STRUCTURAL CHANGES IN THE BT6 ALLOY WELDED JOINTS AFTER HEAT TREATMENT

V.N. Gridnev, V.N. Zamkov, N.P. Kushnareva, V.I. Trefilov, E.V. Turtsevich and A.D. Shevelev

Acadmy of Sciences of the Ukr. SSR, USSR

BT6 alloy (Ti-6Al-4Vl has found a wide application in the fabrication of titanium welded structures. However, welded joints of this alloy do not meet service requirements in a number of cases [l, 2]. Subsequent heat treatment is necessary to improve mechanical properties of joints. With the increase of thickness of welded elements the problem becomes more acute. In this connection the effect of thermoplastic cycle of electron beam welding and subsequent annealing on mechanical properties and structure of welded joints with the aim of selecting the most favourable heat treatment conditions for welded structures is considered in this paper.

Materials and Experimental Procedure

Investigations have been carried out on 55 mm thick alloy plates, cut out of a forging. Welding was carried out at the following parameters: 60 kV accelerating voltage, 430 mA welding current, 65m/h welding velocity. Weld width along the sample height varied in significantly and wa.s 3-3.5 mm. Annea~ing of wel~ing samples wa.s carr~ed out in a vacuum fu~nace at l .io-5-2 .10~ mm Hg residual pressure. Cooling rate was 2-3 deg/min.

Two annealing temperatures have been chosen: annealing in two phase (a+ Sl-region [1] at 750°C for 2 and 48 hours, generally accepted for this types of alloys as well as annealing in S - region at 1020°C for 2 hours.

The results of the paper [3] show, that annealing in the S - region provided the increase of Kie values of Ti-6Al-4V.alloy semifinished products. The evaluation of mechanical properties of the weld metal and welded joints was performed on round standard samples. Fracture toughness (Kiel was esti­mated on 50 mm thick specimens for excentric tension with a notch along the weld axis. ·

Structural investigations were carried out using transmission electron microscope of Y3MB-lOOB type at lOOkV accelerating voltage. Final polishing was carried out in ethanol 'solution of hydrogen chloride and sulphuric acid at the following .concentration of components: H2S04 - 4g-equiv/l, HCl -2.4g-equiv/l. Internal friction (decrement oil and oscillation period were measured on low frequency inverted pendulum by electronic method of recording in the temperature interval 20-900°c in 2.lo-5 mm Hg vacuum at frequency f 2-5 Hz. Decomposition of complex spectrum was performed by means of elementary peaks plotting, corresponding to one time' of relaxation according to the parameter values: temperature value (Tmaxl, height (oimaxl and activation energy (Hl according to relationship

01 = oimax sech [~ (! - ! l] (ll R T Tmax

578 V.N. Gridnev et al.

using tabulated values, given in [4] for a better agreement with the experi­mental curve. Test samples were produced by the centreless polishing method up to 1.2 mm dia with subsequent chemical thinning to 0.8 - l mm. Measure­ments were carried out on 3-5 samples for each state.

Mechanical Properties

Strength of welded joints and weld metal in all the states under investi­gation is equal to that of the base metal (Table 1). At the investigated heat treatment conditions welds and base metal possess the highest plasticity (c and wl after annealing at 750°C. Annealing at 1020°C causes a considerable decrease of weld metal plasticity. At the same time, the decrease of the actual ultimate strength (cbul takes place both in weld and base metal.

Fracture toughness of the base metal in all the states under investiga­tion is about the same. Welds have a lower Kie value. The difference is -35% after 2 hours annealing at 750°C, and -15% after annealing 750°C - 48 h and 1020°C - 2 h.

No noticeable changes of mechanical properties have been observed in the heat-affected zone in comparison with the base metal. As to the above heat treatment conditions the best combination of mechanical characteristics of welded joints, made by electron beam, is ensured by annealing at 750°C for 48 hours.

