Precipitation of Ti5Si3 phase in TiAl Alloys

Materials Science and Engineering A328 (2002) 113– 121 www.elsevier.com/locate/msea

Precipitation of Ti5Si3 phase in TiAl Alloys Fu-Sheng Sun a,b,*, F.H. Sam Froes b b

a Beijing Institute of Aeronautical Materials, P.O. Box 81, Beijing 100095, China Institute for Materials and Ad6anced Processes (IMAP), Uni6ersity of Idaho, Mines Building, Rm 321, Moscow, ID 83844 -3026, USA

Received 2 April 2001; received in revised form 6 June 2001

Abstract The nucleation, growth and interface structure of Ti5Si3 precipitates in Ti52Al48 – 3Si, Ti52Al48 – 3Si2M (M =Cr, V), and Ti52Al48 –3Si2Cr2V (at.%) alloys were studied. It is found that nucleation of Ti5Si3 precipitates depends on the dislocation types, k domain boundary or k/k lamellar boundaries. Most of the Ti5Si3 nucleates at 1/2[101( ] super-dislocation networks, particularly at the extended nodes of [101( ], where stacking faults exist, while some Ti5Si3 precipitates at 1/2[11( 0] dislocations. In addition, most Ti5Si3 are precipitating out heterogeneously at pseudo-twin type k/k *p and 120°-rotational order-fault type k/k *R lamellar boundaries, but no Ti5Si3 is found to precipitate at true-twin type k/k *T lamellar boundaries. The growth mechanism of Ti5Si3 in TiAl is also studied. © 2002 Elsevier Science B.V. All rights reserved. Keywords: TiAl; Ti5Si3; Precipitation; Nucleation; Interface

1. Introduction Si-containing TiAl alloys have been of considerable interest due to their significantly improved mechanical properties [1–4]. Addition of Si to TiAl alloys results in formation of the Ti5Si3 phase. Since the solubility of Si in k is limited at both room- and elevated-temperatures, with increasing Si content, the phase present are h2 + k “ h2 +k+ Ti5Si3 “k +Ti5Si3 in Ti52Al48 – xSi alloys [2]. Addition of 0.5 at.% Si to TiAl results in formation of fine Ti5Si3 particles precipitated at k/k and k/h2 lamellar boundaries. Addition of 2.0 – 6.0 at.% Si to TiAl results in large Ti5Si3 particles or whiskers formation in TiAl via powder metallurgy [5] or ingot metallurgy [6], respectively. In addition, fine Ti5Si3 were found to precipitate at k/k *p lamellar boundaries [7,8]. The physical and mechanical properties of single phase and polycrystalline Ti5Si3 compound were investigated [9 – 12]. Because Ti5Si3 with a D88 structure is a strengthening phase in TiAl, the addition of Si to TiAl not only significantly increases the oxidation resistance of Ti –Al –Nb –Si alloy [13], but also enhances the tensile strength and creep resistance of the Ti – 46Al – 1Cr –0.2Si (at.%) alloy [3,4]. * Corresponding author. E-mail address: [email protected] (F.-S. Sun).

The objective of the present study was to investigate the precipitation behavior of Ti5Si3 phase in Ti52Al48 – 3Si, Ti52Al48 –3Si2Cr, Ti52Al48 –3Si2V, and Ti52Al48 – 3Si2Cr2V (at.%) alloys, and the effects of dislocation types, k/k boundary types and k/k domain boundaries on the nucleation of Ti5Si3. The growth mechanisms of Ti5Si3 in the k phase or at k/k boundaries was also studied.

2. Experimental procedure TiAl alloys with nominal compositions of Ti52Al48 – 3Si, Ti52Al48 –3Si2Cr, Ti52Al48 –3Si2V, and Ti52Al48 – 3Si2Cr2V (at.%) were prepared by non-consumable electrode arc melting in an argon atmosphere, with each ingot being inverted and melted four times in order to obtain homogeneous compositions. High purity Ti sponge, Al (99.99%), Cr (99.99%) and Al –V, Al –Si master alloys were used. The ingots were HIP’d (Hot isostatic pressed) at 1200 °C/200 MPa for 3 h and then heat-treated at 1200 °C for 12 h. Samples (about 0.2 mm thick) were cut from the ingots and mechanically polished to a thickness of about 60 mm, and then electropolished using a standard twin-jet polishing method. The morphologies of the Ti5Si3 were examined

0921-5093/02/$ - see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 6 7 8 - 1

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on a H800 transmission electron microscope equipped with a double-tilt specimen holder (9 45°). For HRTEM, a JEM 2010 with a Link ISIS energy dispersive spectrometer (EDS) at 200 KV was used. A JEOL

JXA 8600 microscope equipped with electron probe microanalyzer (EPMA) was utilized to analyze the matrix compositions. X-ray diffraction was carried out at 35 kV and 25 mA using Cu Kh radiation at a scan rate of 0.2° per min.

