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Microstructure of Interpass Rolled Wire + Arc Additive Manufacturing Ti-6Al-4V Components

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Mechanical property anisotropy is one of the issues that are limiting the industrial adoption of additive manufacturing (AM) Ti-6Al-4V components. To improve the deposits’ microstructure, the effect of high-pressure interpass rolling was evaluated, and a flat and a profiled roller were compared. The microstructure was changed from large columnar prior b grains that traversed the component to equiaxed grains that were between 56 and 139 lm in size. The repetitive variation in Widmansta ̈ tten a lamellae size was retained; however, with rolling, the overall size was reduced. A ‘‘fundamental study’’ was used to gain insight into the microstructural changes that occurred due to the combination of deformation and deposition. High-pressure interpass rolling can overcome many of the shortcomings of AM, potentially aiding industrial implementation of the process.
Microstructure of Interpass Rolled Wire + Arc
Additive Manufacturing Ti-6Al-4V Components
Mechanical property anisotropy is one of the issues that are limiting the industrial adoption of
additive manufacturing (AM) Ti-6Al-4V components. To improve the deposits’ microstructure,
the effect of high-pressure interpass rolling was evaluated, and a flat and a profiled roller were
compared. The microstructure was changed from large columnar prior bgrains that traversed
the component to equiaxed grains that were between 56 and 139 lm in size. The repetitive
variation in Widmansta
¨tten alamellae size was retained; however, with rolling, the overall size
was reduced. A ‘‘fundamental study’’ was used to gain insight into the microstructural changes
that occurred due to the combination of deformation and deposition. High-pressure interpass
rolling can overcome many of the shortcomings of AM, potentially aiding industrial
implementation of the process.
DOI: 10.1007/s11661-015-3172-1
ÓThe Author(s) 2015. This article is published with open access at
ADDITIVE manufacturing (AM) is a fabrication
technique in which a structure is created by depositing
successive layers. This approach enables substantial
material savings, compared to subtractive techniques
such as machining: in the aerospace sector, buy-to-fly
ratios (the ratio of the weight of the initial workpiece to
the one of the finished part) can be as high as 30,
in AM, this can be potentially reduced to around 1.5 or
The high cost of titanium production and
machining has been the main motivation for the large
number of investigations into AM methods.
Wire + arc additive manufacturing (WAAM) uses
metal in the form of wire in combination with an arc,
and can be based on either Tungsten inert gas (TIG),
or metal inert gas welding.
The microstruc-
tures of Ti-6Al-4V produced by WAAM can be either
¨tten or martensitic which depends not on the
cooling rate of the first cycle after the material is
deposited, but on the peak temperature and cooling rate
when the subsequent layers are deposited.
The prior b
grains are equiaxed near the substrate; however, further
away epitaxial growth results in large columnar grains
that traverse the deposited layers. These grains grow
opposite to the heat flow and are highly textured, leading
to anisotropic mechanical properties.
High-pressure rolling, originally developed for weld-
is a local mechanical stretching method in which a
load is applied with a moving roller. If the load is
sufficient to compress plastically the bead in the normal
direction, a plastic stretching will occur in the rolling
direction, thus decreasing the longitudinal residual
When applied to steel WAAM structures,
rolling resulted in a reduction of the grain size, due to
the enhanced recrystallization that occurred with the
deposition of the subsequent layer on the plastically
deformed component.
Rolling of Ti-6Al-4V WAAM
structures was initially presented,
and will be
discussed in depth in this paper.
The experiments were performed on a custom-made
rolling rig, equipped with a Lincoln Electric Invertec
V310-T AC/DC TIG power supply. A schematic view
of the setup is shown in Figure 1(a) (the X,Y,andZ
directions are defined in this figure). The parameters
for the pulsed TIG process, which are presented in
Table I, produced a wall width of 6 mm. Aerospace
grade 5 Ti-6Al-4V welding wire was provided by VBC
Group; its chemical composition was taken from the
material certificate and is shown in Table II.
A. Evaluation of Strain and Microstructure
These experiments were used to characterize the
microstructure of WAAM samples that were rolled
between depositing layers. Baseplates which were
405-mm long, 60-mm wide, and 6-mm thick were
clamped by screws along each side of the plate
(Figure 1(b)). Six walls were built: a ‘‘control’’ left in
the as-deposited condition; three samples that used a
‘‘profiled’’ roller with loads of 50, 75, and 100 kN, and
two that used a roller with a flat profile and loads of 50
Senior Lecturer, and STEWART W. WILLIAMS, Professor, are with the
Welding Engineering and Laser Processing Centre, Cranfield University,
Building 46, Bedfordshire MK43 0AL, U.K. Contact e-mail: f.martina@ JONATHAN MEYER, Research Engineer, is with Airbus
Group Innovation, 20A1 Building, New Filton House, Bristol, BS99 7AR,
Manuscript submitted October 22, 2014.
and 75 kN. Rolling loads were chosen on the basis of
previous work on steel,
and considering the higher
strength of Ti-6Al-4V. A sample was produced with a
100 kN rolling load; however, it fractured at the
interface between the wall and the substrate after rolling
the seventh layer.
The ‘‘profiled’’ roller’s shape approximately con-
formed to the profile of the deposit as shown in
Figure 1(c). Both rollers were made of case-hardened
H13 tool steel and a rolling speed of 3 mm s1was used.
Layer deposition and rolling application were alter-
nated, and the part was allowed to cool to room
temperature, before rolling was applied. Deposition
started 20 mm from the end of the baseplate and
stopped 15 mm from the other end (Figure 1(b)), giving
a total wall length of 370 mm. Rolling began and ended
35 mm from the ends of the deposit.
The layer height (LH) from the baseplate was mea-
sured with a digital vernier at three points labeled M1,
M2, and M3 (lhi;j) which are indicated in Figure 1(b),
Deposition direction
Single axis manipulator
TIG torch
Hydraulic cylinder
Load cell
Roller fork
Linear bearing
Roller fork
Section A-A A
Wire feeder
Deposition start
Rolling start
Rolling end
Deposition end
50 mm
roller radius
3.6 mm
30 mm
(b) (c)
Fig. 1—Experimental setup for rolling investigation: (a) schematic of experimental setup, (b) details of base plate including holes for clamping,
and (c) dimensions of the roller.
