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High temperature precipitation hardening in a rapidly quenched AlTiNi alloy II. Precipitate characterisation
.•
MATERIALS SCIENCE & ENGINEERING
EL SEVIER
Materials Science and Engineering A221 (1996) 22-32
A
High temperature precipitation hardening in a rapidly quenched A1-Ti-Ni alloy II. Precipitate characterisation J.F. Nie*, B.C. Muddle Department of Materials Engineering, Monash University, Clayton, Victoria, 3168, Australia
Received 1 May 1996; revised 1 July 1996
Abstract Changes in the morphology and structure of second phase particles, formed during isothermal ageing of rapidly quenched At-6Ti-I.5Ni (wt.%) alloy at temperatures in the range 300-500°C, have been examined using a combination of conventional amplitude contrast imaging and electron microdiffraction. The precipitation-hardening response is initially attributable to a fine-scale distribution of coherent particles of metastable LI 2 phase and the orientation relationship between the metastable precipitate and c~-aluminium matrix is such that (001)LlJ/(001)~, [100]L~J/[100]~. During isothermal ageing, the metastable precipitates evolve with a transitional, three-dimensional cross-like morphology defined by three orthogonat sets of ellipsoidal elements forming in pairs. The precipitates comprise variants of a series of one-dimensional tetragonal superlattices that vary in structure from L12 to body-centred tetragonal D02s; the c axis of the tetragonal cells is parallel to the major axis of each ellipsoidal segment and the orientation relationship is such that the principal axes of the ordered product superlattices are parallel to those of the matrix phase. The structural transformation from metastable L12 to equilibrium D0_,z phase, via an intermediate D0z3 structure can be modelled assuming aperiodic shear displacements of 1/21110](001) within one-dimensional clusters of L12 unit cells, to create a range of one-dimensional tetragonal superlattices with differing c parameters. Keywords: Precipitation hardening; Amplitude contrast imaging; Electron microdiffraction;Tetragonal superlattice
1. Introduction In the preceding contribution [1], the decomposition of a rapidly quenched A I - 6 T i - 1 . 5 N i (wt.%) alloy, during isothermal ageing in the temperature range 300500°C, has been demonstrated to comprise initial dissolution of dispersed primary particles of metastable ternary phase [2], followed by uniform fine-scale solid state precipitation throughout a-aluminium matrix phase. This precipitation reaction gives rise to significant precipitation hardening, with maximum hardness values ( ~ 1 7 5 kg mm -2) comparable with those achievable in conventional high strength age-hardening aluminium alloys [3]. The temperature range of this ageing response is well in excess of that typical of conventional hardenable alloys (130-200°C) [3] and
* Corresponding author.
this, combined with the intrinsic thermal stability of the precipitate phase(s) and precipitate dispersions involved, suggests that if these microstructures could be reproduced in components of useful scale, then low density ternary alloys of appropriate Ti:Ni ratio might sustain superior mechanical properties at elevated temperatures (100-200°C). The precipitation reaction in the A 1 - T i - N i alloy has been demonstrated [1] to commence with homogeneous nucleation and growth of coherent fine particles of metastable ordered cubic L12 phase. However, from early stages of precipitation there is diffraction evidence of structural changes within these particles and, during the course of ageing, the particles develop an unusual three-dimensional cross-like morphology associated with a crystal structure intermediate between the Llz phase and the ordered tetragonal (D022) structure of equilibrium ~-Als(Ti,Ni). Precipitation culminates in the formation of a nano-scale ( < 100 m~a) dispersion of 0921-5093/96/$15.00 © 1996- - Elsevier Science S.A. All rights reserved PH S0921-5093(96)10468-8
J.F. Nie, B.C. Muddle /Materials Science and Engineering A221 (1996) 22-32
spheroidal particles of the equilibrium phase. It is the purpose of the present paper to report the results of a detailed study, using diffraction contrast imaging and electron microdiffraction, of the changes in morphology and structure during this precipitation sequence.
2. Experimental procedures Rapidly quenched ribbons of A1-6Ti-I.5Ni alloy, prepared by free-jet melt spinning [2], were cut into segments ~ 5 cm in length, sealed in vycor tube under a partial pressure of argon, and aged isothermally in a salt bath for up to 2400 h in the temperature range 300-500°C (_+ 2°C). Assessment of the age-hardening response of the alloy was described in detail in Part I [11. Samples for electron microscopy were punched mechanically from the ribbon and thinned to perforation by twin-jet electropolishing in a solution of 40% acetic acid, 30% orthophosphoric acid, 20% nitric acid and 10% water at 11 V, 0.2 A and ambient temperature. All specimens were examined in a Philips EM420 transmission electron microscope, operating at 120 kV.
3. Experimental results Decomposition of the rapidly quenched A1-Ti-Ni alloy involved initial dissolution of primary particles of a metastable ternary phase followed immediately by homogeneous nucleation and growth of coherent free particles of metastable ordered cubic L12 phase [1]. While the initial precipitates had the L12 structure, from very early stages of solid state precipitation there was diffraction evidence of continuous structural changes within the precipitate particles during ageing and these structural changes were accompanied by the development of an unusual three-dimensional cross-like morphology during particle growth. Evidence of the structural changes accompanying particle growth is reproduced in Fig. 1, which shows a sequence of <00t>~ selected area electron diffraction (SAED) patterns recorded from samples aged for varying times at 400°C. During initial ageing, the intensity of reflections from the metastable L12 phase increased significantly. However, after 5 h at 400°C, Fig. l(b), streaks of diffuse intensity were associated with the L12 reflections, extending parallel to <100>~. With extended ageing, subsidiary intensity maxima were observed to develop within these streaks and the separation of these maxima about the sites of the L12 reflections was observed to increase with ageing time, Fig. l(c,d). To account for the complex electron diffraction patterns of Fig. 1, the morphology and structure of the cross-like precipitates formed in samples aged for dif-
23
ferent periods at 400°C and 500°C were studied in greater detail. A series of two-beam images of these precipitates, with the electron beam (B) approximately parallel to <001>~, are provided in Fig. 2. In those micrographs obtained with diffracting vector g = + [200L, Fig. 2(a,b), images of the cross-like precipitate consist of two orthogonat sets of paired ellipsoidal lobes. That pair of ellipsoids with major axes extending perpendicular to g appear as black/black lobes separated by a line of no contrast perpendicular to g, while the second pair of ellipsoids appear continuous with a substructure of finely-spaced, parallel fringes perpendicular to g. When images of the same particles were viewed with g = ___[020], Fig. 2(c,d), the contrast pattern was reversed and the pair of ellipsoids which appeared as black/black lobes for g = + [200]~ were found to be continuous with a substructure of fine fringes, while the pair initially exhibiting fine internal fringes now exhibited black/black lobe contrast with the line of no contrast perpendicular to g = _+ [020L. In both cases, the internal fine fringes were only visible in that pair of ellipsoids with major axes extending parallel to g = <002>~, and in each case the fringes were arranged perpendicular to g. The apparent spacing of the fringes varied along the major axis of the ellipsoids. Fig. 3 compares two-beam images of a cross-like precipitate observed with B parallel to (a) <011 >~ and (b) <001>~, with g=[200]~ in each case. In <011>~ projection, Fig. 3(a), the lobes of strain contrast perpendicular to g are reduced in projected length, imply-
Fig. 1. Selected area electron diffraction patterns recorded parallel to <001)~ for specimens of A1-6Ti-I.5Ni alloy aged (a) 1 h, (b) 5 h, (c) 24 h, and (d) 240 h at 400°C. Note progressive change in the distribution of diffracted beams from the precipitate phase(s) with increase in ageing time.