Electron Microscopic Investigations

Weld metal due to a high cooling rate has an a'-structure (Fig. 1), similar to the one obtained by tempering of the alloy from S-region. Along with considerable density of dislocations (Fig. la), rather uniformly distri·­buted through the crystals, the presence of a large number of thin strips attracts one's attention (Fig. lb). It is known that 12 different orien­tations of a(a') - phase can be formed in one S - grain. This, together with the inner structure· and phase dispersi ty mentioned above, complicates the interpretation of microdiffraction patterns of weld metal. Nevertheless, microdiffraction (Fig. le) shows that thin strips,. seen in Fig. lb in a' -phase are twins. Indeed, when the twinning plane is parallel to electron beam, the twinning axis will be parallel to the reciprocal lattice plane, corresponding to electron-diffraction photograph [5], and the diffraction pattern of the twin is obtained by rotation by 180 deg. round the twinning axis. The electron-diffraction photograph (Fig. le) obviously consists of 2 diffraction patterns. One diffraction pattern is obtained out of the other by rotation by 180 deg. round the axis (lOl), which is in the reciprocal lattice plane with the zone axis (101).

Fig. ld gives a theoretical electron-diffraction photograph, estimated in paper [6]. In this case (101) plane is the twinning plane, which is typical for twinning of a- and a'-crystals. At the same time a presence of a large number of dislocations and of elastic stress fields, connected with them, creates rather high internal stresses in the weld volume.

A structure typical for the base metal and consisting of a-phase crystals, the cross-section size of which reaches several microns and of narrow S-phase crystals between them is shown in Fig. 2a. Unlike the weld metal, the dis­locations observed here are located, as a rule, in the area adjacent to the

BT6 ALLOY WELDED JOINTS 579

a/6-phase interface and are practically oriented normal to the interface, This must create an increased level of microstresses in the interfaee r.egion. On a/6-boundary one·can see a special contrast, a detailed study of which has been carried out over the last· years, with respect to peculiarities of the (a + 6)-titanium alloys destruction. At first, there was such an opinion that this is an area with a high dislocation density or a hydride layer [7].

The final structure of the layer was found by means of electron micro­scopy analysis in the papers [7-9.l. It is shown that the intermediate a(6-layer is observed both in the form of a monolithic interphase layer with a fee structure, and of a striated interlayer, they often being simultaneously present on the a and 6 -phase boundaries. The interlayers do not grow in the process of i,sothermal annealing, but only do so at rather slow cooling. As for the structure of the striated layer, it was at first identified as a hexagonal a-phase with orientation different from the main a-phase [7], and, finally, as that consisting of twinning and not-twinned fee ares [8]. Besides, Margolin H. and his collaborators [9J have shoWn that the interphase zone is not only a layer with a special crystallography, but also a region with a high dislocation density. It is still to be clarified what role does the interphase layer play in the distruction of (a + 6)-alloys. However, it has been shown [10], that treatment of Ti-6Al-4V alloy which led to for­mation of structures completely equivalent from.the point of view of volu­metric content, sizes and shape as well as to different types and sizes of interlayers ensured the yield point change from 68.3 to 93 kgf/mm2 and ulti­mate strength change from 96.5 to 104 kgf/mm2, o - from 12 to 16.2%. In this connection it appears interesting to follow the propagation of a crack in the base metal, the crack being presented in thin foil (Fig. 2b).

A twin formed in the base metal while approaching the a/ 6 -boundary, creates a stress field here, As can be seen, the latter can not relax by means of plastic deformation, and leads to crack formation in 6-phase. The crack propagates along 6-phase and a-twins. Heat treatment of the alloy at 750°C for 2 hours does not change the base metal structure. Annealing of welded joint at 750°C for 48 h and at 1020°C for 2 h leads to the formation of fine-dispersed precipitations in the base metal a-phase. After 1020°C the cross-section size of 6-phase particles is also increased (Fig 3a). At the same time, microcracks are observed which are completely loc.alized in 6-phase (Fig. 3b). Stresses in the top of the' crack are relaxed in a-phase by slipping.

In weld metal after heat treatment at 1020°C for 2 h (a + 6)-structure is formed, it being similar to that of the base metal, but having a very· wide interphase layer (Fig. 3c). Micro-diffraction analysis has shown that Burgers orientation relation is observed between a- and 6-crystal lattices, and the investigations of interphase layer image in various reflexes, if we accept the authors [7, 8) point of view, show the presence of 2 phases here simultaneously. One of them has a fee structure, the other, prevalent in the interphase layer, may be described as a h9P phase with orientation, differing from the main a-phase.

Internal Friction

Temperature dependence curves 01 (Fig. 4a) consist of the background and of peaks. To separate these components of damping the background values in the 300 - 700°C range were extrapolated by the exponent [12J, and analysis was performed for peaks lying in the temperature range up to 550 - 6oo 0 c.

