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Influence of LPBF-Surface Characteristics on Fatigue Properties of Scalmalloy®


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Laser powder bed fusion (LPBF) has indisputable advantages when designing new components with complex geometries due to toolless manufacturing and the ability to manufacture components with undercuts. However, fatigue properties rely heavily on the surface condition. In this work, in-process surface parameters (three differing contour parameter sets) and post-process surface treatments, namely turning and shot peening, are varied to investigate the influence of each treatment on the resulting fatigue properties of LPBF-manufactured specimens of the aluminium–magnesium–scandium alloy Scalmalloy®. Therefore, metallographic analysis and surface roughness measurements, as well as residual stress measurements, computer tomography measurements, SEM-analyses, tensile and fatigue tests, along with fracture surface analysis, were performed. Despite the fact that newly developed in-process contour parameters are able to reduce the surface roughness significantly, only a minor improvement in fatigue properties could be observed: Crack initiation is caused by sharp, microscopic notches at the surface in combination with high tensile residual stresses at the surface, which are present on all in-process contour parameter specimens. Specimens using contour parameters with high line energy show keyhole pores localized in the subsurface area, which have no effect on crack initiation. Contours with low line energy have a slightly positive effect on fatigue strength because less pores can be found at the surface and subsurface area, which even more greatly promotes an early crack initiation. The post-process parameter sets, turning and shot peening, both improve fatigue behaviour significantly: Turned specimens show lowest surface roughness, while, for shot peened specimens, the tensile residual stresses of the surface radially shifted from the surface towards the centre of the specimens, which counteracts the crack initiation at the surface.
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Influence of LPBF-Surface Characteristics on Fatigue Properties
of Scalmalloy®
Jens Musekamp 1, * , Thorsten Reiber 2, Holger Claus Hoche 1, Matthias Oechsner 1, Matthias Weigold 2and
Eberhard Abele 2
Citation: Musekamp, J.; Reiber, T.;
Hoche, H.C.; Oechsner, M.; Weigold,
M.; Abele, E. Influence of LPBF-
Surface Characteristics on Fatigue
Properties of Scalmalloy®.Metals
2021,11, 1961.
Academic Editor: Marco Mandolini
Received: 27 October 2021
Accepted: 29 November 2021
Published: 6 December 2021
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Attribution (CC BY) license (https://
1Center for Engineering Materials, State Materials Testing Institute Darmstadt (MPA) Chair,
Institute für Materials Technology (IfW), Technical University of Darmstadt, Grafenstraße 2, 64293 Darmstadt,
Germany; (H.C.H.); (M.O.)
2Institute of Production Management, Technology and Machine Tools (PTW), Technical University of
Darmstadt, Otto-Berndt-Straße 2, 64287 Darmstadt, Germany; (T.R.); (M.W.); (E.A.)
*Correspondence:; Tel.: +49-6151-16-24890
Laser powder bed fusion (LPBF) has indisputable advantages when designing new com-
ponents with complex geometries due to toolless manufacturing and the ability to manufacture
components with undercuts. However, fatigue properties rely heavily on the surface condition. In
this work, in-process surface parameters (three differing contour parameter sets) and post-process
surface treatments, namely turning and shot peening, are varied to investigate the influence of each
treatment on the resulting fatigue properties of LPBF-manufactured specimens of the aluminium–
magnesium–scandium alloy Scalmalloy
. Therefore, metallographic analysis and surface roughness
measurements, as well as residual stress measurements, computer tomography measurements, SEM-
analyses, tensile and fatigue tests, along with fracture surface analysis, were performed. Despite the
fact that newly developed in-process contour parameters are able to reduce the surface roughness
significantly, only a minor improvement in fatigue properties could be observed: Crack initiation is
caused by sharp, microscopic notches at the surface in combination with high tensile residual stresses
at the surface, which are present on all in-process contour parameter specimens. Specimens using
contour parameters with high line energy show keyhole pores localized in the subsurface area, which
have no effect on crack initiation. Contours with low line energy have a slightly positive effect on
fatigue strength because less pores can be found at the surface and subsurface area, which even more
greatly promotes an early crack initiation. The post-process parameter sets, turning and shot peening,
both improve fatigue behaviour significantly: Turned specimens show lowest surface roughness,
while, for shot peened specimens, the tensile residual stresses of the surface radially shifted from the
surface towards the centre of the specimens, which counteracts the crack initiation at the surface.
fatigue; aluminium alloys; stress measurements; post processing; laser powder bed fusion
1. Introduction
Laser powder bed fusion (LPBF) belongs to the group of additive manufacturing
(AM) processes in which a three-dimensional component is produced by joining volume
elements and building up thin surfaces layer by layer [
]. By manufacturing with LPBF,
complex geometries such as thin walls or internal grid structures which have an almost
100% density can be realized, allowing a component to be optimized both in terms of weight
and function [
]. This new freedom in design benefits a wide range of the manufacturing
sector to improve and develop new products e.g., optimized valves and pumps [
], new
tools, and optimized turbine blades or automotive applications [
], respectively. Especially
for aluminium alloys, there is a high need for sustainability, efficiency increase, and cost
reduction in the automotive, aerospace, and aircraft industries [
]. Moreover, there is a
Metals 2021,11, 1961.
Metals 2021,11, 1961 2 of 19
need to find alternatives for titanium alloys in terms of the strength–width ratio [
], which
contributes much, along with the design rules of LPBF.
To fulfil the requirements of high strength aluminium alloys, Scalmalloy
was devel-
oped especially for manufacturing with the LPBF-process [
]. The aluminium–magnesium–
scandium alloy has high tensile strength along with high elongation at break compared
to other high strength aluminium alloys [
]. Al
) and AlMgO phases act as
effective nucleating agents during solidification of the melt [
] and are responsible for the
high mechanical properties obtained with additional precipitation of these phases during
precipitation hardening heat treatment (325
C, 4 h) [
]. The infill parameters were
optimized with regard to the density to achieve up to 99.8% dense specimens [1619].
Apart from a high density, the surface zone, which has high roughness in the LPBF-
manufactured state [
], and in which an increased pore density may be present [
exerts a significant influence on the fatigue behaviour [
]. The roughness is caused by
three characteristic effects: the irregular geometric solidification of the melt pool due to
the process dynamics [
], powder particles adhering to the surface [
], and the
staircase effect [
]. The latter only comes into play with inclined surfaces. On the one
hand, irregular melt tracks occur due to the balling effect at low line energy E
, which is
determined by the ratio of laser power P and scanning speed v. This effect is caused by the
minimization of the surface energy of the melt, which is driven by its surface tension [
On the other hand, irregular melt tracks can result in an increased line energy owing to the
highly dynamic Marangoni convection [
]. It must be taken into account when only using
an infill parameter set that the outline of vertical and inclined surfaces is defined by the
melt track fronts of numerous individual melt tracks depending on the layer rotation angle.
Roughness values for vertical walls in the case of aluminium alloys are at 8
25 µm [3133].
Murakami et al. [
] states that on AM-materials in general, the crack initiation more
often starts at the surface due to surface roughness or defects at the surface, than at internal
The high roughness and defects in the surface zone are considered to act as notches
and thus contribute to the reduction and scattering of the fatigue strength [
]. Therefore,
the surface zone of safety-relevant components is usually removed by a turning post-
processing [
], which results in a considerable expense of time and costs. Wagener
et al. [
] showed that on LPBF-AlSi10Mg especially, the near-surface porosity, which
is brought in by contour scans in overhanging regions, lowers the fatigue properties
significantly. Moreover, notches, which are located near the surface (subsurface porosity)
exhibit a high influence on the fatigue properties.