Table I. Deposition Parameters
Wire feed speed 1.6 m min1
Travel speed 4.5 mm s1
Peak current 150 A
Background current 70 A
Average current 110 A
Pulse duration 0.05 s
Frequency 10 Hz
Gas flow rate 10 L min1
Trailing shield gas flow rate 20 L min1
Electrode to workpiece distance 3.5 mm
Table II. Chemical Composition of Ti-6Al-4V Wire Used in the Experiments
Ti Al V Fe O C N H TOE Y Others
Bal. 6.08 4 0.18 0.16 0.035 0.011 0.0017 <0.2 <0.001 <0.05
before and after rolling; iindicates the layer number,
and jthe point of measurement. The height of each layer
(LHi), and mean LH (LH) were calculated from:
LH ¼1
16 X
The first four layers were excluded from the calculation
of the overall LH due to the thermal effect of the
baseplate which reduced the deposit wall width and
increased the LH.
The average total engineering strains (yand z)
introduced by rolling were calculated in terms of the
fractional difference between the unrolled and rolled
wall widths (for y) and LHs (for z), respectively. The
semi-automatic nature of the manipulating system,
combined with the manual start and stop of the welding
arc, resulted in small inaccuracies in the length of the
deposits. Consequently, the change in the deposits’
length (x) could not be assessed.
Cross-section specimens were extracted from the XY,
XZ and YZ planes, at the mid point of each wall. These
sections were mounted in resin, ground, polished and
etched with a solution of hydrofluoric acid for optical
microscopy imaging. Images taken from the center of
the specimens extracted from the YZ plane were used
for prior grain size measurement, which was done with
the the linear intercept method.
For each image, five
measurements were taken for three directions (0, 45, and
90 deg line orientations) giving a total of 15 measure-
ments per sample.
The thickness of alamellae in the samples produced
with the profiled roller was measured using five scanning
electron microscope (SEM) images according to the
method explained in Tiley et al.
Each image corre-
sponded to a specific location within a generic band, as
Fig. 2—Semi-automatic method to measure the thickness of the alaths, as described in Tiley et al.
(a) shows as an example the original image
taken from location C of the specimen rolled at 75 kN. (b) shows five lines of known length with 0 deg inclination. (c) shows the lines broken
where they intercept aphase boundaries. (d) shows the full set of lines oriented with 0, 45, 90, and 135 deg inclinations.
described in Martina et al.
Prior to the SEM investi-
gation, the samples were re-polished for eight minutes at
100 kg/m2and 80 rpm. SEM images were taken using a
backscatter detector, a voltage of 20 keV, and a spot size
of four (Figure 2(a)). The images were processed with
Adobe Photoshop CS4.
For each image, five lines
were drawn to give five sets of measurements for a
particular direction (Figure 2(b)). The lines were divided
where they intercepted the boundaries of aphase laths
(Figure 2(c)). This was repeated for four directions (0,
45, 90, and 135 deg orientations) giving 20 measure-
ments per location (Figure 2(d)), in accordance with
ASTM E1382,
to take into account the strong
anisotropy of the microstructure. The length of the
intercepts across each line was determined using the
ImageJ software
and its plugin Measure Roi PA.
Finally, the alamellae thickness was determined
as 1=1:5ð1=kÞmean
where kmean is the mean intercept
B. Fundamental Study
The fundamental study was used to understand
separately the factors that influence the microstructural
Table III. Average Engineering Strains (Pct)
Profiled Roller Flat Roller
50 kN 7.9 8.1 4.4 6.8
75 kN 18.2 17.5 15.0 20.3
Fig. 3—Optical microscopy images of cross sections taken from YZ plane from (a) control sample, (b) sample rolled at 50 kN, and (c) sample
rolled at 75 kN with the profiled roller; and (d) sample rolled at 50 kN and (e) sample rolled at 75 kN with the flat roller. Note the difference in
the components’ height and width.
changes, occurring as a result of rolling and deposition.
Baseplates were 250-mm long, 60-mm wide, and 6-mm
thick, and were clamped by screws along each side of the
plate. In this study, four linear walls of 20 layers each
were deposited without any interpass rolling. Subse-
quently, the following treatments were applied: for two
walls, only the last layer was rolled with loads of 50 and
75 kN, respectively; for the other two, only the last layer
was rolled at 50 and 75 kN, respectively, after which an
additional layer was deposited. Only the profiled roller
was used for this investigation. These four walls were
sectioned, and critical points in terms of microstructural
changes were identified. The sample rolled at 75 kN with
a subsequent layer deposited was repeated: before
depositing the 21st layer, two R-type thermocouples
were spot welded into the bottom of /3.2 3mmdeep
holes, 4.5 and 4.9 mm below the top. These holes
corresponded to 5.6 and 6 mm below the expected top of
the 21st layer. Finally, the same experiment was
repeated a further two times to validate the temperature
measurements 6 mm below the surface; one measured
the temperature at four points with R-type thermocou-
ples, and one measured it at four more points with
K-type thermocouples.
A. Strain
Average engineering strains are shown in Table III.
The strains for the two directions were very similar for
the samples rolled with the profiled roller; however, the
strain was greater in the transverse direction (y) for the
sample rolled with the flat roller.
B. Microstructure
The cross-sectional microstructures from the five
samples are shown in Figure 3. All the samples were
slightly narrower near the baseplate due to the different
heat flow in this region.
The control sample demonstrates large, columnar
prior bgrains, which grew epitaxially from the baseplate
toward the top of the sample (Figures 3(a) and 4). The
thickness of these grains ranged from 1 to 3 mm and
could traverse the whole height of the deposit.
In the rolled samples, there was a significant reduction
in the prior bgrain size which decreased with increasing
rolling load. High-magnification optical microscopy
images taken from the three main sectioning planes
(YZ,XZ and XY) are shown in Figures 5through 7.
Prior bgrain size was measured using images from the
YZ cross-sectional plane and is reported in Table IV.
