Open Journal of Metal, 2011, 1, 1-11
doi:10.4236/ojmetal.2011.11001 Published Online September 2011 (http://www.SciRP.org/journal/ojmetal)
Characterization of the β Phase Decomposition in
Ti-5Al-5Mo-5V-3Cr at Slow Heating Rates
Quentin Contrepois, Marc Carton, Jacqueline Lecomte-Beckers
Aerospace and Mechanical Engineering Department, Metallic Materials Science Unit,
University of Liège, Chemin des chevreuils 1, Belgium
Received July 12, 2011; revised August 22, 2011; accepted September 8, 2011
Abstract
The influence of slow heating rates: 2, 5, 10 and 30˚C/min (0.033, 0.083, 0.166 and 0.50˚C/s) on the β phase
decomposition of Ti-5Al-5Mo-5V-3Cr (Ti-5553) during continuous heating were characterized by differen-
tial scanning calorimetry (DSC) analysis, light microscopy, scanning electron microscopy, X-ray diffraction
and hardness testing. Starting microstructure was the β phase obtained by heating the Ti-5553 above the Tβ
temperature and a water quench. Results show that heating rate has a significant impact on the precipitation
mechanisms and on the β→α transformation in this range of heating rates. The main formation of α precipi-
tates occurs between 500 and 600˚C at all heating rates tested. A heating at 2˚C/min produces very fine and
homogeneously distributed α plate precipitates which have nucleated on the nanometer size ωiso precipitates.
The ωiso precipitates between 350 and 400˚C. At higher heating rates 10, 15 or 30˚C/min, the amount of pre-
cipitation of ωiso is lower so an additional formation of nanometer size precipitates occurs between 450 and
500˚C. It is supposed that both precipitates act as nucleation sites for α phase precipitation. The resultant mi-
crostructure consists in a fine intragranular distribution of α precipitates and a coarser precipitation of α at
the grain boundaries. It is shown that the precipitation of ωiso phase retards or prevents the precipitation of
nanometer size precipitates occurring between 450 and 500˚C. This cannot be generalized to all the β-metast-
able titanium alloys since Ti-LCB does not exhibit the same heating rate dependence on DSC curves.
Keywords: β Metastable Titanium Alloy, ω Phase, Precipitation, Microstructure, DSC
1. Introduction
Ti-5Al-5Mo-5V-3Cr, or Ti-5553, is a high-strength β-
metastable (or β) titanium alloy comparable to the old
Russian β-metastable alloy VT-22. Ti-5553 was primar-
ily designed for high-strength forging applications on
account of improved properties and deep hardenability
over large thickness; moreover, it exhibits less sensitivity
to forming variables compared to Ti-10Al-2Fe-3Al and
thus offers processing benefits [1]. A wide variety of str-
ength, ductility and fracture toughness values can be ach-
ieved depending on thermomechanical processing and
heat treatment with these types of alloys [2,3] making Ti-
5553 potentially useful for various thick aerospace appli-
cations where weight reduction must be achieved (for
example in the low-pressure compressor or the landing
gear). Ti-LCB (Low Cost Beta Titanium) is another β
metastable alloy used in mechanical applications; it is
very attractive since the more costly β stabilizing ele-
ments are replaced.
The mechanical properties of titanium alloys are con-
trolled by their microstructure, in particular the size,
shape and distribution of the α phase in the β-matrix [3].
Therefore, a good understanding of the alloy behavior
during thermomechanical treatments is closely related to
the knowledge of phase transformations or precipitations
and the effect of temperature on the microstructure evo-
lution.
This paper investigates Ti-5553 and to a lesser extent
Ti-LCB solid β phase decomposition and precipitations
using differential scanning calorimetry (DSC) during
various slow heating rates. Phases are revealed by X-ray
diffraction (XRD) analysis in addition to light micros-
copy (LM), scanning electron microscopy (SEM) and
hardness testing. The starting microstructure is a fully
β-matrix free of α phase, resulting from a heating above
the Tβ transus temperature and a water quench (WQ).
