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Ordered ω phases in a 4Zr–4Nb-containing TiAl-based alloy
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Acta Materialia 56 (2008) 1689–1700 www.elsevier.com/locate/actamat
Ordered x phases in a 4Zr–4Nb-containing TiAl-based alloy Z.W. Huang * School of Materials Science and Engineering, Southwest Jiaotong University, Chengdu, Sichuan 610031, PR China Department of Metallurgy and Materials, University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Received 28 September 2007; received in revised form 5 December 2007; accepted 7 December 2007 Available online 1 February 2008
Abstract A considerable amount of B2 phase with a cellular morphology is retained in a 4Zr–4Nb-containing TiAl-based alloy. Heterogeneous precipitation of ordered x from B2 is found to occur readily after HIPping: B2 ? x with the B82-structure in cell regions and B2 ? x with the D88-structure in cell-wall regions. Congregated D88-x domains and particles form as a network surrounding the well-developed B82-x cells. The heterogeneous formation of different x variants is caused by a heterogeneous distribution of Zr + Nb elements across B2, which plays an important role in stabilizing vacancies and promotes the formation of D88-x. Fine D88-x particles are also observed to precipitate from the B82-x cell matrix after ageing at 700 °C for 1000 h, showing a transformation path of b ? B2 ? B82-x ? D88-x for the aged cells. The heterogeneous formation of a D88-x network and B82-x cells is found to be detrimental to ductility and fatigue strength. A very brittle fine-grained TiAl alloy is produced as a result. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Titanium aluminides; TEM; Phase transformation; Heterogeneous nucleation of phase; Mechanical properties
1. Introduction In the development of high strength TiAl-based alloys, alloy design based on major alloying of the Ti-Al binary system has been applied to achieve overall improvements in properties. Alloys such as Ti–44Al–8X–1B, where X represents combinations of Nb, Zr, Ta and Hf, have been chosen for detailed studies [1,2]. These refractory elements are able to increase both high-temperature creep resistance and oxidation resistance. However, they are all b-phase stabilizers. A high level of such elements added to TiAl alloys may cause varied amounts of b-phase to be retained, thereby introducing an (a + b) phase field for the alloy system. This is expected to facilitate thermomechanical processing at elevated temperatures, since a grain growth can be restricted, particularly when the volume fraction of the bphase is high [3]. On the other hand, increasing evidence has indicated that the stabilized b-phase in highly alloyed *
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TiAl alloys can be retained to room temperature. The retained b has always been observed as an ordered B2 structure which can readily decompose into the x phase during cooling to room temperature [1,2,4–7]. The formation of x-type phases has been extensively investigated in transition-metal-based body-centred cubic (bcc) alloys and is well documented in the literature [8– 27]. It is generally accepted that the transformation of b(B2) to x phase is a displacive process in which one pair of {2 2 2}b planes collapses to the intermediate position along one h1 1 1ib direction, leaving the next plane unaltered [10,11,17]. The {2 2 2}b plane collapse could be incomplete (partial collapse), in which case the pair of {2 2 2}b planes does not collapse into the double layer to produce an ideal hexagonal P6/mmm structure. Such a partial collapse produces a trigonal P3m1 structure [17,20–22,24]. The product phase was termed diffuse x if exhibiting diffuse x scattering/reflections, which deviates from the ideal crystalline x maxima along the line through (0 0 0) and (2 2 2)b [12,28–30], or termed a crystalline x if exhibiting discreet diffraction maxima at the position of
1359-6454/$34.00 Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2007.12.013
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ideal crystalline x lattice. The crystalline x can inherit B2 chemical order to become an ordered x, although it was still considered as an intermediate phase in the course of structural change [19–21]. More stabilized x-type products with further chemical ordering and fully crystalline lattice structures such as B82 and D88 were reported in various transition-metal-based alloys containing varied percentages of Al [18–21,25–27]. These x-type products are transformed either directly from parent b(B2) or from metastable intermediate x by a superimposition of displacive and replacive ordering reactions. The two x-type products often form after a prolonged ageing at elevated temperature and are therefore closer to the equilibrium composition and structure and always fully three-dimensionally developed in morphology. c-TiAl-based alloys of engineering importance are normally based on Ti–(44–49 at.%) Al with appropriate combination of alloying elements. Al was reported to increase the x-phase stability by lowering the valence electron density of B2, while the transition metals added may increase the valence electron density of B2 [31]. It would be very interesting to know how the B2 ? x transformation would be affected under the interaction between Al and transition metals. Research into x formation and transformation within multicomponent TiAl-based alloys has rarely been reported [4–7]. Detailed investigation is therefore needed for a clear understanding of the x phase in such types of TiAl alloys. The work should include the stability of the x phase at the early stage of formation, the lattice structure, the transformation path, the size and morphology of the precipitate, the transformation path during prolonged ageing at elevated temperature and the effects on mechanical properties. The x formation is well known to result in brittleness in conventional b-Ti alloys and other transition-metal-based bcc alloys [8,9,13,32]. The co-presence of B2 and x was also reported to be detrimental to high-Nb-containing TiAl-based alloys [33]. However, more research is needed to understand clearly how and to what degree the retained b(B2) and precipitated x affect the mechanical properties of TiAl alloys, especially when the volume fraction of B2 + x is relatively large (say >8%) and exists at grain boundaries and at triple junctions, which is often the case in highly alloyed TiAl alloys [34]. The present work was carried out on a cast ingot of cTiAl-based alloy containing 4Zr and 4Nb. The alloy was subjected to a hot isostatic pressing (HIPping). The HIPped billet was cooled slowly to room temperature. The as-HIPped alloy contained multiple phases at room temperature and a significant amounts of B2 and x phases were observed [1].
process at the IRC (Interdisciplinary Research Centre), Birmingham University. The cast ingot had been doublemelted and then HIPped at 1260 °C under a pressure of 150 MPa for 4 h at Southwest Jiaotong University, Chengdu. Microstructures were examined by back-scattered electron (BSE) microscopy using a JEOL 840A scanning electron microscope (SEM) operating at 20 kV. This was performed to reveal the size, shape and distribution of the constituent phases both before and after HIPping. Image analysis was carried out on BSE micrographs to quantify the size and volume fraction of constituent phases and structures using a mean linear-intercept method. The measurements were conducted on a minimum of 20 micrographs that were randomly chosen. The total area investigated was 5.4 mm2. Thin foils for transmission electron microscopy (TEM) were prepared by a twin-jet polishing with an electrolyte of 5 vol.% perchloric acid, 30 vol.% butan-1-ol and 65 vol.% methanol operating at 30 V and at a temperature of 25/30 °C. TEM was carried out using a JEOL 2000 FX microscope operating at 200 kV. Energy dispersive X-ray (EDX) microanalysis was performed on selected microstructural features using the TEM which was fitted with a LINK analytical EDX microanalyzer. The EDX analysis on each microstructural feature was performed at six locations. For all measurements conducted, mean values are presented with errors representing the uncertainty of the statistical analysis. The errors were assessed using 95% confidence limits, based on 2s/n1/2, where s is the standard deviation and n is the number of measurements. Tensile tests were conducted at both room temperature and 700 °C on a Zwick testing machine in air, at a strain rate of 0.7 104 s1. The tensile data presented are an average of three samples for room temperature tests and an average of two samples for 700 °C tests. S–N fatigue tests were carried out at room temperature in air. Specimens were tested in pure bending using an electromagnetic resonance testing machine at a frequency of 80 Hz under a stress ratio of R = 0.1 (where R = rmin/rmax; rmin and rmax are, respectively, the minimum and maximum stresses applied over the fatigue cycle). Testpieces of dimensions 80 10 10 mm3 were prepared using electro-discharge machining. Conventional grinding and polishing to 0.25 lm diamond were used to prepare the surface. A major loading span of 60 mm and a minor loading span of 20 mm were used. 3. Results and discussion
2. Experimental
3.1. Microstructure after ingot casting and subsequent HIPping
The alloy chosen for the present study was Ti–44Al– 4Zr–4Nb–0.2Si–1B (alloy 4–4–1, at.%). It was grain refined by adding 1% of boron and supplied as a cast ingot that was made using a PACH (plasma arc cold hearth) melting
Fig. 1a and b shows the multiphase microstructure of cast ingot alloy 4–4–1 before and after 4 h/1260 °C HIPping, respectively. The alloy after ingot-casting exhibited a near-lamellar microstructure, mixed with retained b(B2)
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Table 1 The volume fraction and size of the microstructural constituents in alloy 44-1 after ingot casting and 1260 °C HIPping
Size (lm) Volume fraction (%)
Fig. 1. Back-scattered SEM images showing the microstructure of (a) cast ingot and (b) after 4 h/1260 °C HIPping.