Horizontal bands were observed in the macrostructures
of all the samples, one being produced with each
deposited layer (Figures 3and 8(a)). In addition, colum-
nar prior bgrains were observed in the top ca. 2 mm of all
samples. Furthermore, in the top layers of all samples,
martensite was found, which is evidenced by a character-
istic needle-like microstructure
and is shown in
Figure 8(b) for the sample rolled at 75 kN with the
profiled roller.
The rest of the sample had a Widmansta
microstructure (Figure 8(c)). SEM images taken from
the YZ sectioning plane at locations A, B, C, D, E (see
Figure 8(a)) are shown in Figures 9through 11, for the
control and specimens rolled at 50 and 75 kN, respec-
tively. The thickness of the aphase lamellae, measured
from these images, is plotted in Figure 8(d).
There are two differences that need to be considered:
the difference due to the location within the band; and the
difference due to the rolling load. With respect to the first,
the variation is obvious being well outside the confidence
interval of the mean indicated in Figure 8(d), and the
thickness of the alamellae was larger near the top of the
band (point E) and smaller near the bottom of the band
(point A). With respect to the second, the difference is less,
particularly for locations C, D, and E. To assess this, the p
values comparing the statistical significance of the differ-
ent rolling loads are provided in Table V.
The main difference between the microstructures
produced by the flat and profiled rollers is the size of
the prior bgrains along the sides of the samples. While
the grain size was fairly uniform across the sample with
the profiled roller (Figures 3(b) and (c)), the sample
produced with the flat roller had grains that were
approximately five to ten times bigger along the sides
Fig. 4—Optical microscopy images of control sample cross sections
taken from (a)XZ and (b)XY sectioning planes.
(Figures 3(d) and (e)). This is highlighted in high-mag-
nification optical microscopy images taken from the
edges of the samples, from the XY sectioning plane and
collected in Figure 12.
The microstructures of the samples from the funda-
mental study are shown in Figure 13. Figures 13(b) and
(d) show the samples whose last layer was rolled only;
due to process noise, it is impossible to determine
exactly the extent of the plastic deformation induced by
the rolling step. However, it is possible to identify the
shift in the Zcoordinates of the top band (highlighted
by the red lines), as well as a modest reduction in height.
Apart from the deformation induced by the roller, the
appearance of these samples was relatively unchanged.
The samples that had a subsequent layer deposited
(Figures 13(a) and (e), high-magnification optical micro-
scopy images shown in Figure 14) had a significantly
different microstructure. There were three different
regions which are labeled with (1), (2), and (3) in
Figure 14. (1), closer to the top of the deposit, exhibited
columnar prior bgrains aligned with the Zaxis. These
grains grew from grains within the region (2), located
below, which was characterized by a refined equiaxed
microstructure. Region (3) had long columnar prior b
grains that were identical to those observed in the
control. The microstructure of the aphase within these
regions was also significantly different. Within regions
(1) and (2), the microstructure was predominantly
martensitic, while in region (3), the microstructure was
the Widmansta
¨tten microstructure that is observed in
the bulk material. The boundary between regions (2)
and (3), namely the recrystallization boundary, was
located 4.8 and 5.6 mm from the top surface, respec-
tively, for the 50 and 75 kN specimens with a subsequent
deposited layer (Figures 14(a) and (b)). For the latter,
the peak temperatures at the recrystallization boundary
and at the first horizontal band in the deposited
microstructure are shown in Figure 15. The recrystal-
lization boundary had a peak temperature of 1053 K
(780 °C) and a cooling rate from 973 K to 673 K
(700 °C to 400 °C) of 8.4 K/s; the first band, after
removing the three outliers at 862 K, 1263 K, and
1341 K (589 °C, 990 °C, and 1068 °C) had a peak
temperature of 1013 K (740 °C) and a similar cooling
rate from 973 K to 673 K (700 °C to 400 °C) of 7.4 K/s.
In Figure 15, the temperature distribution of the first
band was based on six-point averages calculated every
five seconds. Error bars were calculated in the same way.
A. Prior bGrains
The grain refinement that is achieved with the
combination of rolling and deposition is a key finding.
Fig. 5—High-magnification optical microscopy images of cross sections taken on the YZ sectioning plane from (a) profiled roller, 50 kN; (b)
profiled roller, 75 kN; (c) flat roller, 50 kN; and (d) flat roller, 75 kN. Prior bgrain boundaries are highlighted for reader’s convenience.
Understanding the mechanism that causes this is com-
plex: the final microstructure is a consequence of the
combination of deformation from rolling and the
subsequent heat treatment from the deposition process.
To the authors’ knowledge, this kind of microstructural
change in titanium (static recrystallization) has not been
reported elsewhere.
Investigations of static recrystallization on titanium
are mostly limited to commercially pure titanium
where the kinetics of the transformation are not
complicated by the presence of a second phase. This
is an example of ‘‘classic’’ recrystallization where
rolling introduces stacking faults, points defects, dis-
locations, and twins which provide the driving force
for recrystallization when the material is heated above
the recrystallization temperature. The recrystallization
temperature in commercially pure titanium is 921 K to
942 K (646 °C to 669 °C), well below the btransus
temperature of 1155 K (882 °C).
There is a single
investigation on static recrystallization of Ti-6Al-4V
which investigates recrystallization up to temperatures
of 1153 K (880 °C).
There is considerably more
work on dynamic recrystallization of Ti-6Al-4V which
occurs during hot deformation. In Seshacharyulu
et al.,
recrystallization of the bphase was observed
at 1373 K (1100 °C) and relatively low strain-rates of
0.01 s1.
In the fundamental study, the temperature at the
recrystallization boundary (5.6 mm from the top, see
Figure 15) was 1053 K (780 °C), below the b-transus
temperature and the bphase recrystallization tempera-
ture mentioned in Seshacharyulu et al.
Since errors in
this measurement can occur due to precise placement of
the thermocouple which is exacerbated by the steep
thermal gradient, the measurement was repeated nine
times to improve confidence in this finding. The signif-
icant influence of the rolling load on the position of the
recrystallization boundary (see Figures 14(a) vs (b))
suggests that it is dependent on the amount of strain
rather than the peak temperature. Although they
modeled rolling of friction stir welds, Colegrove et al.
showed that the higher the rolling load, the deeper the
deformation, which appears to strongly influence the
recrystallization of the material.
