Effect of Single and Duplex Stage Heat Treatment on the Microstructure and Mechanical Properties of Cast Ti–6Al–4V Alloy
A.-H. Hussein Faculty of Engineering, Cairo University
R. Reda (&) A. Nofal Central Metallurgical R&D Institute
, P.O. Box 87, Helwan,
The effect of single and duplex stage heat treatment on the microstructure and mechanical properties of cast Ti-6Al-4V alloy was studied. The single stage heat treatment (SSHT) involved solution annealing at 935 C for 10 min followed by water quenching, while the duplex stage heat treatment (DSHT) involved solution annealing at 935 C for 10 min, furnace cooling to 600 and 700 C, followed by isothermal holding for 30 min and subsequent water quenching. The properties characterization was conducted using microstructure investigation along with tensile, hardness, and Charpy impact tests. The impact fracture surfaces were observed using Scanning Electron Microscope. After SSHT, an increase in the tensile strength and hardness at the expense of the tensile elongation and impact toughness was recorded as compared with the as-cast alloy. This is attributed to the formation of a brittle a0 martensite phase in the microstructure after quenching from 935 C. On the other hand, DSHT enhances the percent elongation and impact toughness over the as-cast alloy and SSHT. The formation of a0 martensite-free microstructure as well as controlling the composition of the different microstructural constituents during DSHT optimizes the mechanical properties.
The requirements for outstanding mechanical properties
can justify the growing use of Ti-alloys in applications
such as the aerospace, automotive, medical, marine,
chemical industries, and sporting goods sectors . Due
to the significant raw-material costs, the manufacture of
near net-shape of Ti-alloy components is an attractive field.
Casting technologies enable manufacturing of complex
shapes and large parts [13, 5, 8], such as: cast frames for
aircraft engines, compressor casings, cast fan frames,
exhaust gas pipes of auxiliary gas turbines, connecting
rods, intake and outlet valves, and rim screws. Titanium
castings have been used also in biomedical and dental
applications, e.g., cast hip joint stems as well as for crowns
and bridges. Heat treatment is used for improving the
properties of the casting parts through microstructural
control. The properties of Ti-alloys can be significantly
changed through processing as well as heat treatment [1, 4,
The most commonly used Ti-alloy is the two phase
(a ? b) alloy, Ti6Al4V. The existence of the a/b
transformation means that a variety of microstructures and
property combinations can be achieved in the alloy through
heat treatment, thus permitting the adaptation of properties to
specific applications [8, 9]. For example, solution treatment
in b or at high temperature in a ? b range, followed by aging
for long periods, increases tensile strength and hardness at
the expense of tensile elongation and impact toughness [1,
8]. After such heat treatment, the alloy shows brittle failure
when subject to impact. On the other hand, annealing at low
temperature in a ? b range for long time followed by
furnace or air cooling enhances the tensile elongation and
impact toughness at the expense of tensile strength and
hardness . Therefore, the objective of this work is to
examine the effect of different heat treatment cycles, aiming
at optimizing the mechanical properties in order to improve
the performance of the Ti6Al4V-casting parts in their
applications. This aim was achieved through conducting a
number of heat treatment cycles and characterizing the
properties of the alloy after these treatments using
microstructural investigation in addition to tensile, hardness, and
Charpy impact tests. The fracture surfaces of the impact
specimens were also investigated.
The material used in this study was cast Ti6Al4V alloy.
The chemical composition of the studied alloy is shown in
Heat treatment cycles conducted throughout this study
are shown in Fig. 1. The single stage heat treatment
(SSHT) involved heating the samples to 935 C, and
isothermal holding for 10 min, followed by water quenching.
The duplex stage heat treatment cycles (DSHT) involved
furnace cooling from 935 C, after 10 min holding, to 600
and 700 C followed by isothermal holding for 30 min,
before water quenching down to room temperature.
Computerized furnace with a controlled atmosphere was used
Table 1 Chemical composition of the studied alloy
for heat treatment. Heat treatments were carried out in an
argon environment at a flow rate of 200 CFH and 1 bar.
The as-cast and heat-treated specimens were prepared by
standard metallographic techniques, which consist of
polishing and etching in an etchant composed of 10% HNO3, 5%
HF, and 85% distilled water. After etching, the specimens
were examined with an optical microscope and using
backscattered electron imaging in a scanning electron microscope
along with energy-dispersive spectrometry (EDS) to analyze
the different phases. X-ray diffraction analysis with CuKa
radiation as well as image analysis was used for quantitative
measurement of the volume fraction of the different phases.
After heat treatment and prior to mechanical testing,
specimens were hand-ground progressively to achieve a
smooth surface. The surface oxide layer formed during the
heat treatments was completely removed by polishing.
Tensile specimens were prepared in accordance with
ASTM E8, with a specimen length of 130 mm, gage length
of 30 mm, and gage diameter of 6 mm. Uniaxial tensile
tests were carried out at room temperature at a strain rate of
Standard Charpy V-notch impact specimens were
prepared in accordance with ASTM E23 standard specifications.
