy
Mariusz Kro´ l 0 1
0 Institute of Engineering Materials and Biomaterials, Faculty of Mechanical Engineering, Silesian University of Technology , Konarskiego 18a St., 44-100 Gliwice , Poland
1 & Mariusz Kro ́l
The light as-cast Mg-9Li-1.5Al alloys were manufactured and modified by 0.2 mass% Zr, commercial 0.2 mass% TiBor and 0.2 mass% AlSr master alloys. The thermal derivative analysis using UMSA platform was utilised to characterise a crystallisation process. Samples were heated up to 700 C and then freely cooled down (* 0.6 C s-1) to ambient temperature in order to simulate the natural cooling of casts. Dilatometry study was used to identify changes in solid state. The relative elongation of unmodified and modified alloys was measured in the temperature range from 20 to 400 C, with a heating rate 1.0 C s-1. The effects of Zr, TiBor and AlSr content on the microstructure of analysed magnesium alloys were investigated. Evaluation of microstructure was identified by light microscope, scanning electron microscopy, X-ray diffraction and energy-dispersive X-ray spectroscopy. The results showed that the addition of TiBor reduced the grain size of Mg-9Li-1.5Al cast alloy from 930 to 530 lm, while the addition of AlSr master alloy reduced the grain size to 480 lm. Moreover, an addition of TiBor and AlSr simultaneously reduced the grain size to 430 lm. The addition of Zr causes a reduction in grain size to 630 lm. The addition of grain refinement causes changes in crystallisation process and variations in the coefficient of linear thermal expansion (CTE).
Magnesium-lithium alloys; Thermal expansion; Thermal derivative analysis; Grain refinement
Following the World War II, Mg–Li-based alloys, as an
emerging light structural metallic material, have been
presumed to be extensively used in aerospace and civilian
areas, such as transport, electronics packaging and sports
manufacturing for their high specific strength and specific
stiffness, excellent damping and electromagnetic shielding
In the beginning, the analysis of Mg–Li-based alloys and
composites including applications, microstructures,
mechanical properties and processing technologies was
regularly presented. Furthermore, the development in the
research of intermetallic particles reinforced Mg–Li matrix
composites was studied [
Mg–Li-based alloys are the lightest among the
magnesium alloys and serve as the important structural materials
for aerospace industry. They maintain high cold
formability, and the problems of magnesium–lithium binary
alloys, such as poor resistance to corrosion and low
strength, limit their utilisation .
There are few standard ways to enhance the strength of
metallic materials through decreases in grain size and can
be divided into two groups like thermal and chemical
methods. The thermal method of grain refinement is
superheating, and chemicals are zirconium, carbon, silicon
carbide, manganese, titanium, calcium, strontium,
antimony, cerium additions and Elfinal process [
The additional methods capable of raising the strength
are an addition of alloying elements. The increase in a third
element—for example aluminium—to Mg–Li-based alloys
serves to change the strength relating to intermetallic
particle formation [
The mechanical properties of Mg–Li alloys increase
apparently with the addition of Al content, during the
elongation of alloys reduces dangerously when the Al
content is higher than 6 mass%. Titanium, strontium and
zirconium are also a useful alloying element to improve
Mg–Li properties [
According to the Mg–Li binary phase diagram, if the Li
content is less than around 5.5 mass%, the metal is formed
of a single a(Mg) phase with hcp structure. When the Li
content is within 5.5 and 11 mass%, the hcp a(Mg) phase
coexists with bcc-structured b(Li) phase, while the
structure has only a single bcc-structured b(Li) phase when the
Li content is higher than 11 mass% [
Nucleation of the primary phase is the initial step in the
transition of molten alloys into the solid state. Nucleation
control has been the subject of many studies [
]. In the
case of Mg alloys, a finer grain size increases most
mechanical properties, including yield strength due to the
high Hall–Petch coefficient, corrosion resistance and creep
In this study, the ternary alloys Mg–9Li–1.5Al (mass%)
were manufactured and modified by processed by
commercial Zr, TiBor and AlSr master alloys. The
microstructure evolution, thermal derivative analysis,
dilatation, scanning electron microscopy, X-ray diffraction
and energy-dispersive X-ray spectroscopy techniques were
used to examine the effect of grain refinements.