A higher rolling load also introduced more stored
energy. This explains why for the same deposition
parameters (which corresponded to identical heat treat-
ment conditions) the 75 kN specimens underwent more
recrystallization, and exhibited smaller prior bgrains.
Finally, the results with the flat roller demonstrated a
non-uniform refinement of the prior bgrain size, which
is larger along the side of the walls. As seen in
Figures 3(d) and (e), the flat roller causes a significant
amount of flattening of the deposited material.
Fig. 6—High-magnification optical microscopy images of cross sections taken on the XZ sectioning plane from (a) profiled roller, 50 kN; (b)
profiled roller, 75 kN; (c) flat roller, 50 kN; and (d) flat roller, 75 kN. Prior bgrain boundaries are highlighted for reader’s convenience.
Therefore, the strain induced in the material is likely to
be concentrated around the center of the deposit. Since
there is a link between the strain in the material and the
subsequent prior bgrain size, which is evidenced by the
correlation between the rolling load and prior bgrain
size in Table IV, the concentration of strain in the center
of the deposit causes the greater grain refinement in this
B. aPhase and Banding
The temperature of the point that is associated with
the top band in the microstructure (6 mm from the top
surface) was 1013 K (740 °C). Within the limits of
experimental error, this point appears to be related more
to the adissolution temperature of 1021 K (748 °C),
than to the b-transus temperature as previously sug-
Please note the precise value for these
transition temperatures can vary according to heating
rates and conditions.
The gradient in the alamellae size shown in
Figure 8(d), already observed in plasma
and laser
deposited Ti-6Al-4V structures, is due to the differences
in peak temperatures and cooling rates observed in
different Zcoordinates during the deposition of a layer.
This concept was presented in Martina et al.
and will
be refined further here. As shown in Figure 8(e), a
typical thermal cycle may be simplified into three main
sections: heating (stage 1), a period tPwhen the material
is around the peak temperature TP(stage 2), and cooling
(stage 3). It was suggested there is little difference in
cooling rate in the top 4 to 6 mm of the deposited
structure, i.e., irrespective of the peak temperature and
distance from the heat source, the average cooling rate
on the tail (stage 3) is similar.
This is confirmed by the
experimental measurements shown in Figure 15, where
Table IV. Prior bGrain Sizes
50 kN 75 kN
Average (lm) SD Average ( lm) SD
Profiled roller 125 29.5 89 12.2
Flat roller 139 20.8 56 6.4
Fig. 7—High-magnification optical microscopy images of cross sections taken on the XY sectioning plane from (a) profiled roller, 50 kN; (b)
profiled roller, 75 kN; (c) flat roller, 50 kN; and (d) flat roller, 75 kN. Prior bgrain boundaries are highlighted for reader’s convenience.
Martensite +
Band iBand i
Band i+1
New deposited material
(new martensite)
Martensite +
Coarsened Widmanstätten
(pattern created)
Further coarsening
(pattern retained)
Layer i: Layer i+1:
3 mm
Repetitive α size
5 μm
(1) (1)
tP, stage 2
Top of the deposit
Band i+1
Band i
Stage 1
Stage 3
(b) (c)
(d) (e)
Fig. 8—(a) Microstructure observed before and after the deposition of a new layer, and their locations within the components; (b) optical micro-
scopy image of martensite observed in the top of the sample rolled at 75 kN; (c) scanning electron microscope (SEM) image of the Wid-
¨tten microstructure observed in the rest of the samples; (d) plot of the athickness vs the locations showed in (a) (error bars indicate
95 pct confidence interval of the mean); (e) cooling curves for top of deposit, first band, and area immediately below.
Fig. 9—Scanning electron microscope (SEM) images taken from the control sample to show the increase in the thickness of the aphase. (a) Ta-
ken from Location A; (b) taken from Location B, (c) taken from Location C, (d) taken from Location D, and (e) taken from Location E. The
locations of A to E within each band are shown in Fig. 8(a).
Fig. 10—Scanning electron microscope (SEM) images taken from the sample rolled at 50 kN with the profiled roller, to show the increase in the
thickness of the aphase. (a) Taken from Location A; (b) taken from Location B, (c) taken from Location C, (d) taken from Location D, and (e)
taken from Location E. The locations of A to E within each band are shown in Fig. 8(a).
Fig. 11—Scanning electron microscope (SEM) images taken from the sample rolled at 75 kN with the profiled roller, to show the increase in the
thickness of the aphase. (a) Taken from Location A; (b) taken from Location B, (c) taken from Location C, (d) taken from Location D, and (e)
taken from Location E. The locations of A to E within each band are shown in Fig. 8(a).
the cooling rate at the two locations (7.4 to 8.4 K/s) was
sufficient to produce at least a partially martensitic
Upon deposition of a new layer iþ1 which creates
the band iþ1 (see Figure 8(a)), martensite is produced
in the newly deposited material. Immediately below, the
material deposited during layer iis taken above the b
transus temperature and experiences a cooling rate
sufficient to produce martensite. Closer to, but above the
band iþ1 there is sufficient temperature and time
during stage 2 to transform the martensitic ainto very
fine Widmansta
¨tten. In fact, during annealing, marten-
site decomposes to aþbeither by bparticles at
dislocations, or blayers between aboundaries.
the area between the bands iand iþ1, coarsening of the
very fine Widmansta
¨tten produced during the deposition
of layer ioccurs, and the pattern shown in Figures 8(d)
through 11 is generated. Below band i, no significant
microstructural changes occur, apart from a possible
coarsening of alamellae which retain the repetitive
The addition of deformation introduced a small but
measurable reduction in the overall size of the alamellae
(Figure 8(d)). There are three reasons for this behavior.
