submicrocrystalline structure formation in ti and ti-64 alloy by warm “abc” deformation
TRANSCRIPT
Submicrocrystalline Structure Formation in Ti and Ti-64 Alloy
by Warm “abc” Deformation
S.V. Zherebtsova, S.V. Mironovb, G.A. Salishchevc
Institute for Metals Superplasticity Problems, Khalturin Str 39, Ufa, 450001, Russia
Keywords: titanium, titanium alloy, microstructure refinement, “abc” deformation, superplastic flow.
Abstract. Mechanical behavior and microstructure evolution of Ti and Ti-64 titanium alloy during
warm “abc” deformation has been studied. The “abc” deformation was consisted of successive
compression of a sample along three orthogonal directions. Mechanical behavior of each material was
described by set of successive σ-ε curves combined into cumulative σ-Σε curve. Microstructure of Ti
was found to be refined to a grain size of about 0.4 µm due to formation of deformation-induced
boundaries within initial grains. Although a stage like steady state flow was observed at the
cumulative σ-Σε curve such mechanical behavior was hardly associated with superplastic flow. In
two-phase Ti-64 alloy the structure was found to be refined to a grain size of about 0.4 µm after warm
“abc” deformation due to globularization α- and β-particles following breaking down of α-lamellar
and β- layers. Microstructure refinement of the alloy was associated with softening and superplastic
flow.
Introduction
One of the necessary criterions for implementation of superplastic deformation is a certain
microstructure with grains (or subgrains) smaller than 10-15µm that allows attaining superplastic
deformation at temperatures above 0.4-0.5 of melting temperature [1]. Microstructure refinement to a
grain size less than 1µm (submicrocrystalline (SMC) structure) decreases the temperature of
superplasticity by few hundred degrees compared to conventional microstructures [2]. However
producing of SMC structure in bulk specimens is a difficult engineering and technological problem. A
promising technique for fabrication of SMC bulk billets is “abc” deformation [3] consisting of
successive compression of a sample along three orthogonal directions under isothermal condition.
Microstructure refinement during severe plastic deformation (SPD) has been studied extensively
[4-7]. Refinement of a single-phase materials microstructure is usually associated with development
of deformation-induced boundaries and twin boundaries and increasing their misorientation by
absorption of dislocations [4, 5, 7]. However those investigations were focused on metals with
symmetric cubic crystalline lattice (fcc or bcc). Large deformation of hcp metal was studied
insufficiently; meanwhile limited number of slip system operating and twinning can be key factors
influencing on microstructure evolution in Ti.
During high-temperature deformation lamellar microstructure of two-phase titanium alloys
evolves by means of globularization of each phase of the alloys [8]. Progress toward globularization
briefly might be described as formation of transverse intra-phase boundaries in lamellae and then their
segmentation and spheroidization [2, 8]. However formation of SMC structure requires deformation
at lower temperatures and it is not clear enough if the structure refinement at low temperature similar
to that at high temperature. In addition, it is of interest to compare mechanical behavior of the alloy
with that of single-phase titanium during low-temperature deformation.
This work was aimed at the investigation of microstructure evolution and mechanical behavior of
titanium and two-phase Ti-64 titanium alloy during warm “abc” deformation.
Materials Science Forum Vols. 551-552 (2007) pp 183-188online at http://www.scientific.net© (2007) Trans Tech Publications, SwitzerlandOnline available since 2007/Jul/15
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Experimental Procedures
Commercially pure titanium (wt %: 0.25 Al, 0.15 Fe, 0.07 Si, 0.05 C, 0.12 O, 0.02 N) with a mean
grain size of 35 µm and alpha/beta titanium alloy Ti-64 (wt %: 6.3 Al, 4.1 V, 0.18 Fe, 0.03 Si, 0.02 Zr,
0.01 C, 0.18 O, 0.01 N) after water-quenching following beta-annealing with a mean β-grain size of
250 µm were used in the present work. The alloy had a phase transition temperature (at which
α+β→β) of 995°C.