The influence of pores on fatigue properties can be simulated using the
of Murakami [
]. In [
], a comparison of fatigue properties of forged, castes, wrought
and LPBF-manufactured material regarding the fatigue properties was summarized and
simulated with the
area-concept. The large scatter of the fatigue properties for machined
materials could be significantly reduced if the data are correlated to the surface of the defect
size at the failure origin. The fatigue strength for LPBF-manufactured AlSi10Mg and LPBF-
manufactured Ti6Al4V-specimens can even exceed the fatigue strength of conventional
manufactured specimens if they are machined and stress-relieved. Subsurface porosity,
which can result due to a poor contour scan, can reduce the fatigue properties signifi-
cantly. Moreover, for additively manufactured 316L, the fatigue properties are somewhat
equivalent to those of wrought materials [38].
Maskery et al. [
] investigated the pore size distribution of LPBF-AlSi10Mg and
found out that it follows a Weibull distribution, which can be used for lifecycle-modelling,
determining the crack growth rates.
In [
], it was stated that subsurface porosity and surface roughness decrease the
fatigue properties of AlSi10Mg significantly. Different post-process treatments can reduce
these lowering factors: a vibratory finishing increases the fatigue properties due to lowering
the surface roughness. Sand blasting increases the fatigue properties even more due to the
Metals 2021,11, 1961 3 of 19
induction of compressive residual stresses, although the surface roughness is higher. The
highest increase in fatigue strength was achieved by machining and polishing.
The influence of machining on fatigue properties of LPBF-AlSi10Mg was shown in [
where machining significantly increases the fatigue limit. This was mainly explained with
the decrease in surface roughness, because satellites (powder at the surface) and balling
(agglomerates of molten material) can be identified on as-built specimens, which could
act as crack initiation areas. After machining, some porosity is still present at the surface,
which can cause crack initiation.
Brandão [
] improved the fatigue properties of AlSi10Mg by jet blasting after manu-
facturing, which increases the fatigue properties in the same way as a vibratory finishing.
In [
], the fatigue properties of heat treated and turned Scalmalloy
(laser power =
370 W, hatch distance = 100
m, 30
m layer thickness and 1600 mm/s scanning speed)
have been investigated at a stress ratio of R = 0.1 up to N = 3
cycles. The fatigue
properties show high scatter and a significant correlation with the building orientation:
Specimens in building direction with a hot isostatic pressing (HIP)-treatment exhibit the
lowest fatigue limit of
= 160 MPa (
= 88 MPa), whereas specimens in a diagonal
direction show a fatigue limit up to
= 240 MPa (
= 132 MPa). The crack initiation of
was dependent on local microstructural features around the crack tip. The
main driver for the poor fatigue crack growth behaviour is explained with the fully coherent
) precipitations: In Scalmalloy
, neither grain boundaries nor precipitations
are able to prevent crack growth propagation under cyclic loads.
Compared to the post-processing methods described above, in-process methods such
as contour scans can be performed directly during the LPBF-process in order to save
additional time and costs. When using contour scans, the material is melted in one or
more vectors parallel to the contour of the component. Exposure of the contour can take
place before (pre-contour) or after the infill (post-contour). The process parameters of the
contour scans can be optimized independently of the infill parameters in terms of low
roughness. To achieve a low roughness, increased energy input can be applied. This leads
to more uniform melt tracks and the peaks and valleys between the layers can be decreased
by re-melting [
]. Using two contour scans with a high line energy E
= 1.17 J/mm,
a roughness of Ra = 6.1
m on vertical walls in Scalmalloy
specimens could be achieved.
However, some pores were still present despite the contour scans [
]. The reason for the
occurrence of pores when exposing using high line energies can be the transition from
conduction-mode to keyhole-mode welding [
], which is described in detail in [
]. In
order to avoid the formation of keyhole pores, welding in keyhole mode can be prevented
by reduction of the line energy.
In this work, the influence of surface properties such as roughness and residual stress,
the intrinsic pores, and the subsurface porosity on the fatigue properties for the LPBF-
manufactured Al–Mg–Sc alloy Scalmalloy
are investigated to extend the knowledge of
these influencing factors. This helps to optimise the fatigue strength when designing for
subsequent operational use. Therefore, three contour scanning strategies derived from a
preliminary study of the authors [
] are compared to the as-built surface state without
contour parameters, the shot peened-state, and the turned-state. The properties are exam-
ined using roughness measurements, residual stress measurements, optical Microscopy
(OM), fracture surface analyses by means of Scanning Electron Microscopy (SEM), tensile
tests, and fatigue tests.
2. Materials and Methods
Commercially available Scalmalloy
powder material from Heraeus Additive Man-
ufacturing GmbH with a mean particle size of 20 to 65
m (D10 to D90) was used. The
chemical composition of an exemplary additively manufactured specimen measured by
spark-emission spectroscopy is shown in Table 1.
Metals 2021,11, 1961 4 of 19
Table 1.
Composition of the manufactured specimens measured via spark-emission spectroscopy along with the nominal
values of Scalmalloy®.
- Si Fe Mn Mg Zn Ti V Sc Al
0 0 0.3 4 0 0 0 0.6
0.4 0.4 0.8 4.9 0.25 0.15 0.05 0.8
data data data
Measured 0.04 ±0.01 0.1 ±0.004 0.5 ±0.003 3.95 ±0.05 <0.01 0.08 ±0.003 0.02 ±0.001 0.65 ±0.05
All investigated specimens were produced on an LPBF system EOS M 290 using an
Yb-fibre laser, having a recommended maximum laser power of 370 W and a focus diameter
of 100
m. The maximum build volume of the system is 250 mm
250 mm
325 mm.
In order to keep the oxygen content below 0.1% during the building process, the building
chamber is inerted with argon gas and a laminar argon flow is utilized.
The process strategy time homogenization is a function invented by EOS GmbH and
implemented in the EOSPRINT 2 software [
]. The volume energy E
is determined by
the ratio of laser power P and the product of scanning speed v, hatch distance h and layer
height l
. For fatigue testing, six different surface states were considered, whereas four
states were directly created during the manufacturing without any post-processing. The
as-built state of the surface only consists of the infill parameter set (Table 2). The three
different specimens that have contour parameter sets were manufactured with the infill
parameter set and subsequently the contour parameters, designated as contour A, B, and C
(Table 3). The latter represent potential parameter sets developed in a preliminary study
by the authors applying DoE with the objective to reduce the roughness of vertical walls
without introducing additional pores into the contour area [
]. The factors with the largest
influence on the surface roughness and density were identified as the scan speed v, the
laser power P, and the percentage overlap of the contour scan with the infill oci (Figure 1).
Table 2. Used infill parameter sets for Scalmalloy®, applying the strategy time homogenization.
Volume energy, EV[J/mm3]110
Line energy, EL[J/mm] 0.23
Laser power, P [W] 370
Scan speed, v [mm/s] 1600
Hatch distance, h [mm] 0.07
Layer height, lz[mm] 0.03
Exposure pattern Stripes
Stripe width [mm] 7
Layer rotation 67
Table 3.
In-process contour parameter sets with different line energy E
, laser power P, scan speed v,
and percentage overlap of the contour scan with the infill oci used in this study [44].
- Type EL[J/mm] P [W] v [mm/s] oci [%]
Contour A Post-cont. 0.9 370 411 50
Contour B Pre-cont. 0.6 370 600 50
Contour C Pre-cont. 0.07 230 3286 50
The results of the development of the contour parameters of the former study are
summarized as follows [44]:
Contour A: the lowest overall roughness of R
< 7
m (R
< 55
m) can be attained by
applying a post-contour with a high line energy in the range of E
= 0.9 J/mm. The
high line energy has a negative effect on the formation of keyhole pores, which start
to form in the E
-range between 0.6 J/mm and 0.75 J/mm. For contour A, a second
contour scan with a lower line energy of E
= 0.55 J/mm
was applied with an offset
to the first contour of 120%, which means a shift towards the inside of the sample.