24
J.F. Nie, B,C. Mz~ddle ~Materials Science and Engineerh~g A221 (1996) 2 2 - 3 2
200
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ing that they define precipitate protrusions that are inclined to the image plane. In addition, the circular core of darker contrast evident in Fig. 3(b) (and the images of Fig. 2) is not observed in (011)~ projection. This implies that the circular core observed in <001)~ images represents the projection of a set of paired ellipsoids with major axes extending parallel to the incident electron beam. Since the majority of precipitates exhibited a two-dimensional cross-like shape in any <001)~ projection, it is thus suggested that the precipitates have a three-dimensional, cross-like morphology comprising three orthogonal sets of ellipsoidal pairs. The alternative of a two-dimensional, cross-like morphology is not supported by the evidence because, if the precipitates were two-dimensional only, then only one third of them on average would be expected to exhibit a cross-like shape in a given (001)~ projection. Fig. 4 shows on-axis bright field and centred dark field (CDF) (001)~ images of a cross-like precipitate (arrowed) in a specimen aged 1 h at 500°C, and the corresponding (001)~ electron microdiffraction pattern recorded with the electron beam embracing the entire precipitate. With B approximately parallel to the zone
axis, both of the orthogonal precipitate variants normal to B are imaged simultaneously, with the substructure defined by fine-scale striations within both variants perhaps most distinctly revealed in the CDF image, Fig. 4(b). These striations or fringes were normal to and irregularly spaced along the major axis of each ellipsoidal segment of the particle, and their distribution varied in each variant and from particle to particle. Meaningful measurements of the fringe spacings proved difficult due to the fine scale of individual precipitates and the obvious variation in spacing. Electron microdiffraction patterns recorded from individual particles parallel to <001)~ varied from particle to particle, and it proved difficult to index these patterns consistently for a single, well-defined crystal structure. In Fig. 4(c), discrete L12 reflections are no longer evident and instead multiple precipitate reflections lie approximately in rows parallel to the <001)~ directions, distributed symmetrically about the positions of the type 1/2{200}~, 1/2{020}~ and t/2{220}~ and with a separation such that they commonly overlapped. When a range of cross-like precipitates was examined, it became evident that the separation of pairs of intensity
J.F. Nie, B.C. Muddle /Materials Science and Eng#~eering A22I (1996) 22-32
25
Fig. 3. Two-beam TEM images of cross-like precipitates in samples aged 1 h at 500°C. Electron beam is approximately paraIlel to (a) (011>;, and (b) (001>~.
maxima about positions of the type 1/2{200}6 and 1/2{020}~ varied from particle to particle. Fig. 5 compares (00I >~ patterns recorded from two separate precipitates and selected as examples because, in contrast to the pattern in Fig. 4(c), they contain only single pairs of diffuse intensity maxima about L12 reflections at 1/2{200}~, 1/2{020}~ and I/2{220}~. The distances separating the paired reflections arrowed in Fig. 5(a) and (b) are distinguishably different and correspond to fractions of 0.19 and 0.24 respectively of the spacing between the (200)~ and (220)6 reflections. The electron microdiffraction patterns recorded from opposing ellipsoidal precipitate segments were effectively identical, while those patterns recorded from orthogonal variants differed by a rotation of 90 ° about the direction of the electron beam. Patterns such as those shown in Figs. 4 and 5 could not be indexed for the metastable cubic L12 phase, but were found to be similar to those expected for a body-centred tetragonal 9023 structure (a = 0.3890, c-- 1.6822 nm, space group I4/mmm) [4]. However, given the variation observed in the precipitate patterns, rarely could an individual pattern be accurately indexed for the D023 structure. The D023 unit cell can be considered to be derived from the ordered cubic Llz structure by a periodic shear displacement of 1/2[110](001)L12 between every two L12 unit cells stacked normal to (001)L12, Fig. 6. The D023 structure is thus commonly described as a one-dimensional L12-derivative superlattice. Schematic [100], [010] and [001] zone axis patterns calculated [5] for the above D023 structure and a pattern superimposing these three orthogonal orientations of the 9023 phase are included in Fig. 7. Comparison of those electron microdiffraction patterns shown in Figs. 4 and
5 with the schematic patterns of Fig. 7 indicates that the experimental patterns recorded from individual particles are similar in fotrn to the composite pattern comprising three orthogonal variants of the D023 structure, although clearly not compatible in the detail of the spacing of diffraction maxima. The match is sufficient to prompt the proposal that the particles can be regarded as being composed of three variants of a tetragonal structure (or structures), which is derived from the L12 structure and is closely related to the tetragonal D0z3 phase. The relationship between the structure and the morphology observed for the cross-like precipitates was determined to be such that the three orthogonal sets of paired ellipsoids corresponded to three variants of the tetragonal phase and that the major axis of each ellipsoidal precipitate segment, corresponding to the c axis of the tetragonal structure, was parallel to (001 >~. The finely-spaced fi'inges within each ellipsoidal segment were arranged in such a way that they lay perpendicular to the c axis of the tetragonal structure (i.e. parallel to the basal plane). The orientation relationship between precipitate and matrix structures was such that the principal axes of the tetragonal phase were parallel to those of the matrix phase. To account more precisely for the diffraction patterns observed from the cross-like precipitates, it proved necessary to relax the conditions defining the crystal structure of the precipitates. The precipitate was assumed to have a tetragonal superlattice, which was derived from a one-dimensional cluster of L12 unit cells by a periodic shear displacement on the system 1/21t10](001). If the lattice parameter of the L12 unit cell, a, was assumed to be identical to that of the c~-A1
26
J.F. Nie, B.C. Muddle / Materials Science and Engineering A221 (1996) 22-32
b
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Fig. 4. (a) (001)~ bright field image of a cross-like precipitate (arrowed) in a specimen aged 1 h at 500°C, and (b) corresponding centred dark field image of the precipitate using precipitate diffracted beams circled in (c), which shows the corresponding electron microdiffraction pattern.
matrix phase and the interval between the (001)L~2 shear planes was defined as Ma, then the c parameter of the tetragonal superlattice was 2Ma, and M could take any integer value. If the orientation relationship between the tetragonal superlattice and the matrix were
Fig. 5. (a,b) Electron microdiffraction patterns recorded from two separate cross-like precipitates in ribbons aged t h at 500°C; electron beam is parallel to (001)~.
assumed to be such that the principal axes of the superlattice were parallel to those of the cz-A1 matrix, then the diffraction patterns shown in Fig. 8(a)-(c) represent the [100], [010] and [001] variants of the tetragonal superlattice respectively, while the pattern in Fig. 7(d) is a composite obtained by superimposing the variants of Fig. 8(a-c) [6]. In each pattern, the large circles represent diffraction from the c~-A1matrix phase, while the smaller circles represent diffracted beams from the precipitate. Coordinates associated with the small circles indicate positions of the superlattice reflections in the diffraction patterns and are not the Miller indices. In these generalised schematic diffraction patterns, a variation in M will result in a change in the positions of the precipitate reflections, which implies, in turn, that Lla-derivative, tetragonal superlattices with different c parameters will give rise to distinguishably different diffraction patterns. For example, when the 1/2 [110](001)L1o shear displacement takes place between each succesgive (001)Lt2 plane, M is unity and the c axis
27
J.F. Nie, B, C. Muddle / Materials Science and Engineering A221 (1996) 22-32
1/21110]((
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of the resultant superlattice phase is equivalent to 2a, and the distribution of precipitate reflections will be identical to that expected for a D022 superlattice. When a
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Fig. 7. Calculated (a) [100], (b) [010] and (c) [001] electron diffraction patterns for tetragonal D0~3 phase; the pattern in (d) is composed of superimposed patterns (a)-(c). The large circles represent diffraction pattern expected for [001] a-A1 phase, while small circles define diffracted beams from the D023 phase.