580 V. N. Gridnev et al.

At heating up to 200°C the peaks 01 were absent, background values were not higher than 3,10- , that is why these parts of curves are not given.

Let us consider the quantitative composition of peaks in the investigated alloy, taking the curve of the base metal as an example (Fig. 4b, c, d, curve 1). Fig. 4d shows the 01 spectrum profile. Its decomposition by the method given in [4], permits to single out 5 elementary peaks.

The peak, which was for the first time observed in titanium at 450°c and 1 Hz frequency [13] and was connected with the presence of oxygen in it was the most studied. According to [14] the specified maximum is only dis­played in the presence of substitution elements. The latter causes localized distortions of octahedral pores, which are spherically symmetric in pure titanium. This is the cause for relaxation of interstitial atoms with respect to substitution elements in hep lattice in the stress field. According to the data [14] in titanium, containing aluminium - Tmax = 450°c, H = 45 kcal/mol, and vanadium~ 470°C (f = 1 Hz). The appropriate Tmax at the used frequency of 4.3 Hz will be equal to 476 and 506°c. It is indubitable, that III and IV peaks reflect oxygen relaxation with respect to the atoms of aluminium and vanadium in a-phase of the alloys respectively. Hof the peak IV is 47,5 kcal/mol. Tmax values obtained as a result of decomposition are 478 and 515°C. It should be added that the best agreement with the curve is obtained, if we take into consideration the braodening of the peaks 81 [4], which equals 0.8-1 according to data of Fig 4d. Such a broadening was observed for Snoek peaks and b,c.c. metals [15] and is, probably, connected with a distribution of relaxation times.

As for the II peak nature (Tmax = 420°c) a number of assumptions can be made. Miller [16] has presented data on the presence of a peak in iodide titanium with a lower content of interstitial elements, than in [14], which has found an explanation in theoretical works [17, 18]. They showed a possi­bility of relaxation of pairs of oxygen atoms(O-O)in a h.c.p. lattice, in particular, in Ti-0 system [18]. The occurence of 0-0 relaxation in titanium has been experimentally proved in the paper [19]. The authors of this paper also believe, that relaxation within 3 oxygen atom clusters is possible too. Tmax of the peak is 430°C, f = 1 Hz and H = 52.7 kcal/mol. At recalculation for f = 4.3 Hz frequency we obtain Tmax = 462°C, i.e. a partial superimposi­tion on peak III is observed. However, this peak can hardly be observed in a phase with Al content not lower than 6%. In papers [20, 21) a peak was observed which was assigned to Si-0 pairs in a-titanium. Silicon content in the investigated alloy is 0.04 w.t.%, i.e. if such an assumption is admissible, then the appropriate Tmax should be about 350 - 370°C and olmax lower, than the observed value.

We are inclined to believe that II peak is connected with oxygen relax­ation in 8-phase. According to the date [22) molybdenum and chrome do not affect Tmax (-200°C, f = 1 Hz) of the oxygen peak in 8-phase. In connection with this the authors have to come to the conclusion that Ti-0 link is stronger. However, vanadium may lead to a redistribution of oxygen and its localization around vanadium, and this will lead to increase of Tmax. Moreover, taki.ng into account a high content of vanadium in 8-phase one can expect a formation of a number of complexes (V-0, V-0-V) by analogy to Nb-Zr-N system [15), and of peaks with different Tmax respectively. In· this respect, peak 1 is deter­mined by V-0 complexes (Tmax = 370 - 380°C) whereas peak II (Tmax = 420. -430°C) by (V-0-V). The height relationship of the peaks quite correlates with the expected growth of concentration of the latter at increasing the vanadium content in 8-phase. The relation between the peaks and oxygen content in the alloy is confirmed by their growth after heating the samples in the air. ·

BT6 ALLOY WELDED JOINTS 581

The accuracy of determination of Tmax of peak V. ( 550 ,- 570°C) an.d its shape is low, since it is situated in the region of high.values of 01 back­ground. It must be the reason for the fact that in a number of cases the peak width is narrower than the theoretical value. According to (16] in this temperature interval one should expect peaks due to nitrogen in a-titanium.

On curves o1 bends are observed in 620 - 630°c and 680 - .700°C regions (Fig. 4a). In commercial titanium and its oxygen containing alloy grain boundary peaks correspond to thes~ temperatures [13].

