DEFORMATION MECHANISMS IN TITANIUM. EFFECT OF THE IMPURITY RATE

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1 1805 DEFORMATION MECHANISMS IN TITANIUM. EFFECT OF THE IMPURITY RATE M.J. Philippe, 1 aboratoi re de meta 11 urgie, CNAM, Pads, France. C. Esl ing, 1 aboratoi re de meta 11 urgie structura 1 e, Uni vers ite de Metz, France. B. Hocheid, laboratoire de metallurgie, CNAM, Paris, France. introduction Like all other hexagonal materials, titanium has several glide and twin systems likely to operate. The corresponding critical resolved shear stresses (CRSS) vary much with the rates of interstitials. It is usually very difficult to determine these values. This determination requires the use of monocrystals with precise impurity rates, stressed in well defined directions. Moreov~r.these values for the monocrystal cannot be immediately extrapolated to the behaviour of the polycrystals. Thus we prepared polycrystalline samples by a process similar to the industrial one and studied the texture and the behaviour in tension and in compression as well as the deformation mechanisms of these various titanium alloys. Materiel s used Van Arkel titanium was melted by an electric arc and doped in nitrogen, oxygen and iron with respectively 0 to 300 ppm, 2500 ppm and 1000 ppm in weight. These ingots were then cold-rolled down to thickness of 1 mm to obtain a texture similar to that of the industrially produced sheets. The sheets deformed by 50 % were then submitted to vacuum annealing for an hour at 600 C. The non doped material show a grain size of about µ. The interstitial elements (O,N) do not really modify the grain size; the latter decrease weakly as the rate of interstitials increases. On the contrary, the addition of iron produces a rather large decrease of the grain size (Table I). Denomination rates in ppm of weight grain size after 02 Fe N2 annea 1 ing T '\, 50 µ T '\, 50 µ TSO, 3F~ '\, 50 µ - T90 2 4Fe '\, 40 µ Titanium _T Fe '\, 40 µ T Fe '\, 40 µ T90 2 7Fe '\, 30 µ T OFe '\, 20 µ T9024Fe1N '\, 40 µ T90 2 4Fe3N '\, 40 µ Table I The grain size as a function of the impurity rate.

2 1806 Textures The various samples show similar textures, the major component being close to {2115}<10lO>(fig.1). This preferred orientation is so strong that itcan be easily apprehended by a simple pole figures analysis without resorting to full three - dimensional ODF calculation (1-2). RO a) b) Fig. 1 - Texture of the sheets Pole figure {0002} a)-and {10l1} b) However, an increase of the oxygen rate slightly decreases,the tilt of the c axes with respect to the normal to the sheet plane (from 35 to 25 in the cases of our samples). These results are similar to those given by Larson et al (3). As we will see later a high rate of oxygen inhibits the twinning and the texture evolution observed is in agreement with the results of Thornburg (4) showing that the stable orientation of the c axes is at 35 from DN towards the transverse direction when twinning operates and becomes central when it does not. This was confirmed by our results with zirconium (5). Analysis of the plastic range of the tensile curves All the t~nsile curves of titanium show a kink at the yield point. The length of the plateau following this kink increases with the rate of interstitial impurities. Oxygen also increases appreciably the values of the yield stress (Re) and of the fracture stress (Rm). At the same time the maximal uniform elongation (eu) and the total elongation (et) decrease (tab le I I). The best modelling of the stress-strain curve (cr, ) obtained from the experimental points is achieved with a Jaoul-Crussard type representation (6). It is shown in this study that the strain hardening exponent n varies during the deformation. By plotting lnlicr versus ln, it is possible to find out, if n has one or several values, in relation to the slope of the curve (7). All the titanium alloys show at least two or even three stages when uniformely elongated from 0 to 12%. The first stage always corresponds to a relatively high n exponent. This holds whatever the alloy and the orientation of the tensile stress. Sometimes it increases up to 5% elongation \\fiereas it decreases down to 1 or 2% with high purity materials.. On the contrary, the exponent n associated to the second stage is very small or even practically zero. Finally the third stage appears only after 8% of elongation when the alloys are highly doped in oxygen. The hardening coefficient is negative, which indicates that diffuse necking has been reached. (tab 1 e I I I).

3 1807 Denomination et % eu % Re Rm RD TD RD TD RD TD RD TD T TB T502 3Fe T9024Fe 48, T18024Fe 37,1 35, T2502 4Fe , ,4 43, T902?Fe ,2 35,9 130 T9021 OFe ,5 32, T9024Fe 1N2 37, T902 4Fe 3N2 39,5 41, , ~~----~--~---~--~--~- Hv Table II Results of the tensile test in function of the impurity rate. T802 T90 2 4Fe T Fe Denomination T Fe - n n n n Stage I Stage I I Stage III RO TD RD TO RD TD RD TD T90 2 4Fe3N 2 RD TD Table III :. Values of the strain hardening exponent n in the 0-12 % range of elongation. Forming limit diagrams The curves have been drawn from Irsid type simulation tests, only in the limit cases of uniaxial tension and symmetrical biaxial expansion. In the case of uniaxial tension, the best characteristics are obtained with pure or iron-doped materials. The effect of oxygen is well marked in the field of biaxial expansion. An increase of the oxygen rate corresponds to a marked decrease of the major deformations e: 1 e: 2 at necking and fracture. An

