TITANIUM'99: SCIENCE AND TECHNOLOGY

Mechanisms of Martensite Formation and Tempering in Ti(anium Alloys and Their Relationship to Mechanical Property Development

0. M. IVASISHIN *, H. M. FLOWER**, G. LUTJERING ***

• Institute for Metal Physics, Kiev· 252142, Ukraine •• Imperial College of Science, Technology arid Medicine, London SW7 2BP, UK ***·Technical University Hamburg-Harburg, 21071 Hamburg, Germany

ANNOTATION

Present paper is aimed to elucidate the mechanisms of formation of orthorhombic niartensites which occur over certain composition ranges in· titanium alloys. The experimental results available are well consistent with the explanation that the formation of the orthorhombic structure is being due to _the.elastic distortions imposed on the h.c.p. lattice as it. decomposes into two constituents having different solute contents and coherently bonded in a two dimensionally modulated structure. Calculations of fuli _interatomic interactions have been employed and predict·· bulk· modulation in · the · basial plane, different for isomorphous and eutectciid binary systems and favorable for the orthorhombic distortion in the former case. Mechanical properties of orthorhombic marterisite in beta isomorphous Ti-7%Mo alloy were determined as functions of the heat treatment condition. Possible explanations of a significant difference in martensite properties in as-quenched and aged conditions are discussed.

Key words: titanium alloys, martensite, crystal lattice, mechanical properties, solid solution, coherent stresses.

1. INTRODUCTION

Martensitic transformation in titanium alloys has been extensively studied for two main reasons. Firstly, a range of crystallography, morphology, and substructure of martensite product is observed in titanium alloys, resulting in an array of martensite related phenomena. Secondly, martensite is an important -microstructural component which may or may not be desirable but must be dealt with at various processing routes where rapid cooling is employed. Despite extensive experimental investigations, some features of martensite transformation in titanium alloys remain unclear. Among these are crystal lattice structure and mechanical properties of a"-martensite which occurs over certain composition ranges in some alloys, alloyed by isomorphous beta-stabilising elements. Orthorhombic symmetry in titanium alloy martensite, first identified in [I], is associated with unusual mechanical behavior (much lower strength and higher ductility, as compared to "normal" hexagonal a'-martensite, in the as­quenched condition followed by a pronounced strengthening during ageing. Unfortunately, the aged a"­martensite is very brittle making martensite type strengthening generally of no practical use. Explanations of

martensite orthorhombicity ( a,::;- * I ) in some titanium· alloys ·by ordering of high-temperature beta-phase [2] or · bv3

by incompleteness of b.c.c. ~ h.c.p. lattice reconstruction [ l, 3] were inconclusive. The mo·st recent explanation [4]-was based on the assumption that macroscale orthorhombicity derives from the coherent stresses arising from modulated decomposition of supersaturated martensite solid solution. The experimental background of this explanation was found in TEM observation of a fine modulated microstructure, supposedly of spinodal type, in martensites, which were identified as orthorhombic ~by X-ray diffraction [5]. Following this hypothesis, it may be suggested, that the crystal structure of titanium martensite, hexagonal or orthorhombic, is related to the stability of martensite solid solutions, specifically alloyed with various elements, with respect to the decomposition via mechanisms which are able to form composition modulations during quenchif!g and/~r following ageing. Calculations of interatomic interactions in substitutional solute atom subsystem to support this suggestion were performed in the present work. To further support the above explanation, experimental data on the crystal structure ofmartensite in as-quenched and aged conditions have been obtained and related to corresponding data on microchemical, inhomogeneity. Lastly, mechanical properties of an alloy with typical a"-orthorhombic martensite in as-quenched and aged conditions have been determined and explained on the base of

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TITANIUM'99: SCIENCE AND TECHNOLOGY

thennodynamic .and microstructural understanding of the nature of a"-martensite. The work was -perfonned under INT AS gr.ant 93-630, the results of which are overviewed in the present paper . . - ' . .. .•

2. CALCULATION ON STABILITY OF HEXAGONAL TITANIUM BASED SUBSTITUTIONAL SOLID SOLUTIONS

A microscopic approach to the calculation of interatomic interaction was used. A dis.ordered Ti-M;e solid so.lution was considered, in which Ti and Me atoms are randomly di*ibuted in an h.c_.p. )attic~ having. averaged parameters a

0 and c

0. Statistical-thennodynamic description of such a solution permits calculation of ~h_e energies

of both strain-induced (indirect) interaction which results from static deformation of the h.c.p. ionic subsystem and electrochemical ( direct) interaction between electron shells of Me ions. Details of the calculation are given in [6]. It should be only noted here that, according to [7], in order to predict the transfonnation behavior of the solid solution, it is sufficient to calculate the Fourier components w pp' (k) of the energy of "mixing" w pp' (r) for the pair

of atoms that occupy the sui:,stitutional sub lattices p and p'., The conclusion on whether the ~olid solution is stable or liab~e to d~composition depends.on:~\'hether the _non-analytic function wPP:(k)_within the first Brillouin zone

(BZ) is positive -or negative: If the latter,. the probable character of transformation can be predicted from the

location o~ t~e- aq_~olute ~~~imum o~ t_he .lower (ro=II-'.') branch of eigenvalues Aro(k) (ro="±") of the llw PP.(k)II . . . . . . . . . - .

