PACS numbers: 61.66.Dk, 61.72.Ff, 61.72.Hh, 62.20.M-, 68.37.Hk, 68.55.jd, 81.05.Bx, 81.10.Fq, 81.20.Hy, 81.40.Ef, 81.40.Jj, 81.40.Np, 81.70.Bt, 83.50.Uv, 83.60.Pq
o.M. iVaSiShin 1, S.V. akhonin 2, D.g. SaVVakin 1, V.a. berezoS 2, V.i. bonDarChuk 1, o.o. StaSYuk 1, and P.e. MarkoVSkY 1
1 G.V. Kurdyumov Institute for Metal Physics, NAS of Ukraine, 36 Academician Vernadsky Blvd, UA-03142 Kyiv, Ukraine
2 Ye.O. Paton Electric Welding Institute, NAS of Ukraine, 11 Kazimir Malevich Str., UA-03150 Kyiv, Ukraine
effect of MicrostrUctUre, DeforMation MoDe anD rate on Mechanical behaVioUr of electron-beaM MelteD ti–6al–4V
anD ti–1.5al–6.8Mo–4.5fe alloys
two commercial costefficient titanium alloys—a lowalloyed α+βti–6Al–4V (mas.%) and a metastable βalloy ti–1.5Al–6.8Mo–4.5Fe melted with a single electronbeam cold hearth melting approach—are employed in a present study as program materi
als. the influence of microstructure formed by means of the subsequent thermo
mechanical and heat treatments on both the mechanical behaviour (evaluated by the deformation energy, UD) when tested using standard methods with different deformation rates and the ballistic resistance of plate materials is investigated. As revealed, the weakest dependence of UD on the strain rate corresponds to ti–6Al–4V alloy with microstructure morphology close to globular one, whereas in the same alloy with a lamellar microstructure (annealed in a singlephase βfield) as well as in ti–1.5Al–6.8Mo–4.5Fe alloy with various microstructures, UD values significantly depend on the rate of deformation. Moreover, only after annealing at a temperature of the twophase α+β field, the UD value for ti–6Al–4V upon threepoint flexure at varying deformation rate is equal to the sum of the corresponding values obtained under tension and compression. During subsequent ballistic tests of plates of differ
ent thicknesses, it is established that the best ballistic impact resistance corresponds to the materials with higher UD values and lower strainrate sensitivity. the dam
age of alloys during all types of testing is discussed in detail in terms of possible mechanisms of deformation and fracture.
Keywords: titanium alloys, microstructure, mechanical properties, mechanical be
haviour, deformation rates, deformation energy, ballistic tests.
introduction
the titaniumbased alloys and coatings are a unique class of structural materials for many applications, and first of all aerospace and military, due to their high specific strength, fatigue endurance, resistances to crack growth and corrosion [1–6]. In some cases, these materials are employed in parts working at high stresses and tough operating condi
tions, particularly, in parts subjected to highrate deformations and high strengths. For instance, similar service conditions will take place in the new generation of civil supersonic aircraft [7, 8] as well as in a wide range of products for military use [9, 10]. Application of titanium alloys as armour material is an extreme case of their use from the view
point of mechanical properties. there is a large number of published up to now works on the ballistic impact behaviour of different titanium al
loys tested with different types of projectiles [9–19], but only few of them paid some attention to different types of microstructure, and no unambiguous relationships between microstructure and ballistic proper
ties have been revealed [15, 18]. At the same time, in some previous works, we showed the determining role of phase composition and micro
structure in the mechanical behaviour of various titanium alloys tested with different deformation rates (quasistatic range) on tension, compres
sion and threepoint bending [20–22]. earlier [23, 24], ti–6Al–4V (wt.%) alloy produced via singlemelt electronbeam with cold hearth melting (ebCh) and then milled to assigned thickness without any final anneal
ing was recognized as the most appropriate titanium armour material from the viewpoint of price/properties (or reliability) ratio.
the goal of this work was to evaluate real possibility to improve mechanical behaviour under deformation with different rates, and, fi
nally, to enhance ballistic properties of two titanium costefficient al
loys ti–6Al–4V and ti–1.5Al–6.8Mo–4.5Fe (wt.%) (socalled ‘low cost beta’) produced with abovementioned melting technique. to achieve this goal, we used various types of subsequent heat treatment, which allowed varying the structural and phase state of the alloys. With the aim to evaluate possibility to obtain the best ballistic impact protection properties, a special surface rapid heat treatment (Srht) was employed for the both alloys that allowed forming gradient microstructure with strengthened surface layer combined with a ductile core [25–27].
Materials and experimental Procedure
two above mentioned alloys were melted with ebCh melting approach in the form of 100 mm diameter and about 200 mm long ingots, using titanium sponge and Al–V master alloy (for ti–6Al–4V) or pure Al, Fe and Mo (for ti–1.5Al–6.8Mo–4.5Fe). the ingots were melted at labora
tory electronic beam unit Ue208 [28] by single melting. real chemical
compositions of melted ingots (averaged for three different locations of ingots) were as follow: (i) ti–5.8Al–3.86V and (ii) ti–1.6Al–6.3Mo–
4.36Fe (wt.%). the ingots were subjected to superficial turning and subsequent 3D pressing with total reduction of about 60% at single
phase βtemperatures: starting temperature was 1100°C and 900°C for the ti–6Al–4V and ti–1.5Al–6.8Mo–4.5Fe, respectively. In more de
tail, the melting and thermomechanical processing of ti–6Al–4V and ti–1.5Al–6.8Mo–4.5Fe alloys were described in [29] and [30], respec
tively. After hot pressing, both alloys were rolled in standard mills at temperatures of twophase α + β field: ti–6Al–4V at 880 °C to the final thicknesses of 23 mm and 11 mm, and ti–1.5Al–6.8Mo–4.5Fe at 750 °C to the thickness of 11 mm. Some rolled plates were annealed at different temperatures of twophase α + β and singleβphase fields with aim to obtain different types of microstructure. the Srht was performed at special induction heating device developed at the G.V. Kurdyumov Insti
tute for Metal Physics of the National Academy of Sciences of Ukraine [31]. All types and regimes of treatments employed in the present study are listed in table 1. their application allowed us to obtain a wide range of structural conditions, which varied in phase composition, type and fineness of microstructure, and to form gradient states with hardened surfaces combined with a ductile core.