Within the framework of the present work, alloys of
magnesium with lithium and aluminium and with Zr, TiBor
and AlSr as grain refinement have been melted, casted and
examined. The raw materials used in these experiments
were magnesium with technical grade (min. 99.5% Mg),
aluminium 3N8 (99.98% Al), lithium (99.9% Li), TiBor,
AlSr and Zr as refinements were utilised. Melting and
casting of alloys were carried out using laboratory vacuum
induction furnace VSG 02 from the company Balzers.
Melts held in a crucible of Al2O3 in shape of ø 60 mm 9
80 mm, using the ceramic material sheath thermocouple
for measuring the temperature of melting and casting
alloys. Melting temperature was approx. 700/720 C and
the melting time approx. 5 min., which, taking into account
the strong bath stirring electrodynamic eddy currents in
enough for the complete homogenisation of the melt. Grain
refinement was introduced at the end of the melting process
from the vacuum containers. After placed of grain
refinement in the alloy, melts kept in the liquid state for 2 min,
followed by the casting. The chemical composition of
analysed Mg–Li–Al alloys and used grain refinements are
given in Tables 1 and 2.
The linear thermal expansion coefficient of the
investigated magnesium–lithium–aluminium alloys was measured
in argon atmosphere using the Bahr 805A/D dilatometer
over a temperature range from ambient temperature to
400 C at heating and cooling rates of 1 C s-1. The
thermocouple type S was utilised to measure changes in
temperatures. Cylindrical samples in shape 10 mm in
length and 4 mm in diameter were used in dilatometry test.
The thermal derivative analysis (TDA) was done on the
prepared cylindrical samples in shape of 18 mm in
diameter and 20 mm using UMSA device [
10, 11, 14
were melted at 700 C in an argon atmosphere. Following
isothermal holding for 90 s, all the melts were solidified
and cooled to ambient temperature in the crucibles with
argon protection in the furnace to minimise the oxidation.
The signal from the thermal derivative analysis was
acquired using a high-speed National Instruments data
acquisition system. The recorded data were analysed in
Fityk software. The cooling curves and corresponding
derivative curves were plotted to define thermal events,
based on the first and second derivative of cooling curve.
The test samples were solidified at an average cooling rate
of approximately 0.6 C s-1 in the range of liquidus and
solidus temperature. The cooling rate was determined using
the following formula:
CR ¼ tsol
where Tliq and Tsol are the liquidus and solidus
temperatures, respectively, and tliq and tsol the times from the
cooling curve that correspond to liquidus and solidus
The Newtonian method was applied in this research to
determine the baseline. This means that the thermal
gradient across the specimen is considered to be zero and that
heat transfer between the casting and the mould occurs
through convection. The Fourier method is another method
used to calculate the baseline in which the positions of two
thermocouples at the beginning and end of the
solidification process are considered. However, it is much more
complicated than the Newtonian method. Besides, a
thermal contraction of the metal occurs during solidification,
the exact positions of the two-thermocouple tips are
difficult to estimate [
]. The base line has been predicted
by the sixth polynomial fitting (dT/dt)BL =
a0?a1T?a2t)BL = a0?a1T?a2T2?a3T3?a4T4?a5T5?a6T6 between the
beginning and the end of solidification in the first
derivative curve. The sixth-order polynomial yields a correlation
coefficient greater than 0.97.
Metallographic specimens were horizontally sliced at
the position that is 10 mm from the top. The as-cast grains
of the etched samples were measured using polarised light
in optical microscope Leica equipped Q-WinTM image
analyser. The grain size was determined by the linear
intercept method at the centre of transverse sections.