Thus, in BT6 alloy the peaks at 470 and 515°C are due to the presence of oxygen: at 560 - 580°C probably of nitrogen in a-phase, whereas the peaks observed at 370°C and 420°C are due to oxygen in S-phase (at f = 4.3 Hz). In 8-phase oxygen is localized near vanadium a,toms.

Let us follow. the change of the above-mentioned peaks in welding joints after heat treatment (Fig. 4b, c, d). o1 curve bf an as-welded joint shows an early rise of damping at which the peaks are almost not displayed (Fig, 4b, curve 2). This correlates quite well with the existence of a deformed a'-phase. Module defect on the heating curve (Fig. 4a, curve 2"1. starts from 560°c, i.e. from this temperature in as-welded joints a noticeable mobility. of structure elements is observed at heating with a 3 deg/min rate. In welds, annealed at 750°C for 2 hours, the lowest peaks (Fig. 4c, d) are observed in · a- and 8-phases (Fig. 4c, d, curves 3). In all the states peak III is the most pronounced one. It is quite understandable, due to the preval~nce of a-phase in the alloy and a high aluminium content in it. Heat treatment of welded joints for 48 hours leads to increase of the oxygen content in 8·-Phase and nitrogen content in a-phase (Fig. 4c, d, curves 4). Maximum gas content according to 01, data (curves 5) is observed in welds after annealing at 1020°c.

It should be noted that we have observed an increase of peaks height in 8-phase, which seem to be connected with oxygen localization near vanadium atoms, both in base and weld metal after natural ageing for a long time (2 years) (Fig. 4c, d, curves 5).

01 background characterizes a temperature dependence of defect mobility, in particular ductile movement of dislocations (12]. Therefore, the more stabilized is the structure, the lower is the background. That is why, the earliest rise of background is observed in as-welded· joints, and the one which is the most displaced into the region of increased temperatures - after annealing at 750°c for 48 hours. The welds treated at 1020°C, for 2 hours in this respect are closer to those subjected to 750°C, 2 hours treatment. And though the differences are not large, the trend of such a regularity is distinctly displayed in a series of samples.

The problem of interaction of interstitial in particular, of oxygen, with dislocations in titanium is ambiguous. Theoretical estimations on the basis of elastic interaction permit ·us to obtain a value of the order of 0.1-0.3 ev. If one takes into consideration strengthening at the expense of oxygen atoms, 1.4 ev. value will be obtained [23], which is an experimentally obtained one. The explanation of the divergence in [24] is based on the conclusion that it is necessary to take into consideration, not only elastic but also, chemical interaction; arid in particular, the number of chemical bindings realized.

The 01 amplitude dependence in titanium is of a linear nature in a wide range of deformation values (Fig. 5a). The slope of the~ straight lines.

582 V.N. Gridnev et al.

being more pronounced in deformed samples, is decreased after· annealing, and it is believed to depend on interstitials concentration on dislocations C (25], and above the temperature of atmosphere condensation (TkL changes according to the well-known Cottrell relation:

l!. C = Coe kT

C0 - concentration of interstitial elements in the solution; U - enthalpy of the binding between solutes and dislocations.

Presenting the results in log ~~l coordinates from T-1 enables to esti­mate the interaction energy of interstitials with dislocations. In iron­based alloys appropriate values for interaction energies have been obtained using this method (26].

The processing of the data shown on Fig. 2a by the above method gives the value of interaction energy equal to 1.4 :t 0.1 ev (Fig. 2b), Tk = 520 -530°C. It should be noted that this value is to be considered as an estimated one. It is feasible that other factors (deformation anisotropy, twinning) can also affect the 81 amplitude Qependence in titanium.

Discussion of Results

It ·appears to be of interest to analyze the whole complex of experimental results presented above to determine their correlation with the mechanical characteristics.