matrix. If. the minimum falls on one of the symmetry points on the BZ surface, homogeneous. atomic ordering can occur through the_ solid. solution. _The locati~ns of the absolute minimum within the BZ predicts a, periodic modulation in .the or_dered. structure, the spatial period and direction of modulation being detennined by the position of the minimum. ·A minimum at the, BZ- centre predicts spinodal decomposition of the solid sol~tion. In Figure 1 the calculated dependencies of Aro(k). along 17 main directions ~e pres~nted_ for_ Ti~Fe ,and Ti-Mo

. A, (k) eV

0,60

0,40

0,20

0, 00 -l';--Ml'---/t-1'"T7"t'rj----f-tt---t7'~m-.,....-!,ci

-0,20

-0,40

r MK rA LH A· ML r HK.

a)

_ A± (k) e_V , 10,00

5,00

-5,00

r MK [A LH A ML r HK

b)

Figure 1. The eigenvalues, A+(k) {upper curves) and A_(k) (bottom curves), of Fourier components'. matrix of

the «mixing» energy of atoms in h.c.p. solid solutions (a) Ti-Fe and (b) Ti-Mo. The values of J\:+(O) 'and A.(O) are

marked (x) and(*), respe~tiv~ly. Points r, A, H, K, L, M corresponds to centre ~d main symmetry points of BZ. . ~. ' . .

alloys. It is seen from_ minima availability _and locations that. calculations predict a~omic ordering with a modulation in the basal 'plane for both alloys. However, substitutional interaction in the eutectoid and isomorphous type alloys is different since the interactiop matrix eigenvalues lie along different directions: rK an.dTM respectively, As was shown in [8, 9] modulations along the rM direction do not change the hexagonal symmetry of the solid solution while that along rK leads to its orthorhombic distortion. This is a first- indication that isomorphous type alloys, in contrast to eutectoid type are prone to form orthorhombic solid solutions if composition modulation occurs. A second consideration derives from the fact that the above calculations relate to the very initial stages of decomposition. While this is continuing, the "electrochemical" contribution is being decreased which may result in the strain induced part.of atomic interaction det\;!nnining the instability type. In this case, for isomorphous type alloys an absolute minimum of interaction energy falls exactly in the BZ centre (Figure 2a) indicating on the possibility of spinodal decomposition of the solid solution, with a modulation direction in the basal plane. This is just the case where coherent stresses between the solute rich and solute lean constituents of the initially homogeneous hexagonal solid solution may provoke an orthorhombic distortion, as it was explained in [4]. Comparison ·of Figures· 2a and 2b· shows that in eutectoid type alloys spinodal decomposition is not allowed; the location· of the minimum in (0, 0, 1/2) might result in a short0 period modulation in the c direction, predicted also in [ l O].

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TITANIUM'99: SCIENCE AND TECHNOLOGY

A,(k) ~V o.so . ,

.-o.so

-1.00

r MK r A LH A ML r HK MK-fA LH A ML r HK

a) . b).

Figure 2. The eigenval~es, A+(k) (upper c·urves) and A.(k) (bottom cµrves), offourier components' matrix of the

strain-induced contribution into «mixing>> energy of atoms in h.c.p. solid solutions (ti) Ti-Fe and (b) Ti-Mo. The u ~ • - - • •• •

values of A+(O) and A.(O) are marked (x) and(*), respectively.

To detennine at which'temperatures solid solutions of a given compos1t10n are· losing sta~ility, a criterion proposed in [7J for cubic solid solutions was modified in [8J for h.c.p.- sohitions:

The results of calculations shown in Figure 3, clearly- indicate· again the significant: difference between the isomorphdus and .eutectoid type alloys. The fonner lose stability at relatively high:temperatures while the latter are thennodynamically stable to such. low temperatures that their decomposition, even if allowed, should, be regarded as impossible because.oflow diffusivity of solute atoms.

·r;K 1600

1200

800

400

1 2

5 6

0.02 0.04 0.06 o.oe 0.10 0.12 CM,

Figure 3. Stability boundaries for the h.c.p. solid solutions: (I) Ti- Ta, (2) Ti-V, (3) Ti-Nb, (4) Ti-Mo, (5) Ti-Fe, and (6) Ti-Cr.

(OIIO] -.1 /(0001] _

t:.21101 .

·i'"": -,--~

Figure 4. Two-dimensional modulated structure in h.c.p. solid solution.

3. X-RAY DIFFRACTION PICTURE OF a" MARTENSITE

An explanation of X-ray diffraction data for a" mart~nsite given in [4J was based on the analysis of the coherent stresses ~ising in the ~wo-dimensional pe-riodic macrolattice of the decomposed. hexagonal solid solutions, in which square cross section. r_ods of solute lean and solute rich phases are alternated with rectangular cross section rods of close to mean composition (Figure 4). In contrast to cubic solid solutions considered in [7J, in which the modulation directions are fixed, in hexagonal martensite the modulation directions (habit planes) [n

1 n

2 OJ and

[ n I n

2 OJ change during the decomposition. Calculations undertaken in [ 4 ]:has shown that macro lattice constituents,-

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TITANIUM'99: SCIENCE AND TECHNOLOGY

fonnerly hexagonal, can gain an orthorhombic or even rhombic distortion leading to a modification of X-ray data in such a way that (hid) diffraction maxima (I= 0, ±2, ±4, ... ) split, as shown in the Figure 5. It is worth noting that although stress induced splitting of each peak is negligible, the "averaging" of the intensities of all the diffraction to the left and right sides from the main diffraction peak of hexagonal phase creates an illusion of an orthorhombic lattice with a significant orthorhombicity calculated on the assumption that all the left-side and all the right-side diffraction intensity define the overall splitting in the orthorhombic doublet. The net result is that this imaginary orthoi"hombicity is an order of magnitude higher than what happens in the reality to the constituents of the decomposed martensite. Certainly, such an orthorhombicity which is often used as a

