Bhale2016

Effect of Deep Cryogenic Treatment on Microstructure
and Properties of AE42 Mg Alloy
Pranav Bhale, H. Shastri, A.K. Mondal, M. Masanta, and S. Kumar
(Submitted December 29, 2015; in revised form June 29, 2016)
The effect of deep cryogenic treatment (DCT) on microstructure and mechanical properties including
corrosion behavior of the squeeze-cast AE42 alloy has been investigated. For comparison, the same has also
been studied on the untreated alloy. Both the untreated and deep cryogenic-treated (DCTed) alloys com-
prised a-Mg and Al4RE phases. Volume fraction of the Al4RE phase in the AE42 alloy reduced gradually
following DCT carried out from 4 to 16 h. Ductility and UTS increase significantly with a marginal increase
in YS of all the DCTed alloys. The improvement was attributed to the dissolution of the brittle Al4RE phase
following DCT. Among the alloys employed, the best tensile properties were obtained for the 16-h DCT alloy
due to its lowest content of the brittle Al4RE phase. Creep resistance of the DCTed alloys was lower than
that of the untreated alloy owing to the presence of less amount of thermally stable intermetallic Al4RE
phase. Wear resistance of the alloy reduces following DCT due to reduced hardness of the DCTed alloys.
The untreated alloy exhibits the best corrosion resistance, whereas poor corrosion resistance of the DCTed
alloys is attributed to the reduced amount of Al4RE phase that fails to built a corrosion resistance barrier.
Keywords AE42 magnesium alloy, creep, corrosion, deep cryo-
genic treatment, tensile, wear
1. Introduction
Magnesium (Mg) alloys are suitable for applications in
automobile and aerospace industries owing to their superior
specific strength compared to other alloys. These alloys also
possess good damping capacity, high thermal conductivity and
good machinability (Ref 1). However, relatively poor ambient
and elevated temperature mechanical properties as well as very
inferior corrosion resistance of Mg alloys restricted their
widespread applications. In order to overcome this problem,
several attempts were made in the recent past. One such
endeavor of improving properties of Mg alloy is to carry out
temperature-controlled treatment where the alloy is heat-treated
or cold-treated. Heat treatment is a common practice followed
from earlier ages, whereas cold treatment is a much newer
method followed to investigate the changes in properties of Mg
alloy. The conventional cold treatment (CT) is the process
where the material is subjected to temperature near to À80 °C.
There is another method known as ÔDeep Cryogenic Treatment
(DCT)Õ which is carried out using liquid nitrogen (tempera-
ture % À200 °C). DCT improves certain properties of material
beyond the improvement achieved by conventional CT (Ref 2).
In general, the basic step involved in CT/DCT, cooling the
specimen gradually to the specified temperature and holding it
for a desired period, and then bringing it back to room
temperature (Ref 3). DCT was carried out extensively on
various types of steels, and significant improvement in
mechanical properties was reported. Bensely et al. (Ref 4)
carried out DCT, conventional heat treatment (CHT) and
shallow cryogenic treatment (SCT) on case carburized EN 353
steel and compared the results. They observed the best wear
resistance in the DCT specimen. Liu et al. (Ref 5) carried out
DCT on high-chromium cast iron. The treated cast iron
exhibited superior bulk hardness and wear resistance as
compared to that of the untreated specimen. The improvement
was attributed to conversion of retained austenite to martensite
as well as precipitation and fine distribution of carbides.
Zhirafar et al. (Ref 6) carried out DCT on 4340 steel and
reported an improved hardness and fatigue strength with
reduced toughness. They concluded these changes were due to
reduction in volume % of retained austenite that transformed to
martensite after DCT. Firouzdor et al. (Ref 7) carried out DCT
on M2 HSS drills and reported an increase in tool life by 77%
with improved wear resistance. DCT in their study facilitated in
carbide formation and also increased the homogeneous carbide
population in martensite matrix. DCT on tool steel carried out
by Rhyim et al. (Ref 8) increased tool life by decreasing
retained austenite as well as by forming nanoscale fine g-
carbide. Thus, the superior properties exhibited by steels
following DCT was attributed to nearly complete transforma-
tion of retained austenite into martensite as well as the
precipitation of submicroscopic carbides. There have been
few studies on the DCT of Mg alloys. Asl et al. (Ref 2) carried
out DCT on AZ91 alloy and observed the changes in the
distribution of b-Mg17Al12 precipitates. They reported an
improved mechanical properties including wear resistance
following DCT. An improvement in hardness and wear
resistance after DCT on AZ91 alloy was reported by Amini
et al. (Ref 9). The improvement was attributed to the changes in
crystal structure after DCT. Jiang et al. (Ref 10) reported the
effect of DCT on microstructure and mechanical properties of
the AZ31 alloy. In their study, tensile strength and hardness of
Pranav Bhale, H. Shastri, and A.K. Mondal, Department of
Metallurgical and Materials Engineering, National Institute of
Technology, Rourkela 769008, India; M. Masanta, Department of
Mechanical Engineering, National Institute of Technology, Rourkela
769008, India; and S. Kumar, Department of Materials Engineering,
Indian Institute of Science, Bangalore 560012, India. Contact e-mails:
mondala@nitrkl.ac.in and ashok.mondal@gmail.com.
