High temperature precipitation hardening in a rapidly quenched AlTiNi alloy I. Precipitation hardening response

MATERIALS SCIENCE & ENGINEERING

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ELSEVIER

Materials Science and Engineering A221 (1996) 11-21

High temperature precipitation hardening in a rapidly quenched AI-Ti-Ni alloy I. Precipitation hardening response J.F. Nie*, B.C. Muddle Department of JVlaterials Engineering, Monash University, Clayton, Victoria 3168, Australia Received 1 May I996; revised 1 July 1996

Abstract A significant precipitation-hardening response has been observed in rapidly quenched A1-6Ti-I.5Ni (wt.%) alloy aged isothermally in the temperature range 300-500°C, and the underlying precipitate microstructures characterised using transmission electron microscopy (TEM). Primary intermetallic dispersoids of cubic ternary phase in as-quenched alloy decompose rapidly during heat treatment and are replaced by uniform precipitation of fine-scale, coherent particles of a metastable L12 phase. These metastable precipitates evolve into a transitional, three-dimensional cross-like morphology and eventually into nano-scale ( < 100 nm) spheroidal particles of equilibrium D0ze phase g-A13(Ti,Ni). The changes in form are accompanied by the development of a series of one-dimensional long period superlattices, culminating in formation of equilibrium b.c.t, phase. Maximum hardness (175 kg m m - 2), which is associated with a dispersion of coherent intermediate precipitates and a minor fraction of d-A13(Ti,Ni), is comparable with that of conventional high strength precipitation-hardening alloys (150-200 VHN). The temperatures of this ageing response, together with the thermal stability of the precipitate phase(s), suggest that low density, rapidly quenched AI-Ti-Ni alloys, with weight ratio Ti:Ni in the range 3:1-4:1, may have potential for applications involving elevated temperatures (150-200°C), where the creep resistance of conventional precipitation-hardened alloys declines rapidly. Keywords: Precipitation-hardened alloys; Temperature; Transmission electron microscopy

1. Introduction

A common approach to the design of high strength, creep resistant aluminium alloys involves production of alloy microstructures comprising a large volume fraction of thermally stable, fine-scale intermetallic particles distributed uniformly in an aluminium matrix [1]. Since conventional dispersion-strengthened aluminium alloys typically overage rapidly at temperatures in excess of 100°C [2], with a corresponding decline in strength, most recent attempts at alloy development have focused on novel alloy systems and non-equilibrium processing techniques such as rapid solidification processing. At the levels of undercooling achieved during rapid solidification, the solubility of alloying elements may potentially be increased and a large volume fraction of fine-scale dispersoids may be generated, either directly * Corresponding author. 0921-5093/96/$15.00 © 1 9 9 6 - Elsevier Science S.A. All rights reserved PII S0921-5093(96)10467-6

from the melt duri_ng rapid quenching~ or from super, saturated cz-A1 solid solution by suitab!e post-solidific~, tion heat treatment [1]. The most successful group of alloys developed using this approach has been that based on the A1-Fe system, with ternary and often quaternary additions [3]. As an alternative, the A1-Ti system is on e of the more promising [4,5], for the titanium is low density aiid has tow solid solubility and diffusivity in aluminium [6,7]. Under conditions of rapid solidification, formation of the equilibrium intermetallic phase A13Ti (b.c.t., D022) is generally suppressed and replaced by a metastable ordered cubic (L12) phase, which is sub-stoichiometric with respect to titanium (,-~ A14Ti) [8] and which has a potentially low lattice misfit with ~-AI matrix phase. However, in binary A1-Ti alloys, the metastable intermetallic phase forms directly from the melt, and is thus relatively coarse in scale (0.1-0.3 gm). In addition, the volume fraction of fine-scale, solid state

