International Journal of Machine Tools and Manufacture 128 (2018) 1–20
Contents lists available at ScienceDirect
International Journal of Machine Tools and Manufacture
journal homepage: www.elsevier.com/locate/ijmactool
Optimisation of process parameters to address fundamental challenges
during selective laser melting of Ti-6Al-4V: A review
H. Shipley, D. McDonnell, M. Culleton, R. Coull, R. Lupoi, G. O'Donnell, D. Trimble *
Department of Mechanical & Manufacturing Engineering, Trinity College Dublin, Ireland
A R T I C L E I N F O
A B S T R A C T
Keywords:
Additive manufacturing
Selective laser melting
Ti-6Al-4V
Process parameters
Process optimisation
Selective Laser Melting (SLM) is an additive manufacturing (AM) technique which has been heavily investigated
for the processing of Ti-6Al-4V (Ti64) which is used in the biomedical, aerospace and other industries. To date the
SLM processing of this material has been inhibited by the requirement of post processes due to three primary
challenges of martensitic microstructures, undesired porosity and residual stresses which are present in the asbuilt state. This work identifies the state of the art in process optimisation which is being used to confront
these challenges in the as-built state with a view to removing the reliance on post processing. Regarding process
optimisation, maximising part density is the primary goal due to the negative influence of pores on fracture and
fatigue properties. To accomplish this, a high energy input is required which results in high cooling rates during
processing. It is these cooling rates which are instrumental in the microstructural evolution and residual stress
production. Accordingly novel methods have been proposed which aim to maintain the necessary high level of
energy input but control the cooling rates to tailor the microstructure and reduce residual stresses. Research gaps
have been identified pertaining to all three of these challenges when considering mechanical properties of as-built
components. Thus in its current state post processes remain critical, however promising techniques in early stage
development provide encouragement going forward.
1. Introduction
Ti64 is the most widely used titanium alloy, accounting for more than
50% of all titanium usage worldwide [1]. This is due to its good stability
at high operating temperatures, high specific strength and good corrosion
resistance properties. Conventionally, Ti64 components have been
manufactured through processes such as powder metallurgy, forging and
casting which cannot easily produce complex shapes and frequently
result in components with poor mechanical properties [1,2]. However,
the disadvantage with Ti alloys has always been cost in comparison to its
alternatives (Table 1). Furthermore Ti64 has been classified as a difficult
to machine metal, thus the cost of extraction is only a small fraction of the
total cost of a component when fabricated using conventional
manufacturing methods [3,4].
In contrast, additive manufacturing (AM) techniques do not have the
same extent of design constraints that limit conventional processes [5].
AM allows a far greater degree of geometrical freedom and material
flexibility enabling mass customisation of parts. Moreover, remaining
unprocessed powder can be reused which, along with savings in time,
energy and other costs can reduce the cost per part substantially [6,7].
Various AM techniques such as SLM, electron beam melting (EBM),
laser engineered net shaping (LENS) and binder jetting (BJG) have been
developed and possess different characteristics. SLM and EBM can be
defined as powder bed fusion processes whereby a metallic powder bed is
fused using an electron/laser beam [9]. LENS can be defined as a
blown-powder metal printing system whereby parts are created through
injecting metal powder into a molten pool created by a high powered
laser beam [10]. Whilst BJG operates using a completely different principle whereby metallic powder particles are fused using a binding agent
followed by applying thermal energy similar to conventional sintering
mechanisms.
Each process has advantages and disadvantages and the choice of
which to use can be application dependent. One instrumental differentiator is the types of materials that can be processed by each system
[11,12]. DebRoy et al. [13] examined the materials that can be processed
by each of the aforementioned methods and SLM proved to be the most
versatile. Accordingly, SLM machines have become popular in industrial
settings which has led to extensive research regarding the use of SLM for
Ti64 processing, particularly for use in the aerospace and biomedical
fields [14–17].
* Corresponding author.
E-mail address: dtrimble@tcd.ie (D. Trimble).
https://doi.org/10.1016/j.ijmachtools.2018.01.003
Received 12 September 2017; Received in revised form 18 January 2018; Accepted 21 January 2018
Available online 31 January 2018
0890-6955/© 2018 Elsevier Ltd. All rights reserved.
H. Shipley et al.
International Journal of Machine Tools and Manufacture 128 (2018) 1–20
Table 1
The cost of titanium in comparison to alternative materials [8].
Form
Steel ($/pound)
Aluminium ($/pound)
Titanium ($/pound)
Ore
Metal
Ingot
Sheet
0.02
0.1
0.15
0.3–0.6
0.01
1.1
1.15
1.0–5.0
0.22
5.44
9.07
15.0–50.0
As stated, SLM is a powder bed fusion AM technique whereby a part is
built by selectively melting areas of powder layers using a laser beam
(Fig. 1) [18,19]. In detail; upon irradiation the powder material is heated
and melts to form a liquid pool, known as the melt pool, which solidifies
and cools down rapidly. After the cross section of the part is scanned, the
building platform is lowered by a pre-defined distance and a new layer of
powder is deposited. This process is repeated until the part is completed.
Due to the high reactivity of Ti alloys, the process needs to be conducted
under an inert argon atmosphere whilst the part is built on a solid substrate to counteract warping of the material due to build-up of thermal
stresses.
Despite the numerous advantages that SLM processing of Ti64 can
provide, there are a number of ongoing challenges associated with this
technique. Firstly, due to the high cooling rates that occur as a result of
the requirement to maximise part density, the microstructure of as-built
SLM fabricated Ti64 components is composed of acicular α0 martensite
[20–24]. As a result SLM processed Ti64 components tend to have high
tensile strength but poor ductility [25]. Secondly, SLM processed Ti64
can suffer from microstructural defects such as balling and porosity
which can greatly affect the fatigue performance of the part [26]. Finally,
residual stresses, which have been shown to have considerable influence
on crack growth behaviour, can occur in as-built components due to the
high cooling rates and temperature gradients present during processing
[26].
There are a myriad of ways to control each of these challenges individually and collectively with the common theme being control of process parameters (Fig. 2). However, given the difference between SLM
systems currently available there are a large number of process parameters which have been identified. Yadroitsev [27] and Rehme [28]
defined 130 and 157 process parameters respectively which can be categorised as; pre-process, in-process and post-process parameters.
In-process and post-process parameters such as laser power, scanning
speed and stress relief regimes have been heavily studied and their effect
on the microstructure, defects and residual stresses formed during SLM
processing of Ti64 are widely accepted.
Similarly, many studies concerning the effect of pre-process
Fig. 2. Illustration of operating parameters studied for SLM processing.
parameters namely powder characterisation have taken place. Generally
these powder characteristics are categorised into three categories; particle morphology, particle chemistry and particle microstructure. Whilst
it is accepted that these will influence the final part quality, a lack of
understanding regarding the effects of initial particle characteristics on
the properties of SLM components remains. One reason behind this is the
number of powder properties that can be altered when optimising a
powder for any given application [29]. Each individual powder property
contributes to the flowability, packing density, optical penetration depth
and thermal conductivity which effect the properties of the produced
parts. Furthermore, the process is further complicated by the powder
flow which can depend upon the apparatus used. Powders that flow in
one machine have been observed to behave differently in others [30–32].
Due to this lack of understanding, the effect of metal powders on the
microstructure of components is rarely discussed whilst the complete
characterisation of powders through morphology, chemistry and microstructure is almost non-existent [29]. Accordingly, examination of the
interaction between powder characteristics and mechanical properties of
SLM processed Ti64 is beyond the scope of this paper. Rather this paper
will focus on the relationship between in-process and post-process parameters and the associated properties of SLM fabricated Ti64 as presented in the state of the art literature.
Regarding in-process parameters, an equation known as the volumetric energy density (1), which describes the average applied energy
per volume of material is used to examine the effects of process parameters during the SLM processing.