Metals 2021,11, 1961 5 of 19
Contour B: the lowest roughness of R
< 9.8
m (R
= 65.2
m) utilizing a contour
parameter set with no significant increase in contour porosity is achieved when using
a pre-contour with E
= 0.6 J/mm and o
= 50%. Isolated larger pores can be detected
in the area of the vertical specimen edge.
Contour C: A reduction of roughness in the range of low line energies E
< 0.14 J/mm
can be achieved by using a pre-contour. The lowest roughness R
= 12.3
m (R
= 78.4
µm) is obtained at EL= 0.07 J/mm in combination with oci = 50%.
Figure 1.
Simplified illustration of the contour and infill vectors based on [
]. Additionally, the
percentage overlap of the contour scan and the infill o
depending on the contour melt pool width is
The individual parameters of the three contour parameter sets applied in this study
are summarized in Table 3:
The entire building platform with the fabricated specimens was subjected to heat
treatment immediately after the building process at 325
C for 4 h with slow cooling to set
maximum quasistatic properties [
]. Table 4gives an overview about all manufactured
fatigue specimens.
Table 4. Overview of all manufactured and tested fatigue specimens.
Name Surface Heat Treatment Building Direction
AB As-built
325 C, 4 h Z
T Turned
SP Shot peened
A Contour A
B Contour B
C Contour C
Figure 2shows the geometry of the manufactured fatigue specimens. All specimens
were manufactured with identical radii in the test areas.
Figure 2.
Manufactured and tested fatigue specimen used in this study. The testing area was applied
with the different surface.
Metals 2021,11, 1961 6 of 19
As-built specimens and specimens with contour parameters were manufactured in
the final geometry (Figure 2). Only the thread of the fixation to clamp the specimen into
the machine was turned afterwards. To avoid crack formation in the overhang region
adjacent to the examination area, manual post-processing was carried out by applying
rotating tools for grinding and rubbering on both sides adjacent to the examination area.
The two post-processing states represent the turned and the shot peened states. Turned
specimens were manufactured as a cylinder with a 10 mm diameter and turned to the final
geometry by reducing the diameter in the testing area by 6 mm. Shot peened specimens
were manufactured as as-built specimens, and shot peened after the heat treatment using
a peening media by Iepco AG Iepconorm A (nutshells, particle size = 200
followed by Iepconorm B-4 (C, Si, Mn, P, S, particle size = 125
m) and Iepconorm
C (ZrO
, SiO
, Al
, particle size = 400
m) at a pressure of 2 bar, respectively. The
blasting process was carried out for 20 s per specimen using peening media and a working
distance of 5 cm. The exact diameter of the test area of the specimens was measured before
performing the tests.
In order to obtain the quasistatic properties and to calculate the X-ray elastic constant
of residual stress measurements, standing (Z) and lying (X) cylinders were manufactured
on the same building platform in addition to the fatigue specimens to manufacture ten-
sile specimens with geometry B6
30 according to DIN 50125:2016 [
] for quasistatic
investigations. Half of the cylinders for the tensile tests were cut from the platform before
heat treatment. Afterwards, all cylinders were turned into the final geometry and the
stress–strain curve was obtained with a universal testing machine H&P 250, using strain
control according to DIN EN ISO 6892:2009-01 [48].
The roughness of the fatigue specimens in the test area was determined with a
Hommel-Etamic Turbo Wave V7.59 on a length of 15 mm with three measurements per
specimen variation. Because the test length was too short for a roughness measurement
according to DIN EN ISO 4288:1998 [
], the measurement direction was reversed and the
specimens were rotated at a certain angle.
The residual stresses were measured with x-ray-diffraction (XRD) using Cr-K
radiation at the {311}-reflex of Aluminium according to DIN EN 15305:2009-01 [50], using
the sin
-method. Therefore, a diffractometer by Stresstech, model Xstress G3R in a
–arrangement, was used. The residual stresses were calculated from seven
-angles from 0
. The X-ray elastic constant used to calculate the
residual stress values was
= 1.9
/N, which was calculated from the tensile
tests and [
]. Initially, the residual stresses were measured in loading direction on three
positions in the testing area of the fatigue specimens. The value was then evaluated with a
cross correlation to gather the measurement uncertainty.
Cross sections for microstructural investigations were prepared to investigate the
microstructure by optical microscopy (OM) using standard metallographic polishing tech-
niques with a final polishing step and 1
m OPS-suspension. OM-specimens were etched
with TIM3-etchtant (water, HCl, HF) and analysed using a LEICA DM4000 microscope.
SEM-analyses were performed with a Zeiss EVO 60 scanning electron microscope.
All manufactured specimens were examined regarding the porosity by 3D-computed
tomography using a GE v|tome|x 240 d with a nanofocus-tube. Only the testing area in
the centre of the specimens was considered in order to optimize the maximum possible
resolution. The achieved voxel-size is 16
m. To reduce beam hardening, an Sn-filter with
0.1 mm thickness was used. After reconstruction, the porosity was analysed with the
vgdefx-algorithm of the porosity module of VGStudioMAX 3.0 to determine the properties
and distribution of the pores. In Figure 3, a reconstructed model of the testing area with
the pore distribution and a coloured pore size scale is shown exemplarily.
The fatigue tests were performed with a Zwick 100 HFP 5100 following the bead string
method to draw the Woehler curves as described in DIN 50,100 [
]. The specimens have a
notch factor of k = 1 and were tested with a load ratio of R = 0.1. Fracture surfaces were
analysed by SEM to gather knowledge about crack initiation and propagation.
Metals 2021,11, 1961 7 of 19
Figure 3.
Exemplary reconstructed model of the testing area of an A-specimen with pore distribution
and coloured pore sizes.
3. Results
3.1. Quasistatic Tests
Table 5shows the results of the tensile tests according to DIN EN ISO 6892-1:2009 [
It is proven that the tensile strength R
and the yield strength R
were significantly
increased by the heat treatment, but the elongation at fracture A was decreased to the
same extent, which corresponds to the results in [
]. The modulus of elasticity did not
change significantly. The increase in yield strength is due to precipitation hardening of the
-precipitations that occur during the heat treatment [
]. It can also be observed
that standing (Z)-specimens drop slightly in strength compared to the lying (X)-specimens,
regardless of the heat treatment. This has also been observed in other LPBF-materials
and can be explained by the Schmid factor, i.e., the orientation of the maximum acting
shear stress to the sliding system [
]. For face-centred cubic materials such as aluminium,
the {111} <110> plane is the sliding system that has the lowest strength: Because there
is a preferred orientation in additively manufactured materials, a larger Schmid factor
can occur in the standing (Z)-specimens compared to the building directions, which
macroscopically manifests itself in a lower tensile strength and yield strength [52].
Table 5. Quasistatic properties of turned tensile specimens with geometry B6 ×30.
Orientation Heat Treatment Rm[MPa] Rp0,2 [MPa] A5 [%] E [GPa]
X As-built 352 ±1 290 ±3 24.5 ±1 70 ±2
Z As-built 349 ±2 269 ±1 22 ±2 70 ±1
X 325 C, 4 h 520 ±2 490 ±2 11.5 ±1 71 ±1
Z 325 C, 4 h 510 ±2 476 ±2 14 ±1 71 ±1
Compared to other AM-materials such as Inconel 718, the fluctuations in the me-
chanical properties depending on the building direction of less than 1% are practically
Figure 4a shows an exemplary overview of the fracture surface of a heat treated tensile
specimen investigated via SEM. In detailed view, the dimples can be observed in the
fracture microstructure, which is characteristic for a transcrystalline forced fracture. The
dimples are less (Figure 4b) or more (Figure 4c) distinct depending on the area observed.