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intermetallic precipitates is low in quenched and aged microstructures and the distribution is inhomogeneous [8-11]. These difficulties have diverted recent attention towards ternary alloys based on the AI-Ti system, and the feasibility of developing potentially useful microstructures has been demonstrated [12-16] in a range of ternary alloys, in which the ternary addition is typically a transition or rare earth element. Of the rapidly solidified A1-Ti based ternary alloys examined thus far, A1-Ti-Ni alloys, with a Ti:Ni weight ratio in the range 3:1-4:1, have perhaps the most interesting as-quenched microstructures, for they not only exhibit a uniform, nano-scale ( < 100 nm) distribution of primary metastable ternary intermetallic phase, but they also retain up to ~ 3.5 wt.% Ti in ~-A1 solid solution [15]. There is thus the potential that decomposition of this highly supersaturated solid solution may generate a high density of fine-scale, solid state intermetallic precipitates after suitable ageing treatments, to complement the primary intermetaltic phase. In the present work, a rapidly solidified AI6Ti-l.5Ni (wt.%) alloy typical of this group has been selected for detailed isothermal ageing studies, and it is the purpose of the present paper to report observation of a significant precipitation-hardening response in this alloy at ageing temperatures (300-500°C) well in excess of those associated with conventional precipitationhardened alloys. The precipitation sequence underlying this hardening response is also characterised, while details of the structural changes occurring within the strengthening precipitates during ageing are presented in a companion paper [17].

2. Experimental procedures Rapidly quenched ribbons of AI-6Ti-I.5Ni alloy, prepared by free-jet melt spinning [15], 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). To assess the age-hardening response of the alloy, the hardness of the ribbons (30-50 gm thickness) was measured as a function of thermal treatment using a nano-indentation instrument (UMIS 2000, C.S.I.R.O. Division of Materials Science and Technology, Australia). This instrument uses a triangular-based diamond pyramid indentor with a face angle of 65.3 °, the load ranges from 1 to 200 raN, and the maximum penetration depth is 2 gm. Measurement involves bringing the indentor to the surface of the sample with a small contact force (0.1 raN) and then monitoring continuously the force and displacement associated with indentation. Hardness is determined as a function of penetration depth [18].

To assess the reliability of the technique, hardness measurements were made on both melt-spun ribbons and bulk samples of annealed, pure aluminium and on bulk samples of a peak-aged, high strength precipitation-hardening aluminium alloy (Weldalite 049TM). Micro-indentation measurements for pure A1 produced average hardness values of 28 and 32 kg m m - 2 (equivalent Vickers hardness, EVH) for thin ribbon and bulk samples respectively, and 214 kg mm -2 for the bulk Weldalite 049 TM sample. These values are to be compared with conventional bulk Vickers hardness numbers of 17 VHN (2.5 kg load) and 197 VHN (5 kg load) for bulk samples of pure A1 and Weldalite 049TM respectively. These preliminary data suggest that hardness values defined by the UMIS 2000 may be systematically ,-~ 15 VHN higher than those determined using standard Vickers hardness testing. Samples for etectron microscopy were punched mechanically from the ribbon and thinned to perforation by twin-jet electropolishing in a solution of 40% acetic acid, 30°./0 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 (TEM), operating at 120 kV.

3. Experimental results 3. I. A g e hardening response

The as-quenched A1-6Ti-1.5Ni alloy ribbons had an average microhardness of 133 kg mm -2 (EVH), as indicated in the results of micro-indentation hardness measurements recorded in Fig. 1. This high initial value of hardness is attributable to a combination of solid solution strengthening associated with ,-~3.5 wt.% Ti retained in supersaturated c~-A1 matrix solid solution [15] and dispersion hardening arising from a uniform, 200 = 190 "~ 180 t70 .~ 160 ;~ 150 140 '~ 130 120 ©

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Ageing Time (h) Fig. i. Microhardnessmeasurementsfor rapidly solidifiedAI-6Ti1.5Ni alloy,agedisothermallyat 300°C, 350°Cand 400°C. Error bars represent standard deviation on average of 40 individual measurements for each sample.