3. Results

3.1. Crystal structure of TiAl and Ti5Si3 phase

Fig. 1. HRTEM images showing the Ti5Si3 phase in the Ti52Al48 – 3Si2Cr2V alloy from [11( 00] direction (a); and [1( 21( 3] direction (b).

Ti52Al48 –3Si, Ti52Al48 –3Si2Cr, Ti52Al48 –3Si2V, and Ti52Al48 –3Si2Cr2V alloys are comprised of k (an orthogonal L10 structure) and Ti5Si3 (a hexagonal D88 structure), with the compositions and lattice parameters of the Ti5Si3 and k phases varying with the addition of Cr and V. The lattice parameters of the k in Ti52Al48 – 3Si2Cr2V were a= 0.4024 nm, c =0.4091 nm, and c/ a = 1.0167, and those of the Ti5Si3, with a composition of 54.3 at.% Ti, 8.32 at.% Al, 28.7 at.% Si, 3.50 at.% Cr and 5.13 at.% V, were a= 0.7517 nm, c= 0.5136 nm, and c/a= 0.6833. Fig. 1 shows a HRTEM image of the Ti5Si3 in Ti52Al48 –3Si2Cr2V along a [11( 00] axis or [1( 21( 3] axis, indicating a 16H stacking sequence. Fig. 2 is a HRTEM image of Ti52Al48 –3Si2Cr, from which a pseudo-twin type k/k *p boundary with an almost straight interface (marked as G in Fig. 2a) and some misfit dislocations (marked as B in Fig. 2a), and a 120°-rotational order-fault type k/k *R boundary consisting of some segments of perfect boundary (marked as P in Fig. 2b) and interface defects (marked as D in Fig. 2b), are observed. Three types of k/k lamellar boundaries are found to exist in Si-bearing TiAl alloys, similar to those observed in TiAl polysynthetically twinned crystals [14,15].

3.2. Nucleation at dislocations

Fig. 2. HRTEM images illustrating k/k boundaries in Ti52Al48 – 3Si2Cr alloys, (a) a pseudo-twin type k/k*p boundary; and (b) a 120°-rotational order-fault type k/k *R boundary.

In Ti52Al48 –3Si2Cr and Ti52Al48 –3Si2V alloys, as shown in Fig. 3, the nucleation of Ti5Si3 in the k phase is heterogeneous, and Ti5Si3 generally precipitates at dislocations. Fig. 1a exhibits some normal dislocations with b=1/2[11( 0], with Ti5Si3 particles precipitated on them (marked with arrow). Detailed observation in Fig. 2a (marked with arrow) shows that a small Ti5Si3 particle nucleates at the end of an normal dislocation with b= 1/2[11( 0], indicating that the normal dislocations are preferential sites for the nucleation of Ti5Si3. Similarly, most disordered alloys such as Al [16], Fe [17,18], Mg [19], and Ni [20] show heterogeneous nucleation at dislocations. For example, Mg2Al3 particles or Mg2Si particles were precipitated on a number of dislocations in Al–10Mg, and Al–0.9Mg–0.43Si alloys. In disordered Al alloys, dislocations with b= 1/2[11( 0] are crystallographically equivalent to those with b =1/ 2[101( ]. These dislocations play the same role in the

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an ordered structure. Consequently, the effect of these two types of dislocations on nucleation of Ti5Si3 is different. Fig. 3c and d show the precipitation of Ti5Si3 at dislocation networks in the k phase in Ti52Al48 –3Si2V. These Ti5Si3 particles nucleate at the extended nodes of dislocation networks (marked with arrow in Fig. 3d), which are superdislocations with b=1/2[101( ] dissociated to partial dislocations (marked S in Fig. 3d). This indicates that Ti5Si3 particles prefer to precipitate at stacking faults of k. Similarly, stacking faults were preferential sites for heterogeneous nucleation in other disordered alloy systems. A previous study showed that TiB2 precipitates nucleated at extrinsic faults of a dislocation network in boron-doped TiC matrix [21]. Ag2Al phase in Al alloy precipitated at extended dislocation nodes in a dislocation network [22]. The present work shows that the majority of Ti5Si3 particles are precipitated at those dislocation networks consisting of superdislocations with b=1/2[101( ] and their dissociated partial dislocations. The number of Ti5Si3 particles precipitated at 1/2[11( 0] normal dislocation is more less than that at 1/2[101( ] superdislocations. This will be discussed in section IV.