Ev ¼
P
v⋅h⋅t
(1)
Where; P is the laser power, v is the scanning velocity, h is the hatch
distance and t is the powder layer thickness. Therefore considerable
research has been conducted on the influence of these parameters to
optimise the microstructure, process defects or residual stresses for a
variety of materials [19,20,23,33–47]. Furthermore alternative parameters such as powder bed temperature, focal offset distance and
inter-layer time have been studied to optimise mechanical performance
in the as-built condition [22,46,48]. However, at present post process
heat or thermomechanical treatment is considered essential to reduce
residual stresses, close undesired pores and transform the microstructure
from the as-built α0 martensite to a αþβ structure in order to improve
mechanical properties.
Fig. 1. Schematic of typical SLM machine.
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International Journal of Machine Tools and Manufacture 128 (2018) 1–20
[55–57].
Another common defect is the presence of spherical and/or sharp
crack-like pores in the volume of SLM fabricated components (Fig. 4).
Reports on the mechanism of pore formation are limited with some researchers focusing on the quality of the feedstock but most relying on
unsubstantiated assumptions concerning the detailed mechanisms that
occur during melting and solidification in SLM [58,59]. The sharp
crack-like pores have been ascribed to insufficient energy input or balling
whilst the spherical pores are generally reported as a result of gas
entrapment, denudation around the melt pool or as a result of the keyhole
effect [23,24,40,53,60–64].
Porosity is critical for SLM processed Ti64 parts. It has detrimental
effects on fracture properties and exerts the largest influence on fatigue
performance as cracks initiate from internal pores and propagate radially
outwards [36,65]. Leuders et al. [26] examined the fatigue behaviour of
SLM processed Ti64 in the as-built, heat treated and hot isostatic pressed
(HIPed) conditions. Employing HIP treatment reduced the pore size
below the detection limit thus the samples had a theoretical relative
density of 100% whilst the mean density of as-built and annealed samples
was 99.77%. This difference was reflected in the mean fatigue life of the
samples which ranged from 27,000 to 290,000 cycles for the as-built and
heat treated samples whilst none of the HIPed samples failed before
being interrupted at 2 106 cycles.
Similarly Kasperovich et al. [66] examined the fatigue resistance of
SLM processed Ti64 in as-built, annealed and HIP treated conditions.
They noted that the sites with the highest stress concentrations served as
crack initiation sites and the most critical of these were crack like pores
induced by lack of fusion during processing. Their results were similar to
those obtained by Leuders et al. [26]. The mean fatigue life of samples in
the as built condition was 2.3 103 to 5.6 103 cycles. Annealing
samples at 700 and 900 C did not lead to any significant change in
porosity, thus no significant improvement in fatigue life was observed. In
contrast, samples which were HIP treated exhibited significant
improvement with fatigue lives ranging from 1.5 10 to 3 105 cycles
which was comparable to the fatigue life obtained in a reference sample
of wrought Ti64.
Additionally, some HIP treated samples were subjected to polishing to
remove any surface defects which may act as stress raisers. These samples
were tested at various amplitudes during fatigue testing ranging from
200 to 600 MPa. The samples tested in the as built state displayed traces
of breaks which initiated at the rough outer surface and consequently
failed between 1 104 and 1 105 cycles. In contrast the samples that
were polished demonstrated considerably longer fatigue life (>105 cycles), whilst two machined samples tested at 350 MPa remained unbroken following 10 106 cycles. Despite the influence of polishing to the
surface of components, internal pores remain crucial in fatigue behaviour. Accordingly, maximising density is the primary objective when
selecting process parameters for SLM processing of Ti64.
Several authors have reviewed aspects of the SLM process [5,49,50].
Murr et al. [36] and Beese & Carroll [51] compared the microstructure
and mechanical behaviour of SLM processed Ti alloys to conventional
subtractive and other additive manufacturing techniques respectively.
Zhang et al. [52] evaluated the use of SLM processed Ti alloys for
biomedical applications whilst Kasperovich et al. [53] studied the effect
of process parameters on porosity formation in Ti64. Yet these studies
focus their reviews on the three main challenges individually rather than
collectively. However producing parts in the as-built state which have
comparable if not superior mechanical properties to those of subtractive
manufacturing processes is the ultimate goal of the AM community. To
realise this a unified understanding of and approach to addressing these
challenges is required. Thus, the aim of this research is to concurrently
examine the state of the art processes used in SLM fabricating of Ti64
which seek to address the microstructure, undesired porosity and residual stress concerns which are currently present in the literature.
To accomplish this, the manuscript presents the effect of process
parameters on various aspects of SLM processed Ti64 components in a
sequential manner. The first section primarily addresses the porosity
frequently present in SLM, in terms of why it's important and must be
maximised as well as examining methods of quantifying this porosity
from an in-process parameter perspective. Following porosity is the
microstructure section. The primary influence on the microstructure of
SLM processed Ti64 is the cooling rate of the process which is dictated by
the parameters which are selected primarily to maximise the porosity.
Various aspects of the microstructure are discussed concerning how
process parameters effect microstructural evolution and how to change
the common α0 martensitic structure into an equilibrium αþβ structure
for certain applications. Finally residual stresses are considered. Similar
to the microstructure it is the cooling rate and thermal gradients caused
by parameters chosen to maximise the part density which are instrumental in causing the frequently observed internal stresses in SLM
fabricated Ti64 components. The effects of these parameters on residual
stresses is considered as well as examining alternative process parameters
that could be used to remove residual stresses yet maintain maximum
part density.
2. Porosity
Due to the full melting mechanism employed, SLM is prone to melt
pool instability which, along with poorly chosen process parameters, can
result in microstructural defects and porosity [20,26,36,54]. Two main
defect types dominate SLM processing of Ti64, balling and undesired
porosity. The balling phenomenon is a common defect observed during
SLM and causes the deposition of the following layer to be impeded
which in turn leads to bad layer deposition, cracking or even process
failure (Fig. 3). This transpires when molten material does not wet the
underlying substrate well due to high surface tension differences generated as a result of variations in thermal properties across the melt pool
Fig. 3. SEM images of balling observed following processing
of commercially pure Ti at (a) high energy density and (b)
low energy density as adapted from Ref. [20].
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International Journal of Machine Tools and Manufacture 128 (2018) 1–20
Fig. 4. (a) Sharp crack-like pores and (b) spherical pores
commonly observed post SLM processing of Ti64, as adapted
from Ref. [53].
parts with a density of only 33.33 J/mm3.
Within the current state of the art, it is evident that there is variability
in part density obtained given the same energy densities. For example
Fig. 6 demonstrates that for SLM processed Ti64, different levels of
porosity can be obtained given the same energy density [53,66,71].
Accordingly, the validity of using the energy density variable as a means
of process characterisation has recently been questioned. Prashanth et al.
[72] questioned whether the energy density variable properly represents
the effective energy transferred to the powder bed in Al-12Si samples.
They noted that important process parameters such as laser diameter,
hatch style and others are disregarded. Similarly Bertoli et al. [73]
examined the limitations of the energy density as a means of process
characterisation in 316 L stainless steel. They noted that the same energy
density value can be obtained using significantly different process parameters. These process parameters have varying influence on porosity,
thus a comparison by energy density alone can be misleading and
insufficient [24,38,53,60,68,70,74].
2.1. Quantification of density
2.1.1. Energy density
A common method to attempt to quantify the presence of defects
within components is using the energy density variable. Kruth et al. [55],
Hauser [67] and Olakanmi et al. [50] have ascribed the presence of
balling to a high energy density in their studies of various metal powders.
In contrast Gu et al. [20] and Song et al. [68] reported balling at low
energy densities during their studies of commercially pure titanium
(CP-Ti) and Ti64 respectively. Similarly, Kasperovich et al. [53], Han
et al. [69] and Cunningham et al. [70] used energy density to characterise the presence of pores during SLM processing.
