H. LIU et al.: CHARACTERIZATION OF AZ31B MAGNESIUM ALLOY JOINTS WELDED WITH A Nd:YAG LASER
487–492
CHARACTERIZATION OF AZ31B MAGNESIUM ALLOY JOINTS
WELDED WITH A Nd:YAG LASER
KARAKTERIZACIJA ZVARNEGA SPOJA MAGNEZIJEVE ZLITINE
AZ31B, IZDELANEGA Z Nd:YAG LASERJEM
Hongtao Liu
1
, Jixue Zhou
1
, Yanfei Chen
2
, Tao Li
2
, Xuebing Jiao
3
, Yuansheng Yang
4
, Tao Lin
2
,
Kaiming Cheng
1
1
Shandong Key Laboratory for High Strength Lightweight Metallic Materials, Advanced Materials Institute, Shandong Academy of Sciences,
Jinan 250014, China
2
Shandong Engineering Research Center for Lightweight Automobiles Magnesium Alloys, Advanced Materials Institute, Shandong Academy
of Sciences, Jinan 250014, China
3
College of Material Science and Engineering, Shandong Jianzhu University, Jinan 250101, China
4
Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
kelekm73@163.com
Prejem rokopisa – received: 2017-10-26; sprejem za objavo – accepted for publication: 2018-02-15
doi:10.17222/mit.2017.180
In the present study, an AZ31B magnesium alloy was welded using a low-power pulsed Nd:YAG laser attached to a six-axis
welding robot. The welding metallurgical characteristics, including microstructure, porosity formation, texture evolution and
microhardness distribution were systematically investigated using a scanning electron microscope (SEM), electron backscatter
diffraction (EBSD) and a microhardness tester. The transverse cross-sectional microstructure of the AZ31B-alloy joint was
composed of the base metal (BM), columnar-grain zone (CGZ) and equiaxed-grain zone (EGZ) in the center of the fusion zone
(FZ). Some pores were observed in the FZ and believed to be related to the pre-existing pores in the BM. Especially, it was
noted that many columnar grains observed in the transverse cross-section of the FZ were almost perpendicular to the
solid-fusion boundary. Nevertheless, a large number of fine equiaxed grains formed on the top surface of the FZ and exhibited
obvious random orientations. Different grain-growth mechanisms found on the transverse cross-section and the top surface of
the weld bead were empirically investigated. Owing to apparent grain coarsening, the lowest microhardness value (~ 43 HV)
was detected in the center of the FZ, which was evidently lower than that of the BM (~ 68 HV).
Keywords: AZ31B magnesium alloy, Nd:YAG laser welding, microstructure, pores, texture, electron backscatter diffraction
Avtorji {tudije so magnezijevo zlitino AZ31B varili s pulznim Nd:YAG laserjem nizke mo~i, name{~enem na {estosnem
varilnem robotu. Sistemati~no so izvedli metalur{ko karakterizacijo varjenja, vklju~no z analizo mikrostrukture, analizo tvorbe
poroznosti in razvoja teksture s pomo~jo vrsti~ne elektronske mikroskopije (SEM) v kombinaciji s sipanjem povratnih
elektronov (EBSD). Porazdelitev mikrotrdote po preseku zvara so dolo~ili z merilnikom mikrotrdote. Posnetki mikrostrukture v
pre~nem preseku zvarnega spoja zlitine AZ31B so pokazali, da je le ta sestavljen iz osnovne kovine (BM; angl.: base metal),
cone stebri~astih kristalnih zrn (CGZ; angl.: columnar grains zone), cone enakoosnih zrn (EGZ; angl.: equiaxed grains zone) in
centralne staljene cone (FZ, angl.: fusion zone). V FZ so opazili tudi nekaj por in avtorji verjamejo, da je to posledica
predhodno obstoje~ih por v BM. [e posebej poudarjajo opa`anje, da se mnoga stebri~asta zrna nahajajo v pre~nem preseku FZ ,
ki so skoraj pravokotna na fazno mejo trdno-talina. Kakorkoli, na vrhu povr{ine FZ se nahaja veliko {tevilo tvorjenih drobnih
enakoosnih kristalnih zrn z o~itno naklju~no orientacijo. Avtorji prispevka empiri~no razpravljajo v ~lanku tudi o razli~nih
mehanizmih rasti zrn med pre~nim presekom in vrhom povr{ine varilne posteljice (bazen~ka staljene kovine). Zaradi jasne rasti
zrn so najmanj{o mikrotrdoto (~ 43 HV) izmerili v centru FZ, kar je o~itno precej manj kot v BM (~ 68 HV).