Fig. 3. TEM bright field micrographs showing the Ti5Si3 phase precipitated at normal dislocations with b= 1/2[11( 0] in Ti52Al48 – 3Si2Cr alloy (a and b); and the precipitates are also present at 1/2[101( ] superdislocation networks (c); superdislocation extended nodes; and (d) in Ti52Al48 –3Si2V alloy.

nucleation of particles. However, dislocations with b= 1/2[11( 0] in TiAl are normal dislocations, while those with b= 1/2[101( ] are superdislocations, since TiAl has

Fig. 4. TEM micrographs revealing the precipitation of Ti5Si3 particles at different k/k boundaries, such as true-twin type k/k *, T pseudotwin type k/k *, 120°-rotational order-fault type k/k *R lamellar p boundaries, in Ti52Al48 – 3Si 2Cr2V alloy.

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3Si2Cr2V. At some true-twin type k/k *T lamellar boundaries (marked T in Fig. 4a), no Ti5Si3 particle was precipitated, similar to the precipitate-free zone in Al alloys [23]. However, at other k/k lamellar boundaries which are pseudo-twin type k/k *p boundaries (marked P in Fig. 4a), numerous elongated Ti5Si3 particles are precipitated along these boundaries at regular interval. In addition, a needle-like Ti5Si3 is found to precipitate at k/k boundaries that are 120°-rotational order-fault type k/k *R lamellar boundaries (marked R in Fig. 4a), which are often observed in V-bearing TiAl alloys. Fig. 4b exhibits a Ti5Si3 particle precipitated at a k/k boundary, where some boundary dislocations exist. It is also found that numerous Ti5Si3 particles precipitate at the boundaries of k ordered domains that consist of 1/2[101( ] dislocations (marked D in Fig. 4a), and there is a specific angle between the growth direction of Ti5Si3 and the boundaries. A similar dependence of nucleation rate on the types of boundaries has also been found in Al alloys [24]. In Al –4Cu alloys, numerous CuAl2 precipitates were found to nucleate at low-angle boundaries and high-angle random boundaries, whereas very few CuAl2 particles precipitated at a high-angle boundary with a misorientation near a coincidence site relation. The precipitation behaviors of CuAl2 are dependent upon boundary energy and the orientation between precipitates and adjoining grains. In contrast, the nucleation and growth behavior of Ti5Si3 in TiAl is more complex, since various ordered boundaries, such as true-twin type k/k *, T pseudo-twin type k/k *, p and 120°-rotational order-fault type k/k *R lamellar boundaries, are present in TiAl. The type of boundary significantly affects the precipitation behavior of Ti5Si3 particles in TiAl alloys. Fig. 5. HRTEM images showing Ti5Si3/k interfaces in Ti52Al48 – 3Si2Cr alloy (a and c), and that in Ti52Al48 –3Si2V alloy (b).

Fig. 6. HRTEM image exhibiting the detailed morphology of k/Ti5Si3 interface with ledges.

3.3. Nucleation at k/k boundary Fig. 4 shows the heterogeneous nucleation and growth of Ti5Si3 particles at k/k boundaries in Ti52Al48 –

3.4. Diffusion growth of Ti5Si3 HRTEM images of the Ti5Si3 precipitates in Ti52Al48 –3Si2Cr and Ti52Al48 –3Si2V are shown in Fig. 5. A rectangular Ti5Si3 particle, precipitated in k, is shown in Fig. 5a, with relative straight k/Ti5Si3 interfaces, but with a curvature at the upper edge. Close examination of the k/Ti5Si3 interface is shown in Fig. 6. Some ledges occur along the interface at regular interval (marked with arrows in Fig. 6), with the distance between two adjacent ledges almost the constant (about 10.4 nm). Further observation (Fig. 7a) shows that the ledge height is equal to the height of four lays of (111)k planes. As shown in Fig. 7b (marked with arrows), misfit dislocations at a semicoherent interface and big ledges at Ti5Si3 are observed in the interface corresponding to the upper edge of the Ti5Si3 particle in Fig. 5a. This indicates that the growth of these Ti5Si3 particles is controlled by a ledge mechanism [25]. Selected area electron diffraction (SADP) confirms the orienta-

