Scripta Materialia 134 (2017) 6–10
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Scripta Materialia
journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Effect of Zirconium addition on crack, microstructure and mechanical
behavior of selective laser melted Al-Cu-Mg alloy
Hu Zhang, Haihong Zhu ⁎, Xiaojia Nie, Jie Yin, Zhiheng Hu, Xiaoyan Zeng
Wuhan National Laboratory for Optoelectronics, Huazhong University of Science and Technology, Wuhan, Hubei 430074, PR China
a r t i c l e
i n f o
Article history:
Received 3 February 2017
Received in revised form 20 February 2017
Accepted 20 February 2017
Available online 15 March 2017
Keywords:
Selective laser melting
Aluminum alloys
Zirconium
Microstructure
Mechanical properties
a b s t r a c t
Selective laser melting (SLM) was used to fabricate Al-Cu-Mg alloy and Zirconium-modified Al-Cu-Mg alloy components. The hot-cracking phenomena during the SLM process was significantly reduced by the grain refining effect caused by the addition of Zr. Compared to the Al-Cu-Mg part, the Zr-modified Al-Cu-Mg part with the
ultrafine grain exhibits increased yield strength (446 ± 4.3 MPa) and ultimate tensile strength (451 ±
3.6 MPa). These findings provide a basis for innovative alloy design for SLM.
© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Selective laser melting (SLM), one of the emerging additive
manufacturing (AM) technologies, has several advantages by allowing
to produce structures with complex geometry, reducing materials
usage and cost, and achieving high cooling rate (105 K/s) during processing [1]. SLM of aluminum (Al) has been heavily studied, as Al alloys
are widely used in the aerospace and automotive industries due to its
low density and the high thermal conductivity [2]. However, processing
Al alloys with SLM can be challenging because of its poor flowability and
high reflectivity along with high thermal conductivity, which can lead to
high levels of porosity and cracking [3].
In recent years, a few Al-Si based cast alloys have been processed
successfully and crack free by SLM [2,4–6]. The high Si-content, ranging
from 7 to 20%, offers good fluidity and reduced solidification shrinkage,
thus facilitating the fabrication of sound components. A few works have
been devoted to SLM of high strength wrought Al alloys. However, due
to their poor weldability, the application of SLM in these alloys is significantly limited due to the formation of hot cracking. For instance, the
SLM of Al 6061 and Al 7075 have shown unacceptable levels of cracking
as discussed by Fucher et al. [7] and Kaufmann et al. [8]. Recently, Montero Sistiaga et al. [9] proved that the addition of Silicon to Al 7075 powder eliminates the cracking during SLM, but no tensile properties data
was reported.
For the SLM of Al-Cu-Mg (2xxx) alloys, our previous work has
showed that dense and crack free Al-4.24Cu-1.97Mg-0.56Mn parts
with good tensile properties can be successfully fabricated [10]. However, the process still suffer from its high susceptibility to solidification
⁎ Corresponding author.
E-mail address: zhuhh@hust.edu.cn (H. Zhu).
http://dx.doi.org/10.1016/j.scriptamat.2017.02.036
1359-6462/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
cracking when processed with high laser scanning speed. Here, for the
first time, we demonstrate that the addition of Zr to the SLM of Al-CuMg powder can reduce and eliminate the cracks that are formed during
SLM. Accordingly, the influence of Zr addition on the microstructural
and mechanical properties of SLM fabricated Al-Cu-Mg alloy are presented and discussed.
99.0% purity Zr powder with a non-spherical shape and a mean particle size of 8.8 μm, and gas atomized Al-Cu-Mg powder with spherical
morphology and an average particle diameter of 36 μm was used in
this experiment. The nominal chemical composition (wt%) of the powder was 4.24 Cu, 1.97 Mg, 0.56 Mn and aluminum (balance). The Zr and
Al-Cu-Mg powder mixture with 2% wt% Zr particles were blended by
mechanical mixing in an argon atmosphere for 4 h.
Details regarding SLM equipment and procedure have been addressed in our previous publication [11]. Based on a series of preliminary experiments, the SLM parameters were adopted as follows: laser
power (P) 200 W, hatch spacing (h) 90 μm, layer thickness (t) 40 μm
and a zigzag scan pattern with 90° rotation between the adjacent layers.
Cubic samples (5 × 5 × 8 mm3) were produced with various scanning
speed (v) of 5, 10, 15 and 20 m/min to assess the porosity and crack
density.
The relative density of SLM fabricated samples was evaluated by
image analysing of 8 images using Image-Pro Plus software. Metallographic samples were prepared according to standard procedures, and
etched with Keller's reagent. The microstructure was characterized
using a Nikon EPIPHOT 300 optical microscope (OM) and a FEI Nova
Nano SEM 450 scanning electron microscope (SEM). The crystallographic orientation of the samples was investigated using electron
backscattered diffraction (EBSD), which was performed on an FEI Sirion
H. Zhang et al. / Scripta Materialia 134 (2017) 6–10
200 instrument equipped with a TSL/EDAX system. Phase analysis was
performed by X-ray diffraction (XRD) using a D3290 PANalytical X'pert
PRO with Cu Kα radiation. Tensile testing specimens were designed according to ASTM B557M-10 standard. The tests were performed with a
crosshead velocity of 2 mm/min. Three samples were tested for each
process condition.