Table 1. Treatments and microstructural states
## type of treatment regime of treatment Microstructure ti–6Al–4V
1 hot deformation 3D hot pressing + rolling
at 880 °C laminated, not recrystal
lized, unstable 2 Annealing
in the α + βfield 880 °C, 1 h,
furnace cooling laminated, not recrystal
lized, stabilized 3 Annealing
in the βfield 1100 °C, 1 h, furnace cooling
Coarsegrained lamellar 4 Surface rapid
heat treatment 80 kWt, 66 khz, 1000 A,
3 mm/s + aging 550 °C, 6 h Gradient: hardened sur
face (UtS ≥ 1300 MPa [29]), laminated core ti–1.5Al–6.8Mo–4.5Fe
5 hot deformation 3D hot pressing + rolling at
750 °C laminated, not recrystal
lized, unstable 6 Annealing
in the α + βfield 750 °C, 1 h,
furnace cooling elongated βgrains, globu
lar α, stabilized 7 Quenching
from βfield 900 °C, 0.5 h,
water quenching Singlephase metastable β 8 Surface rapid
heat treatment 80 kWt, 66 khz, 1000 A, 2.5 mm/s + aging 538 °C, 6 h
Gradient: hardened sur
face (UtS ≥ 1550 MPa [30]), laminated α core
Microstructure of alloys after different treatments was examined with light microscopy (lM) and scanning electron microscopy (SeM).
the SeM was also employed for investigation of fracture surfaces of specimens after mechanical and ballistic tests. tensile properties were studied in accordance with the AStM e8M standard on specimens with gage length of 25 mm and gage diameter of 4 mm at INStroN3376 unit at different strain rates (quasistatic regimes). threepoint flexure tests were done on bar samples with dimensions of 60×8×8 mm at IN
StroN 8802 unit. All specimens for mechanical tests were cut in roll
ing direction, and, at least, three specimens were tested for each micro
structural state and deformation rate. ballistic tests were performed in the certificated laboratory at the Ivan Chernyakhovsky National Uni
versity of Defence of Ukraine using ballistic barrels and projectiles list
ed in table 2. the bullets had different calibres (diameters) and steel cores (nonhardened soft (lPS) with hardness comparable to the hard
ness of tested titanium alloys, i.e. about 300–350 hV, or thermally har
dened (PP and b32) with hardness above 750 hV). Different bullet masses and speeds allowed varying specific kinetic energy during tests.
ballistic tests were performed in accordance with Ukrainian Standards b41032002 and 39752000.
results
Microstructure
typical microstructures of the alloys after different processing routes are presented in Figs. 1 and 2. As seen from Fig. 1, hot deformation of cast ti–6Al–4V alloy did not lead to recrystallization, and microstruc
ture was characterized by primary βgrains elongated in rolling direc
tion with fine αlamellas grown inside from grain boundaries (Fig. 1, a).
the thickness of these ‘flattened’ grains varied within the range of 20–100 µm due to intense and deep plastic deformation of initial cast structure, in which the βgrain size reached several tens (or even hun
dreds) of millimetres [29]. Further annealing at 880 °C (this tempera
Table 2. Characteristics of projectiles used in ballistic tests
## Designation bullet
Caliber, mm bullet
Mass, g type of core Average
speed, m/sec Specific kinetic energy, J/mm2
1 PP 5.45 3.61 hardened steel,
ogival shape 915 ± 18 65 ± 4
2 lPS 7.62 9.62 Soft steel,
ogival shape 840 ± 27 78 ± 5 3 b32 7.62 10.4 hardened steel,
ogival shape 845 ± 20 82 ± 5
ture cor responds to the α + β field) did not changed both grain and in
tragrain microstructure in asrolled material (compare (Fig. 1, a and b), but some βgrain boundaries (obviously, those with specific crystallo
graphic orientation) were decorated by lamellar (or platelike) αphase.
this allows concluding that the density of defects introduced by rolling was not enough to initiate recrystallization even on subsequent anneal
ing at temperatures of α + β field. the increase in annealing tempera
ture up to 1100 °C (singlephase βfield for this alloy) naturally caused intense βgrain growth, and relatively coarse and inhomogeneous βgrain microstructure was formed with grain size varying from 50 to 500 µm due to inheritance of nonuniformity of initial microstructure after an
nealing (Fig. 1, c).
After Srht, the alloy had layered structure (Fig. 1, d), in which nearsurface upper layers had thickness of about 4 mm and consisted of
Fig. 1. Microstructure of ti–6Al–4V alloy in different states: a — asrolled, b — after annealing at 880 °C for 2 h, c — after anneal
ing at 1100 °C for 1 h, and d — after Srht with lM
βtransformed microstructure formed as a result of α + β → β trans
formation during rapid heating and subsequent β → α (α′) + β decomposi
tion upon cooling [32].