JMEPEG ÓASM International
DOI: 10.1007/s11665-016-2238-6 1059-9495/$19.00
Journal of Materials Engineering and Performance

the alloy were improved following DCT. This improvement
was due to the generation of ÔFrame-likeÕ twinning during DCT,
which changed the grain orientation.
A summary of the above literature revealed that DCT on Mg
alloy is beneficial for improving mechanical properties.
Although DCT on steels was carried out extensively, the
number of studies on Mg alloys is limited. Therefore, in the
present investigation, an attempt has been made to improve
mechanical properties and corrosion behavior of the AE42
alloy developed for automobile powertrain and aerospace
applications.
2. Experimental Procedure
The AE42 Mg alloy was used in the present investigation
and its chemical composition is shown in Table 1. The alloy
was fabricated by direct squeeze-casting. In this process, the
AE42 alloy melt (at 770 °C) was poured manually into a die
preheated to 250 °C. An upper ram of a vertical hydraulic press
squeezed it with a speed of 10 mm/s. The liquid solidified
under a pressure of 100 MPa in about 20 s. The squeeze-cast
block was allowed to cool to ambient temperature. Specimens
having dimension of 15 mm diameter and 10 mm thickness for
corrosion tests, 10 mm diameter and 5 mm thickness for wear
tests and 10 mm * 10 mm * 10 mm for impression creep tests,
and microstructural analysis were machined from the cast block
by wire electro-discharge machining. Deep cryogenic treatment
(DCT) of all the fabricated specimens was carried out by
soaking the specimens into liquid nitrogen (À200 °C) inside a
vacuum flask for 4, 8, and 16 h and then allowed to cool to
ambient temperature (%25 °C) after removal from the flask.
Microstructural observation of the specimens was carried out
using optical microscopy (OM) (Model: ZEISS AxioVision
I10) and field emission scanning electron microscopy (FESEM)
(Model: FEI Sirion XL30) furnished with energy-dispersive x-
ray spectroscopy (EDS). The specimens for microstructural
examination were prepared by standard metallographic tech-
nique. For etching, a solution containing 10 mL acetic acid,
100 mL ethanol, 6 ml picric acid and 20 ml of distilled water
was used. The presence of phases in the specimens was
confirmed by x-ray diffraction (XRD) (Model: PAN analytical,
DY-1656) using CuKa (k = 1.5418 A˚ ) technique. The surfaces
of all the specimens were slightly grinded with silicon carbide
paper of different grades (up to 2500 grit) using deionized
water as coolant. The specimens were then cleaned in alcohol
and dried. Electrochemical corrosion tests were carried out at
ambient temperature (%25 °C) in aqueous 0.5 wt.% NaCl
solution, at neutral pH. A three-electrode corrosion cell
(333 mL) including an Ag/AgCl (3 mol/L KCl) reference
electrode, a counter electrode of Pt and the specimen as
working electrode was employed. Potentiodynamic polarization
scan was performed from À200 mV relative to the free
corrosion potential using a scan rate of 0.2 mV/s and the test
lasted for about 30 min. Corrosion rate was calculated from the
obtained Tafel curve. The current density was determined at the
point of intersection of the Tafel slope obtained from the
cathodic branch of polarization curve with the vertical line
passing through the corrosion potential. Corrosion rate can be
determined from the cathodic branch of the curve knowing
exposed specimen surface area and Tafel equation (Ref 11).