and a(a2) and equiaxed c. As reported previously, the phase transformation sequence for the 4Zr–4Nb-containing TiAl alloy was [2]: L!Lþb!b!aþb!aþbþc!aþc There was no single a phase field during the solidification process. After HIPping at 1260 °C, the alloy still exhibited a near-lamellar microstructure, consisting of lamellar colonies (70 vol.%) and equiaxed B2, a2 and c grains (30 vol.% in total). The lamellar colonies were measured to be 70 lm (±11.2 lm) in size. The retained B2 (11.5 ± 3.7 lm in size) and retained a2 (5.2 ± 1.2 lm in size) grains were observed mainly around lamellar colonies and c grain boundaries. The volume fractions and mean sizes (with the error range) of all the microstructural constituents determined by image analysis of BSE micrographs are summarized in Table 1. The microstructure after HIPping, in which considerable amounts of equiaxed B2, a2 and c grains remain with a2 + c lamellar structure, makes the 4–4–1 alloy quite different from other highly alloyed and grain-refined c-TiAl alloys such as Ti–44Al–8Nb–1B, Ti–44Al–4Nb–4Hf–1B and Ti–44Al–4Nb–4Ta–1B, all of
Lamellar colonies
Equiaxed a2
Equiaxed c
b(B2)
70.9 ± 11.2 71.2 ± 10.7
5.2 ± 1.2 4.2 ± 0.8
19.3 ± 4.8 15.7 ± 4.5
11.5 ± 3.7 8.9 ± 2.5
which show a fully lamellar microstructure, although small amounts of b are retained. Fig. 2a, a bright field (BF) TEM micrograph, shows all the major microstructural constituents obtained in 4–4–1 alloy after 1260 °C HIPping. The a2 + c lamellae (labelled as L) are the major constituents in this alloy. The corresponding selected area diffraction patterns (SADPs) from the a2 + c colonies are shown in Fig. 2b. A considerable amount (9 vol.%) of retained b-phase was also observed. The retained high temperature phase was always located at lamellar colony boundaries. The reason why the retained b segregates around lamellar colony boundary was explained in a previous paper, based on the b ? a phase transformation in which the solute redistribution factor k0 = Ca/ Cb > 1 [34]. Unlike the usual monolithic morphology of retained b observed in other highly alloyed TiAl-based alloys, such as Ti–46Al–5Nb–1W and Ti–44Al–8Nb–1B [6,35], the retained b(B2) in the present alloy exhibits a cellular morphology, consisting of many small cells (1–2 lm in size), between which are narrow cell-wall regions. Compared to the relatively smooth cell surface, the cell-wall regions look rough, owing to a congregation of small domain and particle clusters in narrow areas. SADPs taken from the cell structure show strong reflections corresponding to the ordered B2 lattice and always contain x reflections (Fig. 2c). The cell structure is therefore labelled as B2 + x in Fig. 2a. It is important to note that the B2 + x cell structure always keeps a general orientation relationship with their neighbouring a2 + c lamellar structure. This clearly indicates that both the transformed a2 + c lamellae and the retained B2 + x originate from a single b grain. Therefore, the general orientation relationships among the four phases shown in Fig. 2a can be defined as f1 1 1gc ==f0 0 0 1ga2 ==f1 1 0gB2 ==f1 1 2 0gx and h1 1 0ic ==h1 1 2 0ia2 ==h1 1 1>B2 ==h0 0 0 1ix A small amount of equiaxed a2 was also observed at room temperature. The diffraction pattern taken from an equiaxed a2 grain (Fig. 2d) shows that the equiaxed a2 has the same orientation as the a2 in the lamellar structure, indicating that the a2 is the remnant of the a grain after ordering and eutectoid transformation a ? a2 ? a2 + c. The retained a2 therefore shows the same lattice structure and orientation as its transformed counterpart. This clearly indicates that the solid phase transformation to a2 + c is
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Fig. 2. (a) TEM BF (bright field) micrograph showing a multiphase microstructure consisting of a2 + c lamellae (L), cellular B2 + x, equiaxed a2 and c phases. (b–e) The selected-area diffraction patterns taken from the indicated lamellae, B2 + x cell structure, retained a2 and c grain, respectively, with B ¼ h1 1 2 0ia2 ==h1 1 0ic in (b), B ¼ h1 1 1iB2 ==h0 0 0 1ix in (c), B ¼ h1 1 2 0ia2 in (d) and B ¼ h1 1 4ic in (e).
not complete in this highly alloyed TiAl alloy. Equiaxed c grains were found to be distributed extensively in the asHIPped alloy since the single a-phase field is replaced by a (a + b + c) phase field in alloy 4–4–1 [2]. Some of the equiaxed c grains were small in size and located close to retained B2 and a2, as indicated in Fig. 2a. They were independent of neighbouring a2 + c lamellae as well as
equi- axed B2 and a2 grains in terms of the orientation relationship (Fig. 2e). 3.2. Ordered x phases after HIPping The formation of x from B2 has been described as a displacive process in which one pair of (2 2 2)B2 planes
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collapses to the intermediate position along a h1 1 1iB2 direction [10,11]. Four variants of x are expected to form from the parent B2, corresponding to the interplanar collapse along each of the four h1 1 1iB2 directions. Fig. 3a is a centred dark field (CDF) image taken from the B2 + x cell structure with a 1 0 1 0x reflection, corresponding to a single variant of x (see Fig. 3b). All the illuminated cells therefore belong to one rotational variant of the x phase. It is observed that these x cells are actually composed of sub-micron-sized domains and these domains are separated by internal interfaces. The pure set of x diffraction spots along one of the h1 1 2 0ix zone axes shown in Fig. 3b indicates that the x inside the smooth cells/domains is fully transformed from B2. The precipitation of x throughout the retained B2 blocks was found not to be uniform in terms of the size and distribution of the precipitates. Fig. 4a and b, using both BF and CDF images, reveal that the rough cell-wall regions (e.g. B and D) are in fact composed of many small x domains and particle clusters. These congregated x domains and particles formed as a network surrounding the relatively smooth x cells (e.g. cell C). SADPs along the h1 1 0i>B2 direction obtained from the rough region B (Fig. 4c) show strong reflections corresponding to B2, indicating that some B2 was left in the narrow cell-wall regions after the B2 ? x transformation. The finer diffraction maxima in the complicated pattern correspond to various variants of the precipitated x. In contrast to the complex condition of the cell-wall region, individual x cells are single crystal in structure. As shown in Fig. 4d and e, the SADP taken either from cell A or cell C shows a diffraction pattern of one variant only: ½1 1 0 2x for cell A and ½1 1 2 0x for cell C. The micrographs shown in Fig. 4 clearly indicate that the x phase in the as-HIPped alloy 4–4–1 was precipitated
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non-uniformly across retained b blocks. While the individual x cells which showed a single crystalline structure had a size range of 1–2 lm, the small x domains and particles were <100 nm. These small domains and particles congregated into narrow areas between the relatively smooth x cells, co-existed with untransformed B2 and were connected to form a network across the retained grain. Fig. 5a–c shows a set of SADPs (zone axes using cubic indices are [1 0 0]B2, [1 1 1]B2 and [1 1 2]B2, respectively) obtained specifically from individual smooth x cells in the as-HIPped 4–4–1 alloy. The x cells do not generate any diffuse scatterings in their diffraction patterns. Discreet and sharp x spots without any diffraction streaking and any shift along the line through (0 0 0) and (2 2 2)B2 can be observed, suggesting that the precipitated x phase had already reached a mature, fully crystalline condition after HIPping. This is different from conventional Ti-based alloys [22,28–30], in which the maxima corresponding to the x phase exhibited short streaks and shift from the ideal crystalline position. This is also different from the Nb + Wcontaining TiAl-based alloy studied previously. The diffraction maxima corresponding to an early type of x phase in this alloy showed short streaks perpendicular to the ½2 2 2B2 and ½2 2 2B2 directions, although no shift of diffraction maxima was observed [6]. It appears that once the high-temperature b-phase in the Zr + Nb-containing TiAl alloy cooled down to room temperature after HIPping, not only was the collapse of the double layers in B2 structure complete, leading to the formation of an ideal crystalline x, but also the precipitated x phase was already three-dimensionally developed, as imaged in the CDF micrographs of Figs. 3a and 4b. Further TEM work on the x cells was conducted to determine which lattice structure and space group the
Fig. 3. (a) TEM CDF image taken with a 1 0 1 0x reflection ðB ¼ h1 2 1 0ix Þ showing only one x variant in the cell structure. (b) The corresponding diffraction pattern from this variant. Note that the x cells were divided into submicron domains by internal interfaces.
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Fig. 4. (a) TEM BF and (b) CDF images showing the x domains and particle clusters precipitated favourably in the rough cell-wall regions B and D, forming a network around the relatively smooth x cell C. (c) SADPs taken from region B along the beam direction of [1 1 0]B2. (d) SADP from cell A ðB ¼ h1 1 0 2ix Þ and (e) from cell C ðB ¼ h1 1 2 0ix .
x belongs to: trigonal with P3m1 or B82 hexagonal with P63/mmc. The smallest diffraction aperture was used on a single x cell to produce a pure set of x spots in one of h1 1 2 0ix zone axes. Tilting along a row of 0 0 0lx reflections was performed to allow this row to be excited, with the aim of eliminating roots for double diffraction. The insert in Fig. 5c shows a row of excited microdiffraction spots obtained from a single x cell, revealing the distinction of odd reflections. The distinction of odd 0 0 0lx reflection suggests that the x-phase in these smooth cells had transformed into a hexagonal structure rather than a trigonal symmetry. The formation of the more stabilized hexagonal structure was probably caused by the slow cooling which is a normal practice after HIPping. Fig. 5d–f shows another set of SADPs (same zone axes [1 0 0]B2, [1 1 1]B2 and [1 1 0]B2), obtained specifically from the rough cell-wall regions composed of small x domains and particle clusters in the same alloy. Analysis of the two sets of SADPs indicates significant differences, especially in the zone axes [1 0 0]B2 and [1 1 1]B2, where extra spots exist at 1/3{1 1 0}B2 and 2/3{1 1 0}B2. The SADPs in the [1 0 0]B2 and [1 1 1]B2 directions are therefore more complicated than those taken from the smooth cell regions. Those extra reflections in Fig. 5d and e look like superlattice spots. The corresponding microstructure of small x
domains and particle clusters, as shown in Fig. 6a and b, is distinctly different from smooth x cells and is imaged in the CDF (Fig. 6b) separately from those smooth x cells, using the reflections in Fig. 6c which are additional to those showing in Fig. 6d and e (i.e. those extra spots at 1/ 3{1 1 0}B2 and 2/3{1 1 0}B2 in Fig. 6c). It has been established that the x diffraction patterns from the rough cellwall regions correspond to a D88 structure [21,25,26], whereas those x patterns from smooth cells correspond to a B82 structure [18–20]. From the SADPs in Fig. 5d–f, it was found that the lattice parameters of the D88 unit cell are related to the B2 structure as: aD = 61/2aB2 and cD = 31/2aB2 (subscript ‘‘D” denotes D88-type x phase). For the B82-type x-phase, the lattice parameters are: aB = 21/ 2 aB2 and cB = 31/2aB2 (subscript ‘‘B” denotes B82-type x phase). The B82 structure was reported to be an ideally crystalline x-derivative formed by a complete collapse of one pair of {2 2 2}B2 planes, the double layers are therefore not rumpled [18–21]. As confirmed in Fig. 5c, the x-type B82 phase shows a hexagonal symmetry (P63/mmc space group) rather than a trigonal symmetry with P3m1 space group. From the SADPs in Fig. 5d–f, the D88-type structure is also an ideally crystalline x-derivative formed by the same complete collapse of the {2 2 2}B2 planes. The orientation relationship between the D88-type x and parent B2 can be established:
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Fig. 5. (a–c) SADPs taken from the individual w cells, compared to those in (d–e) taken from the rough cell-wall regions. Note those additional superlattice spots in (d–e). The zone axes using cubic indices are: (a and d) h0 0 1i, (b and e) h1 1 1i and (c and f) h1 1 2i, respectively.
