Effect of withdrawal rate on microstructure and lattice misfit of a Ni3Al based single crystal superalloy

Journal of Alloys and Compounds 592 (2014) 164–169

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Effect of withdrawal rate on microstructure and lattice misfit of a Ni3Al based single crystal superalloy Cheng Ai, Shusuo Li ⇑, Heng Zhang, Lei Liu, Yue Ma, Yanling Pei, Shengkai Gong School of Materials Science and Engineering, Beihang University, No. 37 Xueyuan Road, Beijing 100191, PR China

a r t i c l e

i n f o

Article history: Received 13 September 2013 Received in revised form 21 December 2013 Accepted 31 December 2013 Available online 10 January 2014 Keywords: High-temperature alloys Microstructure Liquid diffusion coefficient Lattice misfit

a b s t r a c t A newly developed Re-containing Ni3Al based single crystal superalloy was prepared by seeding technique. We investigated the effect of withdrawal rate (1–400 lm/s) on the microstructure and the c/c0 lattice misfit. With increasing withdrawal rate, the as-cast microstructure changed from planar to cellular and then to dendritic, accompanied by the significant refinement of dendrite arm and c0 phases. The morphology of c0 phases also transformed in response to different withdrawal rates, which was ascribed to the influence of the c/c0 lattice misfit. In association with the evolution of the microstructures, the liquid diffusion coefficients of Ni and Ni3Al based superalloys were estimated by Kurz–Fisher model. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction Ni based superalloys have been widely used in various aerospace turbine blades due to their excellent comprehensive properties [1–3]. In order to further improve high temperature mechanical properties and decrease density, Ni3Al based superalloys were developed based on Ni based superalloys in the 1980s, which had the most important characteristic of high Al content [4,5]. Recently, a Ni3Al based single crystal superalloy IC6SX was reported to show good high temperature stress rupture properties [6,7]. The as-cast microstructures in both Ni based and Ni3Al based superalloys are significantly influenced by the solidification parameters, such as the temperature gradient, withdrawal rate and cooling rate [8–11]. On one hand, the dendrite arm spacings and c0 phases can be refined when produced at higher temperature gradient and withdrawal rate [8–15]. On the other hand, the morphology of c0 phases also changes with increasing withdrawal rate, but the effect of withdrawal rate on the morphological evolution of c0 phases was different according to different alloys. c0 phases changed from cubic to sphere in Alloy D [13], but changed from sphere to cubic in NASAIR 100 [15] with increasing withdrawal rate. The reason for these different behaviors of morphological evolutions is still unclear. The as-cast microstructure is also influenced by the material constants, for example, the dendrite arm spacing kd / liquid diffusion coefficient D0:25 [16,17]. The liquid diffusion coefficient is a L ⇑ Corresponding author. Tel.: +86 10 82314488; fax: +86 10 82338200. E-mail addresses: [email protected] (C. Ai), [email protected] (S. Li). 0925-8388/$ - see front matter Ó 2014 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.12.262

very important thermophysical property of superalloys and greatly useful in the simulation of dendritic growth [18], grain structure [19], microsegregation, macrosegregation and freckle formation [20] in directional solidification of superalloys. Nevertheless, to the authors’ knowledge, the determinations of DL in previous literatures were mostly based on experience, such as SRR99 (DL = 1.3  109 m2/s [17]), CMSX-2 (DL = 3.6  109 m2/s [20]), PWA1484 (DL = 2  109 m2/s [21]), CMSX-4 (DL = 3.6  109 m2/s [22], DL = 1  109 m2/s [23]), René N5 (DL = 5  109 m2/s [24]), and the investigation on the influence of individual alloying elements on DL is limited. Using the correlation between the microstructures and diffusion coefficient with respect to different withdrawal rate (Kurz–Fisher model), a more accurate estimation of DL is hopeful to be achieved. A new Re containing Ni3Al based SC superalloy was designed and examined in this report. The influence of withdrawal rate on as-cast microstructure was characterized and the effect of misfit on the morphological evolution of c0 phases with increasing withdrawal rate was also discussed. Besides, the liquid diffusion coefficients of Ni and Ni3Al based SC superalloys were estimated based on Kurz–Fisher model.

