Microstructure evolution and mechanical properties at high temperature of selective laser melted AlSi10Mg

Abstract

In this study, the microstructure and tensile properties of selective laser melted AlSi10Mg at elevated temperature were investigated with focus on the interfacial region. In-situ SEM and in-situ EBSD analysis were proposed to characterize the microstructural evolution with temperature. The as-fabricated AlSi10Mg sample presents high tensile strength with the ultimate tensile strength (UTS) of ∼450 MPa and yield strength (YS) of ∼300 MPa, which results from the mixed strengthening mechanism among grain boundary, solid solution, dislocation and Orowan looping mechanism. When holding at the temperature below 200 °C for 30 min, the microstructure presents little change, and only a slight decrement of yield strength appears due to the relief of the residual stress. However, when the holding temperature further increases to 300 °C and 400 °C, the coarsening and precipitation of Si particles in α-Al matrix occur obviously, which leads to an obvious decrease of solid solution strength. At the same time, matrix softening and the weakness of dislocation strengthening also play important roles. When the holding temperature reaches to 400 °C, the yield strength decreases significantly to about 25 MPa which is very similar to the as-cast Al alloy. This might be concluded that the YS is dominated by the matrix materials. Because the softening mechanism counteracts work hardening, the extremely high elongation occurs.

Full text

Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
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at
ScienceDirect
Journal
of
Materials
Science
&
Technology
j
o
ur
nal
homepage:
www.jmst.org
Research
Article
Microstructure
evolution
and
mechanical
properties
at
high
temperature
of
selective
laser
melted
AlSi10Mg
Y.
Caoa,b,
X.
Lina,b,∗,
Q.Z.
Wanga,b,
S.Q.
Shia,b,
L.
Maa,b,
N.
Kanga,b,∗,
W.D.
Huanga,b
aState
Key
Laboratory
of
Solidification
Processing,
Northwestern
Polytechnical
University,
Xi’an
710072,
China
bKey
Laboratory
of
Metal
High
Performance
Additive
Manufacturing
and
Innovative
Design,
MIIT
China,
Northwestern
Polytechnical
University,
Xi’an
710072,
China
a
r
t
i
c
l
e
i
n
f
o
Article
history:
Received
27
March
2020
Received
in
revised
form
23
April
2020
Accepted
24
April
2020
Available
online
9
July
2020
Keywords:
Selective
laser
melting
AlSi10Mg
Microstructure
In-situ
EBSD
High
temperature
tensile
property
a
b
s
t
r
a
c
t
In
this
study,
the
microstructure
and
tensile
properties
of
selective
laser
melted
AlSi10Mg
at
elevated
temperature
were
investigated
with
focus
on
the
interfacial
region.
In-situ
SEM
and
in-situ
EBSD
anal-
ysis
were
proposed
to
characterize
the
microstructural
evolution
with
temperature.
The
as-fabricated
AlSi10Mg
sample
presents
high
tensile
strength
with
the
ultimate
tensile
strength
(UTS)
of
∼450
MPa
and
yield
strength
(YS)
of
∼300
MPa,
which
results
from
the
mixed
strengthening
mechanism
among
grain
boundary,
solid
solution,
dislocation
and
Orowan
looping
mechanism.
When
holding
at
the
tem-
perature
below
200 ◦C
for
30
min,
the
microstructure
presents
little
change,
and
only
a
slight
decrement
of
yield
strength
appears
due
to
the
relief
of
the
residual
stress.
However,
when
the
holding
temperature
further
increases
to
300 ◦C
and
400 ◦C,
the
coarsening
and
precipitation
of
Si
particles
in
␣-Al
matrix
occur
obviously,
which
leads
to
an
obvious
decrease
of
solid
solution
strength.
At
the
same
time,
matrix
softening
and
the
weakness
of
dislocation
strengthening
also
play
important
roles.
When
the
holding
temperature
reaches
to
400 ◦C,
the
yield
strength
decreases
significantly
to
about
25
MPa
which
is
very
similar
to
the
as-cast
Al
alloy.
This
might
be
concluded
that
the
YS
is
dominated
by
the
matrix
materials.
Because
the
softening
mechanism
counteracts
work
hardening,
the
extremely
high
elongation
occurs.
©
2020
Published
by
Elsevier
Ltd
on
behalf
of
The
editorial
office
of
Journal
of
Materials
Science
&
Technology.
1.
Introduction
Additive
manufacturing
(AM)
is
one
of
the
emerging
technolo-
gies
that
can
produce
the
parts
with
complex
internal
structures
from
a
three
dimensional
computer
aided
design
(CAD)
model
in
a
layer
by
layer
strategy
[1,2].
Selective
laser
melting
(SLM)
technique
is
an
important
powder
bed
fusion
AM
technique
for
the
metallic
materials,
which
allows
the
desired
components
to
be
fabricated
with
the
high
strength
and
stiffness.