ti–1.5Al–6.8Mo–4.5Fe alloy after hot deformation had laminated microstructure without elongated βgrains (Fig. 2, a; initial cast micro
Fig. 2. Microstructure of ti–1.5Al–6.8Mo–4.5Fe alloy in different states: a — after rolling and annealing at 750 °C for 2 h; b — after quenching from 900 °C into water (metastable βstate); (c–e) — after Srht, where c — overall view, d — surface layer, e — core. here, a, d, e — SeM; b, c — lM
structure of this alloy was discussed in detail earlier [30]). Subsequent annealing at 750 °C, 2 h (twophase α + β field for this alloy) caused par
tial transformation of intragrain αphase from lamellar to globular one (Fig. 2, b). It is important to underline that, in both cases, βgrain boundaries were decorated by lamellar αphase, which is generally con
sidered as a structural element that embrittles titanium alloys [2]. the application of Srht to this alloy, like in above mentioned case, also led to formation of gradient microstructure (Fig. 2, c): the top layer (about 40 mm) undergone α + β → β transformation and consisted of βtrans
formed grains with diameter of 20–40 µm (Fig. 2, d), whereas deeper layers had initial fine globularlike microstructure (Fig. 2, e).
standard Mechanical tests
In present work, we did tensile and threepoint flexure tests, while for the discussion and evaluation of the influence of deformation mode and rate on the mechanical behaviour of studied alloys we also used earlier results of compression tests [22].
tensile tests
the mechanical behaviour of both alloys under tension in dependence on their microstructure and tension rate was studied and discussed in de
tail in the previous work [21], including ti–6Al–4V and ti–1.5Al–
6.8Mo–4.5Fe alloys produced by standard vacuum arc (VAr) melting and multistep deformation routes. It is necessary to underline once again that, in the present study, we employed singlemelt ebCh alloy as start
ing material and subjected it to thermomechanical treatments described in the second section (‘Materials and experimental Procedure’).
Fig. 3. typical engineering stress–strain curves of ti–6Al–4V alloy tested at diffe
rent strain rates after 1 h annealing at 880 °C (a) and 1100 °C (b)
Ti–6Al–4V Alloy. typical en
gineering stress–strain curves of ti–6Al–4V alloy were character
ized by sharp transition from elas
tic to plastic deformation (Fig. 3) that is not typical for this alloy ob
tained by VAr melting [2, 29]. Af
ter α + β annealing (Fig. 1, b), the material had better combination of strength and ductility (Fig. 3, a) as compared to annealing in the single phase βfield (Fig. 3, b) which resulted to coarsegrained lamellar micro
structure (Fig. 1, c). Increase in strain rate led to reduction of elonga
tion for both structural conditions, but, in the latter case of coarse
grained lamellar microstructure, this reduction was bigger. After Srht, the material was not tensile tested because it was impossible to machine cylindrical samples adequately to obtain symmetrical twolayered mate
rial. however, in the earlier works, it was shown that surface treatment of this alloy allows improving tensile properties up to the level similar to that after bulk rapid heat treatment when tested with some standard tension rate [27].
Calculated dependences of UD on tensile strain for different strain rates (Fig. 4) showed essential difference between these two structural states. First, it should be underlined that, in the material annealed at the temperatures of twophase field, the UD value was higher and had less pronounced dependence on strain rate as compared to the alloy an
nealed in the singlephase field.
Ti–1.5Al–6.8Mo–4.5FeAlloy. the influence of tension rate on me
chanical behaviour of this alloy has been studied in detail in our previ
ous works [20, 21, 34]. In the present study, we took as a program material a newly melted by singlemelt ebCh alloy which was subjected to hot pressing and rolling as described in ref. [30]. this program mate
rial had specific microstructure features (Fig. 2) and crystallographic texture as compared to earlier studied tIMetAllCb alloy of the same composition, but melted by standard VAr technology [2] and subjected to much heavier total deformation. typical engineering stress–strain curves for our material are presented in Fig. 5. It is interesting to note that, in both principally different conditions—twophase α + β fine glob
ular (Fig. 2, a) and singlephase metastable β (Fig. 2, b), this alloy had very similar stress–strain curves at the same strain rates (compare curves 1 and 2 on Fig. 5, a, with curves 1 and 3 on Fig. 5, b). this fact Fig. 4. Dependences of deformation en
ergy on tension rate for ti–6Al–4V alloy in different structural conditions: an
nealing at 880 °C (1) and 1100 °C (2)
is unexpected and may be explained by very fine microstructure of the alloy annealed at 750 °C.
the dependences of UD on strain rate, calculated by tensile curves for these two microstructural con
ditions, are presented in Fig. 6. they also showed similar decrease in deformation energy with deformation rate. Such dependence is rather typical for this alloy [20, 21, 34], and was explained by factors: (i) high density of interphase α/β boundaries, which are sites of plastic defor
mation localization and nucleation of pores and cracks [34, 35] in an
nealed condition, and (ii) by specific crystallographic texture in metast
able condition, as it will be shown later.
A comparison of two materials from the viewpoint of dependences of deformation energy on strain rate allows concluding as follows below.
the ti–1.5Al–6.8Mo–4.5Fe alloy in both annealed α + β and quenched for metastable βphase states, despite higher initial (at 10−3 s−1) levels of UD, is more sensitive to deformation rate as compared to ti–6Al–4V alloy annealed in the twophase α + β field.
Fig. 6. Deformation energy vs. strain rate for ti–1.5Al–6.8Mo–4.5Fe alloy in two conditions: 1 — annealed in two
phase α+βfield; (2 — quenched from βfield (metastable singlephase condi
tion)
Fig. 5. typical engineering stress–strain curves of ti–1.5Al–6.8Mo–4.5Fe alloy tested at various strain rates in different microstructural states: a — after anneal
ing 750 °C for 1 h; b — after water quenching from 900 °C
three-Point flexure
Ti–6Al–4V Alloy. the overviews of specimens after threepoint flexu
ral tests are presented in Fig. 7. As seen, the material with coarse
grained lamellar microstructure cracked on earlier stages of plastic de
formation (specimens 1 and 2), while, in the specimen with globular microstructure, crack appeared on later stage of plastic flow when de
formed at lower bending rate (specimen 3).
the flexural stress–strain curves illustrate essential superiority of the material annealed in the twophase α+β field over the coarsegrained lamellar condition in both strength and especially ductility: the lami
nated microstructure after α+β annealing had approximately 6 times larger plastic deformation (compare Fig. 8, a and b). Gradient microstructure after Srht (Fig. 1, d) led to specific changes in flexure properties of the alloy: the curves look like a combination of the first two versions—
strength level approached to the α+β annealed condition, while strains were intermediate between these two structural states (Fig. 8, c).
