- Email: [email protected]
The influence of partitioning on the growth of intragranular α in near-β Ti alloys
Journal of Alloys and Compounds 643 (2015) 212–222
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
The influence of partitioning on the growth of intragranular a in near-b Ti alloys Tong Li a,⇑, Mansur Ahmed b, Gang Sha a,c, Rongpei Shi d, Gilberto Casillas e, Hung-Wei Yen a,f, Yunzhi Wang d, Elena V. Pereloma b,e, Julie M. Cairney a,⇑ a
Australian Centre for Microscopy and Microanalysis and School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia School of Mechanical, Materials & Mechatronic Engineering, University of Wollongong, NSW 2522, Australia School of Materials Science and Engineering, Nanjing University of Science and Technology, Jiangsu 210094, China d Department of Materials Science and Engineering, The Ohio State University, 2041 College Road, Columbus, OH 43210, USA e Electron Microscopy Centre, University of Wollongong, NSW 2519, Australia f Department of Materials Science and Engineering, National Taiwan University, Taipei 10617, Taiwan b c
a r t i c l e
i n f o
Article history: Received 17 March 2015 Accepted 20 April 2015 Available online 27 April 2015 Keywords: Titanium alloy Microstructure Growth kinetics Atom probe tomography (APT) Transmission electron microscopy (TEM)
a b s t r a c t We report on partitioning of alloying elements during the formation of fine intragranular a plates in a Ti-55521 alloy after thermo-mechanical processing (TMP) and isothermal ageing at 923 K. The microstructures were characterised using atom probe tomography and high-resolution transmission electron microscopy. The partitioning of Mo, V and Al are strongly affected by their diffusivities and their mutual interaction. This leads to a deviation of the measured contents of alloying elements in the two phases from the predicted equilibrium values. The alloying elements at the broad faces and tips of a plates were found to exhibit different pile-up and segregation behaviours, which is thought to affect the lengthening and thickening kinetics of the a plates. As a result, the aspect ratio of a plates decreased rapidly with increasing ageing time. This study suggests that careful selection of alloying elements could be an effective way in controlling the growth anisotropy of a plates and thus a + b microstructures in near-b Ti alloys. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Titanium alloys exhibit excellent properties such as high specific strength, excellent toughness and corrosion resistance, and are important materials in aerospace [1] and biomedical applications [2]. Ti-5553 (Ti–5Al–5V–5Mo–3Cr–0.5Fe, all in wt.%) is a near-b Ti alloy recently developed for high strength airframe components [3]. Ti-55521 (Ti–5Al–5Mo–5V–2Cr–1Fe in wt.%), a slight modification of Ti-5553, has been specifically developed for low cost blended powder metallurgy processing by small variations in Cr and Fe contents to accelerate sintering without any loss in Mo equivalence [4]. A good balance between strength and ductility of a near-b Ti alloy in the solution-treated-and-aged condition with a precipitation-hardened microstructure depends on the volume
⇑ Corresponding authors at: Australian Centre for Microscopy & Microanalysis, Madsen Building F 09, The University of Sydney, NSW 2006, Australia. Tel.: +61 293517548. E-mail addresses: [email protected] (T. Li), [email protected] (J.M. Cairney). http://dx.doi.org/10.1016/j.jallcom.2015.04.143 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
fraction, size, morphology and distribution of the a precipitates, which control the strength, and on b grain size that determines the ductility [5]. Thermo-mechanical processing (TMP) has also been utilised to engineer the microstructure to achieve optimal mechanical properties. A range of a phase morphologies can form in different types of Ti alloys (near a, a/b and near b), including allotriomorphic a (grain boundary a), colonies of Widmanstätten a side-plates that nucleate at grain boundaries and intragranular a in the interior of the b matrix [6–8]. Generally intragranular a precipitates develop into disc-shape during growth due to their orientation relationship with the matrix, coherency state, and the associated anisotropy in the interfacial energy [9], elastic strain energy [10] and growth kinetics. The tips of a plates are thought to have an incoherent a/b interface and the broad faces are thought to be either coherent or semi-coherent [11,12]. Lengthening and thickening of precipitates have been investigated by Aaronson et al. [11,13,14] and is believed to follow a ledge mechanism, where lengthening occurs by the addition of atoms at the incoherent risers of ledges at the tips of the plates and thickening results from the movement of growth ledges across the broad semi-coherent interfaces [15]. For thickening, the lattice rearrangement is slower than
213
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
the diffusion of alloying elements (i.e. interface-controlled). Ivasishin et al. [16,17] found that the heating rate has a strong effect on the a/b microstructure. A slow heating rate (0.25 K/s) results in a fine-grain b microstructure (10 lm) with a uniform dispersion of fine intragranular a precipitates that provide high strength (1.5 GPa). These fine precipitates are thought to form as a result of the formation of isothermal x particles, which act as nucleation sites for a via the phase transformation route b ? b + x ? b + a. An increase in the heating rate (20 K/s) leads to the formation of coarser and non-uniformly distributed a directly via b ? b + a transformation route, resulting in poorer mechanical properties. Alloying elements are of crucial importance for phase formation and microstructure development in Ti alloys. Al and O, being a stabilisers, prefer to segregate to the a phase, while V, Mo, Cr and Fe, being b stabilisers, are prone to diffuse into the b matrix [12,18,19]. The diffusivity of alloying elements differs substantially, e.g., O diffuses at least five orders of magnitude faster than the substitutional elements. A large difference in diffusivity will influence the distribution of alloying element at the moving interfaces during the growth of a precipitates. For example, analogous to the b to a phase transformation in Ti alloys, the kinetics of austenite to ferrite phase transformation in steels can be controlled by either fast carbon diffusion or slower substitutional element diffusion. Because C diffuses nearly 6 orders of magnitude faster than substitutional elements, the transformation mechanism can be described by several models, such as constrained paraequilibirum (PE) model [20,21], local equilibrium with partitioning (LE-P) model or local equilibrium with negligible partitioning (LE-NP) model [22,23]. In addition, interaction among the alloying elements may further complicate their partitioning behaviours. Enomoto and Yoshida [24] have studied the partition behaviour of proeutectoid a in Ti–X1–X2 alloys (X1 = Al or V, and X2 = Cr and Fe) between 773 K and 973 K with scanning transmission electron microscopy (STEM). They found that the transformation mode is not massive nor martensitic, but were unable to confirm whether it is local or para equilibrium at the growing interface due to the fact that the transmission electron microscopy (TEM) is unable to measure the interface composition of a three-dimensional feature accurately. Atom probe tomography (APT) is a powerful technique that is able to provide quantitative three-dimensional information about the distribution of alloying elements [25,26]. Nag et al. previously revealed pronounced partitioning of alloying elements between the a and b phases in Ti-5553 by using APT [27,28], and measured the composition of a that nucleated from x particles formed upon 30 min of heat treatment at 827 K [27–30]. However, the partitioning of alloying elements along different growth directions, the interface chemistry, and the interactions among alloying elements have not been measured [27,28]. Therefore, to date, there is a lack of systematic research addressing the effect of partitioning and segregation on the growth kinetics of a in Ti alloys. Experimental information at a fine length scale is required to derive transformation kinetics, and understanding the interactions among alloying elements is essential to achieve ‘bottom up’ alloy design. In this study APT is coupled with transmission electron microscopy (TEM) to investigate the composition and growth kinetics of intragranular a precipitates in Ti-55521 after TMP and isothermal heat treatment via slow heating, and to establish the role of a and b stabilisers in different stages of the phase transformation. Partitioning of Al, V, Mo, Cr, Fe and O between a and b is characterised in details to reveal how the alloying elements behave at different interfaces, i.e., along thickening and lengthening direction of the precipitates. The aim is to gain deeper insight into the thermodynamic nature of the phase transformation, the interaction among different alloying elements and their influence on the growth kinetics and, hence, morphology of a precipitates.
2. Materials and methods A Ti–5Al–5Mo–5V–2Fe–1Cr (wt.%) alloy (Ti–8.8Al–2.5Mo–4.7V–1.8Cr–0.85Fe (at.%)), was prepared by a powder metallurgy technique [4]. After sintering, the samples were thermo-mechanically processed (TMP) with a 3500 Gleeble thermo-mechanical simulator. A schematic diagram illustrating the heat treatment is shown in Fig. 1. The samples were heated to 1223 K at 10 K/s, held for 2 min, then cooled to 1173 K at 35 K/s where 25% deformation was applied, and again cooled to 1073 K at the same rate and held for 10 min and deformed to 60% reduction. The samples were cooled to ambient temperature at a rate of 10 K/s. They were then encapsulated in a vacuum to prevent high temperature oxidation and placed into a furnace for ageing. The samples were heated from room temperature to 923 K at the slow heating rate of 0.25 K/s and aged for 2 min, 60 min or 240 min. The Ti-55521 alloy exhibited an excellent combination of high strength, at 1.1 GPa, and a superior ductility of 15% [31]. The alloy composition (at.%) determined by spectroscopic chemical analysis is provided in Table 1. Microstructures were observed using scanning electron microscopy (SEM), transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM). SEM observations were conducted using a JEOL-JSM 7001F Field Emission Gun-Scanning Electron Microscope (FEG-SEM) operating at 15 kV and a 10 mm working distance. Thin foils for TEM were prepared using twin jet electropolishing with an electrolyte containing 5–7% perchloric acid in methanol. TEM observations were carried out on a JEOL JEM-2011 at an accelerating voltage of 200 kV. STEM images were acquired with a probe-corrected JEOL ARM200F operating at 200 kV, with a cold field emission gun. High-angle annular dark field (HAADF) images were acquired with 50 and 180 mrad inner and outer collection angles respectively, while BF images used 11 mrad collection angles. Both images were taken with a dwell time of 38 ls, and convergence solid angle of 25 mrad, resulting in a probe current of 35 pA. Measurement of the thickness and length of the intragranular a were obtained using ImageJ [32]. APT specimens were electropolished in a solution containing perchloric acid (5%), 2-butoxyelethnol (35%) and methanol at 30 V and 233 K. The electropolished APT tips were further sharpened in a Zeiss Auriga FIB until the radius of the tips was less than 100 nm. APT specimens were characterised by a Cameca LEAP 4000X-SI™ instrument in the laser pulsing mode at a specimen temperature of 60 K with a target evaporation rate of 5 ions per 1000 pulses, a pulse rate of 250 kHz, a laser energy of 80 pJ and a laser spot size of approximately 2 lm. The wavelength of the UV laser was 355 nm. APT data were reconstructed and analysed using IVAS 3.6.6™ software. Three or four APT datasets were collected for each condition. The error bars for the concentration in this study have been calculated by pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ðc ð100 cÞ=N , where c is the concentration (in at.%) and N is the total number atoms [25].
3. Results 3.1. Microstructural evolution SEM images of the Ti-55521 alloy after TMP, and also after subsequent ageing for 60 min and 240 min at 923 K, are shown in Fig. 2. The a phase is dark and b is light grey. Fig. 2a reveals that colonies of 1–2 lm thick a plates are present in the b matrix after
Fig. 1. Schematic diagram of heat treatment.