The aim of this study is to show the particular β phase

2 Q. CONTREPOIS ET AL.
decomposition of the Ti-5553 alloy during slow heating
rates in the range of 2–30˚C/min and to put in relief the
importance of heating rate when studying these types of
alloys. Ivasishin et al. [4,5] have shown the importance
of the heating rate on various Ti alloys including the
predecessor of Ti-5553, VT-22, and showed that differ-
ent precipitation mechanisms occurred by varying the
heating rate until the ageing temperature. At slow heating
rates (15˚C/min or 0.25˚C/s), they observed a fine, pla-
te-like α precipitation and at very high heating rates
(1200˚C/min or 20˚C/s) a coarse plate-like α microstruc-
ture. They attributed this difference to the precipitation
of the ω isothermal phase (ωiso) which acts as nucleation
sites for the α precipitates in addition to the classical mi-
crostructural items such as grain boundaries, dislocations
or precipitates. The main difference with our study is the
difference in heating rates analysed: 15 and 1200˚C/min
in the Ivasishin studies and 2, 5, 15 and 30˚C/min for
ours, plus the studied alloy Ti-5553 which is slightly di-
fferent from VT-22. In addition, in a recent study, Wain
et al. [6] characterized a refinement of the α precipitation
in Ti-5553 by using either a low-temperature pre-ageing
at 250˚C or a sufficiently slow heating rate up to the final
ageing temperature when starting from a β phase micro-
structure and according to these authors, the mechanism
responsible for such a refinement is also an enhanced α
nucleation on fine ωiso precipitates. Two forms of ω
phase are differentiated: athermal ω (ωath) which appears
without diffusion during cooling from the high tempera-
ture β phase field and isothermal ω (ωiso) which precipi-
tates during low temperature ageing.
Influence of heating rate on β phase decomposition
was also studied in the β-metastable titanium alloys
Ti-B19 (Ti-3Al-5Mo-5V-4Cr-2Zr) [7] by electrical resis-
tivity measurements and Ti17 (Ti-5Al-4Mo-4Cr-2Zr-2Sn)
[8] by in situ high energy XRD. It was shown in Ti-B19
[7] that ωiso precipitates between 280 and 350˚C during a
6˚C/min heating but not during 60 and 600˚C/min heat-
ing. Moreover the precipitation of the orthorhombic α”
phase was evidence between 350 and 450˚C at 6˚C/min.
Similarly, two different β phase decompositions are evi-
dence in Ti17 [8]: at low heating rate 30˚C/min:
β (250˚C) → ω + β (393˚C) → α” + β (490˚C) → α” + α
+ β (590˚C) → β + α (633˚C)
and at high heating rate 300˚C/min:
β (300˚C) → α” + β (528˚C) → α” + α + β (604˚C) → α
+ β (632˚C)
(Temperatures in brackets correspond to the temperatures
of analyses).
It is then shown that the precipitation of ωiso is sup-
pressed at 300˚C/min heating rate.
Presence of ω precipitates in Ti-5553 has already been
reported in detail in previous papers [9,10,11,12,13]. Jo-
nes et al. [13] for example clearly showed, using in situ
high energy XRD in a synchrotron that the ωath phase
appears during rapid cooling from above the β transus
temperature in Ti-5553 as much as 20% in volume. The
role played by the ω phase in the α precipitation in this
type of alloy is of interest and the possible mechanisms
have been discussed previously [9,10,11,12].
Ohmori et al. [9] have discussed the role of ω precipi-
tates on the nucleation of the HCP α and orthorhombic
α” phases in a β-metastable titanium alloy (Ti-9.87V-
1.78Fe-3.20Al). Thanks to transmission electron micros-
copy (TEM) analysis they showed that α laths nucleate at
the ω / β interface and grow into both the β matrix and
the ω particles at temperatures above 350˚C. Nag et al.
[12] investigated the role of ω precipitates on the intra-
granular nucleation of α precipitates in Ti-5553, it is
thought to arise from two possible causes: 1- nucleation
near the ω / β interface by local rejection of Al which is
an ω destabilizer and a strong α stabilizer [14]; 2-dis-
placive transformation within the core of the ω precipi-
tates, leading to the formation of α platelets [10]. The
importance of ω / β misfit in the α nucleation mecha-
nisms is reminded, when there is a large misfit between
the ω and β phases, α precipitates nucleate on ledges and
misfit dislocations at ω / β interfaces. But Ti-5553 is
thought to exhibit a low ω / β misfit so the role of ω on
the nucleation of α precipitates is not clearly understood.
