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The β→ω transformation during room temperature aging in rapidly solidified Ti-6Al-4V alloy
Pergamon
ScriptaMetallurgicaet Materialia,Vol. 31, No. 11, pp. 1519-1524, 1994 Copyright© 1994 ElsevierScienceLtd Printed in the USA. All rights reserved 0956-716X/94 $6.00 + 00
THE [5--->coT R A N S F O R M A T I O N DURING ROOM TEMPERATURE AGING IN RAPIDLY SOLIDIFIED Ti-6A1-4V ALLOY Z. Fan Department of Materials Science and Engineering University of Surrey, Guildford, Surrey, GU2 5XH, UK ( R e c e i v e d May 12, 1994) ( R e v i s e d J u l y 8, 1994)
Introduction Since its first discovery by Frost et al [1] in aged Ti-Cr binary alloys, the co-phase has received extensive study, initially because of its deleterious effects on mechanical and physical properties, but later also due to the interests in the transformation mechanism. It is well known that the ~-phase can form by either quenching from high temperature (athermal co-phase) or aging at low temperature (isothermal co-phase) [2,3]. The ideal co-phase has a hexagonal structure, and a well defined orieuntation relationship with the parent bcc [5-phase, as first reported by Silcock [4]: {Ill]B//(0001)~, [ 110] I~//[ 1120] ~. Other structures of the co-phase have also been reported in the literature, for example, tfigonal co, ordered co, Zr2Al-type co and Ni2Al-type co (see review in Ref [5]). Recently there has been a renewed interest in co-phase transformation due to the development of aluminide-based intermetallic materials, where an ordered B2 phase can transform into an ordered co-phase during low temperature aging [5]. Experimental observation of the co-phase in commercial Ti-6AI-4V alloy has been very rare [6-8]. In fact, the occurrence of the co transformation in this alloy has been excluded by the argument that the concentrations of ct-stabilisers (A1 and O) in the [5-phase are sufficient to suppress the co-phase formation [6,7]. There appears to have been one report of the occurrence of the co-phase in quenched and aged Ti-6AI-4V alloy containing 0.116 wt % Fe and 0.15 wt % O [8]. In this paper we will report the experimental observation of the 13~co transformation during room temperature aging of a rapidly solidified Ti-6A14V alloy. Experiments The master Ti-6A1-4V alloy was supplied by DRA (UK), which contains 6.43 A1, 4.02 V and 0.19 Fe (in wt.%). Melting and melt-spinning were performed in a MARKO's Model 5T melt spinner in a stainless steel chamber with a high vacuum/inert gas (argon) atmosphere. The alloy was melted in a water cooled copper hearth using a non-consumable tungsten electrode under a high purity argon atmosphere. The chemically homogeneous melt was then delivered at a controlled rate to contact the circumferential surface of a Mo wheel rotating at 2500 rpm by which the melt is rapi~lly solidified as long fibres, which have a crescent-shaped cross-section, typically 100~300 I.tm in width and 40~100 I.tm in thickness. The melt spun fibres were comminuted into finer particles with particle size less than 200 gin. The comminuted alloy powders were then consolidated by hot-isostatic-pressing (HIPping) at 900°C and 300 MPa for 2 hours, and furnace cooled to room temperature. Further details for preparation of the materials under this study were given elsewhere [9]. Room temperature aging was carried out in air for about a year. Thin foils for TEM study were prepared in a "TENUPOL" unit using an electrolyte of 5 vol % perchloric acid in methanol. The polishing temperature was around -40°C, the voltage 30V, the current about 30 ~tA. TEM observation was performed on a JEOL 200CX STEM under an accelerating voltage of 200 kV. EDX analysis was performed on a JEOL JXA8600 EPMA system for phase composition. E x p e r i m e n t a l Results Microstructure of an as-HIPned Ti-6A1-4V Alloy The results from chemical analysis of the Ti-6A1-4V alloy after different processing steps are tabulated in Table 1. The oxygen level has increased after the melt-spinning process, especially after the comminution process, from 2050 ppm in the master alloy to 3100 ppm in the comminuted alloy powder. The oxygen content in the as-HIPped alloy is just slightly higher than 3100 ppm due to the good vacuum level in the HIPping process. 1519
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The microstructure of the consolidated Ti-6A1-4V alloy is show in Fig 1 by a low mag TEM micrograph. The a-phase has an equiaxed morphology with an average grain diameter about 6 gm, while the l~-phase has an irregular shape and is located at (x-grain boundaries and (x-grain triple points. The volume fraction of the [3-phase is about 0.08. The (x-grains are single crystals in which there is no sign of any transformation. The chemical compositions of the (x- and l-phases obtained by EDX analysis are presented in Table 2. These results indicate that c~-phase is rich in A1 and contains oractically no Fe, and that there is an enrichment of V and Fe in the 13-phase. In addition, a compositional variation in the p-phase has also been observed during the EDX analysis. The larger [3-grains located at et-grain triple points seem to contain less t3-stabiliser, while the smaller ~-grains located at co-grain boundaries contain more 13-stabilisers. TABLE. 2. EPMA results of phase composition in the TAB. 1. Results of impurity analysis of Ti-6AI-4V consolidated Ti-6A1-4V alloy (in wt %). alloy after different processing steps (in wt.%). Process Master allo), Melt-spun fibre Comminutedpowder
C 0.01 0.015 0.020
N2 0.0075 0.009 0.014
02 0.205 0.230 0.310
Phase (x ~
T1 Bal. Bal.