Table 1 Alloy composition (at.%) of Ti-55521.
at.%
Al
Mo
V
Cr
Fe
O
9.3
2.4
4.2
1.7
0.89
0.74
214
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
38 ± 6% for the samples aged at 923 K for 60 min and reaches 50 ± 3% after 240 min. TEM specimens were analysed along the [1 1 0] b axis in order to view the intragranular a plates edge-on [34]. The habit plane of the a phase, as shown in Fig. 3a, is ð11; 11; 13Þ b and the growth direc 5 [35]. After TMP (Fig. 3b), HAADF-STEM reveals that tion is ½3 3 fine intragranular a plates with a high aspect ratio already exist in some areas, e.g. the plate labelled 1 which is 300 nm long and 10 nm thick. Some intragranular a, labelled 2, has a relatively low aspect ratio. The average thickness of intragranular a plates measured from Fig. 3b is 10.4 ± 4.4 nm. Aberration-corrected high resolution HAADF-STEM, in Fig. 3c, taken from a specimen after TMP, shows an [0 0 0 1]-orientated intragranular a plate within the [1 1 0]-orientated b. The well-developed hexagonal close-packed (hcp) crystal has a semi-coherent interface with ledges visible at the broad face. In addition, ellipsoidal athermal x (2–3 nm) was observed, with faint reflections at 1/3 and 2/3 {1 1 2}b positions in the selected area diffraction patterns (Fig. 3d), and an orientation 0x. After ageing at 923 K for 2 min, the relationship [1 1 0]b//½1 1 2 HAADF-STEM image in Fig. 3e shows that the size of the intragranular a plates has grown to 30 nm in thickness and 200 nm in length in a basket weave form, with a Burgers orientation relation 1 1 0a//½1 1 1b [34]. After ageing ship, i.e. (0 0 0 1)a//(1 0 1)b and ½2 for 60 min, Fig. 3f shows that, the size of the intragranular a phase increases to 80 nm in thickness and 300 nm in length. After ageing for 240 min the intragranular a phase takes the form of coarse plates, 120 nm in thickness and 400 nm in length (Fig. 3g). The length, thickness and the aspect ratio of length to thickness were measured from Fig. 3(b, e–g) from twenty a plates in different TEM images, and the results are summarised in Fig. 3h. During isothermal ageing, excluding the data point at TMP, the length curve in Fig. 3h was best fit by a power-law curve represented by L = 163t0.15 (R2 = 0.99), and the width plot was best fit by a power-law curve represented by W = 27.8t0.26 (R2 = 0.99). The lengthening and thickening rates at various ageing times were estimated simply by the average length and width divided by time, as listed in Table 2. It is evident that a dramatic drop in the aspect ratio of intragranular a plates (from 12 ± 6.5 to 5.4 ± 2.2) is observed after ageing at 923 K for 2 min. This gradually decreases to 3.8 ± 0.9 after ageing for 60 min and levels off at 3.3 ± 0.6 after 240 min. 3.2. Elemental partitioning
Fig. 2. SEM images of Ti-55521 (a) after TMP, (b) aged for 60 min and (c) aged for 240 min at 923 K.
TMP. After ageing at 923 K for 60 min and 240 min, Fig. 2b and c shows that, besides the a colonies, which become coarser as ageing time increases, very thin, homogeneously-distributed intragranular a is present within the b phase. The relatively coarse a after TMP is predominantly allotriomorphic a along b grain boundaries, which could be residual grain boundaries during a b-solution treatment at 1223 K, or recrystallized grain boundaries that form during isothermal holding at 1073 K (below the b transus 1093 K). The b grains became elongated during the deformation at 1073 K, and thus the allotriomorphic a phase is curved. The volume fraction of a, measured by XRD in our previous study [33] (both allotriomorphic and intragranular) is 22 ± 3% after TMP, increases to
Representative atom probe analysis results from a specimen processed by TMP are shown in Fig. 4. a stabilisers (Al and O) and b stabilisers (V, Mo, Cr and Fe) are prone to diffuse to the a and b phases, respectively [18], so a Fe iso-concentration surface at 0.4 at.%, delineates the a phase (Fe-lean). The combined atom maps of Al, O (Fig. 4b) and Mo, V, Cr and Fe (Fig. 4c) confirm that the a phase regions are low in Mo, V, Cr and Fe. Two relatively long a plates with a thickness of 15 nm and a length over 100 nm, as well as four relatively thin a plates, 4–10 nm in thickness and 40 nm in length are observed in Fig. 4c. From the SEM images, we know that the coarse allotriomorphic a is approximately 1– 2 lm, (Fig. 2), so the a plates observed in the APT data must be the thin intragranular a phase observed in the TEM images in Fig. 3. A 1D concentration profile obtained by using an analysis Table 2 Lengthening and thickening rate at different ageing time. Time (min)
dL/dt (m/s)
dW/dt (m/s)
2 60 240
1.5 ± 0.7 109 8.3 ± 1.9 1011 2.6 ± 0.7 1011
2.7 ± 0.7 1010 2.3 ± 0.7 1011 8.0 ± 2.4 1012
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
215
Fig. 3. (a) Schematic diagram of an intragranular a precipitate with its habit plane, growth direction and the Burgers orientation relationship, (b) HAADF-STEM image of Ti-55521 after TMP showing thin intragranular a plates in b matrix along [1 1 0] b zone axis, (c) aberration-corrected high resolution HAADF-STEM image taken long [1 1 0] b zone axis showing a coherent interface between a precipitate and b after TMP, (d) dark field image showing x phase and its selected area diffraction pattern 0 0]x, (e) HAADF-STEM image of Ti-55521 along [1 1 0] b zone axis, B = [1 1 0]b//[1 1 after aging at 923 K for 2 min along [1 1 0] b zone axis, (f) HAADF-STEM image of Ti55521 after aging at 923 K for 60 min [1 1 0] b zone axis, (g) bright field image after aging at 923 K for 240 min and (h) a plot of width, length and aspect ratio versus aging time measured from (b, e–g).
value of Al in the parent b (8.8 at.%). The O content is 2.2 ± 0.1 at.%. Average Mo, V, Cr and Fe concentrations in the thin intragranular a plates are 1.5 ± 0.2 at.%, 2.3 ± 0.1 at.%, 0.22 ± 0.04 at.% and 0.10 ± 0.04 at.%, respectively. Additionally, in an atom probe data that contained only a (most likely allotriomorphic a), the composition is (10.3 ± 0.2) at.% Al–(0.09 ± 0.01) at.% Mo–(1.3 ± 0.1) at.% V–(0.13 ± 0.01) at.% Cr–(0.06 ± 0.01) at.% Fe–(4.0 ± 0.2) at.% O. Compared to the intragranular a, the allotriomorphic a has significantly higher levels of partitioning of alloying elements, especially Mo, V and Al which diffuse relatively slow. To describe the equilibrium shape of intragranular a [36], Furuhara et al. [34,37] proposed three interphase boundaries: 1 1 1 1 3b, side facet, and broad face (parallel to the habit plane 1 tip (perpendicular to the lattice invariant line [3 3 5]b) as sketched in Fig. 3a. Our experimental evidence shows that the four thin intragranular a plates are approximately disc-shaped (in Fig. 4a– c and the video in the supplementary material). For this reason, in this study, partitioning was considered only at the broad face and the tips. APT reconstructions from the sample aged for 2 min are shown in Fig. 5a–b (and the video in the supplementary material). A 0.4 at.% Fe iso-concentration surface reveals coarsening of the intragranular a. Fig. 5b is the same reconstruction shown in Fig. 5a, but rotated 90° anticlockwise. A Mo-rich region is observed in the vicinity of the tip of the precipitate in the b phase, highlighted by the Mo iso-concentration surface at 5.7 at.% (the nominal composition of Mo in b is 2.5 at.%). Two analysis volumes were placed perpendicular to the tip and broad face regions in Fig. 5a and b (with cross-sections of 60 90 nm and 87 64 nm respectively). 1D concentration profiles, in Fig. 5c, show that a significant build-up of Mo (6.1 ± 0.1 at.% at maxima) is present in the b phase at the a/b interface at the tip region (Fig. 5a), and Al pile-up (12.5 ± 0.1 at.% at maxima) is present in the a phase, on the opposite side of the same interface. A slight V build-up is also observed at the tip, and nearly no build-up for Cr and Fe. Comparatively, the pronounced build-up behaviour of Mo and Al were not observed at the broad face region, as shown in Fig. 5d (highlighted by the red1-dotted line indicating the Mo concentration plateaus at around 4.1 at.%). The interface at the tip is classified as a ‘high build-up’ interface, and the broad face as a ‘low build-up interface’. Similar high build-up and low build-up interfaces were achieved for the sample aged for 60 min. Two representative APT reconstructions are shown in Fig. 6a and c, where Fe (0.4 at.%) and Mo (5.7 at.%) iso-concentration surfaces are displayed. A high build-up of Mo is observed in region 1 in Fig. 6a and low (no) build-up of Mo in region 2 in Fig. 6c. Note that high build-up of Mo is also present at the tip in Fig. 6c. The solute distribution in region 1 and 2 are plotted in Fig. 6b and d obtained by placing analysis cylinders 1 and 2 (20 nm and 18 nm in diameter respectively) perpendicular to the two types of interfaces. After ageing for 240 min, only high build-up interfaces were observed in all atom probe datasets due to the further coarsening of a precipitates. An example is shown in Fig. 6e, with the corresponding 1D concentration profile (21 nm diameter) shown in Fig. 6f. The Mo and V build-up length extends over a longer distance (>25 nm, which was the length of the dataset), making it appear less pronounced compared to the samples aged for 2 min and 60 min. Another dataset from the specimen aged at 923 K for 240 min provides evidence that by this stage the a plates have grown considerably (see supplementary material). To compare various ageing conditions, Fig. 7 combines the concentration profiles in Figs. 5 and 6 for different ageing times across
cylinder (24 nm in diameter) perpendicular to one of four parallel smaller a plates (in Fig. 4b), as shown in Fig. 4d, reveals that the Al content in a is 9.0 ± 0.1 at.%, roughly equivalent to the nominal
1 For interpretation of color in Fig. 5, the reader is referred to the web version of this article.
216
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
Fig. 4. (a) 3D-APT reconstruction of Ti-55521 after TMP, with isoconcentration surfaces at 0.4 at.% Fe, cross-sectional APT reconstructions (in x–z view) of (b) Al, O and TiO, (c) Mo, V, Cr and Fe from the region highlighted by the orange rectangular box in (a), and (d) 1D concentration profiles from the cylinder across the a plate in the rectangular region in (b).