Firstly, it could be a consequence of the larger number
of grain boundaries. In the case of basketweave or
Table V. PValues of Student’s TTests Performed on the Size Distribution of the Measured aLamellae Thickness of a Specific
Location Against the Same Corresponding to Another Sample
Control vs 50 kN <0.001 <0.05 <0.05 <0.001 <0.001
Control vs 75 kN <0.001 <0.001 <0.001 <0.001 <0.05
50 kN vs 75 kN <0.001 <0.001 0.14 0.74 0.14
Fig. 12—High-magnification optical microscopy images of cross sections taken on the XY plane from (a) profiled roller, 50 kN; (b) profiled roll-
er, 75 kN; (c) flat roller, 50 kN; and (d) flat roller, 75 kN. These images are taken on the edge of the samples to highlight the different grain size
produced by the two types of rollers for the same load. Prior bgrain boundaries are highlighted for reader’s convenience.
¨tten microstructures, acolonies nucleate at
the grain boundaries and grow perpendicular to
When acolonies collide with each other, and
therefore, cannot grow further, additional colonies start
nucleating perpendicular to the side of those which have
already developed. In the case of the rolled samples,
there were more aphase nucleation sites due to the
larger number of prior bgrains. Therefore, lamellae
growth is hindered by their own competition, hence the
overall size of the lamellae decreases with the increasing
rolling load. However, there is a large difference in the
scale of the prior bvs the alamellae thickness (100 lmvs
<1lm, respectively) so the influence of this may be
limited. An alternative explanation is that deformation
influences the transformation of martensitic ato Wid-
¨tten providing more nucleation sites for the trans-
formation which leads to a finer microstructure. Finally,
recrystallization happened at temperatures comparable to
those used to achieve a fully equiaxed microstructure
during plates production, in which case the aphase
equilibrium volume fraction is large enough to stimulate
the growth of aphase from the deformed lamellae.
At the top of the deposit, there was a region with
columnar prior bgrains which extended approximately
2 mm (see optical microscopy images in Figures 3and
14). Although not measured, the boundary with the
equiaxed material underneath may correspond closely to
the region that was molten during the deposition of the
last layer; the solidifying grains grow epitaxially from
the material underneath.
C. Practical Aspects
High-pressure interpass rolling was performed after
the part was allowed to cool down to room temperature.
The time taken for the part to cool down and the rolling
process itself will affect the productivity of WAAM.
However, depositing large parts will extend the cooling
time between layers lessening its impact on productivity.
In addition, various techniques may be considered for
improving process productivity including depositing
multiple parts at once, using multiple processing heads,
i.e., one for deposition and another for rolling, and
cryogenic cooling. Therefore, there are a number of
(a) (b) (c) (d) (e)
Fig. 13—Optical microscopy images from the fundamental study samples rolled with loads of (a) 75 kN with subsequent deposited layer, (b)75
kN, (c) control, (d) 50 kN, and (e) 50 kN with a subsequent deposited layer.
options to efficiently implement interpass rolling on a
production system.
Finally, with regard to the flat roller, this solution is
of particular interest because it has a number of
practical advantages over the profiled roller: firstly, the
shape of the roller is independent of the deposit
geometry so a single roller could be used for a variety
of wall widths; secondly, most real WAAM parts have
intersecting features which cannot be rolled easily with
the profiled roller.
1. Rolling induced significant prior bgrain refinement,
a reduction in the overall thickness of aphase
lamellae, and a modification of the microstructure
from strongly columnar to equiaxed. This is due to
the recrystallization that occurred when the previ-
ously deformed layer was heated during the depo-
sition of the next layer.
2. A fundamental study was used to understand the
microstructural changes that occurred during the
process. The last layer of an ‘‘unrolled’’ control sample
was rolled after which a new layer was deposited. The
size of the recrystallized region was influenced by the
load and hence the extent of deformation in the
material. In addition, the temperature of the recrystal-
lized region boundary and first microstructural band
were 1053 K and 1013 K (780 °Cand740°C),
respectively, below the btransus and very close to the
adissolution temperature of 1021 K (748 °C).
3. The flat roller, which has significant practical
advantages, provided similar reductions in prior b
grain size and may be the preferred choice for
commercial exploitation of the process.
The authors would like to thank Mr. Flemming
Nielsen and Mr. Brian Brooks for their help during the
experimental work. Mr. Andrew Dyer’s and Dr. Xian-
wei Liu’s assistance during the optical and scanning
(a) (b)
Fig. 14—Optical microscopy images from the fundamental study samples rolled with loads of (a) 50 kN with subsequent deposited layer, and (b)
75 kN with subsequent deposited layer. (1) indicates the top area with columnar grains, (2) the recrystallized area, and (3) the bulk of the mate-
rial which did not undergo any recrystallization.
0 102030405060
Temperature (K)
Time (s)
First bandRecrystallization boundary
Fig. 15—Thermal history at 5.6 mm (recrystallization boundary) and
6.0 mm (first band) below the top of the sample rolled with 75 kN
with a subsequent deposited layer.
electron microscopy images analysis was greatly appre-
ciated. Jack Donoghue provided valuable assistance for
the grain size analysis. The financial support from the
Engineering and Physical Sciences Research Council
under Grant No. EP/K029010/1 and Airbus Group
Innovations is acknowledged. Enquiries for
access to the data referred to in this article should be
directed to
This article is distributed under the terms of the Crea-
tive Commons Attribution 4.0 International License
(, which per-
mits unrestricted use, distribution, and reproduction in
any medium, provided you give appropriate credit to the
original author(s) and the source, provide a link to the
Creative Commons license, and indicate if changes were
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... Although WAAM has many advantages, it also has certain defects, such as oxidation, delamination, high residual stress, deformation, cracking, porosity, and surface finish, etc. [5,[7][8][9]. To improve forming quality and mechanical properties, a wide range of ancillary processes were proposed by researchers including heat treatment [10][11][12][13], inter-pass cold rolling [14][15][16][17], inter-pass cooling [18], molten pool oscillation [19], and ultrasonic vibration. ...