h1 1 1hB2 ==h0 0 0 1iD and f1 1 0gB2 ==f1 0 1 0gD This is different from the orientation relationship between the B82-type x and parent B2, which has previously been described [18,20] as h1 1 1iB2 ==h0 0 0 1iB and f1 1 0gB2 ==f1 1 2 0gB The different orientation relationship plane pair between D88 x B2 and B82 x B2 suggests that chemical ordering occurs in the {000l}x basal planes. Moreover, an ordered array of vacancies was reported to be present in the base planes of the D88-type x structure. The main difference between the B82 and D88 structures lies in whether or not the Wyckoff positions 2(b) are occupied. Two vacancies for 18 atom positions per unit cell are required for the formation of the vacancy-ordered D88 structure [21,25,26]. Based on this, the present observation of the D88-type x-phase is indicative of the retention of a high concentration of vacancies in the retained B2 phase region. However, there is no direct evidence for vacancy retention in the present TiAl alloy, although high equilibrium concentrations of vacancies in ordered bcc structures have been reported [36–38]. Tewari and co-workers [26] found that the D88 structure can readily form in Zr3Al–Nb alloys, but not in b alloys close to Zr3Al compositions. Based on this fact, they suggested that Nb had an important role to play in stabilizing
vacancies, hence promoting the D88 structure. However, a high percentage of Nb alone did not promote the formation of the D88 structure in TiAl-based alloys. For instance, the formation of D88-type x phase in Ti–45Al– 10Nb was found to be much less significant [5] than in the currently studied Ti–44Al–4Nb–4Zr–0.2Si–1B and no D88-type x phase was detected in Ti–44Al–8Nb–1B alloy, even after long-term ageing at 700 °C [35]. The present results obtained from the 4Zr + 4Nb-containing TiAl alloy tend to suggest that the combination of Zr and Nb has played a more important role in stabilizing vacancies in the retained B2 structure, therefore promoting the formation of the D88-type x structure. The present TEM examination reveals that the formation of the D88-type x domains and particles occurs only in rough cell-wall regions, which form as a network surrounding the relatively smooth B82-type x cells. This appears to indicate that the retention of vacancies occurs in a non-uniform fashion in the retained b(B2) blocks, therefore promoting the formation of the vacancy-related D88-type x in the narrow cell-wall regions as a result. The retention of vacancies in the smooth cells is assumed to be less significant since only the B82-type x phase formed there. Consequently, the retained b(B2) blocks exhibit a picture of heterogeneous x-precipitation in terms of precipitate type, size and distribution. Considering vacancy
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Fig. 6. (a) TEM (a) BF and (b) CDF image taken using a superlatice reflection (which is an extra spot either at 1/3{1 1 0}B2 or at 2/3{1 1 0}B2 in (c), showing the x domains and particle clusters in the rough cell-wall regions, decorating the smooth x cells. SADPs along [0 0 0 1]x//[1 1 1]B2 directions taken from (c) x domains and particle clusters, (d) smooth x cell A and (e) smooth x cell B, respectively.
involvement, there is a question that needs to be answered: where do these vacancies come from? The alloy studied was not subjected to any quench treatment, only cooling during the ingot-casting process, using the plasma arc cold hearth technique. It is possible that these vacancies were introduced during solidification through the cold Cu hearth, which resulted in a fast cooling of the alloy. It is also possible that these vacancies are structural vacancies formed within the B2 phase with more being formed in the Nb + Zr-enriched regions. TEM–EDX microanalysis was conducted to determine the chemical composition of the major phases indicated in Fig. 2a. The results are listed in Table 2. It is interesting to note that the rough cell-wall regions containing D88-type x show higher concentration of Zr and Nb than the smooth cell regions consisting of B82-type x: 24% more Zr and 10% more Nb are found to partition in the rough cell-wall regions than in the smooth cell regions. The segregation of Zr + Nb refractory elements on such a fine-scale (1– 2 lm across) has not been reported previously in TiAl.
The mechanism responsible for this fine-scale segregation is not understood clearly. Spinodal decomposition was suggested for the nanometer-scale phase separation observed in rapidly solidified Zr–Al and Zr3Al–Nb alloys [18,26]. On the other hand, microsegregation of refractory elements did occur in highly alloyed TiAl alloys due to dendrite coring effect, which generated a white-coloured network enriched in refractory elements as a result of columnar dendrite arm growth [34,39]. However, the segregation scale was bigger than the present case: the distance between neighbouring dendrite cores was normally 10–20 lm. Banerjee and Cahn [18] reported a fine-cell structure in their alloys. The formation of a fine-cell structure (0.5 lm in size) in a Zr–Al alloy was caused by a local cellular solidification, which pushed high-melting-point solutes to cell boundaries. No matter which mechanism was in control, the retained b(B2) in the 4Zr–4Nb TiAl alloy was divided into a cell structure, with Zr and Nb being segregated into the cell wall regions. As a consequence of the microsegregation, preferential formation of the D88-type x occurred.
Table 2 TEM-EDX results showing the composition of the microstructural features in alloy 4–4–1 after 1260 °C HIPping Microstructure features
Ti
Al
Nb
Zr
Si
Equiaxed c Equiaxed a2 Cells in retained b(B2) Cell-walls in retained b(B2)
44.6 ± 1.5 57.1 ± 1.5 55.5 ± 2.3 51.6 ± 2.3
46.7 ± 2.2 34.7 ± 1.6 33.8 ± 2.1 36.2 ± 1.5
4.2 ± 0.3 4.3 ± 0.1 6.3 ± 0.5 6.9 ± 0.2
4.4 ± 0.4 3.6 ± 0.1 4.1 ± 0.3 5.1 ± 0.5
0.1 ± 0.1 0.3 ± 0.1 0.3 ± 0.1 0.2 ± 0.1
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The enrichment in Zr and Nb is found to be coupled with slightly higher Al but lower Ti in the rough cell-wall regions than in the smooth cell regions. From the EDX analysis, it is also noted that Zr tends to concentrate into equiaxed c grains, rather than into the retained a2, while Nb is preferentially concentrated into retained b(B2) (both rough and smooth regions) rather than a2 and c. These results are consistent with those reported by Kainnuma et al. [40]. The fact that the selective formation of the vacancyrelated D88-type x coincides with the Zr + Nb enrichment in the cell-wall regions would encourage one to consider that there is a strong relationship between Zr + Nb atoms and vacancies. In fact, Zr was recognized to be a good vacancy-trapper. The binding energy between a Zr atom and a vacancy in aluminium was measured to be 0.24 ± 0.02 eV [41], while the binding energy between an Nb atom and a vacancy in an aluminium-rich alloy was measured to be 0.18 ± 0.02 eV [42]. Both values are high enough to allow Zr and Nb atoms to trap vacancies, leading to the formation of Zr/Nb–v (vacancy) pairs. Consequently, Zr/Nb–v pairs rather than Zr/Nb atoms would occupy Ti sites, providing the vacancies to meet the stoichiometric requirement for the formation of the D88 structure. Therefore, it is reasonable to believe that the higher concentration of Zr + Nb in the narrow cell-wall regions is the necessary composition which not only ensures the
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retention of b phase, but also guarantees the capture of vacancies required for the formation of the vacancy-related D88-type x. Following the aforementioned TEM examination and EDX analysis, it now becomes clear that the cellular morphology of retained b(B2) observed in 4–4–1 alloy is simply caused by the heterogeneous formation of a different type of x, which is controlled by a Zr + Nb-enriched network. The small D88-type x domains and particles precipitated favourably in rough cell-wall regions, while the formation of B82-type x dominated individual smooth cells. The solid b(B2) block was therefore subdivided into individual cells/ domains. 3.3. Ordered x phases after long-term ageing The occurrence of a different x-type phase (D88 structure), which is apparently a more ordered derivative of the B82 structure, has been reported previously in Ti–Al– Nb and Zr–Al–Nb systems [5,21,25,26]. Bendersky et al. [20] and Tewari et al. [26] have discussed the evolution of different x-type phases from the bcc phase and have constructed symmetry trees to describe the symmetry relationship among all the observed transformation paths. According to the work on the Zr3Al–Nb ternary system carried out by Tewari et al., the x-type B82 structure was
Fig. 7. TEM (a) BF and (b) CDF micrographs taken from alloy 4–4–1 after 1000 h ageing at 700 °C showing the precipitation of fine D88-type x particles from the B82-x cell matrix. (c) SADPs taken from the region A correspond both to the D88-x and to B82-x phases (d) SADP from region B corresponds to the B82-x phase only.