Investigation of microstructure in the zones of intensive plastic deformation allowed clarifying the features of the influence of micro
structure on mechanical behaviour during flexural testing (Fig. 9). In the case of coarsegrained lamellar microstructure, main crack propa
gated from the bottom surface (where tension stress on flexure was maximal) almost normally to this surface with rather small branching (Fig. 9, a). the starting short crack was nucleated and grew initially along a packet of αlamellas (Fig. 9, b). At further, propagation crack was deviated on boundaries of neighbouring αphase packets with dif
ferent crystallographic orientations (Fig. 9, c).
At the same time, after annealing at temperature of twophase field, there were no coarse equiaxed βgrains and intragrain αlamellas, and the main crack initially grew at an angle of 45° to the bottom surface and changed the direction of growth when achieved length of approxi
mately 3 mm (Fig. 9, d). there was no noticeable effect of microstructure on crack propagation; however, at higher magnifications plenty of pores were observed around crack on the etched surface of the specimens (Fig. 9, e, f ). the specimens after Srht were not completely fractured at max
imal bend angle, and cracks stopped at distance of 2–3 mm from the thermally strengthened layer (Fig. 9, g): final cracks were shorter as
Fig. 7. overviews of ti–6Al–4V spe
cimens after (1, 2) βannealing (1100 °C, 1 h), and (3, 4) α + β annealing (880 °C, 2 h), tested for threepoint bending with rates (1, 3) 10−3 s−1 and (2, 4) 10−1 s−1
compared to other structural conditions for the same rates and degrees of deformation. In the ‘basic’ nonhardened layer, the main crack sev
eral times changed direction and propagated mainly along the bounda
ries of neighbouring αphase packets (Fig. 9, h), while on the last stage the crack crossed both these packets and individual αlamellas (Fig. 9, i).
Ti–1.5Al–6.8Mo–4.5Fe Alloy. At higher flexure rate, all ti–1.5Al–
6.8Mo–4.5Fe specimens, in both α + β annealed and as quenched from single βphase conditions, were broken on the stage of elastic deforma
tion (Fig. 10). only metastable β specimens tested at lower flexure rate had nonzero ductility; and this behaviour is opposite to ti–6Al–4V al
loy. Investigation of fracture surfaces of annealed specimens allowed distinguishing three separate zones: 1 — crack nucleation and initial growth, 2 — intensive growth zone, and 3 — final fracture (Fig. 11, a).
the overview of the first zone (indicated by arrow on Fig. 11, a, b) re
vealed that crack appeared in subsurface layer and initially grew by ductile mechanism. however, when crack length reached about 30 µm, the character of crack growth changed to a brittle one, and layered frac
ture surface was observed in this zone (2 in Fig. 11, a and c).
Fig. 8. Flexural stress–strain cur
ves after threepoint bending of ti–6Al–4V specimens after а — annealing at 1100 °C (coarsegrained lamellar microstructure), b — an
nea ling at 880 °C, and c — Srht.
Fle xure rates: 10−3 s−1 (1) and 10−1 s−1 (2)
Fig. 10. Flexural stress–strain curves for ti–1.5Al–6.8Mo–4.5Fe alloy in α + + β annealed and metastable β states tested with different rates
Fig. 9. Microstructure of ti–6Al–4V alloy after annealing 1 h at various tempera
tures and threepoint flexural testing: (a–c) 1100 °C (coarsegrained lamellar state), (d–f) — 880 °C, and (g–i) — after Srht. SeM data
Fracture surfaces of the specimens in singlephase metastable β state did not have evident signs of crack nucleation sites (Fig. 11, d).
both initial zone of crack growth (perpendicular to the outer surface of specimen) and zone of intensive propagation (parallel one) had the same character characterized by mixture of ductile and brittle (cleavage) ar
eas (Fig. 11, d) that was earlier observed in this alloy upon tensile tests [34] and explained by specific orientations of βgrain boundaries rela
tively to direction of applied tension stress. on the boundaries of cleav
age areas (presumably, grain boundaries), a relief similar to lamellar intragrain microstructure was observed (Fig. 11, f ). In the zone of final fracture, character of surface was ductile (Fig. 11, f ).
Summarizing the results of tensile and threepoint flexural tests, it is possible to conclude that only ti–6Al–4V alloy in annealed α + β state demonstrated ductile behaviour in both types of tests. this is obviously Fig. 11. Fracture surfaces of ti–1.5Al–6.8Mo–4.5Fe specimens in (a–c) α + β annealed and (d–g) metastab
le β states, after threepoint flexure with rate 10−1 с−1. Arrow (а) shows crack nucleation site. SeM data
due to comparatively large area of uniform plastic deformation (Figs. 3, a, and 8, b), while in another structural condition (Fig. 3, b) this alloy, as well as ti–1.5Al–6.8Mo–4.5Fe alloy (Figs. 5 and 10) underwent ear
ly localization of plastic deformation, that caused much earlier failure.