These findings are a characteristic behaviour for ductile materials, such as aluminium
alloys, according to VDI 3822 Part 2 [56].
Metals 2021,11, 1961 8 of 19
Figure 4.
) Overview of an exemplary fracture surface of a standing (Z) and heat treated tensile specimen. (
) Exemplary
detailed view with less dimples. (c) Exemplary detailed view with dimples.
3.2. Metallographic and Porosity Analysis
In Table 6the surface roughness as well as the measured overall porosity is shown. It
can be stated that AB has the highest surface roughness and T the lowest. All specimens
with post-processing parameters T and SP as well as AB do not differ significantly in
their overall porosity, because they were manufactured only with the infill parameter
set. Figure 5provides overviews of the cross-sections of the test area from each fatigue
specimen examined by OM. The overall porosities in the infill of the specimens do not
differ significantly. The high surface roughness due to adhering powder particles and
agglomerations at the surface of the as-built specimen without contour scan (Figure 5a)
is clearly visible. The shot peening process (Figure 5b) removes the powder particles and
agglomerations to a high extent resulting in a smoother surface. The highest reduction
of the surface roughness was achieved by turning, as can be obtained from Table 6and
Figure 5c.
Table 6. Results of the roughness measurements and overall porosity measurement via CT.
Name Ra
Overall Porosity (CT)
Largest Pore
Distance of Largest Pore
to Surface
AB 17 ±2 101 ±7 0.007 ±0.004 0.0012 0.174
T 1 ±0.2 5 ±0,8 0.005 ±0.003 0.0007 0.160
SP 8 ±2 43 ±11 0.003 ±0.001 0.00151 0.164
A 5 ±2.5 23 ±11 0.252 ±0.05 0.00302 0.112
B 6 ±0.3 30 ±7 0.019 ±0.001 0.00133 0.096
C 11 ±0.1 63 ±4 0.012 ±0.003 0.00184 1.498
In general, the specimens built with contour parameters result in a smoother surface
in comparison to AB (Table 6). Among all contour parameters, specimen A exhibits the
highest overall porosity. In the etched state in Figure 5d, on the right hand side, it can be
noticed that the large pores are located in the contour area. Contour parameter C leads to
the highest surface roughness (Table 6and comparison of Figure 5d–f). It is noticeable that
in the etched state of specimens with contour parameter set A and B, microstructural effects
due to the contour scans are clearly visible. For specimens with contour parameter set A,
the smoother surface can be explained by the performed post-contour scan, which results in
the characteristic contour ring near the surface (Figure 5d, right hand side). Specimens with
contour parameter set B exhibit the same visible contour parameter ring in the cross section
in Figure 5e (right hand side), which is slightly thinner than in A. In contour parameter
C, the line energy of the pre-contour was significantly lower than the line energy of the
infill parameter: The overlap of o
= 50% (Table 3in comparison to Table 2) was enough to
completely re-melt the surface area, resulting in no defined contour parameter ring in the
etched state of cross sections on specimens with contour parameter set C (Figure 5f).
Metals 2021,11, 1961 9 of 19
Figure 5.
Overview of the metallographic cross sections of the (
) as-built-surface AB for comparison and the two post-
process-surface states (
) shot peened SP, (
) turned T, (
) contour parameter A, (
) contour parameter B, (
) contour
parameter C.
In Figure 6the distribution of the pores evaluated on the reconstructed CT-scans of the
specimens in dependence of the distance of the surface, is illustrated. It can be stated that on
specimens built with contour parameters A and B, the majority of the pores and the pores
with larger volume are located in the area near the surface, which is caused by the contour
scan. Only on specimens with contour parameter set C are the pores evenly distributed
throughout the specimen. This indicates that the line energy of the contour parameter set A
is high enough that keyhole-welding-mode described in [
] occurs and forms around
pores with larger volumes. Although the near-surface porosity is significantly decreased
with contour parameter set B, the amount of porosity still indicates that the reduction of
the line energy was not sufficient to prevent pores. However, on contour parameter set
C, the contour line energy was so low that the infill parameter set re-melted the contour
scan, which prevents keyhole-pore formation, but results in a significantly higher surface
roughness due to the adhesion of not fully melted powder to the surface.
3.3. Residual Stress Analysis
In Figure 7a, residual stress–depth profiles in the measuring area of the AB- and post-
processed (turning and shot-peening) fatigue specimens can be observed. It is noticeable
that the residual stresses of the AB-specimen change from low tensile residual stresses
at the surface to nearly no residual stresses with increasing depth. This behaviour was
also observed and simulated in [
] and can be explained by the higher cooling rate at the
surface: Due to the higher thermal conductivity of the material than the ambient air, inner
areas of the specimens cool down slower, because the heat flows through the centre of the
specimen into the building platform, which results in higher residual stresses at the surface
due to more rapid cooling than the centre. The material has no time to assume the most
Metals 2021,11, 1961 10 of 19
energetically favourable status, causing residual stresses to build up on the surface. The
residual stress–depth curve of the turned specimen shows a small area of residual stresses
at the surface, probably due to the induced material deformation owing to the turning
process (a). The shot peened specimen shows high compressive stresses at the surface.
The zero crossing of the residual stresses curve was shifted approximately 400
m into the
subsurface area.
Figure 6.
Pore size distribution of the manufactured specimens with differing contour parameters
dependent of the distance to the surface.
Figure 7.
) Residual stress–depth profiles of as-built (AB), shot peened (SP) and turned (T) spec-
imens. (
) Residual stress–depth profile of contour parameters A, B and C. The curves call the
Metals 2021,11, 1961 11 of 19
In Figure 7b, the residual stress–depth profile of specimens with contour parameters A,
B, and C, and the surface area of the etched square sections, is shown. Contour parameter
A shows a first maximum at approximately 40
m near the surface, similar to the AB-
specimens, which can be explained by the effect described above for the AB-specimens. At
approximately 350
m depth, a second maximum of approximately 60 MPa tensile residual
stresses were identified, which is most likely induced by the post-contour parameters: The
re-melting of the material results in a locally rapid cooling at the interface of the contour to
the bulk, which results in high residual stresses. The re-melted surface area can be seen
in the etched cross section above the graph in Figure 7b, green. The heat treatment after
the manufacturing process was not sufficient to remove those residual stresses. Specimens
built with contour parameter set B show the same second peak at approx. 350
m, but
the induced residual stresses are significantly lower in comparison to contour parameter
A. This leads to the assumption that a lower line energy induces fewer residual stresses.
Likewise, the contour on the B-specimens in the cross section in the etched state is visible
(Figure 7b, purple). Specimens built with contour parameter C do not show the second
peak. In comparison to A and B, only a pre-contour-parameter was used. The area of the
contour was completely remelted, which is the reason why no contour can be made visible
in the square section of specimens with contour parameter set C (Figure 7b, yellow) in
comparison to A and B.
3.4. Fatigue Investigations
The fatigue results of the specimens with differing surface roughness due to turning
(T) and shot peening (SP) in comparison to the as-built state without contour (AB) are
summarized in Figure 8. It is evident that all fatigue tests reveal a high scatter, even the
turned specimens. The fatigue strength limit for AB is around 40 MPa with a steep time
strength area in comparison to shot peened and turned specimens. The fatigue strength
area was more than doubled with shot peening SP or turning T, which is a similar result
in comparison to AlSi10Mg [
]. The fatigue strength of SP exhibits the highest fatigue
strength of all test series in this work and even exceeds the fatigue strength of turned
specimens. This behaviour indicates a significant influence of surface roughness, residual
stresses, and porosity on fatigue behaviour.