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3.4. Mechanical behaviour of alloy 4–4–1 The tensile properties of the 4–4–1 alloy, tested at both room temperature and 700 °C, are shown in Fig. 8. The 4– 4–1 alloy shows a relatively high tensile strength since its grain size is effectively refined. The 0.1% proof stress is higher than the 600 MPa level at room temperature, but the alloy is found to be very brittle. All the ductility values obtained at both room temperature and 700 °C are <0.5%. The value of the elongation to failure at room temperature is too low to allow even a 0.2% proof stress value to be measured. It is well known that the formation of x may cause brittleness in conventional b-Ti alloys owing to its inherent brittleness which results from restricted slip [9,13]. However, the ductility obtained is significantly lower than that in a Ti–44Al–8Nb–1B alloy which had a similar grain size but lower volume fraction (3 vol.%) of B2 + x phases [34,35]. This tends to suggest that the significant retention of ordered B2 (9 vol.%) and the heterogeneous formation of ordered x phase from it may be responsible for the high degree of brittleness of the 4Zr– 4Nb-containing TiAl alloy. The heterogeneous precipita-
PS 0.1%
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found to be a more stable phase than the x-type D88 structure in the Zr3Al–Nb system [26]. Their work showed that ageing at 950 °C for 600 s can actually convert the preexisting D88 particles into B82 particles. However, the work on a Ti–37.5Al–20Nb alloy carried out by Bendersky et al. revealed a reverse transformation path, in which the x-type D88 precipitates were found as spherical (ellipsoidal) particles in a B82 matrix after prolonged ageing (18 days) at 700 °C [21]. The present work on the 4Zr–4Nb-containing TiAl alloy confirms the transformation path suggested by Bendersky et al. [20]. Fig. 7a and b show BF and CDF micrographs taken from a sample which was subjected to ageing at 700 °C for 1000 h. Significant precipitation of very fine particles occurred in the smooth B82-x cell matrix after this long-term ageing. The precipitates were distributed in high density and were very fine in size (5–10 nm). The diffraction pattern in Fig. 7c was taken from region A, as indicated in Fig. 7a. Besides the reflections from the B82-x phase, reflections from the D88-x phase – the same as those observed in cell-wall regions – are observed. This clearly indicates that these fine D88-x particles were transformed from the preexisting B82-x matrix (i.e. B82 ? D88). Therefore, the structure transformation path in the retained b cells after ageing at 700 °C for 1000 h can be summarized as: b ? B2 ? B82 ? D88. Interestingly, it was found that some D88-x precipitate-free zones (PFZs) were present. The PFZs were normally located along b(B2) grain boundaries and at cell interfaces. The corresponding SADP taken from one of the PFZs (B in Fig. 7a) still showed a B82-x structure (see Fig. 7d). This behaviour was attributed to the easier escape of vacancies in the near-boundary regions during ageing since the formation of D88-x structure necessitates the retention of vacancies.
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100 0
0
Fig. 8. (a) Room temperature and (b) 700 °C tensile results for alloy 4–4–1 after 0 h (as-HIPped) and 1000 h ageing at 700 °C.
tion of ordered and fully mature x, especially the congregated D88-type x domains and particles formed as a network surrounding the relatively smooth B82-x cells, is assumed to be very detrimental to the ductility of the TiAl-based alloy. Fig. 8 also demonstrates that prolonged ageing for 1000 h at 700 °C causes little change to the strength and ductility of the thermally exposed TiAl alloy. This seems to suggest that the subsequent conversion of well dispersed D88-type x fine particles in the B82-type x matrix does not further negatively influence the ductility. S–N fatigue curves produced at room temperature are shown in Fig. 9 for alloy 4–4–1. The fatigue limit, rFL, is typically defined by the value of maximum stress, rmax, at run-out (P107 cycles). The most important feature related to the near-lamellar alloy is that it always exhibits relatively flat S–N curves. The difference in loading stress between the life at run-out (P107 cycles) and the life measured for failure at 104 cycles is only 10 MPa for the TiAl alloy, indicative of a high degree of sensitivity to the stress applied. The heterogeneous formation of the D88-type x particle network and B82-x cells has been found to provide favoured sites for microcrack initiation under cyclic loading. Fig. 10, a BSE micrograph taken from near the fracture region of an S–N fatigue sample (as-HIPped alloy 4– 4–1) tested at rmax = 480 MPa, shows that many second-
Z.W. Huang / Acta Materialia 56 (2008) 1689–1700
Maximum stress (MPa)
550
500
450
400 0h 1000h 350 1e+3
1e+4
Denotes run out 1e+5
1e+6
1e+7
1e+8
Number of cycles to failure Fig. 9. S–N curves of alloy 4–4–1 after 0 h (as-HIPped) and 1000 h ageing at 700 °C.
Fig. 10. BSE image taken from the S–N fatigue sample of as-HIPped alloy 4–4–1 tested at rmax = 480 MPa, showing that the secondary cracks near the fracture (bottom) occur mainly on segregated B2 + x blocks.
ary cracks occur in these B2 + x blocks during testing, whereas neighbouring a2 + c lamellae tend to deform. It has also been found that prolonged ageing at 700 °C for 1000 h does not further detrimentally influence the brittle nature of the alloy through the widespread production of fine T88-type x precipitates inside the cells. As seen in Fig. 9, a rFL value of 450 MPa was obtained for alloy 4–4–1 in the as-HIPped condition, while a rFL value of 460 MPa was obtained for the alloy after ageing at 700 °C for 1000 h. 4. Conclusions 1. A considerable amount of b-phase (9 vol.%) was retained in a Ti–44Al–4Zr–4Nb–0.2Si–1B ingot alloy after HIPping. The retained b was segregated at colony boundaries, ordered to B2 structure and showed a characteristic cellular morphology.
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2. Heterogeneous precipitation of x phase from B2 was found to occur readily after HIPping: B2 ? x with the B82-structure in cell regions and B2 ? x with the D88structure in cell-wall regions. Congregated D88-x domains and particles formed as a network surrounding the well-developed B82-x cells. 3. EDX analysis has shown that the heterogeneous formation of x phase throughout B2 blocks is caused by a heterogeneous distribution of Zr + Nb elements in the B2 cell structure. The combined Zr and Nb enrichment in the cell-wall regions is thought to have played an important role in stabilizing vacancies, promoting the vacancy-related D88-type x structure. 4. A B82-x ? D88-x transformation was found to occur inside the cell regions after ageing for 1000 h at 700 °C. Fine dispersed D88-type x particles precipitated extensively in the B82-x matrix. The structure transformation path in the retained b cells after ageing at 700 °C for 1000 h can be summarized as: b ? B2 ? B82-x? D88-x. 5. The 4Zr–4Nb-containing TiAl alloy is very brittle even though its grain size is much refined. The heterogeneous precipitation of ordered fully mature x, especially the congregated D88-type x domains and particles which form as a network across the retained grain, is assumed to be detrimental to the ductility of this alloy. 6. The precipitation of fine-dispersed D88-type x particles in the B82-x matrix after ageing does not further detrimentally influence the alloy. 7. The D88-x and B88-x phases that transformed from B2 showed not only different transformation paths and lattice structures, but also different chemical compositions, precipitate sizes and local distributions. Despite the differences between the two types of x phases, both the x structures form readily once ordered B2 transforms from retained b and they exhibit a mature and well crystalline lattice structure in the highly alloyed c-TiAlbased alloy. Acknowledgements This research was sponsored by the National Natural Science Foundation of China under Project No. 50471046. The author is also grateful to the Department of Metallurgy and Materials, The University of Birmingham, UK, for some experimental support. References [1] [2] [3] [4]
Cheng TT, Loretto MH. Acta Mater 1998;46:4801–19. Cheng TT, Willis MR, Jones IP. Intermetallics 1999;7:89–99. Kim Y-W. IRC workshop on TiAl, Birmingham, 1–3 May, 1996. Zhang B, Wang J, Wan X, Chen W. Scripta Metall Mater 1994;30:399–404. [5] Song H, Meng X, He L, Ye H, Zhang W, Chen G. Acta Metall Sinica 2001;37:804–9. [6] Huang ZW, Voice W, Bowen P. Mater Sci Eng A 2002;329– 331:435–45.
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[7] Huang ZW, Voice W, Bowen P. Scripta Mater 2003;48:79–84. [8] Frost PD, Parris WM, Hirsch LL, Doig JR, Schwartz CM. Trans Am Soc Metals 1954;46:231–56. [9] Silcock JM. Acta Metall 1958;6:481–93. [10] Hatt BA, Roberts JA. Acta Metall 1960;8:575–84. [11] Fontaine De D. Acta Metall 1970;18:275–9. [12] Dawson CW, Sass SL. Metall Trans 1970;1:2225–33. [13] William JC, Hickman BS, Markus HL. Metall Trans 1971;2: 1913–9. [14] Kuan TS, Sass SL. Acta Metall 1976;24:1053–9. [15] Pynn R. J. Phys. F: Metal Phys 1978;8:1–13. [16] Georgopoulos P, Cohen JB. Acta Metall 1981;29:1535–51. [17] Sikka SK, Vohra YK, Chidambaram R. Prog Mater Sci 1982;27:245–310. [18] Banerjee S, Cahn RW. Acta Metall 1983;31:1721–35. [19] Strychor R, Williams JC, Soffa WA. Metall Trans 1988;19:225–34. [20] Bendersky LA, Boettinger WJ, Burton BP, Biancaniello FS, Shoemaker CB. Acta Metall Mater 1990;38:931–43. [21] Bendersky LA, Burton BP, Boettinger WJ, Biancaniello FS. Scripta Metall Mater 1990;24:1541–6. [22] Chang CP, Loretto MH. Phil Mag A 1991;63:389–406. [23] Ezakp H, Morinaga M, Kato M, Yukawa N. Acta Metall Mater 1991;39:1755–61. [24] Shao G, Miodownik AP, Tsakiropoulos P. Phil Mag A 1995;71: 1389–408. [25] Banerjee S, Tewari R, Mukhopadhyay P. Prog Mater Sci 1997;42: 109–23.