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tion relationship of the Ti5Si3 and k phases as (111( )k // (0002)Ti5Si3 and [154]k //[1( 21( 0]Ti5Si3. Another type of k/Ti5Si3 interface in Ti52Al48 – 3Si2V is shown in Fig. 5b, with straight k/Ti5Si3 interfaces along two sides of the Ti5Si3 particles (marked as S), where no ledge exists. Here the k/Ti5Si3 interface (marked as H) is spherical, indicating another growth mechanism as most rounded Ti5Si3 particles have. The orientation relationship of the Ti5Si3 and k phases is found to be (101( )k//(101( 0)Ti5Si3, and [141]k //[1( 21( 6]Ti5Si3. A Ti5Si3 particle is precipitated at a k/k boundary of Ti52Al48 –3Si2Cr, as shown in Fig. 5c, characterized by almost straight k/Ti5Si3 interfaces along both sides of the Ti5Si3 particle (marked as M and N). The orientation relationship of the Ti5Si3 and k phases is (101( )k// (101( 0)Ti5Si3, indicating an incoherent interface. This implies that the growth of these needle-like Ti5Si3 precipitated at a k/k boundary is controlled by interfacial diffusion, similar to the grain boundary allotriomorphs of CuAl2 phase found in Al– 4Cu alloys [26].

4. Discussion

4.1. Heterogeneous precipitation beha6ior of Ti5Si3 This work indicates that the nucleation of Ti5Si3 in TiAl is heterogeneous, suggesting that the free energy

Fig. 7. HRTEM image showing (a) some ledges; and (b) misfit dislocations at a semicoherent Ti5Si3/k interface.

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for Ti5Si3 homogeneous nucleation is high. In contrast, the precipitation of Ti3AlC and Ti3AlN particles in TiAl was homogeneous [27]. After quenching from 1150 °C, needle-like Ti3AlC and Ti3AlN phase were precipitated during aging at 900 °C. Heterogeneous precipitation of Ti5Si3 also occurred in another Si-bearing k alloy [7]. To understand the difference between the heterogeneous precipitation of Ti5Si3 in k and the homogeneous nucleation of Ti3AlC and Ti3AlN in k, a simplified model for a homogeneous semicoherent nucleation [28] will be considered. The free energy of nuclei formation is estimated as following: DG ok = 16y(|c + |st)3/3(DGw + V)2

(1)

where DG is the free energy for nuclei formation, DGw is the chemical driving force for nucleation, |st is the structural energy of a semicoherent boundary, |c is the chemical energy of a semicoherent boundary, and V is any dilatational strain energy not relieved by the interfacial dislocations. Semicoherent k/Ti5Si3 interfaces were found in the present study. Unlike the situation in disordered alloys, the chemical energy of k/Ti5Si3 semicoherent interfaces (|c) plays an important role in the precipitation of Ti5Si3. The energy of coherent boundary |c is calculated by the following equation [29]. o k

|c = [DE oNsZs(C−C o)2]/N oZl

(2)

where Ns is the atoms/unit area of interface, Zl is the lattice coordination number, Zs is the number of AB bonds per atom of A or B in the interface, DE o is the heat of solution (per gram atom) of B in an infinitely dilute solution of B in A, N o is Avogadro’s number, and C-C o is the difference in composition between the nucleus and matrix. The chemical energy of k/Ti5(SiAl)3 interfaces depends on the bond types and number Zs. For the phase transformation of TiAl“ Ti5(SiAl)3, the number of TiSi and AlSi bonds increase, and the number of TiAl, TiTi and AlAl bonds decrease. In contrast to disordered alloys, this increases the chemical interface energy. The effect of the difference between the Ti5(SiAl)3 nuclei and the k matrix compositions (C-C o) on the chemistry of the semicoherent k/Ti5(SiAl)3 energy boundary is significant. With increasing C-C o, the chemical energy of the k/Ti5(SiAl)3 semicoherent boundary significantly increases. Table 1 shows the compositions of k and Ti5Si3 in TiAl alloys. With respect to the Si and Al contents in Ti5Si3 and k, (C-C o)Si and (C-C o)Al are about 24.0–28.7 and 35.0– 40.0 at.%, respectively, which increases the chemical energy of k/Ti5(SiAl)3 semicoherent boundaries (|c). In contrast, the (C-C o)c and (C-C o)Al for TiAl“ Ti3AlC phase transformation are about 15.0–20.0 and 15.0–