ballistic impact tests
Ti–6Al–4V Alloy: 11 mm Thick Plate. the 11 mm thick plate of ti–6Al–4V alloy was subjected to ballistic tests with two types of pro
jectiles, which differed in calibre, mass, state of steel core, speed of bullets, and therefore had different specific kinetic energies (see ## 1 and 2 in table 2). the overviews of tested plates with different micro
structural states are shown in Fig. 12. As seen, in asrolled (without final heat treatment) condition, the plate was pierced by 7.62 calibre lPS bullet, and withstood a bullet of 5.45 calibre (Fig. 12, a, b). the behaviour of plate annealed at twophase α + β field temperature was very similar (Fig. 12, c and d), while, after annealing at singlephase β temperature (with resulting coarsegrained lamellar microstructure,
βannealing at 1100 °C for 1 h, (e, f ) — after α + β annealing at 880 °C for 2 h, and (g, h) — after Srht (surface strengthening). here, front (a, c, e, g) and back (b, d, f, h) — views are presented. A — results of impact by bullets #1, B — results of impact by bullets #2
Fig. 12. overview of 11 mm thick ti–6Al–4V pla tes after ballistic tests with front (a, c, e, g) and back (b, d, f, h) views for dif
ferent conditions: (a, b) — as r olled, (c, d) — after
Fig. 13. SeM images for half sections of shot targets (a, d, g, j) and typical micro
structure of ti–6Al–4V alloy (11 mm thick plate) after ballistic tests: (a–c) — as rolled, (d–f) — annealed at 880 °C for 1 h, (g–i) — annealed at 1100 °C, and (j–l) — after Srht. PP bullets (#1 in table 2)
Fig. 1, c), the plate was pierced by bullets of both types (Fig. 12, e, f ).
In other words, 11 mm thick plates of ti–6Al–4V alloy with both lami
nated microstructural states showed rather good resistance against pen
etration of bullets with average specific kinetic energy ≈65 J/mm2, and could not stand against bullets with energy ≈78 J/mm2. At the same time, the plate with βtransformed coarsegrained lamellar microstruc
ture was not able to resists against even bullets with lower energy.
the most promising result for 11 mm thick plate was obtained in asrolled condition with subsequent special Srht; this plate was not pierced by both types of bullets (Fig. 12, g and h) that is in a good agree
ment with the results of threepoint flexure tests (Figs. 8, c, and 9, h, i).
In other words, the surface strengthening treatment allows obtaining a condition with higher ballistic resistance, and the present results are apparently better than the data obtained in J. Fanning’s work [9]. the Fig. 14. SeM images for half sections of shot targets (a, d, g) and typical microstructure of ti–6Al–4V alloy (11 mm thick plate) after ballistic tests: (a–c) — asrolled, (d–f) — annealed at 880 °C for 1 h, (g, h) — annealed at 1100 °C for 1 h. lPS bullets (#2 in table 2)
Fig. 15. overview of ti–6Al–4V alloy (23 mm thick plates) after ballistic tests with projectile #3 (table 2, bullets b32) for different conditions: (a, b) — after βannealing at 1100 °C for 1 h), (c, d) — after α + β annealing at 880 °C for 2 h, and (e, f ) — after Srht
peculiarities of plastic flow were studied on the sections cut from the sites impacted by bullets (Fig. 13).
In the asrolled specimen (Fig. 13, a), separate adiabatic shear bands (ASb) with thickness of about 10–12 µm were observed on the lateral sides of the bullet channel (Fig. 13, b), while, in the lower part of the crater, a βtransformed microstructure was found (Fig. 13, c; the depth of this βtransformed layer was about 2 mm). this may be explained by intensive local heating of the material above betatransus, i.e. an essen
tial part of kinetic energy transformed into a heat. Annealing at the temperature of α + β field made material ‘weaker’; it was pierced by this type of projectile (Fig. 13, d). Several ASb were observed near the mid
dle section of the bullet channel (Fig. 13, e), and, in some places, pores appeared inside the ASbs (Fig. 13, f ). After annealing at the tempera
ture of βfield, the material was also pierced, and coarse cracks appeared around the bullet channel where the bullet punched material (Fig. 13, g). In upper part of the channel, no evidences of plastic deformation were found (Fig. 13, h), while, in the bottom zone, multiple ASbs in
clined at an angle of more than 45° were observed (Fig. 13, i). After Srht, the plate was not pierced by bullet (Fig. 13, j), and ASbs paral lel to the surface were found near the middle section of the channel (Fig. 13, k) as well as in the ‘deep bottom’ — on the distance more than 3 mm from the channel tip (Fig. 13, l).
Figure 14 illustrates similar features of plates’ microstructures af
ter shooting by lPS bullets (#2 in table 2). First of all, it should be mentioned that all microstructures (except after surface strengthening by Srht) were pierced by this relatively highenergy projectile (see
table 1), and bullet channels looked like punching out a ‘cork’ (Fig. 14, a, d, g). Due to high localization of plastic deformation, the flow oc
curred in the material adjacent to the bullets channels, that led to the formation of cracks (Fig. 14, b, c), and only in some cases ASbs were observed (Fig. 14, f ). besides, it should be mentioned that, in most duc
Fig. 16. SeM images for half sections of shot tar
gets (a, d, g) and typical microstructure of ti–6
Al–4V alloy (23 mm thick plate) after ballistic tests:
(a–c) — after βannealing at 1100 °C for 1 h, (d–f) — after α + β annealing at 880 °C for 2 h, and (g–i) — after Srht
tile condition (after annealing at α + β temperature), some traces of plastic flow were found close to the bullet channel (Fig. 14, e), while, in less ductile one (after annealing at βtemperature; Fig. 14, h), fracture took place without any visible changes in microstructure. At the same time, in Srhttreated plate, which was not pierced (Fig. 12, g, h), the traces of deformation were similar to those observed for projectile #1 (Fig. 13, k, l).