Figure 8.
Results of the fatigue tests with R = 0.1 for the post-process treatments turning (T) and
shot peening (SP) in comparison to specimens with as-built surface (AB) gathered with the bead
string method acc. to DIN 50100 [
]. Run-throughs are marked with an arrow. Fatigue results of the
in-process treatments of contour parameter A, B, and C in comparison to turned specimens (T) and
as-built surface specimens (AB). The lines are only for visualisation.
The as-built specimens without contour parameter sets (AB) have the highest surface
roughness of all tested specimens, which results in the lowest fatigue strength. In Figure 9a
the exemplary fracture surface of a specimen without contour parameters was investigated
Metals 2021,11, 1961 12 of 19
by SEM. A fatigue fracture with ductile residual force fracture according to VDI 3822 Part
2 [
] can be obtained. Fracture origins were located by fracture paths running towards
them, which are caused by the crack propagation. Multiple fracture origins (exemplary
view in Figure 9b) are located at the circumference of the surface, which indicates that
the notches due to the high roughness promote the crack initiation, which was stated by
Murakami [
] as a characteristic property of AM-materials in general. The multiple cracks
then unite and reduce the surface area, which overall results in a poor fatigue strength.
Figure 9.
) Exemplary overview of a fracture surface of an as-built specimen AB. (
) AB detailed view of a crack initiation
area, (
) exemplary overview of a fracture surface of a shot peened specimen SP, (
) SP detailed view of the crack initiation
area, (
) exemplary overview of a fracture surface of a turned specimen T, (
) T detailed view of a crack initiation area. All
fracture surfaces inside one charge are of similar character, regardless of the stress amplitude. R = 0.1.
Metals 2021,11, 1961 13 of 19
With shot peening, the surface with lower roughness and the compressive residual
stresses are improving the fatigue properties significantly (SP in Figure 8), which is a
similar behaviour as found by AlSi10Mg [
]. In Figure 9c, an exemplary overview of an
SP fracture surface investigated by SEM can be seen. Unlike AB, only one specific fracture
origin on every tested specimen is present, which can exemplarily be seen in detail in
Figure 9d. It is noticeable that the fracture origin is located in all investigated specimens
approximately 100
m below the surface. This indicates that the shot peening increases
the resistance against crack initiation due to the induction of compressive residual stresses.
Pores or oxides in the subsurface area, where the compressive stresses are lower or even in
the tensile range, then act as crack initiation areas (residual stresses in Figure 7and fracture
surface in Figure 9c). The ruptured surface area, which can be seen in Figure 10e, does
not exhibit microscopic or macroscopic notches in comparison to AB. These findings agree
with [
], who showed the same effect on LPBF-IN718, where the shot peening shifts the
positive residual stresses from the surface towards the inside of the specimens. In the case
of IN718, this effect even lowers the fatigue properties due to higher notch sensitivity.
Figure 10. Cont.
Metals 2021,11, 1961 14 of 19
Figure 10.
Exemplary SEM-images of the surfaces before the fatigue test of (
) as-built (AB), contour parameter sets (
) A,
(c) B, (d) C, post process improvements, (e) shot peened (SP), and (f) turned (T) specimens.
Turned specimens exhibit the lowest surface roughness of all tested specimens. Only
the slight grooves of the machining induce recurring peaks and valleys together with the
inner porosity or imperfections that were converted to open porosity at the surface due to
the machining process. The resulting surface area can be seen in Figure 10f, showing an
exemplary imperfection (Figure 10f, right hand side). In comparison to shot peened specimens,
the overall surface roughness is lower (Table 6), but a certain amount of tensile residual stresses
at the surface was induced due to the turning process (Figure 7). In Figure 9e, an exemplary
SEM-image of the fracture surface of a turned specimen can be seen. Even at lower stress
amplitudes, the crack initiation area is always located at the surface. The cracks initiate
on small imperfections at the surface, such as little breakouts or open porosity, which
can exemplarily be seen in Figure 9f. Aboulkhair et al. [
] investigated the same crack
initiation areas on turned fatigue specimens for material AlSi10Mg. The open porosity and
imperfections at the surface act as notches and are therefore the most likely crack initiation
areas. Additionally, the induced residual stresses due to the machining process increases
the probability that the crack initiates at surface defects even further.
In comparison to the post-process parameters tested in this work, the in-process
contour parameters only increase fatigue strength by about 5% to 20%.
Contour parameter A has a similar time strength area as AB and an only 3 MPa higher
fatigue strength. B and C show slightly higher fatigue strengths. Specimens with contour
parameters B and C have the same fatigue strength of approximately 48 MPa, but the time
strength graph of C is steeper.
It is evident that the developed contour parameters only slightly increase the fatigue
strength in comparison to those of as-built specimens without a contour of about 3 MPa to
8 MPa. Despite the significant reduction of surface roughness with contour parameter A
in comparison to AB, the fatigue strength is only increased by 3 MPa. Nevertheless, the
remaining notches at the surface are sufficient to act as crack initiation areas; in Figure 10a,
SEM-images of the surface of a specimen AB can be obtained. It can be investigated that
the high roughness and powder agglomerations (spherical material) lead to large notches.
On the surface of specimen A (Figure 10b), not as many powder agglomerations or as
much roughness can be investigated, therefore it can be seen that the surface is littered with
microscopic notches, which have a high notch factor, which lowers the fatigue strength of
the material [59].
In Figure 11a,b, an exemplary fracture surface of A can be seen. The subsurface
porosity present in this specimen, resulting from the contour scan with high line energy
(Figure 5d,e and Figure 6), does not act as crack initiation areas. This leads to the assumption
Metals 2021,11, 1961 15 of 19
that sharp, microscopic notches together with higher tensile residual stresses at the surface,
in comparison to the subsurface area (Figure 7), have a significantly higher impact on
fatigue strength than subsurface pores at R = 0.1.
Metals 2021, 11, x FOR PEER REVIEW 16 of 20
(a) (b)
(c) (d)
(e) (f)
Figure 11. (a) Exemplary overview of a breaking surface of A. (b) Detailed view of a crack initiation area of A, (c) exemplary
overview of a breaking surface of B, (d) detailed view of the crack initiation area of B, (e) exemplary overview of a breaking
surface of C, (f) crack initiation area of C. All breaking surfaces inside one charge are of similar character, regardless of the
stress amplitude. R = 0.1.
In Figure 11c, an exemplary overview of the fracture surface of B can be seen. Similar
to turned specimens in the detail view in Figure 11d, it can be gathered that crack initiation
was at an open pore at the surface of the specimen, which indicates that the high porosity
at the surface and subsurface in comparison to C (Figure 6) lowers the fatigue strength. In
the exemplary overview (Figure 11e) and detailed view (Figure 11f) of the fracture surface
Figure 11.
) Exemplary overview of a breaking surface of A. (
) Detailed view of a crack initiation area of A, (
) exemplary
overview of a breaking surface of B, (
) detailed view of the crack initiation area of B, (
) exemplary overview of a breaking
surface of C, (
) crack initiation area of C. All breaking surfaces inside one charge are of similar character, regardless of the
stress amplitude. R = 0.1.
Metals 2021,11, 1961 16 of 19
In Figure 11c, an exemplary overview of the fracture surface of B can be seen. Similar
to turned specimens in the detail view in Figure 11d, it can be gathered that crack initiation
was at an open pore at the surface of the specimen, which indicates that the high porosity at
the surface and subsurface in comparison to C (Figure 6) lowers the fatigue strength. In the
exemplary overview (Figure 11e) and detailed view (Figure 11f) of the fracture surface of
C, it can be obtained that the crack initiation area was at a notch due to the high roughness.