Their detailed study leads to the conclusion that the β to
α transformation and a compositional partitioning of al-
loying elements occur in sequential steps, resulting in a
mixed-mode displacive-diffusional transformation. In-
deed, their experimental results indicate that a thermally
activated diffusive process leads to the initial partitioning
of alloying elements during coarsening of the ω precipi-
tates, leading to the creation of favorable sites, margin-
ally enriched in Al, for the nucleation of the α precipi-
tates. Subsequently, these precipitates nucleate by a pri-
marily displacive process, with a composition marginally
enriched in Al as compared with the parent β matrix.
These precipitates then grow via a coupled displacive-
diffusional process with the diffusional partitioning of
the alloying elements being rate-controlling. It is very in-
teresting to note that: - after an annealing at 350˚C (2h)
they detect a coarsening of ω precipitates and nucleation
of α precipitates, - after an annealing at 400˚C (2h) there
is less ω phase than at 350˚C but more α precipitates, and
finally after an annealing at 600˚C they observe a growth
and coarsening of α precipitates. They unfortunately do
not mention any heating rate, which leads us to think that
the samples were introduced into the furnace preheated
to the desired temperature. However, we will see in our
study that the thermal route have a high impact on the
precipitation mechanisms, and especially the heating rate.

3 Q. CONTREPOIS ET AL.
This overview shows that the presence of ω precipi-
tates is of great importance for α nucleation and that the
heating rate can have a strong impact on the ω phase for-
mation so that we propose to analyze the role played by
heating rate during the β phase decomposition of a
β-metastable alloy.
2. Materials and Experimental Procedures
Ti-5553 was provided by Timet Ugine (France), it was
produced in the form of a forged billet of 200mm in di-
ameter, the chemical composition of Ti-5553 is presen-
ted in Table 1. The conventional microstructure delive-
red by the producer is an α+β bimodal microstructure
(see Figure 1-(a)), consisting of large globular primary α
precipitates and smaller secondary α platelets in a β- ma-
trix. This microstructure is achieved by a two-step ageing
heat treatment at 820˚C for 2 h (i.e. below the Tβ transus)
and at 630˚C for 8 h. This state is called “bimodal”. Ti-
5553 has been solution heat treated at a super-transus
temperature of 865˚C (i.e. above the Tβ transus) over 30
min and water quenched in order to get a fully β-meta-
stable matrix free of α phase (but the presence of ωath is
expected). This state is termed the “β-solutionized” state
(see Figure 1-(b)). Similar solution heat treatment was
conducted with one other β-metastable titanium alloy:
the Ti-LCB at 810˚C for 30min and water quenched, in
order to compare the chemical composition effect. The
Ti-5553 β-solutionized samples were heated at two dif-
ferent rates: 2˚C and 30˚C/min using a resistance heating
furnace at four different temperatures: 300, 400, 500 and
600˚C and then water quenched.
Differential scanning calorimetry (DSC) is a thermo-
analytical technique in which the difference in the amo-
unt of heat required to increase the temperature of a sam-
ple and a reference is measured as a function of tem-
perature. When the sample undergoes an endothermic or
an exothermic physical transformation, heat flux is mea-
sured and compared to the reference which must have a
well-defined heat capacity over the range of temperatures
scanned. The difference of heat flux enables the detec-
tion of transformations such as solid phase transforma-
tions. Usually, the temperature program for a DSC ana-
lysis is designed in such a way that the sample tempera-
ture increases linearly as a function of time. DSC analy-
ses were performed on a Netzsch DSC 404C device at
four different heating rates: 2, 5, 10 and 15˚C/min, on
Ti-5553 and Ti-LCB samples in the initial β-solutionized
states (respectively 865˚C and 810˚C for 30min then wa-
ter-quenched). The heating rates 15˚C/min should have
been chosen instead of 30˚C/min for the heat treatment
of samples in the resistance heating furnace, but this fur-
nace wasn’t able to correctly heat at a constant rate of
Table 1. Chemical composition of Ti-5553 and Ti-LCB alloys.