Ti–6Al–4V Alloy: 23 mm Thick Plate. the overviews of specimens shot with projectiles b32 are shown in Fig. 15. As seen, the material after annea ling at the temperature of singlephase βfield with coarse
grained lamellar microstructure was completely pierced, and the bullet was stuck in the plate (Fig. 15, a, b). In two other structural states after annealing at α + β temperature (Fig. 15, c, d) and after Srht (Fig. 15, e, f ), the plates were not penetrated through, and the ‘bulge’ effect was noticeab ly less pronounced in the case of Srht. these observations are in a good agreement with the results obtained for 11 mm thick material after the same treatments, and they confirm the positive influence of both annealing at temperatures below the βtransus and surface strength
ening treatment.
(c, d, g, h) #2, and (e, f ) #3 for two conditions: (a–f ) — after α + β annealing at 750 °C for 2 h and (g, h) — after Srht streng thening
Fig. 17. overview of 11 mm thick plates of ti–1.5Al–6.8Mo–4.5Fe alloy after ballistic tests with projectiles (a, b) #1,
In more detail, the microstructure of the plates near the bullet channels is shown in Fig. 16. After annealing at 1100 °C, the material with coarsegrained lamellar microstructure was punched through (Fig.
16, a). there were no noticeable traces of plastic deformation near the edges of bullet channel (Fig. 16, b); separate ASbs were found near the bottom of the channel (where the bullet has already lost most of its en
ergy), and these ASbs end by surface cracks (Fig. 16, c).
After annealing in twophase α + β field, the material had another type of βgrain structure and was also pierced through; however, the bullet was completely stopped inside the plate (Fig. 16, d). Near the top of the bullet channel, some deformation traces were found (Fig. 16, e), and, in the bottom, ASbs and cracks were present as well (Fig. 16, f).
Fig. 18. half sections of shot targets (a, e, g), de
tails of fracture surfaces and typical microstruc
tu re of 11 mm thick pla
tes of ti–1.5Al–6.8Mo–
4.5Fe alloy after ballis
tic tests with projectiles
#1 (a, b), #2 (c, d, e, f ), and #3 (g, h) for two conditions: after α + β annealing at 750 °C for 2 h (a–d) and after Srht strengthening (e–h). here, the SeM data are presented
After Srht, the bullet did not penetrate through the plate, and a bulge of about 2–2.5 mm height formed on the rear side (Fig. 16, g). the ASbs were observed both on the top (Fig. 16, h) and bottom (Fig. 16, i) sections that indicates on the localization of plastic flow in both top (thermally hardened) and bottom (ductile) layers. In other words, hardened by Srht, layer had relatively good reserve of general ducti
lity, despite its higher strength, and its level of ductility was rather close to that of the material annealed at temperatures of α + β field, whereas the alloy with coarsegrained lamellar microstructure had too low properties.
Fig. 19. Deformation energy vs. deformation rate at different types of tests for a — ti–6Al–4V annealed at temperature of singlephase β field, b — ti–6Al–4V an
nealed at temperature of twophase α + β field, c — ti–1.5Al–6.8Mo–4.5Fe quenched from temperature of singlephase β field (finegrained metastable β state), d — ti–1.5Al–6.8Mo–4.5Fe annealed at temperature of twophase α + β field. Data for compression tests were taken from [23]. open ‘star’ symbols on (a, b) indicate re
sults with complete piercing of plates. Solid ‘star’ symbols (b, d) indicate results of ballistic tests without piercing
Ti–1.5Al–6.8Mo–4.5Fe. typical examples of 11 mm thick plates of al
loy ti–1.5Al–6.8Mo–4.5Fe in different microstructural states after bal
listic tests are presented in Fig. 17. First of all, it should be underlined that, after annealing in twophase α + β field, the material was not punched through by bullets of #1 type, however, some surface cracks appeared (Fig. 17, a and b), while higher specific kinetic energies (bul
lets #2 and #3) caused cracking (Fig. 17, c–f ). Application of Srht allowed somewhat to improve ballistic resistance of the alloy against projectiles #2 (Fig. 17, g and h), whereas bullets #3 still punched through the plates.
In more detail, the fracture surfaces of the plates after ballistic tests are shown in Fig. 18. After annealing in α + β field and impact by projectile #1, the bullet channel had relatively ductile relief (Fig. 18, a, b), while impact by more powerful projectile #2 led to completely brittle fracture (Fig. 18, c and d). Another situation was observed in specimens after Srht: despite impact with projectile #2 also led to brittle frac
ture (Fig. 19, g and h), projectile #1 caused only crater in the surface strengthened layer (Fig. 18, e and f ).
Discussions
Under threepoint flexure, the upper part of sample (where bending force is applied) is compressed and the opposite side is stretched, with a zeroload point located approximately in the middle between them.
hence, in the first approximation, we can assume that this type of test is a combination of stretching and compression tests. thus, it can be assumed that the energy of deformation UD, which is spent on the three
point bend, should be approximately equal to the sum of these energies for two tests—tension and compression. experimental data for compres
sion, obtained with the same program materials, were taken from the work [24]. the dependences of UD on deformation rate for the same ma
terials and different methods of tests are represented in Fig. 19.
As can be seen from Fig. 19, this assumption is rather accurate for the case of ti–6Al–4V annealed in twophase α+β field (880 °C, Fig. 19, a). It should be also mentioned that similar situation was observed in the same alloy with globular type of microstructure, and in commercial purity titanium regardless of microstructure [26]. At the same time, ti–6Al–4V and ti–1.5Al–6.8Mo–4.5Fe alloys annealed in singlephase βfiled (coarsegrained lamellar microstructure) have different behav
iour (Fig. 19, b–d). earlier in [26], this fact was explained by the influ
ence of other factors, including structural and crystallographic texture.
Also, as it was established in ref. [21], lamellar or needlelike αparticles play an important role in mechanical behaviour of titanium alloys de
formed with different rates, because their tips are places where defor
mational defects concentrate during plastic flow, leading to the forma
tion of pores and then cracks.