The X-ray qualitative and quantitative microanalysis
and the analysis of a surface concentration of cast elements
in the examined magnesium–lithium–aluminium
unmodified and modified alloys have been made on the scanning
electron microscope ZEISS SUPRA 35 with a system
EDAX XM4 TRIDENT consisting of spectrometer EDS,
WDS and EBSD (20 kV, 15 mm of work distance and
30 lm of aperture).
Phase composition was determined by the X-ray
diffraction technique utilising the X’Pert apparatus
including a cobalt lamp with 40 kV voltage. The
calculation was done by angle range of 2H: 30 –110 . The
The analysis of coefficient of linear thermal expansion of
Mg–9Li–1.5Al (Table 3 and Fig. 1) has validated the fact
that CTE depends on temperature. Untreated Mg–9Li–
1.5Al alloy has CTE about 29.4 10-6 K-1, and with
further increase in temperature to 250 C, CTE increases to
34.3 10-6 K-1. Further increase in temperature to 400 C
caused a decrease in CTE to 31.9 10-6 K-1. Analysis of
the heating and cooling dilatometric curves of analysed
materials modified by TiBor and AlSr is characterised by a
linear decrease in linear expansion coefficient as a function
of temperature. Moreover, based on changes in elongation
during heating and cooling cycle, it was found that shape of
graphs has a linear relationship that means no transitions in
solid state occur.
The phase composition of the investigated alloys treated
by grain refinements and untreated was examined by XRD
technique. The XRD profile of unmodified alloy
demonstrates that there are two phases—b(Li) phase (matrix
phase, a solid solution of magnesium in body-centred cubic
(bcc) lithium lattice) and a(Mg) phase (solid solution of
lithium in hexagonal close-packed (hcp) magnesium
lattice), which is faithful with the observation using a light
microscope. No peaks of Al including phase are noticed
that suggests most of Al exists in a solid solution state.
However, with the addition of Zr, new peaks correspond to
the Zr phases. Furthermore, with the addition of TiBor and
AlSr, no additional peaks were observed.
Optical micrographs of Mg–9Li–1.5Al cast alloy and
after modification by Zr, TiBor and AlSr are given in
Fig. 2. It can be observed that the structure of the analysed
alloys is fundamentally formed of light grey a(Mg) phase
and dark grey b-Li phase, which is in consistent with the
Mg–Li binary system [
]. Most of the a(Mg) phase is
elongated ribbon-like, where b(Li) phase fills the area
between a(Mg) phase grains, separated by boundaries
between two phases. In the duplex phase matrix, the ratio
of a and b phases is 40/60. There also exists some
particlelike (Fig. 2e—marked in white circle) phase in the alloy.
The granular-dispersive phase of the intermetallic
compounds is randomly spread in the a and b matrices. Darker
particles have most probably resulted from the aluminium
addition in the form of g(LiAl) phase as suggested in
8, 9, 19, 20
]. SEM analysis of marked circles shows high
concentration of Al, indicating that it is g(LiAl) phase.
Moreover, the addition of grain refinements effects in the
conversion of the elongated-striped a(Mg) phase with a
blocky structure and rounded corners. Moreover, optical
metallography indicated fully dense alloys.
Analysis of grain size of analysed alloys after thermal
derivative analysis cooled at a rate of 0.6 C s-1 appointed
by the method of linear interception length shows that the
grain refining appearance is different for each alloy (Figs. 3
and 4). Once Zr is added, there is a significant reduction in
grain size. The Mg–9Li–1.5Al alloy was refined from 930
to 640 lm by 0.2 mass% Zr. The results showed that the
addition of TiBor master alloy decreased the grain size of
Mg–9Li–1.5Al cast alloy to 500 lm, while the addition of
AlSr master alloy reduced the grain size to 480 lm.
Moreover, an addition of TiBor and AlSr simultaneously
reduced the grain size to 430 lm. The relatively small
grain size of Mg–9Li–1.5Al?0.2TiBor?0.2Sr is not
obviously reflected in its hardness related to Mg–9Li–
1.5Al?0.2TiBor. It can be noted that addition of AlSr
refiner slightly increases in hardness. In general, the
hardness did not increase with TiBor and Zr addition
Furthermore, investigated alloys were characterised by
SEM equipped with EDS and computer-controlled imaging
system. The SEM images of analysed Mg–Li–Al alloys
point out dark grey and grey areas relating to the b(Li)
phase and a(Mg) phase, respectively, as shown in Fig. 6a.