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value of ~ 170 EVH after 720 h. With continued ageing, the microhardness began to decline gradually but remained substantially higher than that of the as-quenched alloy even after 2400 h. 3.2. Microstructural evolution &¢ring ageing 3.2.1. Isothermal ageing at 400°C

Fig. 2. (a) Transmission electron micrograph typical of the microstructure of rapidly solidifiedAI-6Ti-I.5Ni alloy,and (b) SAED pattern from similar region. fine-scale distribution of primary intermetallic phases throughout the ribbon thickness. The microstructure of the as-quenched alloy shown in Fig. 2 has been characterised in detail in [15]. The dispersed primary phases are novel metastable intermetallics; the major cuboidal phase is ternary, face-centred cubic (f.c.c.) (Fm3c, ao = 2.42 nm) and the minor phase orthorhombic (ao = 1.80, bo = 1.40, co = 2.20 nm). The thermal stability of the rapidly quenched microstructure was examined by carrying out isothermal ageing treatments in the temperature range 300-500°C. As shown in Fig. 1, these treatments generated a strong age-hardening response. For heat treatments at 400°C, the microhardness rose rapidly to a maximum of ~ 170 EVH after 5 h and then declined. It remained, however, approximately equivalent to that of the as-quenched alloy after 240 h exposure at 400°C. For isothermal ageing at 350°C, the average microhardness increased to ~ 175 EVH after 240 h and remained higher than that of the as-quenched alloy after 720 h. With isothermal ageing at 300°C, the average microhardness increased steadily with increasing ageing time, reaching a

The primary intermetallic dispersoids of the metastable cubic and orthorhombic phases in the asquenched microstructure proved to be thermally unstable, and they decomposed rapidly in the initial stages of ageing at 400°C. Fig. 3(a) shows the microstructure typical of a specimen aged 1 h at 400°C. Those local regions of darker contrast correspond to the original primary intermetallic dispersoids, but the cuboidal shape is no longer evident and the contrast variation within these regions suggests decomposition of the particles. This contrast pattern varied with orientation of the thin foil specimen, suggesting that the contrast was now primarily the result of local strain fields. The series of two beam images shown in Fig. 4 reveal lines of no contrast through those regions formerly occupied by primary particles and these lines of zero contrast are invariably nolrnal to the operating diffraction vector, implying a particle or cluster of similarly-oriented fine particles now coherent with the matrix phase. Decomposition of the metastabte primary phase was confirmed using electron diffraction. Selected area electron diffraction (SAED) patterns from regions containing these particles, Fig. 3(b), no longer revealed any evidence of the ternary cubic phase, but a single set of weak additional reflections. As confirmed in electron microdiffraction patterns recorded from the location of individual particles, Fig. 3(c), these additional reflections could be indexed consistently according to an ordered cubic Lla structure with a lattice parameter of 0.4 nm. Based on these diffraction patterns, it is suggested that the metastable ternary phase has been replaced by solid state intermetallic precipitates of metastable L12 phase. These precipitates were inferred to share an identity orientation relationship with the ~-A1 matrix phase: (001)LIJ/(001)~, [100]L1,//[100L, but it was difficult to establish the scale and morphology of individual L12 precipitates at this stage of ageing, since they were densely clustered and surrounded by strain fields. In addition to the coherent precipitation of the metastable L12 phase at those locations originally occupied by primary dispersoids of the metastable ternary phase, an extremely fine precipitate structure was also observed in the central regions of some C-A1 grains and in areas surrounding the decomposed particles. In the bright field electron micrograph shown in Fig. 3(a), there appears a very fine modulated contrast throughout most of the matrix phase. In the regions outlined,