(wt%) Al Mo V Cr Fe
Ti-55535 5 5 3 0.3
Ti-LCB1.5 6.8 4.5
(a) (b)
Figure 1. Microstructures of Ti-5553: (a). SEM micrograph
in as-received bimodal state (820˚C - 2 h/630˚C - 8 h) con-
sisting of globular primary alpha in fine lamellar α+β ma-
trix; (b). LM micrograph of β-solutionized Ti-5553 (865˚C -
30 min/WQ).
15˚C/min. We assume that 15˚C/min and 30˚C/min lead
to comparable microstructures in this paper.
XRD analyses were performed on a Siemens D-5000
X-ray diffractometer using the Cu-Kα ray of wavelength
1.5406 Å. Samples were mechanically polished using the
following steps: SiC papers (320 to 1200), diamond paste
(9 and 3 µm), colloidal silica suspension (OPS with
H2O2). The microstructures were revealed by chemical
etching with 5% HNO3-5% HF in a water solution (kno-
wn as Kroll etchant).
Electronic SEM micrographs were carried out with an
XL30 ESEM-FEG. Hardness was measured as Vickers
hardness with a 10 kg load on an EMCO-TEST device.
3. Results
3.1. DSC Curves of Ti-5553 in β-Solutionized
State (865˚C - 30 min/WQ)
Samples of β-solutionized Ti-5553 were analyzed by
DSC. The thermal cycles consisted of heating at constant
rates from room temperature to 930˚C to achieve a final
temperature above the Tβ transus. In order to character-
ize a possible kinetic effect on the phase transformations,
four different heating rates were tested: 2, 5, 10 and
15˚C/min. The DSC curves obtained during heating are
shown in Figure 2. These curves are complex and the
exothermic or endothermic nature of the transformation
peaks is not easy to identify. The particular difficulty in
DSC analysis is to correctly define a baseline which en-
ables detection of the heat flux characteristic of phase we
can determine three different “steps” or “areas”, called A,

Q. CONTREPOIS ET AL.
4
Figure 2. DSC curves of β-solutionized Ti-5553 (865˚C - 30 min/WQ) at: (a). the four heating rates; (b). 15˚C/min, (c).
10˚C/min; (d). 5˚C/min; (e). 2˚C /min.
B and C in Figure 2, for all the heating rates. transfor-
mations. If we consider a baseline approximately hori-
zontal at a DSC value close to the start of the curves,
The first area A is not clear; it can either be an exother-
mic reaction zone located at about 200˚C or an endo-
thermic reaction ranging from about 220˚C to 350˚C, we
think this is an endothermic peak. The second area B is a
series of exothermic peaks ranging from about 350˚C to
575˚C (labeled 1 to 3 on the curves) and the third area C
is an endothermic peak ranging from about 600˚C to
860˚C. As the Tβ transus is known to be near 865˚C, we
assume that the endothermic reaction (area C) between
about 650˚C to 860˚C is the α/β transformation (or the
α/β phase field).
The effect of heating rate is demonstrated by a shift in
the exothermic peak positions towards higher tempera-
tures for increasing heating rates, as shown in Figure
2-(a). The higher the heating rate is, the higher the tem-
peratures of the different peaks are. This is an effect ob-
served in thermally activated reactions. Another more
interesting effect of the heating rate is the amplitude
variation of the different exothermic peaks between 350
and 600˚C (area B). At heating rates of 10 and 15˚C/min,
we clearly observe three different exothermic peaks

5 Q. CONTREPOIS ET AL.
whereas at the heating rate of 2˚C/min, the peak 2 does
not exist. This peak 2 is slightly visible at 5˚C/min but it
is clearly less important in comparison to the two higher
heating rates. The peak 1 is the main exothermic peak at
the two lower heating rates but it seems to decrease at the
expense of the peak 2 as the heating rate increases. This
variation indicates a change in the exothermic mecha-
nisms occurring in this temperature range which is linked
to the heating rate. However, the relative amplitude of
the peak 3 is not affected by the variations in heating rate.
