Sc/Zr microalloying on strength-corrosion performance synergy of wire-arc directed energy deposited Al-Mg

Abstract

WADED Al-Mg and Al-Mg-Sc-Zr aluminium alloy thin-wall components were produced. Heat treatment (350°C aging for 4 h) was further conducted. The effects of Sc/Zr addition on porosity, microstructure, mechanical properties and corrosion behaviour were investigated. With Sc/Zr adding, pore area fraction decreased from 0.208 to 0.005% and grain size decreased from 84.34 to 19.54 μm. In WADED Al-Mg-Sc-Zr component, microstructure showed refined equiaxed grains. Mechanical properties were improved with 360 MPa ultimate tensile strength (UTS), 185 MPa yield strength, and 23.31% elongation. After heat treatment, UTS increased to 388 MPa with good ductility (22.53%) maintained. The grain refinement and Al3(Sc, Zr) particles were the key factors for mechanical property improvement. The corrosion behaviour was improved with introduced Sc/Zr and further optimised after heat treatment. The corrosion exhibited layered characteristics due to the distribution of primary Al3(Sc, Zr) while minor-size secondary Al3(Sc, Zr) along grain boundaries improved the corrosion behaviour.

Full text

Sc/Zr microalloying on strength-corrosion performance synergy of wire-arc
directed energy deposited Al-Mg
Yubin Zhou
a
*, Zewu Qi
a,b
*, Baoqiang Cong
a
, Zihao Jiang
a
, Xinyi Cai
a
, Hongwei Li
a,c
, Jiaqi Zhang
a
,
Chaofang Dong
d
, Xing He
d
, Sanbao Lin
b
, Xiaoyu Cai
b
and Bojin Qi
a
a
School of Mechanical Engineering and Automation, Beihang University, Beijing, People’s Republic of China;
b
State Key Laboratory of
Advanced Welding and Joining, Harbin Institute of Technology, Harbin, People’s Republic of China;
c
Beijing Hangxing Machinery
Manufacturing Co. Ltd, Beijing, People’s Republic of China;
d
Beijing Advanced Innovation Center for Materials Genome Engineering, Key
Laboratory for Corrosion and Protection (MOE), Institute for Advanced Materials and Technology, University of Science and Technology
Beijing, Beijing, People’s Republic of China
ABSTRACT
WADED Al-Mg and Al-Mg-Sc-Zr aluminium alloy thin-wall components were produced. Heat
treatment (350°C aging for 4 h) was further conducted. The eects of Sc/Zr addition on
porosity, microstructure, mechanical properties and corrosion behaviour were investigated. With
Sc/Zr adding, pore area fraction decreased from 0.208 to 0.005% and grain size decreased from
84.34 to 19.54 μm. In WADED Al-Mg-Sc-Zr component, microstructure showed refined equiaxed
grains. Mechanical properties were improved with 360 MPa ultimate tensile strength (UTS), 185
MPa yield strength, and 23.31% elongation. After heat treatment, UTS increased to 388 MPa
with good ductility (22.53%) maintained. The grain refinement and Al
3
(Sc, Zr) particles were the
key factors for mechanical property improvement. The corrosion behaviour was improved with
introduced Sc/Zr and further optimised after heat treatment. The corrosion exhibited layered
characteristics due to the distribution of primary Al
3
(Sc, Zr) while minor-size secondary Al
3
(Sc,
Zr) along grain boundaries improved the corrosion behaviour.
ARTICLE HISTORY
Received 24 February 2024
Accepted 16 May 2024
KEYWORDS
Wire-arc directed energy
deposition; Al-Mg-Sc-Zr
alloy; heat treatment;
microstructure and
mechanical properties;
corrosion behaviour
1. Introduction
Wire-arc directed energy deposition (WADED) is a
technology that adopts the arc heat source to melt
filling metal wire and directly deposits metallic struc-
tures. Compared with alternative additive manufactur-
ing (AM) techniques, WADED has notable advantages
of high eciency, economical cost and capacity for
large-scale integrated manufacture of aluminium alloy
structures [1–5]. Al-Mg aluminium alloy components
were extensively used in aerospace industry because
of their lightweight application and excellent strength-
plastic synthesis [6]. Many scholars have proven the Al-
Mg alloys fabricating feasibility with WADED process
[7,8]. Meanwhile, the dimensional accuracy optimisation
and defect elimination in WADED process was also con-
ducted by numerous works [9–13]. However, there are
still shortcomings in WADED Al-Mg aluminium alloys,
like porosity and insucient strength (<300 MPa) [14].
Moreover, with the growing requirements of aerospace
aluminium, the mechanical properties and durability
need to be considered simultaneously. Several attempts
have been conducted to solve these problems, including
process parameter optimisation [15], post-treatments
[16,17] and alloy element adding [18,19]. Among these
methods, the microalloying of Sc/Zr elements to
modify Al-Mg alloys has attracted wide attention and
become a key issue in the new-generation high-per-
formance aluminium alloy research field [20,21].
The introduction of Sc/Zr into Al-Mg alloy was often
achieved by laser-based powder bed fusion (PBF-LB)
technology in the early stage. Previous research
reported the Sc/Zr addition could improve the printabil-
ity of normal Al-Mg alloy [22]. Zhou et al. [23] reported
the Zr addition into Al-Mg alloy can eectively inhibit
hot cracks and reduce defect dimension. The main
impact of Sc/Zr can be attributed to two types of
Al
3
(Sc, Zr) precipitates. The primary Al
3
(Sc, Zr) precipi-
tates served as nuclei for grain refinement and increased
© 2024 The Author(s). Published by Informa UK Limited, trading as Taylor & Francis Group
This is an Open Access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use,
distribution, and reproduction in any medium, provided the original work is properly cited. The terms on which this article has been published allow the posting of the Accepted
Manuscript in a repository by the author(s) or with their consent.
CONTACT Zewu Qi qizewu@buaa.edu.cn School of Mechanical Engineering and Automation, Beihang University, Beijing 100191, People’s
Republic of China State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, People’s Republic of China;
Bojin Qi qbj@buaa.edu.cn School of Mechanical Engineering and Automation, Beihang University, Beijing 100191, People’s Republic of China
*These authors contributed equally to this work.
VIRTUAL AND PHYSICAL PROTOTYPING
2024, VOL. 19, NO. 1, e2358981 (21 pages)
https://doi.org/10.1080/17452759.2024.2358981

resistance to heat tears while solidifying [24]. The
secondary Al
3
(Sc, Zr) precipitating in aging treatment
provided precipitation strengthening and further
enhanced the strength [22,25]. Wang et al. [10,20,21]
discovered that fine grains formed at the molten pool
boundary (MPB) due to the primary Al
3
(Sc, Zr) precipitat-
ing at fusion boundary. Columnar grains existed inside
the molten pool because the majority of Sc/Zr were
trapped in the matrix [26]. Qin et al. [27] thought the
Sc/Zr elements in coarse columnar grains were supersatu-
rated, and secondary Al
3
(Sc, Zr) preferentially formed in
this area throughout aging treatment. The disparate dis-
tribution of primary and secondary precipitates resulted
in dual heterogeneous structures in PBF-LB counterparts.
The attempts to add Sc/Zr into WADED aluminium
alloys have been proven feasible by several studies.
Dong et al. [28] introduced Sc into WADED Al-Zn-Mg-
Cu-Sc alloy, the average grain diameter reduced to 59.6
μm from 123 μm. Guo et al. [29] fabricated Al-Zn-Mg-
Cu-Sc-Zr counterparts using cold metal transfer (CMT)
technology and strength could reach 600 MPa after T6
heat treatment. Ren et al. [19,30] reported that Sc addition
could significantly improve the mechanical properties in
as-deposited conditions. Xia et al. [31] further studied
the microstructure evolution in WADED Al-Mg-Sc com-
ponent and confirmed the grain refinement and property
enhancement (with 335.58 MPa ultimate tensile strength
and 22.74% elongation). The Sc/Zr adding not only had
eects on microstructure but also inuenced mechanical
properties simultaneously. However, considering the
layered characteristics and dynamic thermal cycling
caused by WADED process [32,33], the microstructure
and properties of the materials prepared by WADED are
dierent from those of the components manufactured
by other processes. The comprehensive properties of
WADED Al-Mg-Sc-Zr alloy have not been systematically
elaborated. Studying the Sc/Zr adding eects on
WADED Al-Mg component is important to provide
support for Al-Mg-Sc-Zr WADED process optimisation.
Besides, the durability of large-scale WADED aero-
space Al-Mg components in the practical service
environment is also required to be evaluated. According
to previous research, the galvanic corrosion between
precipitation and matrix is the main eect factor on cor-
rosion behaviour after Sc/Zr introduction [34,35]. Zhang
et al. [36] found the primary Al
3
(Sc, Zr) precipitates
acting as micro-cathodic gathered along the MPBs in
PBF-LB components, and corrosion pittings initiated
along MPB. However, diverse viewpoints exist on the
impact of secondary Al
3
(Sc, Zr) precipitates on corrosion
behaviour. Deng et al. [37] reported the corrosion behav-
iour of Al-Zn-Mg-Sc-Zr sheet was enhanced at 120°C as
aging time prolonging because the active corrosion
path was weakened and the dissolving rate was
decreased by the coarsen and dispersed grain boundary
precipitates (GBPs). However, Zhang et al. [34] reported
the secondary Al
3
(Sc, Zr) precipitates along grain bound-
aries (GBs) in PBF-LB counterparts caused more severe
corrosion, and only precipitate size below 10 nm was ben-
eficial for corrosion resistance. Zhou et al. [38] thought the
size and density of secondary Al
3
(Sc, Zr) particles in PBF-
LB counterparts were mostly responsible for the corrosion
behaviour, and the growth of precipitates during the
aging process would weaken the corrosion resistance.
As mentioned above, there is still no conclusion about
the eects of secondary Al
3
(Sc, Zr) on corrosion behav-
iour, thus more factors are needed to be considered com-
prehensively to illustrate the corrosion mechanism.
Inspired by the above works, the process and
dynamic thermal cycle of WADED are dierent from
those in metallurgical forming and laser-based powder
bed fusion, which result in distinctive microstructure,
mechanical properties and corrosion behaviour. In
addition, the heat treatment process for WADED Al-
Mg-Sc-Zr has not been determined. WADED Al-Mg and
WADED Al-Mg-Sc-Zr components were fabricated and
heat treatments were conducted in this paper. Investi-
gations were conducted on the porosity, microstructure,
mechanical properties and corrosion behaviour. This
paper aims to comparatively study and systematically
reveal the eects of Sc/Zr addition on the properties of
WADED Al-Mg alloy. This is beneficial to provide
support for Al-Mg-Sc-Zr WADED process optimisation
and potential utilisation in aerospace industry.
2. Materials and methods
2.1. Material preparation
WADED Al-Mg and Al-Mg-Sc-Zr components are fabri-
cated on 5A06 aluminium alloy substrate (300 mm ×
150 mm × 10 mm) with single-pass and multi-layer
Table 1. Compositions of wire and WADED components.
Alloys
Compositions (wt. %)
Mg Fe Mn Si Sc Ti Zr Al
Al-Mg wire 6.52 0.08 0.73 0.01 – 0.14 – Bal.
Al-Mg-Sc-Zr wire 6.71 0.01 0.10 0.01 0.30 0.03 0.13 Bal.
Al-Mg-Sc-Zr components 6.3 0.01 0.09 0.01 0.29 0.03 0.14 Bal
2 Y. ZHOU ET AL.