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Fig. 3. (a) Transmissionelectron micrograph showing typical microstructure of ribbon aged I h at 400°C, (b) corresponding (001)~ SAED pattern, and (c) electron microdiffractionpatterns from locations occupied originallyby primary ternary phase. The intense reflectionsare from (001)= matrix phase, and the weaker reflectionscorrespond to ordered cubic L12 phase. fine-scale striations defined by the modulation in contrast are aligned approximately normal to the operating diffraction vector. After only 1 h at 400°C the contrast modulation was on such a fine-scale and of such low intensity that it could be detected only for certain orientations of the specimen. Electron microdiffraction patterns were recorded from such areas with extended exposures in an attempt to reveal whether a structural modification could be associated with the contrast modulation. The structural change was best revealed in the electron microdiffraction patterns recorded with the electron beam parallel to (001)~ directions. The additional weak reflections could again be indexed for metastable, ordered cubic L12 phase. Based on these initial observations and those following microstructural evolution at later stages of ageing, the extremely fine structure is inferred to represent the very early stages of homogeneous coherent precipitation of the metastable L1, phase. With increased ageing time at 400°C, decomposition of the supersaturated c~-A1 solid solution became more pronounced. In specimens aged 5 h at 400°C, a uniform distribution of fine-scale, coherent L12 precipitates was evident in most grains, as shown in the bright field and ( - g , 3g) weak beam centred dark field (CDF) images of Fig. 5. Since the reflections from the metastable intermetallic precipitates were extremely weak, the

CDF image was obtained using common reflections of L12 and e.-A1 phases. As the precipitates remained coherent with the c~-A1 matrix phase it was found that, with appropriate tilting of the specimen, the particles or particle strain fields could be imaged in bright contrast. The microstructure contained a high density of closelyspaced fine particles and, although the contrast in the matrix was revealed using common reflections from the L12 and cz-A1 phases, it is suggested, based on the diffraction evidence from such regions, that it arises from coherent L12 precipitates. From such images, it was estimated that the average particle size or, perhaps more correctly, the extent of the particle strain field at this stage of ageing was approximately 5 nm. The intensity of reflections from the metastable L12 phase was increased significantly in (001)~ zone axis patterns from the sample aged 5 h at 400°C, Fig. 5(c). In addition, very weak streaks of intensity were observed to emerge in the pattern. These streaks were associated with the reflections from the metastabte L12 phase, and extended parallel to the (100)~ directions of the c~-A1. Their presence was the first evidence of a structural modulation occurring within the metastable L12 precipitates, and this modulation is described in more detail in part II [17]. After ageing for 24 h at 400°C, precipitation was more pronounced and local coarsening of the fine-scale,

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Fig. 4. Two-beam transmission electron micrographs showing the matrix strain contrast around intermetallic particles in sample aged 1 h at 400°C. Electron beam is approximately parallel to <011>~.

metastable L12 precipitates enabled the coherent structure of individual precipitates to be confirmed. The series of two-beam images recorded approximately parallel to an (001>.~ axis and shown in Fig. 6, reveal the twin-lobe dark contrast typical of coherent, spheroidal or cuboidal precipitates, with the line of zero contrast separating the lobes invariably normal to the operating g = <200>~ reflection. At this stage of ageing, the extent of the average precipitate strain field was estimated to be ~ 10 nm. It is also noteworthy that in local areas of the microstructure, particularly those containing an initially high density of primary particles, the solid state precipitates began to take on a cross-like contrast when observed in <001>~ projection, Fig. 7(a). The strain contrast associated with individual precipitates appeared to comprise separate pairs of twin black lobes distinguished by orthogonal lines of zero contrast perpendicular to

<100>~ directions. In addition, those sites initially occupied by primary cuboidal particles in the asquenched ribbon were now revealed more clearly to comprise dense aggregates of fine-scale cross-like precipitates. In the corresponding SAED patterns, Fig. 7(b), there were now prominent streaks of diffuse intensity parallel to <001>~ through 1/2{200}~, 1/2{020}~ and 1/2{220}~. positions, i.e. positions corresponding to the {100}Lla, {010}L12 and {ll0}L, 2 reflections respectively. There was also now clear evidence of distinguishable intensity maxima within these streaks and straddling the L12 reflections at the positions 1/2{200}~, 1/2{020}~ and 1/2{220}~. These reflections represent evidence of the emergence of a one-dimensional, long period superstructure within the initially L1 a particles and a detailed analysis of this structure will be presented in part II [17].