Severe plastic deformation has been attained by means of successive compressions of a sample
along three orthogonal directions. The method is defined as “abc” deformation (Fig. 1a) [3]. The
initial dimensions of samples were 16×18×20 mm3 for both materials. The temperatures of the “abc”
deformation were 400°C for Ti and 550°C for Ti-64, respectively. Prior to each rotation, a prismatic
shape was restored to the sample by machining the curved faces. The initial strain rate and the true
strain per deformation step were 10-3s-1 and ~0.4, respectively. The true strain was calculated as
ln(h0/h), where h0 and h are initial and final height, respectively. Compression tests were performed
on a SCHENK dynamometer in air. Surface deformation relief at different steps of the “abc”
deformation was studied after additional compression of the samples by ε = 15%. The strain rate
sensitivity was estimated through strain rate changing tests at constant temperature. Apparent
activation energy was obtained as the slope of an Arrhenius plot (the log strain rate vs. 1/T).
The microstructure was examined by using an optical microscope, a JEOL JEM-2000EX
transmission electron microscope (TEM) and a JEOL JSM-840 scanning electron microscope (SEM).
Crystallographic analysis of structure was curried out by using EBSD-technique on a SEM LEO-440
Stereoscan.
Results and Discussion
Microstructure Evolution and Mechanical Behavior of Ti during “abc” Deformation.
Mechanical behavior of Ti under “abc” deformation at temperature of 400°С and strain rate of 10-3s-1
in terms of stress versus cumulative strain (S-Σε) is shown in Fig. 1b. The yield stress increases
noticeably for the first increments of the deformation while further straining just slightly increases
stress level. So that a stage likes steady state flow can be recognized on the final stages of the
cumulative S-Σε curve.
Microstructure evolution of Ti during the “abc” deformation is shown in Fig. 2. The initial stage of
plastic flow is associated with increasing of the dislocation density and formation of dislocation walls,
subboundaries and twins. Deformation induced dislocation boundaries slices initial grains making a
lamellar-type microstructure (Fig. 2a), which is typical for severely deformed metals [5]. Turning the
sample during the “abc” deformation involves new slip systems that leads to dividing of lamellae
(Fig. 2a). At the end of the “abc” deformation the microstructure consisted of equiaxed
grains/subgrains with a mean size of about 0.3 µm (Fig. 2b). Based on the EBSD analysis the number
of grain boundaries with the misorientation higher than 15° (high-angle boundaries) increases with
the strain from 10% after first increment to almost 70% at the end of the deformation (after Σε=6.2).
Misorientation distribution of Ti after “abc” deformation to Σε≈6.2 is shown in Fig. 2c. Twinning on
the {101} >< 612 system was found to be operating mainly at the initial stage of the “abc”
deformation. With strain increasing twinning at the {102} >< 211 and {113} >< 312 systems
becomes more active. The fraction of twins changes during the deformation from 1.3% after first
increment to 2.7% after Σε=6.2.
Changes in deformation relief correspond in general to those in the microstructure. At small strains
microshear bands and twins are observed within the initial grains. At larger deformation slip traces
being almost parallel each other with the length of about few hundred micrometers (while the grain
size is around 35µm) form at the surface (Fig. 3b). At strain above Σε = 2.7 two systems of parallel
bands with an angle of about 90° between them are revealed in the structure (Fig. 3c). The final
184 Superplasticity in Advanced Materials
structure after Σε = 6.2 has a smooth wavy relief however some intersecting bands are still observed
(Fig. 3d).
a b
Fig. 1. Schematic sketch of “abc” deformation (a) and cumulative S-Σε curve for “abc”
deformation of Ti at temperature of 400°C and strain rate of 10-3s-1 (b).