Charpy impact tests were done using a 150-J capacity
machine at room temperature. 5 9 9 55 mm3 Charpy
Fig. 1 Single and duplex stage
heat treatment cycles
specimens with a 45 V notch and 2 mm deep with a
0.25mm root radius were used. After impact testing, fracture
surfaces of the notched specimens were carefully observed
by scanning electron microscope (SEM) to investigate the
fracture mode and crack propagation behavior.
The average bulk Vickers hardness (Hv30) of the
specimens was measured using an applied load of 30 kg,
loading time of 15 s and 100 lm/s indenter speed,
according to ASTM E92. Six readings were taken on each
sample and the average of them is reported.
Results and Discussion
Microstructure of the as-cast Ti6Al4V alloy is shown in
Fig. 2. The bright regions in the optical micrograph
(Fig. 2a) correspond to a-phase, forming a typical
Widmansttaten structure, whereas thin dark regions between
Fig. 2 Microstructure of the as-cast Ti6Al4V alloy: (a) optical
micrograph shows thin areas of dark b-phase between bright lamellae
a-phase (9200); and (b) backscattered SEM image; light and dark
regions are b- and a-phases, respectively
Table 2 Volume fractions (%) of the different phases
Fig. 3 Microstructure of Ti6Al4V alloy after SSHT: (a) optical
micrograph shows a microstructure consisting of bright a-phase, gray
b-phase, and black a0-martensite phase (9200), and (b) backscattered
SEM image reveals the acicular morphology of a0-martensite phase
along with bright retained b- in a-phase matrix
a-plates are b-phase. In the Widmanstatten microstructure,
a-phase is formed along prior b-grain boundaries, and
colonies of lath-type b and a lamellar structure are present
inside prior b grains. In the backscattered SEM image
(Fig. 2b), the bright areas represent b-phase and the dark
regions are a-phase. b-Phase was measured by XRD to be
10%, as presented in Table 2.
The a ? b transformation temperature for Ti6Al4V
alloy used in this study was determined to be 987 C using
heat-flux differential scanning calorimetry (DSC). This
result is in agreement with previous data [1, 4, 6, 8, 11].
The microstructure of Ti6Al4V alloy after single stage
heat treatment (SSHT) from 935 C is shown in Fig. 3.
Optical micrograph of the microstructure after SSHT
(Fig. 3a) consists of a mixture of a0 martensite and b
structures with a plates formed inside and at prior b-grain
boundaries. Table 2 presents the volume fraction of the
different phases after SSHT. Many authors [1, 8, 11] pointed
out that the a-phase cannot be differentiated from a0
martensite phase by x-ray diffraction measurements because the
inter-planar spacing in the two structures are nearly the same.
Consequently, the obtained x-ray diffraction data can only
confirm the presence of a stable b-phase in quantities high
enough to be detected. Therefore, b-phase was determined
using XRD analysis while the bright a-phase was measured
using image analysis and the remained a0 martensite phase
was a rough estimation of the supplement equal to 100.
Backscattered SEM image of the alloy after SSHT (Fig. 3b)
reveals the morphology of the acicular a0 martensite phase
along with bright areas of the retained b-phase in a dark
matrix of a-phase. The presence of a0 martensite phase after
quenching from high temperature in a ? b range is widely
reported by several authors [14, 6, 8, 11].
Microstructures after duplex stage heat treatment (DSHT)
at 600 and 700 C are shown in Fig. 4. In these conditions,
the microstructure consists of a and b-phases. The
microstructure after DSHT appears more homogenous compared
with the as-cast condition. There is no a0 martensite phase
formed after DSHT. Controlled cooling followed by
isothermal holding at low temperature promotes the growth of a
plates, thus enriching the b phase with b stabilizers. The
enriched b-phase has a lower beta-transus (Ms) temperature
and, hence, a lower tendency to transform to a0 so b-phase
remains at room temperature as retained b .
The microstructure after DSHT at 600 C (Fig. 4a)
contains lower amounts of retained b-phase than at 700 C
(Fig. 4b), as reported in Table 2. This is in agreement with
previous results [2, 4, 11] which stated that as the solution
treatment temperature in a ? b range decreases, the
aphase fraction increases at the expense of b-phase. On the
other hand, b-stabilizer concentration (V%) in b-phase
Fig. 4 Optical micrographs of Ti6A14V alloy after DSHT at:
(a) 600 C and (b) 700 C (9200)
Table 3 Chemical composition (%) of a- and b-phases after DSHT
Table 4 Mechanical properties of Ti6Al4V alloy
Yield Strength (MPa)
Impact toughness (J)
increases with decreasing temperature. Table 3 presents the
composition of a- and b-phases after DSHT. As illustrated,
%V is higher after DSHT at 600 than at 700 C.