[200]11 a

[200] ,. [110] [130] ,. I . ()I HCP . Cl

t [200]u:1" [130]111' ' [130]11 [130]111•• [2001111' [200]1

""' ""' ""' ""' ""' ""' ""' :ii!E :ii!E :ii!E :ii!E :ii!E :ii!E . ~ :ii!E :ii!E ::; ::; ::; ::; . ::; ::; :ii!E :ii!E ::; :!I: :ii!E :ii!E I I ·I ::; :ii!E !:' ::; j;-

31° 32°

[200 ]11 b

[200]., C110lecr· C130la'' '° C200llzz" C130l111•· ' [130]11 [130]111•• [200h11' [200lz C 130]1

112 ~ ii ii ii ii ii ii ~· I ::; ::; ::; ~ :ii!E :ii!E I ji II ~ iiili !1i1 :ii!E ::; ii ii -iiri !ii I ~. i >

.· 31° 32° 8

Figure 5. Predicted splitting of(! JO] diffraction peak in a" martensite (a) Ti-6wt%Mo, and (b) Ti-8wt%Mo.

characteristic of a'' martensite, has nothing to do .with reality; it can serve only as a measure of its nonunifonnity. The compositional dependencies of martensite lattice parameters which exhibit an increase of the a/b·.fj ratio with a solute content increase·are well documented [l; 3]. However, this can be explained using the.assumption that increase in solute content should facilitate the compositional modulation in the solid solution and therefore, amplify the stresses, which, in turn, move apart, in according to [4], the diffraction maxima. In this sense,

the orthorhombicity value (a/b/j ratio) depends not only on the_solute content in the as-quenched mart~nsite (mean solute content) but also on a further development of the decomposition upon ageing. This point has been shown quite conclusively in [ 11, 12]. Moreover, the data available suggest that such a development of thennally activated composit.ional modulations which has already begun during quenching, predictably rotates the habit plane of modulated structure, again being equivalent in this to the increase of mean solute content in martensite (Figure 6) [12]. On the other hand, tempera~e is able to affect the key point of the process - the coherent

I 8

'::I j

_.;..-- C octh

ah.c.p.

1,05

1,04

1,03

1,02

1,01

1,00 . 0,0 0,1 0,2

~ • quenched

2h aged

.. •

8 h aged

0,3 · 0,4 0,5 0,6 0,7 0,8 0,9 1,0

~ · t.CMo• arb. units ·· ·

Figure 6, Scheme which. shows that· ageing is, Figure 7. Relationship between the martensite "ortho­equivalent to the solute content increase in changing rhombicity" and the difference in Mo content in the en­"ortorhombisity" of martensite [ 12]. riched and depleted constituents ofTi-7wt%Mo alloy [ 11 ].

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TITANJUM'99: SCIENCE AND TECHNOLOGY

stresses responsible for the lattice distortions. It has been shown in [ 11] that direct correlation between alb .fj ratio value and compositional inhomogeneity of the solid solution, estimated with SAXS techniques, is observed only while coherency is maintained (Figure 7). Long times and/or high temperatures (but still within the single­phase a'' range) lead to the gradual relaxation of coherency stresses, resulting in a lowering alb .fj ratio value, while compositional modulations are continuing to increase in amplitude. Thus, the data available suggest that what is considered to be an unique characteristic of crystal lattice of a" martensite depending only on the solute content, in reality is subject to the influence of at least two more factors related to the amplitude of compositional modulation and the ability to keep the coherency inside the modulated structure. This implies that great care must be taken in understanding ·the precipitation behavior in martensite since conclusions on the type of precipitation and precipitate/matrix interaction are very often arrived at from the X-ray generated results [13, 14). It follows from Figure 5 that a discrete type of diffraction splitting should take place, in the sense that angle positions of additional· diffraction maxima should correspond to three orthorhombically or rhombically distorted phases of

a

13 = 24° 60

(>

6 50

4-0

3,:J.

a 20

1.0

061 20 20

a

13 =-180° 60 ~C) '

I q v

-~l so I

O', 'f

C' 'I' ·a 4-0

$ 30 -C\J

20

1 0 r\i

(;,1 62 63 64 ti5 061

20 Iq

LC>

lf " " "" a-

a . 13 = 201°

~-----~~-

so

60

U)

u', M ,.- . a U) [")

" "' ::::- 20

1 0 ,...,

61 62 63 64 65

28 Figure 8. X-ray diffraction intensity distribution for as-quenched Ti-7wt%Mo in 60<20<66 interval at some fixed~

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TITANIUM'99: SCIENCE AND TECHNOLOGY

· Iq 160

140

·120

' 1.00

, ·.ao .. 60

40

60 61 62 63 64 65 66

Figure 9:- Averaged by alpha and beta X-ray diffraction intensity distribution in (I) as-quenched and (2) aged Ti-7wt%Mo.

different chemical compos1t1ons. An attempt was done [ 15, 16] to study the diffraction picture of a" martensite. in more detail by means ·of full orientational techniques which allow a four­dimen~ional data set on diffraction intensity distribution to be obtained depending on three angles defining the sample orientation: alpha. and beta show the sample ·position regarding the sample holder while theta detennines its orientation to th~ goniometer axis. At