A1 6.53 3.24 . (3.08-3.89)
V 3.09 13.18 10.15-14.25
Fe 0.00 1.61 (1.24-2.01)
The detailed TEM study of the metastable ~-phase decomposition in as-HIPped Ti-6A1-4V alloy by the present authors has been presented elsewhere [10]. It was found that during the continuous cooling the metastable 13-phase can decompose by different modes, depending on the actual composition of the [~-phase. Less enrichment of V and Fe in the [3-phase will favour a direct formation of the equilibrium (x-phase from the 13-matrix, while the greaterenrichment of V and Fe can lead to the spinodal decomposition of the metastable [5-phase, resulting in the [3-stabiliser lean 15'-phase and the further enrichment of V and Fe in the I~-matrix. Upon further continuous cooling, the isothermal m-phase can form from the spinodally decomposed 13'-phase[10].
F I G 1. TEM BF image of the rapidly solidified Ti-6A14V alloy HIPoed at 900°C, 300 MPa for 2 hours, and furnace cooledto room temperature. The c~ and 13 phases are indicated by c~and 13in tile micrograph,
FIG 2. TEM BF image of a [~-grain showing ~ e typical features of the spinodal decomposition, p---)lJ+[3. The electron beam is close to [ 1]-0]13,and the dtrection for the compositional modulation is <] 11 >[~.
The detailed substructure of a 13-grain in an as-HIPped Ti-6AI-4V alloy is presented in Fig 2, which shows the typical features of a two-phase structure resulted from a spinodal decomposition reaction, 13---)[3+13',where 13 and [3' are coherent with each other [11]. In this composition modulated two-phase microstructure, the if-phase is lean of [3-stabilisers (V and Fe), while the l-matrix is further enriched by 13-stabilisers. The direction of the compositional modulation is < 111 >B [ 10]. SAD patterns from the spinodally decomposed [$-phase are shown in Fig 3. Besides the strong bcc reflections, diffuse scattering along a variety of directions is present in these SAD patterns, which is characteristic of the initial stage of the ~---)co transformation and is known as diffuse co scattering [12]. Occasionally, discrete spots corresponding to the co reflections can
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be observed in SAD patterns taken from such 13-grains. An example is given in Fig 4 where the electron beam is parallel to [ 110]~. The co reflections in Fig 4 are shifted systematically along the c axis of the co unit cell (parallel to. <11.1>~ directions). from the Bragg co positions, i.e., 1/3{222} 6 and 2/3{222}[~ for (0001)o~ and (0002)o in the case of onentatlon variant cob to positions of 0.36{222}~. and 0.64{222}~. This systematic shift of the co reflections has also been observed in other Tibase alloys, such as Ti-Fe, Ti-Mn and Ti-Cr, and has recently been discussed in detail by Sinkler and Luzzi [13].
F I G 3. SAD patterns taken from the spinodall2¢ decomposed [3-grains showing the diffuse co scattering which is characteristic of the early stage of the [3 to co transtormation. (a) [011113 zone axis pattern; (b) []-33][~zone axis pattern.
F I G 4. [110]13 zone axis pattern taken from a spinodall~, decomposed 13-grain showing the discrete co spots which are shifted from the Bragg co positions ( 1 / 3 { 2 2 2 } B and 2/3{222}[~) to positions of 0.36{222}[~ and 0.64{222}[3.
F I G 5. TEM BF image showing the appearance of a 13grain after aging at room temperature for a year. The fine cubic co parucles can be seen clearly in the [5-matrix.