depositing strategy. Al-6.5Mg and Al-6.7Mg-0.3Sc-0.13Zr
wires with a 1.2 mm diameter were used as filler material.
Table 1 lists the compositions of wires and components
which were measured by inductively coupled plasma-
atomic emission spectrometry (ICP-AES, Agilent 7700-
CE). The WADED processing illustration is shown in
Figure 1, and process parameters are listed in Table 2.
Heat treatment (aging at 350°C for 4 h with air cooling)
was further conducted for WADED Al-Mg-Sc-Zr and the
component was referred to WADED Al-Mg-Sc-Zr-HT.
2.2. Microstructure and mechanical
characterisation
The sampling illustration of WADED components is shown
in Figure 2. The samples in the centre section were used for
investigating porosity, microstructure, microhardness and
corrosion behaviour characteristics. To prepare samples
for the horizontal and vertical tensile tests, each side was
chosen. The central specimens were systematically
detected using an optical microscope (OM) (ZEISS Scope
A1), X-ray diraction (XRD) (D/Max2500pc), following
grinding and mechanical polishing. The micro-scale
three-dimensional computed tomography system (3D-
CT, NanoVoxel-2000, developed by Sanying Precision
Instruments, China) was utilised to identify internal pore
defects in ϕ 3.5 mm × 5 mm. The samples etched by
Keller’s reagent were detected by OM, scanning electron
microscopy (SEM) (JEOL JSM-7900F), and energy disper-
sive spectrometry (EDS). After electrolytic polishing, the
central sections were investigated using 0.5 µm step
length electron back-scattered diraction (EBSD). Follow-
ing ion milling and mechanical thinning, the specimens
were detected using transmission electron microscopy
(TEM) (JEM-2100) to achieve the composition and struc-
tural information of precipitations.
Microhardness was measured by Vickers (FALCON
500) with a 200 g load for 15 s at 0.5 mm intervals in a
7 mm × 7 mm region. The tensile properties were exam-
ined using the universal testing equipment (Instron
5565) with a strain rate of 1 mm/min. The fracture speci-
mens were further analysed by SEM.
The thermodynamic simulation was performed using
Thermo-Calc simulation based on the TCAL6 database.
For Al-6.7Mg-0.3Sc-0.13Zr and Al-6.5Mg-0.73Mn alloys,
elemental segregation during rapid solidification was
Figure 1. Diagrammatic representation of WADED process and the as-deposited WADED Al-Mg-(Sc-Zr) components.
Table 2. WADED processing parameters.
Parameter Value
Current I
+
&I
−
(A)
t
EN
&t
EP
(ms)
140 &140(A)
8 & 2
Wire feed speed (m/min) 2.4
Traveling speed (mm/min) 300
Ce-W electrode diameter (mm) 4.0
Arc length (mm) 5
Flow rate of argon shielding gas (L/min) 15
Interval time between two layers(s) 60
Building path One way
VIRTUAL AND PHYSICAL PROTOTYPING 3

assessed using the Scheil–Gulliver model. The phase
diagram was tested with TCAL6 thermodynamic database.
2.3. Corrosion test
Electrochemical experiments were conducted in 3.5wt %
NaCl solution at ambient temperature with CHI-660e
electrochemical workstation. The platinum plate and a
saturated calomel electrode (SCE) respectively served as
a counter electrode (CE) and reference electrode (RE).
All samples were mechanically ground up to 2000 grits
to obtain a working electrode (WE), and the WE surface
with 3 mm × 3 mm was checked to ensure no obvious
pores in the vision field. Cathodic potentiostatic polaris-
ation was carried out at −0.5 V relative to the open
circuit potential (OCP) for 5 min to eliminate eects of oxi-
dation film, then stabilised for 2400 s. With a frequency
range of 100 kHz to 10 mHz, electrochemical impedance
spectroscopy (EIS) was measured at corrosion potential.
The commercial programme ZSimpwin was used to
evaluate the experimental data. At a scanning rate of
0.1667 mV/s, the potentiodynamic polarisation was
started at −0.15 V
OCP
, and the scanning was halted
when current reached 10
−2
A/cm
2
. The immersion exper-
iment was conducted in 3.5 wt % NaCl solution at
ambient temperature for 24 h, SEM and EDS were used
to observe corrosion feature details.
3. Results
3.1. Microstructure
The 3D-reconstructed morphology, 2D micropore distri-
bution on building direction (BD) × traveling direction
(TD) plane, and micropore size statistics detected by 3D
X-ray computed tomography are displayed in Figure 3.
Most of the micropores in all samples are near-spherical,
thus it can be inferred hydrogen pore is the main form of
pores [39]. It can be observed a large number of micro-
pores distribute in WADED Al-Mg with a porosity of
0.208% and the maximum diameter of 168.72 μm. The
quantitative statistical result indicates the average diam-
eter is 15.71 ± 9.65 µm. The number fraction of micropores,
whose diameter is less than 30 µm, is 85.63%. In micropore
projection on BD × TD plane, a large number of micropores
aggregating in a specific layer leads to layered pore distri-
bution, which is also reported in the as-deposited WADED
aluminium alloys [40]. After Sc/Zr adding, the number and
size of micropores decreased significantly with an
improved porosity of 0.005%. The largest micropore only
has a diameter of 82.46 µm and the average pore size
decreased to 8.20 ± 8.95 µm. Consequently, the layered
pore distribution is obviously optimised. In particular, it is
essential to note that the number fraction of the micropore
diameter less than 10 and 30 µm is 54.16 and 95.21%,
respectively, indicating the pores in WADED Al-Mg-Sc-Zr
are mainly minor-size. However, WADED Al-Mg-Sc-Zr-HT
exhibits a porosity of 0.045%. The average diameter is
14.41 ± 10.62 µm which is higher than the as-deposited
counterpart. The maximum pore diameter also increases
to 117.85 µm and the micropore number fraction less
than 10 and 30 µm is 34.86 and 86.75%, respectively.
This could be attributed to the heat treatment process
making the micropores gather and grow [41,42]. Although
the layered pore distribution in Sc/Zr-added component
tends to be obvious after aging, it still presents an
improved state compared with WADED Al-Mg. The layer-
distributed micropores will inuence component
synergy performance.
Figure 4 displays the XRD patterns of WADED com-
ponents. The apparent peaks in WADED Al-Mg com-
ponent are identified as α-Al and no other obvious
precipitation phase is identified. The main peak intensity
of Sc/Zr-added components increased, which could be
attributed to the overlapped peak of α-Al and Al
3
(Sc,
Zr) [31,43]. Moreover, the minor peak of Al
3
(Sc, Zr)
Figure 2. Schematic illustration of sampling position and tensile sample dimensions.
4 Y. ZHOU ET AL.