the 1/21110](001)i_h shear displacement takes place on every second (00!)u, plane, M = 2 and the c axis of the resu!tant superlatticd phase is equal to 4a. In this case, the distribution of precipitate reflections will be the same as that expected for a D023 superlattice and the patterns will be equivalent to those shown in Fig. 7. When M -- 3, the spacing between those pairs of precipitate reflections distributed symmetrically about the locations (1,0,0), (0,1,0) and (1,1,0) is reduced and, with further' increases in M, these pairs of reflections approach merger at these locations. When M is infinite, the diffraction patterns are identical to those expected for the L12 phase. Comparison of experimental electron microdiffraction patterns with those calculated on the basis of the above model indicated that the observed patterns could be acc6unted for by assuming that individual precipitate variants had structures comprising domains of one or more L12-derivative, tetragonal superlattices with varying c parameters. Those patterns from separate particles in Fig. 5(a,b) could be modelled effectively assuming structures that combined antiphase domains with M = 2 and M = 3 in appropriate proportions to give rige to fractional values of M of 2.63 and 2.08 respectively. For : other particles, microdiffraction patterns recorded from individual ellipsoidal segments could differ, even within a single precipitate. Fig. 9(a,b) provides an example of two distinguishably different elec-
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J.F. Nie, B.C. Muddle / Materials Science and Engineering A22l (1996) 22-32
28
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Fig. 8. Schematic electron diffraction patterns for a I-D long-period tetragonal superlattice with c = 2M; (a)-(c) represent [100], [010] and [001] zone axis patterns respectively, and (d) is composed of superimposed patterns (a)-(c). The large circles represent (00t)~ reflections, and small filled circles correspond to tetragonal supertattice reflections. The small open circles represent potential additional reflections for a tetragonat superlattice with M ~ 3. Coordinates associated with the small circles indicate co-ordinates of tetragonal superlattice reflections and not the Miller indices [6].
tron microdiffraction patterns recorded from two separate, perpendicular segments of a single cross-like precipitate. In this case, neither of these two patterns could be indexed assuming a single tetragonal superlattice,
Fig. 9. (a,b) Electron microdiffraction patterns recorded from two perpendicular ellipsoidal segments of a single cross-like precipitate in a specimen aged 1 h at 500°C; electron beam is parallel to (001)~.
even if it were assumed to comprise domains with varying rational values of M and thus exhibit an average value of c. However, these patterns could be matched if the precipitate reflections were regarded as comprising direct and double diffraction from two average tetragonal supertattices with different c parameters. The present evidence thus suggests that growth of the initially spheroidal L12 precipitates into a three-dimensional cross-like form is associated with the development of an effectively continuous series of metastable tetragonal superlattices, culminating in the formation of an ordered D023 structure, as a rational intermediate structure between the initial Lt2 phase and the equilibrium D0= phase. However, no evidence was found of a cross-like precipitate having the equilibrium D0= structure. No clear observations of the final transition to the D0a2 phase were obtained, but by the time the equilibrium phase was established, the microstructure comprised uniformly distributed three-dimensional clusters of spheroidal particles, Fig. 10(a), Selected area electron diffraction patterns from such microstructures, Fig. 10(b), could, as shown schematically in Fig. 10(c), be indexed convincingly assuming three orthogonal variants of the tetragonal D0= structure, combined with multiple diffraction involving a combination of precipitate and matrix planes. This would tend to imply that the transition from the metastable tetragonal intermediate structures to the final equilibrium structure takes place in-situ as part of a continuous transition from the initial L12 structure.
4. Discussion The L12 precipitates, which form first from supersaturated AI-6Ti-1.5Ni alloy during elevated temperature ageing, are coherent with the matrix and the strain contrast observed in two-beam amplitude contrast images (e.g., Fig. 6 of [1]) indicates that the strain fields associated with individual precipitates are spherically symmetric [7]. The principal axes of the precipitate structure are parallel to those of the matrix c~-A1 and, with continued ageing, growth of these particles involves the development of independent protrusions parallel to (100)~. Three-dimensional cross-like forms, similar to those observed in the aged AI-6Ti-I.5Ni alloy, have been reported previously [8-10], but not fully characterised. The present study demonstrates that these precipitates comprise three orthogonal sets of paired ellipsoids, and that each ellipsoidal segment remains coherent with the matrix phase. As shown in Fig. 2, the strain contrast observed within systematic two-beam images implies that the strain fields associated with individual lobes of the three-dimensional array remain approximately spherically symmetric.
J.F. Nie, B.C. Muddle /Materials Science and Eng#wering A221 (1996) 22-32
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Fig. 10. (a) Transmission electron micrograph recorded from regions containing particles of equilibrium D022 phase in a sample aged 24 h at 500°C, (b) corresponding (001)= SAED pattern, and (c) calculated (001)~ SAED pattern expected for direct and double diffraction from aluminium matrix phase containing three variants of the D0a2 structure. The large circles represent reflections expected for c~-Al phase, small filled circles represent primary D022 reflections, and the open circles are those arising from double diffraction involving a combination of e-Al and D022 diffracting planes.
29
The structure of these coherent precipitates is complex, as they comprise three orthogonal variants of a series of metastable, L12-derivative, tetragonal superlattices, which may be distinguished in terms of a variation in their effective c parameters. The relationship between the morphology and structure is such that the three orthogonal sets of paired ellipsoids correspond to the three variants of each tetragonal phase, with the major axis of each ellipsoidal segment parallel to the c axis of the tetragonal phase and @01)~. The cubic c~-A1 matrix phase has a crystal point group m3m of order 48, while the L12-derivative, tetragonal superlattices have a crystal point group 4/mmm of order 16. For the observed orientation relationship, the intersection point group [1t] defined by common symmetry elements is 4/mmm and of order 16, and symmetry thus requires that there be 48/16 or 3 crystallographic variants of the precipitate phase in a given matrix orientation, in good agreement with experimental observations. The structure(s) of these precipitates may be modelled successfully by assuming that they comprise onedimensional combinations of a series of tetragonal superlattices, each of which may be considered as being derived from basal plane shear displacements within one-dimensional clusters of L12 unit cells. Electron microdiffraction patterns from individual precipitate segments frequently imply the existence of more than one superlattice unit cell within a given segment, while, within a given particle, the combinations of unit cells required to account for observed diffraction patterns may vary from segment to segment. Invariably these tetragonal unit cells share a common c axis and differ significantly only in the period of the repeat unit along this axis (i.e., in lattice parameter c), defined by the parameter M, Fig. 6. The common c axis appears to become the preferred growth direction for a given variant. To demonstrate the effectiveness of the proposed structural model, it is to be noted that, if schematic diffraction patterns corresponding to the three orthogohal variants of the permitted tetragonal superlattices are superimposed, then the resultant pattern provides an accurate simulation of the conventional SAED patterns observed from regions of specimens containing all such phases. As an example, the schematic electron diffraction pattern, formed by superimposing patterns expected of superlattices with rational values of M = 4, 3, 2, and 1 is shown in Fig. l l(a). The precipitate superlattice reflections lie approximately in rows parallel to @01)~, distributed symmetrically about locations of the type {1"00}~, {010}~ and {110}~, and there is excellent agreement betweer~ the form of this pattern and that obtained experimentally from a sample aged 1 h at 500°C, Fig. 1 l(b). Minor variations in the positions of precipitate reflections about the locations {100}~, {010}~ and {110}~ can be attributed to small variations
30
J.F. Nie, B.C. Muddle / Materials Science and Engineering A221 (1996) 22-32
in the superlattice c parameter, arising fi'om the combination on a fine-scale of two or more superlattices to define an effective value of M that is irrational. The influence of the value of M in defining the position of the superlattice reflections is defined in the schematic Fig. 8, while examples of individual patterns which imply irrational values of M are provided in Fig. 9. Fractional values of M imply that the ordered crystal comprises antiphase domains of differing periodicities, combined in such a manner that the apparently irrational value of M is to be interpreted as an average domain size [12].
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Fig. 11. Comparison of (a) calculatedpattern formedby superimposing patterns expectedof superlattices with M ---4, 3, 2 and I, with (b) observed (001)~ SAED pattern recorded from sample aged 1 h at 500°C.