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Wire and arc additive manufacturing (WAAM) is a competitive technique, which enables the fabrication of medium and large metallic components. However, due to the presence of coarse columnar grains in the additively manufactured parts, the resultant mechanical properties will be reduced, which limits the application of WAAM processes in the engineering fields. Grain refinement and improved mechanical properties can be achieved by introducing ultrasonic vibration. Herein, we applied ultrasonic vibration to the WAAM process and investigated the effects of wire feed speed, welding speed, and ultrasonic amplitude on the weld formation and grain size during ultrasonic vibration. Finally, a regression model between the average grain size and wire feed speed, welding speed, and ultrasonic amplitude was established. The results showed that due to the difference in heat input and cladding amount, wire feed speed, welding speed, and ultrasonic amplitude have a significant influence on the weld width and reinforcement. Excessive ultrasonic amplitude could cause the weld to crack during spreading. The average grain size increased with increasing wire feed speed and decreasing welding speed. With increasing ultrasonic amplitude, the average grain size exhibited a trend of decreasing first and then increasing. This would be helpful to manufacture parts of the required grain size in ultrasonic vibration-assisted WAAM fields.
... With the welding speed increasing, the anisotropy increases instantly to 23 MPa and 103 MPa, respectively. Previously, it was reported that grain morphology could make a difference to the anisotropic property [20,21], which is consistent with the results of this study. Therefore, even though the specimens have undergone hammer peening treatment, a lower welding speed remains an optimal welding parameter needed to control the grain structure and reduce the anisotropic property. ...
As a significant parameter required for welding, welding speed can be used to control the grain structure during the wire-arc additive manufacture of Ti64 alloy. Herein, the prior-β grain morphology and anisotropic tensile strength were examined to investigate the impact of welding speed after hammer peening treatment. Meanwhile, SYSWELD was applied to conduct a numerical analysis of the weld pool morphology and thermal history at different welding speeds. According to the experiment results, equiaxed prior-β grains developed in the faster welding speed specimen (F) and mixture prior-β grains (equiaxed and columnar) took shape in the middle and slow welding speed specimen (M and S), which was consistent with the anisotropic property as observed in the fast welding speed specimen with lower anisotropic tensile strength than others. Meanwhile, simulation results demonstrated that a faster welding speed led to a lower temperature gradient (G) and a higher solidification rate (R) resulted in a larger number of equiaxed grains compared to the slower welding speed specimen.
... The lines were repeated approximately every 2 mm. A similar methodology was already deployed by other authors [22,23]. X-Ray Diffraction (XRD) analyses were conducted with the aim to identify the phases in the specimens and to estimate the relative crystallographic parameters. ...
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This work deals with the effect of different heat treatments on directed energy deposition (DED)-produced Ti-6Al-4 V samples. Annealing treatments at 1050 °C followed by different cooling rates were conducted to allow a complete recrystallization of the microstructure and remove the columnar prior-β grains, thus increasing the overall isotropy of the material. An agine treatment at 540 °C was also performed for further microstructural stabilization. The microstructures, textures and mechanical properties were then assessed. Due to the heat treatments, greatly differing microstructures were achieved in an equiaxed grain morphology. However, a “grain memory” effect was detected which resulted in the grains size increasing along the height of the samples. This effect was correlated to the intrinsic prior-β grain width variation along Z on the as-printed specimens, typical of the DED technology. Electron backscatter diffraction analyses proved that the intensity of the preferential directions increased after the heat treatments, likely due to the crystallographic variant selection mechanisms taking place when the samples cool down from the annealing temperature. This effect is also influenced by the significant difference in terms of prior-β grains sizes between the heat-treated and the as-printed specimens. To sum up, a complete homogenization of the material via heat treatment above the β-transus temperature proved to be challenging. In fact, the data suggest that the intrinsic texture-related anisotropy granted by the manufacturing process is very difficult to be eliminated.
... However, this approach is difficult to be applied in some alloys with a narrow solidification interval and/or suppressed constitutional supercooling, like Ti6Al4V. Recently, some additional intermediate processes, like rolling and high-intensity ultrasonic treatment, are also employed to control the prior-β grain structures in DED Ti6Al4V [6,25], but inevitably rising both manufacturing cost and time. ...
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In metal additive manufacturing (AM), long columnar grains along the building direction due to highly directional thermal gradients often lead to a strong texture and severe anisotropy in mechanical properties of the deposits, which significantly impair part qualification and targeted applications. Here, a novel depositing strategy by periodically alternating processing parameters is designed to partially preserve the equiaxed grains resulted from the columnar-to-equiaxed transition at the top of each deposited layer and, at the same time, to interrupt the epitaxial growth of columnar grains in AM titanium alloy. With the help of the competitive growth of the new grains and the potential coarsening effect during subsequent thermal-cycles, a microstructure of full equiaxed prior-β grains in Ti6Al4V fabricated by laser directed energy deposition (DED) is finally obtained without either using any auxiliary equipment or adjusting alloy chemistry like previous researches. This further contributes to a superior mechanical property and a remarkable reduction on both the crystallographic textures and the property anisotropy.
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Aerospace industries invest a significant amount of resources to meet one common goal, that is, to make the aircraft fly. To keep down its 'buy to fly ratio', researchers have been working hard to introduce additive manufacturing (AM) technique for producing aerospace components. AM technologies are also now being used in major parts of an aircraft like fuel nozzles, turbofan blades, compressor-turbine blades, suspension wishbone, air ducts, etc., due to its just in time production with less complexity, direct tooling, and higher customer satisfaction with significant cost reduction including interior design. Nowadays, aerospace industries face problems meeting the deadline for delivering the aircraft components and replacement parts while maintaining certification standards. The wire arc additive manufacturing (WAAM) technique, one of the AM processes, can fabricate large metallic components with some reduction in lead time. WAAM process can build near net shape parts with high material deposition rate and efficiency while keeping the equipment and feedstock cost and material wastage minimal. This review paper summarizes the latest advancement on wire arc additive manufacturing of titanium and its alloy based on the aerospace application. Titanium and its alloys are used at a large scale in aircraft airframe structures and engine parts due to its high strength-to-weight ratio, excellent corrosion-resistant, high creep and fatigue resistance at an elevated temperature. It has been studied that the mechanical and metallographic properties of titanium and its alloy can be enhanced by using the WAAM process, and it is suited for aerospace applications. The paper will review the challenges like porosity, delamination, residual stress, crack propagation, anisotropic behavior, oxidation, etc., associated with the WAAM process on titanium alloys and propose recommendations for reducing the defects during the WAAM process.