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Micrographs typical of the microstructure in ribbons aged for 240 h at 400°C and the corresponding (00t)~ SAED pattern are presented in Fig. 8. Coarse precipitate aggregates were no longer evident at the initial sites of primary ternary phase particles in this sample, and

the microstructure now contained a relatively high density of finely dispersed, coherent precipitates throughout most of the ~-A1 grains, Fig. 8(a). The distribution of these coherent precipitates was generally uniform, although a dendritic pattern to the distribution, with the core approximately in the grain centre, was detected in some grains, perhaps reflecting the initial pattern of dendritic solidification of individual crystals. As determined from weak beam centred-dark field images, the average range of individual particle strain fields was 20 nm in diameter. Individual coherent metastable precipitates were now observed to have a well-defined three-dimensional cross-like shape, when projected parallel to (001)~, Fig. 8(b). The extension of diffuse intensity, about the 1/2{200}~, 1/2{020}~ and 1/2{220}~ positions and parallel to (100)~, became more pronounced in the (001)~ SAED pattern at this stage, Fig. 8(@ Each diffuse streak contained subsidiary intensity maxima which were adjacent to and distributed symmetrically about the 1/2{200}~, 1/2{020}~ and 1/2{220}~ positions. There were no longer single discrete diffraction spots at the 1/2{200}~, 1/2{020}~ and 1/2{220}~ type positions, but groups of partially superimposed reflections which could be attributed to direct and double diffraction from a mixture of L12 and D022 phases. The emergence of the distributed intensity maxima about these locations could be interpreted on the basis of the development of precipitate structures intermediate between the initial L12 structure and the equilibrium D0= phase. In addition, diffraction maxima at positions of the t y p e 1/4{420}~ and 1/4{240}~ could be attributed to direct and double diffraction from the equilibrium D0= phase (b.c.t., a = 0.385 nm, c = 0.861 nm). While the majority of precipitate particles retained an L1 a structure or a structure intermediate between L12 and the equilibrium phase, a significant fraction of particles of the equilibrium D0= phase were detectable at this stage of ageing. These equilibrium particles could be distinguished readily from those of the metastable phase on the basis of size and shape; they were generally spheroidal in form and much coarser (50-100 nm diameter) than the metastable coherent particles. A more detailed analysis of the changes in precipitate structure accompanying isothermal ageing forms the subject of part II [17]. 3.2.2. Isothermal ageing at 300°C

Fig. 5. (a) Bright field, and (b) corresponding weak-beam centred dark field ( - g , 3g) micrographs showing uniform precipitation of fine-scale, coherent particles in a sample aged 5 h at 400°C; electron beam is approximately parallel to (110)~. (c) (001)~ SAED pattern recorded from region such as in (a).

A similar sequence of solid state precipitation was observed in samples of rapidly-solidified ribbon aged at temperatures above and below 400°C, with the kinetics of precipitation varying with temperature in the anticipated manner. The microstructure typical of ribbon aged 24 h at 300°C is presented in Fig. 9(a). The cuboidal shape of the primary intermetallic particles is no longer evident, and the Sites of the original particles

J.F. Nie, B.C. Muddle / Materials Science and Eng#wering A22I (t996) I 1 - 2 I

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Fig. 6. Two-beam bright field images showing the coherent nature of individual, finely-dispersed precipitates in alloy aged 24 h at 400°C. Electron beam is approximately parallel to (001)~.