Because lithium is invisible in SEM image, b-phase is
darker. Grain refinement of analysed magnesium–lithium–
aluminium alloy by 0.2 mass% Zr causes appearance to
Fig. 3 Microstructure of the grain size for Mg–9Li–1.5Al base alloy at various addition of grain refinements: a Mg–9Li–1.5Al, b Mg–9Li–
1.5Al?0.2Zr, c Mg–9Li–1.5Al?0.2TiBor, d Mg–9Li–1.5Al?0.2Sr, e Mg–9Li–1.5Al?0.2TiBor?0.2Sr, Nomarski contrast
alloys, respectively. Containing Sr, an element with a
larger atomic number, bone-like bright intermetallic
compounds (indicated as 3) are situated along the phase
boundary as well as across a-phase. EDS point analysis,
which was done at location 3, demonstrates the high
concentration of Mg, Sr and Mg, Al, Sr, respectively.
Considering Al dissolved in the matrix, it can be concluded that
these compounds can be related to Al4Sr or Mg2Sr as
suggested in [
]. Shape and location of recognised phases
in analysed alloys could be directly related to the order of
nucleation in the intermetallic compounds.
The solidification pathways of as-cast (Fig. 7a) and after
modification by 0.2 mass% Zr (Fig. 7b) and 0.2 mass%
TiBor (Fig. 7c) and 0.2 mass% AlSr (Fig. 7d) master
alloys have been investigated by thermal derivative
analysis. Figure 7a shows typical cooling curves and its first
derivative curves which were used to determine
characteristic points during solidification. For the alloy with 0.2
mass% Zr, only five well-defined points are observed, i.e.
at 597.7 C (1-TN-nucleation temperature), 578.9 C
(2TG-maximum temperature of growth), 526.1 C
(4-TSOLsolidus temperature), 503.3 C (5-TS-a?b solvus
temperature) and 473.3 C (6-TER-end of reactions) as shown in
Fig. 7b. No more exothermic peaks were found after
modification by Zr that can be corresponded to compounds
of Al3Zr as suggested in [
]. It must be noted that the
melting temperature of Al3Zr intermetallic compound
according to the literature  is 1580 C. Results from the
thermal derivative analysis (Table 4) present that the
addition of 0.2 mass% Zr does not change the nucleation
temperature, i.e. 596.5 C and, however, decreases the
solidus temperature from 549.1 C of investigated
magnesium alloys. For both analysed alloys, no significant
changes in temperature in other reactions like TG, TS and
TER were observed. The solidification time for the cooling
rate 0.6 C s-1, as the difference of the times, at which the
liquidus and solidus temperatures occur in Mg–9Li–1.5Al
and Mg–9Li–1.5Al?0.2Zr, amounts to Dt = 106 and
134 s, respectively.
Addition of TiBor and AlSr causes to appearance new
well-defined exothermic peak (marked as point 3 in
Fig. 7c–e) probably becomes from nucleation of
metastable intermetallic phase g(LiAl) with B2 structure
8, 9, 20
] 3-Tg(LiAl) observed at 552.6 C (Mg–9Li–
1.5Al?TiBor), 559.1 C (Mg–9Li–1.5Al?AlSr) and
546.8 C (Mg–9Li–1.5Al?0.2TiBor?0.2Sr), however,
more studies must be done.