2. Experimental The experimental alloy used in this work was a Ni3Al based SC superalloy with the nominal composition of Ni–7.8Al–10.5Mo–2Cr–3Ta–1.5Re (wt.%). The master alloys were cut into £9*155 mm rods and the h0 0 1i orientation seeds were cut into £9*35 mm rods. Directional solidification experiments were carried out by using the bottom-seeding technique in a modified liquid metal cooling (LMC) apparatus and Ga-In liquid metals were used as coolant. The heating temperature was 1615 °C, and single crystals were produced at 7 withdrawal rates (V): 1, 10, 50,

C. Ai et al. / Journal of Alloys and Compounds 592 (2014) 164–169 100, 200, 300 and 400 lm/s (namely SC1, SC10, SC50, SC100, SC200, SC300 and SC400 sample). The temperature gradient at the solid/liquid interface was measured to be 170 K/cm. The cross-sections and longitudinal sections of SC samples were cut at a distance of 90 mm from the bottom of seed to avoid the influence of seed on the microstructure feature, and subsequently etched with 4 g CuSO4, 20 ml HCl and 20 ml H2O. The microstructure was characterized by an Olympus BX51M optical microscopy (OM) and a Quanta 200F field emission-scanning electron microscopy (FE-SEM). The Rigaku SmartLab X-ray diffractometer (XRD) with Cu Ka radiation was used to determine the lattice constants of c and c0 phases on the cross-section of the SC1 sample. In order to measure the chemical compositions of c and c0 phases, a JEM-2100 transmission electron microscopy (TEM) with Oxford INCA energy dispersive spectrometer (EDS) was used, and at least 5 points were respectively detected on c and c0 phases. The TEM samples were cut from the cross-section of experimental alloys and electrochemically thinned in a solution of 15 ml perchloric acid and 85 ml alcohol at 20 °C. The quantitative analysis of microstructures was carried out in the SISC ias8 system. Primary dendrite arm spacings were measured by triangle method [25], and at least 100 primary dendrite arms and 50 secondary dendrite arms were measured for each specimen. The mean dimension of c0 phases was determined by r = (A/N)0.5, where A is the total area of c0 phases and N is the number of c0 phases in the SEM picture.

165

DT = TL  TS (TL and TS is liquids temperature and solidus temperature of experimental alloy [16]), k is equilibrium segregation coefficient and can be estimated as k  (TNi  TL)/(TNi  TS) [26] (TNi is the melting point of pure Ni). Table 2 shows kd, G, V, DT and k of different Ni and Ni3Al based single crystal superalloys. As the relationship between temperature gradients, withdrawal rates and primary cellular arm spacings were not well matched with the Kurz–Fisher model [21], only the data of primary dendrite arm spacings, temperature gradients and withdrawal rates are listed in Table 2, and DL was calculated by the following method: Take the logarithm of both sides of Eq. (1), then

lnkd ¼ ln4:3 þ 0:25lnðCDT=kÞ þ 0:25lnDL  0:5lnG  0:25lnV

ð2Þ

Further, Eq. (2) was changed into

y¼xþb

ð3Þ

3. Results and discussions 3.1. As-cast microstructure and liquid diffusion coefficient Fig. 1 shows the OM images of cross-sectional and longitudinal sectional microstructures at different withdrawal rates. Planar, cellular and dendritic microstructures were observed at the withdrawal rate of 1 lm/s, 10 lm/s and 50–400 lm/s, respectively. The undercooling near the solid/liquid interface appears and then increases with increasing withdrawal rate [16]. Thus the as-cast microstructure of experimental alloy gradually evolved from planar (Fig. 1a) to cellular (Fig. 1b and e), coarse dendritic (Fig. 1c) and fine dendritic (Fig. 1d and f). The values of primary dendrite/ cell arm spacings (k1) and secondary dendrite arm spacings (k2) at different withdrawal rates are listed in Table 1. The relationship between kd and DL, G, V followed Kurz–Fisher model [16]:

kd ¼ 4:3ðCDL DT=kÞ

0:25

G0:5 V 0:25

ð1Þ 7

where C is Gibbs–Thomson coefficient and C = 10 km [21], DL is liquid diffusion coefficient, DT is solidification range and

where y = lnkd, x = 0.5 ln G0.25 ln V and b = ln 4.3 + 0.25 ln(CDT/k) + 0.25 ln DL. Then, a linear fitting of equation into the data of kd with G and V was performed and the value of b was obtained. Thus the value of DL was calculated in terms of its relation with b. Fig. 2 shows the liquid diffusion coefficient (DL) of Ni and Ni3Al based SC superalloys, among which the liquid diffusion coefficients of 1st and 2nd Ni based superalloys were at the magnitude of 109 m2/s and in accord with previous literature [17,20–24]. The concentrations of Al and Re elements are also exhibited in Fig. 2. Al and Re has the fastest and slowest interdiffusion coefficient between pure Ni, respectively [27], and therefore the concentration of these two elements should be closely related with DL. Comparing DL of different generation Ni based SC superalloys, it was found that with increasing Re content (from 1st to 2nd and 3rd generation), DL decreased. In addition, as shown in Fig. 2, the liquid diffusion coefficients of Ni3Al based SC superalloys were one order of magnitude higher than those of Ni based SC superalloys, reached 108 m2/s magnitude, indicating the strongly positive correlation between Al content and the liquid diffusion coefficient.

Fig. 1. OM images of as-cast microstructures at different withdrawal rates: cross-section: (a) 1 lm/s, (b) 10 lm/s, (c) 50 lm/s, (d) 400 lm/s and longitudinal section, and (e) 10 lm/s, (f) 400 lm/s.

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Table 1 Values of primary dendrite/cell arm spacings (k1) and secondary dendrite arm spacings (k2) at different withdrawal rates (V). V (lm/s)

10

50

100

200

300

400

k1 (lm)

277.4

217.3

171.9

163.5

155.0

150.9

k2 (lm)

–

65.6

58.0

47.0

42.7

37.2

Table 2 Primary dendrite arm spacing (kd), temperature gradient (G), withdrawal rate (V), solidification range (DT) and equilibrium segregation coefficient (k) of different Ni and Ni3Al based SC superalloys. Classification

Alloy

V (lm/s)

G (K/cm)

kd (lm)

DT (K)

k

References

Ni3Al based SC superalloys

IC6SX This work

25–100 50–400

50 170

376–482 151–217

26 31

0.65 0.71

[6,7] [28]

1st generation Ni based SC superalloys

AM3 DD3

50–600 30–200

360 45

51–116 186–291

41 35

0.69 0.74

[14,29] [30,31]

2nd generation Ni based SC superalloys

CMSX-4 Alloy B

19–41 50–500

30–60 250

227–315 61–103

40 51

0.63 0.49

[4,22] [32,33]

3rd generation Ni based SC superalloys

CMSX-10 Alloy D

14–97 50–500

20 250

188–390 46–88

32 45

0.62 0.51

[3,22] [13,33]

number to small size and large number with increasing withdrawal rate. In this study, the morphology of c0 phases changed from sphere to cube with increasing withdrawal rate, which was same with the research on NASAIR 100 [15], but was opposite to the research on Alloy D [13]. The reason of this difference behavior would be further discussed in Section 3.3. As shown in Table 3, the size difference of c0 phases between cellular/intercellular and dendrite/interdendritic region was mainly caused by dendritic microsegregation of Al element during solidification [13,14,28]. Higher Al content led to a higher degree of supersaturation, which promoted c0 phases to precipitate from c phases earlier and then resulted in higher growth rate, longer growth time and consequently larger c0 phases. Fig. 2. Liquid diffusion coefficients (DL), Al and Re contents (in at.%) of different Ni and Ni3Al based single crystal superalloys.