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Microstructure after A~in~, After aging at room temperature for a year, the general ct+~l microstructure was not changed. However, the substructure of the p-phase has been altered. A TEM BF image of a [3-grain after room temperature aging is presented in Fig 5, where the free co precipitates can be seen clearly. Electron diffraction on such ~-grains indicated that B~o~ transformation has occurred at an advanced stage, as evidenced by the discrete spots from the co reflections in SAD patterns. Two examples of such SAD patterns are shown in Fig 6, with the electron beam being parallel to [110]~ (Fig 6a) and []13]B (Fig 6b) respectively. It is very interesting to note that after room temperature aging the systematic shifts of co spots in the as-HIPped state (Fig 4) have been substantially reduced, these co spots in Fig 6 appear at the positions very close to 1/3{222}[~ and 2/3[222 }[~, which correspond to the Bragg co reflections.
F I G 6. SAD patterns taken from I]-grains after aging at room temperature for a year showing the discrete spots corresponding to the Bragg co reflections, ta) [110115 zone axis pattern; (b) []'15115zone axis pattern.
F I G 7. TEM DF image using (000t)co reflection showing the detailed morphology of the to-phase after aging at room temperature for a year.
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A TEM dark field image using the (0001)o~ reflection is shown in Fig 7. The m-phase appears in the ~-matrix as fine particles with a diameter around 5 to 10 nm. In addition, the detailed TEM dark field observation revealed that the co particles tend to form parallel strings in the 13-matrix, as evidenced in Fig 8 by the comparison between the BF and corresponding DF TEM images of a 13-grain, where the electron beam is close to [110]~. This result indicates that the to-phase is transformed from the spinodally decomposed [~'-phase which is lean of [5-stat~itisers.
FIG 8. TEM BF image of a 13-grain after aging at room temperature for a year (a) and its corresponding DF image (b) showing the fine cubic co particles which tend to form parallel strings. In both (a) and (b), the electron beam is close to [ll0]~. Discussion The occurrence of the m-phase in as-HIPped Ti-6A1-4V alloy can be mainly attributed to the following three factors. Firstly, the sufficient enrichment of 13-stabilisers in the 13-phase makes the co-phase formation thermodynamically favourable in competing with the martensitic transformation [14]. EDX analysis showed that V and Fe concentration in the retained 13phase are more than 13.18 and 1.61 wt %, respectively (Table 2). The effectiveness of Fe as a 13-stabiliser is about 4 times as much as that of V on a weight percent basis. Therefore, the equivalent V concentration in the retained ~phase is over 20 wt %, which is sufficient to suppress the martensitic transformation [14-16]. This allows the m-transformation to replace the martensitic transformation during the continuous cooling. Secondly, the slow cooling rate from the HIPping temperature can create a favourable kinetic condition for the co-phase formation. Recently, Moffat and Larbalestier [17] studied the effect of cooling rate from the annealing temperature to room temperature on the metastable 13-phase decomposition mode in Ti-Nb binary alloys. They found that precipitation of the m-phase was favoured by a slow cooling rate, in contrast, the formation of cO' martensite is favoured by a higher cooling rate. The present results in the Ti-6A1-4V alloy support the above observation in Ti-Nb alloys. The cooling rate dependence of the mode of metastable 13-phase decomposition can be understood in terms of the role played by diffusion during the decomposition process. Martensitic transformation is diffusionless, therefore, diffusion has no influence on it, while the isothermal m-phase transformation involves nucleation and growth, a slower cooling rate thus favours the m-phase formation by allowing the extent of diffusion required by the co transformation process to be completed. Finally, the formation of a solute lean l~'-phase through the spinodal decomposition of the metastable 13phase favours the co-phase formation in terms of compositional depletion of 13-stabilisers and lattice plane shift through the coherency strain. It has been recently confirmed by atom probe analysis that isothermal co precipitates are depleted in all the alloying elements, both tx- and 13-stabilisers [18]. Therefore, the formation of a solute lean 13'-phase will reduce the amount of diffusion required by the formation of the m-phase compared with its direct formation from the parent 13-phase. In addition, the coherency strain between 13and 13' phases will make both bcc structures distorted towards a trigonal structure by shifting the { 111 }1~ planes. Such a trigonal lattice is an intermediate structure between the bcc 13 and the ideal co (hexagonal) structure. However, the continuous cooling may not create a thermodynamic and kinetic condition which is favourable enough to complete the 13'---~co transformation to form the crystalline co-phase, and hence the m-phase can only be observed by electron