Fig. 5. (a) 3D-APT reconstruction of Ti-55521 after TMP and aging at 923 K for 2 min with iso-concentration surfaces at 0.4 at.% Fe and 5.7 at.%, Mo, (b) the same APT reconstruction as in (a) rotating 90° anti-clockwise, (c) 1D concentration profiles along the analysis cube placed in the tip region in (a), and (d) in the broad face region in (b).
the high and low build-up interfaces, with the position of the interface aligned for each case. The data for the high build-up interfaces is shown on the left hand side of the figure. No build-up of O was observed, whereas a continuous Al build-up was observed in the a phase. Mo and V exhibit pronounced build-up in b near the interface at all ageing conditions, together with a depletion of Al. The concentration gradient of Mo is the most significant, followed by the Al and V. Cr shows a gradual build-up after ageing for 2 min and 60 min but the build-up disappears for samples aged for
240 min. There is no significant build-up for Fe at any ageing condition. Comparatively, at the low build-up interface, a much less pronounced build-up of Mo is observed, and no build-up is seen for other alloying elements. 3.3. Chemical analysis in both phases at interfaces A summary of the average composition in the intragranular a and b, as well as the local concentration at the interface, for the
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
217
Fig. 6. Two reconstructed APT datasets from Ti-52221 after TMP and ageing at 923 K for 60 min isoconcentration surfaces at 0.4 at.% Fe and 5.7 at.%, Mo, showing a high build-up in region 1 in (a) and a low build-up in region 2 in (b), (c–d) 1D concentration profiles along the analysis cubes placed in region 1 and 2 respectively, and (e) an 3DAPT reconstruction of Ti-52221 after TMP and 240 min aging with isoconcentration surfaces at 0.4 at.% Fe and 5.7 at.%, Mo, and (f) 1D concentration profile along the analysis cylinder placed in the rectangular box in (e).
two interface types (high and low build-up) is provided in Fig. 8 (values are listed in a Table in the Supplementary Material). The average compositions (squares) were measured in locations far away from the interfaces. The interface concentrations (diamonds for the high build-up interfaces and triangles for the low build-up interfaces) were taken at the position of the dashed vertical lines in Fig. 7. Thermo-Calc has been used to predict equilibrium compositions at 923 K based on the nominal alloy composition (Ti–8.8Al–2.5Mo–4.7V–1.8Cr–0.85Fe in at.%), and the result is plotted as dashed lines in Fig. 8. In terms of average compositions (squares), in the a phase (Fig. 8a–f), the O, in Fig. 8a, as a strong a stabiliser, decreases slightly with ageing time. Al, also a strong a stabiliser, which has a similar concentration in both a and b after TMP, increases to 11.2 ± 0.1 at.% after ageing for 2 min at 923 K, which is equal to the equilibrium value. For longer ageing time, the average Al concentration becomes higher than the equilibrium values. The Mo and V concentration decreases in the a phase dramatically after 2 min of ageing, and remains nearly unchanged with further ageing. There is no significant variation of Cr and Fe in the a phase after TMP and ageing. Compared to the equilibrium values, Mo and Fe have lower contents than equilibrium concentrations, while
V and Cr are supersaturated in the a phase (especially V). In the b phase (Fig. 8a1–f1), the average Al concentration drops quickly in the b phase after ageing for 2 min, slightly decreases after ageing for 60 min and remains constant with further ageing, but does not reach equilibrium. The average Mo, V and Fe concentrations increase after ageing at 923 K for 2 min, and then remain constant with further ageing. The average concentration of Mo and Fe is approximately close to equilibrium, while that for V is significantly lower than the equilibrium value. A slight increase in the Cr concentration is observed, and it nearly reaches equilibrium after ageing for 240 min. At the high build-up interface (diamonds in Fig. 8), the concentration of Al in the a phase increases with ageing time to higher than equilibrium values (Fig. 8b). In the b phase, there is always a significant difference between composition at high build-up interface and the average composition for Mo and V (Fig. 8c1 and e1), but not for Al, Fe and Cr. The Mo (Fig. 8c1) content at the high build-up interface is above the equilibrium value while V is below it. At the low build-up interface (triangles), only Al and Mo (Fig. 8b1 and c1) show a gentle deviation from the average composition. The concentrations of V, Cr and Fe at the low build-up interfaces in b are the same as average composition in b.
218
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
Fig. 7. (a–f) O, Al, Mo, V, Cr and Fe concentration profiles at a high build-up interface, which is assumed to be diffusion-controlled interface, for samples aged at 923 K for 2 min, 60 min and 240 min, (a1–f1) O, Al, Mo, V, Cr and Fe concentration profiles at a low build-up interface, which is assumed to be interface-controlled, for samples aged at 923 K for 2 min and 60 min. The position of the interface has been aligned. The same bin size of 0.6 nm was used in all the concentration profiles.
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
219
Fig. 8. Average concentration (far-field concentration), concentration at low build-up and high build-up interfaces in (a–f) a phase, and (a1–f1) b phase after TMP and ageing for 2 min, 60 min and 240 min for 923 K.
We note that although trajectory aberrations in APT data can affect data from interfaces [25,26], this effect leads to blurring of interfaces, not differences in the extent of partitioning. The clear difference between the high and low build-up interface suggest the results presented here are due to the nature of interface itself, and not APT artefacts.
(11–16%) and high strength (1179–1304 MPa) [31]. The formation of evenly distributed fine intragranular a is thought to improve the strength and fracture toughness, therefore partitioning behaviour during the formation of intragranular a is the main focus of this study.
4.1. Nucleation of intragranular a 4. Discussion The bimodal structure of coarse allotriomorphic a and fine intragranular a is responsible for the combination of good ductility
After TMP, both coarse allotriomorphic a and fine intragranular a were observed (Figs. 2a and 3b–c). The allotriomorphic a forms along b grain boundaries during isothermal holding at 1073 K.
220
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
Table 3 Diffusivity of alloying elements in a-Ti and b-Ti at 923 K [43–45]. Diffusivity (m2/s)
Al
Mo 16
a-Ti
3.92 10 1.0 1017
b-Ti
V 19
6.25 10 5.2 1017
4.3 10 7.3 1016
Table 4 Simulated equilibrium composition with the absence of one element.
a phase Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium
(no (no (no (no (no
Fe) Cr) V) Mo) Al)
(no (no (no (no (no
Fe) Cr) V) Mo) Al)
b phase Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium
Cr 19
Al
Mo
V
Cr
Fe
10.5 10.5 9.4 10.2 – 10.8
0.75 1.4 0.78 – 0.11 0.88
0.009 0.001 – 0.009 0.079 0.007
0.003 – 0.008 0.007 0.003 0.032
– 0.019 0.093 0.019 0.012 0.29
Al
Mo
V
Cr
Fe
6.1 5.7 5.2 5.9 – 6.0
5.2 4.5 13.2 – 4.8 4.7
11.9 13.0 – 14.4 8.5 11.0
4.6 – 12.8 5.5 3.6 4.3
– 2.0 0.31 2.2 1.6 1.6
The nucleation of the fine intragranular a occurs during cooling from 1073 K to room temperature. In the higher temperature regime of cooling, fine a may nucleate at dislocations, as thin intragranular a plates have previously been found to precipitate at dislocations at temperatures above 873–923 K [38]. In addition, the formation of x (Fig. 3d) during cooling provides potent heterogeneous nucleation sites for intragranular a that forms during the slow heating (at a heat rate of 0.25 K/s). Duerig et al. found a plates nucleated from dislocations to be non-uniformly distributed with a high aspect ratio [39,40], like plate 1 in Fig. 3b. In the present study, ageing results in a uniform distribution of fine intragranular a (Fig. 2b and c), indicating that the dominant nucleation site during isothermal ageing is unlikely to be dislocations, but rather fine x precipitates. After TMP, the composition of the intragranular a in Fig. 4 deviates significantly from the composition of allotriomorphic a, which suggests that the intragranular a is far from equilibrium. A classical nucleation process should result in a region in the parent phase near equilibrium values, although the Gibbs-Thompson effect may lead to some variation [41]. Thus the nucleation of intragranular a is thought to follow a non-classical nucleation mode [27,42]. 4.2. Growth of intragranular a 4.2.1. Growth mode The intragranular a plates grow by a ledge mechanism, where migrating ledges (Fig. 3c) accumulate at the (incoherent) tips of the plates. In the early stage of a growth (after TMP), partitioning has already taken place (Fig. 4d), indicating growth via a mixed-mode process. During isothermal ageing, high build-up is consistently observed at the tips and low/no build-up interfaces are consistently observed at the broad faces (Fig. 7). The lengthening of a plates at the tips is thought to be diffusion-controlled, and the thickening along the broad face interface-controlled. At the diffusion-controlled interface (tip), the rate controlling factor is the volume diffusion of the slowest diffusing elements, as listed in Table 3 [43–45], i.e. Mo, V in b and Al in a. At the interface-controlled interface the rate controlling factor is the ledge growth, leaving sufficient time for the solutes to diffuse away from the interface.
Fe 16
7.33 10 1.1 1015
O 13
1.22 10 2.2 1013
4.0 105 1.4 105
Table 5 Chemical driving force for growth in various ageing conditions measured from the average composition in b phase and the concentrations at the high build-up interface in the b phase. Growth driving force (J)
2 min
60 min
240 min
In b (or low build-up interface) High build-up interface
38.95 1.94
24.05 0
23.88 0
4.2.2. Solute segregation at the tip The Mo concentration at the high build-up interface is 25% higher than the predicted equilibrium value, while the concentration of V at the same interface is 25% lower (Fig. 8c1 and e1) than equilibrium value. The high Mo concentration at the tip could be attributed to solute segregation, which occurs in addition to elemental partitioning. It is known there is a potential well of alloying elements at the a/b interface (interphase region) [46–48], which can promote segregation. In our case, Mo is speculated to be able to alter the interfacial energy most significantly, and therefore Mo is thought to segregate. Accordingly, the fact that the V content is considerably below equilibrium in the same region (Fig. 8e1) could be explained by the diffusion of V being impeded by the strong Mo interfacial segregation. A high Mo accumulation has been observed around the allotriomorphic a and its side plate in Ti-5553 in [28], but the Mo segregation has not been previously reported at the phase boundary. 4.2.3. Solute portioning in both phases At the early stages of a growth, the concentration of Al (a stabiliser) in a was found to be similar to that in the parent b phase (Fig. 4d). However, after TMP only, the O concentration in a is 2.0 ± 0.3 at.% (0.6 ± 0.06 at.% in b), which is actually slightly higher than that observed after ageing. The partitioning of b stabilisers (Mo, V, Cr, Fe) in Fig. 4d is relatively pronounced compared to nearly no partitioning of Al, despite the diffusion rates of Mo and V being similar to Al at 923 K. This suggests that, although Al is assumed to be a strong a stabiliser, it is not particularly effective at the early stage of a formation, while the fast diffusion of O and partitioning of b stabilisers are kinetically and thermodynamically favoured. This has not been reported previously. During further growth upon ageing, the Al content in a increases, and O content in a decreases (Fig. 8a and b). This suggests that Al is increasingly active as a stabiliser as ageing progresses, and the continuous partitioning of Al reduces the O partitioning. Additionally, a pronounced supersaturation in V (Fig. 8e) and depletion in Mo (Fig. 8c) in were observed in a, which is most likely due to the diffusion of V from a to b being impeded by the high Mo content at the diffusion-controlled interface. To further evaluate the interactions among the alloying elements thermodynamically, the equilibrium concentrations of the remaining alloying elements when a given element is removed from the Ti–8.8Al–2.5Mo–4.7V–1.8Cr–0.85Fe alloy were calculated by Thermo-Calc and the results are summarised in Table 4. In the a phase, the absence of V causes a decrease in Al. It is therefore expected that an increase in V will introduce more Al, which is consistent with our observations (Fig 8b and e). In b, the absence of Fe or Cr does not change the concentrations of the other elements
221
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222 Table 6 Atomic mobility of alloying elements in b matrix after TMP, in b after ageing for 240 min at 923 K, at high build-up interface in b after ageing for 240 min at 923 K. Atomic mobility in b (m2/s K) TMP 240 min High build-up interface (240 min)
Al
Mo 19
4.2 10 6.8 1019 6.7 1019
V 22
6.0 10 8.2 1023 3.2 1023
much compared to their equilibrium values. In the absence of V, the concentrations of both Mo and Cr increase dramatically. If there is no Mo in the system, V should increase from 11 at.% to 14.4 at.%. This again suggests that Mo and V have a strong influence on each other (and other alloying elements), while Fe, Cr and Al do not affect the equilibrium concentration in b phase significantly (for Fe, Cr, it may be due to the small addition). We note that O was not included in the phase equilibrium by Thermo-Calc, which could be a possible reason for the deviation of composition in a measured by APT from the calculated values, O influences the composition of a significantly. However, this is not expected to influence the general trends discussed above.
