Microstructural studies on Alloy 693

Journal of Nuclear Materials 453 (2014) 91–97

Contents lists available at ScienceDirect

Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat

Microstructural studies on Alloy 693 R. Halder a, R.S. Dutta a, P. Sengupta a,⇑, I. Samajdar b, G.K. Dey a a b

Materials Science Division, Bhabha Atomic Research Centre, Mumbai 400 085, India Dept. of Metall. Engg. & Mater. Sci., Indian Institute of Technology Bombay, Mumbai 400 072, India

a r t i c l e

i n f o

Article history: Received 1 March 2014 Accepted 2 July 2014 Available online 9 July 2014

a b s t r a c t Superalloy 693, is a newly identified ‘high-temperature corrosion resistant alloy’. Present study focuses on microstructure and mechanical properties of the alloy prepared by double ‘vacuum melting’ route. In general, the alloy contains ordered Ni3Al precipitates distributed within austenitic matrix. M6C primary carbide, M23C6 type secondary carbide and NbC particles are also found to be present. Heat treatment of the alloy at 1373 K for 30 min followed by water quenching (WQ) brings about a microstructure that is free from secondary carbides and Ni3Al type precipitates but contains primary carbides. Tensile property of Alloy 693 materials was measured with as received and solution annealed (1323 K, 60 min, WQ) and (1373 K, 30 min, WQ) conditions. Yield strength, ultimate tensile strength (UTS) and hardness of the alloy are found to drop with annealing. It is noted that in annealed condition, considerable cold working of the alloy can be performed. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction Immobilization of high level nuclear wastes within inert glass matrices is considered to be the best available options today [1– 5]. In general, (a) induction heated metallic melter pot and (b) Joule heated ceramic melter pot are used for this purpose [6,7]. However these technologies suffer from faster degradation of Superalloy 690 (high Cr, low Fe containing Ni based austenitic alloy) components, e.g. thermowells, feeders, electrodes, pour spout assemblies, metallic process pot, etc., during service. Experimental studies under partially simulated service conditions [8–16] suggest that the following physicochemical changes at the Alloy 690/borosilicate melt interfaces might play crucial roles in the material failures; (i) secondary Cr carbide (Cr23C6 type) precipitation along grain boundaries and at triple point junctions, (ii) Cr depletion within austenitic matrix close to interface, (iii) intergranular attack and incorporation of waste glass melt pool components through diffusion and redistribution along these openings, (iv) formation of Cr2O3 and glassy layers (containing needle shaped Ni2CrO4, and cubic NiCr2O4 phases) at the interface.

⇑ Corresponding author. E-mail address: [email protected] (P. Sengupta). http://dx.doi.org/10.1016/j.jnucmat.2014.07.005 0022-3115/Ó 2014 Elsevier B.V. All rights reserved.

Such pre-mature failures not only cause unprecedented shut down of vitrification plant operations but also generate huge pile-up of hazardous non-compactable radioactive metallic wastes. To avoid such situations, ‘preventive management strategy’, either through development of diffusion barrier coating on Alloy 690 or through its compositional modification appear to be more appropriate ones. Towards the first option, the authors have already successfully developed (i) Ni-YSZ composite [17] and (ii) Ni-aluminide [18] coatings on the laboratory scale specimens. However, considering the harsh environment experienced by the Superalloy within vitrification furnaces, it is thought worthwhile to explore the alloy modification route as well [19,20]. Superalloy 693, is a newly developed ‘high-temperature corrosion resistant alloy’ which contains more than about 2.5 wt% Al than its predecessor, i.e. Alloy 690 (Table 1). The present study describes the microstructure and mechanical properties of Superalloy 693 prepared through a combination of ‘vacuum induction melting’ and ‘vacuum arc melting’ routes. Gallium (Ga) isotopes are commonly produced within spent nuclear fuels during neutron irradiation and subsequent decay of radio-nuclides such as U235, U238, Th232, Pu239, Am241 and Np237 [21]. Since Ga is known to cause stress associated failures in aluminium alloys [21] and as Alloy 693 contains 2.5 wt% of Al it was felt prudent to examine the possible effect of low energy Ga ion irradiation on the alloy. Further, as during immobilization of high level nuclear wastes, the vitrification furnace components are being exposed to ‘ionic damage mechanisms’ so Ga ion irradiation is also used for simulating the expected service conditions.