· constant theta pole figures can be generated from these data. At constant Iq

. the shape of reciprocal lattice points can be detennined, at fixed alpha or beta the spatial intensity distr.ibution is seen. Averaging the data by alpha and/or beta

leads to the common "theta-two theta" diffraction picture. The experimental details of the techniques used are given elsewhere [ 17]. Some examples of cross-sections at fixed beta values for Ti-7%Mo alloy in as-quenched and aged conditions are presented in Figure 8, while in Figure 9 the averaged by alpha and beta dependencies lq

(0) are shown. The last are looking as those corresponding to orthorhombic. symmetry, which are much better resolved for the aged condition. However, the local analysis of intensity distribution on Figure 8 shows a significant nonunifonnity, with maxima in a wide angle range instead of some fixed angles. This leads to the impression that as-quenched .martensite contains a continuous set of states having different alb .Jj ratio (including those with alb./:3 =!). On subsequent aging, diffraction maxima become more distinctive evidencing some transfonnations within this spectrum of states. Having in mind that a deviation of alb .Ji from unity is associate with a microcompositional induced stresses, it might be suggested that in polycrystalline material, various martensite plates differ in the level of microchemical,inpomogeneity, and this might be due to a different temperature of fonnation within the martensite transfonnation range. X-ray experiments has shown some non­unifonn distribution in diffract~on intensity also in alloys with solute content below the range where distinctive a" martensite fonns. Common theta-two theta analysis gives a typical h.c.p. diffraction picture in this case, for instance, for Ti-4\Vt°/o Mo alloys, however, at some orientations of the specimen states with alb.Ji :;t:l were observed. It can be concluded that a' ~ a" transition derives from a gradual increase in fraction of distorted microvolumes, this being detennined by thennodynamic (composition/temperature) and kinetic (time/temperature) factors. · ·

4. MECHANICAL PROPERTIES OF a"-MARTENSITE

The results of the tensile and HCF fatigue test of the Ti-7wt% Mo with 100% a" martensite structure in as­quenched and aged (450°C for 2 to !Oh: conditions where transfonnations in martensite are limited to the development of compositional modulations) are shown in Table 1. It can be seen that yield strength increased with ageing drastically by factor of about 2.5. The strengthening should be certainly related to the compositional modulations and resulting high coherency stresses. Strength increase leads to a corresponding increase in HCF (resistance to fatigue crack nucleation). A slight decrease in HCF with ageing time increase from 2 to 1 Oh indicates that the tendency for planar slip, i.e., slip band fonnation increases with the development of modulated structure at longer ageing, leading to easier crack nucleation. Tensile plasticity (elongation 8 and fracture strain EF) is reducing· after ageing to a highly detrimental level, a not un.expected result given the strengthening

T bl I M h . I a e ec amca prope rt' fT" 7wl°/cM 1es o 1- 0 0

cro.2 UTS O"f El,% E cr w1I cro.2 Condition MP a MP a MP a

EF GP a

0"107

As-quenched 534 780 1149 26.5 0.84 72 320 0.56 . 1067-1267 1156-1335 1156-1335 1.0-2.1 0.02-0.04 105-109 5?0 0.43-0.51 Aged 450°C, 2h

Aged 450°C, I Oh 1306 1353 1353 0.39 0.02 106 500 0.40 * A- significant scatter in prope'rties was observed in this conditions.

82

TITANIUM'99: SCIENCE AND TECHNOLOGY

. -

8968_ 20KV X370 100~M

mechanism. A very interesting result is the low modulus of elasticity for the as-quenched condition (72 GPa) increasing with ageing to I 06-107 GPa which is typical for the h.c.p. a-Ti.

Information regarding the microstructural reasons of the embrittlement was gained from fracture surface observations [ 15]. These revealed that the only visible difference between ductile (as-quenched) and brittle (aged) conditions is the increasing fraction of fracture surface exhibiting fracture (ductile on a macroscale!) along large martensite plate boundaries (Figure 10). It can be concluded that the large martensite plates are deforming first due to the strength difference with the neighboring smaller plates. Strain incompatibility leads to void formation at the boundaries of the large martensitic plates at low macroscopic strains inducing a low macroscopic ductility. This conclusion is supported by observations of the broken fatigue specimen surfaces which showed that fatigue cracks also nucleated at the large plate boundaries due to their preferential plastic deformation. This tendency was facilitated by increasing age hardening, as it can be seen from decreasing ratio of fatigue strength to yield strength (see Table 1 ).

Obviously, in order to tailor the balance of the tensile, fatigue, and fracture toughness properties of titanium alloys with a" martensite microstructure in favor of those related to the plasticity (but with a high strength level), the key reason of plasticity deterioration should be ameliorated: large martensite plates are to be avoided; plates are to be as uniform in size as possible. It has proved to be very difficult to achieve this using conventional methods of heat treatment. Primary martensite plates grow very fast, their sizes are often limited only by prior beta grain sizes while subsequent generations of plates gradually decrease in size due to the barrier function of previously formed plates. Besides grain boundaries, additional microstructural elements may be invoked to refine the martensite plates. One possibility is to introduce subgrain boundaries into the high-temperature beta phase, for instance, by beta deformation, capable of stopping primary plate growth. However, the most common and easiest to control method of refining the martensite is to form in the beta phase, by means of rapid heating, a microchemical inhomogeneity which would thermodynamically restrict the growth of martensite plates with sizes that correlate with the space distribution of alloying elements after rapid heating [ 18, 19]. With this approach, the beta grain size, which is also an important factor in tailoring the properties, and martensite morphology can be treated independently allowing a wider range of microstructures in titanium alloys [ 18, 19].