Cr 19
2.7 10 1.2 1019 8.8 1021
Fe 21
4.3 10 8.0 1022 4.3 1022
O 20
6.9 10 5.0 1021 2.3 1021
– –
measured for samples after TMP and aged for 240 min (Table 6). The atomic mobility of all solute elements decreases in the b matrix after ageing. For the slow diffusing element, Mo, the atomic mobility is further reduced at the high build-up interface compared to that in the low build-up interface. In addition, the atomic mobility of V at the high build-up interface (240 min) is dramatically decreased than that in b (240 min), confirming the speculation that the diffusion of V is impeded by the Mo segregation as described previously. Both thermodynamic and kinetic effects retard the migration of a/b interface in the lengthening direction, but not the thickening direction, where little pile-up behaviour occurs. Therefore, the overall effect is that the aspect ratio is reduced and the a plates become stubby.
4.3. Morphology of intragranular a 5. Conclusions The previous phase field studies on the growth of a precipitates in Ti–6Al–4V at 1073 K [49,50] have predicted that the lengthening rate is about 10 times faster than the thickening rate, and the lengthening rate should follow a linear law while the thickening follows a parabolic law. However, we observe that the aspect ratio of the intragranular a plates drops significantly after ageing for only 2 min at 923 K, followed by a continuous decrease with further ageing (Fig. 3h). At 2 min, the lengthening rate is roughly five times higher than the thickening rate, while at 240 min, the lengthening rate (2.7 ± 0.7 1011 m/s) is nearly the same as the thickening rate (8.0 ± 2.4 1012 m/s). A drop in the aspect ratio of fine intragranular a phase in Ti–7Mo–3Al–3Nb–3Cr and T i–5Al–5Mo–5V–3Cr alloys was also observed in previous studies [51,52]. In the phase field studies, local equilibrium is assumed to be maintained at the a/b interface. In fact, a deviation from the local equilibrium is observed at the a/b interface in the present study, which may explain the discrepancy between modelling and experimental observations. For a precipitate of a given volume, the aspect ratio is determined by the interplay between strain energy anisotropy and interface energy anisotropy [36]. However, for a growing precipitate, the growth anisotropy may play a dominant role. The growth kinetic is determined by a combination of chemical driving force for growth and the mobility of the interface. The chemical driving force for growth at different ageing conditions in b (equivalent to the low build-up interfaces due to similar compositions) and at the high build-up interfaces were calculated in Thermo-Calc, as listed in Table 5. A dramatic decrease in chemical driving force is found at the high build-up interfaces compared with the low build-up interfaces during isothermal ageing. It could be due to that the pronounced build-up of Mo in b and Al in a at the tip locally stabilises both b and a, leading to a decrease in the chemical driving force for the growth. On top of this, the chemical driving force would be further dissipated by Mo segregation at the tip. The growth driving force at the high build-up interfaces becomes negligible, suggesting the composition at the high build-up interface lie in the single b phase region. Secondly, from the kinetic view, the interface mobility could be decreased by the slow diffuser, Mo, at the tip, similar to a solute-drag effect at grain boundary [53] or phase boundaries reported in steels [47]. The atomic mobility of each alloying element at each type of interface were recalculated in DICTRA by using the corresponding compositions
To understand the growth kinetics, solute partitioning around intragranular a plates has been characterised by using APT and TEM after TMP and isothermal ageing of the alloy at 923 K for various times. 1. The a plates do not form at an equilibrium composition, and are therefore believed to nucleate via a non-classical mechanism. 2. In the early stage of a growth, Al is not an effective a stabiliser. The fast diffusion of O and partitioning of b stabilisers play a key role. As ageing time increases the partitioning of Al proceeds, which has a negative effect on O partitioning in a. 3. APT reveals a high build-up of the slow diffusing elements, Mo, V, Al, at the tip and low/no pile-up at the broad face of an intragranular a plate. 4. A high level of Mo at the tip is caused by the coupling between Mo segregation and pile-up due to slow diffusion. This high Mo segregation has a strong effect on the diffusion of V from a to b, causing low V content at the interface and high V in the a phase. 5. As aging progresses, the lengthening kinetic drops dramatically to close to the thickening kinetic. This is attributed to Mo segregation at the tip decreasing the chemical driving force and interface mobility for lengthening but not for thickening. This leads the formation of the stubby intragranular a plates. The empirical experimental information presented about the interactions between the alloying elements in both the a and b phases and at a/b interfaces is expected to provide useful and important guidance for the modelling of transformation kinetics and for the design of Ti alloys. Acknowledgements This work was supported the Australian Research Council (DP120100206), and with technical support and scientific input from the Australian Microscopy and Microanalysis Research Facility (AMMRF). R. Shi and Y. Wang would like to acknowledge the support by the US NSF DMREF program under Grant No. DMR-1435483. The authors acknowledge the use of JEOL ARM200F (ARC LIEF LE120100104) at UOW Electron Microscopy Centre. The authors are grateful to Prof. O.M. Ivasishin and Dr. D. Savvakin, IMP, NAS Ukraine, for supplying the material.
222
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
The authors would like to thank Dr. Hao Chen at Tsinghua University for extremely valuable discussions. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom.2015.04. 143. References [1] G. Lutjering, J.C. Williams, Titanium, Springer, Berlin, 2007. [2] M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, Prog. Mater Sci. 54 (2009) 397–425. [3] R.R. Boyer, R.D. Briggs, J. Mater. Eng. Perform. 14 (2005) 681–685. [4] O.M. Ivasishin, V.M. Anokhin, A.N. Demidik, D.G. Savvakin, Key Eng. Mater. 188 (2000) 55–62. [5] O.M. Ivasishin, R.V. Teliovich, Mater. Sci. Eng. A 263 (1999) 142–154. [6] S. Semiatin, V. Seetharaman, I. Weiss, JOM 49 (1997) 33–39. [7] I. Weiss, S. Semiatin, Mater. Sci. Eng. A 243 (1998) 46–65. [8] I. Weiss, S. Semiatin, Mater. Sci. Eng. A 263 (1999) 243–256. [9] E.A. Brener, V.I. Mel’nikov, Adv. Phys. 40 (1991) 53–97. [10] M. Fleck, C. Hüter, D. Pilipenko, R. Spatschek, E.A. Brener, Philos. Mag. 90 (2010) 265–286. [11] C. Laird, H. Aaronson, Acta Metall. 17 (1969) 505–519. [12] C. Zener, J. Appl. Phys. 20 (1949) 950–953. [13] H. Aaronson, C. Laird, K. Kinsman, Phase Transformations, ASM, Metals Park, 1970. pp. 313–396. [14] K.R. Kinsman, E. Eichen, H.I. Aaronson, Metall. Trans. A 6 (1975) 303–317. [15] G. Spanos, R. Masumura, R. Vandermeer, M. Enomoto, Acta Metall. Mater. 42 (1994) 4165–4176. [16] O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Mater. Sci. Eng. A 405 (2005) 296–305. [17] O.M. Ivasishin, P.E. Markovsky, Y.V. Matviychuk, S.L. Semiatin, C.H. Ward, S. Fox, J. Alloys Comp. 457 (2008) 296–309. [18] G. Welsch, R. Boyer, E. Collings, Materials Properties Handbook: Titanium Alloys, ASM international, Materials Park, 1993. [19] T. Furuhara, T. Maki, Mater. Sci. Eng. A 312 (2001) 145–154. [20] M. Hillert, Introduction to paraequilibrium, in: Internal Report, Swedish Institute of Metals Research, Stockholm, 1953. [21] M. Hillert, Metall. Trans. A 6 (1975) 5–19. [22] D.E. Coates, Metall. Trans. 3 (1972) 1203–1212. [23] J. Kirkaldy, Can. J. Phys. 36 (1958) 907–916. [24] M. Enomoto, T. Yoshida, ISIJ Int. 31 (1991) 767–774.