The specimens were cleaned ultrasonically after corrosion tests
and observed under microscope with and without removal of
the corrosion products. Corrosion products from the tested
specimen surfaces were removed using a solution of 20%
chromic acid and about 1% silver nitrate following ASTM G1
standard. The specimens were immersed in the solution kept in
a glass, and the container was heated to 45 °C for 1 min in
normal atmosphere. Finally, the specimens were rinsed in
deionized water and dried with ethanol. Hardness of all the
specimens was measured using Vickers microhardness tester
(Model: LecoTM
248AT, USA) using 100 g load applied for
10 s. The tensile tests were carried out as per ASTM E8M-03
standard using a universal testing machine (Model: INSTRON
5967) at ambient temperature. At each condition, two speci-
mens were tested and the average value was reported as tensile
properties. Fracture surfaces of the broken tensile specimens
were examined under a FESEM. The creep tests were carried
out at a temperature of 200 °C and stress of 210 MPa using
impression creep technique. Dry sliding wear tests were carried
out in accordance with the ASTM G-99 standard employing a
ball-on-plate wear testing machine (TR-208 M1). The machine
consists of an EN32 steel disk of diameter 100 mm (hardness
RC 58). Wear tests were carried out at a disk speed of 20 rpm
using a load of 20 N for 10 min. During wear test, wear depth
of the specimens was recorded as a function of sliding distance.
Both the disk and the specimens were cleaned thoroughly with
acetone before and after each test to avoid contamination. Each
test was repeated at least twice to check the reproducibility, and
average value is reported here as wear depth. Micrographs of
Table 1 Chemical composition of the AE42 alloy
Element Al Zn Mn Si Ce La Nd Pr Th Be Mg
wt.% 3.9 0.01 0.3 0.01 1.2 0.6 0.4 0.1 0.6 0.001 Balance
Fig. 1 XRD patterns obtained from the untreated and DCTed AE42
alloys

the worn-out surfaces after wear tests were examined under a
FESEM.
3. Results and Discussion
3.1 Microstructural Characterization
The XRD patterns obtained from the untreated and deep
cryogenic-treated (DCTed) AE42 alloys are shown in Fig. 1. It
is evident from the figure that all the alloys consist of primary
Mg (a-Mg) peaks along with the peaks confirming to Al4RE
phase. The Al11RE3 and Al4RE phases are considered to be
same and are usually represented as Al4RE phase in the
literature (Ref 12). No peak was observed for the b-Mg17Al12
phase which is generally present in Mg-Al alloys. Thus, the b-
Mg17Al12 phase (M. P. 437 °C) is completely suppressed in the
RE-containing AE42 alloy, which is favorable for its better
creep resistance. In the present investigation, the RE was Ce-
rich misch metal. The electronegativity values of Mg, Al, Ce
and La are 1.31, 1.61, 1.12 and 1.10, respectively. The
difference in electronegativity between Al and Ce/La is more
than that between Mg and Al. Therefore, Al had more tendency
to form Al4RE phase rather than b-Mg17Al12 phase. Similarly,
Mg-RE phase formation was also not favored. Powell et al.
(Ref 12) and Mondal et al. (Ref 13, 14) too reported the
presence of the same phases in the AE42 alloy. The peaks
corresponding to Al4RE phase were present in all the alloys;
however, their intensity in all the DCT alloys was weaker as
compared to that of the as-cast AE42 alloy. Further, the
intensity gradually decreased with the increase in DCT time,
which indicated a reduction in the amount of Al4RE phase in
the DCT alloys. The reduction in the Al4RE phase was the
highest in the alloy subjected to DCT for 16 h. A similar
observation, i.e., reduction in the eutectic b-Mg17Al12 phase in
AZ91 alloy following DCT, was reported by Asl et al. as well
(Ref 2). Besides a-Mg and Al4RE, there was no evidence of
formation of new phase following DCT.
Figure 2(a)-(d) displays optical micrographs of all the
squeeze-cast AE42 alloys with and without DCT. All the
alloys exhibited dendritic microstructures with the second
phase distributed at the interdendritic regions. The dendritic
morphology in all the microstructures remains almost same.