Coming back to the mechanical properties of as-quenched martensite, it is worth noting the fact that martensite in which the crystal lattice is already well distorted by coherent stresses has very low yield strength and elastic modulus. This implies a specific deformation mechanism, at least on an initial stage. A possible explanation can derive from a high instability of an as-quenched condition which was characterized earlier as a wide spectrum of states with various alb .J3 ratios. This condition changes, as it was mentioned, on ageing (see Figure 9). However, it might be suggested that applied stresses are also affecting this, at first in a reversible way (low modulus), then irreversibly (low yield strength). An alternative explanation could be based on a mechanism relating to the sliding of the boundaries of twins inside the martensite plates or that of plates themselves, as occurs in some alloys exhibiting shape memory or superelasticity properties. However, such a mechanism would imply a deformation at constant stresses while tensile experiments showed rather increasing stress type of deformation. And again, this needs a possibility of nonuniform deformation to be considered, with gradual involvement of martensite plates which differ in properties. Additional investigations are necessary for in-depth understanding of this very interesting phenomenon in a mechanical behavior of a'' martensite.

SUMMARY

1. Calculations of full interatomic interactions predict bulk modulation in the basal plane of h.c.p. solid solution, different for isomorphous and eutectoid binary systems and favorable for the orthorhombic distortion in the former case. Isomorphous systems are undergoing a spinodal mode of decomposition when the interactions are strain dominated. lsomorphous alloys are losing the stability at much higher temperature as compared to the eutectoid alloys.

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TITANIUM'99: SCIENCE AND TECHNOLOGY

2. X-ray and SAXS experimental results available, except for rather continuous than discrete character of X-ray diffraction peak splitting, are well consistent with- the"explanation that the formation of the orthorhombic structure is being due to the elastic distortions imposed on h.c.p. solid solution as it decomposes into two constituent having different solute contents and coherently _bonded in a two dimensionally modulated structure.

3. The as-quenched yield stress and elastic modulus of orthorhombic martensite in beta isomorphous Ti-7wt%Mo alloy were very low but increased very sharply on ageing at 450°C. No phase transformations are involved at ageing employed, so strengthening can be attributed only to the further development of the modulated structure of the as-quenched condition. Possible way to improve the strength/ductility balance of martensitic structures is to make martensite more uniform and finer.

ACNOWLEDGEMENT

This work was supported by the INT AS foundation (grant 93:630) . .

LITERATURE

[I]. BAGARYATSKY YU.A. TAGUNOVA T.V. NOSOVA G.I. Metastable Phases in Titanium Alloys with Transitive Elements. - Problemi Metalovedeniya I Fiz.i Met., Metallurgizdat. Moscow. 1958. P. 210-234. -

[2]. NIKOLIN 8.1. Multilayered structures and polytypism in metal alloys. Naukova Dumka publisher. Kiev. 1984.256p

[3]. BAGARYATSKY YU.A. NOSOVA G.I. TAGUNOVA T.V. General Regularities of Metastable Phases Formation it Titanium Alloys. II Doklady AN SSSR, 122. 1958, P. 593 - 599.

[4). IVASISHIN O.M. KOSENKO N.S. On the Nature of Orthorhombic Symmetry of Martensite in Titanium · Alloys. Proc. ofSymp. Titanium '92, Science and technology. - San Diego (USA). - 1992. - P. 721-727.

[5]. DA VIS R. FLOWER H.M. WEST D.R.F. · Martensitic Transformations in Ti-Mo alloys J. Mater: Sci. -1979. ·- 14, No. 4. - P. 712 - 722.

[6]. IVASISHIN O.M. KOSENKO M.S. SHEVCHENKO S.V. TATARENKO V.A. AND TSINMAN C.L. Strain-Induced and "Electrochemical" Interact_ions of P-Stabilizing Substitutional atoms in H.C.P.-Titanium Alloys. Met. Phys. Adv. Tech. 1998, 17. P. 13-29.

[7). KHACHATURIAN A.G. Theory of Structural Transformations in Solids. - New York: John Wiley & Sons. Inc.1983. 574p.

[8]. _IVASISHIN O.M. KOSENKO M.S. TATARENKO V.A. SHEVCHENKO S.V. TSYNMAN C.L. Strain­. Induced and "Electrochemical" Interactions of P-Stabilizing Substitutional atoms in H.C.P.-Titanium Alloys. Metalofizika I Novejshie Thec;hnologii, 1995, 17, No. 9, P. 61 - 68.

[9). JORO'vKOV M.V. Redact. mag. l~estiya. Vuziv. Fizika. Tomsk. 1991. Deponent. at VINITI 1991, 1023-. 891. 51- P.

[I OJ IV ASISHIN O.M., TIMOSHEVSKII A.N .. The Influence of the Alloying by p-stabilizing Elements on the Titanium Electronic Structure. Proc. of Symp. Titanium '95, Science and technology. - Birmingham (UK). -1995. - P. 2470-2477.

[11) IVASISHIN O.M., USTINOV A.I., KOSENKO N.M. Effect of Concentrational Inhomogeneties on a"­Martensite Parameters in Ti-8.0wt%Mo _Alloy. Proc. of Symp. Titanium '95, Science and technology. -Birmingham (UK). - 1995. - P. 2236-2241..