In
general,
the
use
of
the
high
power
laser
with
small
beam
diameter
and
high
scanning
speed
leads
to
a
rapid
thermal
cycles
in
the
deposit
with
high
heating
and
cooling
rate
(about
103-108
◦C/s)
[3,4].
Therefore,
the
SLMed
sample
always
presents
ultrafine
micro
substructures
and
high
mechanical
strength
[5,6].
However,
the
significant
high
thermal
stress
is
also
easy
to
cause
undesirable
residual
stress/strain,
such
as
distortion
and
cracks
in
the
deposit.
∗Corresponding
authors
at:
State
Key
Laboratory
of
Solidification
Processing,
Northwestern
Polytechnical
University,
Xi’an
710072,
China.
E-mail
addresses:
xlin@nwpu.edu.cn
(X.
Lin),
nan.kang@nwpu.edu.cn
(N.
Kang).
AlSiMg
is
widely
utilized
in
SLM
technologies
[3,7–10],
in
which
the
microstructure
is
characteristic
as
a
network
of
␣-Al
(FCC)/Si
eutectic
along
the
inter–dendrite
of
the
supersaturated
primary
␣-
Al.
Given
to
the
rapid
solidification
induced
fine
microstructural
feature,
the
SLMed
AlSi10Mg
presents
much
higher
yield
strength
(YS,
about
300
MPa
[11–15])
than
that
of
the
casting
one
of
about
170
MPa
[16].
In
order
to
further
optimize
the
microstructure
and
residual
stress,
the
post
heat
treatment
was
always
employed.
According
to
the
work
of
Fousová
et
al.
[17],
when
performing
the
artificial
aging
(AA)
at
294 ◦C,
the
original
Si
network
still
remained
at
the
beginning
30
min
and
then
was
partially
broken
in
2
h
[18].
Takata
et
al.
[13]
found
that
the
YS
decreased
from
about
280
MPa
of
as-fabricated
sample
to
180
MPa
after
stress
relieving
at
300 ◦C
for
2
h.
Iurrioz
et
al.
[19]
reported
that
owing
to
the
lack
of
Si
in
the
matrix
and
Ostwald
ripening
effect,
the
ultimate
tensile
strength
(UTS)
decreased
significantly
with
the
increase
of
solution
temper-
ature
from
450 ◦C
to
550 ◦C.
It
should
be
indicated
that
most
of
the
previous
works
focused
on
the
room
temperature
(RT)
tensile
properties
of
as-SLMed
and
heat
treated
samples.
The
high
specific
power
output
of
combustion
engine
[20]
requires
a
new
generation
of
aluminum
alloy
which
proposes
https://doi.org/10.1016/j.jmst.2020.04.066
1005-0302/©
2020
Published
by
Elsevier
Ltd
on
behalf
of
The
editorial
office
of
Journal
of
Materials
Science
&
Technology.

Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
163
Table
1
Chemical
composition
of
investigated
AlSi10Mg
alloy
powder
(wt%).
Al
Si
Mg
Fe
Mn
Zn
Ti
Cu
O
Balance
9.34
0.34
0.38
0.15
0.052
0.036
<0.03
0.046
Fig.
1.
Micrograph
of
AlSi10Mg
powder
used
in
this
study.
required
properties,
that
is,
high
tensile
strength
at
elevated
temperature.
This
is
a
very
important
application
in
automotive
industries
[21].
Uzan
et
al.
[22]
studied
the
tensile
and
creep
prop-
erties
of
AlSi10Mg
specimens
built
by
SLM
in
the
range
of
25−400
◦C.
They
indicated
the
UTS
of
AlSi10Mg
specimens
continuously
decreases
with
the
increase
of
temperature.
Moreover,
the
creep
behavior
illustrated
that
SLMed
AlSi10Mg
can
be
reckoned
as
an
aluminum
matrix
composite
reinforced
by
sub-micro
Si
particles.
Its
plastic
deformation
during
creep
was
governed
by
dislocation
movement
in
primary
aluminum
grains.
Unfortunately,
to
the
best
of
authors’
knowledge,
mechanical
properties
and
microstructure
of
SLMed
sample
at
elevated
temperature
haven’t
been
in-situ
anal-
ysis
in
the
meantime.
So,
the
margin
in
their
relationship
still
needs
to
be
filled
up
for
the
possible
high-temperature
theory
and
appli-
cation.
In
present
work,
in-situ
investigation
was
performed
to
under-
stand
the
effect
of
the
temperature
on
the
microstructure
evolution
and
tensile
behavior
of
SLMed
AlSi10Mg
alloy.
The
precipitation
and
coarsening
of
Si
particles
as
well
as
orientation
of
␣-Al
matrix
with
the
temperature
were
also
characterized
in
detail
in
order
to
further
determine
its
high-temperature
mechanical
behavior.
2.