[25] M.K. Miller, A. Cerezo, M. Hetherington, G. Smith, Atom Probe Field Ion Microscopy, Oxford Science Publications, Oxford, 1996. [26] B. Gault, M.P. Moody, J.M. Cairney, S.P. Ringer, Atom Probe Microscopy, Springer, New York, 2012. [27] S. Nag, R. Banerjee, R. Srinivasan, J.Y. Hwang, M. Harper, H.L. Fraser, Acta Mater. 57 (2009) 2136–2147. [28] S. Nag, R. Banerjee, J. Hwang, M. Harper, H. Fraser, Philos. Mag. 89 (2009) 535– 552. [29] H.P. Ng, A. Devaraj, S. Nag, C.J. Bettles, M. Gibson, H.L. Fraser, B.C. Muddle, R. Banerjee, Acta Mater. 59 (2011) 2981–2991. [30] S. Nag, Y. Zheng, R.E.A. Williams, A. Devaraj, A. Boyne, Y. Wang, P.C. Collins, G.B. Viswanathan, J.S. Tiley, B.C. Muddle, R. Banerjee, H.L. Fraser, Acta Mater. 60 (2012) 6247–6256. [31] M. Ahmed, T. Li, G. Casillas, J.M. Cairney, D. Wexler, E.V. Pereloma, J. Alloys Comp. 629 (2015) 260–273. [32]
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
The influence of partitioning on the growth of intragranular a in near-b Ti alloys Tong Li a,⇑, Mansur Ahmed b, Gang Sha a,c, Rongpei Shi d, Gilberto Casillas e, Hung-Wei Yen a,f, Yunzhi Wang d, Elena V. Pereloma b,e, Julie M. Cairney a,⇑ a
Australian Centre for Microscopy and Microanalysis and School of Aerospace, Mechanical & Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia School of Mechanical, Materials & Mechatronic Engineering, University of Wollongong, NSW 2522, Australia School of Materials Science and Engineering, Nanjing University of Science and Technology, Jiangsu 210094, China d Department of Materials Science and Engineering, The Ohio State University, 2041 College Road, Columbus, OH 43210, USA e Electron Microscopy Centre, University of Wollongong, NSW 2519, Australia f Department of Materials Science and Engineering, National Taiwan University, Taipei 10617, Taiwan b c
a r t i c l e
i n f o
Article history: Received 17 March 2015 Accepted 20 April 2015 Available online 27 April 2015 Keywords: Titanium alloy Microstructure Growth kinetics Atom probe tomography (APT) Transmission electron microscopy (TEM)
a b s t r a c t We report on partitioning of alloying elements during the formation of fine intragranular a plates in a Ti-55521 alloy after thermo-mechanical processing (TMP) and isothermal ageing at 923 K. The microstructures were characterised using atom probe tomography and high-resolution transmission electron microscopy. The partitioning of Mo, V and Al are strongly affected by their diffusivities and their mutual interaction. This leads to a deviation of the measured contents of alloying elements in the two phases from the predicted equilibrium values. The alloying elements at the broad faces and tips of a plates were found to exhibit different pile-up and segregation behaviours, which is thought to affect the lengthening and thickening kinetics of the a plates. As a result, the aspect ratio of a plates decreased rapidly with increasing ageing time. This study suggests that careful selection of alloying elements could be an effective way in controlling the growth anisotropy of a plates and thus a + b microstructures in near-b Ti alloys. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Titanium alloys exhibit excellent properties such as high specific strength, excellent toughness and corrosion resistance, and are important materials in aerospace [1] and biomedical applications [2]. Ti-5553 (Ti–5Al–5V–5Mo–3Cr–0.5Fe, all in wt.%) is a near-b Ti alloy recently developed for high strength airframe components [3]. Ti-55521 (Ti–5Al–5Mo–5V–2Cr–1Fe in wt.%), a slight modification of Ti-5553, has been specifically developed for low cost blended powder metallurgy processing by small variations in Cr and Fe contents to accelerate sintering without any loss in Mo equivalence [4]. A good balance between strength and ductility of a near-b Ti alloy in the solution-treated-and-aged condition with a precipitation-hardened microstructure depends on the volume
⇑ Corresponding authors at: Australian Centre for Microscopy & Microanalysis, Madsen Building F 09, The University of Sydney, NSW 2006, Australia. Tel.: +61 293517548. E-mail addresses: [email protected] (T. Li), [email protected] (J.M. Cairney). http://dx.doi.org/10.1016/j.jallcom.2015.04.143 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
fraction, size, morphology and distribution of the a precipitates, which control the strength, and on b grain size that determines the ductility [5]. Thermo-mechanical processing (TMP) has also been utilised to engineer the microstructure to achieve optimal mechanical properties. A range of a phase morphologies can form in different types of Ti alloys (near a, a/b and near b), including allotriomorphic a (grain boundary a), colonies of Widmanstätten a side-plates that nucleate at grain boundaries and intragranular a in the interior of the b matrix [6–8]. Generally intragranular a precipitates develop into disc-shape during growth due to their orientation relationship with the matrix, coherency state, and the associated anisotropy in the interfacial energy [9], elastic strain energy [10] and growth kinetics. The tips of a plates are thought to have an incoherent a/b interface and the broad faces are thought to be either coherent or semi-coherent [11,12]. Lengthening and thickening of precipitates have been investigated by Aaronson et al. [11,13,14] and is believed to follow a ledge mechanism, where lengthening occurs by the addition of atoms at the incoherent risers of ledges at the tips of the plates and thickening results from the movement of growth ledges across the broad semi-coherent interfaces [15]. For thickening, the lattice rearrangement is slower than
213
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
the diffusion of alloying elements (i.e. interface-controlled). Ivasishin et al. [16,17] found that the heating rate has a strong effect on the a/b microstructure. A slow heating rate (0.25 K/s) results in a fine-grain b microstructure (10 lm) with a uniform dispersion of fine intragranular a precipitates that provide high strength (1.5 GPa). These fine precipitates are thought to form as a result of the formation of isothermal x particles, which act as nucleation sites for a via the phase transformation route b ? b + x ? b + a. An increase in the heating rate (20 K/s) leads to the formation of coarser and non-uniformly distributed a directly via b ? b + a transformation route, resulting in poorer mechanical properties. Alloying elements are of crucial importance for phase formation and microstructure development in Ti alloys. Al and O, being a stabilisers, prefer to segregate to the a phase, while V, Mo, Cr and Fe, being b stabilisers, are prone to diffuse into the b matrix [12,18,19]. The diffusivity of alloying elements differs substantially, e.g., O diffuses at least five orders of magnitude faster than the substitutional elements. A large difference in diffusivity will influence the distribution of alloying element at the moving interfaces during the growth of a precipitates. For example, analogous to the b to a phase transformation in Ti alloys, the kinetics of austenite to ferrite phase transformation in steels can be controlled by either fast carbon diffusion or slower substitutional element diffusion. Because C diffuses nearly 6 orders of magnitude faster than substitutional elements, the transformation mechanism can be described by several models, such as constrained paraequilibirum (PE) model [20,21], local equilibrium with partitioning (LE-P) model or local equilibrium with negligible partitioning (LE-NP) model [22,23]. In addition, interaction among the alloying elements may further complicate their partitioning behaviours. Enomoto and Yoshida [24] have studied the partition behaviour of proeutectoid a in Ti–X1–X2 alloys (X1 = Al or V, and X2 = Cr and Fe) between 773 K and 973 K with scanning transmission electron microscopy (STEM). They found that the transformation mode is not massive nor martensitic, but were unable to confirm whether it is local or para equilibrium at the growing interface due to the fact that the transmission electron microscopy (TEM) is unable to measure the interface composition of a three-dimensional feature accurately. Atom probe tomography (APT) is a powerful technique that is able to provide quantitative three-dimensional information about the distribution of alloying elements [25,26]. Nag et al. previously revealed pronounced partitioning of alloying elements between the a and b phases in Ti-5553 by using APT [27,28], and measured the composition of a that nucleated from x particles formed upon 30 min of heat treatment at 827 K [27–30]. However, the partitioning of alloying elements along different growth directions, the interface chemistry, and the interactions among alloying elements have not been measured [27,28]. Therefore, to date, there is a lack of systematic research addressing the effect of partitioning and segregation on the growth kinetics of a in Ti alloys. Experimental information at a fine length scale is required to derive transformation kinetics, and understanding the interactions among alloying elements is essential to achieve ‘bottom up’ alloy design. In this study APT is coupled with transmission electron microscopy (TEM) to investigate the composition and growth kinetics of intragranular a precipitates in Ti-55521 after TMP and isothermal heat treatment via slow heating, and to establish the role of a and b stabilisers in different stages of the phase transformation. Partitioning of Al, V, Mo, Cr, Fe and O between a and b is characterised in details to reveal how the alloying elements behave at different interfaces, i.e., along thickening and lengthening direction of the precipitates. The aim is to gain deeper insight into the thermodynamic nature of the phase transformation, the interaction among different alloying elements and their influence on the growth kinetics and, hence, morphology of a precipitates.
2. Materials and methods A Ti–5Al–5Mo–5V–2Fe–1Cr (wt.%) alloy (Ti–8.8Al–2.5Mo–4.7V–1.8Cr–0.85Fe (at.%)), was prepared by a powder metallurgy technique [4]. After sintering, the samples were thermo-mechanically processed (TMP) with a 3500 Gleeble thermo-mechanical simulator. A schematic diagram illustrating the heat treatment is shown in Fig. 1. The samples were heated to 1223 K at 10 K/s, held for 2 min, then cooled to 1173 K at 35 K/s where 25% deformation was applied, and again cooled to 1073 K at the same rate and held for 10 min and deformed to 60% reduction. The samples were cooled to ambient temperature at a rate of 10 K/s. They were then encapsulated in a vacuum to prevent high temperature oxidation and placed into a furnace for ageing. The samples were heated from room temperature to 923 K at the slow heating rate of 0.25 K/s and aged for 2 min, 60 min or 240 min. The Ti-55521 alloy exhibited an excellent combination of high strength, at 1.1 GPa, and a superior ductility of 15% [31]. The alloy composition (at.%) determined by spectroscopic chemical analysis is provided in Table 1. Microstructures were observed using scanning electron microscopy (SEM), transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM). SEM observations were conducted using a JEOL-JSM 7001F Field Emission Gun-Scanning Electron Microscope (FEG-SEM) operating at 15 kV and a 10 mm working distance. Thin foils for TEM were prepared using twin jet electropolishing with an electrolyte containing 5–7% perchloric acid in methanol. TEM observations were carried out on a JEOL JEM-2011 at an accelerating voltage of 200 kV. STEM images were acquired with a probe-corrected JEOL ARM200F operating at 200 kV, with a cold field emission gun. High-angle annular dark field (HAADF) images were acquired with 50 and 180 mrad inner and outer collection angles respectively, while BF images used 11 mrad collection angles. Both images were taken with a dwell time of 38 ls, and convergence solid angle of 25 mrad, resulting in a probe current of 35 pA. Measurement of the thickness and length of the intragranular a were obtained using ImageJ [32]. APT specimens were electropolished in a solution containing perchloric acid (5%), 2-butoxyelethnol (35%) and methanol at 30 V and 233 K. The electropolished APT tips were further sharpened in a Zeiss Auriga FIB until the radius of the tips was less than 100 nm. APT specimens were characterised by a Cameca LEAP 4000X-SI™ instrument in the laser pulsing mode at a specimen temperature of 60 K with a target evaporation rate of 5 ions per 1000 pulses, a pulse rate of 250 kHz, a laser energy of 80 pJ and a laser spot size of approximately 2 lm. The wavelength of the UV laser was 355 nm. APT data were reconstructed and analysed using IVAS 3.6.6™ software. Three or four APT datasets were collected for each condition. The error bars for the concentration in this study have been calculated by pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ðc ð100 cÞ=N , where c is the concentration (in at.%) and N is the total number atoms [25].
3. Results 3.1. Microstructural evolution SEM images of the Ti-55521 alloy after TMP, and also after subsequent ageing for 60 min and 240 min at 923 K, are shown in Fig. 2. The a phase is dark and b is light grey. Fig. 2a reveals that colonies of 1–2 lm thick a plates are present in the b matrix after
Fig. 1. Schematic diagram of heat treatment.