Another moment to be emphasized is the following: ti–1.5Al–
6.8Mo–4.5Fe in metastable singlephase βcondition had high ductility that resulted in high UD values under tension and compression, while for threepoint flexure this parameter was rather high only for rela
tively slow deformation rate (Fig. 19, c). Perhaps, this is a result of probable decomposition of metastable βphase during the complex mode of deformation (and possible precipitation of embrittling particles of ωphase, especially, taking into account local heating inside ASbs formed at high deformation rates). regarding this alloy in the twophase α + β state, UD values for threepoint flexure were much lower than for ten
sile and compression tests. this is obviously due to the peculiarities of microstructure (the presence of thin lamellar particles of αphase whose tips play abovementioned role of stress concentrators) and the crystal
lographic texture formed upon rolling, which may be unfavourably ori
ented relative to the applied load.
the results of present study and previous data allow us to conclude that the UD values of investigated alloys generally are determined by their phase composition, microstructure, and structural and crystallo
graphic texture, which, in turn, determine the balance between strength and ductility, and this balance depends on deformation mode and rate too. thereby, relatively lowalloyed ti–6Al–4V alloy with globular (equiaxed) microstructure has rather high values of ductility. however, hardening of this alloy by heat treatment cannot ensure high strength;
in the best case after special rapid heat treatment (bulk), UtS does not exceed 1300–1350 MPa [22, 27]. At the same time, metastable βalloy ti–1.5Al–6.8Mo–4.5Fe can be thermally strengthened above 1500 MPa, but high content of βalloying elements and relatively low betatransus (temperature of α + β0 → β transformation [3]) do not allow to form equi
axed microstructure with αglobules of, at least, 2–3 µm in diameter and relatively low (as in ti–6Al–4V alloy) density of interphase α/β boundaries which are preferable sites for pore and crack nucleation dur
ing plastic flow [21]. these features caused low ductility of thermally strengthened ti–1.5Al–6.8Mo–4.5Fe alloy, which may be improved by means of special treatment based on heavy cold deformation followed by rapid recrystallization for forming fine βgrains (below 5–7 µm) [22, 36]. Such technological approach was employed for rod materials, but its application for processing of plate material (taking into account the possibility of technological equipment) seems quite problematic. As shown in refs. [22, 37], the best combination of high strength (above 1800 MPa) and good ductility can be obtained by rapid heat treatment of such titanium alloy as Vt23 (highalloyed martensitic class alloy), where the content of βalloying elements is ‘intermediate’ between ti–
6Al–4V and ti–1.5Al–6.8Mo–4.5Fe. In this case, it is possible to expect that UD will be higher, including high deformation rates.
of course, the results of standard mechanical tests are difficult to compare with the results of ballistic ones. Nevertheless, employing such rather versatile parameter as UD and calculating (with some simpli
fications) its values for ballistic tests basing on the speed of the bullet, its calibre and the thickness of tested plate, we tried to fulfil this com
parison. though this approach is rather arbitrary, it allows to assess the ballistic resistance of materials considering the relation between UD values for the same deformation rates obtained in ballistic tests and, if possible, by other standard methods. In other words, if UDballistic >
> UDStandard, material will be pierced, and if UDballistic < UDStandard, material will resist against ballistic impact. An attempt to make such a compari
son is presented in Fig. 19, a, b, d, where the results of ballistic tests are plotted by ‘star’ symbols. First of all, it should be noted that an increase in strain rate by four orders of magnitude reduces UD levels also by several orders. Despite the significant difference between UD levels of studied alloys at low deformation rates, at rates 8·104–8.2·104 (these values were calculated for ballistic tests), deformation energy varied within the narrow range of 5.7–7.4 J/mm3, and difference be
tween pierced and not pierced microstructural states was also very small. It should be kept in mind that, in this comparison, we did not take into account the influence of all other factors besides kinetic en
ergy (e.g., the different hardness of bullets and test material). by the way, our results showed some superiority of alloys after annealing in twophase α + β field and Srht over the same materials with βtrans
formed lamellar microstructure.
conclusions
the obtained results allow drawing the following conclusions.
(1) the studied alloys obtained by ebMCh melting approach with following hot 3D pressing and rolling had better properties (both stand
ard mechanical at all deformation rates and modes, and ballistic im pact properties) after annealing at temperatures of twophase α + β field.
In this condition, ti–6Al–4V alloy had microstructure with flat β
grains elongated in rolling direction and fine αlamellas inside these grains, while ti–1.5Al–6.8Mo–4.5Fe alloy had fine globulartype microstructure.
(2) Annealing in singlephase βfield resulted in coarsegrained la
mellar microstructures and led to a significant decrease in all mecha
nical properties due to essential drop of ductility.
(3) Special Srhtreatment, which formed thermally strengthened surface layer in combination with ductile core, allowed achieving the
best level of ballistic impact properties as compared to other micro
structural conditions studied in the present work.
(4) Good antiballistic properties of ti–6Al–4V alloy can be ex
plained by a unique combination of properties (of about 1000 MPa) with high reserve of ductility (including high uniform plastic deformation in notcoarsegrained lamellar state). this properties balance can be ob
tained even after singlemelt ebCh melting followed by hot deforma
tion (rolling) without any final annealing. At the same time, ti–1.5Al–
6.8Mo–4.5Fe alloy, despite of ability to obtain very high strength (not less than 1500 MPa), is characterized by relatively poor ductility, and first of all, small stage of uniform plastic deformation with early start of localized plastic flow, especially, at increase of deformation rate.
(5) For a higher efficiency of surface hardening treatment in order to form a gradient microstructure with increased antiballistic proper
ties, attention should be paid to alloys with an intermediate (between ti–6Al–4V and lCb alloys) content of alloying elements. these alloys have two benefits: (i) a welldeveloped globular microstructure can be formed by annealing at α + β temperature; (ii) surface rapid heat treat
ment allows achieving strength levels ≥1500 MPa.