Furthermore, the energy density variable has been commonly used as
a means to define a process window for fully dense components. Han
et al. [69] and Kasperovich et al. [53] defined process windows of
120–202 J/mm3 and 83–120 J/mm3 respectively to produce Ti64 components with a density greater than 99.9% (Fig. 5). Moreover, Kasperovich et al. [53] stated an energy density of 117 J/mm3 should be used to
produce fully dense components and this value closely correlated with
that of 120 J/mm3 reported by Attar et al. [19,40]. In contrast, Cunningham et al. [70] reported a far wider process window, observing parts
with densities greater than 99.9% for energy densities ranging from
48.61 to 194.44 J/mm3 whilst Gong et al. [71] fabricated fully dense
2.1.2. Process parameters
Individual process parameters such as; laser power, scanning speed,
powder layer thickness, hatch distance, powder bed temperature and
focal offset distance have been examined regarding their effect on
porosity. Of these; hatch distance which is the distance between the
centre lines of two successive laser scans, is determined to have the least
impact. In their studies on SLM processed Ti64 Kasperovich et al. [53]
and Han et al. [69] observed a variation in porosity of less than 1% for a
450% and 42% increase in hatch distance respectively while all other
process parameters remained constant. However, pore formation in the
boundary regions has been observed when small hatch distances
(60 μm) were used which may be linked to overheating during the laser's return movement [24,53]. Kasperovich et al. [53] also studied the
effect of laser focus and established a low plateau region with low
porosity of less than 0.25% for focus values from 5 to 2 mm after which
a steep incline in porosity towards 1.25% was observed. This corresponds
to the effect of focus offset on porosity during EBM as reported by Gong
et al. [38].
With regards to layer thickness, many commercial machines keep a
constant powder layer thickness thus it remains relatively unexplored
with regards to process optimisation. Xu et al. [22] managed to produce
samples with a density greater than 99.5% for layer thicknesses between
30 and 90 μm. Qiu et al. [58] concluded that the overall porosity and the
size of pores increase continuously with increased layer thickness. As
shown in Fig. 7 there is little effect on porosity for layer thicknesses
between 20 and 40 μm. Furthermore at 60 μm the top surface displayed
open pores thus increasing porosity considerably for layer thicknesses
greater than 60 μm.
It is well accepted that the stability, dimensions and behaviour of the
Fig. 5. Process window defined by Han et al. [69] relating relative part
density and energy density.
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International Journal of Machine Tools and Manufacture 128 (2018) 1–20
Fig. 6. Analysis of literature demonstrating the variance
in porosity for parts processed at the same energy density
[24,62,66,70,71,75].
fabricate Ti64 ELI samples with varying levels of porosity. Similarly Song
et al. [68] used scanning speed and laser power to define a process
window for SLM processed Ti64 (Fig. 8(b)). They determined that the
high energy input in Zone I would yield cracks, Zone II would produce
fully dense components whilst Zone III would result in balling due to melt
pool instability. Whilst similar conclusions are reported by both sets of
authors, the parameters that define their zones are substantially different
(Figs. 8 and 9).
To illustrate the insufficiency of using only the laser power and
scanning speed to define an optimal processing window, we have
collated data from across the literature where authors have reported the
processing parameters and porosity achieved in Ti64 components [22,24,
46,53,62,66,70,75–77]. These data points are illustrated in Fig. 9 and are
overlaid with the process zones reported by Song et al. [68] and Gong
et al. [38,60]. It is evident that processing with low laser power values at
all scanning speeds results in parts with poor density. Scanning speeds
under 200 mm/s appear to be sparsely studied but show discouraging
signs from the available literature. Thus it is between these limits that the
uncertainty appears.
The fully dense zone reported by Song et al. [68] sits inside the
over-heating zone and “marginal parameter” zone reported by Gong et al.
[38,60]. Furthermore comparing it to the wider literature, parts with
varying levels of porosity including some with <99% density are present.
Examining the fully dense zone reported by Gong et al. [38,60] it appears
to be better correlated with the additional literature. Apart from one
outlying data point, the minimum density reported for components
processed in that zone is 99.5% although the vast majority of the points
in that region have a density of less than 99.9%. What neither study
accounted for, but what shows very encouraging results is processing at
higher laser powers. Of the data available, the average density of components processed under 190 W is 97.63% whilst those processed at or
above 190 W is 99.83%. Currently, no process zone has been defined at
these laser powers and more studies are required to fully understand the
density of parts processed in this range.
Fig. 7. Influence of powder layer thickness on porosity as adapted
from Ref. [58].
melt pool determine the extent of porosity. Thus, it can be inferred that
laser power and scanning speed which have the greatest effect on the
melt pool, will therefore have the maximum influence on porosity [24,
53]. Scanning velocities from 100 to 4250 mm/s and laser powers from
40 to 400 W have been examined and their impact on porosity of Ti64
components assessed [20,24,38,53,58,60]. Han et al. [69] achieved
components with density greater than 99% for scan speeds of between
400 and 1100 mm/s. Gong et al. [38] reported less than 1% porosity for
velocities from 600 to 1600 mm/s whilst Qui et al. [65] reported 99.9%
density for scanning velocities up to 2600 mm/s. However, it is incorrect
to consider these parameters independently. The window with which
fully dense parts can be manufactured is a function of the relationship
between scanning velocity and laser power rather then a result of each
individually.
Gong et al. [38,60] composed a process window based on this relationship, from which porosity classifications can be made (Fig. 8 (a)).
They concluded that Zone I parameters would produce fully dense
components. Zone OH parameters should be avoided as the heat produced cannot be conducted away immediately. Zone II and III parameters, which are referred to as “marginal parameters”, can be used to
2.2. Discussion
Many studies use the term “fully dense” without clarifying a numerical value to which this refers. ASTM F3001 regarding additive
manufacturing of Ti64 ELI with powder bed fusion states that components should not contain cracks, defects, discontinuities, foreign material, inclusion, imperfections or porosity detrimental to the usage of the
component. However, there is no specification regarding the level of
density required in the as-built state. Eylon and Froes [78] defined fully
dense as having a density of 99% or greater. However, from the work of
Leuders et al. [26] it is clear that huge variability in fatigue
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International Journal of Machine Tools and Manufacture 128 (2018) 1–20
Fig. 8. Process windows relating scanning speed and
laser power to porosity as defined by (a) Gong et al. [38,
60] and (b) Song et al. [68].
Fig. 9. A meta-analysis assessing the relationship between scanning velocity and laser power as reported in
the literature. The data discovered in the literature is
compared to the process windows as reported by Gong
et al. [38,60] and Song et al. [68].
characteristics is possible between 99 and 100%. Thus, further investigations into the minimum required density for SLM processed Ti64
components are required.
Examining SLM process parameters individually is erroneous. Much
uncertainty surrounds the effect of individual parameters due to their
interdependent relationship. This has led to the adoption of the two
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International Journal of Machine Tools and Manufacture 128 (2018) 1–20
crystalline structure.
methods of combing these process parameters in an attempt to understand their effect on porosity. However, it is clear that neither of the two
commonly used methods for densification characterisation are sufficient
in their current state. Neither energy density nor the relationship between scanning speed and laser power contain enough detail about the
laser-matter interaction in order to attempt to predict density. Whilst
energy density is more detailed than the scanning speed-laser power
relationship, it does not account for process parameters such as laser
type, laser spot size, powder bed temperature or focal offset distance,
each of which can have an effect on the final density in the as-built
condition.
Resulting from the uncertainty regarding (a) minimum density
tolerable and (b) component densities achieved for different operating
parameters, HIP treatment is considered necessary for cyclically loaded
Ti64 components processed by SLM. Whilst HIP treatment has proven
effective and has been reported to close almost all pores in SLM processed
Ti64 thus increasing the fatigue life by almost an order of magnitude, it
adds significant time to the production cycle [26,65,66]. Therefore there
is a desire to produce components in the as-built state with densities that
can provide fatigue characteristics which are similar to if not superior to
those of wrought and conventionally processed Ti64.
EL ¼
Laser Power ðPÞ
Scanning Speed ðvÞ
(2)
Although these studies reported altering the microstructure through
changing process parameters, the fact remains that the concept behind
these techniques lies with controlling the cooling rate. Yet, during SLM
processing of Ti64 the cooling rate is dictated by parameters chosen to
maximise density. Thus, despite attempts to reduce the cooling rate the
microstructure remains martensitic for SLM processed Ti64 in the asbuilt state.
The α0 martensitic microstructure produced during SLM processing of
Ti alloys is contained within elongated prior β grains which grow
epitaxially through successive layer depositions [20–24]. These α0
structures consist of closely spaced interfaces, separating neighbouring
laths along with a high density of dislocations which results in a more
effective barrier against dislocation movement during deformation when
compared to α structures. Thus these α0 structures produce components
with strength and microhardness properties which are frequently greater
than those observed in cast or wrought Ti64 [19,69].