Klju~ne besede: magnezijeva zlitina AZ31B, varjenje z Nd:YAG (Neodim:Itrij-Aluminijev granat) laserjem, mikrostrukture,
pore, teksture, sipanje povratnih elektronov
1 INTRODUCTION
Being promising structural metallic materials, mag-
nesium alloys have gained increased interest in the
automotive and aircraft industries for their light weight,
good electromagnetic shielding, and excellent castabi-
lity.
1–3
However, the poor formability and cold process-
ability of magnesium alloys restrict their widespread
engineering applications. This is because magnesium has
a hexagonal close-packed (HCP) crystal structure that
cannot activate enough slip systems at room temperature.
Therefore, the high-performance welding technologies
are essential for further expanding the application fields
of Mg alloys. Various welding technologies have been
applied to weld Mg alloys, such as tungsten-inert-gas
(TIG) welding,
4,5
laser-beam welding (LBW),
6,7
laser-arc
hybrid welding,
8
and friction-stir welding (FSW).
9
Among these welding methods, LBW is suitable to
weld magnesium alloys for its remarkable advantages,
including low residual stress, high-power density and
large depth-to-width ratio.
10,11
After the laser welding,
the solidification in the fusion zone (FZ) produces signi-
ficant microstructural changes. The reliability of a
welded joint depends strongly on the microstructural
changes in the FZ and the base material (BM) after the
welding.
12
Nevertheless, limited work regarding the mi-
crostructure including the porosity origin, grain-growth
Materiali in tehnologije / Materials and technology 52 (2018) 4, 487–492 487
UDK 67.017:669.721.5:621.791.318 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 52(4)487(2018)

mechanism and texture evolution after the LBW of mag-
nesium alloys has been reported.
Therefore, laser welding of AZ31B Mg alloys was
performed in the present study, using a low-power pulsed
Nd:YAG laser attached to a six-axis welding robot. The
solidification microstructure, grain growth, porosity
formation, texture evolution and microhardness distri-
bution were systematically investigated to understand the
metallurgical characteristics of an automatic-laser-
welded AZ31B alloy joint.
2 EXPERIMENTAL PART
In the present study, 2-mm-thick as-rolled AZ31B
Mg alloy sheets were selected for laser welding; their
chemical composition in weight percent is listed in
Table 1. Before the laser welding, the AZ31B alloy
sheets were cleaned with acetone to remove grease and
then polished with SiC papers to remove oxides. As
shown in Figure 1, a low-power pulse Nd:YAG laser
attached to a six-axis robot was employed for the
welding research. The laser butt welding was performed
on an AZ31B specimen under the conditions of 600 W
laser power, 9.2 mm/s welding speed, and 600 μm
focused beam spot. During the laser-welding process, the
argon shielding gas was ejected to protect the top and
bottom surfaces at 15 L/min and 20 L/min, respectively.
Table 1: Chemical composition of the AZ31B sheets, in mass
fractions (w/%)
Al Zn Mn Fe Si Cu Mg
3.15 1.00 0.29 <0.003 <0.01 <0.001 Balance
After the laser welding, the welded joint was aged at
473 K for 1 h. The microstructures of weld samples were
examined with optical microscopy (OM) and scanning
electron microscopy (SEM) after mechanical polishing
and subsequent etching in a corrosive solution of 4.2 g
picric acid, 70 mL ethanol, 10 mL glacier acetic and 10
mL deionized water. SEM observations were conducted
using a Zeiss EVO MA10 SEM in the second electron
detection mode at 20 kV.