One of the important factors, determining crack developments, is the possibility of stress relaxation, depending on the plasticity of a- and 8-phase components. The results of electro.n microscope and internal friction investigations indicate a higher plasticity of a-phase in the alloy. Indeed, the nature of amplitude dependence 81 shows that at microdeformation (E = lo-3 - lo-5) the a-phase is the first to deform. At the heat-treatment conditions studied the plasticity of a-phase remains high enough. Therefore, the observed changes of strength and plasticity should be chiefly determined by the 8-phase and the state of interphase layers. The decrease of plasticity and strength of the base metal after annealing at 1020°C is connected with 8-phase embrittlement and the appearance of interphase layers. Such heat­treatment leads to the increase of oxygen content in the alloy , which is registered by the peak increase in a- and 8-phases. It is believed, that as a result of the stabilizing heat treatment oxygen in (a + 8)-area should be mostly concentrated in a-phase, where its stability is much higher. However, an increased content of vanadium in 8-phase, which possesses a high affinity to oxygen, provides for the preservation of a considerable amount of oxygen in a solid solution, thus causing 8-phase embrittlement.

As has been shown, a-phase deformation occurs not only in the form of sliding, but of twinning. Three cases of stress relaxation have been observed on interphase boundaries.

1. Slip lines, which have appeared in a-phase, freely passed the adjacent 8-plates, causing a shear in them.

2. Stresses in the retarded twin top did not relax by crystallographic shear, but caused crack formation in 8-phase.

3. A crack, formed in 8-phase, is completely localized within this phase boundaries, since stresses in its tip are relaxed in a-region.

BT6 ALLOY WELDED JOINTS 583

This phenomena can be accounted for in.connection with the presence or absence of an interphase layer on a- and 6-phase interface. Thus, at its absence plastic deformation is easily transmitted from a- to 6-phase, not leading to destruction. But even a thin interlayer between a- and 6-phase appears to be a serious obstacle for slip propagation: stresses, which have been accumulated here, appear to be enough for destruction of a more brittle neighbouring 6-phase. From our point of view, this determines decrease of plasticity from 12 to 4% after weld heat-treatment at 1020°C, when a struc­ture with a very wide a/6 interphase layer and 6-phase with an increased oxygen content are formed after transition into the 6-area. An optimum propert~es combination is observed after annealing of welded joints at 750°C for 48 h. This may be connected with .a further plasticizing of a-phase in connection with the formation of fine-dispersed precipitations. The precipi­tated phase, probably, is an intermetallic compound of titanium and aluminum. A possibility of such an identification quite correlates with the results of paper [27]. The authors of this paper managed to show, that microhetero­geneity is already observed in solution at 2% Al content, and at 5% and higher Al content a2-phase of Ti3Al type is precipitated. The yield point of Ti3Af compound according to [28] data is 50.4 kgf/mm2 (for TiAl - 42 kgf/ mm2 ), this being lower than in BT6 alloy a-phase •. A frequently observed pre­valence of phase precipitation in the region, adjacent·to a/6 -interphase boundaries, is accounted for by a predominant enrichment of these zones by aluminum in the process of its redistribution during decomposition.

References

[ l]. Metallurgy and technology of welding of titanium and its alloys. (S.M. Gurevich ed., K., Naukova dumka, 1979, 300 pp.).

[ 2]. A. Dukro. TA6V titanium alloy electron beam welding. In col.: "Titanium. Metal science and technology. Trans. of 3rd Metallurgy Conference on Titanium". M., VILS, 1978, v.2, pp.321-330. .

[ 3]. I.V. Holl, K. Hammond. Fracture toughness, strength and microstructure of (a + 6)-titanium alloys. In col: Titanium. Metal science and tech­nology. Trans. of the 3rd International Conference on Titanium. M, VILS, v.l, pp. 351-356. '

[ 4]. A.S. Nowick, B.S. Berry. Lognormal Distribution Function for Describing Unelastic and Other Relaxation processes. J. Research Develop., 1961, 5(4), pp. 297-320.

[ 5]. P. Hirsh, A. Hovi, P. Nikolson, D. Pashly, M. Wellan. Diffraction geo­metry on twin structures. In: Electron microscopy of thin crystals. M., Mir, 1968. pp. 153-160.

[ 6]. A.B. Notkin, L.M. Utevskii, A.S. Feinborn. Analysis of electron micro­graJns of twin hexagonal crystals of a-phase in titanium alloys. In col. : Application of X-ray and raster type electron microscopy in metal science. Materials of the seminar. MDNTP, M., 1976, pp. 3-12.

[ 7]. C.G. Rhodes and J.C. WilliaJns. Observation of an interface phase in the a/6-boundaries in titan.ium alloys. Met Trans., 6A, N.8, pp. 1670-1671.