The effect of different heat treatment cycles on the
mechanical properties of Ti6Al4V alloys is presented in Table 4.
The tensile strength increased after SSHT with a reduction in
the tensile elongation compared with the as-cast alloy. This is
attributed to the formation of a brittle a0 martensite phase after
quenching from 935 C. After DSHT at 600 C, the tensile
strength slightly decreases with an increase in the tensile
elongation compared with the as-cast alloy, while DSHT at
700 C slightly increases the ultimate tensile strength as well
as tensile elongation over the as-cast alloy.
Improvement in the tensile elongation after DSHT is
attributed to the formation of a martensite free
homogenous microstructure. The increase in the ultimate tensile
strength after DSHT at 700 C over 600 C is due to the
higher volume fraction of retained b-phase at higher
temperature, as presented in Table 2. This result is in
agreement with Rhodes et al.  who stated that the presence
of the retained b-phase in the microstructure of Ti6Al4V
alloy improve its mechanical properties. Rhodes et al. 
reported also that b-phase enhances the strength by solid
solution strengthening by vanadium. Therefore, the main
strengthening mechanism in DSHT conditions is the solid
solution strengthening effect of b-phase.
Yield strength value after DSHT at 700 C is lower than
at 600 C in contrast to the increased value of ultimate
tensile strength. This can be interpreted by the higher V
content in b-phase after DSHT at 600 C over 700 C.
Rhodes et al.  reported that the increase in V content of
b-phase enhances the yield strength.
Table 4 presents also the effect of SSHT and DSHT on the
hardness values of Ti6Al4V alloy. The highest hardness
value is associated with the formation of a0 martensite phase
after SSHT. DSHT enhances the hardness values compared
with the as-cast alloy. Hardness is higher after DSHT at
700 C than at 600 C. This is attributed to the higher
volume fraction of b-phase at 700 C. b-Phase increases the
hardness by solid solution strengthening effect of vanadium.
The effect of SSHT and DSHT on the Charpy impact
toughness is reported in Table 4. The impact toughness
results change in the same trend as tensile elongation
results. The impact toughness after SSHT decreases
compared with as-cast alloy; as a result of formation of brittle
a0 martensite phase after quenching from 935 C. On the
other hand, the impact toughness increases after DSHT at
600 C and decreases after DSHT at 700 C. This is
attributed to the lower volume fraction of the strengthening
phase, i.e., the retained b-phase, at 600 C.
Figure 5 shows SEM fractographs of some specimens after
impact testing. The fracture surface of the as-cast
microstructure (Fig. 5a) was a mixed mode fracture consisting
primarily of ductile dimples with some cleavage fracture.
Fig. 5 SEM micrographs reveal the fracture surface of Ti6A14V
alloy: (a) cast alloy (b) SSHT; arrows refer to the microcracks, and
(c) DSHT/600 C
After SSHT, the amount of dimples decreased as compared
with the as-cast structure, in addition to the presence of
many microcracks (arrowed), as shown in Fig. 5b. This
indicates that the brittle a0 martensite phase plays a role in
the fracture behavior of SSHT samples.
For the duplex condition (DSHT/600 C), the energy
consumption is based on both the comparatively ductile
behavior of the homogenous microstructure as well as
crack deflection along and within the lamellar grains, see
Fig. 5c. The fracture surface reveals more homogenous
distribution of finer dimples after DSHT. The
Widmanstatten structure has large colonies with a lamellar structure
consisting of a-phase with strong b-phase, and these
colonies have different lamellar directions. Crack branching
and zigzaging occurred when cracks propagated, which
increased resistance to crack propagation and increased the
toughness, compared to the SSHT samples .
The effect of SSHT and DSHT on the mechanical
properties of the as-cast Ti6Al4V alloy was studied.
After SSHT, an increase in the tensile strength and
hardness occurs at the expense of the tensile elongation
and impact toughness compared with the as-cast alloy.
This is attributed to formation of a brittle a0 martensite
phase after quenching from 935 C.
There is an improvement in the tensile elongation after
DSHT over the as-cast alloy which may be attributed
to the formation of homogenous microstructure free
from brittle a0 martensite phase after quenching from
low temperature (600 and 700 C).
The increase in the ultimate tensile strength after
DSHT at 700 C over 600 C is due to the higher
volume fraction of retained b-phase at higher
temperature that causes soli-strengthening effect. On the other
hand, yield strength after DSHT is lower at 700 C
than at 600 C; this is attributed to the higher V
content of the strengthening phase at 600 C.
DSHT enhances the hardness values. This indicates
that higher fraction of retained b-phase after DSHT as
compared with the as-cast alloy plays an important role
in the alloy strengthening.
Fracture surface of the as-cast microstructure show
mixed mode of fracture. Lower amount of dimples along
with many microcracks characterize the fracture surface
after SSHT. After DSHT, the fracture surface reveals
more homogenous distribution of finer dimples.