Precipitation Reactions in Age-Hardenable Alloys During Laser Additive Manufacturing
Precipitation Reactions in Age-Hardenable Alloys During Laser Additive Manufacturing
We describe and study the thermal profiles experienced by various agehardenable alloys during laser additive manufacturing (LAM), employing two different manufacturing techniques: selective laser melting and laser metal deposition. Using scanning electron microscopy and atom probe tomography, we reveal at which stages during the manufacturing process desired and undesired precipitation reactions can occur in age-hardenable alloys. Using examples from a maraging steel, a nickel-base superalloy and a scandiumcontaining aluminium alloy, we demonstrate that precipitation can already occur during the production of the powders used as starting material, during the deposition of material (i.e. during solidification and subsequent cooling), during the intrinsic heat treatment effected by LAM (i.e. in the heat affected zones) and, naturally, during an ageing post-heat treatment. These examples demonstrate the importance of understanding and controlling the thermal profile during the entire additive manufacturing cycle of age-hardenable materials including powder synthesis.
ERIC A. JA¨ GLE ,1,4 ZHENDONG SHENG,1,2 LIANG WU,1 LIN LU,1
JEROEN RISSE,3 ANDREAS WEISHEIT,3 and DIERK RAABE1
Many classes of alloys owe their high strength to
the presence of finely dispersed second phase
particles (i.e. phases different from the matrix phase).
Since they form by precipitation phase
transformations, they are called precipitates, and the materials
featuring them, precipitation-strengthened alloys,1
Examples are most Al alloys,2–4 many Ni-based
alloys5–7 and some steels.8–11 During conventional
processing, these materials undergo two subsequent
heat treatments. First, in a homogenization
treatment in the single-phase region of the phase
diagram, i.e. at relatively high temperature, all
elements are brought into solid solution. This is
followed by a rapid quenching, designed to limit or
fully suppress any precipitation during cooling.
Precipitates occurring after quenching from the
solutionized state are called primary precipitates.
Subsequently, the material is annealed at a lower
temperature, where the remaining solutes that are
now in a supersaturated state are allowed to
precipitate. In this age-hardening step, the desired
fine dispersion of particles forms. These particles
are called secondary precipitates and are typically
only up to a few nanometres in size.
During laser additive manufacturing (LAM), the
same alloy experiences a very different, unique
thermal profile as it is subjected to the subsequent
processing steps (see Fig. 1).12 First, the alloy needs
to be rendered into powder form in an atomization
process. For this, the material is melted and the
liquid is atomized via gas or water into small
droplets that cool down quickly and thus assume a
solid powder particle state.13 This cooling from the
melt is very quick and precipitation is normally
suppressed. Therefore, an additional
homogenization heat treatment step is usually deemed
unnecessary. During LAM, the powder is re-heated to
temperatures above the melting temperature.
Heating to and cooling from this temperature are also
very quick, because of the small size of the heat
source (the laser beam), which generates a melt pool
that is small in size relative to the underlying
substrate or the underlying layers, respectively
(‘‘self-quenching effect’’). Again, precipitation is
typically suppressed during this step. Additionally
to the melting and deposition temperature ‘‘spike’’
or ‘‘peak’’, the already-deposited material
experiences additional peaks as more material is deposited
(in an adjacent track, or an overlying layer, i.e. the
adjacent material is in the heat-affected zone (HAZ)
of the melt pool). This series of subsequently
imposed temperature peaks can be referred to as
‘‘intrinsic heat treatment’’. The maximum
temperatures during the peaks are at first quite close to the
melting temperature (and indeed some portions of
every layer are re-melted during the deposition of
the ensuing layer), but the peak temperature drops
quickly from layer to layer. Due to insufficient heat
conduction to the substrate and/or an installed
heating system, the base temperature between the
peaks may also be elevated. Finally, an ageing heat
treatment analogous to conventional processing
may be performed.
During these processing steps, either desired or
undesired precipitation may occur during (at least)
four of the different processing steps (cf. the
numbers in Fig. 1):
During atomization, either if the quenching
rate is not high enough to suppress
precipitation, or when the time spent in the liquid state
is not sufficient to completely dissolve
preexisting, coarse precipitates.
During the actual LAM processing, i.e. during
cooling from the liquid state after deposition.
Again, this might occur if the cooling rate is
not fast enough.
During the intrinsic heat treatment, i.e.
during temperature peaks.
During the regular ageing heat treatment
applied to the final part.