Acknowledgments. Parts of this work were done within the frames of the projects ## III0913, 60/13c, and 60/16п funded by the Na
tional Academy of Sciences of Ukraine, and supported by the NAto grant G5030 from the Program ‘Science for Peace and Security’.
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in final version, July 26, 2018
О.М. Івасишин 1, С.В. Ахонін 2, Д.Г. Саввакін 1, В.О. Березос 2, В.І. Бондарчук 1, О.О. Стасюк 1, П.Є. Марковський 1
1 Інститут металофізики ім. Г. В. Курдюмова НАН України, бульв. Академіка Вернадського, 36, 03142 Київ, Україна
2 Інститут електрозварювання ім. Є. О. Патона НАН України, вул. Казимира Малевича, 11, 03150 Київ, Україна
ВпЛИВ мІКРОСтРуКтуРИ, СпОСОбу тА шВИДКОСтИ ДЕфОРмАЦІї НА мЕхАНІчНу пОВЕДІНКу СтОпІВ Ti−6Al−4V І Ti−1,5Al−6,8Mo−4,5Fe, ВИгОтОВЛЕНИх ЕЛЕКтРОННО-пРОмЕНЕВИм тОпЛЕННЯм
На прикладі двох промислових економно леґованих титанових стопів, — мало
леґованого α + β-стопу ti–6Al–4V (мас.%) і метастабільного βстопу ti–1,5Al–
6,8Mo–4,5Fe, — виготовлених одноразовим електроннопроменевим топленням з проміжною ємністю, вивчено вплив формованої за подальших (термомеханічно
го та термічного) оброблянь мікроструктури на механічну поведінку (виражену через енергію деформації UD) при випробуваннях з різними швидкостями де
формації та балістичну стійкість. Встановлено, що найменша залежність UD від швидкости деформації відповідає стопу ti–6Al–4V, відпаленого за температури двофазної α + β-области, в той час як цей же стоп із пластинчастою мікрострук
турою (піс ля відпалювання в однофазній βобласті), як і стоп ti–1,5Al–6,8Mo–
4,5Fe у різних структурних станах, характеризується інтенсивним пониженням вели чини UD з ростом швидкости деформації. Окрім того, тільки відпалений за тем ператури двофазної α + β-области матеріял після випробувань на триточковий вигин характеризувався тим, що для кожної швидкости деформації величина UD дорівнювала сумі аналогічних характеристик, одержаних випробуваннями на розтягання та стискання. При наступних балістичних випробуваннях пластин цих стопів різної товщини було встановлено, що найліпша балістична стійкість відповідає матеріялам, які характеризуються великими значеннями UD та мен
шою її чутливістю до швидкости деформації. Детально обговорюється пошко
дження стопів при випробуваннях різних типів з точки зору ймовірних механіз
мів деформації та руйнування.
Ключові слова: титанові стопи, мікроструктура, механічні властивості, механічна поведінка, швидкості деформації, енергія деформації, балістичні випробування.
О.М. Ивасишин 1, С.В. Ахонин 2, Д.Г. Саввакин 1, В.А. Березос 2, В.И. Бондарчук 1, А.А. Стасюк 1, П.Е. Марковский 1
1 Институт металлофизики им. Г. В. Курдюмова НАН Украины, бульв. Академика Вернадского, 36, 03142 Киев, Украина
2 Институт электросварки им. Е. О. Патона НАН Украины, ул. Казимира Малевича, 11, 03150 Киев, Украина
ВЛИЯНИЕ мИКРОСтРуКтуРы, СпОСОбА И СКОРОСтИ ДЕфОРмАЦИИ НА мЕхАНИчЕСКОЕ пОВЕДЕНИЕ СпЛАВОВ Ti−6Al−4V И Ti−1,5Al−6,8Mo−4,5Fe,
пРИгОтОВЛЕННых эЛЕКтРОННО-ЛучЕВОЙ пЛАВКОЙ
На примере двух промышленных экономно легированных титановых спла
вов, — малолегированного α + βсплава Тi–6Al–4V (масс.%) и метастабильного βсплава ti–1,5Al–6,8Mo–4,5Fe, — приготовленных однократной электроннолу
чевой плавкой с промежуточной ёмкостью, изучено влияние формируемой при
последующих (термомеханической и термической) обработках микроструктуры на механическое поведение (выраженное через энергию деформации UD) при ис пы таниях с разными скоростями деформации и баллистическую стойкость.
Установлено, что наименьшая зависимость UD от скорости деформации соответ
ствует сплаву ti–6Al–4V с близкой к глобулярной морфологией микрострукту
ры, в то время как этот же сплав с пластинчатой микроструктурой (после от
жига в однофазной βобласти), как и сплав ti–1,5Al–6,8Mo–4,5Fe в различных структурных состояниях, характеризуется интенсивным снижением величины UD с ростом скорости деформации. Кроме того, только отожжённый при темпе
ратуре двухфазной α + β-области материал после испытаний на трёхточечный изгиб характеризовался тем, что для каждой скорости деформации величина UD была равна сумме аналогичных характеристик, полученных испытаниями на растяжение и сжатие. При последующих баллистических испытаниях пластин этих сплавов различной толщины было установлено, что наилучшая баллисти
ческая стойкость соответствует материалам, характеризуемым большими значе
ниями UD и меньшей её чувствительностью к скорости деформации. Подробно обсуждается повреждение сплавов при испытаниях разных типов с точки зрения вероятных механизмов деформации и разрушения.
Ключевые слова: титановые сплавы, микроструктура, механические свойства, механическое поведение, скорости деформации, энергия деформации, баллисти
ческие испытания.
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