These properties are satisfactory for applications such as components
for aircraft landing gear where a minimum ultimate tensile strength
(UTS) of 855 MPa and yield strength (YS) 758 MPa are required yet the
required ductility sits at only 6% [8]. However, in many applications
such as biomedical implants, martensitic microstructures are undesirable
as they have poor ductility (<10%), favour intergranular failure and
demonstrate significant anisotropic mechanical behaviour [22]. According to ASTM F13-12a regarding the use of Ti64 ELI for surgical
implant applications and ASTM F2924-14 regarding powder bed fusion
of Ti64, the microstructure must contain a mix of the α and β phases and
facilitate a minimum elongation of 10% [82,83]. Thus, it is considered
necessary to transform the α0 microstructure into an equilibrium αþβ
microstructure to prevent anisotropy and poor fatigue performance [48,
84].
3. Microstructure
Resulting from thermomechanical processing Ti alloys can attain an
equiaxed, lamellar or bi modal microstructure, each of which possess
different mechanical characteristics. Optimising in-process and postprocesses enables tuning of the microstructure in SLM processed Ti alloys. The desired microstructure and resulting mechanical properties is
dependent on the desired application of the part. For parts requiring high
strength properties such as aerospace components, a martensitic microstructure may be ideal whilst applications such as biomedical implants
may require increased fatigue performance thus necessitating an equilibrium structure.
Mechanical properties of lamellar microstructures, as frequently
observed during SLM processing, is dependent on the β grain size, α
lamellae thickness, α lamellae size and α colony size which are greatly
affected by cooling rate [79]. Due to the full melting mechanism inherent
to the SLM process along with the process parameters selected to maximise density, the cooling rate during SLM processing of Ti64 is of the
order of 103–108 K/s. A lath-type martensite is observed throughout this
range with finer acicular martensite morphology present for cooling rates
above 105 K/s [10]. Thus, controlling the cooling rate during solidification is the most commonly used method of microstructure control [80].
Various authors have examined the influence of process parameters
on the cooling rate and subsequent microstructure formation during SLM
processing of Ti alloys. Do & Li [81] and Han et al. [69] examined the
effect of laser energy input on the microstructure of Ti64. Both studies
observed martensitic structures at all levels of energy input tested. Do &
Li [81] noted that an increase in the energy density will decrease the
cooling rate and lead to an increased lath size within the martensitic
structure. Han et al. [69] determined that with an increase in energy
density, the width of individual α0 and the spacing between them will
decrease whilst the width of the columnar grain will increase.
Similarly Attar et al. [10] examined the effect of linear energy density
(2) on microstructural formation during SLM processing of CP-Ti and also
observed the formation of a α0 structure. They examined the effect of
thermal cycles on the martensitic formation during SLM processing. They
noted that martensite structures formed during early stages of printing
can continue to grow during subsequent thermal cycles whilst finer
martensite can be observed towards the end of the SLM process. In a
different approach Huang et al. [80] proposed the introduction of electromagnetic vibrations into the SLM process in order to control the
microstructure via the magnetic flux density. Their study concluded that
increasing the magnetic flux density increases induction heating which
can lead to improved grain growth and the formation of a coarse
3.1. Heat treatment
Due to the priority given to achieving maximum density during
process parameter selection the most heavily studied method of
martensite decomposition is post processing heat treatment through
annealing and hot isostatic pressing (HIP) [23,25,26,64,66]. Several
authors have demonstrated that the microstructural transformation that
occurs as a result of annealing is comparable to that obtained by HIP
treatment. Kasperovich et al. [66], Qiu et al. [64] and Leuders et al. [26]
reported both processes transform the as built α0 martensite into αþβ
structures. Concerning these heat treatment processes, the final microstructure will be determined by the relationship between maximum
temperature, residence time and cooling rate [25].
3.1.1. Temperature
Heat treatment of Ti64 can be divided into sub-transus heating in the
αþβ field and super-transus heating in the β field (Fig. 10). As early as
1982 Rosenberg et al. [85] suggested that annealing at temperatures high
in the αþβ field, approximately 70 C below the beta transus temperature, provides an excellent combination of fracture toughness and
ductility in Ti alloys. More recently, Wu et al. [86] used scanning electron
and optical microscopy as well as microhardness testing, to examine heat
treatments ranging from 300 to 1020 C on SLM processed Ti64. Below
600 C minimal change from the as-built structure was observed. Between 750 and 990 C the acicular structure degenerated and the α volume fraction of the platelets decreased as the temperature increased.
Whilst above 1000 C the original prior β grains from the as-built
microstructure transformed into large equiaxed β grains.
Similarly, Vrancken et al. [25] studied the effect of heat treatment on
SLM processed Ti64 ELI and observed an increase in the β fraction with
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the αþβ regime with a slight increase as temperatures increased into the β
regime (Fig. 12). They attributed this lack of improvement in ductility to
process defects which resulted in failure of the component.
Fig. 13 shows a more uniform reaction to a temperature increase
between the studies for yield strength values. In both cases an increase in
temperature has caused a decrease in the yield strength of the sample.
This can be attributed to the increased grain size of the αþβ structures in
comparison to the α0 martensite structures (Figs. 11(d) and 16(b)).
3.1.2. Residence time
Residence time can be described as the duration a sample is held at
the maximum temperature of a heat treatment process. Plaza et al. [91]
examined the effect of sub-transus heat treatment on the microstructure
and mechanical properties of Ti64 (Table 2). They annealed samples to a
variety of temperatures in the αþβ field which were furnace cooled to
760 C and then air cooled. They established that the microstructure of
samples heat treated in the αþβ field consists of α grains, whose size
depends on temperature and duration of treatment. Comparing samples 2
and 4 (Table 2) that were annealed at the same temperature for different
durations, it is evident that increased residence time resulted in increased
grain size and consequently ductility (Fig. 14).
Likewise, Vrancken et al. [25] reported that grain size for both α and β
phase have a tendency to increase as residence time is increased. Due to
the larger range of temperatures studied, they were able to determine
that residence time begins to have a limited effect for higher temperatures in the αþβ field as the high α content below this will hinder the β
grain growth. Fig. 15 shows two samples heat treated at 940 C for 2 and
20 h respectively. It can be observed that there is limited growth in the α
grain size, although the α phase did begin to form into an equiaxed
structure as indicated by the arrows in Fig. 15 (b). However the finer
lamellar structure produced during SLM is expected to increase the time
taken to achieve a fully equiaxed structure drastically when compared to
more traditional processes such as forging.
In contrast, residence time has a greater influence on microstructural
evolution when samples are treated above the β transus. This increased
influence results from the rapid grain growth that takes place above the β
transus temperature. Furthermore, at these temperatures the α colony
size is limited by β grain size meaning that for longer residence times as
the β grain size increases, larger α colonies are possible [25]. Lütjering
Fig. 10. Phase diagram for Ti64 showing the sub-transus αþβ and supertransus β fields which are crucial for microstructural evolution [87].
an increase in temperature (Fig. 11(a) and (b)). Thus α fraction was
decreased from 87% at 780 C to 23% at 950 C forming an αþβ structure. Following sub transus heat treatment prior β grains were easily
observed, however after heat treatment above the β transus the prior β
grains are no longer present indicating extensive grain growth (Fig. 11(c)
and (d)). Similar results were obtained by Sercombe et al. [88], Gil et al.
[89], Sallica-Leva et al. [90] and Vilaro et al. [23] in their studies of SLM
processed Ti alloys.
According to Sallica-Leva et al. [90] the degree of martensite
decomposition will determine the balance between mechanical strength
and ductility in heat treated components. Theoretically, as the martensite
is decomposed into an αþβ structure and the grain size increases as the
temperature is increased, ductility should improve whilst yield strength
and UTS values will decline. Experimental results published by Vrancken
et al. [25] and Sallica-Leva et al. [90] concurred with this theory as they
observed sharp increases in ductility values as the temperature was
increased (Fig. 12). In contrast, Vilaro et al. [23] observed little change in
ductility as the temperature was increased towards the β transus within
Fig. 11. Microstructure of SLM processed Ti64 in the as
built condition (a) following sub transus heat treatment
(b) and (c) and super transus heat treatment (d). Note, the
α phase appearing light and the β phase dark [25].