Modification of Mg–9Li–1.5Al alloy by AlSr master
alloy causes slight decreases in nucleation temperature to
582.7 C, but strongly reduces a solidus temperature to
535.3 C in accordance with the unmodified alloy. The
addition of TiBor and AlSr10 master alloys causes
white granular particles of the intermetallic compounds
randomly spread in the a and b matrix (Fig. 6b). EDS point
analysis carried out on particles (indicated as 3) to
determine the chemical composition shows that the mixture
consists of high concentration of Zr, Al and Mg elements,
which the most probably resulted in the form of Al3Zr
phases as suggested in [
]. With the addition of TiBor
master alloy as grain refinement to the Mg–9Li–1.5Al
alloy, no other phases are observed in microstructure
Figure 6d, e shows the SEM images and EDS analysis
of the Mg–9Li–1.5Al?Sr and Mg–9Li–1.5Al?TiBor?Sr
Fig. 6 SEM micrographs with EDS analysis in labelled points of Mg–Li–Al alloys: a Mg–9Li–1.5Al, b Mg–9Li–1.5Al?0.2Zr, c Mg–9Li–
1.5Al?0.2TiBor, d Mg–9Li–1.5Al?0.2Sr, e Mg–9Li–1.5Al?0.2TiBor?0.2Sr
decreases in nucleation temperature and solidus
temperature to 579.32 and 520 C, respectively.
The addition of TiBor and AlSr master alloys increases a
crystallisation time to 132 s for Mg–9Li–1.5Al?TiBor, 146
s for Mg–9Li–1.5Al?TiBor and 143 s for Mg–9Li–
1.5Al?0.2TiBor?0.2Sr. Moreover, an addition of Sr
causes a reduction in a period between the solidus temperature
and solvus temperature. Based on thermal derivative
analysis, it was found that only addition of TiBor and AlSr
simultaneously causes a decrease in temperature of
maximum growth of primary b(Li) phase from 581.3 to
In general, no exothermic peaks, corresponding to
precipitation of Mg2Sr, with melting temperature 680 C
], were observed during analysis of crystallisation
process of Mg–Li–Al alloys containing Sr as grain
refinement. This may be the causes that the amount of Mg2Sr
intermetallic compound is very low, below the level of
detection of the method.
The crystallisation pathway of investigated unmodified
and modified magnesium alloys, based on thermal
derivative analysis, microstructure investigation, SEM
observation and study of binary and ternary systems [
], may be
(solvus line in Mg–Li binary system) causing forming of
a(Mg) and b(Li) according to reaction b(Li)?
Thermal derivative analysis can be implemented to
registration melting and cooling processes of Mg–Li–
The effects of addition of 0.2 mass% Zr, 0.2 mass% TiBor
and 0.2 mass% AlSr grain refiners on the crystallisation
process during solidification of the Mg–9Li–1.5Al alloy
such as TN, TG, TSOL, TS and TER, linear coefficient of
thermal expansion, microstructure and hardness were
studied. The results are summarised as follows:
At a 0.6 C s-1 cooling rate, the solidification time of
Mg–9Li–1.5Al alloy is increased by adding grain
In analysed material, an addition of Zr does not change
TN, TG, TS and TER; however, decreases TSOL. The
addition of TiBor and AlSr causes a reduction in
nucleation and solidus temperature.
Analysis of the heating and cooling dilatometric curves
of analysed materials modified by TiBor and AlSr is
characterised by a linear reduction in coefficient of
linear thermal expansion as a function of temperature.
No transitions in the solid state occur during heating
and cooling cycles.
The as-cast Mg–9Li–1.5Al alloy comprises a(Mg) and
b(Li) phases. AlSr addition results in the formation of
new intermetallic compound distributed along the
phase boundary as well as across a-phase. EDX
analysis shows that the precipitated compounds are
Al4Sr or Mg2Sr.
The addition of Zr, TiBor and AlSr in level 0.2 mass%
reduces the grain size of analysed magnesium alloy.
The strongest effect in reduction in grain size (grain
size decreased almost twice) was achieved when TiBor
and AlSr simultaneously were used.
Based on hardness measurements, it was found that the
addition of AlSr refiner slightly increases hardness.
The hardness did not increase with TiBor addition.
Acknowledgements This work was financed by the Ministry of
Science and Higher Education of Poland as the statutory financial grant
of the Faculty of Mechanical Engineering SUT.
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