Figure 3. The porosity and statistic results of (a
1
)–(a
3
) WADED Al-Mg, (b
1
)–(b
3
) WADED Al-Mg-Sc-Zr, (c
1
)–(c
3
) WADED Al-Mg-Sc-Zr-HT
components.
Figure 4. XRD patterns of WADED Al-Mg, Al-Mg-Sc-Zr and Al-Mg-Sc-Zr-HT components.
VIRTUAL AND PHYSICAL PROTOTYPING 5

occurred at 64.5° and 77.5° in WADED Al-Mg-Sc-Zr-HT,
which relates to the secondary Al
3
(Sc, Zr) precipitating
in aging process.
The OM (Figure 5) and SEM (Figure 6) morphologies
are conducted to reveal the microstructure character-
istics of WADED components, and the composition of
the corresponding phases are listed in Table 3. Each
layer has alternate inner-layer and inter-layer regions
because of the layer-by-layer depositing mechanism
[41]. Figure 5a shows the homogeneous microstructure
of WADED Al-Mg component composed of coarse
equiaxed grains (CEG). For WADED Al-Mg component
in Figure 6a
1
–a
4
, the slight grain size dierence results
in its layered characteristic not being obvious. From
the corresponding EDS result and the previous studies,
the large-size phases distributed along the GB can be
identified as β-Al
3
Mg
2
and Al
6
(Mn, Fe) [14,19]. For
WADED Al-Mg-Sc-Zr component in Figure 5b, the grain
shows a smaller size and fine equiaxed grains (FEG) at
MPB leading to layered characteristics obvious, which
is dierent from the microstructure in the PBF-LB com-
ponent consisting of coarse columnar grains and fine
grains [44,45]. As Figure 6b
1
–b
4
shows, the CEG exists
in inner-layer region while the FEG exists in inter-layer
region. Sc/Zr-riched particles aggregated in inter-layer
regions can be identified as Al
3
(Sc, Zr) particles [46].
For WADED Al-Mg-Sc-Zr-HT sample in Figure 5c and
Figure 6c
1
–c
4
, the GBs tend to be semi-continuous. The
changed microstructure characteristics after Sc/Zr micro-
alloying would result in performance dierences.
TEM morphologies in Figure 7 show characteristics of
Al
3
(Sc, Zr) particles. As shown in Figure 7a
1
–a
3
, the
primary Al
3
(Sc, Zr) particles serve as nucleation sites and
often exist at the GB of FEG [21,27]. Selected area electron
diraction (SAED) images along the [110] zone axis show
Al
3
(Sc, Zr) particles with an L1
2
ordered structure and a
{110} crystal planar space compatible with the α-Al matrix.
Since the cooling rate during the forming process deter-
mines the primary Al
3
(Sc, Zr) size [47], the primary Al
3
(Sc,
Zr) precipitates in WADED Al-Mg-Sc-Zr component (10
3
-
10
5
K/s cooling rate) have a larger radius compared with
that in PBF-LB counterparts (∼10
6
K/s cooling rate) [20,48].
Figure 7b
1
and b
2
show the particles with small
radii are secondary Al
3
(Sc, Zr) precipitates [38]. The
precipitates along the GBs result in a narrower precipi-
tate-free zone (PFZ) compared with that in WADED
Figure 5. OM morphology of (a) WADED Al-Mg, (b) WADED Al-Mg-Sc-Zr and (c) WADED Al-Mg-Sc-Zr-HT components.
6 Y. ZHOU ET AL.

Figure 6. SEM morphology and element mapping of (a
1
)–(a
4
) WADED Al-Mg, (b
1
)–(b
4
) WADED Al-Mg-Sc-Zr and (c
1
)–(c
4
) WADED Al-
Mg-Sc-Zr-HT components.
VIRTUAL AND PHYSICAL PROTOTYPING 7

Al-Mg-Sc-Zr component. The secondary Al
3
(Sc, Zr)
particles generated in aging process are more preferen-
tially generated in the interior region of CEG since Sc/Zr
elements in coarse grains were supersaturated in the
matrix [27]. Both primary and secondary Al
3
(Sc, Zr) precipi-
tates exist in this condition [38]. The secondary precipi-
tates present an average radius of 24.5 nm in WADED
Al-Mg-Sc-Zr-HT. The HRTEM and SAED obtained from
the interface of secondary Al
3
(Sc, Zr) particles and the
matrix show L1
2
structure (Figure 7b
3
). Figure 7b
4
shows
the inverse Fourier transformation image from the area
marked by the yellow rectangle in Figure 7b
3
on the
Table 3. Chemical compositions of phase in Figure 6 (at. %).
Point
Compositions (at. %)
Al Mg Mn Fe Sc Zr
A1 82.6 1.0 11.7 4.7 – –
A2 87.0 11.6 1.2 0.2 – –
A3 84.0 2.3 9.4 4.3 – –
A4 92.5 7.2 0.3 – – –
B1 80.9 2.9 0.1 – 10.6 5.5
B2 87.8 6.9 0.1 – 3.6 1.6
B3 84.1 5.5 – 0.1 7.9 2.4
B4 93.1 6.6 – – 0.2 0.1
C1 92.0 7.6 – – 0.2 –
C2 91.3 7.7 – – 0.9 0.1
C3 77.4 1.1 – – 17.2 4.3
C4 91.7 8.0 – 0.1 0.1 –
Figure 7. TEM images and SAED patterns:(a
1
)–(a
3
) WADED Al-Mg-Sc-Zr, (b
1
)–(b
5
) WADED Al-Mg-Sc-Zr-HT, (b
3
) HRTEM picture of the
matrix and Al
3
(Sc, Zr) interface, (b
4
) dislocations shown by inverse Fourier transformation image from the area marked in (b
3
) on (002)
plane, (b
5
) HRTEM image of Al
3
(Sc, Zr) particle.
8 Y. ZHOU ET AL.

(002) plane. For the precipitates with a larger radius (39.61
nm), several dislocations exist at the interface of the pre-
cipitate and matrix (Figure 7b
4
). However, for the second-
ary Al
3
(Sc, Zr) particle with a small radius (6.11 nm) (Figure
7b
5
), the stacking faults (SFs) occurred. The various pre-
cipitate characteristics in dierent regions might lead to
heterogeneity of service performance.
The EBSD results of WADED components are illus-
trated in Figure 8. Grain refinement with Sc/Zr addition
is apparent based on the inverse pole figure (IPF) and
grain size distribution. The WADED Al-Mg component
shows an average grain diameter of 84.34 ± 29.58 μm.
WADED Al-Mg-Sc-Zr exhibits refined grains with an
average diameter of 19.54 ± 6.26 μm and the proportion
Figure 8. Inverse pole figure EBSD maps, grain size statistical results, grain boundary distribution and pole figures of (a
1
)–(a
4
) WADED
Al-Mg, (b
1
)–(b
4
) WADED Al-Mg-Sc-Zr and (c
1
)–(c
4
) WADED Al-Mg-Sc-Zr-HT components.
VIRTUAL AND PHYSICAL PROTOTYPING 9

of grains smaller than 15 μm accounted for about
23.18%. After aging, WADED Al-Mg-Sc-Zr-HT shows an
average grain diameter of 21.26 ± 6.47 and with about
27.91% grain less than 15 μm. The grains in WADED
Al-Mg-Sc-Zr did not grow obviously after aging due to
the thermal stability was improved with Sc/Zr addition
[49]. The grain size in WADED Al-Mg-Sc-Zr is comparable
to that of powder-based DED counterparts and nearly
half of that in as-cast counterparts [21,47,50]. Moreover,
there is no obvious grain growth after heat treatment.
The low-angle grain boundaries (LAGBs) proportion
with small misorientation angles (0°–15°) in WADED
Al-Mg, Al-Mg-Sc-Zr and Al-Mg-Sc-Zr-HT components
are 11.4%, 4.1%, 7.5%, respectively. The LAGB fraction
decreased with Sc/Zr adding and slightly increased
after aging. As the previous works reported, the percen-
tage of LAGBs enhanced as Al
3
(Sc, Zr) volume fraction
increased [38,51]. The pole figures show all WADED
specimens exhibit weak texture with the maximum
value in WADED Al-Mg, Al-Mg-Sc-Zr and Al-Mg-Sc-Zr-
HT components are respectively 2.534, 1.669 and
1.742, indicating no preferred orientation generated.
3.2. Microhardness
WADED Al-Mg and Al-Mg-Sc-Zr both exhibit homo-
geneous hardness distribution (Figure 9). The hardness
uctuates only in regional areas, which corresponds to
the alternation of FEG to CEG. The average hardness
value is enhanced from 91.7 HV
0.2
to 99.5 HV
0.2
after
Sc/Zr adding. After 350°C -4 h aging, the hardness of
WADED Al-Mg-Sc-Zr component further increases to
120.3 HV
0.2
. This could be attributed to the precipitation
of secondary Al
3
(Sc, Zr) which has a strengthening eect
[43,46].
3.3. Tensile properties
Figure 10 displays tensile curves and statistics results of
ultimate tensile strength (UTS), yield strength (YS) and
uniform elongation of samples. WADED Al-Mg com-
ponents demonstrate UTS of ∼330 MPa, YS of ∼155
MPa and elongation of ∼23.56% along horizontal direc-
tion. With Sc/Zr added, the UTS and YS respectively
increase to 360 and 185 MPa with good ductility
Figure 9. Microhardness of (a) WADED Al-Mg, (b) WADED Al-Mg-Sc-Zr; and (c) WADED Al-Mg-Sc-Zr-HT and (d) the statistic results.
10 Y. ZHOU ET AL.