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diffraction in a form of diffuse co scattering. Aging at room temperature will further enhance the thermodynamic and kinetic conditions by allowing the diffusion required by the transformation. It is expected that at room temperature, the diffusion of V, A1 and Fe in bcc titanium is very slow. Therefore, a much longer time is necessary. Nevertheless, if the aging time at room temperature is long enough, the crystalline co-phase can be formed from the [~-stabiliser lean [~'-phase. The similar phenomenon has also been observed by the same author in a rapidly solidified Ti-7.5Mn-0.5B alloy after aging at room temperature [19]. Furthermore, the reduction of the systematic shifts of the co-reflections from the Bragg positions after aging at room temperature may indicate the existence of the followingtwo processes during the room temperature aging. Firstly, it suggests that the double plane collapsing required by thep--)co transformation in this alloy has not been completed after continuous cooling and can be completed gradually during the room temperature aging. Secondly, accompanying the double plane collapsing process, there is a depletion process of [3-stabilising elements (V and Fe) from the co-phase during the room temperature aging. As has been discussed by Sinkler and Luzzi [13], the magnitude of the co maxima shifts in SAD patterns depends on both the solute concentration and the identity of the solute. The amount of the co maxima shift increases with increasing solute concentration beyond the limit for co-phase formation. In addition, the amount of the co maxima shift also increases for a given solute concentration as the group number (in the Periodic Table) of the solute element increases. Therefore, the reduction of the systematic shifts of co-reflections suggests that a depletion process of V and Fe from the cophase is also present during the aging. Concluding Remarks Durin the continuous cooling of the Ti-6A1-4V alloy the metastable p-phase can decompose spinodally to form the ~stabil~er lean 13'-phase and the ~-stabiliser rich ~-phase. Upon further continuous cooling, the isothermal co-phase can form from the spinodally decomposed lY-phase. However, the thermodynamic and kinetic conditions created by the continuous coolin" may not be enough to complete the [~"--)co transformation to form the crystalline co-phase, and hence the co-phase can only b~ observed by electron diffraction in a form of diffuse co scattering. Aging at room temperature will further enhance me thermodynamic and kinetic conditions to such an extent that the crystalline co-phase can be formed from the 13-stabiliser lean [3'-phase, if the aging time is long enough. It is also possible that a depletion process of V and Fe from the co-phase occurs during the room temperature aging. Acknowledgements The author thanks Dr. L. Chandrasekaran for his help with orocessing the Ti-6A1-4V alloy, Professor A. P. Miodownik and Dr. P. Tsakiropoulos for their useful discussions. Ttils wor[¢ has been supported by the MOD, UK. References [1] P. D. Frost, W. M. Parris, L. L. Hirsch, J. R. Doig and C. M. Schwartz: Trans. ASM, 46, 231 (1954). [2] J. C. Williams: in Titanium Science and Technology, R. I. Jaffee and H. M. Burte eds, Plenum Press, New York, Vol 3, 1973, p1433, 1973. [3] S. K. Sikka, Y. K. Vohra and R. Chidambaram, Progress in Materials Science, 27, 145 (1982). [4] J. M. Silcock: Acta Met., 6, 481 (1985). [5] C. P. Chang and M. M. Loretto: Philos. Mag., 63A, 389 (1991). [6] J. C. Williams and M. J. Blackburn: Trans. ASM, 60, 373 (1967). [7] J. C. Williams, B. S. Hickman and D. H. Leslie: Met. Trans., 2, 477 (1971). [8] A. Lasalmonie and M. Loubradou: J. Mat. Sci., 14, 2589 (1979). [9] Z. Fan, A. P. Miodownik, L. Chandrasekaran and M. Ward-Close: J. Mat. Sci., 29, 1127 (1994). [10] Z. Fan and A. P. Miodownik: J. Mat. Sci., in the press, 1994. [11] J. W. Cahn: Trans. Met. Soc. AIME, 242, 166 (1968). [12] J. C. Williams, D. de Fontaine and N. E. Paton: Metall. Trans., 4, 2701 (1973). [13] W. Sinlder artd D. E. Luzzi: Acta Metall. Mater., 42, 1249 (1994). [14] N. A. Vanderpuye and A. P. Miodownik: in The science, technology and applications of titanium (Jaffee and Promisel eds.), Pergamon Press, p719, 1970. [15] P. Pietrokowsky and P. Duwez: Trans. AIME, 194, 627 (1952). [16] T. Sato, S. Hukai and Y. C. Huang: J. Aust. Inst. Met., 5, 149 (1960). [17] D. L. Moffat and D. C. Larbalestier: Met. Trans., 19A, 1677 (1988). [18] L. Hadjaj, A. Menand and C. Martin: quoted by N. E. Paton and H. L. Fraser: in Proceeding of Sixth Worm Conf. On Titanium, P. Lacombe, R. Tricot and G. Beranger eds., Societe Francaise de Metallurgie, France, Vol 3, p1469, 1988. [19] Z. Fan: Mat. Sci. Tech., in the press, 1994.