Material
and
methods
2.1.
Samples
preparation
Table
1
presents
the
chemical
composition
of
the
AlSi10Mg
powder
with
an
average
particle
size
of
30
␮m
used
in
this
study,
which
is
shown
in
Fig.
1.
Samples
were
fabricated
using
a
commercial
machine
(Solutions
SLM
280)
under
protective
argon
atmosphere
(oxygen
content
inferior
to
0.2
vol.%),
with
the
recip-
rocating
scanning
strategy.
The
substrate
plate
was
preheated
to
150 ◦C
before
and
during
the
SLM
procedure.
The
dimension
of
cube
deposited
samples
is
12
mm
×
12
mm
×
70
mm.
The
SLM
processing
parameters
as
addressed
in
Table
2.
The
cube
samples
were
cut
off
without
any
stress
relief
pro-
cedure.
Dimension
of
the
as-built
specimens
for
in-situ
scanning
electron
microscope
(SEM)
and
in-situ
electron
backscatter
diffrac-
tion
(EBSD)
is
10
mm
×
1
mm
×
5
mm,
as
presented
in
Fig.
2(a).
The
tensile
specimens
of
SLMed
AlSi10Mg
alloy
were
machined
from
cube
samples
according
to
GB/T
4338−2006
standard,
which
Table
2
SLM
processing
parameters.
Laser
power
Scanning
speed
Hatching
distance
Beam
radius
Layer
thickness
340
W
1600
mm/s
100
␮m
70
␮m
30
␮m
is
shown
in
Fig.
2(b).
The
SEM
and
EBSD
samples
were
polished
and
then
etched
by
Keller
reagent
and
ion
beam
respectively.
2.2.
Microstructure
characterization
The
in-situ
microstructure
characterizations
were
conducted
using
Zeiss
Gemini
500
field
emission
scanning
electron
microscopy
with
15
kV
accelerating
voltage,
at
the
temperature
of
100 ◦C,
200 ◦C,
300 ◦C,
400 ◦C,
respectively.
The
SEM
pictures
were
taken
within
30
min
after
reaching
the
testing
temperature.
The
in-
situ
orientation
analyses
were
carried
out
on
the
same
instrument
equipped
with
an
electron
EBSD
detector
using
step
sizes
of
0.53
␮m
at
200 ◦C
and
400 ◦C,
respectively.
The
EBSD
information
was
recorded
with
20
kV
accelerating
voltage
at
RT
and
after
exposure
at
high
temperature
for
about
30
min.
2.3.
Tensile
tests
Tensile
tests
were
performed
with
MTS
DDL200
system
at
the
temperature
of
RT,
100 ◦C,
200 ◦C,
300 ◦C,
400 ◦C,
respectively.
Three
specimens
were
tested
at
least
at
each
testing
temperature.
The
loading
direction
is
parallel
to
the
XY
plane
(vertical
to
the
building
direction
Z),
as
shown
in
Fig.
2.
Prior
to
tensile
testing,
the
tensile
specimens
were
heated
up
to
the
pre-set
temperature
and
held
for
30
min.
The
crosshead
velocity
was
set
as
0.5
mm/min.
3.
Results
and
discussion
Fig.
3
shows
the
microstructure
at
the
inner
and
boundary
of
molten
pool
in
as-built
AlSi10Mg
sample.
As
presented
in
Fig.
3(a),
unlike
the
as-cast
microstructure
of
AlSi10Mg
alloy
where
the
large
rod-like
and
needle-like
Si
particles
(about
150
␮m)
[23]
precipi-
tate
in
the
Al
matrix,
the
fiber
eutectic
structure
with
the
fraction
of
about
50
%
is
formed
in
the
interdendrite
of
␣-Al.
Due
to
the
different
local
thermal
histories
at
the
inner
and
boundary
of
the
molten
pool,
the
microstructure
of
SLMed
AlSi10Mg
parts
is
hetero-
geneous.
This
heterogeneity
is
characterized
by
the
transformation
of
microstructure
from
the
finer
columnar
␣-Al
dendrites
(aver-
age
diameter
of
∼400
nm)
in
the
molten
pool
(see
in
Fig.
3(b))
to
the
coarser
␣-Al
dendrites
(average
diameter
of
∼200
nm)
at
the
molten
pool
boundary
(see
in
Fig.
3(c)).
The
fractions
of
fiber
eutec-
tic
at
inner
and
boundary
of
molten
pool
are
about
52
%
and
43
%,
respectively.
The
microstructural
images
of
the
as-built
and
the
correspond
heated
samples
at
100 ◦C,
200 ◦C,
300 ◦C
are
shown
in
Fig.
4.
The
microstructural
characteristics
from
both
boundary
and
inner
regions
after
holding
at
100 ◦C
and
200 ◦C
for
30
min
(Fig.
4(d,
e,
k,
j))
seem
to
be
the
same
with
as-built
one
(Fig.