Table 1 Alloy composition (at.%) of Ti-55521.
at.%
Al
Mo
V
Cr
Fe
O
9.3
2.4
4.2
1.7
0.89
0.74
214
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
38 ± 6% for the samples aged at 923 K for 60 min and reaches 50 ± 3% after 240 min. TEM specimens were analysed along the [1 1 0] b axis in order to view the intragranular a plates edge-on [34]. The habit plane of the a phase, as shown in Fig. 3a, is ð11; 11; 13Þ b and the growth direc 5 [35]. After TMP (Fig. 3b), HAADF-STEM reveals that tion is ½3 3 fine intragranular a plates with a high aspect ratio already exist in some areas, e.g. the plate labelled 1 which is 300 nm long and 10 nm thick. Some intragranular a, labelled 2, has a relatively low aspect ratio. The average thickness of intragranular a plates measured from Fig. 3b is 10.4 ± 4.4 nm. Aberration-corrected high resolution HAADF-STEM, in Fig. 3c, taken from a specimen after TMP, shows an [0 0 0 1]-orientated intragranular a plate within the [1 1 0]-orientated b. The well-developed hexagonal close-packed (hcp) crystal has a semi-coherent interface with ledges visible at the broad face. In addition, ellipsoidal athermal x (2–3 nm) was observed, with faint reflections at 1/3 and 2/3 {1 1 2}b positions in the selected area diffraction patterns (Fig. 3d), and an orientation 0x. After ageing at 923 K for 2 min, the relationship [1 1 0]b//½1 1 2 HAADF-STEM image in Fig. 3e shows that the size of the intragranular a plates has grown to 30 nm in thickness and 200 nm in length in a basket weave form, with a Burgers orientation relation 1 1 0a//½1 1 1b [34]. After ageing ship, i.e. (0 0 0 1)a//(1 0 1)b and ½2 for 60 min, Fig. 3f shows that, the size of the intragranular a phase increases to 80 nm in thickness and 300 nm in length. After ageing for 240 min the intragranular a phase takes the form of coarse plates, 120 nm in thickness and 400 nm in length (Fig. 3g). The length, thickness and the aspect ratio of length to thickness were measured from Fig. 3(b, e–g) from twenty a plates in different TEM images, and the results are summarised in Fig. 3h. During isothermal ageing, excluding the data point at TMP, the length curve in Fig. 3h was best fit by a power-law curve represented by L = 163t0.15 (R2 = 0.99), and the width plot was best fit by a power-law curve represented by W = 27.8t0.26 (R2 = 0.99). The lengthening and thickening rates at various ageing times were estimated simply by the average length and width divided by time, as listed in Table 2. It is evident that a dramatic drop in the aspect ratio of intragranular a plates (from 12 ± 6.5 to 5.4 ± 2.2) is observed after ageing at 923 K for 2 min. This gradually decreases to 3.8 ± 0.9 after ageing for 60 min and levels off at 3.3 ± 0.6 after 240 min. 3.2. Elemental partitioning
Fig. 2. SEM images of Ti-55521 (a) after TMP, (b) aged for 60 min and (c) aged for 240 min at 923 K.
TMP. After ageing at 923 K for 60 min and 240 min, Fig. 2b and c shows that, besides the a colonies, which become coarser as ageing time increases, very thin, homogeneously-distributed intragranular a is present within the b phase. The relatively coarse a after TMP is predominantly allotriomorphic a along b grain boundaries, which could be residual grain boundaries during a b-solution treatment at 1223 K, or recrystallized grain boundaries that form during isothermal holding at 1073 K (below the b transus 1093 K). The b grains became elongated during the deformation at 1073 K, and thus the allotriomorphic a phase is curved. The volume fraction of a, measured by XRD in our previous study [33] (both allotriomorphic and intragranular) is 22 ± 3% after TMP, increases to
Representative atom probe analysis results from a specimen processed by TMP are shown in Fig. 4. a stabilisers (Al and O) and b stabilisers (V, Mo, Cr and Fe) are prone to diffuse to the a and b phases, respectively [18], so a Fe iso-concentration surface at 0.4 at.%, delineates the a phase (Fe-lean). The combined atom maps of Al, O (Fig. 4b) and Mo, V, Cr and Fe (Fig. 4c) confirm that the a phase regions are low in Mo, V, Cr and Fe. Two relatively long a plates with a thickness of 15 nm and a length over 100 nm, as well as four relatively thin a plates, 4–10 nm in thickness and 40 nm in length are observed in Fig. 4c. From the SEM images, we know that the coarse allotriomorphic a is approximately 1– 2 lm, (Fig. 2), so the a plates observed in the APT data must be the thin intragranular a phase observed in the TEM images in Fig. 3. A 1D concentration profile obtained by using an analysis Table 2 Lengthening and thickening rate at different ageing time. Time (min)
dL/dt (m/s)
dW/dt (m/s)
2 60 240
1.5 ± 0.7 109 8.3 ± 1.9 1011 2.6 ± 0.7 1011
2.7 ± 0.7 1010 2.3 ± 0.7 1011 8.0 ± 2.4 1012
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
215
Fig. 3. (a) Schematic diagram of an intragranular a precipitate with its habit plane, growth direction and the Burgers orientation relationship, (b) HAADF-STEM image of Ti-55521 after TMP showing thin intragranular a plates in b matrix along [1 1 0] b zone axis, (c) aberration-corrected high resolution HAADF-STEM image taken long [1 1 0] b zone axis showing a coherent interface between a precipitate and b after TMP, (d) dark field image showing x phase and its selected area diffraction pattern 0 0]x, (e) HAADF-STEM image of Ti-55521 along [1 1 0] b zone axis, B = [1 1 0]b//[1 1 after aging at 923 K for 2 min along [1 1 0] b zone axis, (f) HAADF-STEM image of Ti55521 after aging at 923 K for 60 min [1 1 0] b zone axis, (g) bright field image after aging at 923 K for 240 min and (h) a plot of width, length and aspect ratio versus aging time measured from (b, e–g).
value of Al in the parent b (8.8 at.%). The O content is 2.2 ± 0.1 at.%. Average Mo, V, Cr and Fe concentrations in the thin intragranular a plates are 1.5 ± 0.2 at.%, 2.3 ± 0.1 at.%, 0.22 ± 0.04 at.% and 0.10 ± 0.04 at.%, respectively. Additionally, in an atom probe data that contained only a (most likely allotriomorphic a), the composition is (10.3 ± 0.2) at.% Al–(0.09 ± 0.01) at.% Mo–(1.3 ± 0.1) at.% V–(0.13 ± 0.01) at.% Cr–(0.06 ± 0.01) at.% Fe–(4.0 ± 0.2) at.% O. Compared to the intragranular a, the allotriomorphic a has significantly higher levels of partitioning of alloying elements, especially Mo, V and Al which diffuse relatively slow. To describe the equilibrium shape of intragranular a [36], Furuhara et al. [34,37] proposed three interphase boundaries: 1 1 1 1 3b, side facet, and broad face (parallel to the habit plane 1 tip (perpendicular to the lattice invariant line [3 3 5]b) as sketched in Fig. 3a. Our experimental evidence shows that the four thin intragranular a plates are approximately disc-shaped (in Fig. 4a– c and the video in the supplementary material). For this reason, in this study, partitioning was considered only at the broad face and the tips. APT reconstructions from the sample aged for 2 min are shown in Fig. 5a–b (and the video in the supplementary material). A 0.4 at.% Fe iso-concentration surface reveals coarsening of the intragranular a. Fig. 5b is the same reconstruction shown in Fig. 5a, but rotated 90° anticlockwise. A Mo-rich region is observed in the vicinity of the tip of the precipitate in the b phase, highlighted by the Mo iso-concentration surface at 5.7 at.% (the nominal composition of Mo in b is 2.5 at.%). Two analysis volumes were placed perpendicular to the tip and broad face regions in Fig. 5a and b (with cross-sections of 60 90 nm and 87 64 nm respectively). 1D concentration profiles, in Fig. 5c, show that a significant build-up of Mo (6.1 ± 0.1 at.% at maxima) is present in the b phase at the a/b interface at the tip region (Fig. 5a), and Al pile-up (12.5 ± 0.1 at.% at maxima) is present in the a phase, on the opposite side of the same interface. A slight V build-up is also observed at the tip, and nearly no build-up for Cr and Fe. Comparatively, the pronounced build-up behaviour of Mo and Al were not observed at the broad face region, as shown in Fig. 5d (highlighted by the red1-dotted line indicating the Mo concentration plateaus at around 4.1 at.%). The interface at the tip is classified as a ‘high build-up’ interface, and the broad face as a ‘low build-up interface’. Similar high build-up and low build-up interfaces were achieved for the sample aged for 60 min. Two representative APT reconstructions are shown in Fig. 6a and c, where Fe (0.4 at.%) and Mo (5.7 at.%) iso-concentration surfaces are displayed. A high build-up of Mo is observed in region 1 in Fig. 6a and low (no) build-up of Mo in region 2 in Fig. 6c. Note that high build-up of Mo is also present at the tip in Fig. 6c. The solute distribution in region 1 and 2 are plotted in Fig. 6b and d obtained by placing analysis cylinders 1 and 2 (20 nm and 18 nm in diameter respectively) perpendicular to the two types of interfaces. After ageing for 240 min, only high build-up interfaces were observed in all atom probe datasets due to the further coarsening of a precipitates. An example is shown in Fig. 6e, with the corresponding 1D concentration profile (21 nm diameter) shown in Fig. 6f. The Mo and V build-up length extends over a longer distance (>25 nm, which was the length of the dataset), making it appear less pronounced compared to the samples aged for 2 min and 60 min. Another dataset from the specimen aged at 923 K for 240 min provides evidence that by this stage the a plates have grown considerably (see supplementary material). To compare various ageing conditions, Fig. 7 combines the concentration profiles in Figs. 5 and 6 for different ageing times across
cylinder (24 nm in diameter) perpendicular to one of four parallel smaller a plates (in Fig. 4b), as shown in Fig. 4d, reveals that the Al content in a is 9.0 ± 0.1 at.%, roughly equivalent to the nominal
1 For interpretation of color in Fig. 5, the reader is referred to the web version of this article.
216
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
Fig. 4. (a) 3D-APT reconstruction of Ti-55521 after TMP, with isoconcentration surfaces at 0.4 at.% Fe, cross-sectional APT reconstructions (in x–z view) of (b) Al, O and TiO, (c) Mo, V, Cr and Fe from the region highlighted by the orange rectangular box in (a), and (d) 1D concentration profiles from the cylinder across the a plate in the rectangular region in (b).