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Fig. 12. Relationship between heat treatment temperature and ductility as adapted
from Vilaro et al. [23], Sallica-Leva et al.
[90] and Vrancken et al. [25].
Fig. 13. Influence of heat treatment temperature on yield stress as adapted from
Vilaro et al. [23] and Vrancken et al. [25].
be inferred that heat treatment above the β transus temperature should
have a short residence time to prevent excessive α colony growth and
consequently improve mechanical properties.
Table 2
Mechanical properties obtained from various αþβ heat treatment strategies as adapted
from Ref. [91].
Sample
T ( C)
t (h)
YS (MPa)
UTS (MPa)
Elongation (%)
1
2
3
4
5
6
915
930
930
930
945
960
4
1
2
4
4
4
950
901
938
938
931
908
989
964
979
979
969
952
15
14.2
16.1
15.8
15.5
15.5
3.1.3. Cooling rate
Vrancken et al. [25] examined the effect of cooling rate on the heat
treated microstructure of SLM processed Ti64. Due to the high α fraction
for temperatures low in the αþβ field, the influence of the cooling rate on
microstructural evolution is minimal. This is demonstrated by the comparable α needle size in samples cooled by air and furnace cooling as well
as water quenching (Table 3). As the heat treatment temperature is
increased, the β fraction increases and single α grains can grow to a larger
extent. At these temperatures larger needle sizes are obtainable at low
cooling rates such as those produced by furnace cooling (Table 3). Thus,
for temperatures approaching and beyond the β transus, cooling rate
becomes the most important parameter determining the primary α
et al. [92] concluded that α colony size is the most important microstructural parameter in determining mechanical properties. They noted
that decreasing α colony size lead to improved yield stress, ductility,
crack nucleation resistance for both high and low cycle fatigue (HCF &
LCF) and micro-crack propagation resistance properties. Therefore, it can
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Fig. 14. Variance in α grain size resulting from (a) 1 h at
930 C and (b) 4 h at 930 C as adapted from Ref. [91].
Fig. 15. α plate size after sub transus heat treatment for
(a) 2 h and (b) 20 h with the arrows showing where some
α has started to globularise. Note, alpha is lighter and β is
darker [25].
Fig. 16. α0 morphology (as adapted from Ref. [23])
exhibited in SLM processed Ti64 which has been heat
treated followed by water quenching. The final
morphology is dependent on heat treatment temperature
whereby an equiaxed structure (a) and a columnar
structure (b) can be observed for super and sub transus
heating respectively.
gradual decomposition of the α0 martensite into an αþβ structure as the
temperature increased in the αþβ field. However for annealing temperatures around the β transus, α0 needles originating from the β phase are
present resulting in a so called α-Widmanst€atten structure. Thus, given
that a martensitic microstructure must be avoided, water quenching
should be discounted entirely. Therefore, the possible cooling strategies
to optimise mechanical properties are furnace cooling or air cooling
followed by furnace cooling.
Table 3
α needle sizes obtained for various cooling rates as observed by Vrancken et al. [25].
Cooling Method
Water Quenching
Air Cooling
Furnace Cooling
Cooling Rate (K/sec)
1500
500
0.17
Needle Size
850 C
950 C
1.16 0.13 μm
1.22 0.09 μm
1.27 0.13 μm
1.48 0.14 μm
1.57 0.21 μm
2.23 0.12 μm
morphology [92,93].
For high cooling rates such as those in water quenching both
Vrancken et al. [25] and Vilaro et al. [23] observed a new form of α0
martensitic microstructure following heat treatment above the β transus.
They observed a shearing mechanism followed by nonthermal nucleation
which resulted in equiaxed β grains in contrast to the columnar β grains
observed at lower treatment temperatures (Fig. 16). This corresponds to
the findings of Ahmed and Rack [94] in their study of phase transformations during cooling in αþβ Ti alloys. They observed a comparable
transformation in a conventionally processed Ti64 bar whereby the β
phase was transformed into an α0 martensite structure after water
quenching.
Regarding slower cooling techniques, furnace cooling produces a
lamellar αþβ structure following heat treatment in both the αþβ and the
β fields. In contrast, the influence of air cooling appears to be heavily
dependent on the maximum temperature. Vilaro et al. [23] observed a
3.2. In-situ martensite decomposition
In-situ or in-process martensite decomposition is an alternative to
heat treatment to improve mechanical properties of laser fabricated
components in the as-built state. In the past several authors [95–98]
studied the microstructural evolution as well as the possibility of in-situ
martensite decomposition during direct laser fabrication (DLF) processing of Ti64. DLF processing involves focusing a laser beam to melt a
stream of metallic powder deposited by a powder jet, which solidifies to
form a fully dense layer. This processes shares a myriad of similarities
with SLM processing regarding as-built microstructures. Thus it is prudent to consider results published using this technique in the study of
SLM processing.
Kelly and Kampe [97,98] characterised the microstructural evolution
during multi-layer DLF processing of Ti64. They established that the
microstructural transformation through the β transus temperature for a
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single layer will transform the β phase by a diffusionless or
diffusion-controlled process to an α0 or αþβ phase dependent upon the
cooling rate. Similar to SLM they observed a diffusionless transformation
for high cooling rates (>410 C/s) whilst lower cooling rates (<410 C/s)
resulted in diffusion-controlled transformations. Moreover, for successive layer depositions, they established that the microstructure of the nth
layer will be dependent upon the microstructure of the previous layer,
the maximum temperature of the current thermal gradient, the residence
time and the cooling rate.
For example, if the structure of the previous layer is α0 and the current
thermal cycle does not surpass the β transus then the α0 structure will be
maintained as the residence time at elevated temperatures in the αþβ
regime is insufficient to decompose the α’. Correspondingly Crespo and
Vilar [96], in their study of DLF processing, reported that deposition of
new layers generates new thermal cycles and as a consequence previously
deposited material undergoes additional phase transformations. They
established that tempering occurred in previously deposited layers due to
the successive addition of new layers and furthermore that lower idle
times lead to an increase in the workpiece temperature. However neither
of these effects were sufficient to reduce the cooling rate below the
martensite critical cooling rate of 410 C/s. In contrast, lowering the scan
speed sufficiently enabled a diffusion-controlled phase transformation
from the β phase to an αþβ phase as lower scanning speeds lead to longer
laser/matter interaction times and thus lower thermal gradients.
Xu et al. [22] were the first to realise in-situ decomposition in Ti64
using SLM. They achieved a transformation from a martensitic structure
into an ultrafine lamellar αþβ structure by optimising SLM process parameters. In particular focal offset distance (FOD), a crucial process
parameter that controls the amount of energy delivered to the powder
per unit area, was optimised (Fig. 17). They concluded that a complete
transformation from a fully acicular α0 martensite to an ultrafine lamellar
structure is possible with a reduction in FOD from 4 to 0 mm. However,
they noted that optimisation of a single processing variable i.e. FOD is
insufficient for martensite decomposition but rather the proper combination of parameters is required.
Regarding this, it was established that energy densities of 33.74 J/
mm3 served to increase martensite retention whilst energy densities of
50.62 J/mm3 enabled decomposition given the appropriate FOD. They
noted that layer thickness was decisive in the determination of the
resulting microstructure when other variables are kept constant as well as
determining the range with which FOD favours martensite. This can be
explained by the significant influence layer thickness has on the cooling
rate as well as on the repeated thermal cycles experienced by previously
melted layers. Furthermore, they observed that the last few layers of each
printed sample exhibited an α0 martensitic structure due to the lack of
succeeding thermal cycles.
As expected, samples containing a martensitic microstructure, displayed high strength but relatively low ductility (<9%). However, for
samples that contained an ultrafine lamellar αþβ structure as a result of
in-situ martensite decomposition, high yield strength (1106 6 MPa)
and ductility (11.4 0.4%) values were recorded. From Fig. 18 it can be
seen that these values compare favourably against samples containing an
α0 martensite structure found throughout the literature as well as those
post processed by mill annealing.