Of the above list, in traditional processing only
possibility number 4 would be considered as desired
option for precipitation strengthening, as the
particles that result from too slow cooling (in steps 1 and
2) are either too coarse, or not dispersed
homogeneously throughout the microstructure (e.g.,
concentrated at grain boundaries), or both. Secondary
precipitation during the intrinsic heat treatment is
an interesting possibility for desired precipitation
that would allow shortening or even completely
avoiding subsequent ageing treatment. Primary
precipitates forming during solidification may be
beneficial for grain refinement.
The control of precipitation during additive
manufacturing is important, because solution heat
treatment after LAM, which might be used to
redissolve unwanted precipitates is not desirable or in
some cases even impossible. Solutionizing increases
the cost and complexity of the entire manufacturing
process and subsequent quenching might induce
warpage of the treated part, spoiling the (near-)
netshape nature of the LAM process. In some alloys
such as the supersaturated Al-Sc-alloy discussed
below, solutionizing cannot be used at all to
In this paper, examples of all four types of
precipitation occurring in LAM-produced alloys
are given, after a brief description of the common
experimental techniques. The alloys under
investigation are a maraging steel, an Al-Sc alloy and a
Laser Additive Manufacturing
Two different LAM processes were employed in
this work, namely selective laser melting (SLM) and
laser metal deposition (LMD).12 In SLM, a thin
powder layer is spread on a substrate plate using a
powder distribution system and selectively melted
using a scanning laser beam with a diameter in the
range of 100 lm to 1000 lm. After each layer, the
build platform is lowered and new powder is
applied. Typical layer heights DL are 20 lm to
150 lm, laser powers PL are in the range of 50 W to
1000 W and scanning speeds, vS, may exceed
2000 mm/s. In LMD, on the other hand, powder is
fed into the melt pool, generated by a laser beam, by
a carrier gas being conducted through a co-axial or
off-axial nozzle. The entire deposition head
consisting of laser optics and powder-feeding nozzle is then
moved relative to the fixed substrate plate creating,
track by track and layer by layer, a bulk volume.
Typical layer heights are 200 lm to 1000 lm, i.e.
much larger than in SLM, but travel speeds are two
orders of magnitude slower than in SLM (tens of
mm/s). The beam diameter can vary in a wide range
(100 lm to several 1000 lm). Depending on the
beam size, the laser power employed in LMD ranges
from a few 100 W to several kW. These different
process parameters obviously strongly influence the
time–temperature profile experienced by the
material. Solidification and cooling rates are generally
lower in LMD as compared to SLM (103–105 versus
104–106 K/s). In Table I, the main process
parameters used to produce the materials analyzed in this
work are compiled. The volume energy density, EV,
introduced by the laser beam is defined as:
SVcoalunmsepeeende,rvgSy, d(menmsitsy,1)EV (J mm 3)
Layer height, DL, (lm)
EV ¼ yHvSDL
Here, the hatch distance, yH, defines the distance
between adjacent scan vectors (tracks).
Specimens were deposited using SLM and LMD
in the shape of cubes with side lengths of 10–
20 mm. They were cut along the build direction
(cross-section view) and metallographically
prepared. Most of the analyses presented in this work
were obtained by using SEMs (scanning electron
microscopes; JEOL 6500F and Zeiss Merlin)
equipped with EDS (energy dispersive spectroscopy)
detectors. For the analysis of fine precipitates,
atom-probe tomography (APT) was employed. For
this method, samples are lifted from polished
crosssections by a FIB (focused ion beam)-liftout
technique (using an FEI Helios NanoLab 600i
dualbeam microscope) and also sharpened by FIB
milling. The samples are measured in a Cameca
LEAP 3000 HR X local-electrode atom probe, using
laser pulsing at a repetition rate of 250 kHz, a pulse
energy of 0.4 nJ and a base temperature of 60 K
(with the exception of the Al-Sc-alloy, for which
voltage pulsing at a pulse fraction of 15% is applied
at a base temperature of 40 K).
Precipitates in the Raw Powder: A Supersaturated Al-Sc-Alloy
Al–Sc alloys have been shown to possess an
outstanding combination of strength, ductility and
corrosion resistance.14 A new class of
Al-Sc-(Mg-Zr)alloys called Scalmalloy has been developed for use
in additive manufacturing.3 In this alloy, the Sc
content is well above the maximum equilibrium
solubility of 0.3wt.%. Therefore, to obtain a
homogeneous supersaturation of Sc in the Al matrix, rapid
solidification (and further, rapid cooling) needs to be
employed. It has been shown that this way, the
entire Sc content can be used in the strengthening
precipitation reaction yielding Al3(Sc,Zr) particles.3
We investigated specimens produced by SLM
from two different batches of Scalmalloy powder.