Fig. 1 shows the polished cross section of the SLM-fabricated Al-CuMg and Zr/Al-Cu-Mg cubes. In Fig. 1a–d, the defects of the Al-Cu-Mg
samples are significantly affected by the scanning speed. When the
scanning speed was reduced from 20 m/min to 5 m/min, the irregular
shape pores and cracks were gradually eliminated. Specimens with
scanning speed of 5 m/min have the highest densities (up to 99.8%)
(Fig. 1a). On the other hand, cracks are not observed in all the Zr/AlCu-Mg samples, as displayed in Fig. 1e–h. Apparently, the addition of
Zr prevents the formation of cracks. The morphology of pores changes
from an irregular (Fig. 1h) to a spherical shape (Fig. 1e) with decreasing
scanning speed. In this case, the highest densities of 99.8% for Zr/Al-CuMg mixture was achieved at scanning speed of 15 m/min (Fig. 1g). The
optimal scanning speed for highest relative density increases from
5 m/min to 15 m/min with addition of Zr.
Fig. 2 illustrates microstructures of the Al-Cu-Mg and Zr/Al-Cu-Mg
compositions fabricated at various scanning speed. Fig. 2a and b show
SEM micrographs of the cross sections of Al-Cu-Mg alloy. The microstructure is not uniform throughout the material, but it displays the typical hatch overlapped laser tracks of SLM processing. These arc shaped
melt pools are the result of the heat flow as it flows radially away
from the center. Fig. 2b shows the higher magnification image of the
hatch overlapped region, and reveals cellular solidification mode, with
the phase contrast of grey for supersaturated primary α-Al phase and
white for Cu-rich interdendritic eutectic [12]. Dendritic coarsening is occurring at the melt pool boundaries, where the metal solidification is affected by the accumulation of thermal effects from additional re-melted
layers [13].
With the addition of Zr particles, pronounced structural modification
of Zr/Al-Cu-Mg mixture can be observed (Fig. 2c–f). The scanning speed
significantly influenced the microstructure evolution of Zr/Al-Cu-Mg
compositions. At the scanning speed of 5 m/min (Fig. 2c), which produce 99.8% density of Al-Cu-Mg sample, a uniform ultrafine equiaxed
grain is observed and the melt pools can't be distinguished. At higher
scanning speed of 15 m/min, the microstructure consists of fine grained
microstructure and coarser grains growing along the temperature gradient. From higher magnification images (Fig. 2d and f), some highly homogeneous precipitates were observed. These precipitates was
identified by XRD as Al3Zr particles (Fig. 2g). The repetitive melting
and re-solidification processed occurring during SLM influences the formation of Al3Zr precipitates [14]. These particles are very effective heterogeneous nuclei due to the structural similarity between FCC
7
aluminum and Al3Zr and the small lattice parameter misfit, which are
reported as 4.049 Å, and 4.007 Å, respectively. The presence of Al3Zr
acts as ideal nucleus of α-Al during solidification and results in a significant grain refinement, which has been reported for laser welding Zradded aluminum alloys [15].
Furthermore, the EBSD maps of the Al-Cu-Mg and Zr/Al-Cu-Mg compositions fabricated at different scanning speed are shown in Fig. 3. For
the Al-Cu-Mg alloy, the inverse pole figure (IPF) in Fig. 3a shows that the
microstructure consists mostly of larger columnar grains. There is no
obvious morphology of melt pool, even though the melt pools are clearly visible in SEM images (Fig. 2a). This indicates the solidification occurred with epitaxial growth from the melt pool boundary. In Fig. 3b,
with the addition of Zr, nanostructured and ultrafine equiaxed grains
were formed and the average grain size was 0.8 μm. In Fig. 3c, with
higher scanning speed, the melt pool shapes are more distinguished,
where the melt pool boundaries are decorated by finer grains with average size of 160 nm. The elongated grains with average size of 1.16 μm
are directed along the maximum thermal gradient, towards the center
of the melt pool. The reason for the “biomodal” grain distribution is because the grain boundaries are pinned by small Al3Zr precipitates [16].
Such grain boundary pinning helps to stabilize against grain growth
during the heat impact occurring by the deposition of subsequent
layers. Similar results and explanation were reported by Spierings et
al. [17].
According to Fig. 3d and e, with the addition of Zr, the boundary misorientation distribution shift from low angles to high angles (HABs, misorientation angle larger than 15°). The distribution of the
misorientation angles of SLM-fabricated Zr/Al-Cu-Mg compositions are
randomly oriented (Fig. 3e), as the distribution follows nicely the ideal
distribution for random orientation [18]. This indicates the SLM-fabricated Zr/Al-Cu-Mg compositions possess low anisotropic properties.