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Table 1 Compositions of k and Ti5Si3 phase (at.%) Alloys

Phases

Ti

Al

Si

Ti52Al48–3Si

k phase Ti5Si3 (C-C o) k phase Ti5Si3 (C-C o) k phase Ti5Si3 (C-C o) k phase Ti5Si3 (C-C o)

50.1 9 0.5 62.8 9 0.5 (12.7 9 0.5) 52.9 9 0.5 63.1 9 0.5 (10.2 9 0.5) 47.9 9 0.5 56.6 9 0.5 (8.7 9 0.5) 48.8 9 0.5 54.3 9 0.5 (5.50 9 0.5)

49.2 9 0.5 9.20 9 0.5 (−40.0 9 0.5) 44.8 9 0.5 9.51 9 0.5 (−35.3 9 0.5) 48.8 9 0.5 9.95 9 0.5 (−38.9 9 0.5) 45.9 9 0.5 8.32 9 0.5 (−37.6 9 0.5)

0 28.0 9 0.5 (28.0 9 0.5) 0 24.0 9 0.5 (24.0 9 0.5) 0 26.4 9 0.5 (26.4 9 0.5) 0 28.7 9 0.5 (28.7 9 0.5)

Ti52Al48–3Si2Cr

Ti52Al48–3Si2V

Ti52Al48–3Si2Cr2V

Cr

V

2.32 9 0.2 2.64 9 0.2 (0.32 9 0.2)

2.64 9 0.2 3.50 9 0.2 (0.86 9 0.2)

2.22 9 0.2 7.01 9 0.2 (4.79 9 0.2) 2.67 9 0.2 5.13 9 0.2 (2.46 9 0.2)

Table 2 Mismatch of (111)k and (0002)Ti5Si3 planes Alloys

k phase

d (nm)

Ti5Si3 phase

d (nm)

l

1/l

Ti52Al48–3Si Ti52Al48–3Si2Cr Ti52Al48–3Si2V Ti52Al48–3Si2Cr2V

(111) (111) (111) (111)

0.2334 0.2333 0.2338 0.2336

(0002) (0002) (0002) (0002)

0.2625 0.2612 0.2581 0.2568

0.1109 0.1068 0.0941 0.0904

9.021 9.363 10.63 11.07

Table 3 Mismatch of (001)k and (001)Ti3AlC planes Alloys

k phase

d (nm)

Ti3AlC phase

d (nm)

l

1/l

(Ti51Al49)99.5C0.5

(001)

0.4086

(001)

0.4156

0.0171

58.48

20.0 at.%, respectively, lower than that of the phase transformation of TiAl“Ti5(SiAl)3. The diffusivity of carbon, an interstitial element, in k is higher than that of silicon. Therefore, the chemical energy of k/ Ti5(SiAl)3 interfaces is higher than that of k/Ti3AlC interfaces. The structural energy of the k/Ti5(SiAl)3 semicoherent boundary (|st) has an effect on the free energy for Ti5(SiAl)3 nuclei formation in k. The orientation relationship of the k/Ti5(SiAl)3 semicoherent interfaces is (111( )k //(0002)Ti5Si3 and [154]k //[1( 21( 0]Ti5Si3, while the orientation relationship of the k/Ti3AlC interfaces is [100]Ti3AlC//[100]k and [001]Ti3AlC//[001]k [27]. Table 2 shows the mismatch of (111)k and (0002)Ti5Si3 planes, where the mismatch is defined as l = (d1 −d0)/d0. The mismatch (l) in Ti52Al48 – 3Si, Ti52Al48 – 3Si2Cr, Ti52Al48 –3Si2V, and Ti52Al48 – 3Si2Cr2V are about 0.0904–0.1109 (Table 2), much higher than that of (001)k and (001)Ti3AlC in (Ti51Al49)99.5C0.5 alloy (l= 0.0171 in Table 3). The 1/l in Tables 2 and 3 represents the periodicity of the misfit dislocations at the interfaces. With increasing mismatch (l), the structural en-

ergy of the semicoherent boundary (|st) increases, therefore the structural energy of the k/Ti5(SiAl)3 semicoherent boundary (|st) is much higher than that of k/Ti3AlC semicoherent boundary. Considering the chemical interface energy |c and the structural energy |st in Eq. (1), the free energy for Ti5Si3 nucleus formed in k is much higher than that for Ti3AlC nucleus formed in k, hence, the precipitation of Ti5Si3 in k is heterogeneous, and the precipitation of Ti3AlC or Ti3AlN in k is homogeneous.