The first batch had been atomized by a standard gas
atomization technique (ECKA Granules, Fu¨ rth,
Germany) while the second batch was
manufactured using the electrode induction-melting gas
atomization (EIGA) technique (TLS Technik,
Bitterfeld, Germany). In this crucible-free technique,
the bottom part of a rotating bar is melted by
induction heating and the resulting flow of liquid
metal becomes atomized in a nozzle system using
In otherwise identically SLM-produced
specimens, the material made with EIGA-atomized
powder showed particles of 10–50 lm diameter (see
Fig. 2b for an example) distributed homogeneously.
Analysis by SEM–EDS shows that the precipitates
contain approximately 19 at.% Sc, 7 at.% Zr and
74 at.% Al (with traces of Si and Mn), compatible
with the phase Al3(Sc,Zr). These particles are
already visible in the powder material before
processing by LAM (see Fig. 2a). The size of the
particles as well as the absence of the particles in
the material produced by standard gas atomization
suggest that the particles precipitated during the
EIGA process, or, more likely, were already present
in the material prior to atomization and were not
(fully) dissolved during the comparatively short
time in the liquid state during this process.
Using the measured volume fraction of particles
( 0.3%) and their measured composition, the
amount of Sc removed from the matrix by being
bound inside these large particles can be estimated:
0.09 wt.%. This small amount of scandium accounts
for only a moderate drop in the strength of the
material. Fractographic analysis of tensile test
specimens made from this material reveal that the
large precipitates also do not act as crack initiation
sites and hence have only a small deteriorating
impact on mechanical strength and ductility.
Precipitation during material deposition:
Al-Sc alloy and Ni-base superalloy
Microsegregation occurs during solidification of
most alloys (containing elements with partitioning
coefficient not equal to one), provided the cooling
rates are not high enough to completely trap solute
into the growing crystal.15 We observed
microsegregation both in a maraging steel (18Ni-300, see
next example) and in a Ni-base superalloy (Inconel
738LC ). Evidently, even the relatively fast cooling
rates prevailing in the SLM process (estimated by
simulations to be 104–106Ks 1, depending on
process parameters)16 are not sufficient for (complete)
solute trapping. As the remaining liquid phase
becomes strongly enriched in solutes during
solidification, the formation of intermetallic phases may
become energetically favorable.
In samples of Inconel 738LC produced by SLM,
APT measurements revealed the presence of
carbides (TiC), borides (Cr- and Ni-rich) and an
intermetallic phase (Zr-, Si-, Ta-, and Mo-rich), mostly in
the intercellular and interdendritic regions where
the final liquid film solidified (see Fig. 3).
Interestingly, the boride phase and the intermetallic phase
are always co-located, with the boride often forming
an incomplete shell around the intermetallic phase.
Note that Scheil-type thermodynamic simulations
(using the software ThermoCalc in conjunction with
the database TTNI8) do indeed predict both the
formation of carbide (TiC) and boride (TiB2), but
also predict the formation of the coherent c’-phase
(which is the phase that emerges during
agehardening, but is not found in our experiments)
and other intermetallic phases whose composition
does not match the one observed in our
In samples of Scalmalloy produced by SLM,
various precipitate phases are visible both inside
small grains and along their grain boundaries (see
Fig. 4). APT measurements allow the determination
of the chemical composition of the precipitates,
which are compatible with Al3(Sc,Zr)-, Al6Mn-, and
Mg-rich precipitates. Al6Mn precipitation is known
to occur during the ageing of 5083 alloy, which has a
similar composition as Scalmalloy except for the
addition of Sc and Zr; however, the concentration
observed in the center of the Mg-rich precipitate is
significantly higher (>50 at.%) than would be
expected for the commonly observed Al3Mg2-phase
(b-phase).17,18 Instead, the composition is
compatible with the Mg17Al12-phase.
Due to the presence of the precipitates both inside
the grains and on the grain boundaries, the exact
mechanism of precipitation is not clear. They could
emerge in the solid state during cooling after
solidification or they could be the result of
reheating during the deposition of an adjacent track/
layer (see next example) and afterwards stimulate
nucleation during partial re-melting of the layer as
the next one is deposited above.
Precipitation During Intrinsic Heat Treatment: Maraging Steel
A characteristic feature of LAM processes, in
particular the SLM process, are the high cooling
rates experienced by the processed material owing
to the small size of the melt pool and effective heat
conduction through the substrate and already
deposited material. It can therefore be expected
that no precipitation in the bulk material during
cooling after solidification takes place. This can
indeed be shown by analyzing the material using
statistical analysis of APT datasets.
We investigated a maraging steel (18Ni-300/
1.2709/Bo¨hler V720 ). Maraging steels obtain their
high strength and toughness by a microstructure
consisting of martensite and finely dispersed
intermetallic phases (here: Ni3Ti, Ni3Mo and Fe7Mo68)
that form during a precipitation hardening heat
treatment.9,10,19 In Fig. 5a, the radial distribution
function (RDF) of pairs of titanium atoms is shown.