(∼23.31% elongation) along horizontal direction. For
WADED Al-Mg-Sc-Zr-HT, the horizontal UTS increased
to 388 MPa, and the elongation could still be maintained
at 22.53%.
It is worth noting the mechanical properties along
vertical direction are inferior to horizontal direction of
all samples. A similar phenomenon is often found in
WADED aluminium alloy, which may related to the
pore distribution and the microstructure [1,53]. This
would be discussed in the following part combined
with the fractography. Moreover, all WADED samples
display typical Portevin–Le Chatelier (PLC) eects (serra-
tion in the ow curves). It was also reported in PBF-LB
counterparts in previous works [21]. The PLC eect is
dependent on dislocation density, grain size, strain rate
and temperature [54–56]. Compared with WADED Al-
Mg alloy, the tensile curves of WADED Al-Mg-Sc-Zr com-
ponent transform from smooth ow to serrated ow due
to its finer grains [27]. The reduction of PLC eects in
WADED Al-Mg-Sc-Zr-HT is related to the nano-size sec-
ondary precipitates forming in aging process which
trapped vacancies and reduced the mobility of the Mg
atoms [56]. Figure 10d presents the tensile properties
of WADED Al-Mg-Sc-Zr components in this work com-
pared with other Al-Mg-(Sc-Zr) components by cast,
powder-based DED and WADED methods in previous
works. The statistical results demonstrate that WADED
Al-Mg-Sc-Zr component in this work has an excellent
strength-to-ductility balance.
The side surface fractography of WADED components
is shown in Figure 11. For WADED Al-Mg component
(Figure 11a,b), the coarse β-Al
3
Mg
2
phase distributes
along GBs when Mg reaches 3.0 wt % which weakens
the GB binding force, resulting in crack tendency
under tensile loads [30]. Intergranular fracture is the
main fracture type, leaving continuous phases existing
at the fracture edge. The morphology of WADED Al-
Mg-Sc-Zr components varies along distinct tensile
Figure 10. (a) Engineering strain/stress curves, (b) true strain/stress curves (c) statistic results of WADED Al-Mg, Al-Mg-Sc-Zr and Al-
Mg-Sc-Zr-HT components and (d) comparison between tensile properties in this work and others in Ref. [8,20,21,30,31,33,43,46,52].
VIRTUAL AND PHYSICAL PROTOTYPING 11

directions (Figure 11c,d). In horizontal direction, several
precipitated particles distribute at the fracture edge.
These precipitates can be inferred as Al
3
(Sc, Zr) and
have a relatively high strength to prevent them from
plastically deforming and getting detached from the sur-
rounding matrix in the final tensile stage [9,57]. In verti-
cal direction, the fracture edge is along layer boundaries
[158]. The morphologies of WADED Al-Mg-Sc-Zr-HT
component show features that are similar to those of
the as-deposited counterpart (Figure 11e,f). It can be
concluded inter-layer region is a sensitive region for frac-
ture in vertical tensile direction. In particular, the layered
characteristics of microstructure tend to be more
obvious in Sc/Zr added components (Figure 5), which
might be one of the reasons for mechanical property
anisotropy.
Figure 12 presents the front surface fracture mor-
phologies of WADED components. The fractography of
WADED Al-Mg component in horizontal and vertical
directions exhibits an amount of dimples, indicating a
high-degree local plastic deformation and good ductility.
The dimple size in fractography of as-deposited and aged
WADED Al-Mg-Sc-Zr is similar with that in WADED Al-Mg.
The small particles can be observed in dimples are
Figure 11. Side surface fractography of horizontal and vertical direction of (a), (b) WADED Al-Mg, (c), (d) WADED Al-Mg-Sc-Zr and (e),
(f) WADED Al-Mg-Sc-Zr-HT components.
12 Y. ZHOU ET AL.

Al
3
(Sc, Zr) particles [31]. The components display charac-
teristics of ductile fracture. The coarse particles exist in the
fracture surface of Sc/Zr added sample, which corre-
sponds to the large-size particles in the interlayer region
(Figure 6). The large-size particles have a relatively high
strength, which prevents them from plastically deforming
[9]. Moreover, the layered pore distribution along vertical
direction also leads to mechanical properties lower than
Figure 12. Front surface fractography of (a), (b) WADED Al-Mg, (c), (d) WADED Al-Mg-Sc-Zr and (e), (f) WADED Al-Mg-Sc-Zr-HT
components.
VIRTUAL AND PHYSICAL PROTOTYPING 13

those in horizontal direction [41]. Therefore, the combi-
nation eects of coarse particles distributing along the
interlayer and layered pore distribution lead to the aniso-
tropy of mechanical properties.
3.4. Corrosion behaviour
Figure 13a shows the OCP evolution of each component.
The OCPs range from −0.9 V
SCE
to −0.75 V
SCE
in test sol-
ution. The OCP final value enhanced after Sc/Zr addition
and WADED Al-Mg-Sc-Zr-HT component has the highest
OCP value, indicating the lowest corrosion tendency
[59]. The potentiodynamic polarisation curves (Figure
13b) show a trace of active dissolution. The corrosion
current decreased after Sc/Zr addition and further
reduced after aging, indicating that electrochemical cor-
rosion behaviour was improved.
Figure 14a displays the Nyquist plots of each sample
which contain an impedance loop at high-frequency
range and another one at low-frequency range. It indi-
cates the electron transfer process is responsible for
the electrochemistry reaction at high frequency, and
the mass transfer step occurs at low frequency [60].
Figure 14a demonstrates the equivalent circuit model
used to explain the electrochemical process on alloy
surface, where R
s
stands for the solution resistance, Q
dl
is for the double layer constant phase element, R
ct
rep-
resents the charge transfer resistance, L and R
L
denote
the equivalent inductance and resistance of the cor-
rosion induction [61,62]. The occurrence of induction
indicates the low-resistance Al(OH)
3
film on alloy
surface cannot achieve a stable protective state, and
local corrosion would initiate in the subsequent stage.
It can be seen the impedance loop radius increases
Figure 13. (a) Open-circuit potential and (b) potentiodynamic polarisation curves of WADED Al-Mg, Al-Mg-Sc-Zr and Al-Mg-Sc-Zr-HT
samples in 3.5 wt.% NaCl solution.
Figure 14. (a) Nyquist and (b) Bode plots of WADED Al-Mg, WADED Al-Mg-Sc-Zr and WADED Al-Mg-Sc-Zr-HT components in 3.5 wt.%
NaCl solution.
14 Y. ZHOU ET AL.

with Sc/Zr adding and WADED Al-Mg-Sc-Zr-HT sample
has the largest radius. The corresponding Bode plot
also shows the WADED Al-Mg-Sc-Zr-HT component has
the largest modules at the low frequency, indicating a
better corrosion resistance.
Figure 15 displays the corrosion morphology of each
sample following a 24-h immersion in 3.5 wt % NaCl.
Based on EDS findings, the corrosion tends to initiate
from coarse Al
3
Mg
2
and Al
6
(Mn, Fe) phases in WADED
Al-Mg component. Both precipitates act as micro-
anode phases and tend to be preferentially dissolved
[36,63]. Due to the random phase distribution and the
vague layered microstructure characteristic, it exhibits
uniform corrosion features with large-size pittings. The
corrosion morphology of WADED Al-Mg-Sc-Zr sample
shows layered characteristics and pittings mainly distrib-
uted in inter-layer regions (Figure 15b
1
, 15b
2
). This can
be attributed to the galvanic corrosion caused by the
high potential dierence between Al
3
(Sc, Zr) and
matrix, and the gathering of particles results in pitting
aggregation in inter-layer region [36]. The corrosion
behaviour in inner-layer region is significantly improved
after aging (Figure 15c
1
, c
2
), which could be related to
the secondary Al
3
(Sc, Zr) forming in aging process [38].
The impacts of secondary Al
3
(Sc, Zr) on corrosion behav-
iour will be demonstrated in the discussion.
4. Discussion
4.1. Eects of Sc/Zr adding on microstructure
4.1.1. Eects of Sc/Zr adding on porosity
The previous study reported the pore-inhibition mech-
anism is complicated and related to the material proper-
ties, such as boiling temperature, melt viscosity and
thermal conductivity [23]. In this study, the pore-inhibit-
ing process after Sc/Zr adding is discussed from the
microalloying perspective.
According to Figure 3, due to the high thermal con-
ductivity and heat dissipation speed of aluminium
alloys, the hydrogen not discharged from the molten
pool in time would be retained to form pores in
WADED process, which is the main origination of pores
[64,65]. Pore development can be divided into four
Figure 15. The corrosion morphology and corresponding element mapping of (a
1
), (a
2
) WADED Al-Mg, (b
1
), (b
2
) WADED Al-Mg-Sc-Zr
and (c
1
), (c
2
) WADED Al-Mg-Sc-Zr-HT samples.
VIRTUAL AND PHYSICAL PROTOTYPING 15