[ 8]. C.G. Rhodes ·and N.E. Paton. Formation characteristics of Uxa - 6 -interface phase in Ti-6Al-4V. Met. Trans., 1979, lOA, N.2, pp. 209-216.

[ 9]. H. Margolin, E. Levine, M. Jound. The interface phase in alpha~beta titanium alloys. Met. Trans., 1977, BA, N.2, pp.373-377.

[10]. K.J. Roads, N.E. Paton. Heat-treatment effect on a/6 interphase layer structure in (a + 6) titanium alloys. In col.: Titanium. Metal science and technology. Trans. of the 3rd International Conference on Titanium. 11., VILS, 1978, v.2, pp. 489-499.

584 V. N. Gridnev et al.

[11]. G. Tsviker. Titanium and its alloy. Metallurgia, 1979, pp. 132-136. [12]. G. Schoek, E. Bisogni, J. Shyne. The activation energy of high tempera-

ture friction. Acta Met., 1964, 12, N. 12. · [13]. J.N. Pratt, W.J, Bratina and B. Chalme.rs. International friction in

titanium and titanium-oxygen alloys. Acta Met., 1954, 2, pp. 203-208. [14]. D. Gupta and D. Weinig. Interactions between interstitial and sub­

stitutional aolutes in an h.c.p. lattice. Acta Met., 1962, 10, N.4, pp. 292-298.

[15]. V.N. Gridnev, N.P. Kushnareva, V.N. Minakov, E.V. Turtsevich. The connection between internal friction relaxation spectrum with a struc­ture state in a niobium-based alloy. Metal physics and metal science. 1977, M.3, 44, pp. 575-581.

[16]. Miller D.R. Internal friction of titanium and its alloys. Trans. of AIME, 1962, 224, N.2, pp. 275-281.

[17]. E. Povolo, E.A. Bisogni. Mecha.i-iical relaxation modes of paired point defects in h.c.p. crystals. Acta Met., 1966, 14, N.6, pp. 711-718.

[18]. K.M. Browne. Mechanical relaxation and diffusion of interstitial atoms in h.c.p. metals. Acta Met., 1972, 20, pp. 507-514.

[19]. J.A. Bertin, F. Gacougnolle, S. Sarrazin et.J. de Fouquet. Pie de fro­ttement interieur du a l'oxygene dans le titane de haute purete et dans le titane a 0.02% fer. Nuovo Cimento, 1976, B33, p. 302-307.

[20]. P.A. Bleasdale and D.J. Bacon. Internal friction due Zo inter-stitial­substitutional impurity complexes in some dilute titanium alloys. Nuovo Cimento, 1976, B33, N.l, 308-315.

[ 21]. S. Mishra, M.K. Asundi. Internal friction studies in titanium base alpha and beta alloys. Titanium Sci. and Technol. vol. 2, New-York-London, 1973, pp. 833-903.

[22]. V.N. Gridnev, V.F. Gribin, A.I. Efimov, P.G. Yakoveriko. Effect of alloying and of annealing temperature on titanium alloy structure state. In col.: Titanium. Metal science and technology. Trans. of the 3rd International Conference on Tit'anium. VILS, 1978, v.2, pp. 555-559.

[23]. B.A. Kolachev. Physical metal science of titanium. M., Metallurgia, 1976, 186 p.

[24]. W.R. Tyson. Solution hardening of titanium by oxygen. Scripta Met., 1969, 3, N.12, pp. 917·-922.

[25]. G.A. Beresnev, V.I. Sarrak, NiA. Shilov. Temperature dependence of iron deformation resistance. In col.: Problems of metal science and metal physics. M., Metallurgia, 1968, v.9, pp. 157-165.

[26]. V.G. Gavriljuk, N.P. Kushnareva, V.G. Prokoperiko. Alloying effect on dislocation mobility in iron. Metal physics and metal science, 1976, 42, N.6, pp. 1288-1293.

[27]. L.S. Moroz, I.N. Rasuvaeva, S.S. Ushkov. Peculiarities of aluminium effect on mechanical properties of titanium, In col.: A new construc­tion materials titanium, M., Nauka, 1972, p. 109-114.