The curves are obtained from APT datasets
measured using conventionally produced material,
material produced by SLM8,20 and produced by
LMD. In the conventionally synthesized material in
the as-produced state, i.e. before precipitation
hardening, the RDF assumes a value close to one for all
Ti–Ti pair distances. This means that it is equally
likely to find a Ti atom in the close vicinity of
another Ti atom as it is finding it further away. Also
in as-produced material made by SLM, there is
practically no deviation from unity visible. In
LMDproduced material (taken from the middle of the
sample), however, there is a clear deviation to
values larger than one at small Ti–Ti distances.
Apparently, in this material, the Ti atoms are likely
to already be found in close vicinity to each other
without precipitation heat treatment. This is an
indication of the onset of precipitation, because the
Ti atoms exhibit high mobility (compared to Mo
atoms) and therefore form Ni3Ti precipitates first.
To prove that this clustering (early stage of
precipitation) does not occur during cooling after
solidification but rather during the intrinsic heat
treatment, i.e. re-heating of the material during the
deposition of additional layers, we also investigated
samples taken from the very top of the specimen. In
the uppermost layer, and in particular in the last
track of this layer, no intrinsic heat treatment takes
place. Indeed, the RDF corresponding to this
situation is again nearly equal to one for all pair
distances (cf. the black, dotted line in Fig. 5a).
Precipitation During Ageing Heat Treatment:
Finally, further precipitation reactions occur
during post-manufacturing heat treatments. Even
though some of the initial solute content may not
be available for precipitation any more, as explained
in the previous examples, the remainder of solute
atoms reacts in the same way as in conventionally
produced materials. A potential difference exists in
the different defect density of LAM- and
conventionally produced material. Conventionally
produced material is solution heat treated before
precipitation and hence in recrystallized state,
which is relatively poor in defects (e.g., dislocations,
dislocation cells, and low-angle grain boundaries).
Residual stresses generated during the LAM
process can be released by plastic deformation,
generating a higher density of dislocations. This presence
of dislocations could influence the nucleation rate
and the spatial distribution of nucleation sites. We
observed precipitation in a LMD-produced and
postheat treated (8 h at 480 C, i.e. until peak hardness
is reached) maraging steel and compared the
morphology, number density and chemical composition
of the emerging particles with the ones observed in
a conventionally produced and identically
heattreated specimen (see Fig. 5b). Despite the different
processing routes, no significant differences
between the samples could be observed (note that
in other regards, e.g., in the presence of retained
and reversed austenite, grain size and morphology,
and texture, significant differences between the
samples do exist).
We studied desired and undesired precipitation
reactions before, during and after laser additive
manufacturing in a variety of alloys that gain their
strength by precipitation hardening. In particular,
precipitates emerging during the following steps
Precipitation of Al3(Sc,Zr) during the produc
tion of starting material powder in a
supersaturated Al-Sc alloy.
Precipitation of carbide, boride and a Zr-rich
intermetallic phase during the LAM process,
i.e. during solidification in a Ni-based
superalloy. Similarly, precipitation of Al3(Sc,Zr),
Al6Mn and Mg17Al12 from a supersaturated
Al-Sc alloy during cooling after deposition was
Early stages of Ni3Ti precipitation (Ti–Ti
clustering) during the intrinsic heat treatment
occurring in LMD of a maraging steel.
Precipitation of Ni3Ti, Ni3Mo and Fe7Mo6
during the ageing heat treatment after LAM
in a maraging steel.
These findings reveal that LAM can be used to
produce supersaturated alloys that show
precipitation during age hardening annealing, as intended.
In addition, LAM even provides the potential to
dispense with the need of a post-heat treatment by
exploiting its intrinsic heat treatment. However,
care has to be taken in the production of the powder
and in the choice of process parameters to make
sure that no undesired precipitation reactions occur
in the process chain. Hence, a detailed knowledge of
the thermal profile experienced by the material
before and during LAM, as obtained, e.g., by process
simulations, is necessary.
Open access funding provided by Max Planck
Society (Max Planck Institute for Iron Research).
The authors would like to thank F. Palm, Airbus
Group Innovations, for providing the SLM-produced
Scalmalloy specimens, J. van Humbeeck, KU
Leuven, for providing the SLM-produced maraging
steel specimens and S. Kleber, Bo¨hler Uddeholm
GmbH for providing the conventionally produced
maraging steel (Bo¨hler V720). The support by M.
Kuzmina, P. Ku¨ rnsteiner and S. Ocylok with
experiments is gratefully acknowledged.
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license, and indicate if changes were made.
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