ScriptaMetallurgicaet Materialia,Vol. 31, No. 11, pp. 1519-1524, 1994 Copyright© 1994 ElsevierScienceLtd Printed in the USA. All rights reserved 0956-716X/94 $6.00 + 00
THE [5--->coT R A N S F O R M A T I O N DURING ROOM TEMPERATURE AGING IN RAPIDLY SOLIDIFIED Ti-6A1-4V ALLOY Z. Fan Department of Materials Science and Engineering University of Surrey, Guildford, Surrey, GU2 5XH, UK ( R e c e i v e d May 12, 1994) ( R e v i s e d J u l y 8, 1994)
Introduction Since its first discovery by Frost et al [1] in aged Ti-Cr binary alloys, the co-phase has received extensive study, initially because of its deleterious effects on mechanical and physical properties, but later also due to the interests in the transformation mechanism. It is well known that the ~-phase can form by either quenching from high temperature (athermal co-phase) or aging at low temperature (isothermal co-phase) [2,3]. The ideal co-phase has a hexagonal structure, and a well defined orieuntation relationship with the parent bcc [5-phase, as first reported by Silcock [4]: {Ill]B//(0001)~, [ 110] I~//[ 1120] ~. Other structures of the co-phase have also been reported in the literature, for example, tfigonal co, ordered co, Zr2Al-type co and Ni2Al-type co (see review in Ref [5]). Recently there has been a renewed interest in co-phase transformation due to the development of aluminide-based intermetallic materials, where an ordered B2 phase can transform into an ordered co-phase during low temperature aging [5]. Experimental observation of the co-phase in commercial Ti-6AI-4V alloy has been very rare [6-8]. In fact, the occurrence of the co transformation in this alloy has been excluded by the argument that the concentrations of ct-stabilisers (A1 and O) in the [5-phase are sufficient to suppress the co-phase formation [6,7]. There appears to have been one report of the occurrence of the co-phase in quenched and aged Ti-6AI-4V alloy containing 0.116 wt % Fe and 0.15 wt % O [8]. In this paper we will report the experimental observation of the 13~co transformation during room temperature aging of a rapidly solidified Ti-6A14V alloy. Experiments The master Ti-6A1-4V alloy was supplied by DRA (UK), which contains 6.43 A1, 4.02 V and 0.19 Fe (in wt.%). Melting and melt-spinning were performed in a MARKO's Model 5T melt spinner in a stainless steel chamber with a high vacuum/inert gas (argon) atmosphere. The alloy was melted in a water cooled copper hearth using a non-consumable tungsten electrode under a high purity argon atmosphere. The chemically homogeneous melt was then delivered at a controlled rate to contact the circumferential surface of a Mo wheel rotating at 2500 rpm by which the melt is rapi~lly solidified as long fibres, which have a crescent-shaped cross-section, typically 100~300 I.tm in width and 40~100 I.tm in thickness. The melt spun fibres were comminuted into finer particles with particle size less than 200 gin. The comminuted alloy powders were then consolidated by hot-isostatic-pressing (HIPping) at 900°C and 300 MPa for 2 hours, and furnace cooled to room temperature. Further details for preparation of the materials under this study were given elsewhere [9]. Room temperature aging was carried out in air for about a year. Thin foils for TEM study were prepared in a "TENUPOL" unit using an electrolyte of 5 vol % perchloric acid in methanol. The polishing temperature was around -40°C, the voltage 30V, the current about 30 ~tA. TEM observation was performed on a JEOL 200CX STEM under an accelerating voltage of 200 kV. EDX analysis was performed on a JEOL JXA8600 EPMA system for phase composition. E x p e r i m e n t a l Results Microstructure of an as-HIPned Ti-6A1-4V Alloy The results from chemical analysis of the Ti-6A1-4V alloy after different processing steps are tabulated in Table 1. The oxygen level has increased after the melt-spinning process, especially after the comminution process, from 2050 ppm in the master alloy to 3100 ppm in the comminuted alloy powder. The oxygen content in the as-HIPped alloy is just slightly higher than 3100 ppm due to the good vacuum level in the HIPping process. 1519
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The microstructure of the consolidated Ti-6A1-4V alloy is show in Fig 1 by a low mag TEM micrograph. The a-phase has an equiaxed morphology with an average grain diameter about 6 gm, while the l~-phase has an irregular shape and is located at (x-grain boundaries and (x-grain triple points. The volume fraction of the [3-phase is about 0.08. The (x-grains are single crystals in which there is no sign of any transformation. The chemical compositions of the (x- and l-phases obtained by EDX analysis are presented in Table 2. These results indicate that c~-phase is rich in A1 and contains oractically no Fe, and that there is an enrichment of V and Fe in the 13-phase. In addition, a compositional variation in the p-phase has also been observed during the EDX analysis. The larger [3-grains located at et-grain triple points seem to contain less t3-stabiliser, while the smaller ~-grains located at co-grain boundaries contain more 13-stabilisers. TABLE. 2. EPMA results of phase composition in the TAB. 1. Results of impurity analysis of Ti-6AI-4V consolidated Ti-6A1-4V alloy (in wt %). alloy after different processing steps (in wt.%). Process Master allo), Melt-spun fibre Comminutedpowder
C 0.01 0.015 0.020
N2 0.0075 0.009 0.014
02 0.205 0.230 0.310
Phase (x ~
T1 Bal. Bal.