4.2. Effect of dislocation types on the nucleation and growth of Ti5Si3 Most Ti5Si3 particles nucleate at 1/2[101( ] superdislocations, and some Ti5Si3 particles nucleate at 1/2[11( 0] normal dislocations. One explanation of this phenomenon is a decrease of strain energy for Ti5Si3 nucleation at dislocations. Considering the nucleation of Ti5Si3 at dislocations, a dimensionless parameter is proposed [30]: x= DGwvb 2/2y 2| 2

(3)

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where DGw is the chemical driving force, v is the shear modulus, b is the Burgers vector of dislocations, and | is the strain energy. For a k/Ti5Si3 semicoherent interface, |=|c + |st, where |st is the structural energy, and |c is the chemical energy of the semicoherent boundary. When x \1, Ti5Si3 may nucleate at dislocations spontaneously, when x B 1, a nucleation barrier for Ti5Si3 nucleation exists, and if 0.4Bx B0.7, most Ti5Si3 particles precipitate at dislocations. With increasing x, the nucleation of Ti5Si3 at dislocations becomes easier. The presence of dislocations decreases the structural energy of the semicoherent boundary |st, and provides nucleation sites for Ti5Si3 due to the increase of x. Normal dislocations with b =1/2[11( 0] are usually present as single dislocations in k, in contrast superdislocations with b = 1/2[101( ] occur as dislocation pairs or dislocation networks (Fig. 3). In addition, the Burgers vector of 1/2[11( 0] normal dislocation is smaller than that of 1/2[101( ] superdislocations, as shown in Table 4. With respect to the structural energy of semicoherent boundary |st, the dislocation networks of 1/2[101( ] superdislocations are more effective in decreasing the structural energy than 1/2[11( 0] normal dislocations. The dislocation networks of 1/2[101( ] superdislocations have a larger influence in decreasing in the chemical energy |c than 1/2[11( 0] normal dislocations. Ti5Si3 particles were often precipitated at the extended nodes of 1/2[101( ] superdislocations (Fig. 3c and d). Previous investigations showed that the dissociation of superdislocations with Burgers vector b =[011] into Shockley type partial dislocations with b =1/6[121] and 1/6[1( 12] was as following [31– 33]: [011]= 1/6[1( 12]+SISF +1/6[121] + APB + 1/6[1( 12] +CSF + 1/6[121]

(4)

In addition to Shockley type partials, three types of planar faults such as APB, CSF and SISF are present in k phase. The present work confirms that [101( ] superdislocations are dissociated into superpartials (Fig. 3d). Moreover, the relationship of ECSF \ESISF \EAPB was found with respect to the stacking fault energy of CSF, SISF and APB [31]. Ti5Si3 particles are readily precipitated at the extended nodes of the [101( ] superdislocations, where stacking faults exist. Ti5Si3 has a hexagonal (hcp) D88 structure, and k has an orthogonal

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L10 structure (close to the fcc structure). The extrinsic stacking faults (CSF) of k are close to the structure of Ti5Si3 (hcp), therefore Ti5Si3 nucleates at the extrinsic stacking faults in k, where (0002)Ti5Si3//(111)k, and a four-atomic layer of Ti5Si3 forms in k (Fig. 7a). The precipitation of Ti5Si3 at extrinsic stacking faults in k decreases the chemical energy of k/Ti5Si3 semicoherent boundary. Therefore, 1/2[101( ] superdislocations are the preferential sites for Ti5Si3 nucleation.