Those diffracted beams recorded in Fig. l l(b), but not reproduced in the schematic solution of Fig. 1 l(a), may be accounted for by multiple diffraction involving matrix and precipitate crystals. For clarity of presentation, potential contributions from multiple diffraction have been omitted from the schematic solution of Fig. l l(a). However, the pattern of reflections to be expected from double diffraction when the precipitate phase has the equilibrium D022 structure is shown schematically in Fig. 10. Given that the D0a2 structure may be regarded as a one-dimensional superlattice with M = 1, a similar pattern of double diffraction is observed for those long period structures with M > 1 and this pattern readily accounts for the doubling of those direct precipitate reflections distributed parallel to (001)~ about the locations {100}~, {010}~ and {110}~. Different transformation sequences have been proposed for solid state precipitates forming in binary A1-Ti alloys during extended isothermal heat treatment [8-10]. It has been reported [10], for example, that the metastable Llz phase transforms directly to the equilibrium tetragonal D022 phase with prolonged ageing. Alternatively, the precipitation sequence has been suggested to involve either a displacive series of transitions L12--+D023--->D022, or a series of diffusional transformations Llz--+/3" ~ D022, where /3" represents a Cubic transition phase that has not been well-defined [8]. The present experimental evidence suggests that during continued ageing, the metastable L12 phase transforms to the equilibrium D022 phase via the formation of a series of metastable transition supertattices. The structural change accompanying the continuous transition from L12 to D022 can be modelted successfully assuming aperiodic shear displacements of 1/ 21110](001) within one-dimensional clusters of L12 unit cells. Fig. 6(a) shows the atomic arrangement in an ordered cubic L12 unit cell with an A3B composition. To simplify description of the displacive transformation from the L l z to DOzz, only three of those Llz-derivative, tetragonal superlattices, with M = 3, M = 2 (D023) and M = 1 (D022) are presented in Fig. 6(b-d) respectively. As illustrated schematically in Fig. 6(b), a tetragonal superlattice with M = 3 can be derived from an Llz structure by a periodic shear displacement of 1/ 21t10](001) between every three Lla unit cell clusters. Once this superlattice phase is generated, a subsequent shear displacement of 1/21110](002), Fig. 6(c), can result in a D023 superlattice ( M = 2). Furthermore, an additional shear displacement of 1/21110](001) within the resultant D023 structure gives rise to a D02z supertattice ( M = 1). The products of the periodic shear displacements described above are those superlattices with integer values of M, i.e., those expected to comprise uniformlyspaced planar antiphase domain boundaries (APBs). However, it also appears possible that aperiodic shear
J.F. Nie, B.C. Muddle /llgaterials Science and Engineering A221 (t996) 22-32
displacements of 1/21110](001) may occur in one-dimensional stacks of L12 unit cells. As shown in Fig. 9, models based on such aperiodic shear displacements are required to account for those precipitate diffraction patterns which imply fractional values of M. The existence of aperiodic displacements is also supported by the nature of the substructure within the cross-like precipitates. Those striations normal to the c axis of the superlattices are consistent with planar defects on the (001) basal plane that might be interpreted either as stacking faults associated with the t/21110](001) displacive transitions or as APBs separating superlattices of differing c parameters. The irregular spacing of these striations supports the notion of aperiodic shear displacements. However, direct verification of the nature of the substructure of these precipitates would require high-resolution electron microscopy, and there remains a need for further work on this aspect. Although the structural changes within the precipitates can be modelled by simple shear displacements, it is likely that the process of structural transformation is influenced by the diffusion of solute atoms. This suggestion is supported by the observation that the structural transition from L12 to D022 was observed to proceed more rapidly in precipitates formed at those sites originally occupied by metastable primary intermetallic particles and thus locally solute rich. There is evidence [8] that the metastable L12 phase that forms in rapidly-solidified A I - T i alloys has a Ti content of ~ 20 at.% Ti and is thus sub-stoichiometric in Ti with respect to fully ordered A13Ti. This implies that the structural transformation from L12 to D022 may also involve a significant increase in solute content. It might thus be inferred that the transformation from LI_, to a superlattice with a large M occurs when the titanium (and nickel) concentration reaches a critical magnitude, and that the subsequent transformations from the superlatrice with a large M to a superlattice with a small M, and from a superlattice with a low M to the equilibrium D022 phase might be associated with increasing solute concentration. If such a deduction is assumed correct, then the faster structural transformation in precipitates formed at locations of decomposed particles may be interpreted as a result of higher local solute concentration at those locations. The critical solute concentration, required for the nucleation of L12 phase and for subsequent structural transformation from the L12 to D022, could be achieved more readily in these locations due to relatively short range diffusion of solute atoms. For precipitates forming within the matrix phase, there exists a relatively low local solute concentration, and thus it takes longer to achieve, by long range diffusion, the critical magnitude required for the nucleation of L12 phase and for subsequent structural transformation(s). Since it is difficult to determine the actual composition of the precipitates in the various stages of
31
ageing due to the fine-scale of precipitates and the small differences in composition expected, such a suggestion could not be verified experimentally. However, the variation of ordering of superlattices as a function of composition has been observed and reported in a number of alloys, including Cu3+_~Pd [13] and A13_xTi 1+.~ [14].
5. Conclusions (i) The precipitation-hardening response in rapidly solidified A I - 6 T i - I . 5 N i (wt.%) alloy is initially attributable to a uniform title-scale distribution of coherent particles of the metastable L12 phase. The orientation relationship between the metastable precipitate and ~-aluminium matrix is such that: (001)L12// (001)~, [1001L12//[100]~. (ii) During isothermal ageing in the temperature range 300-500°C, the metastable precipitates evolve into a transitional, three-dimensional cross-like morphology comprising three orthogonal sets of ellipsoidal pairs. These cross-like precipitates comprise three orthogonal variants of a series of one-dimensional tetragonal superlattices that vary in structure from L12 to body-centred tetragonal D023. The c axis of the tetragonal superlattices is parallel to the major axis of each ellipsoidal segment and the orientation relationship is such that the principal axes of the ordered precipitate superlattices are parallel to those of the matrix phase. (iii) The structural transformation from metastable L12 to equilibrium D022 phase, via an intermediate D023 structure can be modelled assuming aperiodic shear displacements of t/21110](001) within one-dimensional clusters of El 2 unit cells. The fine-scale striations within each ellipsoidal segment of the cross-like precipitates are parallel to the (001) basal planes of the tetragonal cells and are interpreted to be evidence of planar defects arising from a displacive structural transformation, i.e., stacking faults or antiphase boundaries between tetragonal domains with differing c parameters.
Acknowledgements This research was supported by the Australian Research Council. One of the authors (JFN) acknowledges gratefully the support of a Monash Graduate Scholarship.
References [1] J.F. Nie and B.C. Muddle,Mater. Sci. Eng., A22I (1996) 11-21. [2] J.F. Nie and B.C. Muddle, Mater. Sci. Eng., A215 (1996) 92.
32
J.F. )Vie, B.C. Muddle / Materials Science and Engineering A221 (t996) 22-32
[3] I.J. Polmear, Light Alloys, 3rd edn., Edward Arnold, 1995. [4] E.V. Kozlov, N.M. Kormin and N.M. Matveyeva, Izv. Acad. Naulc. Set'. Metall., 5 (1979) 150. [5] M.E. Schlienger, J.T. Stanley and H.L. Fraser, Diffract Microdev Software, Version 1.2. [6] M.J. Marcinkowski and L. Zwell, Acta MetaIL, ti (1963) 373, [7] M.F. Ashby and L.M. Brown, Phil. Mag., 8 (1963) 1083. [8] A. Majumdar, Ph.D. Thesis, Department of Materials Engineering, Monash University, Clayton, Australia, 1989. [9] T. Ohashi and R. Ichikawa, J. Jpn Inst. Light Metals, 27 (1977) 105.
[t0] S. Hori, H. Tai and Y. Narita, in S, Steeb and H. Warlimont (eds.), Rapidly QuenchedMetals, Elsevier Science Publishers B.V., Amsterdam, 1985, p. 911. [11] J.W. Cahn and G. Katonji, in H.L Aaronson, D.E. Laughlin, R.F. Sekerka and C.M. Wayman (eds.), Int. Conf. on Solid-Solid Phase Transformations, TMS-AIME, Warrendale, PA, 1982, p. 3. [12] K. Fujiwara, J. Phys. Soe. Jpn, I2 (1957) 7. [13] D. Broddin, G. Van Tendeloo, J. Van Landuyt, S. Amelinckx, R. Portier, M. Guymont and A. Loiseau, Phil Mag., 54A (1986) 395. [14] R. Miida, Jpn. J. Appt. Phys., 25 (1986) 1815.