stages: formation, growth, detachment and escape of
bubbles [66]. When T < 250°C the Al (solid) reacts with
H
2
O vapour to form Al(OH)
3
and H
2
which experience
‘adsorption-diusion-dissolution’ in molten pool; when
T > 400°C, the Al(OH)
3
decomposes to H
2
O; T > 600°C,
Al (liquid) react with H
2
O to form [H] that could directly
dissolve into Al (l), which is considered as one of the
main forms of hydrogen in Al alloys.
The classic Scheil module of thermal-calc software is
used to elaborate the eects of solidification path on
pore formation. As shown in Figure 16a, the Al-Mg-Sc-
Zr alloy system displays a more narrow solidification
range in contrast to Al-Mg alloy system. According to
the pore formation procedure described above, Sc/Zr
adding shortens liquid phase-existence time, which
restricts the [H] dissolving stage (above 600°C). The
stage between 400 and 600°C where the Al(OH)
3
decom-
poses to H
2
O is also restricted after Sc/Zr adding, which
further limits the supplement of [H]. Moreover, due to
the competitive relationship between bubble growth
and grain refinement, the grain refinement promoted
by Al
3
(Sc, Zr) particles which act as heterogeneous
nucleation sites promote can eectively inhibit bubble
growth ([67]; [21]). The combination eects of acceler-
ated solidification range and grain refinement lead to
low porosity in WADED Al-Mg-Sc-Zr.
4.1.2. Eects of Sc/Zr adding on layered
characteristics
The phase diagram in Figure 16b shows the primary
Al
3
(Zr, Sc) with D0
23
crystal structure precipitated early
in the first solidification stage and then transformed
into L1
2
crystal structure (Figure 8). The origination of
the primary Al
3
(Sc, Zr) in WADED Al-Mg-Sc-Zr component
can be divided into two types. The metallurgical process
in the molten pool produces the first one, and the other is
particles from the filler wire which generated during the
wire preparation process and exists with relatively larger
sizes [68]. The primary Al
3
(Sc, Zr) particles from the metal-
lurgical reaction preferentially precipitated along MPB
and inter-layer region [21], and those from the filler wire
gravitationally collected at the bottom of the molten
pool and layers. Both types of primary Al
3
(Sc, Zr) make
the grains at MPB and inter-layer regions significantly
refined [69]. In contrast, CEGs exist in inner-layer region
and molten pool interior due to the higher temperature
in these regions making the Sc/Zr supersaturated in the
matrix [27]. The coexistence of FEG and CEG results in
layered characteristics more obvious in WADED Al-Mg-
Sc-Zr component. The phase diagram shown in Figure
16b also indicates that β-Al
3
Mg
2
is inevitably generated
in Al-Mg-Sc-Zr alloy system. However, the refined grains
weaken the aggregation of β-Al
3
Mg
2
along GBs [14].
After aging treatment, the grain boundaries became
unclear and the layered characteristics were weakened.
4.1.3. Eects of Sc/Zr adding on grain
characteristics
In previous works, the columnar grains dominate in
microstructure of PBF-LB Al-Mg-Sc-Zr counterparts due
to solidification conditions favouring planar growth
[44]. However, the microstructure exhibits equiaxed
grains in this work. This could be attributed to the fol-
lowing factors. The columnar to equiaxed transition
(CET) can be expressed by classical solidification theory
[24,45]:
Gm
R=a������������
−4
p
N0
3ln(1 −
w
)
3
􏽳·1
1+m
􏼨 􏼩m
(1)
Figure 16. (a) The solidification path of the WADED Al-Mg and Al-Mg-Sc-Zr alloy predicted by the classic Scheil module and (b) the
phase diagram obtained from the thermal calc software using the bulk composition.
16 Y. ZHOU ET AL.

where G denotes the temperature gradient, R represents
the solidification velocity and N
0
stands for the nuclei
density. The a and m are material-dependent fitting con-
stants, and ϕ represents the fraction of equiaxed grains.
Since the slow cooling rate of deposited metal and
sparse energy distribution in depositing process,
WADED exhibits the characteristics of a low G/R ratio
[31], which is the favourable solidification condition for
the equiaxed grain formation [46]. In addition, the
Al
3
(Sc, Zr) particles acting as heterogeneous nucleation
sites enhanced N
0
, which also promoted equiaxed
grain formation.
4.2. Mechanism for strength enhancement
The contribution factors that improve YS in aluminium
alloy can be interpreted by applying the subsequent
formula ([25]; [22]):
D
s
y=D
s
gb +D
s
ss +D
s
d+D
s
p(2)
Where σ
y
is the YS evolution contains the solid-solution
strengthening (σ
ss
), dislocations hardening (σ
d
), grain
boundary strengthening (σ
gb
) and precipitation
strengthening (σ
p
). For WADED Al-Mg component, the
main strengthen eect originates from the σ
ss
which is
mainly contributed by the Mg content in the alloy.
Based on the EDS shown in Table 3, the Mg contents
in alloy matrix are similar. In addition, no work-harden-
ing treatments are conducted for components. There-
fore, the discussion of the strengthening mechanism is
mainly focused on σ
gb
and σ
p
.
4.2.1. Grain boundary strengthening
The Hall-Petch equation can be used to quantify the
grain refinement strengthening contribution:
s
gb =
s
0+kd−1
2(3)
where σ
0
represents the intrinsic resistance of the
lattice to dislocation motion, and k (0.17 MPa·m
0.5
) is
the constant standing for the relative strengthening
contribution from GBs [31]. The average grain size in
WADED Al-Mg is 84.34 μm. Considering the bimodal
grain structure after Sc/Zr adding, the contribution
factors of FEG and CEG were calculated respectively.
The FEG in WADED Al-Mg-Sc-Zr has an average size
of 11.44 μm and accounts for 23.18% while those of
CEG is 21.19 μm (76.82%). WADED Al-Mg-Sc-Zr-HT
presents the FEG with 11.42 μm (27.91%) and CEG
with 23.85 μm (72.09%). The strength increase by
grain refinement is approximately 21.19 MPa in
WADED Al-Mg-Sc-Zr and 18.80 MPa in WADED
Al-Mg-Sc-Zr-HT.
4.2.2. Precipitation strengthening
With Sc/Zr addition, the Al
3
(Sc, Zr) precipitates take the
main responsibility for σ
p
. The Al
3
(Sc, Zr) precipitates
have pinning eects on dislocation sliding and play a
crucial part in strength enhancement [16]. The Orowan
dislocation bypassing mechanism dominates since the
Al
3
(Sc, Zr) particles in WADED components have rela-
tively larger radii than those in PBF-LB counterparts
([44]; [16]). The impact of precipitation strengthening
can be calculated as [22,31]:
s
Orowan =M0.4Gb
pl
ln 2
r
b
􏼒 􏼓
������
1−v
√(4)
where M represents the Taylor factor (3.06); G denotes
the shear modulus (25.4 GPa); b stands for the Burgers
vector of Al (0.28 nm);
l
refers to the edge spacing
with of the primary Al
3
(Sc, Zr) particles (∼ 961 nm) and
secondary Al
3
(Sc, Zr) particles (∼275 nm); 
r is the mean
radius with about 110.3 nm for the primary Al
3
(Sc, Zr)
particles and about 24.5 nm for the secondary Al
3
(Sc,
Zr) precipitates; v is the Poisson ratio of the alloy
(0.27). With Sc/Zr adding, the Orowan strengthening
mechanism by primary Al
3
(Sc, Zr) particles contributes
approximately 22.41 MPa. After aging, the secondary
Al
3
(Sc, Zr) precipitates contribute about 60.48 MPa for
WADED Al-Mg-Sc-Zr-HT.
According to the results in Figure 10, the theoretical
YS enhancement is about 43.60 MPa, which is slightly
higher than the actual measured values (∼34.21 MPa).
This may be attributed to the obvious layered character-
istic with Sc/Zr adding and the fracture tends to occur at
the heterogeneous structure boundary (Figure 11). The
calculated strength increment for WADED Al-Mg-Sc-Zr
after aging is about 60.48 MPa which is higher than
the experimental results (∼39.93 MPa). This may be
caused by increased porosity after heat treatment
(Figure 3).
4.3. Corrosion mechanism
The dierent corrosion behaviour of each sample could
be attributed to grain characteristics, crystallographic
orientation and precipitates [70,71].
Grain size has a significant impact on corrosion
behaviour, and which has been reported that the cor-
rosion rate tends to decrease as particle size decreases,
[72]. In Figure 14a, the OCPs of WADED Al-Mg-Sc-Zr
and WADED Al-Mg-Sc-Zr-HT components rise rapidly in
testing process, which represents the generation of Al
(OH)
3
film on alloy surface [34]. Since the irregular
atom arrangement and ununiform element distribution
at GBs, the enhanced GB density made the alloy
VIRTUAL AND PHYSICAL PROTOTYPING 17