The large endothermic area labeled C on all the DSC
curves corresponds to the α/β phase field. The Tβ transus
temperature is confirmed to be at about 860˚C. Addi-
tional DSC results on Ti-5553 and Ti-LCB have previ-
ously been reported by Carton et al. [16] with different
starting microstructures, and they obtained different DSC
curves.
3.2. DSC Curves of Ti-LCB in the β-Solutionized
State (810˚C /30 min/WQ)
DSC analyses carried out on β-solutionized Ti-LCB are
shown on Figure 3. Two exothermic peaks are clearly
illustrated at the four heating rates, the first between
Figure 3. DSC curves of β-solutionized Ti-LCB (810˚C – 30 min/WQ) at: (a) the four heating rates; (b) 15˚C /min; (c) 10˚C
/min; (d) 5˚C /min; (e) 2˚C /min.

6 Q. CONTREPOIS ET AL.
300˚C and 375˚C (peak 1) and the second, the most im-
portant one, between 475˚C and 550˚C (peak 3). A less
marked transformation seems to occur in the temperature
range 400˚C - 450˚C, labeled peak 2 on the DSC curves.
Heating rate influence is demonstrated through the shift
in the peak position, but in comparison to the DSC
curves of Ti-5553 in Figure 2, there is no variation in the
relative intensity of the peaks. Heating rate in the range
2˚C - 15˚C/min is not responsible for a significant modi-
fication of the DSC curves, and by extension we can as-
sume that it does not influence the phase transformation
mechanisms in the same way as in Ti-5553.
Prima [15] has studied in details the phase transforma-
tion in Ti-LCB. He carried out in situ resistivity meas-
urements during continuous heating at 2˚C/min. A-
ccording to the variations observed on his resistivity
curves, he divided the heating into four domains and
made hypothesis correlated with experimental observa-
tions: domain I from room temperature to 200˚C corre-
sponds to the dissolution of ωath, domain II from 200 to
350˚C (peak 1 on Ti-LCB DSC) corresponds to the pre-
cipitation of ωiso, domain III from 350 to 480˚C (peak 2
on Ti-LCB DSC) corresponds to the disappearance or
coalescence of ωiso and α precipitation, domain IV at
480˚C (peak 3 on Ti-LCB DSC) corresponds to the for-
mation of α. Comparison between Prima’s results and
our DSC curves leads to the following correlation: peak
1 corresponds to the precipitation of ωiso and peak 3 cor-
responds to the precipitation of α phase.
The comparison between Ti-LCB and Ti-5553, both
β-metastable alloys, could lead us to think that peak 1 on
the Ti-5553 DSC curves corresponds to ωiso formation,
peak 2 to the disappearance or coalescence of ωiso plus α
precipitation and peak 3 to the precipitation of α phase.
3.3. Microstructure Analysis of β-Solutionized
Ti-5553 After Heating to 300˚C, 400˚C,
500˚C and 600˚C at Two Heating Rates of
2˚C and 30˚C/min.
The β-solutionized Ti-5553 samples were heated at two
different heating rates: 2˚C/min and 30˚C/min from room
temperature to 300˚C, 400˚C, 500˚C and 600˚C followed
by water quenching. According to the position of peaks
on the Ti-5553 DSC curves (Figure 2), 300˚C, 400˚C,
500˚C and 600˚C are approximately located between ea-
ch exothermic DSC peak. As a consequence these tem-
peratures are representative for the microstructural evo-
lution displayed by DSC.
XRD investigations were carried out on each sample.
Results are presented in Figure 4 in the range 2θ: 33-43˚,
where we find the main peaks of the α and β phases. It is
important to note that due to the small size of ω precipitates,
they cannot generate characteristic peak on conventional
Figure 4. X-Ray diffraction data for β-solutionized Ti-5553 (865˚C-30 min/WQ) after various heat treatments at two heating
rates: (a). 30˚C/min; (b). 2˚C/min.