[12) IVASISHINO. M., USTINOV A. I., SKORODZIEVSKIIV. S., KOSENKOM. S.,MATVIYCHUK YU. V, AZAMATOVA F. I. Structural and Compositional Changes on Isothermal Aannealing of a"-martensite in Ti-8 wt.% Mo alloy. Scripta Materiala.- 1997, - 37, - P. 883-888.

[13) MOISEEV V.N., GERAS'KOVA V.A. New Study ofTitanium Alloys, Nauka, Moscow, 1965. [14) KOLATCHEV B.A., MAMONOVA F.S., LJASOTSKAJA B.S. Martensitic Decomposition of Alloys of

. Ti-Mo system under Annealing. lzvestija AN SSSR, Metally, 1974, I. P. 200-205. [15) The Mechanisms of Marterisite Formation and Tempering in Titanium Alloys arid their Relationship to

Mechanical Property Development. II INTAS 93-0630. - 1997. - Final Report. [16] IV ASISHIN O.M. KARASEVSKA YA O.P. Fine structure of the X-ray reflections for Ti~5-7Wl°/oMo.

Metalofizika I Novejshie Thechnologii, 1999, 21, No. 4, (to be published). · [17] KARASEVSKA YA O.P. Orientational X-ray experimental method for monocrystal's phase analysis.

Metalofizika I Novejshie Thechnologii, 1999, 21, No. 4, (to be published). [18] GRIDNEV V.N. IVASISHIN O:M. OSKADEROV LP. Physical Backgrounds of the Rapid Heat Treatment

for Titanium Alloys Strengthening. Naukova Dumka pubiisher. Kiev. 1986. 256p. [19] IVASISHIN O.M. TELIOVICH R.V. Potential of rapid heat treatment of titanium alloys and steels. Mat.

Sci. and Eng., 1999, A263, P. 142-154.

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TITANIUM . 99: SCIENCE AND TECHNOLOGY

STRUCTURAL AND PHASE TRANSFORMATIONS IN COLD STRAINED

ULTRA - FINE GRAINED TITANIUM ALLOYS

Drozdova N. A., Pyshmintsevl. Yu., Popov A.A.

Urals State Technical University, Mira 19, Ekaterinburg, 620002, Russia

E-mail:torri@dialup.ustu.ru

· INTRODUCTION

Recent investigations have demonstrated the possibility of a formation of nano and

submicrometer structure in metals and alloys by severe straining. Ultra-fine grained

materials are known to have some advantageous physical and mechanicaLproperties.

Also some features of structure and phase· transformations in the materials are

revealed. The studies of the features and its influence on the mechanical properties are

very important for the developing effective treatment methods [ 1-3 ] . The

extraordinary properties such as increased strength, ductility, low temperature and high

speed superplasticity have commercial significance. Submicrocrystallined titanium

alloys are advance-materials due to first of all high strength and corrosion resistance:

The aim of the presented research is studying of structure and phase transformations in

ultra-fine grained titanium alloys· and developing on this base treatment techniques to

achieve high· strength thermal stable states.

· 1. Methods and materials of investigations

A single a. phase commercial alloy VTl-0 (99.53 mass.%) and · multyphase

commercial alloy VT6 ( 6Al-4V ) were used. The submicrocrystalline structure was

formed by thermal-deformation route including severe deformation. The samples were

subjected to shear straining using Bridgeman anvil technique and the 20 mm die for

equal channel angular (ECA) pressing- [4]. It is possible deform.enhanced materials in

form of thin films· up to highest stain without fracture by first method. As result

limited number of-mechanical-properties may be measured using these small samples.

However numericaLinvestigations [l,5-8] have demonstrated steady results of nano­

scale structure formation in low-ductile metals and alloys by this method· as a

consequence this method was selected for both materials.

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TITANIUM. 99: SCIENCE AND TECHNOLOGY

Second method is not suitable for·materials with limited plasticity resource. But it

is possible to process bulk samples suitable for careful mechanical testing. Thus ECA

pressing was used only for alloy VTl-0. ·.

,;·. -

2. The results and discussion

High dislocated packets of twinned plates· 80-100 nm thickness could be seen

arranged in the direction of deformation in the alloy VT 1-0 relatively low strained by

Bridgeman·metliod.· ·Plates are divided on,fragments by dislocation,.walls.-:.Near,the

coherent boundaries of the plates the density of dislocation is especially high and .they

set up long-range fields of. elastic -stresses:· These stresses are responsible. for -.the

specific-"diffuse". difraction contrast in the image [3, 7]. The electron: difraction patterns

showed individual diffuse reflections whose possitions corresponded to typical t_win

misorientations.

_ · As a strain increases twin boundaries become thinner and bender. At the· same ti111e

in heavily strained specimens a specific structure was formed with an increased number

of individual structure elements ( fragments ) (- fig .1 a) There: were no' deformation

twins,_ which. is evidence of - the formation of the :·new structure · type. , The .equiaxed ·

structural units formed by dislocation and disclinatiori mechanism ·have average size of

80nm. Character feature of structure is a presence approximately 20 % of extra fine

structure with mean grain size of 20-30 nm: The ·difraction patterns from this region

showed numerous extra spots and diffusion background, which implies that material is

partly amorphous [ 9 ] . . -.

From_ the analysis of foregoing, it may be concluded that at lower strain the

dominated mechanism in a "." Ti is twinni_ng and at the large one the disclination . effects

showed up. . ,-:_ -· '·: .