surface more active for reaction acceleration [72]. The
Sc/Zr added samples also show higher final OCP
values, representing the higher quality of Al(OH)
3
film.
Meanwhile, the grain structure also plays an important
role. The film on the heterogeneous structure transition
area is believed to have poor protective quality [62], and
corrosion preferentially initiates from these regions,
especially the inter-layer region where the transmission
of FEG to CEG occurred (Figure 16).
LAGB has been proven to exhibit better corrosion
resistance [73]. However, the GB density of WADED Al-
Mg-Sc-Zr (162.93 mm
−1
) and WADED Al-Mg-Sc-Zr-HT
(172.59 mm
−1
) samples are much higher than that of
the WADED Al-Mg sample (43.32 mm
−1
). Thus the total
LAGB length is more suitable for evaluating the cor-
rosion behaviour [71]. The total LAGB length in
WADED Al-Mg, Al-Mg-Sc-Zr and Al-Mg-Sc-Zr-HT com-
ponents are respectively 1.04, 1.79 and 3.50 mm, with
the densities of 4.06, 6.63 and 12.96 mm
−1
. The LAGB
density is enhanced with Sc/Zr addition and further
improved after aging treatment, which is consistent
with the electrochemical results.
The precipitates also have a significant eect on cor-
rosion behaviour. The precipitates mostly contain β-
Al
3
Mg
2
and Al
6
(Mn, Fe) in WADED Al-Mg alloy. The β-
Al
3
Mg
2
continuously distributing along the GBs is the
main factor for corrosion since it serves as the anodic
to the matrix and dissolves preferentially [63], leading
to corrosion propagation along GBs (Figure 15a
1
, a
2
).
Since the low Fe content in WADED Al-Mg and Al-Mg-
Sc-Zr, the fewer Al
6
(Mn, Fe) only acts as corrosion
feature site due to the anodic potential relative to the
Al matrix and is not the main factor for corrosion behav-
iour. The main factor for corrosion improvement in
WADED Al-Mg-Sc-Zr is the high GB density rather than
the lower Fe content, which weakens the aggregation
of β-Al
3
Mg
2
to improve corrosion along GBs. In addition,
the gathering of the Al
3
(Sc, Zr) particles results in the
regional distribution of pittings in inter-layer region.
For WADED Al-Mg-Sc-Zr-HT, the minor-size secondary
Al
3
(Sc, Zr) along GBs results in a narrow PFZ, which
cuts the corrosion channel and decreases the electro-
chemical dierence between GB and grain interior
[37,38,74]. Therefore, the corrosion behaviour in inner-
layer region where secondary Al
3
(Sc, Zr) preferentially
precipitated was improved.
5. Conclusions
In this study, WADED Al-Mg and WADED Al-Mg-Sc-Zr
components were fabricated, and aging heat treatment
(350°C-4 h aging) was further conducted. The eects of
Sc/Zr addition on porosity, microstructure, mechanical
characteristics, and corrosion behaviour were studied.
The following conclusions can be drawn:
(1) The total pores area fraction of WADED Al-Mg, Al-
Mg-Sc-Zr and Al-Mg-Sc-Zr-HT components were
0.208%, 0.005% and 0.045%, respectively. The
addition of Sc/Zr elements could eectively inhibit
pore formation.
(2) The equiaxed grain can be refined with Sc/Zr
addition (from 84.34 μm in WADED Al-Mg to 19.54
μm in WADED Al-Mg-Sc-Zr) and slightly grew with
heat treatment (21.26 μm). The layered microstruc-
ture characteristics tended to be more obvious
with Sc/Zr adding.
(3) The mechanical properties of WADED Al-Mg-Sc-Zr
alloy are improved with 360 MPa ultimate tensile
strength (UTS), 185 MPa yield strength (YS) and
23.31% elongation than that of WADED Al-Mg
(UTS: 330 MPa, YS: 155 MPa, elongation: 23.56%).
After heat treatment, the UTS further increased to
388 MPa while maintaining good ductility
(22.53%). The strengthening enhancing mechanisms
were mainly grain boundary and precipitation
strengthening.
(4) The corrosion behaviour was improved with Sc/Zr
addition, as the grain refinement weakened phase
aggregation along grain boundary. The Al
3
(Sc, Zr)
caused the matrix nearby dissolution. The corrosion
behaviour was further optimised after heat treat-
ment due to the narrow PFZ along grain boundary
cut the corrosion channel.
Disclosure statement
No potential conict of interest was reported by the author(s).
Funding
This work is supported by National Natural Science Foundation
of China (grant numbers 52105315, 52075022 and U20B2031),
State Key Lab of Advanced Welding and Joining, Harbin Insti-
tute of Technology (AWJ-24M05), Natural Science Foundation
of Beijing Municipality (grant numbers 3222013) and Funda-
mental Research Funds for the Central Universities (YWF-22-
L-607).
CRediT authorship contribution statement
Yubin Zhou: Data curation, Formal analysis, Investigation,
Methodology, Visualisation, Writing – original draft. Zewu Qi:
Investigation, Methodology, Resources, Supervision, Writing –
review & editing, Funding acquisition. Baoqiang Cong: Inves-
tigation, Resources, Supervision, Writing – review & editing.
Zihao Jiang: Visualisation, Writing – review & editing. Xinyi
18 Y. ZHOU ET AL.