200nm200nm200nm
a b c
Fig. 2. Microstructure evolution of Ti during “abc” deformation at 400°C and 10-3 s-1: а – Σε≈2.5
(dark field); b - Σε≈6.2; c – misorientation distribution of Ti after “abc” deformation to Σε≈6.2
Values of the apparent activation energy and the coefficient of strain rate sensitivity defined at the
final strain of Σε = 6.2 were found to be 257 KJ/mol and 0.1, respectively. The value of the activation
energy is in agreement with that obtained by Conrad [9] and indicates that the plastic flow during
warm plastic deformation is controlled by prismatic slip of dislocations.
Evidently that the steady state flow stage on the cumulative σ-Σε curve is not related to
superplastic flow although the material has a very fine-grained microstructure. Some earlier results
[10] show that dynamic strain aging taking place in the same temperature interval (0.3-0.4 of melting
temperature) [9] may suppress superplasticity of the SMC Ti. This assumption is confirmed by
characteristics of superplasticity in SMC titanium at 450°C [10]. Based on the obtained results it can
be concluded that slip and twining are the only modes of plastic flow of Ti during the “abc”
deformation. The steady state flow stage is most probably associated with changing of the
deformation path. On the each next step of deformation some dislocation boundaries scatter to a large
number of movable dislocations decreasing yield stresses of the material (Baushinger effect) [11]. In
this case the strength of the material is controlled by substructure that, probably, leads to weak
dependence of yield stress on strain.
Materials Science Forum Vols. 551-552 185
a b c
d
Microstructure Evolution and Mechanical Behavior of Ti-64 Alloy during “abc” Deformation.
Mechanical behavior of Ti-64 in terms of stress versus cumulative strain (S-Σε) at 550°C and 10-3 s-1
(Fig. 4a) is described by a curve with flow softening following a peak stress during the initial
increments of the “abc” deformation and steady-state flow stage at larger strain. It is of note that first
two steps of the “abc” deformation have their own S-ε curves (within cumulative S-Σε curve) with
flow softening following a peak stress as well. During later steps, the peak stresses of included S-ε
curves decreased and steady-state flow is shown to occur at each of the steps.
Mechanical behavior of the alloy is strongly associated with structure evolution. An increase in
dislocation density at the beginning of deformation induces the stress peaks in the S-Σε curve.
Depending on the initial colony orientation either transverse sub-boundaries forms within in α- and
β-phases by accumulation of dislocations (Fig. 4b) or lamellae bends through the shear strain. The
shear deformation kinked lamellae and induced formation of transverse intra phase boundaries with
misotrientation varying from low- to high-angle [12]. Simultaneously bending of the lamellae rotates
parts of plates toward their lower slip resistance. Reorganization of dislocations in sub-boundaries
and rotation of lamellae (due to both dislocation sliding and micro shear bands activity) to more
favorable orientation make the alloy softer (Fig. 4a).
At further deformation the misorientation of intra-phase sub-boundaries increases to the high angle
range due to absorption of dislocations. Concurrently, the semi-coherent interphase α/β boundaries
transforme to non-coherent ones [13]. Then grooves form on the surface of the α-plates, leading to the
segmentation of the alpha plates (Fig. 4b). The fragmented β-interlayers and α-plates then
spheroidize. Changing of loading directions during the “abc” deformation accelerates globularization
of lamellae because various slip systems operate at each step.
As the globularization process proceeds, plastic flow via grain-boundary sliding (accommodated
by glide and climb of dislocations within the phases) became activated and steady flow stage is
revealed to occur in the S-Σε curve. Strain-rate-sensitivity coefficient measured at different points of
the S-Σε curve was found to be increased from 0.17 at Σε = 0.4 to 0.35 for Σε = 0.9. The value of the
apparent energy activation was found to be 186 kJ/mol, that is in good agreement with the results
obtained in [14] (182 kJ/mol) for SMC Ti-64 under superplastic conditions. All these results confirm
grain boundary sliding at the late stages of the “abc” deformation.