are now defined by paralM bands of light and dark contrast. According to electron microdiffraction patterns recorded from these locations, the primary particles of the metastable ternary phase had been replaced completely by the precipitation of the metastable L12 phase. However, at this stage no evidence was detected of decomposition of the supersaturated c~-A1solid solution in regions remote from these locations. Fig. 9(b) shows a micrograph typical of the microstructure in ribbon aged 240 h at 300°C. The metastable coherent L12 precipitates, formed at those locations originally occupied by the metastable primary particles, were still finely dispersed within the matrix phase. In this case, however, a very fine mottled contrast was now also evident within ~-A1 grains. Electron microdiffraction patterns recorded from such regions, Fig. 9(c), contained weak diffraction spots consistent with the ordered L12 structure and indicated that the very fine modulated contrast could be associated with the very early stages of coherent precipitation of the metastable L12 phase. The lengthy duration of the heat treatment required to initiate detectable solid state precipitation at 300°C is to be emphasised.

3.2.3. Isothermal ageing at 500°C

With ageing at 500°C, the precipitation sequence was accelerated and the microstructure of specimens aged 1 h at 500°C, Fig. 10(a), already contained finely dispersed, cross-like precipitates within the ~-A1 matrix phase. ~This microstructure and the corresponding (001)~ SAED pattern are very similar to those observed for material aged 240 h at 400°C. The diffraction pattern contains precipitate reflections at 1/4{420}~ and 1/4{240}~ positions consistent with those expected for direct and double diffraction from the equilibrium D022 phase, while diffraction maxima at positions of the type 1/2{200}~, 1/2{020}~ and 1/2{220}~ each comprise a group of partially superimposed reflections that may be attributed to direct and double diffraction from a mixture of L12 and D0= phases. Adjacent to these reflections and symmetrically disposed about them, is a series of weaker intensity maxima, distributed parallel to the (100)~ directions. Their presence contributes to the diffuse intensity lying parallel to (100)~ and passing through the 1/2{200}~, 1/2{020}~ and 1/2{220}~ matrix locations. These low intensity reflections could not be successfully indexed on the basis of either direct or

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double diffraction from the L12, D022 o r ~-A1 phases, or some combination of them. They thus imply the presence of a metastable transition phase or phases within the precipitate particles [17]. After ageing for 24 h at 500°C, Fig. l l(a), the cross-like precipitates were replaced completely by coarser spheroidal particles ~ 0.1-0.2 ~tm in diameter. Most of these relatively coarse particles were identified to be of the equilibrium D022 phase 6-A13(Ti,Ni). According to electron microdiffraction patterns recorded from these particles, Fig. l l(b), they invariably maintained a rational orientation relationship with the ~-AI matrix phase which, as shown in the schematic solution of Fig. 11(c), was such that the principal axes of the t e t r a g o n a l D022 phase were parallel to those of the c~ matrix: (001)DoJ/(001),~, [100]DOJ/[100L.

s0 nm Fig. 8. Transmission electron micrographs showing (a) the microstructure typical of ribbons aged 240 h at 400°C, (b) the cross-like form of individual intermediate precipitates in this sample; and (e) corresponding {001 >~ SAED pattern. Images recorded approximately parallel to <001)~.

4. Discussion

Fig. 7. (a) Transmission electron micrograph showing the microstructure typical of ribbons aged for 24 h at 400°C, and (b) corresponding <001)= SAED pattern. Image recorded approximately parallel to

~.