Cai: Investigation, Writing – review & editing. Hongwei Li:
Investigation, Writing – review & editing. Jiaqi Zhang: Investi-
gation, Writing – review & editing. Chaofang Dong: Software,
Writing – review & editing. Xing He: Software, Writing – review
& editing. Sanbao Lin: Writing – review & editing. Xiaoyu Cai:
Writing – review & editing. Bojin Qi: Investigation, Project
administration, Resources, Writing – review & editing.
Data availability statement
Data will be available upon reasonable request.
References
[1] Cong BQ, Cai XY, Qi ZW, et al. The eects of ultrasonic fre-
quency pulsed arc on wire + arc additively manufactured
high strength aluminum alloys. Addit Manuf.
2022;51:102617. doi:10.1016/j.addma.2022.102617
[2] Cao QH, Zeng CY, Qi BJ, et al. Excellent isotropic mechan-
ical properties of directed energy deposited Mg-Gd-Y-Zr
alloys via establishing homogeneous equiaxed grains
embedded with dispersed nano-precipitation. Addit
Manuf. 2023;67:103498. doi:10.1016/j.addma.2023.103498
[3] Yi H, Jia L, Ding JL, et al. Achieving material diversity in
wire arc additive manufacturing: leaping from alloys to
composites via wire innovation. Int J Mach Tools Manuf.
2024;194:104103. doi:10.1016/j.ijmachtools.2023.104103
[4] Wu BT, Shao ZX, Shao DD, et al. Enhanced corrosion per-
formance in Ti-6Al-4 V alloy produced via wire-arc
directed energy deposition with magnetic arc oscillation.
Addit Manuf. 2023;66:103465. doi:10.1016/j.addma.2023.
103465
[5] Wang LW, Wu T, Wang DL, et al. A novel heterogeneous
multi-wire indirect arc directed energy deposition for
in-situ synthesis Al-Zn-Mg-Cu alloy: process, microstruc-
ture and mechanical properties. Addit Manuf.
2023;72:103639. doi:10.1016/j.addma.2023.103639
[6] Lee S, Utsunomiya A, Akamatsu H, et al. Inuence of scan-
dium and zirconium on grain stability and superplastic
ductilities in ultrafine-grained Al-Mg alloys. Acta Mater.
2002;50(3):553–564. doi:10.1016/S1359-6454(01)00368-8
[7] Horgar A, Fostervoll H, Nyhus B, et al. Additive manufac-
turing using WAAM with AA5183 wire. J Mater Process
Technol. 2018;259:68–74. doi:10.1016/j.jmatprotec.2018.
04.014
[8] Tonelli L, Laghi V, Palermo M, et al. Aa5083 (Al-Mg) plates
produced by wire-and-arc additive manufacturing: eect
of specimen orientation on microstructure and tensile
properties. Prog Addit Manuf. 2021;6(3):479–494. doi:10.
1007/s40964-021-00189-z
[9] Cordova L, Bor T, Smit MD, et al. Eects of powder reuse
on the microstructure and mechanical behaviour of Al-
Mg-Sc-Zr alloy processed by laser powder bed fusion
(LPBF). Addit Manuf. 2020;36:101625. doi:10.1016/j.
addma.2020.101625
[10] Wang ZH, Lin X, Kang N, et al. Directed energy deposition
additive manufacturing of a Sc/Zr-modified Al-Mg alloy:
eect of thermal history on microstructural evolution
and mechanical properties. Mater Sci Eng A.
2021;802:140606. doi:10.1016/j.msea.2020.140606
[11] Liu DH, Wu DJ, Wang RZ, et al. Formation mechanism of
Al-Zn-Mg-Cu alloy fabricated by laser-arc hybrid additive
manufacturing: microstructure evaluation and mechan-
ical properties. Addit Manuf. 2022;50:102554. doi:10.
1016/j.addma.2021.102554
[12] Wang Y, Guo W, Xie YK, et al. In-situ monitoring plume,
spattering behavior and revealing their relationship
with melt ow in laser powder bed fusion of nickel-
based superalloy. J Mater Sci Technol. 2023;177:44–58.
doi:10.1016/j.jmst.2023.07.068
[13] Yi H, Wang Q, Zhang WJ, et al. Wire-arc directed energy
deposited Mg-Al alloy assisted by ultrasonic vibration:
improving properties via controlling grain structures.
J Mater Process Technol. 2023;321:118134. doi:10.1016/
j.jmatprotec.2023.118134
[14] Baek M-S, Shah AW, Kim YK, et al. Microstructures, tensile
properties, and strengthening mechanisms of novel Al-
Mg alloys with high Mg content. J Alloys Compd.
2023;950:169866. doi:10.1016/j.jallcom.2023.169866
[15] Dong BL, Cai XY, Lin SB, et al. Microstructures and
mechanical properties of wire arc additive manufactured
5183-Al: inuences of deposition dimensions. CIRP J
Manuf Sci Technol. 2021;35:744–752. doi:10.1016/j.cirpj.
2021.08.014
[16] Li D, Wu YY, Geng ZW, et al. High strength Al−Mg−Sc−Zr
alloy with heterogeneous grain structure and intragranu-
lar precipitation produced by laser powder bed fusion.
J Alloys Compd. 2023;939:168722. doi:10.1016/j.jallcom.
2023.168722
[17] Dai W, Guo W, Xiao J, et al. Tailoring properties of directed
energy deposited Al-Mg alloy by balancing laser shock
peening and heat treatment. J Mater Sci Technol.
2024;203:78–96. doi:10.1016/j.jmst.2024.03.051
[18] Ren LL, Gu HM, Wang W, et al. Eect of Sc content on the
microstructure and properties of Al-Mg-Sc alloys depos-
ited by wire arc additive manufacturing. Met Mater Int.
2021;27(1):68–77. doi:10.1007/s12540-020-00763-0
[19] Ren LL, Gu HM, Wang W, et al. The microstructure and
properties of an Al-Mg-0.3Sc alloy deposited by wire arc
additive manufacturing. Metals (Basel). 2020;10(3):320.
doi:10.3390/met10030320
[20] Wang ZH, Lin X, Kang N, et al. Strength-ductility synergy
of selective laser melted Al-Mg-Sc-Zr alloy with a hetero-
geneous grain structure. Addit Manuf. 2020;34:101260.
doi:10.1016/j.addma.2020.101260
[21] Wang ZH, Lin X, Wang JF, et al. Remarkable strength-
impact toughness conict in high-strength Al-Mg-Sc-Zr
alloy fabricated via laser powder bed fusion additive
manufacturing. Addit Manuf. 2022;59:103093. doi:10.
1016/j.addma.2022.103093
[22] Ma KK, Wen HM, Hu T, et al. Mechanical behavior and
strengthening mechanisms in ultrafine grain precipi-
tation-strengthened aluminum alloy. Acta Mater.
2014;62:141–155. doi:10.1016/j.actamat.2013.09.042
[23] Zhou L, Hyer H, Park S, et al. Microstructure and mechan-
ical properties of Zr-modified aluminum alloy 5083 man-
ufactured by laser powder bed fusion. Addit Manuf.
2019;28:485–496. doi:10.1016/j.addma.2019.05.027
[24] Gaumann M, Bezencon C, Canalis P, et al. Single-crystal
laser deposition of superalloys: processing-microstruc-
ture maps. Acta Mater. 2001;49(6):1051–1062. doi:10.
1016/S1359-6454(00)00367-0
VIRTUAL AND PHYSICAL PROTOTYPING 19

[25] Kendig KL, Miracle DB. Strengthening mechanisms of an
Al-Mg-Sc-Zr alloy. Acta Mater. 2002;50(16):4165–4175.
doi:10.1016/S1359-6454(02)00258-6
[26] Yang KV, Shi YJ, Palm F, et al. Columnar to equiaxed tran-
sition in Al-Mg(-Sc)-Zr alloys produced by selective laser
melting. Scr Mater. 2018;145:113–117. doi:10.1016/j.
scriptamat.2017.10.021
[27] Qin ZH, Kang N, Mansori ME, et al. Anisotropic high cycle
fatigue property of Sc and Zr-modified Al-Mg alloy fabri-
cated by laser powder bed fusion. Addit Manuf.
2022;49:102514. doi:10.1016/j.addma.2021.102514
[28] Dong BL, Xia YH, Cai XY, et al. Addition of Sc in wire-based
directed energy deposition of Al-Mg-Zn-Cu alloy: microal-
loying to refine grains and improve mechanical proper-
ties. Addit Manuf. 2023;67:103494. doi:10.1016/j.addma.
2023.103494
[29] Guo YL, Han QF, Lu WJ, et al. Microstructure tuning
enables synergistic improvements in strength and duct-
ility of wire-arc additive manufactured commercial Al-
Zn-Mg-Cu alloys. Virtual Phys Prototyp. 2022;17(3):649–
661. doi:10.1080/17452759.2022.2048236
[30] Ren LL, Gu HM, Wang W, et al. Microstructure and prop-
erties of Al-6.0Mg-0.3Sc alloy deposited by double-wire
Arc additive manufacturing. 3D Print Addit Manuf.
2022;9(4):301–310. doi:10.1089/3dp.2020.0039
[31] Xia YH, Cai XY, Dong BL, et al. Wire arc additive manufac-
turing of Al-Mg-Sc alloy: An analysis of the eect of Sc on
microstructure and mechanical properties. Mater Charact.
2023;203:113116. doi:10.1016/j.matchar.2023.113116
[32] Dong BL, Cai XY, Lin SB, et al. Wire arc additive manufac-
turing of Al-Zn-Mg-Cu alloy: microstructures and mech-
anical properties. Addit Manuf. 2020;36:101447. doi:10.
1016/j.addma.2020.101447
[33] Dong BL, Cai XY, Xia YH, et al. Eects of interlayer tempera-
ture on the microstructures of wire arc additive manufac-
tured Al-Zn-Mg-Cu alloy: insights into texture responses
and dynamic precipitation behaviors. Addit Manuf.
2021;48:102453. doi:10.1016/j.addma.2021.102453
[34] Zhang H, Gu DD, Dai DH, et al. Inuence of heat treat-
ment on corrosion behavior of rare earth element Sc
modified Al-Mg alloy processed by selective laser
melting. Appl Surf Sci. 2020;509:145330. doi:10.1016/j.
apsusc.2020.145330
[35] Qiu YC, Yang XF, Li JX, et al. The inuence of Sc and Zr
additions on microstructure and corrosion behavior of
AA5182 alloy sheet. Corros Sci. 2022;199:110181. doi:10.
1016/j.corsci.2022.110181
[36] Zhang ZQ, Sun JE, Xia JY, et al. Anisotropic response in
corrosion behavior of laser powder bed fusion Al-Mn-
Mg-Sc-Zr alloy. Corros Sci. 2022;208:110634. doi:10.
1016/j.corsci.2022.110634
[37] Deng Y, Yin ZM, Zhao K, et al. Eects of Sc and Zr micro-
alloying additions and aging time at 120 °C on the cor-
rosion behaviour of an Al–Zn–Mg alloy. Corros Sci.
2012;65:288–298. doi:10.1016/j.corsci.2012.08.024
[38] Zhou ZX, Chen JQ, Wen F, et al. Optimization of heat
treatment for an Al-Mg-Sc-Mn-Zr alloy with ultrafine
grains manufactured by laser powder bed fusion. Mater
Charact. 2022;189:111977. doi:10.1016/j.matchar.2022.
111977
[39] Zeng CY, Cai XY, Qi ZW, et al. Eect of in-situ synthesized
TiC particles on microstructure and mechanical
properties of directed energy deposited AA2219 Al-Cu
alloy. J Alloys Compd. 2023;964:171331. doi:10.1016/j.
jallcom.2023.171331
[40] Yi H, Yang L, Jia L, et al. Porosity in wire-arc directed
energy deposition of aluminum alloys: formation mech-
anisms, inuencing factors and inhibition strategies.
Addit Manuf. 2024;84:104108. doi:10.1016/j.addma.2024.
104108
[41] Qi ZW, Qi BJ, Cong BQ, et al. Microstructure and mechan-
ical properties of wire plus arc additively manufactured
2024 aluminum alloy components: As-deposited and
post heat-treated. J Manuf Process. 2019;40:27–36.
[42] Zhao ZJ, Dong CF, Kong DC, et al. Inuence of pore
defects on the mechanical property and corrosion behav-
ior of SLM 18Ni300 maraging steel. Mater Charact.
2021;182:111514. doi:10.1016/j.matchar.2021.111514
[43] Ren Y, Dong P, Zeng Y, et al. Eect of heat treatment on
properties of Al-Mg-Sc-Zr alloy printed by selective laser
melting. Appl Surf Sci. 2022;574:151471. doi:10.1016/j.
apsusc.2021.151471
[44] Li RD, Wang MB, Li ZM, et al. Developing a high-strength
Al-Mg-Si-Sc-Zr alloy for selective laser melting: crack-inhi-
biting and multiple strengthening mechanisms.
Acta Mater. 2020;193:83–98. doi:10.1016/j.actamat.2020.
03.060
[45] Wang ZN, Lin X, Wang LL, et al. Microstructure evolution
and mechanical properties of the wire + arc additive man-
ufacturing Al-Cu alloy. Addit Manuf. 2021;47:102298.
doi:10.1016/j.addma.2021.102298
[46] Wang ZH, Lin X, Kang N, et al. Laser powder bed fusion of
high-strength Sc/Zr-modified Al–Mg alloy: phase selec-
tion, microstructural/mechanical heterogeneity, and
tensile deformation behavior. J Mater Sci Technol.
2021;95:40–56. doi:10.1016/j.jmst.2021.03.069
[47] Hyde KB, Norman AF, Prangnell PB. The eect of cooling
rate on the morphology of primary Al3Se intermetallic
particles in Al-Sc alloys. Acta Mater. 2001;49(8):1327–
1337. doi:10.1016/S1359-6454(01)00050-7
[48] Kong DC, Dong CF, Wei SL, et al. About metastable cellu-
lar structure in additively manufactured austenitic stain-
less steels. Addit Manuf. 2021;38:101804. doi:10.1016/j.
addma.2020.101804
[49] Zha M, Tian T, Jia HL, et al. Sc/Zr ratio-dependent mech-
anisms of strength evolution and microstructural
thermal stability of multi-scale hetero-structured Al–
Mg–Sc–Zr alloys. J Mater Sci Technol. 2023;140:67–78.
doi:10.1016/j.jmst.2022.09.009
[50] Jiang JY, Jiang F, Zhang MH, et al. Al3(Sc, Zr) precipitation
in deformed Al-Mg-Mn-Sc-Zr alloy: eect of annealing
temperature and dislocation density. J Alloys Compd.
2020;831:154856. doi:10.1016/j.jallcom.2020.154856
[51] Zhang ZQ, Sun JE, Wu JS, et al. Inuence of heat treat-
ment on corrosion behavior of Al-Mn-Mg-Sc-Zr alloy pro-
duced by laser powder bed fusion. J Mater Res Technol.
2023;23:4734–4746. doi:10.1016/j.jmrt.2023.02.029
[52] Geng HB, Li JL, Xiong JT, et al. Geometric limitation and
tensile properties of wire and arc additive manufacturing
5A06 aluminum alloy parts. J Mater Eng Perform. 2017;26
(2):621–629. doi:10.1007/s11665-016-2480-y
[53] Yuan T, Ren XL, Chen SJ, et al. Al-Zn-Mg-Cu alloy with
both high strength and high plasticity fabricated with
wire arc additive manufacturing. Sci Technol Weld
20 Y. ZHOU ET AL.