Figure 3(a)-(d) shows SEM micrographs of both the untreated
and DCTed AE42 alloys. The micrograph of the AE42 alloy
(Fig. 3a) comprised primary Mg (a-Mg) matrix surrounded by
the lamellar eutectic consisting of alternate layers of a-Mg and
Al4RE, as observed at higher magnification. The untreated
alloy exhibited almost polygonal grains of a-Mg with average
grain size of 25 lm. The Al4RE phase comprising bright
contrast was observed along the grain boundaries and triple
points. The detailed microstructure of the as-cast AE42 alloy is
available elsewhere (Ref 12, 13). There was no evidence for the
formation of b-Mg17Al12 phase in all the alloys. Following
DCT, the microstructures of the AE42 alloy undergo visible
changes. The obvious change was the reduced dimension of the
eutectic formed along the grain boundaries, i.e., refinement of
Al4RE phase. It means DCT assisted refining the eutectic phase
at the grain boundaries. It was observed that the eutectic Al4RE
phase gradually dissolved more into the a-Mg matrix with the
Fig. 2 Optical microstructures of the AE42 alloy (a) untreated and DCTed for (b) 4 h, (c) 8 h and (d) 16 h

increase in time of DCT. Accordingly, the volume fraction of a-
Mg increased. The volume fraction of the a-Mg and Al4RE
phases was calculated from SEM micrographs of all the
specimens and is shown in Fig. 4. It is evident that volume
fraction of the Al4RE phase reduced in all the DCT alloys as
compared to the untreated AE42 alloy. Among the DCT alloys,
the highest and the lowest volume fractions of the Al4RE phase
were exhibited by the DCT alloys soaked for 4 and 16 h,
respectively, while the DCT alloy soaked for 8 h exhibited the
intermediate volume fraction of the Al4RE phase. In addition,
the eutectic structure broke down into smaller structures. A
similar observation, i.e., refinement of the precipitates follow-
ing DCT, was also reported by Asl et al. (Ref 2). About 3.0-
5.0 wt.% Al was dissolved interior of grains of the DCT alloys,
whereas it was only about 2.1 wt.% interior of the grains of the
untreated alloy, as confirmed by EDS analysis. Therefore,
increased solubility of Al in a-Mg of the DCT alloys was
attributed to dissolution of the eutectic Al4RE phase.
3.2 Tensile Behavior
Figure 5(a) shows the tensile plots of the squeeze-cast
untreated as well as DCTed AE42 alloys tested at ambient
temperature. The various tensile properties calculated from
Fig. 5(a) are shown in Fig. 5(b). It is observed that the DCT on
AE42 alloy influenced the tensile properties significantly. The
ductility (% elongation) and ultimate tensile strength (UTS) of
all the DCT alloys increased significantly with a marginal
increase in yield strength (YS) (0.2% proof stress) as compared
to that of the untreated AE42 alloy. The values of YS, UTS and
ductility increased gradually with the increase in DCT time.
The increase in ductility following DCT carried out for 4, 8 and
16 h in the alloys was 23.1, 51.1 and 101.8%, respectively,
compared to the untreated AE42 alloy. The best tensile
properties were exhibited by the alloy subjected to DCT for
16 h. Thus, the DCT has the potential to overcome the most
Fig. 3 SEM micrographs of the AE42 alloy (a) untreated and DCTed for (b) 4 h (c) 8 h and (d) 16 h
Fig. 4 Volume fraction of the a-Mg and Al4RE phases in the un-
treated and DCTed AE42 alloy

signiﬁcant problem of lower ductility of Mg alloy. However,
the increase in YS with the increase in DCT time was not very
signiﬁcant. Liu et al. (Ref 15) reported that the AZ91D alloy
exhibited the best combination of tensile properties following
DCT carried out for 24 h.
The improved tensile properties of the DCTed AE42 alloys
could be related to the partial dissolution of the Al4RE phase
that was present along the grain boundaries of the alloys before
carrying out DCT. The amount of eutectic Al4RE phase was the
highest in the untreated AE42 alloy among the alloys employed
in the present investigation. The Al4RE phase is brittle in nature
and prone to fracture owing to higher brittleness. It might cut
apart the matrix reducing the ductility and UTS of the untreated
alloy. However, the volume fraction of the Al4RE phase was
relatively less in all the DCTed alloys as compared to that of the
untreated one, and accordingly, the alloys were less brittle.
With the increase in DCT time from 0 to 16 h, the solubility of
Al and RE in a-Mg increased. Al was added to Mg to improve
strength by solid solution strengthening (Ref 16-18). Further,
the presence of RE elements was more effective than that of Al
in improving strength of Mg by solid solution strengthening
(Ref 19). The increase in strength by solid solution strength-
ening resulted from the presence of both Al and RE in a-Mg
could be the plausible cause for the improved YS and UTS of
Fig. 5 (a) Tensile plots for the AE42 alloys tested at ambient tem-
perature and (b) variation of tensile properties with DCT time
Fig. 6 Fracture surfaces of the broken tensile specimens of the AE42 alloy corresponding to (a) untreated and (b) DCTed for 16 h
0 2000 4000 6000 8000 10000
0.00
0.04
0.08
0.12
0.16
0.20
0 h
Impressiondepth(mm)
Time (s)
4 h
Fig. 7 Impression creep plots of the AE42 alloy corresponding to
untreated and DCTed for 4 h

the DCT alloys. The superior tensile properties of the DCT
alloy subjected to 16 h were attributed to the presence of the
lowest amount of the brittle Al4RE intermetallic phase and the
maximum effect of solid solution strengthening.