The fatigue strength of B and C is nearly similar, because the lower surface roughness
on B is compensated by its higher subsurface porosity (Figure 6) of B in comparison to C:
Either the rough surface or porosity in the surface area act as notches and therefore as crack
initiation areas. Regarding the microscopic investigations on the surface area of specimens
B (Figure 10c) and C (Figure 10d) in comparison to A or AB, no microscopic notches and no
high amount of powder agglomerations can be found. The fatigue strength is not increased
in the same manner as it is for post-process parameters such as turning or shot peening.
In the case of turning, the porosity and roughness of the surface is nearly removed and
in the case of shot peening the compressive residual stresses increase the crack initiation
4. Conclusions
In this work, specimens of Scalmalloy
with differing surface states were manufac-
tured by LPBF in order to investigate the influence of the surface condition and the influence
of in-process optimization strategies and post-process improvements of the surface on the
fatigue behaviour. Therefore, three different contour parameters and a turning and shot
peening process were compared. Metallographic analysis, fracture surface analysis by SEM,
residual stress measurements via XRD, roughness measurements, porosity analysis via CT,
and tensile and fatigue tests were performed. The following findings were achieved:
Compared to the as-built surface state that only had infill parameters, the post pro-
cessing steps using shot peening and turning increases the fatigue strength by a factor
>2. In contrast, the use of contour parameters improves fatigue properties by only
around 10 to 20% in comparison to the as-built state.
Crack initiation on contour-parameter specimens occur on sharp, microscopic notches
at the surface together with high tensile residual stresses. Even subsurface pores
induced by a contour scan in keyhole-welding-mode do not change the crack initiation
area at the surface.
A contour scan with low line energy results in the same fatigue strength as a contour
scan with medium line energy, despite the higher roughness, because the overall
porosity, even in the surface area, is lowest, which results in less supporting crack
initiation areas in combination with the sharp notches.
Shot peened and turned specimens show comparable fatigue behaviours, despite the
higher roughness of shot peened specimens: This disadvantage is compensated by
the induced compressive residual stresses at the surface, which lowers the tendency
for crack initiation at the surface and shifts the crack initiation area into the specimen.
This study shows that a significant improvement of the fatigue strength can only
be achieved due to conventional post-processing surface treatments. The best results
for improving the fatigue strength with in-process parameters could be achieved with a
contour scan with low and medium line energies that are not in the keyhole-welding-
mode. It is important to note that only fatigue tests under a tension swell load of
R = 0.1 were performed. Further investigations have to be conducted in order to re-
veal whether tension–compression load leads to identical results. The subject of the current
research is to investigate whether a poor contour scan along with shot peening might be
counterproductive, because the tensile residual stresses might be superimposed by the
local stresses at subsurface porosity.
Metals 2021,11, 1961 17 of 19
Author Contributions:
Conceptualization, J.M., T.R., and H.C.H.; investigation, J.M. and T.R.;
writing—original draft preparation, J.M.; writing—review and editing, T.R., H.C.H., M.O., M.W.,
and E.A.; funding acquisition, H.C.H.; supervision, M.O. All authors have read and agreed to the
published version of the manuscript.
The research and development project ‘BadgeB’that forms the basis for this publication is
funded within the scope of the “Additive Fertigung—Individualisierte Produkte, komplexe Massen-
produkte, innovatiove Materialien” (Pro Mat 3D) by the Federal Ministry of Education and Research,
funding no. 02P15B154. The project BadgeB is managed by the KIT project management agency
“Projektträger Karlsruhe—Produktion und Fertigungstechnologien”. The authors are responsible for
the content of this publication.
Data Availability Statement:
The data presented in this study are available on request from the
corresponding author.
Conflicts of Interest:
The funders had no role in the design of the study; in the collection, analyses,
or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.
The authors declare no conflict of interest.
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... Despite the mechanical properties that can be achieved with heat treatments, there are some inherent factors of L-PBF technology that can impede components performance, especially from the structural point of view and regarding the fatigue. These factors include: the effects of built direction, post-processing surface treatments, structure of micro-defects, residual stresses, multiaxial stress state and, obviously the surface roughness [11][12][13]. ...
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Thanks to Laser Powder Bed Fusion (L-PBF) technology, SCALMALLOY® was the first aluminum powder material designed for Additive Manufacturing (AM), achieving a fine microstructure with high performance that is comparable to other cast materials. Despite the mechanical properties that can be achieved, there are some inherent factors that can impede components performance (i.e., surface roughness). Parts produced by L-PBF are usually characterized by rough “as-built” surfaces; hence, it is fundamental during the design phase to understand and consider how the quality of surfaces impacts on the part performance. This paper aims to provide a Computer-Aided Engineering (CAE) workflow to design components with different finishing regions in accordance with the functional distinction that exists among them. To achieve this goal, a comparison of the mechanical properties achieved for SCALMALLOY® specimens with and without post-processing is here assessed to fit proper material models for numerical simulation purposes. The material models, built with/from experimental data, are fit to functionally adapt the performance of 3D-printed objects inside CAE simulations like a Functionally Graded Material (FGM). A CAE design workflow is here applied to a case study, suitable to demonstrate how the methodology may support the integrated product–process design of structural parts reducing the cost of post-processing in AM. This approach may mitigate the performance decrease of “as-built” surfaces since the experimental results show a different fatigue endurance limit between the “as-built” and CNC machined specimens about of three times.
Laser Powder Bed Fusion (LPBF) is considered a promising technology for NiTi intermetallic components production. However, different defects such as near-surface porosity inherent to the process are limiting the standardization and widespread of the technology due to the deterioration of the mechanical properties. In this work, a comprehensive study was implemented to reduce the near-surface porosity by the laser polishing of additively manufactured samples from pre-alloyed NiTi powder by LPBF. Different laser polishing process conditions were employed for the treatment of the near-surface region. Results have shown that the most effective treatment aimed at porosity healing occurs at a laser power of 540 W, a scanning speed of 25 mm/s, a hatch spacing of 200 μm, and a laser spot of 500 μm. The phase identification and martensite-austenite transformation temperatures characterization were carried out before and after the laser polishing of the NiTi samples; the influence of the laser polishing on the functional properties of the material was demonstrated. It was shown that the application of required linear energy density for successful healing of the near-surface porosity leads to the local solutionizing effect. Additionally, electrochemical studies were carried out in Hank's solution and demonstrated that the corrosion resistance was improved after laser polishing.Graphical Abstract
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Their high strength-to-weight ratio, good corrosion resistance and excellent thermal and electrical conductivity have exponentially increased the interest in aluminium alloys in the context of laser-based powder bed fusion (PBF-LB/M) production. Although Al-based alloys are the third most investigated category of alloys in the literature and the second most used in industry, their processing by PBF-LB/M is often hampered by their considerable solidification shrinkage, tendency to oxidation, high laser reflectivity and poor powder flowability. For these reasons, high-strength Al-based alloys traditionally processed by conventional procedures have often proved to be unprintable with additive technology, so the design and development of new tailored Al-based alloys for PBF-LB/M production is necessary. The aim of the present work is to explore all the challenges encountered before, during and after the PBF-LB/M processing of Al-based alloys, in order to critically analyse the solutions proposed in the literature and suggest new approaches for addressing unsolved problems. The analysis covers the critical aspects in the literature as well as industrial needs, industrial patents published to date and possible future developments in the additive market.