4(a,
b,
h,
g)).
This
phenomena
is
similar
to
the
work
in
previous
literatures
[17,24].
As
shown
in
Fig.
4(f)
and
(l),
the
microstructure
changes
obviously
with
prior
Si
particles
coarsening
compared
to
the
as-built
(Fig.
4(e)
and
(i)),
when
the
testing
temperature
increased
to
300 ◦C.
It
is
important
to
note
that
a
few
fine
secondary
Si
particles
precipitate
within
the
columnar
␣-Al
grains.
Fig.
5
presents
the
microstructural
evolution
of
SLMed
AlSi10Mg
alloy
at
high
temperature
of
400 ◦C.
It
can
be
seen
that
the
eutec-
tic
network
coarsens
obviously
after
holding
at
400 ◦C
for
15
min.
At
the
meantime,
Si
particles
uniformly
disperse
in
␣-Al
dendritic
arms
with
a
coarse
spheroidal
shape,
as
shown
in
Fig.
5(b)
and
(d).

164
Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
Fig.
2.
Schematic
illustration
of
SLM
scanning
process
and
specimens
for
(a)
SEM
and
EBSD
microstructural
analyses
and
(b)
tensile
test.
Fig.
3.
Microstructures
of
as-built
AlSi10Mg
samples:
(a)
inner
and
boundary
of
molten
pool
and
high
magnification
images
of
(b)
inner
and
(c)
boundary
regions.

Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
165
Fig.
4.
Microstructure
of
as-built
and
exposed
SLM
AlSi10Mg
specimens:
(a),
(b),
(c),
(g),
(h)
and
(i)
as-built
samples;
(d)
and
(j)
at
100 ◦C
for
30
min;
(e)
and
(k)
at
200 ◦C
for
30
min;
(f)
and
(l)
at
300 ◦C
for
30
min
(AB:
as
built
and
HT:
heat
treated).
With
increasing
the
holding
time
to
30
min,
the
Si
particles
coarsen
and
merge
remarkably
along
higher
volume
fraction.
The
eutectic
network
shows
the
remarkable
ripening,
and
collision,
which
result
in
joining
of
eutectic
Si
fibers,
as
shown
in
Fig.
5(e)
and
(f).
In
addi-
tion,
the
reciprocating
annealing
effect
on
inner
of
molten
pool
is
weaker
than
the
boundary
during
the
deposition
process,
so
the
supersaturation
is
relatively
higher
which
leads
more
Si
particles
precipitated
during
the
heat
process,
as
shown
in
Figs.
3(f),
(l)
and
4
(e),
(f).
The
microstructure
evolution
of
SLMed
AlSi10Mg
from
RT
to
400
◦C
is
schematically
illustrated
in
Fig.
6.
In
the
as-built
specimen,
the
microstructure
is
characterized
as
columnar
supersaturation
␣-Al
dendrites
with
Si
particles
surrounded
by
eutectic
network,
as
shown
in
Fig.
6(a).
When
the
holding
temperature
is
lower
than
300 ◦C,
the
microstructure
changes
slightly,
as
shown
in
Fig.
6(b).
The
Si
particles
emerge
in
␣-Al
matrix
and
prior
Si
particles
become
coarser
when
the
holding
temperature
reaches
300 ◦C,
as
presented
in
Fig.
6(c).
The
eutectic
network
changes
slightly
with
an
exposure
of
30
min,
which
is
quite
different
from
the
broken
network
after
heat
treatment
for
2
h
[13].
The
coarsening
and
precipitation
of
Si
particles
become
more
obvious
after
the
exposure
at
400 ◦C
for
15
min
and
30
min,
as
presented
in
Fig.
6(d)
and
(e),
respectively.
Some
finer
Si
particles
grow
up
quickly
after
holding
at
400 ◦C
for
30
min,
which
leads
themselves
to
contact
with
each
other
and
merge.
In-situ
EBSD
testing
was
proposed
for
investigating
the
ori-
entation
evolution
of
␣-Al
matrix
in
the
AlSi10Mg
deposits
with
increasing
the
holding
temperature,
as
shown
in
Fig.
7.
Comparing
with
as-built
sample
(Fig.
7(a)),
the
orientation
color
map
changes
little
after
holding
at
200 ◦C
for
30
min,
as
presented
in
Fig.
7(b).
After
holding
at
400 ◦C
for
30
min,
there
is
a
significant
precipitation
of
Si
phase
in
␣-Al
grains
represented
by
white
color,
as
shown
in
Fig.
7(d).
In
addition,
the
subtle
orientation
changes
of
␣-Al
grains
only
occur
in
some
region
(indicated
as
circle
I,
II
and
III),
compared
with
the
as-built
shown
in
Fig.
7(c).
It
is
obvious
that
{100}
texture
has
developed
in
as-built
samples,
as
shown
in
Fig.