More recently, Xu et al. [48] hypothesised that the additive nature of
the SLM process can be used to control the thermal profile of the preceding solidified layers. Their results demonstrated that with the proper
combination of processing parameters, significant in-situ decomposition
of α0 martensite to a lamellar (αþβ) structure is possible. They opined that
the inter-layer time and the layer thickness were the most influential
parameters in promoting martensitic decomposition and that by tuning
these parameters a required temperature profile could be achieved to
enable in-situ decomposition. A lamellar (αþβ) microstructure was
observed for an inter-layer time of 1 s whilst a mixed microstructure
containing lamellar (αþβ) along with α0 martensite was observed for an
inter-layer time of 10 s (Fig. 19 (a) & (b)). Thus it can be inferred that an
increase in inter-layer time does not favour martensitic decomposition.
Regarding layer thickness an increase from 60 to 90 μm resulted in a
coarser lamellar (αþβ) microstructure, indicative of more significant
martensite decomposition. Moreover, at the thicker value of 90 μm the α
lath width is significantly thinner. This implies that at larger layer
thicknesses, the inter-layer time will have less influence on martensite
decomposition.
Through manipulation of the inter-layer time, they were able to
manipulate both α and β laths. They concluded that increasing the interlayer time from 1 to 10 s decreased the width of the β laths from
63 32 nm to 12 6 nm. This is comparable to the effect on the α needles which were tuneable to a range of 0.15–0.8 μm. This ability enables
the processing of components with a broad range of mechanical properties in the as-built condition. As can be seen in Fig. 20 varying the α lath
width from 0.25 to 0.52 μm can decrease UTS and yield strength by
approximately 7–9% whilst ductility can increase by 2%.
Increased powder bed temperatures can aid in controlling the cooling
Fig. 18. Tensile behaviour for SLM fabricated Ti64 samples with various
microstructures as adapted from Ref. [16].
Fig. 17. Illustration of focal offset distance as adapted from Ref. [99].
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Fig. 19. Different lamellar microstructures of as-built
SLM processed Ti64 for inter layer times of (a) 1 s and
(b) 10 s as adapted from. Note, β is lighter and α is darker
[48].
Fig. 20. Mechanical properties of three asbuilt αþβ structured Ti64 components reported as a function of α lath width [48].
relationship between yield strength, slip length and α colony size. Yield
strength will decrease as the slip length increases which occurs as α
colony size is increased. Therefore, given that an increase in powder bed
temperature causes α colony sizes to grow, it is inferred that a decrease in
yield strength will occur.
Regarding ductility, a 66.2% increase was observed for a powder bed
temperature increase from 100 to 570 C. However, as temperature was
increased passed that to 670 C a decline of 74.7% was observed
(Fig. 22). The cause of the decrease can be extended from an explanation
given by Qian et al. [100] in their study regarding DLF processing of
Ti64. They determined that α needles grow more rapidly as temperatures
approach the β transus. Indeed, as temperatures increased above 570 C
the α needles increased in size whilst some globular α was also observed.
The presence of these microstructural features indicate that the slip
length had increased which results in lower yield strength and ductility
as was observed during testing (Fig. 22.)
rate and reducing thermal gradients during SLM processing. Using this
principle Ali et al. [46] proposed another method of in-situ martensite
decomposition using variable powder bed temperatures. From examining
the microstructures of the highest (770 C) and lowest (100 C) temperatures used in the study, they established that prior β grains were
present throughout the entire temperature range. However within the
prior β grains, the initial martensitic microstructure began to decompose
as temperature was increased (Fig. 21 (a) – (f)).
Although the martensitic decomposition temperature is considered to
be above 600 C, complete martensite decomposition was observed at a
powder bed temperature of 570 C [89]. At that temperature a basketweave αþβ structure containing α colonies with grain boundary β was
observed. The α needle size and the amount of β between the needles
both increased compared to microstructures produced when fabricated at
lower powder bed temperatures. As powder bed temperature increased
further the α needle size and β volume continued to increase (Fig. 21 (g) –
(l)). Moreover, from powder bed temperatures of 670 C α globularisation initiated leading to increased α colony size which as previously
discussed, has the largest effect on mechanical properties for Ti64 according to Lütjering et al. [92].
Examining the mechanical properties for varying powder bed temperatures, they observed high yield and UTS values, characteristic of SLM
processed Ti64. These values remained relatively consistent for powder
bed temperatures from 100 to 670 C (Fig. 22). Passed this, no yield
strength or ductility measurements were possible due to premature
failure which was ascribed to the larger grain sizes generated for higher
powder bed temperatures. Although no further measurements were
possible, it is expected that an increase in powder bed temperature would
lead to a decrease in yield strength. This theory emanates from the
3.3. Discussion
Two main techniques have been studied to control the microstructure
of SLM fabricated components. Annealing and thermomechanical processing (HIP) involve transforming the microstructure from its α0
martensite structure to an αþβ structure after the part has been printed
whilst in-situ martensite decomposition attempts to produce a component with an αþβ structure.
Regarding heat treatment, some general details can be established.
The most important process parameter is the maximum temperature
achieved. Whilst residence time and cooling rate will both influence the
final microstructure, they themselves are dependent on the temperature
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attained during processing. For example, the difference in needle sizes
resulting from air, furnace and water cooling rises from 8.7% to 33%
for maximum temperatures of 850 C and 950 C respectively [25].
Furthermore, it has been discovered that the residence time only has an
influence on the microstructure for temperatures high in the αþβ regime
and beyond [25,91].
However, it is the combination of these parameters that ensures a
transformation of the α0 microstructure and improvement of mechanical
properties. Given that the tensile properties of SLM produced Ti64 are
superior to those of wrought or cast parts, the main objective of post
processes is the improvement of ductility. Given that, many researchers
believe that the optimum heat treatment cycle involves, annealing/
HIPing to a maximum temperature of 800–900 C for 2 h followed by
furnace cooling [23,25,39,90].
Table 4 highlights the effects of the three in-situ martensite decomposition techniques discussed. In-situ decomposition is a new development in the SLM processing field and to date three different methods
have been proposed by different authors, all of which have been successful [22,46,48]. Although quite different approaches, the basis of
these models remains the same in that they involve optimising the process parameters to influence the cooling rate and/or thermal cycles to
tailor the microstructure in the as-built state. Whilst promising, this area
is in its infancy and no validations of these techniques exist. Furthermore
the durability of these methods across multiple platforms needs to be
examined due to the inherent variability that exists within SLM fabricated components.
4. Residual stress
Residual stress can be defined as stresses that remain inside a body
that is stationary and at equilibrium with its surroundings [101]. As early
as 1993, residual stresses were recognised as one of the major flaws in
metal AM [102]. This holds true for laser based processes which are
known to introduce large amounts of residual stress due to large thermal
gradients inherently present in the process. Unmanaged these stresses
result in deformation, reduced resistance to crack formation, reduced
fatigue performance and anisotropic mechanical behaviour [20,47,77,
103,104]. Although residual stresses are heavily studied concerning parts
processed by similar processes such as multi-pass welding. There are very
few papers in the literature, experimental or numerical, concerning the
residual stresses in components processed by SLM [105].
Mercelis & Kruth [104] outlined the method by which residual
stresses occur during SLM. They proposed a two stage mechanism
including; the temperature gradient mechanism and the cool down phase
(Fig. 23). The temperature gradient mechanism induces residual stress
into the material by way of steep temperature gradients which are
formed due to the rapid heating of the upper surface by the laser beam,
followed by the relatively slow heat conduction through the material. As
the expansion of the heated top layer is restricted by the underlying
material, elastic compressive strains are introduced whilst simultaneously the material strength is reduced due to the temperature rise.
During the cool down phase; the top layers shrink as a result of thermal
contraction. This deformation is restricted by the underlying material
thus tensile stresses are introduced on the outer layers and are balanced
by compressive stresses below [103,104].
Fig. 21. As built microstructures for powder bed temperatures from 100 to
770 C as adapted from Ref. [46]. The development of β particles is highlighted by the red arrows throughout. Martensite is visible in samples (a), (c)
and (e) whilst complete decomposition can be seen for (g), (i) and (k).