In so far as it comprises a uniform dispersion of nano-scale ( < 100 nm) particles of equilibrium intermetallic c~-A13(Ti,Ni) in solid solution strengthened aluminium matrix phase, the microstructure of the rapidly solidified alloy aged 24 h at 500°C, Fig. 11, demonstrates the potential to satisfy the aim [15] of producing a thermally stable, dispersion-strengthened alloy in this ternary A I - T i - N i system. Similar microstructures have not proven possible in binary AI-Ti alloys using a similar combination of rapid solidification and post-solidification heat treatment, largely because of the tendency for binary alloys (2-10 wt.% Ti) to solidify containing coarse particles (0.1-0.3 gm) of metastable L12 intermetallic phase as a primary solidification product. However, in ternary alloys, with a weight ratio Ti:Ni in the range 3:1-4:1, formation of this Lla phase

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is suppressed and replaced by a ternary cubic phase which is finer-scale and more uniformly dispersed [15]. During subsequent isothermal heat treatment in the temperature range 300-500°C, the particles of ternary phase prove to be metastable and decompose rapidly on initial ageing to be replaced by dense clusters°of very

,---~

Fig. I0. (a) Transmission electron micrograph showing the microstructure typical of ribbons aged 1 h at 500°C, and (b) corresponding (0015~, SAED pattern.

Fig. 9. Transmission electron micrographs showing the microstructure typical of ribbons aged (a) 24 h at 300°C, and (b) 240 h at 300°C; (c) the {110)~ electron microdiffraction pattern recorded from matrix region of (b) includes weak reflections arising from LI2 structure.

fine-scale, solid-state precipitates of the metastable L12 phase. This process is accompanied eventually by precipitation of coherent, metastable Lt2 particles within the matrix phase, in a distribution that is more dense and uniform than anything that has been demonstrated thus far in binary AI-Ti alloys [9-11]. This uniform distribution of solid state precipitates appears to be attributable to the initially uniform, fine-scale dispersion of primary particles of the metastable ternary cubic phase, which redistribute their solute content to the matrix as they decompose, and to an initially relatively uniform distribution of titanium and nickel in the surrounding matrix phase in the as-quenched alloys. The solute supersaturation is subsequently relieved by a solid state precipitation sequence, which culminates in the formation of the equilibrium /~-A13(Ti,Ni). While it satisfied the initial aim of producing a fine-scale, thelwnally stable dispersion-strengthened microstructure, the microstructure that is shown in Fig. 11 was not the most promising observed in the present work, for it represented an overaged structure with less than optimum mechanical properties. Associated with

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the uniform, fine-scale precipitation of coherent metastable L12 phase, and the subsequent intermediate structures based upon it, was a strong age-hardening response which, whilst conventional in form, occurred within reasonable times only at temperatures in excess of 300°C. This temperature range is much higher than the ageing temperatures for conventional precipitationhardening aluminium alloys (typically 150-200°C) [2]. At the same time, the maximum hardness levels achievable (e.g. ~. 175 EVH after 240 h at 350°C, Fig. 1) are comparable with those of conventional high strength aluminium alloys (typically 150-200 VHN) [2,6], even when allowance is made for the possibility that the micro-indentation measurements may overestimate hardness levels compared with conventional bulk measurements. Maximum hardness in the aged samples was invariably associated with a mixed precipitate dispersion comprising coherent particles having the structure of metastable L12 phase and of intermediate phases having a structure based on LI=, with a small volume fraction of equilibrium D0=2 A13(Ti,Ni). That the L12-based phases precede formation of the equilibrium phase is

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Fig. 11. (a) Transmission electron micrograph showing the microstructure typical of ribbons aged 24 h at 500°C, (b) typical [100] electronmicrodiffractionpattern recordedfromintermetallicparticles in (a), and (c) schematicsolution of (b). Underlinedindicesrefer to [100] c~-A1matrix.