A1 6.53 3.24 . (3.08-3.89)
V 3.09 13.18 10.15-14.25
Fe 0.00 1.61 (1.24-2.01)
The detailed TEM study of the metastable ~-phase decomposition in as-HIPped Ti-6A1-4V alloy by the present authors has been presented elsewhere [10]. It was found that during the continuous cooling the metastable 13-phase can decompose by different modes, depending on the actual composition of the [~-phase. Less enrichment of V and Fe in the [3-phase will favour a direct formation of the equilibrium (x-phase from the 13-matrix, while the greaterenrichment of V and Fe can lead to the spinodal decomposition of the metastable [5-phase, resulting in the [3-stabiliser lean 15'-phase and the further enrichment of V and Fe in the I~-matrix. Upon further continuous cooling, the isothermal m-phase can form from the spinodally decomposed 13'-phase[10].
F I G 1. TEM BF image of the rapidly solidified Ti-6A14V alloy HIPoed at 900°C, 300 MPa for 2 hours, and furnace cooledto room temperature. The c~ and 13 phases are indicated by c~and 13in tile micrograph,
FIG 2. TEM BF image of a [~-grain showing ~ e typical features of the spinodal decomposition, p---)lJ+[3. The electron beam is close to [ 1]-0]13,and the dtrection for the compositional modulation is <] 11 >[~.
The detailed substructure of a 13-grain in an as-HIPped Ti-6AI-4V alloy is presented in Fig 2, which shows the typical features of a two-phase structure resulted from a spinodal decomposition reaction, 13---)[3+13',where 13 and [3' are coherent with each other [11]. In this composition modulated two-phase microstructure, the if-phase is lean of [3-stabilisers (V and Fe), while the l-matrix is further enriched by 13-stabilisers. The direction of the compositional modulation is < 111 >B [ 10]. SAD patterns from the spinodally decomposed [$-phase are shown in Fig 3. Besides the strong bcc reflections, diffuse scattering along a variety of directions is present in these SAD patterns, which is characteristic of the initial stage of the ~---)co transformation and is known as diffuse co scattering [12]. Occasionally, discrete spots corresponding to the co reflections can
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be observed in SAD patterns taken from such 13-grains. An example is given in Fig 4 where the electron beam is parallel to [ 110]~. The co reflections in Fig 4 are shifted systematically along the c axis of the co unit cell (parallel to. <11.1>~ directions). from the Bragg co positions, i.e., 1/3{222} 6 and 2/3{222}[~ for (0001)o~ and (0002)o in the case of onentatlon variant cob to positions of 0.36{222}~. and 0.64{222}~. This systematic shift of the co reflections has also been observed in other Tibase alloys, such as Ti-Fe, Ti-Mn and Ti-Cr, and has recently been discussed in detail by Sinkler and Luzzi [13].
F I G 3. SAD patterns taken from the spinodall2¢ decomposed [3-grains showing the diffuse co scattering which is characteristic of the early stage of the [3 to co transtormation. (a) [011113 zone axis pattern; (b) []-33][~zone axis pattern.
F I G 4. [110]13 zone axis pattern taken from a spinodall~, decomposed 13-grain showing the discrete co spots which are shifted from the Bragg co positions ( 1 / 3 { 2 2 2 } B and 2/3{222}[~) to positions of 0.36{222}[~ and 0.64{222}[3.
F I G 5. TEM BF image showing the appearance of a 13grain after aging at room temperature for a year. The fine cubic co parucles can be seen clearly in the [5-matrix.