4.3. Effect of boundary energy on the precipitation beha6ior of Ti5Si3 The precipitation of Ti5Si3 at k/k boundaries is heterogeneous, depending on the k/k boundary types (Fig. 4a and b), which is associated with nucleation of Ti5Si3 at k/k boundaries. The free energy change on forming an incoherent Ti5Si3 at k/k boundary can be described as following [28]: DG 0r = 2yr 3 DGw (2−3 cos q+ cos3 q)/3 − yr 2 sin2 q|ii + 2yr 2|hi (1−cos q)

(5)

where DGw represents the chemical driving force, |ii is the surface energy of a k/k lamellar boundary, |hi is the surface energy of k/Ti5Si3 boundary, r is the radius of spherical Ti5Si3 particles and q is the angle between k and Ti5Si3. The precipitation of Ti5Si3 at k/k lamellar boundaries depends on the k/k lamellar boundary energy and the k/Ti5Si3 interface energy. With an increase in the k/k lamellar boundary energy, the precipitation of Ti5Si3 at the k/k lamellar boundaries becomes easier. There are three types of k/k lamellar boundaries in Si-bearing TiAl alloys. The first one is a true-twin type k/k *T lamellar boundary, where no planar defect exists. The second one is a pseudo-twin type k/k *p lamellar boundary, where the mismatch between [11( 0]TiAl and [011( ]TiAl causes the formation of planar defects (Fig. 2a). The third one is a 120°-rotational order-fault type k/k *R lamellar boundary, with some planar defects (Fig. 2b). A previous investigation indicated that the surface energies of these three types of k/k lamellar boundaries are different [15], with a ratio of kT:kP:kR = 1:3:2 or 1:7:6, where kT, kP and kR corresponding to the surface energies of true-twin type k/k *, T pseudo-twin type k/k * p and 120°-rotational order-fault type k/k *R lamellar

Table 4 Lattice parameters and burgers vectors of different dislocations in k-TiAl based alloys Alloy (at.%)

Ti52Al48–3Si Ti52Al48–3Si2Cr Ti52Al48–3Si2V Ti52Al48–3Si2Cr2V

k phase c (nm)

a (nm)

b =1/2[11( 0]

b2

b =1/2[101( ]

b2

0.4094 0.4088 0.4119 0.4091

0.4020 0.4018 0.4016 0.4024

0.2843 0.2841 0.2840 0.2845

0.0808 0.0807 0.0806 0.0810

0.2869 0.2866 0.2876 0.2869

0.0823 0.0821 0.0827 0.0823

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Table 5 Spacing parameters of k phase and mismatch of the [1( 01] and [11( 0] planes Alloy

a1 of [1( 01] (nm)

a1 of [11( 0] (nm)

l (nm)

1/l

a2/l (mm)

Ti52Al48–3Si2Cr Ti52Al48–3Si2V Ti52Al48–3Si2Cr2V

0.2866 0.2876 0.2869

0.2841 0.2840 0.2845

0.00880 0.01268 0.00843

113.6 78.8 118.5

0.032 0.022 0.033

boundaries, respectively. The true-twin type k/k *T lamallar boundary is a precipitation-free k/k boundary since it has the lowest boundary energy (Fig. 4a). The majority of the Ti5Si3 particles are precipitated at pseudotwin type k/k *p and 120°-rotational order-fault type k/k *R lamellar boundaries, since there boundary energies are high, thus decreasing the free energy for Ti5Si3 nucleation at these k/k *p and k/k *R lamellar boundaries. The effects of pseudo-twin type k/k *p and 120°-rotational order-fault type k/k *R lamellar boundaries on the precipitation of Ti5Si3 differ. At a pseudo-twin type k/k *p lamellar boundary (Fig. 2a), misfit dislocations occur due to the mismatch of the [11( 0] and [101] planes. Table 5 shows the spacing parameters of the k, and the mismatch of [11( 0] and [101] planes, which the mismatch l =[a1 − a2]/a2, where a1 and a2 represent the spacing of the planes, and ratio a2/l represents the distance between adjacent misfit dislocations. The a2/l at pseudo-twin type k/k *p lamellar boundaries is about 0.020–0.030 mm, and Ti5Si3 particles are precipitated at the misfit dislocations in the pseudo-twin type k/k *p lamellar boundaries at a periodical distance (Fig. 4a). In contrast, planar defects exist at 120°-rotational order-fault type k/k *R lamellar boundaries (Fig. 2b). The planar defects consist of a two-atomic-layer stacking fault (in the hcp structure) at the boundary [14], and favor Ti5Si3 nucleation, therefore needle-like Ti5Si3 particles easily nucleate at these planar defects (Fig. 4a).