7 Q. CONTREPOIS ET AL.
XRD pattern as confirmed by Jones et al. [13]. In the
β-solutionized state and after heating up to 300˚C, only
the (110)β peak is visible. After heating up to 400˚C and
500˚C, we observe at the heating rate of 2˚C/min a wid-
ening and a shift of this peak to higher values of 2θ. At
the heating rate of 30˚C/min, the widening of (110)β is not
observed at 400˚C but it is observed at 500˚C. At both
heating rates, after heating to 600˚C, characteristic peaks of
the α phase appear whilst the β peak is still visible. Between
the β-solutionized state and the 600˚C heating, the position
of the (110)β peak shift from 2θ = 39.1˚ to 2θ = 39.6˚.
At 2˚C/min, XRD data show an evolution of the micro-
structure between 300 and 500˚C. In this temperature
range, no α phase is clearly detected as it is at 600˚C, so
the widening of the (110)β peak could correspond either to
the precipitation of ωiso or to the nucleation of very small α
precipitates on the ω / β interface. Between 500 and 600˚C,
XRD demonstrated the formation of the α phase.
The micrographs in Figure 5 evidence that heating
Ti-5553 from a β-solutionized state up to 600˚C at two
different heating rates of 2 and 30˚C/min leads to differ-
ent microstructures.
At the 2˚C/min heating rate, the response of the alloy
heated up to 600˚C to chemical etching consists of a
Figure 5. LM micrographs of β-solutionized Ti-5553 (865˚C - 30 min/WQ) after various heat treatments at two heating rates:
(a, b, c, d). 30˚C /min and (e, f, g, h). 2˚C /min. Intragranular precipitation is observed on microstructures a) and b) whereas
only microstructures a) exhibit grain boundaries precipitation. Chemical etching leads to a darkening of surface of sample e)
due to α precipitation.

8 Q. CONTREPOIS ET AL.
complete darkening of the surface (Figure 5-(e)) making
its observation by LM difficult. Between 300 and 500˚C,
no change is observed by LM (Figure 5-(f) to (h)).
At the 30˚C/min heating rate, precipitation is not uni-
form, first precipitates appear inside the grains between
400 and 500˚C (Figure 5-(b)) whereas larger precipitates
appear at the grain boundaries between 500 and 600˚C
(Figure 5-(a)). The surface is darkened inside the grains
whereas it is brighter at the grain boundaries. The preci-
pitation occurring between 400 and 500˚C inside the
grains at 30˚C/min (Figure5-(b)) is not observed at
2˚C/min (Figure 5-(f)).
SEM micrographs (Figure 6) confirm LM observa-
tions and reveal additional information. At 500˚C, at a
heating rate of 30˚C/min (Figure 6-(c) and (d)), nano-
meter size intragranular precipitates can be observed at
high magnifications whereas nothing appears at the grain
boundaries. Between 500 and 600˚C, the α precipitates
appear both inside the grains and at the grain boundaries
(Figure 6-(a) and (b)) but with different size. These pre-
cipitates are plate shaped and form “triangle” patterns
that are fine inside the grain and coarser at the grain
boundaries (Figure 6-(a)). These α precipitates observed
at 600˚C are termed “micron size α precipitates” in this
paper and are defined as αM inside the matrix and αGB at
grain boundaries. This difference in size of α precipitates
between grain boundaries and rest the grain arises from
the previous precipitation which takes place between 400
and 500˚C inside grains and not at the grain boundaries.
We do not know the exact chemical and crystallographic
nature of these small precipitates because XRD data do
not prove the presence of α phase. However the micro-
graphies evidence that they are the precursors of the pre-
cipitation of αM precipitates. Thus we choose to call these -
precipitates αMprec.
At a heating rate of 2˚C/min, no precipitates are visi-
ble at the quench temperature of 500˚C (Figure 6-(h)).
The α phase precipitates homogeneously in the matrix
Figure 6. SEM micrographs of β-solutionized Ti-5553 (865˚C - 30 min/WQ) after heat treatments up to 500˚C and 600˚C at
two heating rates: (a-d). 30˚C /min and (e–h). 2˚C /min. (GB = grain boundaries).