92

R. Halder et al. / Journal of Nuclear Materials 453 (2014) 91–97

Table 1 Bulk composition (in wt%) of Superalloys 690 and 693. Element (wt%)

Cr

Fe

Al

Cu

Si

Mn

S

C

Nb

Ti

N

Ni

Alloy 690 Alloy 693 (minimum) (maximum) Used Alloy (XRF analyses)

27.0–31.0 27.0 31.0 29.32

7.0–11.0 2.5 6.0 3.96

0.50 max 2.5 4.0 3.19

0.50 max – 0.5 <0.02

0.5 max – 0.5 0.04

0.5 max – 1.0 0.09

0.01 max – 0.01 <0.002

0.05 max – 0.15 0.097

– 0.5 2.5 1.86

– 1.0 0.42

– – 130 ppm

Rest Rest Rest Rest

cps PET

30000

Matrix Grey: M1 White: M2

CrKα 20000 10000

FeKα CrKβ 0

TiKα 30000

CaKα

NbLα 70000

50000

cps

ppt 1

ppt 2

Sin LIF

CrKα

5000 3000

ppt 3

NiKβ NiKα

1000 0

30000

FeKα

CrKβ

50000

70000

Sin

Fig. 2. X-ray spectrum showing the presence of various elements within precipitate 1.

Fig. 1. Back scattered electron image (with atomic number contrast) showing the three types of precipitates, designated as precipitate 1 (ppt 1), precipitate 2 (ppt 2) and precipitate 3 (ppt 3) within the alloy. Also note that the matrix is constituted of two phases namely grey phase (M1) and white phase (M2).

2. Experimental studies Microchemical analysis of as-received Alloy 693 was carried out using CAMECA SX-100 Electron Probe Micro-Analyser (EPMA). Small coupons were cut from Alloy 693 sheet, ground with emery paper and polished with diamond paste of 1 lm grain size. An acceleration voltage of 20 keV and stabilized beam current of 4 nA and 20 nA were used for electron imaging (secondary electron; SE and back scattered; BSE) and qualitative analysis respectively. Samples for examination under transmission electron microscopy (TEM) were prepared by jet electropolishing technique using an electrolyte containing 90% methanol and 10% perchloric acid by volume. A temperature of about 238 K was maintained during electropolishing with the applied voltage of around 20 V. The electropolishing of each specimen was continued until a tiny hole was formed at the centre of 3 mm disc. To understand the nature of grain-orientation and its size distributions, electron backscatter diffraction patterns (EBSD) of the samples were obtained using FEI Quanta 3D scanning electron microscope system. The scans were taken over an area of 250  400 lm2, using an accelerating voltage at 30 kV and probe current 16 nA. Kikuchi bands were obtained for each of the different phases present within the alloy. In order to obtain some preliminary idea on the response of the alloy to ionic damage mechanisms, small coupons were exposed to 30 keV Gallium (Ga) ions with ion current 0.1 nA for 10 min at ambient temperature under high vacuum, using a focussed ion beam (FIB) gun attached to the system. To study the ‘‘Ga ion damage on microstructure’’ detailed grain boundary mapping of the austenitic matrix (70  125 lm2) was carried out. The net fluence and dose were found out to be 4.29  1015 ions/cm2 and 2.878 KGy

respectively; assuming that all ions fell on the surface where fluence was being calculated. As the experiment was carried out at high vacuum, attenuation through scattering in the medium was neglected. Tensile property measurements were done in as received and solution annealed (1323 K, 60 min, WQ) and (1373 K, 30 min, WQ) conditions each with 20 mm gauge length. Experiments were carried out at ambient temperature using test velocity of 0.25 mm/ min (i.e. 104/s). Hardness of as received and solution annealed alloy coupons was measured using a load of 200 gf and 10 s as dwell time. Micro-hardness values quoted here were obtained over an average of 10 readings. cps

TAP

9000 NbLα 6000 3000 00

30000

50000

70000

Sin

cps PET NbLα

2500

TiKα 1500

NbLβ

CrKα CrKβTiKβ

500 0

30000

50000

70000

Sin

cps LIF

300

TiKα 200

NiKα CrKα

100

TiKβ 0

30000

50000

70000

Sin

Fig. 3. X-ray spectrum showing the presence of various elements within precipitate 2.