Furthermore the size of the α laths and the quantity of grain boundary β can be
seen to increase as the temperature increases. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web
version of this article.)
continued on next column
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Fig. 22. Mechanical properties observed as
a function of varying powder bed temperature as observed by Ali et al. [46].
Table 4
Overview of the effect of in-situ martensite techniques on the as-built microstructure of SLM processed Ti64.
Operating Parameter
Adjustment
Effect on Microstructure
Focal Offset Distance
4 - 0 mm
α’ – lamellar (αþβ)
Ductility
[22]
[48]
Inter-layer Time
10 - 1 s
(αþβ)þ α’ – lamellar (αþβ)
Layer Thickness
60–90 μm
Coarsening of lamellar (αþβ)
100 - 570 C
α’ – αþβ basketweave
670 C
Equiaxed α & increased α colony size
Powder Bed Temperature
Yield Strength
Ref.
[46]
Fig. 23. Two stage mechanism by which residual stresses
occur proposed in Ref. [104].
4.1.1. Scan speed
However models have shown that increasing the scan speed elongates
and lowers the temperature of the melt pool [94,106–110]. Li et al. [108]
conducted a parametric analysis of thermal behaviour during SLM processing of Al6061. Given an increase in scanning velocity from 100 to
400 mm/s they observed a decrease in temperature from 1500 to
1050 C and a thermal gradient decrease from a maximum of 15 to
13.5 C/μm (Fig. 24 (a) & (b)). Similarly given a reduction in scanning
velocity, Vasinota et al. [111] and Manvatkar et al. [112] in their studies
of stainless steel reported a reduction in temperature gradients and
cooling rates respectively.
Predicting residual stresses in as-built components has proven difficult due to highly localised temperatures and rapid temperature cycles
resulting from fluctuating laser power amongst other factors. Thus typically an extensive experimental investigation is required to optimise the
process parameters to minimise residual stress formation in as-built
components. Therefore stress relieving post processes are considered
essential, especially in HCF components, to minimise residual stresses.
4.1. Effect of conventional process parameters
As the primary goal in SLM is always to achieve a high part density it
can be difficult to obtain the influence of individual parameters on residual stress. Furthermore, due to fluctuating laser power during each
scan, melt pool instabilities and dissimilar powder beds; variations in
melt pool size and dimensions are produced. As a result the correlations
between the process parameters and the residual stresses produced are
weak [103].
4.1.2. Laser power
In contrast models have shown that an increase in laser power will
cause an increase in the size of the melt pool and the maximum temperature [94,103,107–109,113]. Li et al. [108] observed an increase in
the size of the melt pool from 64.3 55.8 33.7 to
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Fig. 24. Effect of scanning velocity (a)–(b) and laser power (c)–(d) on temperature and temperature gradient generated during SLM processing of
AlSi10 Mg [108].
209.2 140.4 81.2 μm given an increase in power from 150 to 300 W.
Furthermore the increase in power resulted in an increase in maximum
temperature from 60 to 1800 C and temperature gradient from 10
to 22 C/μm. Similar results were obtained by Loh et al. [109] during
numerical investigations of SLM processing AL6061, they determined an
increase in the laser power from 150 to 300 W or decrease in the scanning
speed from 1400 to 200 mm/s would result in an increase in the melt
pool size.
One such parameter is the effect of downtime between layers which was
investigated by Van Belle et al. [117]. They demonstrated that reducing
the time between layers from 34 to 8 s lead to a more uniform stress
distribution throughout the part. Furthermore, the reduction in downtime also lead to a residual stress reduction of approximately 100 MPa
and 200 MPa for a build height of 5 and 10 mm respectively (Fig. 25).
It is well understood that in SLM, the choice of scan strategy will
affect the build-up of residual stresses. Many authors have reported the
greatest stress is generated parallel to the scanning vector due to the high
thermal gradients generated in comparison to the perpendicular direction [47,77,103,114–116,118]. According to Vrancken [103], the scan
vector length has the maximum influence on residual stress compared to
other process variables, excluding preheating. Parry et al. [47] determined that increasing the scan area size from 1 to 3 mm2 increases the
maximum stresses generated from 189.3 to 305.2 MPa. Similar results
were reported by Gibson et al. [9] who found that increasing the scan
vector length leads to increased residual stress. Accordingly, limiting
scan vectors will reduce the time that passes between the depositions of
two successive tracks. In such circumstances, the heat has not been fully
dissipated and so the second track is deposited on to warm material
which leads to a reduction in thermal gradient [103].
There are a myriad of ways each individual layer can be scanned.
Perhaps the most traditional and the strategy by which newer methods
are compared is the zigzag strategy (Fig. 26 (a)). A frequently employed
strategy by researchers and industry is the island scanning strategy
(Fig. 26 (b)). This approach subdivides the component into smaller areas
which are scanned individually and at random in an attempt to produce a
more even heat distribution [64]. Additionally, within island scanning
each sub section can be regarded as an area for which the scan strategy
can be independently chosen. Vectors of neighbouring islands are regularly scanned perpendicular to one another thus leaving each layer with
tracks scanned in multiple directions. This can be advantageous as there
4.1.3. Hatch distance
Pohl et al. [114] concluded that increasing the hatch spacing from
100 μm to 300 μm reduced the deflection caused by residual stresses by
more than half. They ascribed this to more localised heating when a
lower hatch spacing was used, thus increasing the temperature gradient.
However, their study failed to account for the number of tracks deposited, which is greatly reduced given a 300% increase in hatch spacing
coupled with no consideration for density of the material. Therefore, the
influence of hatch spacing is considered unknown.
4.1.4. Layer thickness
Through models and experimental results it has generally been
established that an increase in layer thickness results in reduced residual
stress. In each case the authors attributed this to decreased cooling rates
that occur as a result of the increased energy input when thicker layers
are utilised [115–117]. For example, Van Belle [117] observed a
decrease in residual stresses from 700 to 200 MPa given an increase in
layer thickness from 20 to 40 μm.
4.2. Alternative parameters
Process parameters aside from those that constitute the energy density equation have been studied with regards to residual stress formation.
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Fig. 25. Stress calculated within SLM processed maraging steel for a build height of (a) 5 mm and (b) 10 mm where the downtime between layers was 8 s for
Support 1 and 34 s for Support 2&3 [117].
similar effect.
Zaeh et al. [116] experimentally measured the effect of scanning
strategy in their study of residual stresses formed during SLM processing
of tool steel. They measured longitudinal and transverse residual stresses
in the horizontal and vertical directions after processing with scanning
vectors in one direction only as well as using the island strategy (Fig. 27).
In each case the maximum stresses occurred for the single scanning direction strategy whilst the minimum stress occurred when the island
strategy was employed. Their findings support the work undertaken by
Kruth et al. [55] which also demonstrated lower stress throughout the
part when an island scanning strategy was employed when compared to
strategies such as the zigzag pattern that employ longer scanning vectors.
Other than the island strategy, fractal and helix scan strategies have
been proposed as alternatives to the more traditional zigzag pattern
(Fig. 26 (c) & (d)). Ma and Bin [119] studied fractal scanning, whereby a
layer of short scan tracks with varying orientation are created. Through
modelling they concluded that the fractal scanning strategy resulted in
only half the residual stress compared to the zigzag pattern. No further
work appears to have been undertaken regarding fractal scanning,
probably due to the complexities in the software implementations which
are outlined in Ref. [120]. The helix scan strategy starts scanning from
the outside of the part and moves inwards or vice-versa. Although this
serves to reduce the scan vector length and alter the vector orientation
within each layer, Nickel et al. [118] observed no difference in the
magnitude of residual stresses when compared to the zigzag pattern.
However, the helix strategy has found use where parts cannot be made
using a zigzag pattern. Qian et al. [121] concluded that the helix scan
strategy is suitable for processing complex models where the curvature
changes a lot and/or where every layer is irregular and inconsistent.