The analysis of structure· changes_ in both states during subsequegt heating showed

two stages-of transformations. On the first one-at temperature below 250 °C the defect

density inside of.coar~e fragments decrease, microtwins becomes thinner. and sharper,

low angle boundaries ·are ,more: perfect. , Small fragments :disappear. -- Electron

diffraction pattern took a point character with a presence of the:texture notices. When

specimens were strained up to highest deformation leveL the structural · elements

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TITANIUM ' 99: SCIENCE AND TECHNOLOGY

a) b)

___ .,4

400nm ... . c) d)

e) f)

Fig, 1 The TEM microstructure image of alloys VTl-0 (a-d) and VT6 (e,f); a- c,e,f - strained by

torsion; d- ECA pressed; a, d- as prepared state; b- annealed at 250 °C ; c- annealed at 3 50 °C; e­

annealed at 200 °C; f- annealed at 450 °C.

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TITANIUM' 99: SCIENCE AND TECHNOLOGY

increased up to .80 nm, finely dispersed and amorphous components disappeared. The ··

continuous electron diffraction patterns broke up into sequence of closely spaced point

reflections and diffusion halo disappeared. This implies that real nanostructured

titanium not posses thermal stability, and the grains are able to growth at more lower

temperatures than in coarse grained alloy ( fig 1 b ). Low temperature annealing makes

grain size distribution more uniform and boundaries becomes sharper with noticeable

high angle contrast. This change may be caused by decreasing in long-range elastic

stress fields on the boundaries.

Second stage of structural transformations in severe deformed material is moving

and redistribution of high angle_boundaries. In situ type of recrystallization is revealed

in severe strained samples annealed at 300 °C. In such a case, the migration of low­

angle boundaries and coalescence into large conglomerates usually takes place. The ·

nucleation on new grains by way of primary . recrystallization is remarkable at

temperatures upper 3-50 °C ( fig. 1 c) . Under this circumstances many new grains

acquired a common boundary, but the initial defect matrix was still here. But at these

temperatures grain size not exceed 150-200 nm that implies that recrystallization did

not cause noticeable grain growth under these conditions. Significant grain growth

caused by increasing of temperature up to 450 °C was observed. Grain.boundaries with

low density of defects formed perfect triple junctions. The results of bending tests and

microhardness measurements of variously treated VTI-0 alloy are given in fig. 2.

The data analysis has demonstrated the importance of the annealing for mechanical

behavior. The rise of the strength and plasticity as a result of keeping at 250 °C is most

significant result. These changes are caused by transformation of deformation

boundaries in more perfect boundaries of annealing with consequent decreasing . of

long-range elastic stresses. Besides very important is the different in microhardness and

yield stress dependencies on an annealing temperature. Thus only yield stress is more

sensitive for low temperature structure transformations.

Studying of alloy VTl-0 structure subjected to ECA pressing at 400.:.425 °C

showed that the polygonization of the ultra-fine grained material ( mean grain size

2-5 mkm) is main process for this technique. As result subgrain structure sized of 0.25-

0. 5 mkm was observed in the samples after deformation (fig. Id). The electron

diffraction pattern analysis has demonstrated low-angle misorientation of neighboring

fragments. Grain structure is not uniform_ As rule largest grains ( 5 mkm) are

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TITANIUM. 99 SCIENCE AND TECHNOLOGY

surrounded by finest one. Besides some nucleation of new grains inside ··Of grains is

observed.

3500 800 .

3000 700

600 2500

ca 500 ~ 2000 E In 400 C CII ·• c e 1500 ... 300

"' 1000 3 200

500 4

100

0 0 0 100 200 300 400 500 600 T,°C

Fig.2 The influence of 2-h annealing on the structure and properties of the VT 1-0 alloy

to the strain of 4: 1 - mean fragment (grain) size; 2 - microhardness; 3 - U.S.;

4-Y.S.

Sharp crystallographic texture was established by usmg X-ray measurements.

Analysis of direct polar figure have demonstrated that mainly prismatic plane ( 100) is

onentated in perpendicular direction to axi's·; of channel . and. basic plane ( 001 ) has

orientated in one plane. Mechanical tests of the samples in· different directions showed

significant mechanical anisotropy ( tablel.). But it is notable that obtained structure

posses higher properties than for coarse grain cold deformed alloy.

Researches of thermal stability of this · structure have· demonstrated beneficial

influence of recovering at temperatures lower 300 °C on ·strength-elongation complex

( table .. 1. ). The ·weak decreasing of tensile stress and · rise of elongation similar to

strain hardening rate at constant strength value were observed under this conditions.·

The observable growth of substructure with the further increase of annealing

temperature has resulted in· gradual decreasing of strength and growth of plasticity.

The inten_sive change is observed attemperature 525-550 °C which causes growth of

the size of a grain owing to collective· recrystallization.

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TITANIUM. 99 SCIENCE AND TECHNOLOGY

" -... :

Mechanical properties of treated VT 1-0 alloy in axial direction Table 1 . -

Straining Annealing Y.S., T.S., Uniform · · ·, Total Remarks

Mode Teinp., °C MP a MP a elong. % elong.%

Equal 1,~ 550 640 4.5 · 15 r· .,,

Channel - 700 780 2.5 · 14 *

Angular 810 880 2.0 -

14 ** -pressing , 450 520 650 6:1 · .. 18

0

450 · 650 · 725 4.2 10 * •.

450 ·700 -· 780 3.5 16 **

ECA - 930 1025 2.5, 17

. pressing+ 200 910 . -

1630 4.5 ·. 19

::J : ·,"

. -Cold 300 850 1040 5.5 18

rolling 450 790 900 13.2 20

50% .. ,· ·,, .. ;

*-perpendicular direction along [100]; **-alo?g,(001].