93

R. Halder et al. / Journal of Nuclear Materials 453 (2014) 91–97

Matrix Element

Weight %

-------

--------

Al(K)

3.31

Cr(K)

33.33

Fe(K)

5.13

Ni(K)

56.65

Nb(K)

1.56

Fig. 4. EDS spectra showing the average composition of as-received Alloy 693 austenite matrix.

3. Results and discussion 3.1. Microstructural studies Back-scattered electron micrograph of as-received Alloy 693 is given in Fig. 1. It shows that the alloy contains three different types of precipitates, which are, (a) spherical/elliptical shaped grey precipitates (precipitate 1), (b) irregular white precipitates (precipitate 2; partially to totally replacing precipitate 1) and (c) fine grained intergranular precipitates (precipitate 3). The matrix is also found to be a mixture of two phases, grey (M1) and white

(M2) phases. Spot X-ray analyses of precipitates 1 and 2, using wave length dispersive spectrometers (WDS), identify them as Cr23C6 (Fig. 2) and NbC (Fig. 3) precipitates respectively. Similar WDS analyses of precipitate 3 and M1–M2 phases were difficult due to their small sizes. Matrix composition was obtained using energy dispersive spectrometer (EDS) and the result (Fig. 4) matches very closely with the bulk compositional analysis of the alloy obtained from X-ray fluorescence (XRF) analysis (Table 1). For better comprehension of the alloy matrix, samples were studied under TEM equipped with energy dispersive spectrometer (EDS), and it was noted that the white M2 phase is essentially ordered Ni3Al (note the selected area diffraction (SAD) pattern in Fig. 5), distributed within austenite matrix (M1). Further, the

020A 010ppt

200

[011]

[001]

Austenite matrix M 6C

Ni3 Al

Ni3Al

200 nm

Fig. 5. Dark-field image of the matrix showing the uniform distribution of fine ordered Ni3Al type precipitates within austenitic matrix of as-received Alloy 693 sample. Inset shows SAD pattern of Ni3Al type (faint spots) phase along with the austenite matrix (bright spots).

Fig. 6. Bright-field image showing the presence of coarse particles within the asreceived Alloy 693 austenite matrix. Inset shows SAD pattern corresponding to M6C type phase. Also note the uniform distribution of fine grained ordered Ni3Al type precipitates within austenitic matrix.

94

R. Halder et al. / Journal of Nuclear Materials 453 (2014) 91–97

(a)

(b)

Absence of Ni3Al

Coarse particles

Dislocations

020

[001]

200 nm

1 µm

Fig. 7. (a) Bright-field image of solution-annealed (1373 K/30 min, WQ) Alloy 693 showing planar arrangement of dislocations and absence of fine, ordered, uniformly distributed Ni3Al type precipitates. Inset shows SAD pattern of austenite matrix (fcc). (b) Bright-field TEM photomicrograph of coarse particles randomly distributed in the austenite matrix of Alloy 693 in welded + annealed condition.

coarse intergranular precipitates have been identified as M6C type (see the SAD pattern in Fig. 6). Studies on alloy 690 on the other hand shows presence of TiN and Cr23C6 type carbides. The later one has a cube–cube orientation relationship such as {1 0 0}||{1 0 0}Cr23C6 and h1 0 0i||h1 0 0iCr23C6. These precipitates are mostly formed along the grain boundaries – coarse and discrete

(a)

Cr23C6

on high angle boundaries and fine – faceted ones on low angle boundaries [22]. Separate experiments were carried out to identify microstructure of solution annealed samples. Based on previous experiences, two different heat treatments, i.e. (a) 1323 K/2 h and (b) 1373 K/ 30 min, followed by water quenching (WQ) were selected.