4.4. Growth mechanism of Ti5Si3 Differing growth behaviors of Ti5Si3 in k were observed. The first one is associated with needle-like Ti5Si3 precipitates at k/k boundaries, with a k/Ti5Si3 incoherent interface (Fig. 4a and Fig. 5c). The growth of the needle-like Ti5Si3 along k/k lamellar boundary is fast, and the growth rate of Ti5Si3 is controlled by the interfacial diffusion of Si at the k/k boundaries, since the interfacial diffusivity of Si is higher than the volume diffusivity of Si in the k matrix, similar to that of CuAl2 in Al alloy [26]. A second type of growth behavior of the Ti5Si3 was shown in Fig. 5b. The orientation relationship of the Ti5Si3 and k phases is incoherent, (101( )k//(101( 0)Ti5Si3, and [141]k //[1( 21( 6]Ti5Si3. The growth of the Ti5Si3 in the k matrix is controlled by the volume diffusion of Si in k.

The growth of Ti5Si3 shown in Fig. 5a is controlled by a linear ledge mechanism. The distance between adjacent ledges is quite uniform at about 10.4 nm (Fig. 6), and the ledge height is also consistent at four (111)k planes (Fig. 7a). In addition, misfit dislocations (marked with an arrow) exist at k/Ti5Si3 semicoherent interfaces (Fig. 7b). This suggests that k“ Ti5Si3 transformation is similar to the fcc“ hcp transformation found in Al–15%Ag alloys [34]. In Si-bearing TiAl alloys, Shockley partial dislocations move on a height of four (111)k planes and to provide the conversion of L10 to D88 structure, since in the c-direction, a unit cell of Ti5Si3 consists of four atomic layers of (0002). Antiphase boundary (APB) and stacking faults (SF) exist in ordered Ti5Si3. If the growth of Ti5Si3 takes place by a ledge mechanism consisting of one atomic plane height, the antiphase boundary or stacking faults would occur at the k/Ti5Si3 interface, and thus increase the interface energy. This indicates that the transformation of an ordered phase by a linear ledge mechanism is different from that of a disordered phase due to the presence of ordered planar defects. A similar phenomenon occurred at the interface between Ti3Al and TiAl phases, with a ledge height of two (111)TiAl planes [14]. The growth of Ti5Si3 by a linear ledge mechanism is controlled by pipe diffusion of Si in the core of misfit dislocations. For the k“ Ti5Si3 transformation, the ledges migrate by the movement of Shockley partials at a diffusion-controlled rate. The pipe diffusivity of Si in the core of misfit dislocations is higher than the volume diffusivity of Si in k, but lower than the interfacial diffusivity of Si at k/k boundaries.

5. Conclusions In Si-bearing TiAl alloys, the precipitation of Ti5Si3 in k is heterogeneous, and the nucleation of Ti5Si3 depends on the dislocation type. The most preferential sites for Ti5Si3 to nucleate are 1/2[101( ] superdislocations, and some Ti5Si3 particles nucleate at 1/2[11( 0] normal dislocations. Ti5Si3 particles are readily precipitated on the extended nodes or dislocation networks of [101( ] superdislocations, where stacking faults exist. The precipitation of Ti5Si3 particles at k/k boundaries is heterogeneous. Comparing three types of k/k lamellar boundaries, a true-twin type k/k *T lamellar boundary is

F.-S. Sun, F.H.S. Froes / Materials Science and Engineering A328 (2002) 113–121

a Ti5Si3 precipitate-free k/k boundary because of its low boundary energy. Ti5Si3 particles are precipitating out extensively at pseudo-twin type k/k *p and 120°-rotational order-fault type k/k *R lamellar boundaries, where planar defects exist. In addition, Ti5Si3 particles are also present at misfit dislocations on the pseudo-twin type k/k *p lamellar boundaries; there misfit dislocations being located at a uniform periodical distance. Needle-like Ti5Si3 particles nucleate at the planar defects of the 120°-rotational order-fault type k/k *R lamellar boundaries. Different growth mechanisms of Ti5Si3 particles in k were observed. The growth of needle-like Ti5Si3 particles at k/k boundaries is controlled by interfacial diffusion of Si, with an incoherent k/Ti5Si3 interface. For most rectangular Ti5Si3 precipitates in k, growth is controlled by a linear ledge mechanism, and the k/ Ti5Si3 interface is semicoherent, with an orientation relationship (1( 11( )k //(0002)Ti5Si3 and [154]k //[1( 21( 0]Ti5Si3. In addition, the growth of rounded Ti5Si3 particles in the k matrix is controlled by volume diffusion of Si, where the Ti5Si3/k interface is incoherent.

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