[26] Tewari R, Mukhopadhyay P, Banerjee S, Bendersky LA. Acta Mater 1999;47:1307–23. [27] Shao G, Tsakiropoulos P. Acta Mater 2000;48:3671–85. [28] Fan Z, Miodownik AP. Acta Mater 1996;44:93–110. [29] Zhang XD, Wiezorek JMK, Baeslack III WA, Evans DJ, Fraser HL. Scripta Mater 1999;41:659–65. [30] Lin DJ, Chern JH, Ju CP. Biomaterials 2002;23:1723–30. [31] Shao, Nguyen-Manh D, Tsakiropoulos P. In: Kim Y-W, Dimiduk DM, Loretto MH, editors. Gamma titanium aluminides. Warrendale (PA): TMS; 1999. p. 155–9. [32] Froes FH, Bomberger HB. J Metals 1985;37:28–37. [33] Chen GL, Zhang WJ, Liu ZC, Li SJ, Kim Y-W. In: Kim Y-W, Dimiduk DM, Loretto MH, editors. Gamma Titanium Aluminides. Warrendale (PA): TMS; 1999. p. 371–80. [34] Huang ZW. Scripta Mater 2005;52:1021–5. [35] Huang ZW, Zhu DG. Intemetallics; in press. [36] Ho K, Quader MA, Lin F, Dodd RA. Scripta Metall 1977;11:1159–63. [37] Van Ommen AH, Waegemakers A, Moleman AC, Schatter H, Baker H. Acta Metall 1981;29:123–33. [38] Ho K, Dodd RA. Scripta Metall 1978;12:1055–8. [39] McCullough C, Valencia JJ, Levi CG, Mehrabian R. Mater Sci Eng A 1992;156:153–66. [40] Kainuma R, Fujita Y, Mitsui H, Ohnuma I, Ishida K. Intermetallics 2000;8:855–67. [41] Ozbilen S, Flower MH. Acta Metall 1989;37:2993–3000. [42] Raman KS. Scripta Metall 1971;5:791–4.
Acta Materialia 56 (2008) 1689–1700 www.elsevier.com/locate/actamat
Ordered x phases in a 4Zr–4Nb-containing TiAl-based alloy Z.W. Huang * School of Materials Science and Engineering, Southwest Jiaotong University, Chengdu, Sichuan 610031, PR China Department of Metallurgy and Materials, University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Received 28 September 2007; received in revised form 5 December 2007; accepted 7 December 2007 Available online 1 February 2008
Abstract A considerable amount of B2 phase with a cellular morphology is retained in a 4Zr–4Nb-containing TiAl-based alloy. Heterogeneous precipitation of ordered x from B2 is found to occur readily after HIPping: B2 ? x with the B82-structure in cell regions and B2 ? x with the D88-structure in cell-wall regions. Congregated D88-x domains and particles form as a network surrounding the well-developed B82-x cells. The heterogeneous formation of different x variants is caused by a heterogeneous distribution of Zr + Nb elements across B2, which plays an important role in stabilizing vacancies and promotes the formation of D88-x. Fine D88-x particles are also observed to precipitate from the B82-x cell matrix after ageing at 700 °C for 1000 h, showing a transformation path of b ? B2 ? B82-x ? D88-x for the aged cells. The heterogeneous formation of a D88-x network and B82-x cells is found to be detrimental to ductility and fatigue strength. A very brittle fine-grained TiAl alloy is produced as a result. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Titanium aluminides; TEM; Phase transformation; Heterogeneous nucleation of phase; Mechanical properties
1. Introduction In the development of high strength TiAl-based alloys, alloy design based on major alloying of the Ti-Al binary system has been applied to achieve overall improvements in properties. Alloys such as Ti–44Al–8X–1B, where X represents combinations of Nb, Zr, Ta and Hf, have been chosen for detailed studies [1,2]. These refractory elements are able to increase both high-temperature creep resistance and oxidation resistance. However, they are all b-phase stabilizers. A high level of such elements added to TiAl alloys may cause varied amounts of b-phase to be retained, thereby introducing an (a + b) phase field for the alloy system. This is expected to facilitate thermomechanical processing at elevated temperatures, since a grain growth can be restricted, particularly when the volume fraction of the bphase is high [3]. On the other hand, increasing evidence has indicated that the stabilized b-phase in highly alloyed *
Tel.: +44 121 414 3447; fax: +44 121 414 7468. E-mail address: [email protected].
TiAl alloys can be retained to room temperature. The retained b has always been observed as an ordered B2 structure which can readily decompose into the x phase during cooling to room temperature [1,2,4–7]. The formation of x-type phases has been extensively investigated in transition-metal-based body-centred cubic (bcc) alloys and is well documented in the literature [8– 27]. It is generally accepted that the transformation of b(B2) to x phase is a displacive process in which one pair of {2 2 2}b planes collapses to the intermediate position along one h1 1 1ib direction, leaving the next plane unaltered [10,11,17]. The {2 2 2}b plane collapse could be incomplete (partial collapse), in which case the pair of {2 2 2}b planes does not collapse into the double layer to produce an ideal hexagonal P6/mmm structure. Such a partial collapse produces a trigonal P3m1 structure [17,20–22,24]. The product phase was termed diffuse x if exhibiting diffuse x scattering/reflections, which deviates from the ideal crystalline x maxima along the line through (0 0 0) and (2 2 2)b [12,28–30], or termed a crystalline x if exhibiting discreet diffraction maxima at the position of
1359-6454/$34.00 Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2007.12.013
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ideal crystalline x lattice. The crystalline x can inherit B2 chemical order to become an ordered x, although it was still considered as an intermediate phase in the course of structural change [19–21]. More stabilized x-type products with further chemical ordering and fully crystalline lattice structures such as B82 and D88 were reported in various transition-metal-based alloys containing varied percentages of Al [18–21,25–27]. These x-type products are transformed either directly from parent b(B2) or from metastable intermediate x by a superimposition of displacive and replacive ordering reactions. The two x-type products often form after a prolonged ageing at elevated temperature and are therefore closer to the equilibrium composition and structure and always fully three-dimensionally developed in morphology. c-TiAl-based alloys of engineering importance are normally based on Ti–(44–49 at.%) Al with appropriate combination of alloying elements. Al was reported to increase the x-phase stability by lowering the valence electron density of B2, while the transition metals added may increase the valence electron density of B2 [31]. It would be very interesting to know how the B2 ? x transformation would be affected under the interaction between Al and transition metals. Research into x formation and transformation within multicomponent TiAl-based alloys has rarely been reported [4–7]. Detailed investigation is therefore needed for a clear understanding of the x phase in such types of TiAl alloys. The work should include the stability of the x phase at the early stage of formation, the lattice structure, the transformation path, the size and morphology of the precipitate, the transformation path during prolonged ageing at elevated temperature and the effects on mechanical properties. The x formation is well known to result in brittleness in conventional b-Ti alloys and other transition-metal-based bcc alloys [8,9,13,32]. The co-presence of B2 and x was also reported to be detrimental to high-Nb-containing TiAl-based alloys [33]. However, more research is needed to understand clearly how and to what degree the retained b(B2) and precipitated x affect the mechanical properties of TiAl alloys, especially when the volume fraction of B2 + x is relatively large (say >8%) and exists at grain boundaries and at triple junctions, which is often the case in highly alloyed TiAl alloys [34]. The present work was carried out on a cast ingot of cTiAl-based alloy containing 4Zr and 4Nb. The alloy was subjected to a hot isostatic pressing (HIPping). The HIPped billet was cooled slowly to room temperature. The as-HIPped alloy contained multiple phases at room temperature and a significant amounts of B2 and x phases were observed [1].
process at the IRC (Interdisciplinary Research Centre), Birmingham University. The cast ingot had been doublemelted and then HIPped at 1260 °C under a pressure of 150 MPa for 4 h at Southwest Jiaotong University, Chengdu. Microstructures were examined by back-scattered electron (BSE) microscopy using a JEOL 840A scanning electron microscope (SEM) operating at 20 kV. This was performed to reveal the size, shape and distribution of the constituent phases both before and after HIPping. Image analysis was carried out on BSE micrographs to quantify the size and volume fraction of constituent phases and structures using a mean linear-intercept method. The measurements were conducted on a minimum of 20 micrographs that were randomly chosen. The total area investigated was 5.4 mm2. Thin foils for transmission electron microscopy (TEM) were prepared by a twin-jet polishing with an electrolyte of 5 vol.% perchloric acid, 30 vol.% butan-1-ol and 65 vol.% methanol operating at 30 V and at a temperature of 25/30 °C. TEM was carried out using a JEOL 2000 FX microscope operating at 200 kV. Energy dispersive X-ray (EDX) microanalysis was performed on selected microstructural features using the TEM which was fitted with a LINK analytical EDX microanalyzer. The EDX analysis on each microstructural feature was performed at six locations. For all measurements conducted, mean values are presented with errors representing the uncertainty of the statistical analysis. The errors were assessed using 95% confidence limits, based on 2s/n1/2, where s is the standard deviation and n is the number of measurements. Tensile tests were conducted at both room temperature and 700 °C on a Zwick testing machine in air, at a strain rate of 0.7 104 s1. The tensile data presented are an average of three samples for room temperature tests and an average of two samples for 700 °C tests. S–N fatigue tests were carried out at room temperature in air. Specimens were tested in pure bending using an electromagnetic resonance testing machine at a frequency of 80 Hz under a stress ratio of R = 0.1 (where R = rmin/rmax; rmin and rmax are, respectively, the minimum and maximum stresses applied over the fatigue cycle). Testpieces of dimensions 80 10 10 mm3 were prepared using electro-discharge machining. Conventional grinding and polishing to 0.25 lm diamond were used to prepare the surface. A major loading span of 60 mm and a minor loading span of 20 mm were used. 3. Results and discussion
2. Experimental
3.1. Microstructure after ingot casting and subsequent HIPping
The alloy chosen for the present study was Ti–44Al– 4Zr–4Nb–0.2Si–1B (alloy 4–4–1, at.%). It was grain refined by adding 1% of boron and supplied as a cast ingot that was made using a PACH (plasma arc cold hearth) melting
Fig. 1a and b shows the multiphase microstructure of cast ingot alloy 4–4–1 before and after 4 h/1260 °C HIPping, respectively. The alloy after ingot-casting exhibited a near-lamellar microstructure, mixed with retained b(B2)
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Table 1 The volume fraction and size of the microstructural constituents in alloy 44-1 after ingot casting and 1260 °C HIPping
Size (lm) Volume fraction (%)
Fig. 1. Back-scattered SEM images showing the microstructure of (a) cast ingot and (b) after 4 h/1260 °C HIPping.