[28]. J.C. Williams, M.J. Brockburn. Ordered alloys. B. Kear, T. Sims, N. Stoloff, J. Westbrook. Claitor's Publ. Div., 1970, p. 425,

Table 1

Mechanical Properties of BT6 Welded Joil:lts

Mechanical Annealing conditions characteris- No aD.llealiJlg 750°c, ·2 h 750°C, 48 h 1020°c, 2 h

tics Base Weld Base Weld Base Weld Bue Weld metal metal metal metal me~al metal. metal metal

6t,.z J kgt/aua2 20 •• 2;! 22· .100 88. ·21 22··28 02 •• 02 00 .. 22 82 •• 86 a2 •• 02 93x 97 89 96 87 90 84 83

6 2 az •• 22 20 .. 26 a2 •• az 21. ·22 a1 •• a2 02 •• a2 Z2· .BJ 81 •• BJ qz 'kgf/mm 89 93 86 93 83 87 81 82

b, % 6,6 •• 13,3 6.0 •• 10.0 a.2 •• 12.1 z.0 •• 10.0 6.6 •• 12.3 9.3 •• 14.3 B.6 •• 1,2.3 J.3 •• 6.7 . 10.3 7.5 10.7 a.9 12.4 12.6 11.7 4.7

'f, % J0 ... 36 12. ·.2.2 21 •• 4z J2 •• 4J 40 •• 46 ,22 •• 64 24 •• J1 14 •• 1z 34 . 24 36· 39 42 42 27 15

G J kgf/mm2 1JJ .. 14Z 112 •• 148 1JO .• 16z 144 •• 166 121 •• 1z2 1,24 •• 222 21 •• 124 26 .. 28 137 129 142 158 163 172 107 97

3/j . ~c, kgf /mm50 •• 1412 240!.2z2 340 •• 400 222 •• 2zo 412 •• 4J2 J.22· ·JZO .222 •• 424 J2J •• J6,2

390 xx 260 370 245 425 353 404 344

· x Mean values acc. to the results of testiag 5 samples are given ill the deaominator. xx Tests were performed on 3 samples.

VI 00 VI

586 V.N. Gridnev et al.

a)

b)

c)

Fig. 1

BT6 ALLOY WELDED JOINTS 587

Fig 1. Electron micrograph& of weld structure : a. dislocations ' b. twins ill weld metal ; c. experimental and d. estimated diffrac­

tion micrograph of twiDDed ~(~ ) - phase cristals.

a) b)

Pig 2. Electron micrograph& of BT6 alloys 1n an ilrl.­tial state : a. - (o< +fa) - alloy structure b. - crack propagation.

d)

588 V.N. Gridnev et al.

-,.----- -~,

Fig J. ElectroA micrograph& of BT6 alloy after welding and 8.DJlealing :

a)

b)

c)

a. dispersion precipitations in o{-phase after anne­aling at 750°0, 48 h.

b. - a crack in j3- phase and streess relaxation in ol- phase after annealing at 1020°0, 2 h '

c. weld structure after annealing 1020°0, 2 h.

6;10' G"'l'· <\·10' zo

120

:IO <\'~

4

18

15

16 400 6()0

r.·c

BO

1~ 10

60-

12

40-

....

Fig. 4a Fig. 4b,c

590 V.N. Gridnev et al.

Fig. 4.

400 500 T,~O

d

d)

Temperature depeadeace of a logarithmic decremeat tf (1-5) ud shear modu.lus (2", 3") of BT6 alley samples iJl the iaitial state (1), of a weld without heat-treatmeat (2) aad welds, aBDealed at 750°0, for 2h (3), 750•0, 48 h (4) 8.Jll.d 1020°0 2h (5). 2'1a Fig.4a - is a oooli.ag curve a weld metal sample, aot heat-treated; Pig.4c - ~ peak spectres; Pig.4d - decomposition of Pig.4c curves; 5•curves ill Fig.4c, d correspond to a sample, kept at room temperature tor two years.

BT6 ALLOY WELDED JOINTS 591

1/, IO"

651!

35

30

25 -60/J'

a)

20

IS 570

10 ~sJO' ..-5rXJ'

5 m· - wo· =

-~20·

5 /U IJ 20 25 {,. JD'

Pg~ I ~1

z\ b)

0 '\. l,J 1.2 1,3 1.4 (5

r'-10~ K

Fig. 5 Amplitude dependence Si , at various temp'eratures (

tJ..f1 of base metal sample a) and 2'l"" change, depending oa T-1 (b) s 1. in initial state;

2. after B.DRealing at 750°c,2m.