readily rationalised on the basis that the dimensions of the L12 unit cell are similar to that of c~-A1, and that the L12 phase and the intermediate phases derived from it remain coherent with the ~-A1 matrix. They thus face a smaller energy barrier to nucleation and acquire a kinetic advantage over the equilibrium phase. That the decomposition of supersaturated e solid solution proceeds only sluggishly, even though the energy barrier to metastable precipitate nucleation is relatively low, the effective level of undercooling is substantial and the level of supersaturation is high, suggests that the kinetics of precipitation are controlled by atomic mobility and reflect the low diffusivities of Ti, and to a lesser degree Ni, in aluminium. With the progress of ageing, there was progressive transformation within and coarsening of the metastable precipitates, accompanied by a decline in hardness. However, these metastable precipitates were quite resistant to coarsening at temperatures as high as 400°C. The microstructure still contained a fine-scale ( < 0.1 tam) dispersion of coherent metastable precipitates after 240 h at 400°C, and the hardness remained higher than that of the as-quenched alloy. At 300°C, precipitation was sluggish and coherent metastable L12 precipitates remained very finely dispersed even after 720 h. Although a decrease in hardness was noted to occur beyond 720 h of treatment, the microhardness remained substantially higher than that of the as-quenched alloy even after 2400 h. The low coarsening rates observed for these precipitates may be attributed to a combination of factors, including the low equilibrium solid solubility of Ti in A1, the low diffusivities of Ti and Ni in A1 at these temperatures and the relatively low interracial energies of the metastable precipitates with respect to the ~-A1 matrix phase. Given that the rates of precipitation and coarsening of both metastable and equilibrium precipitate phases are low at temperatures of 300°C and above, it is to be expected that the microstructures of dispersed metastable particles corresponding to maximum hardness, Fig. 5, and of dispersed nano-scale equilibrium ~-A13(Ti,Ni), Fig. 11, will prove highly thermally stable at lower temperatures. Rapidly solidified A I - T i - N i alloys aged to maximum hardness, at temperatures of 350-400°C, might be expected to exhibit excellent thermal stability and thus retain superior mechanical properties at elevated temperatures of 150-200°C, where conventional precipitation-hardening aluminium alloys are prone to overageing and a loss of useful strength [2]. If the microstructures demonstrated in these thin ribbons can be reproduced in bulk components of useful scale, then low density, rapidly quenched AITi-Ni alloys, with the weight ratio of Ti:Ni in the range 3:1-4:1, may have significant potential for applications involving elevated temperatures.

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5. Conclusions

Acknowledgements

(i) The primary intermetallic dispersoids of ternary cubic phase in rapidly solidified AI-6Ti-I.5Ni (wt.%) alloy are metastable and decompose rapidly in the initial stages of any heat treatment in the temperature range 300-500°C. They are replaced by dense clusters of fine-scale solid state precipitates of a metastable LI z phase and this process is accompanied by uniform homogeneous precipitation of fine-scale, coherent particles of the metastable L12 phase within the ~-A1 matrix phase. There is an identity orientation relationship between the metastable L12 phase and matrix. (ii) During isothermal ageing, the metastable precipitates evolve into a transitional, three-dimensional cross-like morphology and eventually into nano-scale ( < 100 am) spheroidal particles of the equilibrium D022 phase a-A13(Ti,Ni). These changes in form are accompanied by changes in structure, involving the development of a series of one-dimensional long period superlattices and culminating in the formation of the equilibrium b.c.t, phase. These changes in structure are addressed in detail in the companion paper [17]. (iii) Precipitation within rapidly quenched A1-6Ti1.5Ni alloy is accompanied by a strong hardening response, which produces hardness values (175 EVH) comparable with those of conventional high strength precipitation-hardening alloys (150-200 VHN), but at ageing temperatures (300-500°C), well in excess of those employed for conventional aluminium alloys (150-200°C). The temperatures of this response, together with the apparent resistance of the precipitate phase(s) to coarsening at temperatures up to 400°C, suggest that low density, rapidly quenched A1-Ti-Ni alloys with Ti:Ni weight ratio in the range 3:1-4:1, may have significant potential for applications involving elevated temperatures (150-200°C), where conventional precipitation-hardening alloys overage rapidly.

This research was supported by the Australian Research Council. One of the authors (JFN) acknowledges gratefully the support of a Monash Graduate Scholarship.

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