Joining. 2022;28(1):81–88. doi:10.1080/13621718.2022.
2117532
[54] Saad G, Fayek SA, Fawzy A, et al. Serrated ow and work
hardening characteristics of Al-5356 alloy. J Alloys
Compd. 2010;502(1):139–146. doi:10.1016/j.jallcom.2010.
04.119
[55] Hu ZH, Qi Y, Nie XJ, et al. The Portevin-Le Chatelier (PLC)
eect in an Al-Cu aluminum alloy fabricated by selective
laser melting. Mater Charact. 2021;178:111198. doi:10.
1016/j.matchar.2021.111198
[56] Bakare F, Schieren L, Rouxel B, et al. The impact of L1(2)
dispersoids and strain rate on the Portevin-Le-Chatelier
eect and mechanical properties of Al-Mg alloys. Mater
Sci Eng Struct Mater Prop Microstruct Process.
2021;811:141040. doi:10.1016/j.msea.2021.141040
[57] Van Hooreweder B, Moens D, Boonen R, et al. Analysis of
fracture toughness and crack propagation of Ti6Al4 V
produced by selective laser melting. Adv Eng Mater.
2012;14(1-2):92–97. doi:10.1002/adem.201100233
[58] Fu R, Lu WJ, Guo YL, et al. Achieving high strength-duct-
ility of Al-Zn-Mg-Cu alloys via hot-wire arc additive
manufacturing enabled by strengthening precipitates.
Addit Manuf. 2022;58:103042. doi:10.1016/j.addma.2022.
103042
[59] Kong DC, Wang L, Zhu GL, et al. Heat treatment eects on
the metastable microstructure, mechanical property and
corrosion behavior of Al-added CoCrFeMnNi alloys fabri-
cated by laser powder bed fusion. J Mater Sci Technol.
2023;138:171–182. doi:10.1016/j.jmst.2022.08.018
[60] Li N, Dong CF, Man C, et al. Insight into the localized strain
eect on micro-galvanic corrosion behavior in AA7075-T6
aluminum alloy. Corros Sci. 2021;180:109174. doi:10.
1016/j.corsci.2020.109174
[61] Ao M, Liu HM, Dong CF, et al. Degradation mechanism of
6063 aluminium matrix composite reinforced with TiC
and Al2O3 particles. J Alloys Compd. 2021;859:157838.
doi:10.1016/j.jallcom.2020.157838
[62] Ji YC, Dong CF, Kong DC, et al. Design materials based on
simulation results of silicon induced segregation at
AlSi10Mg interface fabricated by selective laser melting.
J Mater Sci Technol. 2020;46:145–155. doi:10.1016/j.
jmst.2020.01.037
[63] Davenport AJ, Yuan Y, Ambat R, et al. Intergranular cor-
rosion and stress corrosion cracking of sensitised AA5182.
Aluminium Alloys Pts 1 and 2. 2006;519-521:641–646.
[64] Da Silva CLM, Scotti A. The inuence of double pulse on
porosity formation in aluminum GMAW. J Mater Process
Technol. 2006;171(3):366–372. doi:10.1016/j.jmatprotec.
2005.07.008
[65] Boeira AP, Ferreira IL, Garcia A, et al. Alloy composition
and metal/mold heat transfer eciency aecting
inverse segregation and porosity of as-cast Al-Cu alloys.
Mater Des. 2009;30(6):2090–2098. doi:10.1016/j.matdes.
2008.08.032
[66] Hou XR, Zhao L, Ren SB, et al. A comparative study on
Al-Mg-Sc-Zr alloy fabricated by wire arc additive manu-
facturing with controlling interlayer temperature and
continuous printing: porosity, microstructure, and mech-
anical properties. J Mater Sci Technol. 2024;193:199–216.
doi:10.1016/j.jmst.2023.12.062
[67] Cong BQ, Ouyang RJ, Qi BJ, et al. Inuence of cold metal
transfer process and Its heat input on weld bead geome-
try and porosity of aluminum-copper alloy welds. Rare
Met Mater Eng. 2016;45(3):606–611. doi:10.1016/S1875-
5372(16)30080-7
[68] Guo XP, Li HJ, Xue P, et al. Microstructure and mechanical
properties of 600 MPa grade ultra-high strength alumi-
num alloy fabricated by wire-arc additive manufacturing.
J Mater Sci Technol. 2023;149:56–66. doi:10.1016/j.jmst.
2022.12.007
[69] Shen JJ, Chen B, Wan J, et al. Eect of annealing on micro-
structure and mechanical properties of an Al-Mg-Sc-Zr
alloy. Mater Sci Eng Struct Mater Prop Microstruct
Process. 2022;838:142821. doi:10.1016/j.msea.2022.142821
[70] Ralston KD, Fabijanic D, Birbilis N. Eect of grain size on
corrosion of high purity aluminium. Electrochim Acta.
2011;56(4):1729–1736. doi:10.1016/j.electacta.2010.09.
023
[71] Gu DD, Zhang H, Dai DH, et al. Anisotropic corrosion
behavior of Sc and Zr modified Al-Mg alloy produced
by selective laser melting. Corros Sci. 2020;170:108657.
doi:10.1016/j.corsci.2020.108657
[72] Ralston KD, Birbilis N, Davies CHJ. Revealing the relation-
ship between grain size and corrosion rate of metals. Scr
Mater. 2010;63(12):1201–1204. doi:10.1016/j.scriptamat.
2010.08.035
[73] Zhang RF, Qiu Y, Qi YS, et al. A closer inspection of a grain
boundary immune to intergranular corrosion in a sensi-
tised Al-Mg alloy. Corros Sci. 2018;133:1–5. doi:10.1016/
j.corsci.2018.01.009
[74] Ahmad Z, Ul-Hamid A, Bj AA, et al. The corrosion behavior
of scandium alloyed Al 5052 in neutral sodium chloride
solution. Corros Sci. 2001;43(7):1227–1243. doi:10.1016/
S0010-938X(00)00147-5
VIRTUAL AND PHYSICAL PROTOTYPING 21