7(e)
and
(g).
After
the
exposure
at
200 ◦C
and
400 ◦C,
the
preferred
orientation
of
the
deposit
is
almost
the
same
as
as-built,
as
presented
in
Fig.
7(f)
and
(h),
respectively.
It
is
worth
noted
that
the
size
and
preferred
orientation
of
columnar
crystals
change
slightly,
which
means
that
there
should
be
no
recrystallization
in
SLMed
AlSi10Mg
alloy
when
the
holding
temperature
is
lower
than
400 ◦C.
The
Kernel
average
misorientation
(KAM)
diagram
is
often
used
to
characterize
the
density
and
distribution
of
dislocations
in
alloys.
Generally,
grains
with
higher
dislocation
density
possess
larger
dif-
ferences
in
local
dislocation
orientation
[7,25,26].
Fig.
8(a)
and
(c)
presents
the
KAM
maps
of
as-built
AlSi10Mg
samples.
The
KAM
dia-
gram
changes
little
after
exposure
at
200 ◦C
for
30
min,
as
shown
in
Fig.
8(b).
When
the
holding
temperature
increases
to
400 ◦C,
the

166
Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
Fig.
5.
Microstructures
of
as-built
and
exposed
SLM
AlSi10Mg
specimens:
(a)
and
(b)
represent
as-built
sample;
(c)
and
(d)
at
400 ◦C
for
15
min;
(e)
and
(f)
at
400 ◦C
for
30
min
(AB:
as
built
and
HT:
heat
treated).
Table
3
Mechanical
properties
of
SLMed
AlSi10Mg
specimens
at
as-built
and
heated
conditions.
Test
temperature
(◦C)
YS
(MPa)
UTS
(MPa)
Elastic
modulus
(GPa)
Elongation
(%)
25
322
±
10
460
±
7
76
±
4
6.94
±
0.85
100
298
±
7
382
±
12
65
±
2
13.14
±
3.05
200
236
±
8
266
±
19
59
±
1
23.73
±
0.38
300
143
±
6
150
±
9
49
±
3
26.56
±
3.10
400
25
±
330
±
4
18
±
2
75.02
±
4.02
region
with
high
misorientation
angle
indicated
in
red
and
green
color
disappears
after
exposure
at
400 ◦C
(represented
as
region
(I,
II
and
III)
in
Fig.
8(c)
and
(d)).
This
phenomena
can
be
regarded
as
an
evidence
of
the
decrement
of
dislocation
density.
True
and
engineering
stress-strain
curves
at
the
temperatures
from
RT
to
400 ◦C
are
shown
in
Fig.
9(a)
and
(b),
respectively.
Then,
the
effects
of
testing
temperature
on
tensile
strength
(UTS
and
YS)
and
elongation
are
presented
in
Fig.
10
and
summarized
in
Table
3.
Owing
to
ultrafine
microstructure
in
the
deposit,
the
YS
and
UTS
of
SLMed
samples
at
RT
are
usually
higher
than
those
obtained
by
conventional
manufacture
processes,
but
elongation
is
lower.
Overall,
the
UTS
approximately
decreases
linearly
from
460
MPa
to
30
MPa
with
the
increase
in
temperature.
YS
decreases
more
rapidly
at
higher
temperature.
At
the
same
time,
the
yield
ratio
increases
gradually
with
the
increase
of
temperature,
and
tends
to
be
1
when
temperature
exceeds
200 ◦C.
Elongation
shows
the
opposite
trend,
gradually
increasing
below
300 ◦C,
and
increasing
sharply
from
26.56
%
to
75.02
%
beyond
300 ◦C.
The
high
strength
of
SLMed
sample
can
be
attributed
to
mainly
three
aspects:
(I)
grain
boundary
strengthening
governed
by
the
Hall-Petch
relationship
[9]:
y=
0+
kd−1
2(1)
where
yis
the
YS,
0is
the
friction
stress
for
dislocation
move-
ment,
d
is
the
average
matrix
grain
size,
and
k
is
the
strengthening
coefficient
depending
on
the
materials;
(II)
solid
solution
strength-
ening
due
to
the
presence
of
alloying
elements,
and
(III)
dislocation
strengthening
through
interaction
of
dislocations
impeding
each
other’s
motion.
In
the
rapidly
solidified
AlSi10Mg,
precipitation
hardening
is
dominated
by
the
Si
dispersoids
and
precipitates,
referred
to
as
Orowan
strengthening
[9].

Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
167
Fig.
6.
Schematic
illustration
of
microstructure
evolution
of
the
SLMed
AlSi10Mg
samples:
(a)
asbuilt
conditions,
(b)
after
exposure
at
100 ◦C
and
200 ◦C
for
30
min,
(c)
after
exposure
at
300 ◦C
for
30
min,
(d)
after
exposure
at
400 ◦C
for
15
min,
(e)
after
exposure
at
400 ◦C
for
30
min.