Mechanical properties and microstructure of Scalmalloy® processed by Selective Laser Melting (SLM) under certified conditions are studied. Two types of samples are considered: as-built and heat-treated at 325 °C for 4 h. The heat treatment does not modify the grain size or the texture. SEM images show a typical bi-modal microstructure comprising equiaxed fine grained zones (FGZ) with ≈ 0.6 μm grain size and coarse grained zones (CGZ) with ≈ 7.5 μm grain size (longitudinal axis). According to electron backscatter diffraction (EBSD) analysis, FGZ have ≈ 0.8 μm equivalent circle diameter (considering grain boundaries those with misorientation >15⁰) and contain areas with and without texture, while CGZ have ≈ 2.7 μm equivalent circle diameter, and most of the grains are oriented with [001] crystal direction along the building direction. On the contrary, the heat treatment produces a stress relief that is evidenced in EBSD by reductions in the low angle grain boundaries (LAGB) content, grain orientation spread (GOS), and dislocation density. Both samples show plastic deformation before fracture and the as-built condition features Portevin-Le-Chatelier (PLC) effect. The heat-treated samples exhibit the highest strength (yield strength = 454 ± 3 MPa; tensile strength = 509 ± 6 MPa) and hardness (165 ± 3 HV) while the lowest percentage of reduction of area (16 ± 6%) and elongation (11 ± 2%). Finally, the electrical conductivity of the material increases after the heat treatment which is associated with the redistribution of Mg-atoms out of the solid solution and the growth of precipitates.
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A selective laser melted Al-4.5Mg-0.6Sc-0.3Zr alloy, also known as Scalmalloy®, was cold-rolled at two different reductions in thickness to observe the effect on microstructure and mechanical properties. Optical microscopy, scanning electron microscopy, tensile testing with digital image correlation, surface roughness testing and microhardness testing was used to examine samples. Hardness increased significantly in the 30% cold-rolled condition – 20.2% in the horizontal plane, 12.5% in the transverse plane, and 18.2% in the longitudinal plane. Surface roughness was reduced to as low as 1.2μm for cold-rolled conditions compared to >15μm for as-built conditions. Yield strengths were less for cold-rolled conditions compared to as-built and ultimate tensile strengths were similar for as-built and 10% cold rolled conditions. However, 30% cold-rolled condition saw a 10.8% increase compared to as built. Strain distribution among samples was uniform for as-built conditions but concentrated in certain areas for cold-rolled conditions. SEM micrographs of fracture surfaces showed some very ductile fractures evident from dimpling. The microstructure was heavily deformed for all cold-rolled conditions in all build directions showing flattened, elongated grains and grain refinement in both the coarse grain and fine grain regions present in Scalmalloy, especially for the 30% cold-rolled condition.
Scalmalloy alloys fabricated by Laser-Powder Bed Fusion (L-PBF) process are extensively used in the field of aerospace, construction, and automotive applications due to its lightweight, corrosion resistant qualities and fine microstructure. But, there was a lack of in-depth study in the nano-level mechanical properties of the Scalmalloy. So, the prime novelty of this research work is to examine the nano-level mechanical behavior with respect to nano-level precipitations due to Sc-Zr addition. In this work, the L-PBFed Scalmalloy was examined for the irregularities and the grain distributions using Optical Microscopy (OM) and Field Emission Scanning Electron Microscopy (FE-SEM) with Energy-Dispersive X-ray spectroscopy (EDX) analysis. The formation of Al3Sc precipitates was identified using X-ray diffraction (XRD) and Transmission Electron microscopy (TEM) analysis. Further, the influence of precipitates on the grain orientation was evaluated through Electron Backscatter Diffraction (EBSD) analysis. The nano-hardness and nano-wear testing was performed on the L-PBFed Scalmalloy. The maximum nano-hardness and reduced modulus of 2.19 GPa and 87.50 GPa was obtained from the average of 400 indentations. The nano-hardness mapping and reduced modulus mapping proved the homogenous dispersion of the precipitation. From the topographic wear tracks in the nano-wear analysis, the average worn-out height of 250.89 nm was observed and the fluctuation in the co-efficient of friction (CoF) confirms the formation of ceramic phase Al3Sc precipitates which resists the worn-out height. Tensile strength of the L-PBFed Scalmalloy was evaluated based on the 3 strain rates (3.3 × 10⁻⁵, 1.6 × 10⁻⁵ and 2 × 10⁻⁴ s⁻¹) and the outcome shows that the Scalmalloy exhibits negative strain rate and serration caused by Portevin-Le Chatelier effect on the Al3Sc precipitates and the Mg content.
Localized corrosion of Scalmalloy® processed by Selective Laser Melting (SLM) under certified conditions is studied and correlated with the intermetallic particles present in the microstructure. Two types of samples are considered: as-built and heat-treated samples. The heat treatment does not modify the pitting corrosion performance. All samples exhibit two groups of pits with different repassivation behaviour. Si-bearing and Fe-bearing phases produce pitting, while Al(Sc,Zr) is not found to be important. The two types of samples are resistant to intergranular corrosion (IGC) according to ASTM G67 thanks to the low content of β phase. However, the heat-treatment slightly increases the IGC-susceptibility.
Conference Paper
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Unconventional structures in materials may be generated by the corresponding choice of alloys in combination with a rapid extraction of heat. Improved melt spinning procedures allow to solute more than 1.0 wt% Sc in the Al-lattice. Therefore a critical cooling rate is necessary, which can only be achieved by ribbon thickness lower than 25 μm. In the next step, the flakes are usually compressed with high pressures and temperatures, whereby the fine microstructures may be damaged due to grain growth. However, that grain growth can be successfully suppressed by nano-scaled Al3(Sc1-x; Zrx) precipitations. Spark Plasma Sintering (SPS) is a short time sintering technique related to hot pressing, which uses a pulsed electrical current for the direct heating of the sinter body. The use of high heating rates and short dwell times can minimise grain growth, which often leads to improved material properties. Important technological benefits, such as the energy efficiency due to short processing time, fewer processing steps, elimination of the need for sintering aids, and near net shape capability, are possible.
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The scandium modified aluminium alloy Scalmalloy® is specifically developed for the use in laser‐based powder bed fusion (PBF‐LB). It is supposed to show potential in the production of lightweight structures due to its high specific strength compared to other aluminium alloys. A limiting factor is the high surface roughness of additively manufactured parts, which has a negative influence on its mechanical properties, especially under cyclic loads. In order to reduce the surface roughness, methods of design of experiments (DoE) are applied to develop contour parameters. Additionally, the formation of pores in keyhole‐mode welding and strategies to reduce the porosity in the contour area are investigated. The surface roughness of vertical walls can be reduced down to Ra < 7 μm using contour scans with a line energy E L >0.9 J mm ⁻¹ but keyhole pores start to form applying E L >0.6– 0.75 J mm ⁻¹ . Two contour parameter sets in different E L ‐ranges are developed that can be used to reduce the surface roughness compared to parameter sets without contour scans, without increasing the porosity in the contour area. Their impact on the mechanical properties has to be further investigated.
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It is often required to add a support structure for overhanging surface in additive manufacturing (AM), especially during laser powder bed fusion (LPBF), which has a considerable impact on the promotion and wide application of this technique. The aim of this study is to evaluate the influence of process parameters on the overhanging surface roughness of AlSi10Mg specimens without a support structure under different build angles, 30°, 45° and 60°, respectively. A design of experiment (DoE) approach was adopted to clarify the correlations between the overhanging surface roughness and the causative the process parameters such as contour laser power, contour scanning speed, infill laser power, infill scanning speed and infill scanning angle. The results showed that the influence of each process parameter on the overhanging surface roughness varies with the change in build angle. Above all, the dominant factor was the contour scanning speed. The results of the comparison show that the contour process parameters are much more influential than the infill process parameters. In addition, the process parameters affected the overhanging surface roughness were on the decrease as the build angle increases. When the build angle was 30°, it was necessary to comprehensively consider the influence of infill scanning angle, contour scanning speed and infill scanning speed to obtain a smooth overhanging surface with a low roughness. However, as the build angle increased to 60°, only the contour scanning speed had a highly significant impact. The results also revealed that apart from the contour scanning speed, the interaction between contour scanning speed and contour laser power should be given priority when the build angle is 30° or less. The lowest roughnesses of overhanging surface with build angles of 30°, 45° and 60° were 10.9 μm, 7.9 μm and 4.22 μm, respectively.