Fig. 5. (a) 3D-APT reconstruction of Ti-55521 after TMP and aging at 923 K for 2 min with iso-concentration surfaces at 0.4 at.% Fe and 5.7 at.%, Mo, (b) the same APT reconstruction as in (a) rotating 90° anti-clockwise, (c) 1D concentration profiles along the analysis cube placed in the tip region in (a), and (d) in the broad face region in (b).
the high and low build-up interfaces, with the position of the interface aligned for each case. The data for the high build-up interfaces is shown on the left hand side of the figure. No build-up of O was observed, whereas a continuous Al build-up was observed in the a phase. Mo and V exhibit pronounced build-up in b near the interface at all ageing conditions, together with a depletion of Al. The concentration gradient of Mo is the most significant, followed by the Al and V. Cr shows a gradual build-up after ageing for 2 min and 60 min but the build-up disappears for samples aged for
240 min. There is no significant build-up for Fe at any ageing condition. Comparatively, at the low build-up interface, a much less pronounced build-up of Mo is observed, and no build-up is seen for other alloying elements. 3.3. Chemical analysis in both phases at interfaces A summary of the average composition in the intragranular a and b, as well as the local concentration at the interface, for the
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
217
Fig. 6. Two reconstructed APT datasets from Ti-52221 after TMP and ageing at 923 K for 60 min isoconcentration surfaces at 0.4 at.% Fe and 5.7 at.%, Mo, showing a high build-up in region 1 in (a) and a low build-up in region 2 in (b), (c–d) 1D concentration profiles along the analysis cubes placed in region 1 and 2 respectively, and (e) an 3DAPT reconstruction of Ti-52221 after TMP and 240 min aging with isoconcentration surfaces at 0.4 at.% Fe and 5.7 at.%, Mo, and (f) 1D concentration profile along the analysis cylinder placed in the rectangular box in (e).
two interface types (high and low build-up) is provided in Fig. 8 (values are listed in a Table in the Supplementary Material). The average compositions (squares) were measured in locations far away from the interfaces. The interface concentrations (diamonds for the high build-up interfaces and triangles for the low build-up interfaces) were taken at the position of the dashed vertical lines in Fig. 7. Thermo-Calc has been used to predict equilibrium compositions at 923 K based on the nominal alloy composition (Ti–8.8Al–2.5Mo–4.7V–1.8Cr–0.85Fe in at.%), and the result is plotted as dashed lines in Fig. 8. In terms of average compositions (squares), in the a phase (Fig. 8a–f), the O, in Fig. 8a, as a strong a stabiliser, decreases slightly with ageing time. Al, also a strong a stabiliser, which has a similar concentration in both a and b after TMP, increases to 11.2 ± 0.1 at.% after ageing for 2 min at 923 K, which is equal to the equilibrium value. For longer ageing time, the average Al concentration becomes higher than the equilibrium values. The Mo and V concentration decreases in the a phase dramatically after 2 min of ageing, and remains nearly unchanged with further ageing. There is no significant variation of Cr and Fe in the a phase after TMP and ageing. Compared to the equilibrium values, Mo and Fe have lower contents than equilibrium concentrations, while
V and Cr are supersaturated in the a phase (especially V). In the b phase (Fig. 8a1–f1), the average Al concentration drops quickly in the b phase after ageing for 2 min, slightly decreases after ageing for 60 min and remains constant with further ageing, but does not reach equilibrium. The average Mo, V and Fe concentrations increase after ageing at 923 K for 2 min, and then remain constant with further ageing. The average concentration of Mo and Fe is approximately close to equilibrium, while that for V is significantly lower than the equilibrium value. A slight increase in the Cr concentration is observed, and it nearly reaches equilibrium after ageing for 240 min. At the high build-up interface (diamonds in Fig. 8), the concentration of Al in the a phase increases with ageing time to higher than equilibrium values (Fig. 8b). In the b phase, there is always a significant difference between composition at high build-up interface and the average composition for Mo and V (Fig. 8c1 and e1), but not for Al, Fe and Cr. The Mo (Fig. 8c1) content at the high build-up interface is above the equilibrium value while V is below it. At the low build-up interface (triangles), only Al and Mo (Fig. 8b1 and c1) show a gentle deviation from the average composition. The concentrations of V, Cr and Fe at the low build-up interfaces in b are the same as average composition in b.
218
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
Fig. 7. (a–f) O, Al, Mo, V, Cr and Fe concentration profiles at a high build-up interface, which is assumed to be diffusion-controlled interface, for samples aged at 923 K for 2 min, 60 min and 240 min, (a1–f1) O, Al, Mo, V, Cr and Fe concentration profiles at a low build-up interface, which is assumed to be interface-controlled, for samples aged at 923 K for 2 min and 60 min. The position of the interface has been aligned. The same bin size of 0.6 nm was used in all the concentration profiles.
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
219
Fig. 8. Average concentration (far-field concentration), concentration at low build-up and high build-up interfaces in (a–f) a phase, and (a1–f1) b phase after TMP and ageing for 2 min, 60 min and 240 min for 923 K.
We note that although trajectory aberrations in APT data can affect data from interfaces [25,26], this effect leads to blurring of interfaces, not differences in the extent of partitioning. The clear difference between the high and low build-up interface suggest the results presented here are due to the nature of interface itself, and not APT artefacts.
(11–16%) and high strength (1179–1304 MPa) [31]. The formation of evenly distributed fine intragranular a is thought to improve the strength and fracture toughness, therefore partitioning behaviour during the formation of intragranular a is the main focus of this study.
4.1. Nucleation of intragranular a 4. Discussion The bimodal structure of coarse allotriomorphic a and fine intragranular a is responsible for the combination of good ductility
After TMP, both coarse allotriomorphic a and fine intragranular a were observed (Figs. 2a and 3b–c). The allotriomorphic a forms along b grain boundaries during isothermal holding at 1073 K.
220
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
Table 3 Diffusivity of alloying elements in a-Ti and b-Ti at 923 K [43–45]. Diffusivity (m2/s)
Al
Mo 16
a-Ti
3.92 10 1.0 1017
b-Ti
V 19
6.25 10 5.2 1017
4.3 10 7.3 1016
Table 4 Simulated equilibrium composition with the absence of one element.
a phase Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium
(no (no (no (no (no
Fe) Cr) V) Mo) Al)
(no (no (no (no (no
Fe) Cr) V) Mo) Al)
b phase Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium Equilibrium
Cr 19
Al
Mo
V
Cr
Fe
10.5 10.5 9.4 10.2 – 10.8
0.75 1.4 0.78 – 0.11 0.88
0.009 0.001 – 0.009 0.079 0.007
0.003 – 0.008 0.007 0.003 0.032
– 0.019 0.093 0.019 0.012 0.29
Al
Mo
V
Cr
Fe
6.1 5.7 5.2 5.9 – 6.0
5.2 4.5 13.2 – 4.8 4.7
11.9 13.0 – 14.4 8.5 11.0
4.6 – 12.8 5.5 3.6 4.3
– 2.0 0.31 2.2 1.6 1.6
The nucleation of the fine intragranular a occurs during cooling from 1073 K to room temperature. In the higher temperature regime of cooling, fine a may nucleate at dislocations, as thin intragranular a plates have previously been found to precipitate at dislocations at temperatures above 873–923 K [38]. In addition, the formation of x (Fig. 3d) during cooling provides potent heterogeneous nucleation sites for intragranular a that forms during the slow heating (at a heat rate of 0.25 K/s). Duerig et al. found a plates nucleated from dislocations to be non-uniformly distributed with a high aspect ratio [39,40], like plate 1 in Fig. 3b. In the present study, ageing results in a uniform distribution of fine intragranular a (Fig. 2b and c), indicating that the dominant nucleation site during isothermal ageing is unlikely to be dislocations, but rather fine x precipitates. After TMP, the composition of the intragranular a in Fig. 4 deviates significantly from the composition of allotriomorphic a, which suggests that the intragranular a is far from equilibrium. A classical nucleation process should result in a region in the parent phase near equilibrium values, although the Gibbs-Thompson effect may lead to some variation [41]. Thus the nucleation of intragranular a is thought to follow a non-classical nucleation mode [27,42]. 4.2. Growth of intragranular a 4.2.1. Growth mode The intragranular a plates grow by a ledge mechanism, where migrating ledges (Fig. 3c) accumulate at the (incoherent) tips of the plates. In the early stage of a growth (after TMP), partitioning has already taken place (Fig. 4d), indicating growth via a mixed-mode process. During isothermal ageing, high build-up is consistently observed at the tips and low/no build-up interfaces are consistently observed at the broad faces (Fig. 7). The lengthening of a plates at the tips is thought to be diffusion-controlled, and the thickening along the broad face interface-controlled. At the diffusion-controlled interface (tip), the rate controlling factor is the volume diffusion of the slowest diffusing elements, as listed in Table 3 [43–45], i.e. Mo, V in b and Al in a. At the interface-controlled interface the rate controlling factor is the ledge growth, leaving sufficient time for the solutes to diffuse away from the interface.
Fe 16
7.33 10 1.1 1015
O 13
1.22 10 2.2 1013
4.0 105 1.4 105
Table 5 Chemical driving force for growth in various ageing conditions measured from the average composition in b phase and the concentrations at the high build-up interface in the b phase. Growth driving force (J)
2 min
60 min
240 min
In b (or low build-up interface) High build-up interface
38.95 1.94
24.05 0
23.88 0
4.2.2. Solute segregation at the tip The Mo concentration at the high build-up interface is 25% higher than the predicted equilibrium value, while the concentration of V at the same interface is 25% lower (Fig. 8c1 and e1) than equilibrium value. The high Mo concentration at the tip could be attributed to solute segregation, which occurs in addition to elemental partitioning. It is known there is a potential well of alloying elements at the a/b interface (interphase region) [46–48], which can promote segregation. In our case, Mo is speculated to be able to alter the interfacial energy most significantly, and therefore Mo is thought to segregate. Accordingly, the fact that the V content is considerably below equilibrium in the same region (Fig. 8e1) could be explained by the diffusion of V being impeded by the strong Mo interfacial segregation. A high Mo accumulation has been observed around the allotriomorphic a and its side plate in Ti-5553 in [28], but the Mo segregation has not been previously reported at the phase boundary. 4.2.3. Solute portioning in both phases At the early stages of a growth, the concentration of Al (a stabiliser) in a was found to be similar to that in the parent b phase (Fig. 4d). However, after TMP only, the O concentration in a is 2.0 ± 0.3 at.% (0.6 ± 0.06 at.% in b), which is actually slightly higher than that observed after ageing. The partitioning of b stabilisers (Mo, V, Cr, Fe) in Fig. 4d is relatively pronounced compared to nearly no partitioning of Al, despite the diffusion rates of Mo and V being similar to Al at 923 K. This suggests that, although Al is assumed to be a strong a stabiliser, it is not particularly effective at the early stage of a formation, while the fast diffusion of O and partitioning of b stabilisers are kinetically and thermodynamically favoured. This has not been reported previously. During further growth upon ageing, the Al content in a increases, and O content in a decreases (Fig. 8a and b). This suggests that Al is increasingly active as a stabiliser as ageing progresses, and the continuous partitioning of Al reduces the O partitioning. Additionally, a pronounced supersaturation in V (Fig. 8e) and depletion in Mo (Fig. 8c) in were observed in a, which is most likely due to the diffusion of V from a to b being impeded by the high Mo content at the diffusion-controlled interface. To further evaluate the interactions among the alloying elements thermodynamically, the equilibrium concentrations of the remaining alloying elements when a given element is removed from the Ti–8.8Al–2.5Mo–4.7V–1.8Cr–0.85Fe alloy were calculated by Thermo-Calc and the results are summarised in Table 4. In the a phase, the absence of V causes a decrease in Al. It is therefore expected that an increase in V will introduce more Al, which is consistent with our observations (Fig 8b and e). In b, the absence of Fe or Cr does not change the concentrations of the other elements
221
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222 Table 6 Atomic mobility of alloying elements in b matrix after TMP, in b after ageing for 240 min at 923 K, at high build-up interface in b after ageing for 240 min at 923 K. Atomic mobility in b (m2/s K) TMP 240 min High build-up interface (240 min)
Al
Mo 19
4.2 10 6.8 1019 6.7 1019
V 22
6.0 10 8.2 1023 3.2 1023
much compared to their equilibrium values. In the absence of V, the concentrations of both Mo and Cr increase dramatically. If there is no Mo in the system, V should increase from 11 at.% to 14.4 at.%. This again suggests that Mo and V have a strong influence on each other (and other alloying elements), while Fe, Cr and Al do not affect the equilibrium concentration in b phase significantly (for Fe, Cr, it may be due to the small addition). We note that O was not included in the phase equilibrium by Thermo-Calc, which could be a possible reason for the deviation of composition in a measured by APT from the calculated values, O influences the composition of a significantly. However, this is not expected to influence the general trends discussed above.