Thus,
the
YS
change
of
SLMed
AlSi10Mg
sample
could
be
described
as
the
sum
of
various
contributions:
y=
GBS +
SSS +
dis +
Orowan(+matrix )
(2)
where
yis
the
change
of
YS
from
RT;
GBS ,
SSS,
dis,
Orowan are
the
contribution
from
grain
boundary
strengthen-
ing,
solid
solution
strengthening,
dislocation
strengthening
and
Orowan
looping
mechanism
strengthening,
respectively.
It
is
wor-
thy
noted
that
matrix is
the
heat
induced
softening
of
matrix,
which
cannot
be
neglected
especially
above
specific
temperature
[27].
In
this
work,
the
YS
slowly
decreases
from
314
MPa
to
234
MPa,
when
the
testing
temperature
increases
from
RT
to
200 ◦C.
And
YS
finally
decreases
sharply
from
236
MPa
to
25
MPa
from
200 ◦C
to
400 ◦C,
as
illustrated
in
Fig.
10.
Interestingly,
the
casting
AlSi10Mg
[16]
and
EN
AC-460,000
specimens
[27]
were
following
a
pattern
which
is
different
from
the
SLMed
one,
i.e.
the
YS
of
as-cast
sample
hardly
changes
when
the
testing
temperature
is
lower
than
200
◦C.
In
the
case
of
SLM,
after
exposure
at
200 ◦C
for
30
min,
the
local
misorientation
slightly
changes.
The
grain
size
and
orienta-
tion
are
almost
invariable
after
exposure
at
200 ◦C,
as
shown
in
Fig.
7,
which
indicates
that
recrystallization
did
not
occur.
Com-
pared
with
as-cast
sample,
one
possible
reason
for
this
obvious
reduction
of
strength
is
the
slipping
behavior
of
dislocation.
There
are
high
residual
stress
[28]
and
high
dislocation
density
[29,30]
in
SLMed
sample.
For
instance,
as
the
temperature
increases,
the
resistance
to
dislocation
slip
decreases.
Dislocation
motion
results
in
partial
internal
stresses
release,
which
reduces
YS
[31].
Thus,
it
can
be
concluded
that
the
dominant
contribution
of
Y
S
reduction
might
result
from
dis when
the
temperature
is
lower
than
200
◦C.
At
relative
high
temperature,
the
residual
stress
concentrates
at
particle/matrix
interface
region
combined
with
the
thermal
activa-
tion
promotes
dislocation
annihilation,
cross-slip
and
climb
[27].
On
the
other
hand,
at
high
temperature,
the
softening
of
Al
matrix
increases
significantly
[27].
Recent
works
[32–35]
illustrated
that
the
deformation
behavior
of
metals
at
elevated
temperature
mainly
depends
on
the
work
hardening
effect
and
dynamic
softening.
In
this
work,
softening
begins
to
counteract
the
effect
of
work
hard-
ening
at
300 ◦C,
which
is
identical
to
the
true
stress
decrease
after

168
Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
Fig.
7.
Orientation
color
maps:
(a)
and
(b)
before
and
after
exposure
at
200 ◦C
for
30
min,
respectively;
(c)
and
(d)
before
and
after
exposure
at
400 ◦C
for
30
min,
respectively;
Pole
figures
of
AlSi10Mg
alloy
samples:
(e)
and
(f)
before
and
after
exposure
at
200 ◦C
for
30
min,
respectively;
(h)
and
(i)
before
and
after
exposure
at
400 ◦C
for
30
min,
respectively.
Table
4
Solid
solubility
of
Si
in
Al
matrix
and
contribution
of
solution
strengthening.
Sample
Solid
solubility
[36]
(wt%)
SSS (MPa)
SSS (MPa)
As-fabricated
7.89
86.8
0
300 ◦C
3.02
33.2
−53.6
500 ◦C
1.67
18.4
−68.4
reaching
the
peak,
as
shown
in
Fig.
9(a).
After
exposure
at
relative
higher
temperature
for
30
min,
the
precipitation
and
coarsening
of
Si
phase
can
also
reduce
the
solid
solution
hardening
and
affect
the
Orowan
strengthening
of
SLMed
AlSi10Mg
sample.
According
our
previous
work
[36],
solid
solubility
in
Al
matrix
after
heat
treated
at
300 ◦C
and
500 ◦C
for
30
min,
as
shown
in
Table
4.
The
solid
solution
strengthening
of
Si
can
be
estimated
by
equation
[37]:
SSS =
KSi(w˛
si)m(3)
where
KSi is
11
MPa
wt%−1and
m
is
1.
As
such,
the
calculated
SSS
at
300 ◦C
is-53.6
MPa,
and
SSS at
400 ◦C
can
be
roughly
estimated
as
−60
MPa.