9 Q. CONTREPOIS ET AL.
and at the grain boundaries between 500 and 600˚C, and
resultant α plates are fine and homogeneously distributed
(Figure 6-(e) and (f)). We did not observe any difference
between the grain boundaries and the centre of grains.
The arrangements of the precipitates are different from
those observed at 30˚C/min; they do not form triangle
patterns but small plates.
Hardness variations with temperature are illustrated in
Figure 7. Hardness indentations were sufficiently small
to differentiate between intragranular and grain boundary
hardness when differences were highlighted on the mi-
crographs at 30˚C/min. The comparison with XRD re-
sults (Figure 4) shows that increase of hardness is con-
comitant with the widening of the β XRD peak. At a
heating rate of 2˚C/min, hardness increases between 300
Figure 7. Vickers hardness of β-solutionized Ti-5553 (865˚C
-30 min/WQ) after heating at two heating rates: 30°C/min
and 2˚C/min.
and 400˚C, maximum is reached at 500˚C and slightly
decreases with the appearance of micron size α precipi-
tates at 600˚C. At a heating rate of 30˚C/min, when no
precipitation is observed, i.e. 300˚C, 400˚C and at grain
boundaries at 500˚C, the hardness remains low; when
precipitation is observed, the hardness increases, i.e. in-
side grains at 500˚C and everywhere at 600˚C. The pre-
cipitation responsible for the increase in hardness at
30˚C/min is visible in LM after chemical etching wh-
ereas this is not the case at 2˚C/min.
4. Discussion
A summary of the results is given in Table 2.
At 2˚C/min the DSC exothermic peak 1 is responsible-
for a widening of the β XRD peak and an increase in
hardness. According to experimental results from other
authors [9,11], this peak could correspond to the trans-
formation of ω to nanometer size α, this is not clear and
it also could simply correspond to the formation of
nanometer size ωiso. We assume the second hypothesis.
Peak 3 corresponds unambiguously to the formation of
micron size α precipitates. The homogeneous distribution
of the micron size α indicates that ωiso precipitates
formed during heating are responsible for the increase in
hardness and act as nucleation sites for the micron size α
formation (DSC peak 3). The following β phase decom-
position is proposed:
β+ωath → β → β+ωiso → β+α
At 30˚C/min, DSC peak 1 is present but the difference
with 2˚C/min arises from the appearance of DSC peak 2.
Peak 2 corresponds to the formation of intragranular
Table 2. Overview of all the experimental results at the two heating rates of 2˚C/min and 30˚C/min.
Heat treatment
Microstructure
at micron scale
(Figure 5 - 6)
DSC
(Figure2)
XRD
(Figure 4)
Vickers
Hardness
(Figure 7)
Supposed nanometer and micron
scale microstructure evolution
600˚C
fine micron size αMatrix
and coarser micron size
αGB
exothermic peak 3 α + β peaks 409
(GB 378)
fine precipitation of micron size
αM inside grains nucleated on ωiso
and αMprec and coarser micron size
precipitation of αGB at GB
500˚C
fine intragranular
precipitation
and PFZ at GB
exothermic peaks 1 and
2
widening of β
peak
406
(GB 281)
ωiso formation (peak 1) then for-
mation of αMprec inside grains
(peak 2) but nothing at GB
400˚C β grains (area A) β peak 258 dissolution of ωath
30˚C
/min
300˚C β grains (area A) β peak 258 dissolution of ωath
600˚C fine micron size α
precipitates exothermic peak 3 α + β peaks 412 fine homogeneous precipitation of α nucleated on ωiso
500˚C β grains exothermic peak 3* widening of
β peak 440 *
400˚C β grains exothermic peak 1 widening of
β peak 398 ωiso formation
2˚C/m
in
300˚C β grains (area A) β peak 264 dissolution of ωath
PFZ = Precipitate-free zone, GB = grain boundaries, IntraG = Intragranular
*According to the position of the 3rd exothermic peak, micron size α should have been observed.