MATERIALS SCIENCE & ENGINEERING
EL SEVIER
Materials Science and Engineering A221 (1996) 22-32
A
High temperature precipitation hardening in a rapidly quenched A1-Ti-Ni alloy II. Precipitate characterisation J.F. Nie*, B.C. Muddle Department of Materials Engineering, Monash University, Clayton, Victoria, 3168, Australia
Received 1 May 1996; revised 1 July 1996
Abstract Changes in the morphology and structure of second phase particles, formed during isothermal ageing of rapidly quenched At-6Ti-I.5Ni (wt.%) alloy at temperatures in the range 300-500°C, have been examined using a combination of conventional amplitude contrast imaging and electron microdiffraction. The precipitation-hardening response is initially attributable to a fine-scale distribution of coherent particles of metastable LI 2 phase and the orientation relationship between the metastable precipitate and c~-aluminium matrix is such that (001)LlJ/(001)~, [100]L~J/[100]~. During isothermal ageing, the metastable precipitates evolve with a transitional, three-dimensional cross-like morphology defined by three orthogonat sets of ellipsoidal elements forming in pairs. The precipitates comprise variants of a series of one-dimensional tetragonal superlattices that vary in structure from L12 to body-centred tetragonal D02s; the c axis of the tetragonal cells is parallel to the major axis of each ellipsoidal segment and the orientation relationship is such that the principal axes of the ordered product superlattices are parallel to those of the matrix phase. The structural transformation from metastable L12 to equilibrium D0_,z phase, via an intermediate D0z3 structure can be modelled assuming aperiodic shear displacements of 1/21110](001) within one-dimensional clusters of L12 unit cells, to create a range of one-dimensional tetragonal superlattices with differing c parameters. Keywords: Precipitation hardening; Amplitude contrast imaging; Electron microdiffraction;Tetragonal superlattice
1. Introduction In the preceding contribution [1], the decomposition of a rapidly quenched A I - 6 T i - 1 . 5 N i (wt.%) alloy, during isothermal ageing in the temperature range 300500°C, has been demonstrated to comprise initial dissolution of dispersed primary particles of metastable ternary phase [2], followed by uniform fine-scale solid state precipitation throughout a-aluminium matrix phase. This precipitation reaction gives rise to significant precipitation hardening, with maximum hardness values ( ~ 1 7 5 kg mm -2) comparable with those achievable in conventional high strength age-hardening aluminium alloys [3]. The temperature range of this ageing response is well in excess of that typical of conventional hardenable alloys (130-200°C) [3] and
* Corresponding author.
this, combined with the intrinsic thermal stability of the precipitate phase(s) and precipitate dispersions involved, suggests that if these microstructures could be reproduced in components of useful scale, then low density ternary alloys of appropriate Ti:Ni ratio might sustain superior mechanical properties at elevated temperatures (100-200°C). The precipitation reaction in the A 1 - T i - N i alloy has been demonstrated [1] to commence with homogeneous nucleation and growth of coherent fine particles of metastable ordered cubic L12 phase. However, from early stages of precipitation there is diffraction evidence of structural changes within these particles and, during the course of ageing, the particles develop an unusual three-dimensional cross-like morphology associated with a crystal structure intermediate between the Llz phase and the ordered tetragonal (D022) structure of equilibrium ~-Als(Ti,Ni). Precipitation culminates in the formation of a nano-scale ( < 100 m~a) dispersion of 0921-5093/96/$15.00 © 1996- - Elsevier Science S.A. All rights reserved PH S0921-5093(96)10468-8
J.F. Nie, B.C. Muddle /Materials Science and Engineering A221 (1996) 22-32
spheroidal particles of the equilibrium phase. It is the purpose of the present paper to report the results of a detailed study, using diffraction contrast imaging and electron microdiffraction, of the changes in morphology and structure during this precipitation sequence.
2. Experimental procedures Rapidly quenched ribbons of A1-6Ti-I.5Ni alloy, prepared by free-jet melt spinning [2], were cut into segments ~ 5 cm in length, sealed in vycor tube under a partial pressure of argon, and aged isothermally in a salt bath for up to 2400 h in the temperature range 300-500°C (_+ 2°C). Assessment of the age-hardening response of the alloy was described in detail in Part I [11. Samples for electron microscopy were punched mechanically from the ribbon and thinned to perforation by twin-jet electropolishing in a solution of 40% acetic acid, 30% orthophosphoric acid, 20% nitric acid and 10% water at 11 V, 0.2 A and ambient temperature. All specimens were examined in a Philips EM420 transmission electron microscope, operating at 120 kV.
3. Experimental results Decomposition of the rapidly quenched A1-Ti-Ni alloy involved initial dissolution of primary particles of a metastable ternary phase followed immediately by homogeneous nucleation and growth of coherent free particles of metastable ordered cubic L12 phase [1]. While the initial precipitates had the L12 structure, from very early stages of solid state precipitation there was diffraction evidence of continuous structural changes within the precipitate particles during ageing and these structural changes were accompanied by the development of an unusual three-dimensional cross-like morphology during particle growth. Evidence of the structural changes accompanying particle growth is reproduced in Fig. 1, which shows a sequence of <00t>~ selected area electron diffraction (SAED) patterns recorded from samples aged for varying times at 400°C. During initial ageing, the intensity of reflections from the metastable L12 phase increased significantly. However, after 5 h at 400°C, Fig. l(b), streaks of diffuse intensity were associated with the L12 reflections, extending parallel to <100>~. With extended ageing, subsidiary intensity maxima were observed to develop within these streaks and the separation of these maxima about the sites of the L12 reflections was observed to increase with ageing time, Fig. l(c,d). To account for the complex electron diffraction patterns of Fig. 1, the morphology and structure of the cross-like precipitates formed in samples aged for dif-
23
ferent periods at 400°C and 500°C were studied in greater detail. A series of two-beam images of these precipitates, with the electron beam (B) approximately parallel to <001>~, are provided in Fig. 2. In those micrographs obtained with diffracting vector g = + [200L, Fig. 2(a,b), images of the cross-like precipitate consist of two orthogonat sets of paired ellipsoidal lobes. That pair of ellipsoids with major axes extending perpendicular to g appear as black/black lobes separated by a line of no contrast perpendicular to g, while the second pair of ellipsoids appear continuous with a substructure of finely-spaced, parallel fringes perpendicular to g. When images of the same particles were viewed with g = ___[020], Fig. 2(c,d), the contrast pattern was reversed and the pair of ellipsoids which appeared as black/black lobes for g = + [200]~ were found to be continuous with a substructure of fine fringes, while the pair initially exhibiting fine internal fringes now exhibited black/black lobe contrast with the line of no contrast perpendicular to g = _+ [020L. In both cases, the internal fine fringes were only visible in that pair of ellipsoids with major axes extending parallel to g = <002>~, and in each case the fringes were arranged perpendicular to g. The apparent spacing of the fringes varied along the major axis of the ellipsoids. Fig. 3 compares two-beam images of a cross-like precipitate observed with B parallel to (a) <011 >~ and (b) <001>~, with g=[200]~ in each case. In <011>~ projection, Fig. 3(a), the lobes of strain contrast perpendicular to g are reduced in projected length, imply-
Fig. 1. Selected area electron diffraction patterns recorded parallel to <001)~ for specimens of A1-6Ti-I.5Ni alloy aged (a) 1 h, (b) 5 h, (c) 24 h, and (d) 240 h at 400°C. Note progressive change in the distribution of diffracted beams from the precipitate phase(s) with increase in ageing time.