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Microstructure after A~in~, After aging at room temperature for a year, the general ct+~l microstructure was not changed. However, the substructure of the p-phase has been altered. A TEM BF image of a [3-grain after room temperature aging is presented in Fig 5, where the free co precipitates can be seen clearly. Electron diffraction on such ~-grains indicated that B~o~ transformation has occurred at an advanced stage, as evidenced by the discrete spots from the co reflections in SAD patterns. Two examples of such SAD patterns are shown in Fig 6, with the electron beam being parallel to [110]~ (Fig 6a) and []13]B (Fig 6b) respectively. It is very interesting to note that after room temperature aging the systematic shifts of co spots in the as-HIPped state (Fig 4) have been substantially reduced, these co spots in Fig 6 appear at the positions very close to 1/3{222}[~ and 2/3[222 }[~, which correspond to the Bragg co reflections.
F I G 6. SAD patterns taken from I]-grains after aging at room temperature for a year showing the discrete spots corresponding to the Bragg co reflections, ta) [110115 zone axis pattern; (b) []'15115zone axis pattern.
F I G 7. TEM DF image using (000t)co reflection showing the detailed morphology of the to-phase after aging at room temperature for a year.
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A TEM dark field image using the (0001)o~ reflection is shown in Fig 7. The m-phase appears in the ~-matrix as fine particles with a diameter around 5 to 10 nm. In addition, the detailed TEM dark field observation revealed that the co particles tend to form parallel strings in the 13-matrix, as evidenced in Fig 8 by the comparison between the BF and corresponding DF TEM images of a 13-grain, where the electron beam is close to [110]~. This result indicates that the to-phase is transformed from the spinodally decomposed [~'-phase which is lean of [5-stat~itisers.
FIG 8. TEM BF image of a 13-grain after aging at room temperature for a year (a) and its corresponding DF image (b) showing the fine cubic co particles which tend to form parallel strings. In both (a) and (b), the electron beam is close to [ll0]~. Discussion The occurrence of the m-phase in as-HIPped Ti-6A1-4V alloy can be mainly attributed to the following three factors. Firstly, the sufficient enrichment of 13-stabilisers in the 13-phase makes the co-phase formation thermodynamically favourable in competing with the martensitic transformation [14]. EDX analysis showed that V and Fe concentration in the retained 13phase are more than 13.18 and 1.61 wt %, respectively (Table 2). The effectiveness of Fe as a 13-stabiliser is about 4 times as much as that of V on a weight percent basis. Therefore, the equivalent V concentration in the retained ~phase is over 20 wt %, which is sufficient to suppress the martensitic transformation [14-16]. This allows the m-transformation to replace the martensitic transformation during the continuous cooling. Secondly, the slow cooling rate from the HIPping temperature can create a favourable kinetic condition for the co-phase formation. Recently, Moffat and Larbalestier [17] studied the effect of cooling rate from the annealing temperature to room temperature on the metastable 13-phase decomposition mode in Ti-Nb binary alloys. They found that precipitation of the m-phase was favoured by a slow cooling rate, in contrast, the formation of cO' martensite is favoured by a higher cooling rate. The present results in the Ti-6A1-4V alloy support the above observation in Ti-Nb alloys. The cooling rate dependence of the mode of metastable 13-phase decomposition can be understood in terms of the role played by diffusion during the decomposition process. Martensitic transformation is diffusionless, therefore, diffusion has no influence on it, while the isothermal m-phase transformation involves nucleation and growth, a slower cooling rate thus favours the m-phase formation by allowing the extent of diffusion required by the co transformation process to be completed. Finally, the formation of a solute lean l~'-phase through the spinodal decomposition of the metastable 13phase favours the co-phase formation in terms of compositional depletion of 13-stabilisers and lattice plane shift through the coherency strain. It has been recently confirmed by atom probe analysis that isothermal co precipitates are depleted in all the alloying elements, both tx- and 13-stabilisers [18]. Therefore, the formation of a solute lean 13'-phase will reduce the amount of diffusion required by the formation of the m-phase compared with its direct formation from the parent 13-phase. In addition, the coherency strain between 13and 13' phases will make both bcc structures distorted towards a trigonal structure by shifting the { 111 }1~ planes. Such a trigonal lattice is an intermediate structure between the bcc 13 and the ideal co (hexagonal) structure. However, the continuous cooling may not create a thermodynamic and kinetic condition which is favourable enough to complete the 13'---~co transformation to form the crystalline co-phase, and hence the m-phase can only be observed by electron