(b)

NbC

Boundaries: Rotation angle Min Max Fraction 2o 5o 0.013 o 5 15o 0.019 15o 180o 0.989

Ni

(c)

Cr23C6

(d) RD

NbC

Boundaries: Rotation angle Min Max Fraction 2o 5o 0.013 o 5 15o 0.014 15o 180o 0.973

Nickel

Fig. 8. EBSD map, grain boundary map for (a) and (b) as received alloy 693 and (c) and (d) those annealed at 1373 K for 30 min.

R. Halder et al. / Journal of Nuclear Materials 453 (2014) 91–97

95

Fig. 9. (a) and (b) CI filtered (CI < 0.1) EBSD map, (c) and (d) inverse pole figures, (e) mis-orientation angle and (f) grain boundary density for annealed Alloy 693 before (the images on left side column wise) and after Ga ion damage (all the images on right side column wise).

A cleaner microstructure was obtained for the sample heat treated at 1373 K for 30 min. Fig. 7 shows representative bright field images of that sample, revealing absence of ordered Ni3Al precipitates within austenite matrix and presence of intergranular M6C

precipitates; often closely associated with planar arrangement of dislocations. Most of the studies on Alloy 690 are mainly focused on the formation of Cr- rich M23C6 type carbides and depletion of Cr on thermal heat treatment [22–24]. Dutta et al. [25] reported

96

R. Halder et al. / Journal of Nuclear Materials 453 (2014) 91–97 Table 3 Hardness values of Alloy 693 samples (WQ: water quenched).

As-Received 1183 MPa

16000 14000

33%

Load (kN)

12000

1323K/60m/WQ

820 MPa

825 MPa

10000

Alloy condition

Hardness value, VHN

As-received

(a) 562 (b) 503 (c) 562.6

1323 K/2 h, WQ

(a) 278.3 (b) 270.1 (c) 270.0

1373 K/30 min, WQ

(a) 248.7 (b) 261.3 (c) 260.3

1453 K/1 h, WQ (as received in annealed condition)

(a) 281.6 (b) 285.0 (c) 281.6

1323 K/2 h, WQ + 70% cold worked

(a) 486.4 (b) 453.6 (c) 487.4 (d) 455.8

1323 K/2 h, WQ + 65% cold worked

(a) 444.5 (b) 428.2 (c) 438.3 (d) 431.7

Deformed

(a) 452 (b) 455

1373K/30m/WQ

808 MPa 369 MPa

8000

52%

6000

60%

Test Temperature 298K -4 340 MPa Test Velocity 0.25 mm/min (~ 10 /sec) Round c/s diameter 4.0 mm Gage Length 20 mm Extensometer 12.5 mm

4000 2000 0 0

2

4

6

8

10

12

Extension (mm) Fig. 10. Graphical representation showing the strength variation of Alloy 693 with solution annealing.

that in case of Alloy 690, solution annealing followed by water quenching results in dissolution of secondary Cr-carbide precipitates. However, small spherical TiC particles (fcc) and larger, faceted TiN particles (hexagonal) remained randomly distributed within the austenitic matrix. Representative EBSD and grain-boundary maps for as received and solution annealed samples are shown in Fig. 8. The scanned area comprises of both large and small grains, and the color coding shows that all possible orientations are present within both the samples with maximum grain boundaries exhibiting rotation angles greater than 15°. It must be mentioned here that, in the present context the term ‘grain’ refers to collection of neighboring pixels having mis-orientation less than 2°. No texture development was observed within the alloy after heat treatment though several twin boundaries were observed. The annealing twins form mainly because of low stacking fault energy and such twin boundaries have been observed in case of alloy 690 as well [21]. The damage structure produced by Ga ion mainly comprises of dislocations, point defects and precipitates [26]. From CI filtered EBSD maps (max CI < 0.1) presented in Fig. 9(a) and (b), it is noted that there is an increase in the mis-indexed regions after ion irradiation, mainly within the grains. Moreover, formation of new orientations close to the original ones are also observed from the inverse pole figures (Fig. 9(c) and (d)). The increase in low angle mis-orientation and increase in grain boundary density Fig. 9((e) and (f)) is possibly a result of lattice strain produced by ion irradiation as well as formation of dislocation networks, point defects, etc. However, after irradiation, the high angle mis-orientation decreased. This could be due to the formation of mis-indexed regions preferably from high angle misorientation. The observations made in present case are typical of a Ga+ ion damaged structure and is also observed in case of fcc Cu [26]. 3.2. Mechanical properties Preliminary mechanical property studies have been carried out with as-received and annealed Alloy 693. Results are shown in Fig. 10, which indicate that with annealing of the samples, yield strength and ultimate tensile strength (UTS) drop drastically. The