24
J.F. Nie, B,C. Mz~ddle ~Materials Science and Engineerh~g A221 (1996) 2 2 - 3 2
200
"200
-n
,II
J
J
J
t 020
:i 5 0 . . , Fig. 2. Two-beam amplitude contrast images of fine-scale, coherent cross-like precipitates in a specimen of A1-6Ti-1.5Ni alloy aged 240 h at 400°C. Electron beam is approximately parallel to <001)=.
ing that they define precipitate protrusions that are inclined to the image plane. In addition, the circular core of darker contrast evident in Fig. 3(b) (and the images of Fig. 2) is not observed in (011)~ projection. This implies that the circular core observed in <001)~ images represents the projection of a set of paired ellipsoids with major axes extending parallel to the incident electron beam. Since the majority of precipitates exhibited a two-dimensional cross-like shape in any <001)~ projection, it is thus suggested that the precipitates have a three-dimensional, cross-like morphology comprising three orthogonal sets of ellipsoidal pairs. The alternative of a two-dimensional, cross-like morphology is not supported by the evidence because, if the precipitates were two-dimensional only, then only one third of them on average would be expected to exhibit a cross-like shape in a given (001)~ projection. Fig. 4 shows on-axis bright field and centred dark field (CDF) (001)~ images of a cross-like precipitate (arrowed) in a specimen aged 1 h at 500°C, and the corresponding (001)~ electron microdiffraction pattern recorded with the electron beam embracing the entire precipitate. With B approximately parallel to the zone
axis, both of the orthogonal precipitate variants normal to B are imaged simultaneously, with the substructure defined by fine-scale striations within both variants perhaps most distinctly revealed in the CDF image, Fig. 4(b). These striations or fringes were normal to and irregularly spaced along the major axis of each ellipsoidal segment of the particle, and their distribution varied in each variant and from particle to particle. Meaningful measurements of the fringe spacings proved difficult due to the fine scale of individual precipitates and the obvious variation in spacing. Electron microdiffraction patterns recorded from individual particles parallel to <001)~ varied from particle to particle, and it proved difficult to index these patterns consistently for a single, well-defined crystal structure. In Fig. 4(c), discrete L12 reflections are no longer evident and instead multiple precipitate reflections lie approximately in rows parallel to the <001)~ directions, distributed symmetrically about the positions of the type 1/2{200}~, 1/2{020}~ and t/2{220}~ and with a separation such that they commonly overlapped. When a range of cross-like precipitates was examined, it became evident that the separation of pairs of intensity
J.F. Nie, B.C. Muddle /Materials Science and Eng#~eering A22I (1996) 22-32
25
Fig. 3. Two-beam TEM images of cross-like precipitates in samples aged 1 h at 500°C. Electron beam is approximately paraIlel to (a) (011>;, and (b) (001>~.
maxima about positions of the type 1/2{200}6 and 1/2{020}~ varied from particle to particle. Fig. 5 compares (00I >~ patterns recorded from two separate precipitates and selected as examples because, in contrast to the pattern in Fig. 4(c), they contain only single pairs of diffuse intensity maxima about L12 reflections at 1/2{200}~, 1/2{020}~ and I/2{220}~. The distances separating the paired reflections arrowed in Fig. 5(a) and (b) are distinguishably different and correspond to fractions of 0.19 and 0.24 respectively of the spacing between the (200)~ and (220)6 reflections. The electron microdiffraction patterns recorded from opposing ellipsoidal precipitate segments were effectively identical, while those patterns recorded from orthogonal variants differed by a rotation of 90 ° about the direction of the electron beam. Patterns such as those shown in Figs. 4 and 5 could not be indexed for the metastable cubic L12 phase, but were found to be similar to those expected for a body-centred tetragonal 9023 structure (a = 0.3890, c-- 1.6822 nm, space group I4/mmm) [4]. However, given the variation observed in the precipitate patterns, rarely could an individual pattern be accurately indexed for the D023 structure. The D023 unit cell can be considered to be derived from the ordered cubic Llz structure by a periodic shear displacement of 1/2[110](001)L12 between every two L12 unit cells stacked normal to (001)L12, Fig. 6. The D023 structure is thus commonly described as a one-dimensional L12-derivative superlattice. Schematic [100], [010] and [001] zone axis patterns calculated [5] for the above D023 structure and a pattern superimposing these three orthogonal orientations of the 9023 phase are included in Fig. 7. Comparison of those electron microdiffraction patterns shown in Figs. 4 and
5 with the schematic patterns of Fig. 7 indicates that the experimental patterns recorded from individual particles are similar in fotrn to the composite pattern comprising three orthogonal variants of the D023 structure, although clearly not compatible in the detail of the spacing of diffraction maxima. The match is sufficient to prompt the proposal that the particles can be regarded as being composed of three variants of a tetragonal structure (or structures), which is derived from the L12 structure and is closely related to the tetragonal D0z3 phase. The relationship between the structure and the morphology observed for the cross-like precipitates was determined to be such that the three orthogonal sets of paired ellipsoids corresponded to three variants of the tetragonal phase and that the major axis of each ellipsoidal precipitate segment, corresponding to the c axis of the tetragonal structure, was parallel to (001 >~. The finely-spaced fi'inges within each ellipsoidal segment were arranged in such a way that they lay perpendicular to the c axis of the tetragonal structure (i.e. parallel to the basal plane). The orientation relationship between precipitate and matrix structures was such that the principal axes of the tetragonal phase were parallel to those of the matrix phase. To account more precisely for the diffraction patterns observed from the cross-like precipitates, it proved necessary to relax the conditions defining the crystal structure of the precipitates. The precipitate was assumed to have a tetragonal superlattice, which was derived from a one-dimensional cluster of L12 unit cells by a periodic shear displacement on the system 1/21t10](001). If the lattice parameter of the L12 unit cell, a, was assumed to be identical to that of the c~-A1
26
J.F. Nie, B.C. Muddle / Materials Science and Engineering A221 (1996) 22-32
b
.
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Fig. 4. (a) (001)~ bright field image of a cross-like precipitate (arrowed) in a specimen aged 1 h at 500°C, and (b) corresponding centred dark field image of the precipitate using precipitate diffracted beams circled in (c), which shows the corresponding electron microdiffraction pattern.
matrix phase and the interval between the (001)L~2 shear planes was defined as Ma, then the c parameter of the tetragonal superlattice was 2Ma, and M could take any integer value. If the orientation relationship between the tetragonal superlattice and the matrix were
Fig. 5. (a,b) Electron microdiffraction patterns recorded from two separate cross-like precipitates in ribbons aged t h at 500°C; electron beam is parallel to (001)~.
assumed to be such that the principal axes of the superlattice were parallel to those of the cz-A1 matrix, then the diffraction patterns shown in Fig. 8(a)-(c) represent the [100], [010] and [001] variants of the tetragonal superlattice respectively, while the pattern in Fig. 7(d) is a composite obtained by superimposing the variants of Fig. 8(a-c) [6]. In each pattern, the large circles represent diffraction from the c~-A1matrix phase, while the smaller circles represent diffracted beams from the precipitate. Coordinates associated with the small circles indicate positions of the superlattice reflections in the diffraction patterns and are not the Miller indices. In these generalised schematic diffraction patterns, a variation in M will result in a change in the positions of the precipitate reflections, which implies, in turn, that Lla-derivative, tetragonal superlattices with different c parameters will give rise to distinguishably different diffraction patterns. For example, when the 1/2 [110](001)L1o shear displacement takes place between each succesgive (001)Lt2 plane, M is unity and the c axis
27
J.F. Nie, B, C. Muddle / Materials Science and Engineering A221 (1996) 22-32
1/21110]((
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of the resultant superlattice phase is equivalent to 2a, and the distribution of precipitate reflections will be identical to that expected for a D022 superlattice. When a
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Fig. 7. Calculated (a) [100], (b) [010] and (c) [001] electron diffraction patterns for tetragonal D0~3 phase; the pattern in (d) is composed of superimposed patterns (a)-(c). The large circles represent diffraction pattern expected for [001] a-A1 phase, while small circles define diffracted beams from the D023 phase.