The micrograph of the fracture surface of the broken tensile
specimen of the untreated AE42 alloy is shown in Fig. 6. The
fracture surface exhibited a distinct quasi-cleavage type of
failure with several cleavage planes (shown by A) and few
cleavage steps (B), which represent brittle transgranular-type
cleavage fracture. River patterns (C) and cleavage tongues (D)
Fig. 8 Variation of microhardness values for all the AE42 alloys
Fig. 9 Variation of wear depth as a function of sliding distance for
the untreated and DCTed AE42 alloys
Fig. 11 Results of potentiodynamic polarization tests carried out on
the untreated and DCTed AE42 alloys
Fig. 10 SEM micrographs of worn surfaces of the AE42 alloy corresponding to (a) untreated and (b) DCTed for 16 h
Table 2 Summary of various parameters calculated from
Tafel plots
DCT
time, h
Corrosion
potential,
Ecorr, mV
Current
density,
Icorr, mA/cm2
Corrosion
rate, mm/year
0 1497.9 0.029 0.68
4 1523.1 0.034 0.77
8 1506.8 0.040 0.91
16 1461.4 0.042 0.95

on the flat facets of the cleavage planes revealed a series of
plateaus and ledges. The cleavage-type fracture started from a
number of cleavage planes leading to cleavage steps. The
network of these steps called river pattern completes the
fracture. With the increase in DCT time, more number of quasi-
cleavage planes (E) was observed as shown in Fig. 6(b). The
presence of microvoids (F) and plastic deformation zone (G)
characterizing ductile failure was relatively less on the fracture
surface of the untreated alloy; however, their population
increased in the DCT alloys with the increase in DCT time.
Many shrinkage voids (H) were observed in the treated alloy, as
shown in Fig. 6(b). The lamellae of the Al4RE (I) phase
(confirmed by EDS) were observed on the fracture surfaces of
the alloys; however, its number reduced with the increase in
DCT time.
3.3 Creep Behavior
The typical impression creep curves (impression depth versus
time) of the squeeze-cast untreated as well as the 4-h DCTed (as a
representative plot for the DCTed alloys) AE42 alloy tested
at 210 MPa stress and temperature of 200 °C are shown in
Fig. 7. Evidently, the DCT of the alloy resulted in varied creep
characteristics. Both the curves exhibited a distinct primary
(transient) creep regime followed by a steady-state (secondary)
creep. The creep tests were conducted using impression creep
technique where the specimen was subjected to compression,
and therefore, the tertiary stage is not observed in the plots. The
impression velocity was calculated from the steady-state region
of the impression depth versus time curves for both the
specimens, i.e., from the second stage in each case. The
untreated and the 4-h DCTed alloys exhibited impression
velocities of 5.4 9 10À6
and 8.6 9 10À6
mmsÀ1
, respectively.
Thus, the steady-state impression velocity has increased by a
factor of 1.6 following DCT at the stress level and temperature
employed in the present investigation. The presence of the
thermally stable intermetallic Al4RE phase (M. P. 1049 °C)
along the grain boundaries and triple points prevent grain
boundary sliding at elevated temperature, and thus, the creep
resistance of the AE42 alloy improves. The presence of
relatively lower volume fraction of the Al4RE in the DCTed
alloys resulted in lower creep resistance of the DCTed alloys as
compared to that of the untreated AE42 alloy. Thus, the DCTwas
not beneficial for improving creep resistance of the AE42 alloy.
3.4 Hardness
The microhardness values of the alloys are shown in Fig. 8.
It is evident from the figure that the untreated alloy exhibited
the highest hardness among the alloys employed in the present
investigation. The hardness values of the DCT alloys decreased
with the increase in DCT time from 4 to 16 h. Thus, the DCT
on AE42 alloy significantly influenced hardness values of the
alloys. The eutectic Al4RE phase is harder as compared to a-
Mg. The amount of eutectic Al4RE phase is the highest in the
untreated AE42 alloy leading to the highest hardness. The
Fig. 12 XRD patterns obtained from corroded specimens of the un-
treated and 16-h DCTed AE42 alloys
Fig. 13 SEM micrographs of corroded surfaces of the AE42 alloy corresponding to (a) untreated and (b) 16-h DCTed

inferior hardness values of the DCT alloys were attributed to
the presence of lower volume fraction of the hard Al4RE phase.