Several designs and prototypes have been proposed for a robotic prosthetic arm using rigid components. At the same time, they can provide the motion but are either inflexible or very expensive. The soft robotic prosthetic arm presented in this case study is designed to meet the low-cost criterion and, at the same time, aid a trans-radial amputee in his daily task without harming himself or the delicate object. The arm is controlled by the amputee’s muscle signals using electromyography (EMG) sensors. For actuation, artificial pneumatic muscles are employed, lowering the actuation force requirement and enhancing gripping by curling on the object. The envision of prosthetic arm design works with high modularity and adapts according to the patient by utilizing 3D printing for its fabrication.
La fabrication d'additive par arcs électriques (WAAM) est en train de devenir la principale technologie de Fabrication Additive (FA) utilisée pour produire des pièces à parois minces de taille moyenne à grande (Ordre de grandeur : 1 m) à un coût moindre. Pour fabriquer une pièce avec cette technologie, la stratégie de planification du trajet utilisée est la 2.5D. Cette stratégie consiste à découper un modèle 3D en différentes couches planes et parallèles les unes aux autres. L'utilisation de cette stratégie limite la complexité des topologies réalisables en WAAM, notamment celles présentant de grandes variations de courbure, et implique plusieurs départs/arrêt de l'arc lors de son passage d'une couche à l'autre. Ceci induit des phénomènes transitoires dans lesquels le contrôle de l'approvisionnement en énergie et en matière est complexe. Dans cette thèse, une nouvelle stratégie de fabrication visant à réduire au minimum les phases de démarrage et d'arrêt de l'arc est présentée. L'objectif de cette stratégie, appelée "Génération de Trajectoire Continue Tridimensionnelle" (GTCT), est de générer une trajectoire continue en forme de spirale pour des pièces minces en boucle fermée. Une vitesse de fil constante couplée à une vitesse de déplacement adaptative permet une modulation de la géométrie de dépôt qui assure un approvisionnement continu en énergie et en matière tout au long du processus de fabrication. L'utilisation de la stratégie 5 axes couplées à la GTCT permet la fabrication de pièces fermées avec une procédure pour déterminer la zone de fermeture optimale, et des pièces sur des substrats non-plans utiles pour ajouter des fonctionnalités à une structure existante. La fabrication de ces pièces avec la GTCT et plusieurs évaluations numériques ont montré la fiabilité de cette stratégie et sa capacité à produire de nouvelles formes complexes avec une bonne restitution géométrique, difficile ou impossible à atteindre aujourd'hui en 2.5D avec la technologie WAAM.
Although wire arc additive manufacturing (WAAM) of Ti alloys using gas metal arc yields a high deposition rate, cathode spots exhibit unstable behavior. In this study, the cold-metal-transfer process and an electrode-negative (EN) polarity were used in the WAAM process of the Ti–6Al–4V alloy. High-speed imaging was performed to investigate the mechanisms to stabilize cathode plasma jets, arc plasma, and molten metal transfer in the EN polarity. Arc plasma and cathode jets had the same direction, and sound cathode jets were shrouded by arc plasma in the EN mode. The metal transfer was also stabilized by balanced plasma formation under the EN mode, and inconsistent wire melting under the electrode-positive (EP) mode was mitigated. In the deposition test using the EN mode, distributed heating of the substrate and depressed molten pool were observed, which resulted in a 21.0% increase in the bead width and 27% decrease in the wetting angle compared with the EP mode. This study demonstrated that the instability of the gas metal arc WAAM process of Ti alloys can be overcome with the EN-mode cold metal arc process.
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Wire + Arc Additively Manufactured components contain significant residual stresses that manifest in distortion. Each layer of an additively manufactured wall was rolled with the aim of reducing the residual stress. Neutron diffraction and contour method measurements show that the residual stresses were reduced – particularly at the boundary with the substrate. The process also reduced distortion, and refined the microstructure which may facilitate implementation on aerospace components.
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The thermal history developed in laser metal deposition (LMD) processes has been shown to be quite complex and results in the evolution of an equally complex microstructure. A companion article (Part I. Microstructural Characterization) discussed the LMD of Ti-6Al-4V, where the resultant microstructure consists of a periodic, scale-graded layer of basketweave Widmanstätten alpha and a banding that consists of colony Widmanstätten alpha. In order to understand the microstructural evolution in Ti-6Al-4V, a numerical thermal model based on the implicit finite-difference technique was developed to model LMD processes. The effect of different laser-scan velocities on the characteristics of the thermal history was investigated using an eight-layer single-line build. As the laser-scan speed decreases and the position within a layer increases, the peak temperature increases. The heating rate and the peak thermal gradient within a deposited layer were shown to follow the same trend as the peak temperature after two layers were deposited on top of the substrate. In general, the laser-scan speed or z-position within a layer did not have a significant effect on the cooling rate. The cooling rate in a newly deposited layer decreases as the number of layer additions increases. Given the predicted temperature vs time profile from the thermal model, the evolution of phase transformations occurring in the deposit is mapped as each layer is deposited. As a result of the thermal cycling imposed by the periodic deposition of material, a characteristic layer, consisting of two regions heated above and below the beta transus, forms in layer n due to the deposition of layer n+1.
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The macro and microstructure of laser-deposited Ti-6Al-4V has been investigated to determine the evolution of unique microstructural features in mutilayer builds. The macro and microstructures exhibited in the build include large, columnar prior-beta grains, a gradient in the individual alpha-lath thickness between the deposited layers, and the presence of layer bands within each layer, except for the last three layers deposited. The layer band consists of a colony Widmanstätten alpha morphology, while the nominal microstructure between layer bands exhibits a basketweave morphology. Optical microscopy, hardness, and composition measurements were used to determine that the layer-band and gradient morphologies are resultant from the complex thermal history the build experiences and not a result of segregation or oxidation. The gradient alpha and layer-band morphologies form in layer n after the deposition of layer n+3.