3.3. c/c0 segregation and lattice misfit 3.2. Morphology and size of c0 phases The SEM images and average sizes of c0 phases at different withdrawal rates are shown in Fig. 3 and Table 3, respectively. Large and spherical c0 phases were observed in the SC1 sample (Fig. 3a). Spherical c0 phases were also found in the SC10 sample (Fig. 3b and c), but their size obviously decreased. The morphology of c0 phases in the SC50 sample was still irregular (Fig. 3d and e), but already had a tendency to change from sphere to cube. Cubic c0 phases were found in dendrite core of the SC200 sample, while spherical c0 phases were still observed in interdendritic region (Fig. 3f and g). The size of c0 phases in the SC400 sample was very small and had a good cubic degree both in dendrite core and interdendritic region (Fig. 3h and i). c0 phases precipitate from c phases after solidification finished, and therefore the precipitation and growth of c0 phases obeyed the rule of solid-state phase transformation. The size of c0 phases was determined by following three factors: (1) the cooling rate (G*V) increased with increasing withdrawal rate, which induced higher degree of supercooling and further led to the reduction of critical nucleation energy and the increment of nucleation ratio (Ic0 ); (2) larger supercooling decreased precipitation temperature of c0 phases and led to lower diffusion rate of atoms and lower growth rate of c0 phases (vc0 ); (3) the growth time of c0 phases (tc0 ) decreased with increasing withdrawal rate [34]. These three influencing factors caused c0 phases to change from large size and small

The partitioning ratios of alloying elements in c and c0 phases of SC1 sample (planar microstructure), cellular region (SC10 sample) and dendritic region (SC50, SC100, SC200, SC300 and SC400 samples) were characterized by partitioning ratios kc/c0 = (c phases composition, in wt.%)/(c0 phases composition, in wt.%). kc/c0 > 1 indicated that the element segregated to c phases and kc/c0 < 1 indicated that the element segregated to c0 phases. The influence of withdrawal rate on kc/c0 is shown in Fig. 4a. Re, Cr and Mo preferred to segregate to c phases, while Ni, Al and Ta preferred to segregate to c0 phases. Among them, Re and Ta had the severest segregation, which is consistent with the results of TEM–EDS [35] and atom probe tomography (APT) [36]. Due to c phases were rich in Re, Cr and Mo elements and c0 phases were rich in Al and Ta elements. Therefore during the growth of c0 phases, the concentration gradient promoted the diffusion of alloying elements between c and c0 phases. Re, Cr and Mo elements diffused from c phases to c0 phases, while Al and Ta elements diffused from c0 phases to c phases. As shown in Fig. 4a, the segregation degree between c and c0 phases was the lowest in SC1 sample, which was caused by the low withdrawal rate leading to long growth time and high growth temperature of c0 phases. Thus the diffusion of alloying elements between c and c0 phases was very sufficient and led to the lowest composition differences and segregation degrees between c and c0 phases. As c0 phases were ordered phases, the diffusion coefficients of elements in c0 phases were far less than those of in c phases

C. Ai et al. / Journal of Alloys and Compounds 592 (2014) 164–169

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Fig. 3. SEM images of c0 phases at different withdrawal rates: (a) 1 lm/s, (b) 10 lm/s, cell core, (c) 10 lm/s, intercellular region, (d) 50 lm/s, dendrite core, (e) 50 lm/s, interdendritic region, (f) 200 lm/s, dendrite core, (g) 200 lm/s, interdendritic region, (h) 400 lm/s, dendrite core, and (i) 400 lm/s, interdendritic region.

Table 3 Average sizes of c0 phases at different withdrawal rates (V). V (lm/s) a

b

c

P /C /D (lm) Pa/ICd/IDe (lm) a b c d e

1

10

50

100

200

300

400

4.53

2.16 1.76

1.31 1.67

0.96 1.33

0.53 0.87

0.46 0.79

0.30 0.59

P: planar. C: cell core. D: dendrite core. IC: intercellular region. ID: interdendritic region.