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TRANSFORMATION1NTi-6AI-4V
Vol. 31, No. 11
diffraction in a form of diffuse co scattering. Aging at room temperature will further enhance the thermodynamic and kinetic conditions by allowing the diffusion required by the transformation. It is expected that at room temperature, the diffusion of V, A1 and Fe in bcc titanium is very slow. Therefore, a much longer time is necessary. Nevertheless, if the aging time at room temperature is long enough, the crystalline co-phase can be formed from the [~-stabiliser lean [~'-phase. The similar phenomenon has also been observed by the same author in a rapidly solidified Ti-7.5Mn-0.5B alloy after aging at room temperature [19]. Furthermore, the reduction of the systematic shifts of the co-reflections from the Bragg positions after aging at room temperature may indicate the existence of the followingtwo processes during the room temperature aging. Firstly, it suggests that the double plane collapsing required by thep--)co transformation in this alloy has not been completed after continuous cooling and can be completed gradually during the room temperature aging. Secondly, accompanying the double plane collapsing process, there is a depletion process of [3-stabilising elements (V and Fe) from the co-phase during the room temperature aging. As has been discussed by Sinkler and Luzzi [13], the magnitude of the co maxima shifts in SAD patterns depends on both the solute concentration and the identity of the solute. The amount of the co maxima shift increases with increasing solute concentration beyond the limit for co-phase formation. In addition, the amount of the co maxima shift also increases for a given solute concentration as the group number (in the Periodic Table) of the solute element increases. Therefore, the reduction of the systematic shifts of co-reflections suggests that a depletion process of V and Fe from the cophase is also present during the aging. Concluding Remarks Durin the continuous cooling of the Ti-6A1-4V alloy the metastable p-phase can decompose spinodally to form the ~stabil~er lean 13'-phase and the ~-stabiliser rich ~-phase. Upon further continuous cooling, the isothermal co-phase can form from the spinodally decomposed lY-phase. However, the thermodynamic and kinetic conditions created by the continuous coolin" may not be enough to complete the [~"--)co transformation to form the crystalline co-phase, and hence the co-phase can only b~ observed by electron diffraction in a form of diffuse co scattering. Aging at room temperature will further enhance me thermodynamic and kinetic conditions to such an extent that the crystalline co-phase can be formed from the 13-stabiliser lean [3'-phase, if the aging time is long enough. It is also possible that a depletion process of V and Fe from the co-phase occurs during the room temperature aging. Acknowledgements The author thanks Dr. L. Chandrasekaran for his help with orocessing the Ti-6A1-4V alloy, Professor A. P. Miodownik and Dr. P. Tsakiropoulos for their useful discussions. Ttils wor[¢ has been supported by the MOD, UK. References [1] P. D. Frost, W. M. Parris, L. L. Hirsch, J. R. Doig and C. M. Schwartz: Trans. ASM, 46, 231 (1954). [2] J. C. Williams: in Titanium Science and Technology, R. I. Jaffee and H. M. Burte eds, Plenum Press, New York, Vol 3, 1973, p1433, 1973. [3] S. K. Sikka, Y. K. Vohra and R. Chidambaram, Progress in Materials Science, 27, 145 (1982). [4] J. M. Silcock: Acta Met., 6, 481 (1985). [5] C. P. Chang and M. M. Loretto: Philos. Mag., 63A, 389 (1991). [6] J. C. Williams and M. J. Blackburn: Trans. ASM, 60, 373 (1967). [7] J. C. Williams, B. S. Hickman and D. H. Leslie: Met. Trans., 2, 477 (1971). [8] A. Lasalmonie and M. Loubradou: J. Mat. Sci., 14, 2589 (1979). [9] Z. Fan, A. P. Miodownik, L. Chandrasekaran and M. Ward-Close: J. Mat. Sci., 29, 1127 (1994). [10] Z. Fan and A. P. Miodownik: J. Mat. Sci., in the press, 1994. [11] J. W. Cahn: Trans. Met. Soc. AIME, 242, 166 (1968). [12] J. C. Williams, D. de Fontaine and N. E. Paton: Metall. Trans., 4, 2701 (1973). [13] W. Sinlder artd D. E. Luzzi: Acta Metall. Mater., 42, 1249 (1994). [14] N. A. Vanderpuye and A. P. Miodownik: in The science, technology and applications of titanium (Jaffee and Promisel eds.), Pergamon Press, p719, 1970. [15] P. Pietrokowsky and P. Duwez: Trans. AIME, 194, 627 (1952). [16] T. Sato, S. Hukai and Y. C. Huang: J. Aust. Inst. Met., 5, 149 (1960). [17] D. L. Moffat and D. C. Larbalestier: Met. Trans., 19A, 1677 (1988). [18] L. Hadjaj, A. Menand and C. Martin: quoted by N. E. Paton and H. L. Fraser: in Proceeding of Sixth Worm Conf. On Titanium, P. Lacombe, R. Tricot and G. Beranger eds., Societe Francaise de Metallurgie, France, Vol 3, p1469, 1988. [19] Z. Fan: Mat. Sci. Tech., in the press, 1994.