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Additive manufacturing allows for a great degree of design freedom and is rapidly becoming a mainstream manufacturing process. However, as in all manufacturing processes, it has its limitations and specificities. Equipping engineers with this knowledge allows for a higher degree of optimization, extracting the most out of this technology. Therefore, a specific part design was devised and created via L-PBF (Laser Powder Bed Fusion) using AlSi10Mg powder. Certain parameters were varied to identify the influence on material density, hardness, roughness, residual stress and microstructures. It was found that on heat treated parts laser pattern strategy is one of the most influential aspects, showing that chessboard and stripes 67° improved outcome; average Ra roughness varied between 8–12 µm, residual stress was higher on vertical surfaces than horizontal surfaces, with the combination of support structures and stripes 67° strategies generating the lowest residual stress (205 MPa on a lateral/vertical face), hardness was non-orientation dependent and larger on samples with chessboard fabrication strategies, while microstructures were composed of α–Al dendrites surrounded by Si particles. The distribution and grain size of the microstructure is dependent on location regarding melt pool and HAZ area. Furthermore, Al–Mg oxides were encountered on the surface, along with pores generating from lack of fusion.
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The family of additive manufacturing techniques has been attracting significant attention of manufacturers and researchers, due to its unrivalled flexibility to fabricate and repair geometrically complex objects. However, material shaping is not sufficient: wide adoption of additive manufacturing can only occur upon the achievement of satisfactory mechanical performance in terms of structural integrity. The present study exploits a wide range of micro-scale experimental techniques to shed light on fatigue failure mechanisms of Laser Metal Deposition IN718 Ni-base superalloy, and to study the effect of shot peening treatment. Thorough microstructural and fractographic analyses revealed the main deformation mechanism associated with twinning during crack propagation, while crack initiation was found to be promoted by both slip system deformation and twinning around microstructural defects, rather than at sample free-surfaces. It was found that precipitates played a major role in determining the deformation mode. It was discovered that in this case-study, shot-peening residual stresses may have a detrimental effect, in view of the presence of the largest defects within a region where tensile residual stress was present. The results presented here aim at improving understanding of failure mechanisms and thus defining future directions of development for manufacture optimisation.
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Additive manufacturing (AM) is expected to be a promising new manufacturing process for components having complex geometry. However, the disadvantage or challenge of AM is the presence of defects and surface roughness which are inevitably produced by the manufacturing process. Without strict and reliable quality control of components regarding defects and surface roughness, we cannot positively admire the advantages of AM as the new technology. In this paper, the method of quantitative evaluation of defects with complicated shapes and configurations, is explained from the viewpoint of small 3D cracks. First of all, the ideal fatigue strength (goal) to be attained by AM is discussed. Use of Fatigue-Grades from 5 (ideal strength) to 1 (lowest grade) is proposed for mutual comparison of fatigue performance of AM materials. The factors which decrease fatigue strength and degrade tensile properties are quantitatively analyzed. It is verified that the fatigue limit of AM materials is determined by the presence of nonpropagating crack emanating from defects. The reasons are made clear for fatigue fracture from surface defects even by tension-compression fatigue test as due to the difference of population of defects and increase in stress intensity factor for surface crack compared to internal cracks. Practical guides will be presented for the fatigue design and development of high quality AM materials, based on the combination of the 3D defect analysis, the statistics of extremes on defects, and the √area parameter model. It is shown that fatigue notch effect in AM materials is influenced by probability of presence of defects at notch root.
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The development of additively manufactured structures and products should preferably be performed in a virtual environment from the first sketch to the printed component. In order to get a high quality product that can be used as safety relevant component under cyclic loading, the fatigue assessment including the structure optimization with respect to the exploitation of the lightweight design potential requires suitable numerical approaches. Therefore, the local component-related material behavior has to be determined and considered. During the AM process multiple influences on the cyclic material behavior, caused for instance by the exposure lighting strategy or surface conditions, lead to mutual interactions between the load and built orientation on the fatigue properties. Caused by industrial needs for short development cycles and reduced numerical effort, the requirements on a digital twin for the fatigue assessment are growing and only satisfying with an adjusted numerical concept with optimized material parameters. In this work, the influences of the manufacturing process on the cyclic stress-strain behavior as well as on the fatigue life will be discussed. Therefore, load- and strain-controlled fatigue test were performed and the results are summarized in order to establish a better understanding on the fatigue behavior of additively manufactured components and parts.
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The fatigue behaviour of an AlSi10Mg alloy processed by laser powder bed fusion (L-PBF) and subjected to different surface finishing processes was investigated paying special attention to the residual defects on the surface and the dominant fatigue failure mechanisms. Roughness measurements and qualitative surface morphology analysis showed smooth surfaces in the case of vibro-finishing and machining followed by polishing. The fatigue performance did not reveal to be directly related to surface roughness, but residual intrusions left on the finished surfaces. Post-mortem analysis showed single-or multiple-crack nucleation from pores opened on the surface, un-melted powders, or spatters considered as typical L-PBF defects. A fatigue limit of 195 MPa for machined and polished samples was obtained by substantial removal of surface and subsurface defects.
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Metal additive manufacturing, despite of offering unique capabilities e.g. unlimited design freedom, short manufacturing time, etc., suffers from raft of intrinsic defects. Porosity is of the defects which can badly deteriorate a part’s performance. In this respect, enabling one to observe and predict the porosity during this process is of high importance. To this end, in this work a combined numerical and experimental approach has been used to analyze the formation, evolution and disappearance of keyhole and keyhole-induced porosities along with their initiating mechanisms, during single track L-PBF of a Ti6Al4V alloy. In this respect, a high-fidelity numerical model based on the Finite Volume Method (FVM) and accomplished in the commercial software Flow-3D is developed. The model accounts for the major physics taking place during the laser-scanning step of the L-PBF process. To better simulate the actual laser-material interaction, multiple reflection with the ray-tracing method has been implemented along with the Fresnel absorption function. The results show that during the keyhole regime, the heating rises dramatically compared to the shallow-depth melt pool regime due to the large entrapment of laser rays in the keyhole cavities. Also a detailed parametric study is performed to investigate the effect of input power on thermal absorptivity, heat transfer and melt pool anatomy. Furthermore, an X-ray Computed Tomography (X-CT) analysis is carried out to visualize the pores formed during the L-PBF process. It is shown, that the predicted shape, size and depth of the pores are in very good agreement with those found by either X-CT or optical and 3D digital microscopic images.
This study investigated the effects of the ‘as-built’ condition on the fatigue properties of an AlSi10Mg aluminium alloy manufactured by Selective Laser Melting (SLM) in different orientations. The results showed a moderate correlation between defects/surface features at the fracture origin and the roughness measurements. On the other hand, an analysis based on the size of features at fracture origin and residual stress just below the surface proved to be successful. Although considering the synergistic effect of different variables increased the scatter in fatigue life data points, a fatigue strength model including short crack effect and residual stresses made it possible to estimate the EIFS close to the real critical features. Results for life prediction models based on EIFS showed that 87% of the estimated number of cycles to failure are within a factor of 0.5-2 with respect to the experimental number of cycles to failure.