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At present, airframes are mainly composed of monolithic components, instead of small parts joined using welding or riveting. Ribs, stringers, spars, and bulkheads can be included in this category. After milling, they are assembled and joined to the aircraft skins, which have also been milled. The aim of these parts is to obtain a good strength-weight ratio, owing to their homogeneity. The milling of a monolithic structural part implies removing up to 95 per cent of the weight from the raw block material. Therefore, the main objective is to achieve the highest removal rate possible. However, conditions required to achieve this (high feed, large depth of cut) in milling imply high cutting forces, which in turn induce part deflection or vibrations in those zones (thin walls and floors) where stiffness is not sufficiently high. These static and dynamic problems often lead to inaccuracy of geometry, roughness, and possible damage to the machine spindle. This paper proposes a working methodology for efficient process planning, based on previous analysis of the static and dynamic phenomena that can occur during high-speed cutting. This methodology provides several steps that can be taken in order to minimize the bending and vibration effects; suggests optimal monitoring methods to detect process instability; and describes the best way to tune the cutting conditions and chip load, by means of simulation at different machining stages. In this way, the reliability of aeronautical production significantly increases. The global approach presented in this paper has been applied to two test pieces and two real parts, which were milled without suffering either static or dynamic problems.
Wire and arc additive manufacturing (WAAM) is a novel manufacturing technique in which large metal components can be fabricated layer by layer. In this study, the macrostructure, microstructure, and mechanical properties of a Ti-6Al-4V alloy after WAAM deposition have been investigated. The macrostructure of the arc-deposited Ti-6Al-4V was characterized by epitaxial growth of large columnar prior-β grains up through the deposited layers, while the microstructure consisted of fine Widmanstätten α in the upper deposited layers and a banded coarsened Widmanstätten lamella α in the lower layers. This structure developed due to the repeated rapid heating and cooling thermal cycling that occurs during the WAAM process. The average yield and ultimate tensile strengths of the as-deposited material were found to be slightly lower than those for a forged Ti-6Al-4V bar (MIL-T 9047); however, the ductility was similar and, importantly, the mean fatigue life was significantly higher. A small number of WAAM specimens exhibited early fatigue failure, which can be attributed to the rare occurrence of gas pores formed during deposition.
With increasing emphasis on sustainability, additive layer manufacturing (ALM) offers significant advantages in terms of reduced buy-to-fly ratios and improved design flexibility. Plasma wire deposition is a novel ALM technique in which plasma welding and wire feeding are combined. In the present work, a working envelope for the process using Ti–6Al–4V was developed, and regression models were calculated for total wall width, effective wall width and layer height. The plasma wire deposition process is able to produce straight walls of widths up to 17.4 mm giving a maximum effective wall width after machining of 15.9 mm, which is considerably wider than competing processes. In addition, for Ti–6Al–4V the deposition efficiency averages 93% and the maximum deposition rate is 1.8 kg/h. Coarse columnar grains of β phase grew from the base during deposition, which transformed into a Widmanstätten structure of α lamellae on cooling. Bands were identified in the deposits, which had a repetitive basket-weave microstructure that varied in size. The strength measured by micro-indentation hardness of 387 HV on average is as much as 12% higher than the substrate. These preliminary results indicate that plasma wire deposition is likely to be a suitable process for the ALM of large aerospace components.
Parts manufactured by Wire and Arc Additive Manufacture (WAAM) can have significant residual stress and distortion, as well as large grain sizes. To overcome these problems, each layer on a linearly deposited steel WAAM part was rolled with either a ‘profiled’ roller, which had a similar shape to the deposited layer, or a ‘slotted’ roller, in which a groove prevented lateral deformation. Both rollers reduced the distortion and surface roughness, but the slotted roller proved more effective – eliminating the distortion. The residual stresses in the rolled WAAM parts were measured and were lower than those in the unrolled control specimen – particularly adjacent to the baseplate. Rolling also induced additional grain refinement when the rolled material was reheated during the subsequent deposition pass. The application of rolling may be a key technology for enabling implementation of WAAM on large-scale structures.
In situ spatially resolved x-ray diffraction (SRXRD) experiments were used to directly observe the heat-affected zone phases present during gas tungsten arc welding of a Ti–6Al–4V alloy. The experiments were performed at the Stanford Synchrotron Radiation Laboratory using a 250 μm diam x-ray beam to gather real-time experimental information about the α−Ti→β−Ti phase transition during weld heating. Six different welding conditions were investigated using SRXRD to experimentally determine the extent of the single phase β-Ti region surrounding the liquid weld pool. These data were compared to predicted locations of the β-Ti phase boundary determined by calculated weld thermal profiles and equilibrium thermodynamic relationships. The comparison shows differences between the experimentally measured and the calculated locations of the β-Ti boundary. The differences are attributed to kinetics of the α−Ti→β−Ti phase transition, which requires superheating above the β-Ti transus temperature to take place during nonisothermal weld heating. Analysis of the results reveal that the transition to β-Ti requires an additional 3.96 s (±0.29 s) of time and 169 °C (±25.7 °C) of superheat above the β-Ti transus temperature to go to completion under an average weld heating rate of 42.7 °C/s. © 2003 American Institute of Physics.
The authors investigate the efficacy of applying rolling pressure along the weld line in thin butt welds produced using friction stir welding (FSW) as a means of controlling the welding residual stresses. Two cases are examined and in each case, comparison is made against the as welded condition. First, for FSW of AA 2024 aluminium alloy, roller tensioning was applied during welding using two rollers placed behind and either side of the FSW tool. Very little effect was seen for the down forces applied (0, 50, 75 kN). Second, for FSW AA 2199 aluminium alloy, post-weld roller tensioning was applied using a single roller placed directly on the FS weld line. In this case, significant effects were observed with increased loading, causing a marked reduction in the longitudinal tensile residual stress. Indeed, a load of just 20 kN was sufficient to reverse the sign of the weld line residual stress. Only slight differences in Vickers hardness were observed between the different applied loads. Furthermore, unlike some methods, this method is cheap, versatile and easy to apply.