Q. CONTREPOIS ET AL.
10
nanometer size precipitates which are visible after che-
iso Mprec iso MGB
n osr
pr
ce of heating rate in the range of
2) Heating 600˚C at
30˚C/min results in a fine intragranular distribution of
m
vents
the
e Walloon Region for fun-
. J. Dashwood, D. Dye and M. Jackson,
anical Processing of Ti-5Al-5Mo-5V-3Cr,”
mical etching. Their exact nature is not known so we call
them αMprec. Peak 3 corresponds to the formation of mi-
cron size α precipitates, which are fine inside the grain
(αM) and larger at the grain boundaries (αGB). Thus, two
possible precursors exist at this heating rate: ωiso and
αMprec, this lead to various mechanisms possible for the α
phase precipitation and the following β phase decompo-
sition is proposed:
β+ωath → β → β+ω → β+α+ω → β+α+α
It is likely that the precipitatiof ωiso phase retard o
events the precipitation of αMprec. The ωiso precipitation
is sufficiently high at 2˚C/min to inhibit the formation of
αMprec corresponding to DSC peak 2. At faster heating
rates (e.g. 10, 15 and 30˚C/min) the exothermic peak 2 is
present between 450 and 500˚C so that the formation of
the finely distributed intragranular precipitation of αMprec
is not inhibited.
Our results clearly show that α nucleation in Ti- 5553
has different possible mechanisms depending on the
heating rate. Clément [18], who studied Ti-5553, was
confronted to the same difficulties in interpreting the
responses of the different techniques he used to study
phase transformations in Ti-5553. He found that differ-
ences in the peaks position may be observed with this
alloy between various techniques so care should be taken
with their interpretation.
Due to the nanometer size of the particles, precise re-
sponses could only be obtained with extensive and sys-
tematic TEM analyses or/and with in situ high energy
XRD. This will be the subject of future studies, espe-
cially to verify if the αMprec precipitates could be an or-
thorhombic α” phase as proposed by Bruneseaux et al.
for the β phase decomposition of Ti17 [8].
. Conclusions 5
The strong influen
2–30˚C/min (= 0.033–0.5˚C/s) on the β phase decompo-
sition of β-solutionized Ti-5553 was established using
DSC analysis. Phase transformations were characterized
by LM and SEM micrographs, XRD and hardness test.
The results obtained are as follows.
1) Heating β-solutionized Ti-5553 up to 600˚C at
2˚C/min of results in a fine and homogeneous distribu-
tion of α precipitates in the matrix. Formation of micron
size α precipitates occurs between 500 and 600˚C. Be-
fore 300˚C, ωath disappear and ωiso form between 300 and
400˚C. The ωiso precipitates act as nucleation sites for α
plates precipitation. Following sequence is proposed:
β+ωath → β → β+ωiso → β+α
β-solutionized Ti-5553 up to
icron size αM precipitates and a coarser precipitation of
micron size αGB at the grain boundaries. The formation of
ωiso is present between 350 and 450˚C but not suffi-
ciently to inhibit the formation of nanometer precipitates
called αMprec which takes place between 450 and 500˚C.
Both ωiso and αMprec precipitates act as nucleation sites for
α precipitation. Following sequence is proposed:
β+ωath → β → β+ωiso →β+αMprec+ωiso →β+αM+αGB
3) The precipitation of ωiso phase retards or pre
precipitation of αMprec.
4) Influence of slow heating rates on precipitation
mechanisms cannot be generalized to all the β-metastable
titanium alloys since Ti-LCB does not exhibit the same
heating rate dependence on DSC curves.
With these experimental observations, it is shown that
the ω precipitates have a strong influence in the β to α
phase transformation in Ti-5553 during a slow heating.
Slow enough heating produces a fine and uniformly dis-
tributed α precipitation. In this precipitation sequence,
the main factor is the time necessary to form a high
enough volume of ωiso precipitates. In an industrial proc-
ess, the slow heating rate can reasonably be replaced by
an isothermal treatment at 350˚C followed by a heating
to 600˚C in order obtain a microstructure with a fine and
homogeneous distribution of α plates. This technique co-
uld be used to design new microstructures in the β-me-
tastable alloys.
6. Acknowledgements
The authors wish to thank th
ding and support.
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