For EBSD mapping and a texture analysis of the
joint, EBSD measurements were performed under the
conditions of 20 kV electron beam, 10 nA probe current,
17 mm working distance, 20 ms integration time and
0.5 μm step size. The EBSD specimens were first mecha-
nically polished, followed by electrolytic polishing in a
commercial AC2 solution at 16 V for 80 s at –15 °C.
Electron backscatter diffraction (EBSD) was done using
SEM (Zeiss EVO MA10) together with a Nordlys-S
detector, a CCD camera, the HKL data-acquisition
software and the Channel 5.0 post-processing-analysis
software package. The obtained EBSD map covered the
BM and the FZ of the welded specimen. To analyze the
texture evolution, pole figures were produced from the
EBSD map of the BM and FZ.
A Huayin HV-1000 type microhardness tester was
used to measure the Vikers microhardness across the
seam after the specimen was polished. A 200-g load and
20-s dwell time were applied for the hardness tests. An
adequate space was retained to avoid interaction effects
between the adjacent indentations, and three measure-
ments were taken to determine the error bars.
3 RESULTS AND DISCUSSION
3.1 Appearance and microstructure
A macro seam appearance is an important evidence
of the welding qualities. As observed in Figure 2,a
smooth, uniform and regular fish-scale pattern of the
seam was observed on the weld surface. The top and
bottom widths of the welding bead were about 2 mm and
1 mm, respectively. Moreover, no welding defects, such
as cracks, incomplete fusion or slag inclusion were
found.
The typical microstructure of the transverse cross-
section of the AZ31B laser weld is shown in Figure 3.
This joint microstructure can be divided into three zones
with different microstructural features, including BM,
columnar-grain zone (CGZ) near the fusion boundary,
and equiaxed-grain zone (EGZ) in the center of the FZ.
H. LIU et al.: CHARACTERIZATION OF AZ31B MAGNESIUM ALLOY JOINTS WELDED WITH A Nd:YAG LASER
488 Materiali in tehnologije / Materials and technology 52 (2018) 4, 487–492
Figure 2: Appearance of the AZ31B laser seam: a) top surface, b)
bottom surface
Figure 1: Nd:YAG laser-equipped six-axis robot

Owing to the low heat input during the laser welding, no
apparent heat-affected zone (HAZ) was observed. In the
CGZ, the ratio of length over width of columnar grains
was about 10:1. The CGZ and EGZ were induced by
different solidification process during the laser-welding
process.
Laser-beam welding has a high, focused heat input.
The heat generated during LBW is rapidly extracted
from the molten fusion zone by the surrounding colder
plate. In the vicinity of the fusion boundary, a relatively
large thermal gradient contributes to the appearance of
columnar grains whereas in the center of the FZ, equ-
iaxed grains form because of a lower temperature gra-
dient.
13
Thereby, a grain morphology transition in the FZ
was observed, staring in the CGZ near the fusion line
and finishing in the EGZ in the center of the FZ (Fig-
ure 3).
3.2 Porosity formation
On Figure 3, it is also worth noting that no hot
cracks are shown in the microstructure of the joint, but
some small pores are observed in the BM and FZ, indi-
cated by the arrows in Figure 3a. The pore formation is
one of the major issues when laser welding magnesium
alloys as pores deteriorate the mechanical properties of
joints, particularly the tensile strength and elongation.
10
The pore formation can be attributed to many factors
such as hydrogen pores,
14
keyhole collapse,
15
initial
pores in the BM,
15,16
and capture gases from the air.
17
Owing to the active chemical properties, an Mg alloy is
prone to form MgO and H
2
through a chemical reaction
with H
2
O from the air. Along with the molten-Mg
solidification, the H
2
solubility sharply decreases in the
solid phase and then pores are formed in the weld bead.