the 1/21110](001)i_h shear displacement takes place on every second (00!)u, plane, M = 2 and the c axis of the resu!tant superlatticd phase is equal to 4a. In this case, the distribution of precipitate reflections will be the same as that expected for a D023 superlattice and the patterns will be equivalent to those shown in Fig. 7. When M -- 3, the spacing between those pairs of precipitate reflections distributed symmetrically about the locations (1,0,0), (0,1,0) and (1,1,0) is reduced and, with further' increases in M, these pairs of reflections approach merger at these locations. When M is infinite, the diffraction patterns are identical to those expected for the L12 phase. Comparison of experimental electron microdiffraction patterns with those calculated on the basis of the above model indicated that the observed patterns could be acc6unted for by assuming that individual precipitate variants had structures comprising domains of one or more L12-derivative, tetragonal superlattices with varying c parameters. Those patterns from separate particles in Fig. 5(a,b) could be modelled effectively assuming structures that combined antiphase domains with M = 2 and M = 3 in appropriate proportions to give rige to fractional values of M of 2.63 and 2.08 respectively. For : other particles, microdiffraction patterns recorded from individual ellipsoidal segments could differ, even within a single precipitate. Fig. 9(a,b) provides an example of two distinguishably different elec-
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J.F. Nie, B.C. Muddle / Materials Science and Engineering A22l (1996) 22-32
28
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Fig. 8. Schematic electron diffraction patterns for a I-D long-period tetragonal superlattice with c = 2M; (a)-(c) represent [100], [010] and [001] zone axis patterns respectively, and (d) is composed of superimposed patterns (a)-(c). The large circles represent (00t)~ reflections, and small filled circles correspond to tetragonal supertattice reflections. The small open circles represent potential additional reflections for a tetragonat superlattice with M ~ 3. Coordinates associated with the small circles indicate co-ordinates of tetragonal superlattice reflections and not the Miller indices [6].
tron microdiffraction patterns recorded from two separate, perpendicular segments of a single cross-like precipitate. In this case, neither of these two patterns could be indexed assuming a single tetragonal superlattice,
Fig. 9. (a,b) Electron microdiffraction patterns recorded from two perpendicular ellipsoidal segments of a single cross-like precipitate in a specimen aged 1 h at 500°C; electron beam is parallel to (001)~.
even if it were assumed to comprise domains with varying rational values of M and thus exhibit an average value of c. However, these patterns could be matched if the precipitate reflections were regarded as comprising direct and double diffraction from two average tetragonal supertattices with different c parameters. The present evidence thus suggests that growth of the initially spheroidal L12 precipitates into a three-dimensional cross-like form is associated with the development of an effectively continuous series of metastable tetragonal superlattices, culminating in the formation of an ordered D023 structure, as a rational intermediate structure between the initial Lt2 phase and the equilibrium D0= phase. However, no evidence was found of a cross-like precipitate having the equilibrium D0= structure. No clear observations of the final transition to the D0a2 phase were obtained, but by the time the equilibrium phase was established, the microstructure comprised uniformly distributed three-dimensional clusters of spheroidal particles, Fig. 10(a), Selected area electron diffraction patterns from such microstructures, Fig. 10(b), could, as shown schematically in Fig. 10(c), be indexed convincingly assuming three orthogonal variants of the tetragonal D0= structure, combined with multiple diffraction involving a combination of precipitate and matrix planes. This would tend to imply that the transition from the metastable tetragonal intermediate structures to the final equilibrium structure takes place in-situ as part of a continuous transition from the initial L12 structure.
4. Discussion The L12 precipitates, which form first from supersaturated AI-6Ti-1.5Ni alloy during elevated temperature ageing, are coherent with the matrix and the strain contrast observed in two-beam amplitude contrast images (e.g., Fig. 6 of [1]) indicates that the strain fields associated with individual precipitates are spherically symmetric [7]. The principal axes of the precipitate structure are parallel to those of the matrix c~-A1 and, with continued ageing, growth of these particles involves the development of independent protrusions parallel to (100)~. Three-dimensional cross-like forms, similar to those observed in the aged AI-6Ti-I.5Ni alloy, have been reported previously [8-10], but not fully characterised. The present study demonstrates that these precipitates comprise three orthogonal sets of paired ellipsoids, and that each ellipsoidal segment remains coherent with the matrix phase. As shown in Fig. 2, the strain contrast observed within systematic two-beam images implies that the strain fields associated with individual lobes of the three-dimensional array remain approximately spherically symmetric.
J.F. Nie, B.C. Muddle /Materials Science and Eng#wering A221 (1996) 22-32
O0
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Fig. 10. (a) Transmission electron micrograph recorded from regions containing particles of equilibrium D022 phase in a sample aged 24 h at 500°C, (b) corresponding (001)= SAED pattern, and (c) calculated (001)~ SAED pattern expected for direct and double diffraction from aluminium matrix phase containing three variants of the D0a2 structure. The large circles represent reflections expected for c~-Al phase, small filled circles represent primary D022 reflections, and the open circles are those arising from double diffraction involving a combination of e-Al and D022 diffracting planes.
29
The structure of these coherent precipitates is complex, as they comprise three orthogonal variants of a series of metastable, L12-derivative, tetragonal superlattices, which may be distinguished in terms of a variation in their effective c parameters. The relationship between the morphology and structure is such that the three orthogonal sets of paired ellipsoids correspond to the three variants of each tetragonal phase, with the major axis of each ellipsoidal segment parallel to the c axis of the tetragonal phase and @01)~. The cubic c~-A1 matrix phase has a crystal point group m3m of order 48, while the L12-derivative, tetragonal superlattices have a crystal point group 4/mmm of order 16. For the observed orientation relationship, the intersection point group [1t] defined by common symmetry elements is 4/mmm and of order 16, and symmetry thus requires that there be 48/16 or 3 crystallographic variants of the precipitate phase in a given matrix orientation, in good agreement with experimental observations. The structure(s) of these precipitates may be modelled successfully by assuming that they comprise onedimensional combinations of a series of tetragonal superlattices, each of which may be considered as being derived from basal plane shear displacements within one-dimensional clusters of L12 unit cells. Electron microdiffraction patterns from individual precipitate segments frequently imply the existence of more than one superlattice unit cell within a given segment, while, within a given particle, the combinations of unit cells required to account for observed diffraction patterns may vary from segment to segment. Invariably these tetragonal unit cells share a common c axis and differ significantly only in the period of the repeat unit along this axis (i.e., in lattice parameter c), defined by the parameter M, Fig. 6. The common c axis appears to become the preferred growth direction for a given variant. To demonstrate the effectiveness of the proposed structural model, it is to be noted that, if schematic diffraction patterns corresponding to the three orthogohal variants of the permitted tetragonal superlattices are superimposed, then the resultant pattern provides an accurate simulation of the conventional SAED patterns observed from regions of specimens containing all such phases. As an example, the schematic electron diffraction pattern, formed by superimposing patterns expected of superlattices with rational values of M = 4, 3, 2, and 1 is shown in Fig. l l(a). The precipitate superlattice reflections lie approximately in rows parallel to @01)~, distributed symmetrically about locations of the type {1"00}~, {010}~ and {110}~, and there is excellent agreement betweer~ the form of this pattern and that obtained experimentally from a sample aged 1 h at 500°C, Fig. 1 l(b). Minor variations in the positions of precipitate reflections about the locations {100}~, {010}~ and {110}~ can be attributed to small variations
30
J.F. Nie, B.C. Muddle / Materials Science and Engineering A221 (1996) 22-32
in the superlattice c parameter, arising fi'om the combination on a fine-scale of two or more superlattices to define an effective value of M that is irrational. The influence of the value of M in defining the position of the superlattice reflections is defined in the schematic Fig. 8, while examples of individual patterns which imply irrational values of M are provided in Fig. 9. Fractional values of M imply that the ordered crystal comprises antiphase domains of differing periodicities, combined in such a manner that the apparently irrational value of M is to be interpreted as an average domain size [12].
f
0 Z ..........
"
.
@
' °
t ..........
O
020
." a
20o
a
Fig. 11. Comparison of (a) calculatedpattern formedby superimposing patterns expectedof superlattices with M ---4, 3, 2 and I, with (b) observed (001)~ SAED pattern recorded from sample aged 1 h at 500°C.