Tensile tests were performed on the Al-Cu-Mg and Zr/Al-Cu-Mg
parts fabricated with optimal processing parameters (Fig. 4a). For the
Zr/Al-Cu-Mg compositions part, it is an interesting finding that the
yield strength ratio approaches 1. The ultimate tensile strength (UTS)
about 451 ± 3.6 MPa and yield strength about 446 ± 4.3 MPa could
be observed. This result is significant higher than the Al-Cu-Mg specimen, whose UTS is 393 ± 20 MPa and yield strength is 253 ±
9.8 MPa. However, the degradation in elongation is distinct, which decreased from 6 ± 1.6% to 2.67 ± 1.1%. Fig. 4b–c illustrates the corresponding fracture surfaces of SLM processed Al-Cu-Mg and Zr/Al-CuMg compositions part. For the Al-Cu-Mg part, the fracture surfaces generally exhibited ductile-type failure, with the presence of the columnar
grains and dimples like features with a relatively large size of 1–2 μm
(Fig. 4b). As for the Zr/Al-Cu-Mg compositions part, refined shallow
dimples (~0.4–0.9 μm) were present in Fig. 4c, indicating failure in a relatively brittle mode. Moreover, the ultrafine precipitated phase was
Fig. 1. Polished cross sections of SLM-fabricated Al-Cu-Mg and Zr/Al-Cu-Mg samples at different scanning speed.
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H. Zhang et al. / Scripta Materialia 134 (2017) 6–10
Fig. 2. SEM images showing typical microstructures of cross sections of Al-Cu-Mg sample fabricated at v = 5 m/min (a, b) and Zr/Al-Cu-Mg sample fabricated at v = 5 m/min (c, d) and v =
15 m/min (e, f), respectively. (g) XRD spectra for Al-Cu-Mg and Zr/Al-Cu-Mg samples fabricated at different scanning speed.
uniformly incorporated in these dimples. From Fig.4c, spherical pores
also can be observed and the porosity content (surface area) is higher
than the results measured by image analysis. Upon straining, the facture
cracks appear and preference to propagate thought these spherical porosity, which results in lower ductility [19].
By combining the XRD, FESEM and EBSD analysis, the influence of Zr
addition on cracking formation can be addressed. The Zr addition can
reduce the crack sensitivity due to the formation of more low melting
temperature phase that could backfill the cracks initiate in the mushy
zone during the last stage of solidification, as observed by Ghaini et al.
[20]. Additional to this effect, the presence of Zr induces the formation
of ultrafine equiaxed grains. These fine grains increase the total grain
boundary surface area within a given volume, which can toughen the
matrix and avoid intergranular cracking [9].
The enhanced strength obtained in the SLM-fabricated Zr/Al-Cu-Mg
compositions (vs. Al-Cu-Mg) can be rationalized based on the
H. Zhang et al. / Scripta Materialia 134 (2017) 6–10
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Fig. 3. EBSD inverse pole figure (IPF) maps of Al-Cu-Mg fabricated at v = 5 m/min (a) and Zr/Al-Cu-Mg sample fabricated at v = 5 m/min (b) and v = 15 m/min (c), respectively. (d, e)
misorientation angle distribution.
Fig. 4. (a) Stress-Strain curve of the Al-Cu-Mg and Zr/Al-Cu-Mg samples fabricated at different scanning speed. SEM images showing typical morphologies of fracture surfaces of the Al-CuMg sample fabricated at v = 5 m/min (b) and Zr/Al-Cu-Mg sample fabricated at v = 15 m/min (c), respectively. The inset shows magnified of the areas enclosed in green squares.
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H. Zhang et al. / Scripta Materialia 134 (2017) 6–10
microstructure developed in the materials, namely, formation of ultrafine grains and fine second-phase precipitates. Indeed, according to
the Hall-Petch law:
k
σ y ¼ σ 0 þ pffiffiffi
d
ð1Þ
where σy is the increase in yield strength due to fine grain strengthening, σ0 is the friction stress (20 MPa), k is a constant (0.17 MPa·m1/2)
[21] and d is the grain size. According to EBSD data, the grain size hardening σy equals to 136 MPa. This calculation is lower than the experimental result (193 MPa). This indicates that the precipitate
strengthening also contributes for improving tensile strength. During
the SLM process, Al3Zr particles are precipitated. Therefore, due to
their pinning effect on grain boundaries, the movement of dislocations
and grain boundaries will be restrained, leading to the higher strength
of Zr/Al-Cu-Mg specimen.
In summary, the addition of 2 wt% Zr to base Al-Cu-Mg powder significantly reduces and eventually eliminates the cracks formation during SLM. The addition of Zr element can promote the formation of
Al3Zr precipitates and grain refinement, which prevents the formation
and propagation of cracks. Unlike the Al-Cu-Mg part featured with columnar grain, the Zr/Al-Cu-Mg compositions displays ultrafine equiaxed
grain. With the addition of Zr, the yield strength increases from 253 ±
9.8 MPa to 446 ± 4.3 MPa, and ultimate strengths rise from 389 ±
20 MPa to 451 ± 3.6 MPa.
This work was supported by the National Natural Science Foundation of China (Grant No. 61475056). The authors would also like to
thank the Analytical and Testing Center of HUST for the XRD, SEM and
EBSD analysis.
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