3.5 Wear Behavior
The variation of wear depth (in lm) as a function of sliding
distance (in mm) is shown in Fig. 9. The untreated alloy
exhibited the lowest wear depth, i.e., the maximum wear
resistance among the alloys employed. The wear depth for the
treated alloys gradually increased with the increase in DCT
time, while the alloy subjected to DCT for 16 h exhibited the
highest wear depth, i.e., lowest wear resistance.
SEM micrographs of worn surfaces of the untreated and 16-
h DCT alloy, as a representative micrograph for the DCT alloys,
are shown in Fig. 10(a) and (b). From the micrographs, it is
evident that several continuous parallel grooves were formed
on the worn surfaces of both the specimens. However, the
grooves observed on the worn surface of the 16-h DCT alloy
were deeper, wider and more in number. The parallel grooves
present on the worn surfaces were due to microploughing,
which confirms the abrasion mechanism of wear. In addition,
the worn surface of 16-h DCT alloy consisted of irregular-
shaped debris that was scattered from the wear track. The lower
hardness of the DCT alloys compared to the untreated alloy
resulted in their lower wear resistance.
3.6 Corrosion Behavior
Figure 11 shows the results of potentiodynamic polarization
tests carried out on all the specimens, and the respective
corrosion potentials, current densities and corrosion rates
determined from it are shown in Table 2. Corrosion rate of
the untreated alloy was the lowest (% 0.68 mm/year), whereas it
was higher for all the DCTed alloys. Corrosion rate of the
treated alloys increased with the increase in DCT time from 4 to
16 h, while the alloy subjected to DCT for 16 h exhibited the
highest corrosion rate (%0.95 mm/year). XRD patterns ob-
tained from the corroded surfaces of the untreated alloy as well
as the 16-h DCTed alloy, as a representative picture for the
DCTed alloys, are shown in Fig. 12. The presence of Mg(OH)2
phase was confirmed as corrosion product. Intensity of the
peaks corresponding to Mg(OH)2 was more for the corroded
specimen of the treated alloy indicating more corrosion in the
DCTed specimen. The a-Mg phase being more anodic com-
pared to Al4RE phase was preferentially corroded. Figure 13
shows the SEM micrographs taken from the corroded surfaces
of the untreated and DCTed alloys. A thin and uniform
corroded film was observed on the surfaces of the untreated
alloy. However, the corroded film on the surface of the DCTed
alloys was relatively discontinuous, irregular and uneven
resulting in poor corrosion resistance of the DCTed alloys.
The dense and continuous network of the intermetallic
Al4RE phase present along the grain boundaries acts as a
corrosion-resistant barrier, which increased the overall corro-
sion resistance of the untreated alloy (Ref 14). However, the
reduced amount of the same could be the plausible cause for the
inferior corrosion resistance of the DCTed alloys.
4. Conclusions
The effect of DCT on the microstructure and mechanical
properties including corrosion behavior of the creep-resistant
squeeze-cast AE42 alloy was investigated. For comparison, the
same was also studied on the AE42 alloy without DCT. The
following conclusions are drawn:
1. Both the untreated and DCTed alloys comprised primary
Mg (a-Mg) and lamellar eutectic consisting of alternate
layers of a-Mg and Al4RE phases.
2. Volume fraction of the Al4RE phase in the AE42 alloy
reduced gradually following DCT carried out from 4 to
16 h.
3. The ductility and UTS of all the DCT alloys increased
significantly with a marginal increase in YS. The
improvement was attributed to the dissolution of the brit-
tle Al4RE phase following DCT. Among the alloys em-
ployed, the best tensile properties were obtained for the
16-h DCT alloy due to its lowest content of the brittle
Al4RE phase.
4. Creep resistance of the DCTed alloys was lower than that
of the untreated alloy owing to the presence of less
amount of thermally stable intermetallic Al4RE phase.
5. The hardness and wear resistance of the AE42 alloy re-
duced following DCT owing to the reduced volume frac-
tion of harder Al4RE phase.
6. The untreated alloy exhibited the best corrosion resis-
tance, whereas the DCT on the alloy deteriorated its cor-
rosion resistance. The poor corrosion resistance of the
DCT alloys was attributed to the reduced amount of
Al4RE phase.
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