and a(a2) and equiaxed c. As reported previously, the phase transformation sequence for the 4Zr–4Nb-containing TiAl alloy was [2]: L!Lþb!b!aþb!aþbþc!aþc There was no single a phase field during the solidification process. After HIPping at 1260 °C, the alloy still exhibited a near-lamellar microstructure, consisting of lamellar colonies (70 vol.%) and equiaxed B2, a2 and c grains (30 vol.% in total). The lamellar colonies were measured to be 70 lm (±11.2 lm) in size. The retained B2 (11.5 ± 3.7 lm in size) and retained a2 (5.2 ± 1.2 lm in size) grains were observed mainly around lamellar colonies and c grain boundaries. The volume fractions and mean sizes (with the error range) of all the microstructural constituents determined by image analysis of BSE micrographs are summarized in Table 1. The microstructure after HIPping, in which considerable amounts of equiaxed B2, a2 and c grains remain with a2 + c lamellar structure, makes the 4–4–1 alloy quite different from other highly alloyed and grain-refined c-TiAl alloys such as Ti–44Al–8Nb–1B, Ti–44Al–4Nb–4Hf–1B and Ti–44Al–4Nb–4Ta–1B, all of
Lamellar colonies
Equiaxed a2
Equiaxed c
b(B2)
70.9 ± 11.2 71.2 ± 10.7
5.2 ± 1.2 4.2 ± 0.8
19.3 ± 4.8 15.7 ± 4.5
11.5 ± 3.7 8.9 ± 2.5
which show a fully lamellar microstructure, although small amounts of b are retained. Fig. 2a, a bright field (BF) TEM micrograph, shows all the major microstructural constituents obtained in 4–4–1 alloy after 1260 °C HIPping. The a2 + c lamellae (labelled as L) are the major constituents in this alloy. The corresponding selected area diffraction patterns (SADPs) from the a2 + c colonies are shown in Fig. 2b. A considerable amount (9 vol.%) of retained b-phase was also observed. The retained high temperature phase was always located at lamellar colony boundaries. The reason why the retained b segregates around lamellar colony boundary was explained in a previous paper, based on the b ? a phase transformation in which the solute redistribution factor k0 = Ca/ Cb > 1 [34]. Unlike the usual monolithic morphology of retained b observed in other highly alloyed TiAl-based alloys, such as Ti–46Al–5Nb–1W and Ti–44Al–8Nb–1B [6,35], the retained b(B2) in the present alloy exhibits a cellular morphology, consisting of many small cells (1–2 lm in size), between which are narrow cell-wall regions. Compared to the relatively smooth cell surface, the cell-wall regions look rough, owing to a congregation of small domain and particle clusters in narrow areas. SADPs taken from the cell structure show strong reflections corresponding to the ordered B2 lattice and always contain x reflections (Fig. 2c). The cell structure is therefore labelled as B2 + x in Fig. 2a. It is important to note that the B2 + x cell structure always keeps a general orientation relationship with their neighbouring a2 + c lamellar structure. This clearly indicates that both the transformed a2 + c lamellae and the retained B2 + x originate from a single b grain. Therefore, the general orientation relationships among the four phases shown in Fig. 2a can be defined as f1 1 1gc ==f0 0 0 1ga2 ==f1 1 0gB2 ==f1 1 2 0gx and h1 1 0ic ==h1 1 2 0ia2 ==h1 1 1>B2 ==h0 0 0 1ix A small amount of equiaxed a2 was also observed at room temperature. The diffraction pattern taken from an equiaxed a2 grain (Fig. 2d) shows that the equiaxed a2 has the same orientation as the a2 in the lamellar structure, indicating that the a2 is the remnant of the a grain after ordering and eutectoid transformation a ? a2 ? a2 + c. The retained a2 therefore shows the same lattice structure and orientation as its transformed counterpart. This clearly indicates that the solid phase transformation to a2 + c is
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Fig. 2. (a) TEM BF (bright field) micrograph showing a multiphase microstructure consisting of a2 + c lamellae (L), cellular B2 + x, equiaxed a2 and c phases. (b–e) The selected-area diffraction patterns taken from the indicated lamellae, B2 + x cell structure, retained a2 and c grain, respectively, with B ¼ h1 1 2 0ia2 ==h1 1 0ic in (b), B ¼ h1 1 1iB2 ==h0 0 0 1ix in (c), B ¼ h1 1 2 0ia2 in (d) and B ¼ h1 1 4ic in (e).
not complete in this highly alloyed TiAl alloy. Equiaxed c grains were found to be distributed extensively in the asHIPped alloy since the single a-phase field is replaced by a (a + b + c) phase field in alloy 4–4–1 [2]. Some of the equiaxed c grains were small in size and located close to retained B2 and a2, as indicated in Fig. 2a. They were independent of neighbouring a2 + c lamellae as well as
equi- axed B2 and a2 grains in terms of the orientation relationship (Fig. 2e). 3.2. Ordered x phases after HIPping The formation of x from B2 has been described as a displacive process in which one pair of (2 2 2)B2 planes
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collapses to the intermediate position along a h1 1 1iB2 direction [10,11]. Four variants of x are expected to form from the parent B2, corresponding to the interplanar collapse along each of the four h1 1 1iB2 directions. Fig. 3a is a centred dark field (CDF) image taken from the B2 + x cell structure with a 1 0 1 0x reflection, corresponding to a single variant of x (see Fig. 3b). All the illuminated cells therefore belong to one rotational variant of the x phase. It is observed that these x cells are actually composed of sub-micron-sized domains and these domains are separated by internal interfaces. The pure set of x diffraction spots along one of the h1 1 2 0ix zone axes shown in Fig. 3b indicates that the x inside the smooth cells/domains is fully transformed from B2. The precipitation of x throughout the retained B2 blocks was found not to be uniform in terms of the size and distribution of the precipitates. Fig. 4a and b, using both BF and CDF images, reveal that the rough cell-wall regions (e.g. B and D) are in fact composed of many small x domains and particle clusters. These congregated x domains and particles formed as a network surrounding the relatively smooth x cells (e.g. cell C). SADPs along the h1 1 0i>B2 direction obtained from the rough region B (Fig. 4c) show strong reflections corresponding to B2, indicating that some B2 was left in the narrow cell-wall regions after the B2 ? x transformation. The finer diffraction maxima in the complicated pattern correspond to various variants of the precipitated x. In contrast to the complex condition of the cell-wall region, individual x cells are single crystal in structure. As shown in Fig. 4d and e, the SADP taken either from cell A or cell C shows a diffraction pattern of one variant only: ½1 1 0 2x for cell A and ½1 1 2 0x for cell C. The micrographs shown in Fig. 4 clearly indicate that the x phase in the as-HIPped alloy 4–4–1 was precipitated
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non-uniformly across retained b blocks. While the individual x cells which showed a single crystalline structure had a size range of 1–2 lm, the small x domains and particles were <100 nm. These small domains and particles congregated into narrow areas between the relatively smooth x cells, co-existed with untransformed B2 and were connected to form a network across the retained grain. Fig. 5a–c shows a set of SADPs (zone axes using cubic indices are [1 0 0]B2, [1 1 1]B2 and [1 1 2]B2, respectively) obtained specifically from individual smooth x cells in the as-HIPped 4–4–1 alloy. The x cells do not generate any diffuse scatterings in their diffraction patterns. Discreet and sharp x spots without any diffraction streaking and any shift along the line through (0 0 0) and (2 2 2)B2 can be observed, suggesting that the precipitated x phase had already reached a mature, fully crystalline condition after HIPping. This is different from conventional Ti-based alloys [22,28–30], in which the maxima corresponding to the x phase exhibited short streaks and shift from the ideal crystalline position. This is also different from the Nb + Wcontaining TiAl-based alloy studied previously. The diffraction maxima corresponding to an early type of x phase in this alloy showed short streaks perpendicular to the ½2 2 2B2 and ½2 2 2B2 directions, although no shift of diffraction maxima was observed [6]. It appears that once the high-temperature b-phase in the Zr + Nb-containing TiAl alloy cooled down to room temperature after HIPping, not only was the collapse of the double layers in B2 structure complete, leading to the formation of an ideal crystalline x, but also the precipitated x phase was already three-dimensionally developed, as imaged in the CDF micrographs of Figs. 3a and 4b. Further TEM work on the x cells was conducted to determine which lattice structure and space group the
Fig. 3. (a) TEM CDF image taken with a 1 0 1 0x reflection ðB ¼ h1 2 1 0ix Þ showing only one x variant in the cell structure. (b) The corresponding diffraction pattern from this variant. Note that the x cells were divided into submicron domains by internal interfaces.