Those diffracted beams recorded in Fig. l l(b), but not reproduced in the schematic solution of Fig. 1 l(a), may be accounted for by multiple diffraction involving matrix and precipitate crystals. For clarity of presentation, potential contributions from multiple diffraction have been omitted from the schematic solution of Fig. l l(a). However, the pattern of reflections to be expected from double diffraction when the precipitate phase has the equilibrium D022 structure is shown schematically in Fig. 10. Given that the D0a2 structure may be regarded as a one-dimensional superlattice with M = 1, a similar pattern of double diffraction is observed for those long period structures with M > 1 and this pattern readily accounts for the doubling of those direct precipitate reflections distributed parallel to (001)~ about the locations {100}~, {010}~ and {110}~. Different transformation sequences have been proposed for solid state precipitates forming in binary A1-Ti alloys during extended isothermal heat treatment [8-10]. It has been reported [10], for example, that the metastable Llz phase transforms directly to the equilibrium tetragonal D022 phase with prolonged ageing. Alternatively, the precipitation sequence has been suggested to involve either a displacive series of transitions L12--+D023--->D022, or a series of diffusional transformations Llz--+/3" ~ D022, where /3" represents a Cubic transition phase that has not been well-defined [8]. The present experimental evidence suggests that during continued ageing, the metastable L12 phase transforms to the equilibrium D022 phase via the formation of a series of metastable transition supertattices. The structural change accompanying the continuous transition from L12 to D022 can be modelted successfully assuming aperiodic shear displacements of 1/ 21110](001) within one-dimensional clusters of L12 unit cells. Fig. 6(a) shows the atomic arrangement in an ordered cubic L12 unit cell with an A3B composition. To simplify description of the displacive transformation from the L l z to DOzz, only three of those Llz-derivative, tetragonal superlattices, with M = 3, M = 2 (D023) and M = 1 (D022) are presented in Fig. 6(b-d) respectively. As illustrated schematically in Fig. 6(b), a tetragonal superlattice with M = 3 can be derived from an Llz structure by a periodic shear displacement of 1/ 21t10](001) between every three Lla unit cell clusters. Once this superlattice phase is generated, a subsequent shear displacement of 1/21110](002), Fig. 6(c), can result in a D023 superlattice ( M = 2). Furthermore, an additional shear displacement of 1/21110](001) within the resultant D023 structure gives rise to a D02z supertattice ( M = 1). The products of the periodic shear displacements described above are those superlattices with integer values of M, i.e., those expected to comprise uniformlyspaced planar antiphase domain boundaries (APBs). However, it also appears possible that aperiodic shear
J.F. Nie, B.C. Muddle /llgaterials Science and Engineering A221 (t996) 22-32
displacements of 1/21110](001) may occur in one-dimensional stacks of L12 unit cells. As shown in Fig. 9, models based on such aperiodic shear displacements are required to account for those precipitate diffraction patterns which imply fractional values of M. The existence of aperiodic displacements is also supported by the nature of the substructure within the cross-like precipitates. Those striations normal to the c axis of the superlattices are consistent with planar defects on the (001) basal plane that might be interpreted either as stacking faults associated with the t/21110](001) displacive transitions or as APBs separating superlattices of differing c parameters. The irregular spacing of these striations supports the notion of aperiodic shear displacements. However, direct verification of the nature of the substructure of these precipitates would require high-resolution electron microscopy, and there remains a need for further work on this aspect. Although the structural changes within the precipitates can be modelled by simple shear displacements, it is likely that the process of structural transformation is influenced by the diffusion of solute atoms. This suggestion is supported by the observation that the structural transition from L12 to D022 was observed to proceed more rapidly in precipitates formed at those sites originally occupied by metastable primary intermetallic particles and thus locally solute rich. There is evidence [8] that the metastable L12 phase that forms in rapidly-solidified A I - T i alloys has a Ti content of ~ 20 at.% Ti and is thus sub-stoichiometric in Ti with respect to fully ordered A13Ti. This implies that the structural transformation from L12 to D022 may also involve a significant increase in solute content. It might thus be inferred that the transformation from LI_, to a superlattice with a large M occurs when the titanium (and nickel) concentration reaches a critical magnitude, and that the subsequent transformations from the superlatrice with a large M to a superlattice with a small M, and from a superlattice with a low M to the equilibrium D022 phase might be associated with increasing solute concentration. If such a deduction is assumed correct, then the faster structural transformation in precipitates formed at locations of decomposed particles may be interpreted as a result of higher local solute concentration at those locations. The critical solute concentration, required for the nucleation of L12 phase and for subsequent structural transformation from the L12 to D022, could be achieved more readily in these locations due to relatively short range diffusion of solute atoms. For precipitates forming within the matrix phase, there exists a relatively low local solute concentration, and thus it takes longer to achieve, by long range diffusion, the critical magnitude required for the nucleation of L12 phase and for subsequent structural transformation(s). Since it is difficult to determine the actual composition of the precipitates in the various stages of
31
ageing due to the fine-scale of precipitates and the small differences in composition expected, such a suggestion could not be verified experimentally. However, the variation of ordering of superlattices as a function of composition has been observed and reported in a number of alloys, including Cu3+_~Pd [13] and A13_xTi 1+.~ [14].
5. Conclusions (i) The precipitation-hardening response in rapidly solidified A I - 6 T i - I . 5 N i (wt.%) alloy is initially attributable to a uniform title-scale distribution of coherent particles of the metastable L12 phase. The orientation relationship between the metastable precipitate and ~-aluminium matrix is such that: (001)L12// (001)~, [1001L12//[100]~. (ii) During isothermal ageing in the temperature range 300-500°C, the metastable precipitates evolve into a transitional, three-dimensional cross-like morphology comprising three orthogonal sets of ellipsoidal pairs. These cross-like precipitates comprise three orthogonal variants of a series of one-dimensional tetragonal superlattices that vary in structure from L12 to body-centred tetragonal D023. The c axis of the tetragonal superlattices is parallel to the major axis of each ellipsoidal segment and the orientation relationship is such that the principal axes of the ordered precipitate superlattices are parallel to those of the matrix phase. (iii) The structural transformation from metastable L12 to equilibrium D022 phase, via an intermediate D023 structure can be modelled assuming aperiodic shear displacements of t/21110](001) within one-dimensional clusters of El 2 unit cells. The fine-scale striations within each ellipsoidal segment of the cross-like precipitates are parallel to the (001) basal planes of the tetragonal cells and are interpreted to be evidence of planar defects arising from a displacive structural transformation, i.e., stacking faults or antiphase boundaries between tetragonal domains with differing c parameters.
Acknowledgements This research was supported by the Australian Research Council. One of the authors (JFN) acknowledges gratefully the support of a Monash Graduate Scholarship.
References [1] J.F. Nie and B.C. Muddle,Mater. Sci. Eng., A22I (1996) 11-21. [2] J.F. Nie and B.C. Muddle, Mater. Sci. Eng., A215 (1996) 92.
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[3] I.J. Polmear, Light Alloys, 3rd edn., Edward Arnold, 1995. [4] E.V. Kozlov, N.M. Kormin and N.M. Matveyeva, Izv. Acad. Naulc. Set'. Metall., 5 (1979) 150. [5] M.E. Schlienger, J.T. Stanley and H.L. Fraser, Diffract Microdev Software, Version 1.2. [6] M.J. Marcinkowski and L. Zwell, Acta MetaIL, ti (1963) 373, [7] M.F. Ashby and L.M. Brown, Phil. Mag., 8 (1963) 1083. [8] A. Majumdar, Ph.D. Thesis, Department of Materials Engineering, Monash University, Clayton, Australia, 1989. [9] T. Ohashi and R. Ichikawa, J. Jpn Inst. Light Metals, 27 (1977) 105.
[t0] S. Hori, H. Tai and Y. Narita, in S, Steeb and H. Warlimont (eds.), Rapidly QuenchedMetals, Elsevier Science Publishers B.V., Amsterdam, 1985, p. 911. [11] J.W. Cahn and G. Katonji, in H.L Aaronson, D.E. Laughlin, R.F. Sekerka and C.M. Wayman (eds.), Int. Conf. on Solid-Solid Phase Transformations, TMS-AIME, Warrendale, PA, 1982, p. 3. [12] K. Fujiwara, J. Phys. Soe. Jpn, I2 (1957) 7. [13] D. Broddin, G. Van Tendeloo, J. Van Landuyt, S. Amelinckx, R. Portier, M. Guymont and A. Loiseau, Phil Mag., 54A (1986) 395. [14] R. Miida, Jpn. J. Appt. Phys., 25 (1986) 1815.