4 1808 enrichment of titanium with iron does not modify the deformations much at necking and at fracture. (Fig. 2) ~2 " 1 2. Fig. 2 : Forming limit diagrams at necking Discussion Deformation mechanisms in uniaxial tension As the texture of these materials is very sharp, we consider the ideal orientation corresponding to the major component. In a uniaxial tension test (in RD or in TD) the orientation of the lattice favours prismatic glide. The Schmid's factor for two of the three prismatic systems amounts to 0.43 for a tension in R.D and to 0.29 in T.D. Micrographies in optical or electronic microscopy show that the glides often share a common slip direction at the beginning of deformation (fig. 3). Fig.3 Optical micrography at the early deformation stage. Fig.4 Cross slip at 7 % of deforrnation.

1809 This can be understood by the narrow orientation of most crystallites. Thus, the glides easily cross the grain boundaries at the very beginning of deformation.

5 1809 This can be understood by the narrow orientation of most crystallites. Thus, the glides easily cross the grain boundaries at the very beginning of deformation. When the deformation rate increases, other slip lines appear within grain~ corresponding to slip on another prismatic plane (fig. 4). Twinning of {1012} or {1122} type appear at all the deformation stages and in all the Ti-alloys, except for those with high rates of interstitial elements ('V2000 ppm). It also appears less frequentely in the T80 2 alloy for whick no iron has been added. Twinning allows for stronger deformations at necking and stronger full deformations (in tension) on the forming limit diagrams. The mechanisms in biaxial expansion In equibiaxial expansion, the initial orientation of the lattice is not favorable to prismatic glide (Schmid's factor amounting to 0.14). However titanium twins during deformation at room temperature ({10l2} and {1122} twin types). Twinning is unusual or even non-existent in the case of samples strongly charged with oxygen, which accounts for the poor formability of these alloys in expansion (5-8). In addition to its contribution to the whole deformation, the reorientation of the twinned lattice allows for the prismatic slip to become active (fig. 5) (corresponding Schmid's factor amounts to 0.3 in the {1012} twin). Fig.5 : Prismatic slip within a {1012} twin. The intervention of twinning and its consequences The twinning process may be divided into 2 stages: first, nucleation and then propagation of the partial dislocations (emissary dislocations (9)) generating the twinned volume. Nucleation is easier with a low stacking fault energy associated to high values of the local stresses. This applies to hexagonal metals with the c/a ratio far from the ideal compactness and with, moreover, some rate of ~~~~!i!~!i2~~! elements deforming the lattice. The effect of an increasing grain size is to increase the true local stress.

1810 Finally, the twin propagation is only possible if the metal or alloy is not excessively charged with interstitial elements such as 0 2, N 2 which hinders the propagation of

6 1810 Finally, the twin propagation is only possible if the metal or alloy is not excessively charged with interstitial elements such as 0 2, N 2 which hinders the propagation of the-part1ai-dislocations and therefore the development of the twin. The above hypotheses are supported by the following experimental observations. There is no twinning in ultra-pure metals. On the contrary, twinning develops well in titanium doped with iron but with little oxygen and nitrogen and when the grain size is relatively high. Finally, twinning develops uneasily or not at all when titanium is highly charged with 02 In the latter case the micrographs show the existence of micro-twins( < 0.1 µ) in the neighbourhood of grain boundaries and these micro-twins have not extended. ~~ jjr 5~~". References (1) (2) (3) (4) (5) (6) (7) (8) (9) Fig.6 Micro-twins in the neighbourhood of a grain boundary. H.J. Bunge "Texture analysis in Materials Science" Butterworth,London (1982). H.J. Bunge, C. Esling (Ed) "Quantitative Texture analysis" S.F.M, Paris and D.G.M, Metallurgy Information New-York, OberUrsel (1982). F.R. Larson, A. Zarkades, P.H. Avery "Titanium Sciences and technology" (1973). D.R. Thornburg, Thesis Carnegie-Mellon University, Pittsburgh (1971). M.J. Philippe, Thesis Universite de Metz, France (1983). C. Crussard, B. Jaoul, Revue Metallurgie 47 (1950) M.J. Philippe, C. Baudet, B. Hocheid, Mem. Sci. Rev. Met. 10 (1982). M.J. i;uillaume, C. Beauvais, B. Hocheid, Icotom 6, Tokyo, Japan (1981). A.W. Sleeswyk, Acta. Met., 10 (1962) 705.

Microstructural and Textural Evolution by Continuous Cyclic Bending and Annealing in a High Purity Titanium

Materials Transactions, Vol. 45, No. 9 (24) pp. 2826 to 2831 #24 The Japan Institute of Metals Microstructural and Textural Evolution by Continuous Cyclic Bending and Annealing in a High Purity Titanium