From the p.oint of view pf practical use of an alloy in ECA .. p~essed state .i~- the . ' . . . ' ~ ~ ~ . ' - ·. . . •, . '

~mportant .· d~velopment of deformatio!l metl)o_qs, keepi11g the achiev~p properties_.

Taking into account, that with h9t_ deformation it is impossible, the basjc attention was _ • t. ,, I > • • • ." • • • • > • ' ' - '' •' •' '• '

given of cold plastic deformation. The .. deformation was carried. out. by ~. way of . . - . . ' ~- .. : - . . . : .

clr_~wing and subseq_u.ent roHi_ng . t9 11 . mm diameter_ rods. TEM ·. d_e~onstrated t_he

elongation of subgrains in a longitudinal direction, inside them it is possible. to observe . -· . - . _. . ... ·. )\. . -.·

f~n!)ation of d\sl9cation. waU~,and separat«? disl99atiorys. X-ray analysis has sho}v.n, that

texture chang~~ by·rolling process. NameJy some misorientation of a basic plane t~ok

place. The_ me~h_anical properties_ of all(?y after colg de~ormati~n ~nd ~ubsequent

annealing are given in table 1 .. The a11alysis _ of results shows, that realization_ 9f • •' ! • • • • ' , • • ; • • ~ \ ',. L. I •, '

deformation and annealing at 300-350 .°C allpws to reach an optimum of strength and ~ . . 's . -· . . '- . . ' •

plasticity corl)plex. Thus, the grain stru~t~rn _ refinement techniques of a by intensiye ' ' • • ," • • '·.' ,:_ • . '",• •. ' -' T • __. ' j,. ' •. • • ,, -

cold. or warm: peform!l,tion, all.owing to achiev~ i.n _ single:-:phase: titanium alloys . high . , -~- - • . • ~ ! . . ' ' • .§ • - - ~ • • - •

strength state with a sufficient level of plasticity ~re developed. .;.i.·

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TITANIUM . 99 SCIENCE AND TECHNOLOGY

However, a greatest interest in studying of influence of intensive deformation takes

place in-field of heterophase alloys with metastable a.'7 and p phases. It is ·connected

that these phases are mechanically unstable under certain.alloying level and the . . .

p!astic deformation initiates both structural and phase transformations [ 10 ]. Most

widespread: circuit of transformation under quenching are : p ._ a." -a.'

transformations. The plastic deformation allows to add in this circuit· variants of

transformation P - a." through· intermediate phases,. with· deformed light thetrohonal

BCC lattice. Most . perspective from this. point of view ire the ·alloys with a. beta

stabilization factor of 0.8-1.3. However, in other alloys, varying modes of thermal

treatment, it is possible to receive · initial structure containing before intensive

deformation:metastable.a." - phase:

' ·. The· ( a.. + a.."} structure was formed by a ·water ·quenching from temperature 900 °C­

in an alloy ,BT6. The X..:ray. and the transmission electron-microscopic. researches

revealed a-transformation of. a". phases in a.' and P during intensive deformation. The

grains a of a phase do not exceed 100 nm, and b the particles, are .located on

boundaries as thin layers. The grain structure· ·of a. - -phase is typical for materials

subjected to' intensive -plastic· deformation. A high level of stresses and increased

dislocation density create rather complex · contrast The aging ·at· temperatures 200-

4000C resulted to''reorganization of boundaries, on which a typical grain boundary

contrast was forming. Practically the grain size did not change (fig. I e:. t). ·The· stability

of a. grain in many respects is provided by a presence of ultra-fine boundary phase

undergoing -phase transformations .and a.-solutio~ ailoying. ·Increase of temperature up

to' 500 °C causes 'natural integration and coagulation bf transformation products and it

reduces ,efficiency of resistance to a movement of borders. As a result, the grains

essentially grow up to I micron and more at long-time aging.

The mechanical properties of the processed alloy are in close . correlation with

transformations in structure (table 2). Combinations of hardness, plasticity and strength

at bend test achieved after aging at 450 °C, 0.5 h., essentially exceed a typical

properties of the alloy in heat treated state. . ' . .

Thus, it is possible to achieve extra high-strength states having significant. thermal

stability in heterophase titanium ·alloys, by intensive plastic deformation using ,·. .

peculiarities of phase trans~or~ati.ons.

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TITANIUM. 99: SCIENCE AND TECHNOLOGY

VT6 alloy microhardness and bend test results Table 2

Treatment Micro hardness, Y.S., Mpa Max. Stress, · Relative MP a MP a

Plasticity* Quenching 900 °C 3100 - - -Quenching 900 °C + severe 4800 1800 2100 1 .

straining + annealing 200 °C, 30 min 4800 1600 2100 1.4 + annealing 450 °C,30 min 4800 1550 1850 2.25 + annealing 500 °C,4 h 3200 1000 1250 2.8

...

*Plasticity is measured as a displacement until maximum loading during bending ·

CONCLUSION

Thus, the structure and phase transformation features in titanium alloys subjected

to severe plastic deformation are studied .. The materials advance mechanical properties '

have been demonstrated. : A grain .size reduction. to nano or submicrometer,scale by

deformation offers several advantages in industrial application of the alloys as a high

strength and ductile material.

ACKNOWLEDGMENTS

The authors are grateful to all colleagues who took . part in discussions. Authors

especially would like to thank· prof R.Z. Valiev and his co-workers for kindly

assistance in --s~vere deformation procedure. This work was supported. by grant of

INTAS No97-1243

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