Table 2 Results of stress rupture tests on Alloy 693 coupons. Temperature (K)

Stress (MPa)

Life (h)

1223 1223 1223

10 20 30

750 687 204

stress rupture test results are given in Table 2, which suggest that for 1223 K applications Alloy 693 has restricted service life. Hardness values of as-received Alloy 693, solution annealed, solution annealed + cold worked, and only deformed samples are given in Table 3. Cold working of 1373 K/30 min and 1323 K/2 h water quenched (WQ) samples could be done up to 65% and 70% respectively without any noticeable edge cracking. The high hardness value of Alloy 693 was attributed to high volume fraction of Ni3Al strengthening fine precipitates distributed within the austenitic matrix. It is evident from the data that hardness values have decreased after annealing treatments and the effect is more pronounced as the annealing temperature is increased. After the solutionizing treatments, the Ni3Al and Cr23C6 precipitates dissolve making the matrix softer. Hence, the dislocations can easily move and deformation can occur easily. This is further supported by the stress–strain curves, which show decrease in yield strength and UTS after treatment. Higher hardness values of 1323 K/2 h, WQ + 70% cold worked sample as compared to 1373 K/30 min, WQ + 65% cold worked sample could be attributed to its higher initial hardness values. The secondary Cr-carbides dissolve on annealing and make it easier for the dislocation to move and cause deformation. The improved mechanical properties of Alloy 693 as compared to Alloy 690 could be due to difference in composition and volume fraction of undissolved precipitates. Small amount of Ni3Al that remained undissolved in Alloy 693 after solutionizing can also contribute to increased hardness and better tensile properties. 4. Summary Superalloy 693, is a newly developed ‘high-temperature corrosion resistant alloy’ which contains more than about 2.5 wt% Al than its predecessor, i.e. Alloy 690. The experimental results obtained from Superalloy 693 prepared through a combination of ‘vacuum induction melting’ and ‘vacuum arc melting’ routes are summarized below, 1. Detailed microstructural characterizations of as-received Alloy 693 revealed the presence of Ni3Al precipitates within austenitic matrix. Additionally, (a) coarser grained,

R. Halder et al. / Journal of Nuclear Materials 453 (2014) 91–97

2.

3.

4.

5.

spherical/elliptical shaped randomly distributed M6C type precipitates, (b) finer grained intergranular M23C6 type precipitates and (c) coarser grained, spherical/elliptical shaped randomly distributed NbC precipitates were found within the alloy. Microstructural characterizations of solution-annealed (1373 K/30 min, water quenched) samples showed much cleaner microstructure, with almost no intergranular and Ni3Al type precipitates. However, presence of M6C type precipitates within the austenitic matrix was still observed. Grain-boundary mapping of as received and solution annealed samples show most of the grains exhibit rotation angles greater than 15°. Upon Ga ion irradiation, enhancements in grain boundary density, and low angle misorientation was noted. Formation of new orientations close to original were also seen. Tensile properties of Alloy 693 materials were measured in as received and solution annealed conditions (1323 K, 60 min, WQ and 1373 K, 30 min, WQ). It is observed that with annealing of the samples the yield strength and ultimate tensile strength (UTS) have dropped drastically. High hardness of Alloy 693 in as received condition is attributable to fine, uniform distribution of ordered Ni3Al type precipitation within austenite matrix. Annealing treatments bring down the hardness values by almost completely dissolving Ni3Al type precipitates. The effect of hardness reduction is more pronounced as the annealing temperature is increased. It is noted that in annealed condition, considerable cold working of the alloy can be performed.

Acknowledgements Authors are grateful to Prof. Stephane Gin (Editor) and anonymous reviewer for their suggestions. Technical discussion with Dr. R. Kishore on mechanical properties assessments is acknowledged. The work was funded by Department of Atomic Energy, Government of India.