The other important reason for a pore formation is
the fact that a molten pool captures the gas from the air
during a laser-welding process. In the current work, both
the top and bottom surfaces were shielded by argon gas
during the laser welding. As mentioned above, the seam
of the AZ31B alloy joint was very narrow, about 2 mm
(Figure 2). Thus, the size of the molten pool was small
during the laser welding. In this case, the argon gas
could shield the molten pool soundly. Therefore, the
pores could not be induced by the hydrogen pores and
entrapment gases from the air in this welding process.
In addition, the stability of the keyhole during laser
welding, which strongly depends on the surface tension
and vapor pressure, is another reason for a pore forma-
tion. Generally, Mg-alloys have a more stable keyhole
H. LIU et al.: CHARACTERIZATION OF AZ31B MAGNESIUM ALLOY JOINTS WELDED WITH A Nd:YAG LASER
Materiali in tehnologije / Materials and technology 52 (2018) 4, 487–492 489
Figure 3: a) Cross-sectional microstructure of the laser-welded AZ31B joint, b) larger magnification of the cross-section
Figure 4: Microstructure of the AZ31B base material

than Al-alloys during a welding process because of their
characteristics, including a much higher vapor pressure
and a lower surface tension. It is shown in Figure 2 that
after the laser welding, no seam collapse was exhibited
in the AZ31B joint. Thereby, an induction of a pore for-
mation due to an unstable keyhole used in this study was
not likely to occur.
As indicated by the arrows in Figures 3a and 4, the
pores were observed in both the BM and the FZ, demon-
strating that the presence of porosity in the FZ can be
attributed to the initial small pores in the BM. During the
laser-welding process, the pre-existing pores in the BM
might grow and coalesce in the molten pool. Similarly,
Zhao and Debroy
15
observed that the pre-existing pores
in the BM coalesced and expanded in a metal weld dur-
ing the alloy welding. Wahba et al.
16
also reported that
the pre-existing pores in the BM were prone to form
pores, which had few chances to escape from the molten
weld pool during the laser welding of Mg alloys.
3.3 Grain growth and texture evolution
Another noticeable feature from Figure 3 is the fact
that well-developed columnar grains are nearly perpen-
dicular to the fusion line in the CGZ. This is attributed to
the epitaxial growth of the grains adjacent to the solid-
liquid line.
18
As the BM could act as the nucleation sub-
strate during the laser welding, the solidification-grain
orientation was strongly affected by the BM grains
adjacent to the solid-liquid line (indicated by double
arrows in the CGZ of Figure 3b).
19
In Mg alloys, both
<1010>-type and a-type axes <1120> directions of the
hcp-crystal lattice are close-packed crystallographic
directions, which are the preferred growth orientations
during a solidification process.
20,21
Furthermore, the
crystal growth within the fusion zone is also significantly
influenced by the heat-conduction direction (the thermal
gradient) during the solidification process and the
maximum value of the thermal gradient is perpendicular
to the fusion boundary.
12
Therefore, the grains whose
close-packed orientation is parallel to the direction of the
maximum thermal gradient, grow more rapidly and
restrict other, less favorably oriented grains.
As shown in Figures 5 and 6, the texture evolution
and microstructure of the weld top surface were anal-
yzed. Contrary to the columnar grains observed in the
cross-section of the FZ (Figure 3), many equiaxed grains
with random orientations were detected on the top
surface of the FZ, as illustrated in Figure 5a. The BM
grains revealed a strong {0001} texture, whose maximum
orientation density was about 7.2 (Figure 5b) while the
maximum grain-orientation density on the top surface of
the FZ was only about 2.6, as shown in Figure 5c, mani-
festing almost random grain orientations.
Compared with the welding-pool inside, the welding
top surface has a lower thermal gradient G and a much
faster solidification rate R because the shielding-gas flow
can carry the heat away rapidly. Thus, a low G/R value
determined the equiaxed-grain formation on the top sur-
face.