The
contribution
of
YS
from
Orowan
strengthening
can
be
roughly
estimated
by
the
following
equation
[38]:
Orowan,Si =
0.7MGmbfsi1/2
rsi
(4)
where
Gm,
fsi and
rsi are
the
shear
modulus
of
the
matrix,
the
frac-
tion
and
mean
diameter
of
nano-Si
particles,
respectively;
b
is
the
Burgers
vector
(∼0.286
nm)
[39].
ϕ
equals
2
[39]
in
this
study.
Mean
diameter
and
fraction
of
nano-Si
after
heat
treated
for
30
min
and
shear
modulus
of
matrix
are
shown
in
Table
5.
The
shear
mod-
ulus
can
be
approximately
calculated
by
the
elastic
modulus
of
AlSi10Mg
SLMed
sample
tested
in
this
study:
Gm=E
2(1
+
)(5)
where
E
and

are
elastic
modulus
and
Poisson’s
ratio
(∼0.35)
[40]
respectively.
As
such,
the
estimated
orowan at
300 ◦C
and
400
◦C
are
5.2
MPa
and-75.2
MPa
respectively.
The
decrease
of
shear
modulus
at
elevated
temperatures
makes
the
enhancement
[41]
of
Orowan
strengthening
that
should
have
been
greatly
reduced
or
negative.

Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
169
Fig.
8.
KAM
maps
of
AlSi10Mg
samples:
(a)
and
(b)
before
and
after
exposure
at
200 ◦C
for
30
min,
respectively;
(c)
and
(d)
before
and
after
exposure
at
400 ◦C
for
30
min.
Fig.
9.
Tensile
curves
of
as-built
AlSi10Mg
alloy:
(a)
true
stress-strain
curves
and
(b)
engineering
stress-strain
curves.
Fig.
10.
Variation
of
tensile
properties
with
temperature
of
as-built
AlSi10Mg
alloy
[16,22,27]:
(a)
YS
and
UTS
and
(b)
elongation
(SR:
stress
relief
heat
treated).

170
Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
Fig.
11.
Fracture
surfaces
of
AlSi10Mg
tensile
specimens:
(a),
(b)
and
(c)
the
graphs
of
testing
temperature
at
RT,
100 ◦C,
200 ◦C,
respectively;
(d),
(e)
and
(f)
the
amplification
graphs
of
(a),
(b)
and
(c),
respectively.c.
Fig.
12.
Fracture
surfaces
of
AlSi10Mg
tensile
specimens:
(a)
and
(c)
the
graphs
of
testing
temperature
at
300 ◦C;
(b)
and
(d)
the
graphs
of
testing
temperature
at
400 ◦C.
Table
5
Mean
diameter
and
fraction
of
nano-Si
and
contribution
of
Orowan
looping
mechanism.
Sample
Mean
diameter
(nm)
Fraction
(%)
Shear
modulus
(GPa)
Orowan (MPa)
Orowan (MPa)
As-fabricated
46
0.08
28.2
106.9
0
300 ◦C
51
0.26
18.2
112.1
5.2
400 ◦C
83
0.46
6.3
31.7
−75.2

Y.
Cao
et
al.
/
Journal
of
Materials
Science
&
Technology
62
(2021)
162–172
171
It
is
interesting
that
the
YS
of
SLMed
Al-Si
alloys
at
400 ◦C
is
very
close
to
the
as-cast,
as
shown
in
Fig.
10.
According
to
the
work
of
Zamani
et
al.
[27],
it
can
be
concluded
that
the
YS
of
Al-Si
alloy
mainly
depends
on
matrix
above
specific
temperature.
Moreover,
with
the
temperature
above
260 ◦C
and
the
strain
rates
of
0.1
s−1,
dynamic
recrystallization
(DRX)
is
the
main
softening
mechanism
in
6063
alloy
[33,42]
and
AA
6061/B4C
composites
[32].
DRX
occurs
in
Al-Si
alloy
more
easily
at
higher
temperature
and
lower
strain
rates.
Therefore,
the
high
elongation
and
yield
ratio
close
to
1
at
400 ◦C
in
tensile
testing
may
be
caused
by
the
competition
between
dynamic
recrystallization
and
work
hardening.
Based
on
the
aforementioned
results
and
discussion,
neither
the
grain
size
nor
Al-Si
eutectic
network
obviously
change
after
expo-
sure
at
high
temperature.
Thus,
the
dis plays
a
dominant
role
in
the
change
of
mechanical
properties
at
100 ◦C
and
200 ◦C.
The
yat
300 ◦C
and
400 ◦C
can
mainly
be
attributed
to
sum
of
dis,
SSS,
Orowan and
matrix.
The
comparison
between
experi-
mental
results
and
references
[16,22,27]
indicates
that
the
YS
of
Al-Si
alloy
at
400 ◦C
mainly
depends
on
matrix.
Fig.