Cr 19
2.7 10 1.2 1019 8.8 1021
Fe 21
4.3 10 8.0 1022 4.3 1022
O 20
6.9 10 5.0 1021 2.3 1021
– –
measured for samples after TMP and aged for 240 min (Table 6). The atomic mobility of all solute elements decreases in the b matrix after ageing. For the slow diffusing element, Mo, the atomic mobility is further reduced at the high build-up interface compared to that in the low build-up interface. In addition, the atomic mobility of V at the high build-up interface (240 min) is dramatically decreased than that in b (240 min), confirming the speculation that the diffusion of V is impeded by the Mo segregation as described previously. Both thermodynamic and kinetic effects retard the migration of a/b interface in the lengthening direction, but not the thickening direction, where little pile-up behaviour occurs. Therefore, the overall effect is that the aspect ratio is reduced and the a plates become stubby.
4.3. Morphology of intragranular a 5. Conclusions The previous phase field studies on the growth of a precipitates in Ti–6Al–4V at 1073 K [49,50] have predicted that the lengthening rate is about 10 times faster than the thickening rate, and the lengthening rate should follow a linear law while the thickening follows a parabolic law. However, we observe that the aspect ratio of the intragranular a plates drops significantly after ageing for only 2 min at 923 K, followed by a continuous decrease with further ageing (Fig. 3h). At 2 min, the lengthening rate is roughly five times higher than the thickening rate, while at 240 min, the lengthening rate (2.7 ± 0.7 1011 m/s) is nearly the same as the thickening rate (8.0 ± 2.4 1012 m/s). A drop in the aspect ratio of fine intragranular a phase in Ti–7Mo–3Al–3Nb–3Cr and T i–5Al–5Mo–5V–3Cr alloys was also observed in previous studies [51,52]. In the phase field studies, local equilibrium is assumed to be maintained at the a/b interface. In fact, a deviation from the local equilibrium is observed at the a/b interface in the present study, which may explain the discrepancy between modelling and experimental observations. For a precipitate of a given volume, the aspect ratio is determined by the interplay between strain energy anisotropy and interface energy anisotropy [36]. However, for a growing precipitate, the growth anisotropy may play a dominant role. The growth kinetic is determined by a combination of chemical driving force for growth and the mobility of the interface. The chemical driving force for growth at different ageing conditions in b (equivalent to the low build-up interfaces due to similar compositions) and at the high build-up interfaces were calculated in Thermo-Calc, as listed in Table 5. A dramatic decrease in chemical driving force is found at the high build-up interfaces compared with the low build-up interfaces during isothermal ageing. It could be due to that the pronounced build-up of Mo in b and Al in a at the tip locally stabilises both b and a, leading to a decrease in the chemical driving force for the growth. On top of this, the chemical driving force would be further dissipated by Mo segregation at the tip. The growth driving force at the high build-up interfaces becomes negligible, suggesting the composition at the high build-up interface lie in the single b phase region. Secondly, from the kinetic view, the interface mobility could be decreased by the slow diffuser, Mo, at the tip, similar to a solute-drag effect at grain boundary [53] or phase boundaries reported in steels [47]. The atomic mobility of each alloying element at each type of interface were recalculated in DICTRA by using the corresponding compositions
To understand the growth kinetics, solute partitioning around intragranular a plates has been characterised by using APT and TEM after TMP and isothermal ageing of the alloy at 923 K for various times. 1. The a plates do not form at an equilibrium composition, and are therefore believed to nucleate via a non-classical mechanism. 2. In the early stage of a growth, Al is not an effective a stabiliser. The fast diffusion of O and partitioning of b stabilisers play a key role. As ageing time increases the partitioning of Al proceeds, which has a negative effect on O partitioning in a. 3. APT reveals a high build-up of the slow diffusing elements, Mo, V, Al, at the tip and low/no pile-up at the broad face of an intragranular a plate. 4. A high level of Mo at the tip is caused by the coupling between Mo segregation and pile-up due to slow diffusion. This high Mo segregation has a strong effect on the diffusion of V from a to b, causing low V content at the interface and high V in the a phase. 5. As aging progresses, the lengthening kinetic drops dramatically to close to the thickening kinetic. This is attributed to Mo segregation at the tip decreasing the chemical driving force and interface mobility for lengthening but not for thickening. This leads the formation of the stubby intragranular a plates. The empirical experimental information presented about the interactions between the alloying elements in both the a and b phases and at a/b interfaces is expected to provide useful and important guidance for the modelling of transformation kinetics and for the design of Ti alloys. Acknowledgements This work was supported the Australian Research Council (DP120100206), and with technical support and scientific input from the Australian Microscopy and Microanalysis Research Facility (AMMRF). R. Shi and Y. Wang would like to acknowledge the support by the US NSF DMREF program under Grant No. DMR-1435483. The authors acknowledge the use of JEOL ARM200F (ARC LIEF LE120100104) at UOW Electron Microscopy Centre. The authors are grateful to Prof. O.M. Ivasishin and Dr. D. Savvakin, IMP, NAS Ukraine, for supplying the material.
222
T. Li et al. / Journal of Alloys and Compounds 643 (2015) 212–222
The authors would like to thank Dr. Hao Chen at Tsinghua University for extremely valuable discussions. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom.2015.04. 143. References [1] G. Lutjering, J.C. Williams, Titanium, Springer, Berlin, 2007. [2] M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, Prog. Mater Sci. 54 (2009) 397–425. [3] R.R. Boyer, R.D. Briggs, J. Mater. Eng. Perform. 14 (2005) 681–685. [4] O.M. Ivasishin, V.M. Anokhin, A.N. Demidik, D.G. Savvakin, Key Eng. Mater. 188 (2000) 55–62. [5] O.M. Ivasishin, R.V. Teliovich, Mater. Sci. Eng. A 263 (1999) 142–154. [6] S. Semiatin, V. Seetharaman, I. Weiss, JOM 49 (1997) 33–39. [7] I. Weiss, S. Semiatin, Mater. Sci. Eng. A 243 (1998) 46–65. [8] I. Weiss, S. Semiatin, Mater. Sci. Eng. A 263 (1999) 243–256. [9] E.A. Brener, V.I. Mel’nikov, Adv. Phys. 40 (1991) 53–97. [10] M. Fleck, C. Hüter, D. Pilipenko, R. Spatschek, E.A. Brener, Philos. Mag. 90 (2010) 265–286. [11] C. Laird, H. Aaronson, Acta Metall. 17 (1969) 505–519. [12] C. Zener, J. Appl. Phys. 20 (1949) 950–953. [13] H. Aaronson, C. Laird, K. Kinsman, Phase Transformations, ASM, Metals Park, 1970. pp. 313–396. [14] K.R. Kinsman, E. Eichen, H.I. Aaronson, Metall. Trans. A 6 (1975) 303–317. [15] G. Spanos, R. Masumura, R. Vandermeer, M. Enomoto, Acta Metall. Mater. 42 (1994) 4165–4176. [16] O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Mater. Sci. Eng. A 405 (2005) 296–305. [17] O.M. Ivasishin, P.E. Markovsky, Y.V. Matviychuk, S.L. Semiatin, C.H. Ward, S. Fox, J. Alloys Comp. 457 (2008) 296–309. [18] G. Welsch, R. Boyer, E. Collings, Materials Properties Handbook: Titanium Alloys, ASM international, Materials Park, 1993. [19] T. Furuhara, T. Maki, Mater. Sci. Eng. A 312 (2001) 145–154. [20] M. Hillert, Introduction to paraequilibrium, in: Internal Report, Swedish Institute of Metals Research, Stockholm, 1953. [21] M. Hillert, Metall. Trans. A 6 (1975) 5–19. [22] D.E. Coates, Metall. Trans. 3 (1972) 1203–1212. [23] J. Kirkaldy, Can. J. Phys. 36 (1958) 907–916. [24] M. Enomoto, T. Yoshida, ISIJ Int. 31 (1991) 767–774.
[25] M.K. Miller, A. Cerezo, M. Hetherington, G. Smith, Atom Probe Field Ion Microscopy, Oxford Science Publications, Oxford, 1996. [26] B. Gault, M.P. Moody, J.M. Cairney, S.P. Ringer, Atom Probe Microscopy, Springer, New York, 2012. [27] S. Nag, R. Banerjee, R. Srinivasan, J.Y. Hwang, M. Harper, H.L. Fraser, Acta Mater. 57 (2009) 2136–2147. [28] S. Nag, R. Banerjee, J. Hwang, M. Harper, H. Fraser, Philos. Mag. 89 (2009) 535– 552. [29] H.P. Ng, A. Devaraj, S. Nag, C.J. Bettles, M. Gibson, H.L. Fraser, B.C. Muddle, R. Banerjee, Acta Mater. 59 (2011) 2981–2991. [30] S. Nag, Y. Zheng, R.E.A. Williams, A. Devaraj, A. Boyne, Y. Wang, P.C. Collins, G.B. Viswanathan, J.S. Tiley, B.C. Muddle, R. Banerjee, H.L. Fraser, Acta Mater. 60 (2012) 6247–6256. [31] M. Ahmed, T. Li, G. Casillas, J.M. Cairney, D. Wexler, E.V. Pereloma, J. Alloys Comp. 629 (2015) 260–273. [32]