[27,37,38]. Therefore c phases enriched elements (Re, Cr and Mo) and c0 phases enriched elements (Al and Ta) would be discussed separately. (1) With increasing withdrawal rate (1–50 lm/s), the time and temperature of the diffusion of Re, Cr and Mo elements decreased, which led to the increment of the composition differences and segregation degrees between c and c0 phases. While at moderate withdrawal rate (50–200 lm/s), the increasing range of segregation degrees became lower. At high withdrawal rate (200–400 lm/s), the segregation degree of Re was near constant, but the segregation degrees of Cr and Mo still slightly increased, which might be caused by the significantly higher diffusion rates of Cr and Mo than Re in c phases [27].

(2) The general trend of the influence of withdrawal rate on segregation degrees of Al and Ta was similar to Re, Mo and Cr. The segregation degree between c and c0 phases firstly increased (1–100 lm/s) and then maintained almost unchanged (100–400 lm/s), but the transitional withdrawal rate was lower than Re, Cr and Mo elements. The possible reason was the low diffusion rates of alloying elements in c0 phases, which made the composition differences of Al and Ta elements between c and c0 phases reached balanced at a lower withdrawal rate. The lattice constants of c and c0 phases can be determined by XRD [39–41] or calculated by empirical formulas in [42,43]:

ac ¼ 3:524 þ 0:179  C Al þ 0:478  C Mo þ 0:700  C Ta þ 0:110  C Cr þ 0:441  C Re ð½42Þ

ð4Þ

ac0 ¼ 3:570 þ 0:208  C 0Mo þ 0:500  C 0Ta  0:004  C 0Cr þ 0:262  C 0Re ð½42Þ

ð5Þ

ac ¼ 3:522 þ 0:221  C Al þ 0:412  C Mo þ 0:693  C Ta þ 0:122  C Cr þ 0:382  C Re ð½43Þ

ð6Þ

ac0 ¼ 3:569 þ 0:097  C 0Mo þ 0:398  C 0Ta þ 0:014  C 0Cr  0:504  C 0Re ð½43Þ

ð7Þ

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C. Ai et al. / Journal of Alloys and Compounds 592 (2014) 164–169

Fig. 4. (a) Effect of withdrawal rate on c/c0 partitioning ratios (kc/c0 ) of experimental alloys, (b) X-ray diffraction pattern, fitted peak and separated peaks of SC1 sample, (c) effect of withdrawal rate on lattice constants of c and c0 phases (ac and ac0 ), and (d) effect of withdrawal rate on lattice misfit measured by TEM–EDS (dTEM–EDS).

where the units of ac and ac0 are 1010 m, Ci (i=Al,Mo,Ta,Cr,Re) and C0 i 0 (i=Mo,Ta,Cr,Re) are the atomic percentages of elements in c and c phases, respectively. In order to verify whether these formulas were suitable for this study, the SC1 sample, which had a uniform c and c0 phases was characterized by XRD, as shown in Fig. 4b. Because the internal stresses from c0 phases, the lattice constant of c phases deviate to two kinds: ac1 (lattice constant of c phases which were under low stress) and ac2 (lattice constant of c phases which were under high stress) [40]. Therefore, the original XRD curve was separated to three peaks, with increasing 2h, the peak represented c1, c2 and c0 , respectively. The lattice constants of c and c0 phases calculated by the above equations and those measured by XRD are shown in Fig. 4c. By XRD result, the lattice constants showed ac2 > ac1 > ac0 , which was in accord with previous literature [39,40]. Meanwhile, in SC1 sample, ac and ac0 calculated from TEM–EDS results by using the formulas (4) and (5) were very close to ac2 and ac0 measured by XRD, which indicated formulas (4) and (5) were more suitable for our alloy. Moreover, as seen in Fig. 4c, with increasing withdrawal rate, ac gradually increased and the trend of increasing became smaller at high withdrawal rate (200–400 lm/s), while the influence of withdrawal rate on ac0 was negligible. The lattice misfit calculated from TEM–EDS results (dTEM–EDS) can be determined by [41]:

dTEM—EDS ¼ 2ðac0  ac Þ=ðac0 þ ac Þ

ð8Þ

Fig. 4d shows the effect of withdrawal rate on dTEM–EDS. dTEM–EDS gradually decreased with increasing withdrawal rate, but the trend of decrement became slower at high withdrawal rate (200– 400 lm/s). With increasing withdrawal rate, the segregation degree between c and c0 phases increased. Therefore, for a superalloy which

had a negative lattice misfit, the lattice misfit turns more negative. Meanwhile, according to previous analysis (Section 3.2), the size of c0 phases significantly decreased with increasing withdrawal rate. Consequently, the effect of withdrawal rate on the morphological evolution of c0 phases depended on the competitive relation between the following two factors: (1) Misfit factor. For a superalloy which had a negative misfit, the absolute value of lattice misfit increased with decreasing lattice misfit, and therefore the strain energy between c and c0 phases was larger and the tendency of forming cubic c0 phases increased to minimize strain energy [34,42,44,45]. (2) Size factor. The size of c0 phase (r) had a significant influence on interfacial energy (r) and strain energy (DGS), and r / r2, DGS / r3. For smaller c0 phases, the specific surface area (superficial area/volume) was larger. Therefore the morphology of c0 phases was mainly depended on the interfacial energy and tended to form spherical c0 phases [34]. As the cellular/dendritic core is the primary phase during solidification, it is more representative to analyze the morphological evolution of c0 phases in cellular/dendritic core. Table 4 shows the withdrawal rate, size of c0 phases (in planar microstructure, cellular/dendritic core in cellular/dendritic microstructure), as-cast microstructure, lattice misfit and morphological evolution of c0 phases in four superalloys. Compared with this work and NASAIR 100, the sizes of c0 phases in Alloy B and Alloy D were significantly smaller and therefore the influence of size factor was stronger and led c0 phases to change from cube to sphere. In contrast, the misfit factor played a more important role in this work and NASAIR 100, so c0 phases changed from sphere to cube with increasing withdrawal rate.

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C. Ai et al. / Journal of Alloys and Compounds 592 (2014) 164–169 Table 4 Withdrawal rate (V), size of c0 phases (r), as-cast microstructure, lattice misfit and morphological evolution of c0 phases in four SC superalloys. Alloy

a b c

V (lm/s)

r (lm)

As-cast microstructure a

b

Lattice misfit

Morphological evolution

c

References

This work

1–400

4.53–0.30

P –C -D

Negative

Sphere–cube

This work

NASAIR 100

1.5–87

2.05–0.30

Pa–Cb–Dc

Negative

Sphere–cube

[15,46]

Alloy B

3.3–500

0.59–0.04

Pa–Cb-Dc

Negative

Cube–sphere

[32,33]

Alloy D

10–500

0.23–0.03

Cb–Dc

Negative

Cube–sphere

[13,33]

P: planar microstructure. C: cellular microstructure. D: dendritic microstructure.

4. Conclusions (1) The as-cast microstructure changed from planar to cellular and then to dendritic with increasing withdrawal rate in our newly developed Ni3Al based single crystal superalloy. Meanwhile, dendrite arm spacing and c0 phases were significantly refined at higher withdrawal rate. (2) The liquid diffusion coefficients of different Ni and Ni3Al based SC superalloys were estimated based on Kurz–Fisher model and it was found that Al content had strong influence on the liquid diffusion coefficient. (3) The lattice misfit between c and c0 phases in planar, cellular and dendritic region were negative and decreased with increasing withdrawal rate, because of severer segregation between c and c0 phases. The enlarged absolute value of lattice misfit was the reason that the c0 phases changed from sphere to cube at higher withdrawal rate.

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