11
presents
the
fracture
surfaces
of
as-built
AlSi10Mg
ten-
sile
samples
at
temperature
from
RT
to
200 ◦C.
It
is
obviously
found
that
there
are
cleavage
planes
and
elongated
dimples
in
the
frac-
ture
surface
at
RT
(Fig.
11(a)
and
(d))
and
100 ◦C
(Fig.
11(b)
and
(e)),
indicating
a
brittle-ductile
mixed
fractural
mechanism.
At
200
◦C,
the
main
morphology
still
presents
the
cleavage
with
equiaxed
dimples
(Fig.
11(c)
and
(f))
instead
of
elongated
one,
which
shows
that
the
failure
mechanism
gradually
transfers
to
ductile
fracture
with
increasing
temperature,
combined
with
the
tensile
curves
in
Fig.
9.
In
addition,
the
size
of
the
dimples
approximately
matches
the
diameter
of
columnar
␣-Al
dendrites.
Fracture
surfaces
at
testing
temperature
of
300 ◦C
and
400 ◦C
are
shown
in
Fig.
12.
The
lager
and
equiaxed
dimples
are
found
in
the
fracture
surfaces
at
300 ◦C
and
400 ◦C
(seen
in
Fig.
12(a)
and
c)).
In
addition,
there
are
several
small
Si
particles
in
the
dimples,
as
shown
in
Fig.
12(b)
and
(d).
Such
observation
also
reported
by
Dela-
haye
et
al.
[43],
they
pointed
out
that
the
dimples
nucleate
at
the
interface
between
Si
particles
and
Al
matrix.
Since
the
nano-sized
Si
particles
can
hinder
the
motion
of
dislocations,
these
disloca-
tions
will
concentrate
around
the
Si
particles,
resulting
in
stress
concentration
[34].
With
the
stress
accumulation
aggravates,
the
micro-voids
will
nucleate
and
grow
at
the
interface
[43].
The
plas-
tic
deformation
and
neck
of
matrix
between
adjacent
voids,
that
is
the
formation
of
dimples,
make
the
voids
to
connect
with
each
other
and
lead
to
the
appearance
of
micro-cracks
[44].
Due
to
the
strong
work
hardening,
the
size
of
dimples
from
RT
to
200 ◦C
are
significantly
smaller
than
that
at
higher
temperature.
Because
of
the
dynamic
softening
mechanism,
the
micro-cracks
formed
by
the
connection
of
dimples
can
more
easily
extend
and
coalesce,
result-
ing
in
larger
dimples
on
fracture
surface,
as
shown
in
Fig.
12(b)
and
(d).
4.
Conclusions
In
this
work
the
microstructural
evolution
as
well
as
the
ten-
sile
properties
at
elevated
temperature
of
SLMed
AlSi10Mg
were
analyzed.
These
are
some
conclusions
can
be
drawn:
(1)
After
heated
at
100 ◦C
and
200 ◦C
for
30
min,
the
microstruc-
ture
is
almost
invariable.
After
heated
at
300 ◦C
for
30
min,
Si
particles
become
coarser
significantly
and
precipitation
of
Si
phase
from
Al
matrix
occurs,
as
well
as
this
phenomenon
is
more
pronounced
at
400 ◦C.
(2)
The
EBSD
results
show
that
after
exposure
at
200 ◦C
and
400
◦C,
the
size
and
orientation
of
grains
hardly
change.
The
KAM
results
present
that
the
density
of
dislocation
seems
unchanged
after
exposure
to
200 ◦C,
but
decreases
after
exposure
at
400 ◦C.
(3)
The
yield
strength
decreases
with
temperature
rising.
Com-
paring
to
as-cast
Al-Si
alloy,
the
reduction
of
YS
at
100 ◦C
and
200 ◦C
is
attributed
to
the
high
dislocation
density
and
internal
stresses.
As
the
temperature
rises,
strength
reduction
is
mainly
attributed
to
sum
of
dis,
SSS,
Orowan and
matrix.
Moreover,
the
YS
at
400 ◦C
seems
depend
only
on
matrix.
In
addition,
the
high
elongation
and
yield
ratio
close
to
1
at
400 ◦C
in
tensile
testing
may
be
attributed
by
the
competition
between
dynamic
recrystallization
and
work
hardening.
(4)
Fracture
surfaces
combined
with
tensile
curves
illustrate
the
failure
mechanism
transfers
from
brittle-ductile
mixed
fracture
to
ductile
fracture.
And
dynamic
softening
mechanism
results
in
quite
larger
dimples
on
fracture
surfaces
at
elevated
temper-
ature.
Acknowledgements
This
work
was
supported
financially
by
the
National
Key
Research
and
Development
Programme
of
China
(Nos.
2016YFB1100602
and
2016YFB1100100).
Thanks
for
many
useful
discussions
form
Miss
J.L.
Lu
and
Mr.
Z.
Feng
in
our
group.
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