Figure 3. SEM images of surface relief on the titanium
sample deformed at 400°C and 10-3 s-1 to: а - Σε = 1.6; b –
Σε = 2.7; c - Σε = 6.2. The deformation axis is vertical.
186 Superplasticity in Advanced Materials
Cumulative Strain Σε
0 1 2 3
True Stress, MPa
400
600
800
a b c
Figure 4. Cumulative S-Σε curve for “abc” deformation of Ti-64 at 550°C and 10-3s-1 (a); TEM
images of Ti-64 alloy deformed to height reduction of 50% (b); to Σε = 3 (c)
Conclusions
1. The “abc” deformation of Ti at 400°C and 10-3 s-1 to a cumulative strain of Σε = 6.2 results in
formation of submicrocrystalline structure with a grain size of ~0.3 µm. Under “abc” deformation the
submicrocrystalline grains are formed as a result of formation and crossing of deformation induced
high angles boundaries created by slip and twinning. At the final steps of deformation the plastic flow
close to a steady state one is observed. The stability of plastic flow does not relate to superplastic flow
and is maintained by accommodation of microstructure to the strain path changing.
2. The “abc” deformation of alpha/beta Ti-64 titanium alloy at 550°С and 10-3 s-1 to a cumulative
strain of Σε = 3 leads to the formation of a homogenous submicrocrystalline structure with a grain size
of ~0.4 µm. A peak flow stress followed by flow softening is associated with globularization of the
structure. Steady-state flow observed on the final stages of the cumulative stress-strain curve occurs
due to superplastic flow. The apparent energy activation was found to be 186kJ/mol that is in agreement
with published data.
Acknowledgements
This work was supported by RFBR Project #05-08-65396, ISTC Project #3184 and Presidium of RAS
Programme П-08 (Subprogram: "Creation of novel structural and functional materials on the basis of
nanotechnology"; coordinator: academician N.P. Lyakishev).
Reference
[1] O.A. Kaibyshev: Superplastisity of Alloys, Intermetallides and Ceramics (Springer Verlag Berlin
Heidelberg, 1992).
[2] G.A. Salishchev, O.R. Valiakhmetov, R.M. Galeev: J. Mater. Sci. Vol. 28 (1993), p. 2903.
[3] S.V. Zherebtsov et al.: Scripta Mater Vol. 51 (2004), p. 1151.
[4] V.V. Rybin, Large plastic strains and fracture of metals (Metallurgy: Moscow, 1986) (in Russin).
[5] D.A. Hughes and N. Hansen: Acta Met. Vol. 45 (1997), p. 3886.
[6] P.J. Hurley and F.J. Humphreys: Acta Mater. Vol. 51 (2003), p. 1096.
[7] M.R. Barnett and F. Montheillet: Acta Mater. Vol. 50 (2002), p. 2296.
[8] S.L. Semiatin, V. Seetharaman and I. Weiss: .Advances in the Science and Technology of Titanium Alloy
Processing (Warrendale, PA: The Minerals, Metals and Materials Society, 1997).
[9] H, Conrad: Prog. Mater. Sci. Vol. 26 (1981), p. 403.
Materials Science Forum Vols. 551-552 187
[10] S.P. Malysheva et al.: The Physi cs of Metals and Metallography Vol. 95 (2003), p. 105.
[11] G.A. Salishchev et al. in: Nanomaterials by Severe Plastic Deformation – NANOSPD2, edited by
M. Zehetbauer and R. Valiev, volume 2, Kluwer Academic Publishers (2002).
[12] BielerT.R., Semiatin S.L., Int. J. Plast., 18 (2002), p. 1165
[13] M.I.Mazurski and G.A. Salishchev: Phys. Stat. Sol. (b) Vol. 188 (1995), p. 658.
[14] A.V. Sergueeva et al.: Scripta Mater. Vol. 43 (2000), p. 825.
188 Superplasticity in Advanced Materials