11
Owing to the fast solidification rate, the grains on
the top surface of the FZ did not have sufficient time to
grow up, and the texture was hardly influenced by the
BM grains, having the close-packed crystallographic
directions of (<1010> and <1120>) near the fusion boun-
dary. Consequently, a large number of fine equiaxed
grains that had no preferred crystalline orientation were
formed on the top surface of the welding seam. Some
researchers
11,22,23
claimed that a high thermal gradient
would contribute to the growth of columnar grains, while
a relatively low thermal gradient could promote the
formation of equiaxed grains. These findings are in good
agreement with the results observed in Figures 3 and 5a.
Su et al.
21
studied the electron-beam-welding (EBW)
behavior of Mg-Al-based alloys and found that no
apparent columnar grains were formed, while many fine
equiaxed grains were observed in the Mg-alloy welding
bead. Similarly, they also attributed the absence of co-
H. LIU et al.: CHARACTERIZATION OF AZ31B MAGNESIUM ALLOY JOINTS WELDED WITH A Nd:YAG LASER
490 Materiali in tehnologije / Materials and technology 52 (2018) 4, 487–492
Figure 6: Top-surface microstrucutre of the AZ31B joint seam
Figure 5: a) Top-surface EBSD map of the laser welded AZ31B joint
seam, b) {0001} pole figure of the BM, c) {0001} pole figure of the
FZ

lumnar grains to a high cooling rate (a fast solidification
rate) during the EBW of Mg-alloys.
3.4 Microhardness distribution
The microhardness profile across the AZ31B joint is
illustrated in Figure 7, showing that the microhardness
changes remarkably between the BM and FZ. The maxi-
mum hardness was located in the BM and it was about
68 HV. The lowest hardness value (~ 43 HV) was de-
tected in the center of the FZ, and it was evidently lower
than that in the BM. This variation in the microhardness
can be attributed to the microstructural changes in the
BM and FZ.
According to the Hall-Petch relationship,
24
grain
coarsening can reduce the hardness of an alloy. In
contrast, the Orowan hardening mechanism indicates that
a large number of hard precipitate particles within a
matrix can improve the hardness of an alloy.
25
As shown
in Figure 3, the grain coarsening of  -Mg was observed
in the FZ of the AZ31B welded joints. Furthermore, a
very small amount of precipitates appeared in the
AZ31B alloy joints. Consequently, the grain-coarsening
mechanism dominated the microhardness distribution in
this research.
4 CONCLUSIONS
Laser welding of an AZ31B Mg alloy was conducted
using a low-power pulsed Nd:YAG laser. The metallur-
gical characteristics of the AZ31B alloy joint were
systematically investigated in the present work. Owing to
different thermal gradients, a grain-morphology tran-
sition was observed in the FZ, taking place from the
CGZ near the fusion line towards the EGZ in the center
of the FZ. Some small pores were observed in the FZ
and believed to be caused by the pre-existing pores in the
BM. An epitaxial growth of the columnar grains in the
FZ, which were nearly perpendicular to the solid-liquid
boundary, was observed in the cross-section of the FZ.
This was caused by the maximum temperature gradient
and the close-packed crystallographic directions
(<1010> and <1120>) of the Mg alloys. Being different
from the columnar grains in the welding pool, fine
equiaxed grains with random orientations were formed
on the top surface of the welding bead due to a relatively
high cooling rate and a low G/R value on the top surface.
Because of apparent grain coarsening of the FZ, the
minimum hardness value (~ 43 HV) was detected in the
center of the FZ and it was clearly lower than that of the
BM (~ 68 HV).
Acknowledgement
This work was supported by the Natural Science
Foundation of Shandong Province, China (grant no.
ZR2018BEE013), the National Key Research and Deve-
lopment Plan of China (grant no. 2017YFB0103904,
2016YFB0701202), the Shandong Province Key
Research and Development Plan, China (grant no.
2017CXGC0404, 2016ZDJS02A09), and the Youth
Foundation of Shandong Academy of Sciences, China
(grant no. 2016QN015).
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