Finally, it is universally accepted that the use of preheating during
SLM reduces residual stress [46,110,111,116,122,123]. Abe et al. [124]
and Aggarangasi et al. [123] suggested another laser could be used to
locally pre-heat the powder. Vora et al. [125] successfully reduced residual stresses in their study of the aluminium alloy AlSi12 by preheating
the powder bed. Similarly Tang et al. [125] discovered that preheating
the powder thus re-heating each layer during EBM eliminated cracks in a
Fig. 26. Illustration of (a) zigzag (b) island (c) fractal and (d) helix
scan strategies.
is no major stress build up in one direction and so the anisotropy of SLM
fabricated components can be reduced. Alternatively the scan pattern can
be rotated, generally by 90 , between subsequent layers to produce a
16
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International Journal of Machine Tools and Manufacture 128 (2018) 1–20
Fig. 27. Longitudinal stresses measured in (a) horizontal and (b) vertical directions of tool steel processed by SLM with different scanning strategies [116].
TiAl alloy. More recently, Ali et al. [46] observed an 88.3% reduction in
residual stress in Ti64 ELI as a result of preheating the powder bed. As can
be seen from Fig. 28 increasing the preheating temperature and leads to a
decrease in residual stresses. This is attributed to the associated reduction
of the temperature gradient.
Table 5
Summary of the influence of alternative parameters on residual stresses in the as-built state.
4.3. Discussion
Although the effects of the conventional process parameters on residual stresses are well accepted, the correlations between residual stress
formation and these parameters values are weak. Thus leading to the
investigation of alternative parameters in an attempt to reduce residual
stress (Table 5). Preheating is considered the most effective method to
control the residual stresses in the as-built state. Secondary is layer
thickness for which an increase will result in decreased residual stresses.
However Vrancken [103] observed inconclusive results with differences
between different batches of parts with the same thickness in a
comprehensive study of Ti64 components. Additionally scan strategies
have been comprehensively studied and provide a myriad of combinations in order to alter the residual stresses in the as-built state. However,
this ties into a wider discussion on design for additive manufacture
Operating Parameter
Adjustment
Downtime between
layers
Scan area
34–8 s
1–3 mm
[47]
Scan vector length
Increase
[9]
Scanning direction
Unidirectional to
Multidirectional
100–470 C
[55]
Preheating powder bed
Residual
Stress
Ref.
[117]
2
[46]
(DFAM) which is not patient to discuss here.
Further confusion surrounding process parameters arises from the
effect of laser power and scanning speed. Vrancken [103] reported that
laser power had the smallest effect on residual stress in comparison to
either laser scan speed or layer thickness. Though the results presented
by Li et al. [108] demonstrates that given a 400% increase in laser power,
temperature is only reduced by 30% whilst a 200% increase in laser
power results in a 3000% increase in temperature. However a coherent
Fig. 28. Effect on residual stress of a preheated powder bed temperature for SLM
processed Ti64 ELI as adapted from
Ref. [46].
17
H. Shipley et al.
International Journal of Machine Tools and Manufacture 128 (2018) 1–20
thermal conductivity of Ti64, it could be expected that the effect of a
preheated powder bed would diminish as a component increases in the
z-direction, however this remains to be investigated. Furthermore, their
study is unique in overcoming these three fundamental problems
simultaneously. Thus, the reliance upon post process treatments will
continue until further studies can validate reliable methods to simultaneously overcome these challenges in the as-built state.
analysis of these parameters in relation to residual stress is difficult to
achieve due to the need to maximise density. For example, the study
conducted by Pohl et al. [114] utilised a hatch distance of 300 μm which
would be unsuitable to maximise density and so little is known about its
effect.
Although the effect of process parameters on residual stresses has
been reported, conclusions are deduced from empirical and theoretical
studies relating to a number of materials including Ti alloys, Al alloys,
stainless-steel and metallic-glass composites amongst others [77,104,
107,108]. However SLM is a thermal process, thus the thermal properties
of the material used will dictate the temperature gradients formed during
processing. For example, the temperature interval over which residual
stresses can form is limited by the melting temperature of the material
and the rate at which strain is developed whilst the material is cooling
which is determined by its thermal expansion. Thus any variation between the material properties of two materials that effect their temperature gradient leads to different residual stress values [103]. Therefore in
order to minimise residual stresses in the as-built state and remove the
need for post-processing, parameter sets must be developed for each
material. However, this is a slow and expensive process which has been
identified as a block in the development of powder-bed AM processes
[126].
Furthermore, additional uncertainty regarding residual stresses
emerges from the method of measurement. Authors have measured residual stresses in components through a variety of methods such as the
hole drilling, layer removal, contour, part deformation and x-ray and
neutron diffraction methods [103,117]. These techniques vary in level of
accuracy, destruction and sample volume which can result in confusion
around results reported.
6. Future research
In theory advantages of SLM processing of Ti64 include; production to
near-net shape, high design freedom for complex geometries, locally
configured part characteristics and reduced lead times until production.
However, disadvantages such as anisotropic mechanical behaviour due
to porosity and microstructure, as well as shrinkage due to residual
stresses must be considered. To date a vast number of studies have taken
place regarding SLM processing of Ti64, yet there are still significant
research gaps which need to be investigated before the full extent of these
advantages can be realised.
1 Prashanth et al. [72] and Bertoli et al. [73] have questioned the energy density metric for aluminium and stainless steel alloys respectively. However, due to its widespread use for process
characterisation and the importance placed upon porosity, due to its
importance for anisotropic mechanical behaviour and fatigue performance for SLM processed Ti64, it is imperative that the validity of
the energy density metric is investigated for Ti64.
2 The emergence of in-situ martensite decomposition into the SLM field
is fascinating and should be explored further. The ability to alter the
microstructure during processing from an α0 to an αþβ microstructure
will reduce or remove the anisotropic mechanical behaviour. Thus
enabling printing of components which have mechanical properties
suitable for their intended application in the as-built state which has
the potential to cut substantial time and thus cost from the production
process. However, in-situ decomposition has merely been proven as a
concept and needs to be further investigated regarding both method
of decomposition as well as the effect of such a method throughout
the part as the geometry changes during printing.
3 Remembering that the properties of as-built components is a function
of the relationship between process parameters and not that of a
single parameter selection, further investigations regarding the optimisation of porosity, microstructure and residual stresses simultaneously should be conducted. As demonstrated by Ali et al. [46],
controlling the cooling rate to enable in-situ decomposition may also
have the effect of reducing residual stresses in fully dense components
and thus optimise the three primary challenges simultaneously.
5. Conclusion
The process parameters utilised during SLM processing have been
heavily researched in relation to the microstructure, defects and residual
stresses produced as a consequence. However, what has largely been
neglected is the fact that these parameters are interdependent. Thus the
properties of as-built components is a function of the relationship between these process parameters and not that of a single parameter
selection.
Regarding microstructural composition, martensite structures exhibit
certain mechanical properties such as high YS and UTS which are far
greater than those produced through conventional manufacturing
methods [19]. Whilst this is advantageous for some applications, many
applications require increased ductility than that observed in the as-built
state of SLM processed Ti64. Therefore the α0 martensite must be
decomposed into an αþβ structure. Regarding this evolution from α0
martensite to αþβ structures, further development of in-situ martensite
decomposition techniques, especially as specimens are scaled up, is
required to remove the need for heat or thermomechanical treatment.
Likewise for microstructural defects a more in depth understanding of
(a) what level of porosity is tolerable and (b) what process parameters
can achieve these levels of porosity, is required. The use of energy density
and limited process zones currently utilised for porosity prediction are
insufficient given the complexity of the process. Furthermore, whilst the
method of residual stress formation is widely accepted the parameters
which lead to these stresses are not well understood. Disagreements
between authors regarding the effect of individual process parameters
are common due to the priority given to maximising density as well as the
dissimilar materials and testing methods which are utilised. Thus
comprehensive studies regarding the reliance of residual stress formation
on process parameters during the SLM processing of Ti64 are required.
Individually; martensite has been decomposed, defects minimised
and residual stresses relieved in the as-built state. Although Ali et al. [46]
recently overcame these three challenges simultaneously during SLM
processing of Ti64, the scalability of their method is questionable due to
the size of the specimens investigated ð10x10x10 mmÞ. Due to the poor
Acknowledgements
This work was funded by Science Foundation Ireland (SFI) through
the Advanced Materials and BioEngineering Research (AMBER) centre
located in Trinity College Dublin and DePuy Synthes located in Ringaskiddy, Co. Cork.
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