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Fig. 4. (a) TEM BF and (b) CDF images showing the x domains and particle clusters precipitated favourably in the rough cell-wall regions B and D, forming a network around the relatively smooth x cell C. (c) SADPs taken from region B along the beam direction of [1 1 0]B2. (d) SADP from cell A ðB ¼ h1 1 0 2ix Þ and (e) from cell C ðB ¼ h1 1 2 0ix .
x belongs to: trigonal with P3m1 or B82 hexagonal with P63/mmc. The smallest diffraction aperture was used on a single x cell to produce a pure set of x spots in one of h1 1 2 0ix zone axes. Tilting along a row of 0 0 0lx reflections was performed to allow this row to be excited, with the aim of eliminating roots for double diffraction. The insert in Fig. 5c shows a row of excited microdiffraction spots obtained from a single x cell, revealing the distinction of odd reflections. The distinction of odd 0 0 0lx reflection suggests that the x-phase in these smooth cells had transformed into a hexagonal structure rather than a trigonal symmetry. The formation of the more stabilized hexagonal structure was probably caused by the slow cooling which is a normal practice after HIPping. Fig. 5d–f shows another set of SADPs (same zone axes [1 0 0]B2, [1 1 1]B2 and [1 1 0]B2), obtained specifically from the rough cell-wall regions composed of small x domains and particle clusters in the same alloy. Analysis of the two sets of SADPs indicates significant differences, especially in the zone axes [1 0 0]B2 and [1 1 1]B2, where extra spots exist at 1/3{1 1 0}B2 and 2/3{1 1 0}B2. The SADPs in the [1 0 0]B2 and [1 1 1]B2 directions are therefore more complicated than those taken from the smooth cell regions. Those extra reflections in Fig. 5d and e look like superlattice spots. The corresponding microstructure of small x
domains and particle clusters, as shown in Fig. 6a and b, is distinctly different from smooth x cells and is imaged in the CDF (Fig. 6b) separately from those smooth x cells, using the reflections in Fig. 6c which are additional to those showing in Fig. 6d and e (i.e. those extra spots at 1/ 3{1 1 0}B2 and 2/3{1 1 0}B2 in Fig. 6c). It has been established that the x diffraction patterns from the rough cellwall regions correspond to a D88 structure [21,25,26], whereas those x patterns from smooth cells correspond to a B82 structure [18–20]. From the SADPs in Fig. 5d–f, it was found that the lattice parameters of the D88 unit cell are related to the B2 structure as: aD = 61/2aB2 and cD = 31/2aB2 (subscript ‘‘D” denotes D88-type x phase). For the B82-type x-phase, the lattice parameters are: aB = 21/ 2 aB2 and cB = 31/2aB2 (subscript ‘‘B” denotes B82-type x phase). The B82 structure was reported to be an ideally crystalline x-derivative formed by a complete collapse of one pair of {2 2 2}B2 planes, the double layers are therefore not rumpled [18–21]. As confirmed in Fig. 5c, the x-type B82 phase shows a hexagonal symmetry (P63/mmc space group) rather than a trigonal symmetry with P3m1 space group. From the SADPs in Fig. 5d–f, the D88-type structure is also an ideally crystalline x-derivative formed by the same complete collapse of the {2 2 2}B2 planes. The orientation relationship between the D88-type x and parent B2 can be established:
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Fig. 5. (a–c) SADPs taken from the individual w cells, compared to those in (d–e) taken from the rough cell-wall regions. Note those additional superlattice spots in (d–e). The zone axes using cubic indices are: (a and d) h0 0 1i, (b and e) h1 1 1i and (c and f) h1 1 2i, respectively.
h1 1 1hB2 ==h0 0 0 1iD and f1 1 0gB2 ==f1 0 1 0gD This is different from the orientation relationship between the B82-type x and parent B2, which has previously been described [18,20] as h1 1 1iB2 ==h0 0 0 1iB and f1 1 0gB2 ==f1 1 2 0gB The different orientation relationship plane pair between D88 x B2 and B82 x B2 suggests that chemical ordering occurs in the {000l}x basal planes. Moreover, an ordered array of vacancies was reported to be present in the base planes of the D88-type x structure. The main difference between the B82 and D88 structures lies in whether or not the Wyckoff positions 2(b) are occupied. Two vacancies for 18 atom positions per unit cell are required for the formation of the vacancy-ordered D88 structure [21,25,26]. Based on this, the present observation of the D88-type x-phase is indicative of the retention of a high concentration of vacancies in the retained B2 phase region. However, there is no direct evidence for vacancy retention in the present TiAl alloy, although high equilibrium concentrations of vacancies in ordered bcc structures have been reported [36–38]. Tewari and co-workers [26] found that the D88 structure can readily form in Zr3Al–Nb alloys, but not in b alloys close to Zr3Al compositions. Based on this fact, they suggested that Nb had an important role to play in stabilizing
vacancies, hence promoting the D88 structure. However, a high percentage of Nb alone did not promote the formation of the D88 structure in TiAl-based alloys. For instance, the formation of D88-type x phase in Ti–45Al– 10Nb was found to be much less significant [5] than in the currently studied Ti–44Al–4Nb–4Zr–0.2Si–1B and no D88-type x phase was detected in Ti–44Al–8Nb–1B alloy, even after long-term ageing at 700 °C [35]. The present results obtained from the 4Zr + 4Nb-containing TiAl alloy tend to suggest that the combination of Zr and Nb has played a more important role in stabilizing vacancies in the retained B2 structure, therefore promoting the formation of the D88-type x structure. The present TEM examination reveals that the formation of the D88-type x domains and particles occurs only in rough cell-wall regions, which form as a network surrounding the relatively smooth B82-type x cells. This appears to indicate that the retention of vacancies occurs in a non-uniform fashion in the retained b(B2) blocks, therefore promoting the formation of the vacancy-related D88-type x in the narrow cell-wall regions as a result. The retention of vacancies in the smooth cells is assumed to be less significant since only the B82-type x phase formed there. Consequently, the retained b(B2) blocks exhibit a picture of heterogeneous x-precipitation in terms of precipitate type, size and distribution. Considering vacancy
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Fig. 6. (a) TEM (a) BF and (b) CDF image taken using a superlatice reflection (which is an extra spot either at 1/3{1 1 0}B2 or at 2/3{1 1 0}B2 in (c), showing the x domains and particle clusters in the rough cell-wall regions, decorating the smooth x cells. SADPs along [0 0 0 1]x//[1 1 1]B2 directions taken from (c) x domains and particle clusters, (d) smooth x cell A and (e) smooth x cell B, respectively.
involvement, there is a question that needs to be answered: where do these vacancies come from? The alloy studied was not subjected to any quench treatment, only cooling during the ingot-casting process, using the plasma arc cold hearth technique. It is possible that these vacancies were introduced during solidification through the cold Cu hearth, which resulted in a fast cooling of the alloy. It is also possible that these vacancies are structural vacancies formed within the B2 phase with more being formed in the Nb + Zr-enriched regions. TEM–EDX microanalysis was conducted to determine the chemical composition of the major phases indicated in Fig. 2a. The results are listed in Table 2. It is interesting to note that the rough cell-wall regions containing D88-type x show higher concentration of Zr and Nb than the smooth cell regions consisting of B82-type x: 24% more Zr and 10% more Nb are found to partition in the rough cell-wall regions than in the smooth cell regions. The segregation of Zr + Nb refractory elements on such a fine-scale (1– 2 lm across) has not been reported previously in TiAl.
The mechanism responsible for this fine-scale segregation is not understood clearly. Spinodal decomposition was suggested for the nanometer-scale phase separation observed in rapidly solidified Zr–Al and Zr3Al–Nb alloys [18,26]. On the other hand, microsegregation of refractory elements did occur in highly alloyed TiAl alloys due to dendrite coring effect, which generated a white-coloured network enriched in refractory elements as a result of columnar dendrite arm growth [34,39]. However, the segregation scale was bigger than the present case: the distance between neighbouring dendrite cores was normally 10–20 lm. Banerjee and Cahn [18] reported a fine-cell structure in their alloys. The formation of a fine-cell structure (0.5 lm in size) in a Zr–Al alloy was caused by a local cellular solidification, which pushed high-melting-point solutes to cell boundaries. No matter which mechanism was in control, the retained b(B2) in the 4Zr–4Nb TiAl alloy was divided into a cell structure, with Zr and Nb being segregated into the cell wall regions. As a consequence of the microsegregation, preferential formation of the D88-type x occurred.
Table 2 TEM-EDX results showing the composition of the microstructural features in alloy 4–4–1 after 1260 °C HIPping Microstructure features
Ti
Al
Nb
Zr
Si
Equiaxed c Equiaxed a2 Cells in retained b(B2) Cell-walls in retained b(B2)
44.6 ± 1.5 57.1 ± 1.5 55.5 ± 2.3 51.6 ± 2.3
46.7 ± 2.2 34.7 ± 1.6 33.8 ± 2.1 36.2 ± 1.5
4.2 ± 0.3 4.3 ± 0.1 6.3 ± 0.5 6.9 ± 0.2
4.4 ± 0.4 3.6 ± 0.1 4.1 ± 0.3 5.1 ± 0.5
0.1 ± 0.1 0.3 ± 0.1 0.3 ± 0.1 0.2 ± 0.1
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The enrichment in Zr and Nb is found to be coupled with slightly higher Al but lower Ti in the rough cell-wall regions than in the smooth cell regions. From the EDX analysis, it is also noted that Zr tends to concentrate into equiaxed c grains, rather than into the retained a2, while Nb is preferentially concentrated into retained b(B2) (both rough and smooth regions) rather than a2 and c. These results are consistent with those reported by Kainnuma et al. [40]. The fact that the selective formation of the vacancyrelated D88-type x coincides with the Zr + Nb enrichment in the cell-wall regions would encourage one to consider that there is a strong relationship between Zr + Nb atoms and vacancies. In fact, Zr was recognized to be a good vacancy-trapper. The binding energy between a Zr atom and a vacancy in aluminium was measured to be 0.24 ± 0.02 eV [41], while the binding energy between an Nb atom and a vacancy in an aluminium-rich alloy was measured to be 0.18 ± 0.02 eV [42]. Both values are high enough to allow Zr and Nb atoms to trap vacancies, leading to the formation of Zr/Nb–v (vacancy) pairs. Consequently, Zr/Nb–v pairs rather than Zr/Nb atoms would occupy Ti sites, providing the vacancies to meet the stoichiometric requirement for the formation of the D88 structure. Therefore, it is reasonable to believe that the higher concentration of Zr + Nb in the narrow cell-wall regions is the necessary composition which not only ensures the
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retention of b phase, but also guarantees the capture of vacancies required for the formation of the vacancy-related D88-type x. Following the aforementioned TEM examination and EDX analysis, it now becomes clear that the cellular morphology of retained b(B2) observed in 4–4–1 alloy is simply caused by the heterogeneous formation of a different type of x, which is controlled by a Zr + Nb-enriched network. The small D88-type x domains and particles precipitated favourably in rough cell-wall regions, while the formation of B82-type x dominated individual smooth cells. The solid b(B2) block was therefore subdivided into individual cells/ domains. 3.3. Ordered x phases after long-term ageing The occurrence of a different x-type phase (D88 structure), which is apparently a more ordered derivative of the B82 structure, has been reported previously in Ti–Al– Nb and Zr–Al–Nb systems [5,21,25,26]. Bendersky et al. [20] and Tewari et al. [26] have discussed the evolution of different x-type phases from the bcc phase and have constructed symmetry trees to describe the symmetry relationship among all the observed transformation paths. According to the work on the Zr3Al–Nb ternary system carried out by Tewari et al., the x-type B82 structure was
Fig. 7. TEM (a) BF and (b) CDF micrographs taken from alloy 4–4–1 after 1000 h ageing at 700 °C showing the precipitation of fine D88-type x particles from the B82-x cell matrix. (c) SADPs taken from the region A correspond both to the D88-x and to B82-x phases (d) SADP from region B corresponds to the B82-x phase only.
Z.W. Huang / Acta Materialia 56 (2008) 1689–1700
3.4. Mechanical behaviour of alloy 4–4–1 The tensile properties of the 4–4–1 alloy, tested at both room temperature and 700 °C, are shown in Fig. 8. The 4– 4–1 alloy shows a relatively high tensile strength since its grain size is effectively refined. The 0.1% proof stress is higher than the 600 MPa level at room temperature, but the alloy is found to be very brittle. All the ductility values obtained at both room temperature and 700 °C are <0.5%. The value of the elongation to failure at room temperature is too low to allow even a 0.2% proof stress value to be measured. It is well known that the formation of x may cause brittleness in conventional b-Ti alloys owing to its inherent brittleness which results from restricted slip [9,13]. However, the ductility obtained is significantly lower than that in a Ti–44Al–8Nb–1B alloy which had a similar grain size but lower volume fraction (3 vol.%) of B2 + x phases [34,35]. This tends to suggest that the significant retention of ordered B2 (9 vol.%) and the heterogeneous formation of ordered x phase from it may be responsible for the high degree of brittleness of the 4Zr– 4Nb-containing TiAl alloy. The heterogeneous precipita-
PS 0.1%
TS
El%
700
665
652
634
621
1.5
600 500
1
400 300
0.5
200 100
0.2
0.15
0h
1000 h
0
0
b
PS 0.2%
TS
El%
2
800 700 600
635
627 563
1.5
554
500 1
400 300 200
Elongation to failure (%)
2
800
0.44
0.42
0h
1000 h
0.5
Elongation to failure (%)
a Tensile stress (MPa)
found to be a more stable phase than the x-type D88 structure in the Zr3Al–Nb system [26]. Their work showed that ageing at 950 °C for 600 s can actually convert the preexisting D88 particles into B82 particles. However, the work on a Ti–37.5Al–20Nb alloy carried out by Bendersky et al. revealed a reverse transformation path, in which the x-type D88 precipitates were found as spherical (ellipsoidal) particles in a B82 matrix after prolonged ageing (18 days) at 700 °C [21]. The present work on the 4Zr–4Nb-containing TiAl alloy confirms the transformation path suggested by Bendersky et al. [20]. Fig. 7a and b show BF and CDF micrographs taken from a sample which was subjected to ageing at 700 °C for 1000 h. Significant precipitation of very fine particles occurred in the smooth B82-x cell matrix after this long-term ageing. The precipitates were distributed in high density and were very fine in size (5–10 nm). The diffraction pattern in Fig. 7c was taken from region A, as indicated in Fig. 7a. Besides the reflections from the B82-x phase, reflections from the D88-x phase – the same as those observed in cell-wall regions – are observed. This clearly indicates that these fine D88-x particles were transformed from the preexisting B82-x matrix (i.e. B82 ? D88). Therefore, the structure transformation path in the retained b cells after ageing at 700 °C for 1000 h can be summarized as: b ? B2 ? B82 ? D88. Interestingly, it was found that some D88-x precipitate-free zones (PFZs) were present. The PFZs were normally located along b(B2) grain boundaries and at cell interfaces. The corresponding SADP taken from one of the PFZs (B in Fig. 7a) still showed a B82-x structure (see Fig. 7d). This behaviour was attributed to the easier escape of vacancies in the near-boundary regions during ageing since the formation of D88-x structure necessitates the retention of vacancies.
Tensile stress (MPa)
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100 0
0
Fig. 8. (a) Room temperature and (b) 700 °C tensile results for alloy 4–4–1 after 0 h (as-HIPped) and 1000 h ageing at 700 °C.
tion of ordered and fully mature x, especially the congregated D88-type x domains and particles formed as a network surrounding the relatively smooth B82-x cells, is assumed to be very detrimental to the ductility of the TiAl-based alloy. Fig. 8 also demonstrates that prolonged ageing for 1000 h at 700 °C causes little change to the strength and ductility of the thermally exposed TiAl alloy. This seems to suggest that the subsequent conversion of well dispersed D88-type x fine particles in the B82-type x matrix does not further negatively influence the ductility. S–N fatigue curves produced at room temperature are shown in Fig. 9 for alloy 4–4–1. The fatigue limit, rFL, is typically defined by the value of maximum stress, rmax, at run-out (P107 cycles). The most important feature related to the near-lamellar alloy is that it always exhibits relatively flat S–N curves. The difference in loading stress between the life at run-out (P107 cycles) and the life measured for failure at 104 cycles is only 10 MPa for the TiAl alloy, indicative of a high degree of sensitivity to the stress applied. The heterogeneous formation of the D88-type x particle network and B82-x cells has been found to provide favoured sites for microcrack initiation under cyclic loading. Fig. 10, a BSE micrograph taken from near the fracture region of an S–N fatigue sample (as-HIPped alloy 4– 4–1) tested at rmax = 480 MPa, shows that many second-
Z.W. Huang / Acta Materialia 56 (2008) 1689–1700
Maximum stress (MPa)
550
500
450
400 0h 1000h 350 1e+3
1e+4
Denotes run out 1e+5
1e+6
1e+7
1e+8
Number of cycles to failure Fig. 9. S–N curves of alloy 4–4–1 after 0 h (as-HIPped) and 1000 h ageing at 700 °C.
Fig. 10. BSE image taken from the S–N fatigue sample of as-HIPped alloy 4–4–1 tested at rmax = 480 MPa, showing that the secondary cracks near the fracture (bottom) occur mainly on segregated B2 + x blocks.
ary cracks occur in these B2 + x blocks during testing, whereas neighbouring a2 + c lamellae tend to deform. It has also been found that prolonged ageing at 700 °C for 1000 h does not further detrimentally influence the brittle nature of the alloy through the widespread production of fine T88-type x precipitates inside the cells. As seen in Fig. 9, a rFL value of 450 MPa was obtained for alloy 4–4–1 in the as-HIPped condition, while a rFL value of 460 MPa was obtained for the alloy after ageing at 700 °C for 1000 h. 4. Conclusions 1. A considerable amount of b-phase (9 vol.%) was retained in a Ti–44Al–4Zr–4Nb–0.2Si–1B ingot alloy after HIPping. The retained b was segregated at colony boundaries, ordered to B2 structure and showed a characteristic cellular morphology.
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2. Heterogeneous precipitation of x phase from B2 was found to occur readily after HIPping: B2 ? x with the B82-structure in cell regions and B2 ? x with the D88structure in cell-wall regions. Congregated D88-x domains and particles formed as a network surrounding the well-developed B82-x cells. 3. EDX analysis has shown that the heterogeneous formation of x phase throughout B2 blocks is caused by a heterogeneous distribution of Zr + Nb elements in the B2 cell structure. The combined Zr and Nb enrichment in the cell-wall regions is thought to have played an important role in stabilizing vacancies, promoting the vacancy-related D88-type x structure. 4. A B82-x ? D88-x transformation was found to occur inside the cell regions after ageing for 1000 h at 700 °C. Fine dispersed D88-type x particles precipitated extensively in the B82-x matrix. The structure transformation path in the retained b cells after ageing at 700 °C for 1000 h can be summarized as: b ? B2 ? B82-x? D88-x. 5. The 4Zr–4Nb-containing TiAl alloy is very brittle even though its grain size is much refined. The heterogeneous precipitation of ordered fully mature x, especially the congregated D88-type x domains and particles which form as a network across the retained grain, is assumed to be detrimental to the ductility of this alloy. 6. The precipitation of fine-dispersed D88-type x particles in the B82-x matrix after ageing does not further detrimentally influence the alloy. 7. The D88-x and B88-x phases that transformed from B2 showed not only different transformation paths and lattice structures, but also different chemical compositions, precipitate sizes and local distributions. Despite the differences between the two types of x phases, both the x structures form readily once ordered B2 transforms from retained b and they exhibit a mature and well crystalline lattice structure in the highly alloyed c-TiAlbased alloy. Acknowledgements This research was sponsored by the National Natural Science Foundation of China under Project No. 50471046. The author is also grateful to the Department of Metallurgy and Materials, The University of Birmingham, UK, for some experimental support. References [1] [2] [3] [4]
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