97

References [1] M.I. Ojovan, W.E. Lee, New Developments in Glassy Nuclear Wasteforms, Nova Science Publishers, 2007. [2] C.P. Kaushik, R.K. Mishra, P. Sengupta, D. Das, G.B. Kale, K. Raj, J. Nucl. Mater. 358 (2006) 129. [3] R.K. Mishra, P. Sengupta, C.P. Kaushik, A.K. Tyagi, G.B. Kale, K. Raj, J. Nucl. Mater. 360 (2007) 143. [4] R.K. Mishra, V. Sudarsan, P. Sengupta, R.K. Vatsa, A.K. Tyagi, C.P. Kaushik, D. Das, K. Raj, J. Am. Ceram. Soc. 91 (2008) 3903. [5] P. Sengupta, S. Fanara, S. Chakraborty, J. Hazard. Mater. 190 (2011) 229. [6] P. Sengupta, C.P. Kaushik, G.K. Dey, Immobilization of high level nuclear wastes: the Indian scenario, in: M. Ramkumar (Ed.), On a Sustainable Future of the Earth’s Natural Resources, Springer-Verlag, 2013. [7] K. Raj, K.K. Prasad, N.K. Bansal, Nucl. Eng. Des. 236 (2006) 914. [8] I.W. Donald, Waste Immobilization in Glass and Ceramic Based Hosts, Wiely, 2010. [9] P. Sengupta, J. Nucl. Mater. 411 (2011) 181. [10] V. Kain, P. Sengupta, P.K. De, S. Banerjee, Metal. Mater. Trans. 36A (2005) 1075. [11] P. Sengupta, J. Mittra, G.B. Kale, J. Nucl. Mater. 350 (2006) 66. [12] P. Sengupta, C.P. Kaushik, R.K. Mishra, G.B. Kale, J. Am. Ceram. Soc. 90 (2007) 3085. [13] P. Sengupta, N. Soudamini, C.P. Kaushik, R.K. Mishra, D. Das, G.B. Kale, Jagannath, K. Raj, B.P. Sharma, J. Nucl. Mater. 374 (2008) 185. [14] P. Sengupta, C.P. Kaushik, G.B. Kale, D. Das, K. Raj, B.P. Sharma, J. Nucl. Mater. 392 (2009) 379. [15] K.J. Imrich, D.F. Bickford, G.G. Wicks, R.C. Hopkins, Ceram. Trans. 107 (2000) 384. [16] D.F. Bickford, A. Applewhiteramsey, C.M. Jantzen, J. Am. Ceram. Soc. 73 (1990) 2896. [17] P. Sengupta, D. Rogalla, H.W. Becker, G.K. Dey, S. Chakraborty, J. Haz. Mater. 192 (2011) 208. [18] R.S. Dutta, C. Yusufali, B. Paul, S. Majumdar, P. Sengupta, R.K. Mishra, C.P. Kaushik, R.J. Kshirsagar, U.D. Kulkarni, G.K. Dey, J. Nucl. Mater. 432 (2013) 72. [19] J.H. Hsu, J.W. Newkirk, C.W. Kim, C.S. Ray, R.K. Brow, M.E. Schlesinger, D.E. Day, Corros. Sci. 75 (2013) 148. [20] D. Zhu, C.W. Kim, D.E. Day, J. Nucl. Mater. 336 (2005) 47. [21] D.F. Wilson, E.C. Beahm, T.M. Besmann, J.H. DeVan, J.R. DiStefano, U. Gat, S.R. Greene, P.L. Rittenhouse, B.A. Worley, ORNL/TM-13504, 1997. [22] Y.S. Lim, J.S. Kim, H.P. Kim, H.D. Cho, J. Nucl. Mater. 335 (2004) 108. [23] J.J. Kai, M.N. Liu, Scripta Metall. 23 (1989) 17. [24] S. Xia, B. Zhou, W. Chen, Met. Trans. A 21A (1990) 2097. [25] R.S. Dutta, A. Labo, R. Purandare, S.K. Kulkarni, G.K. Dey, Met. Trans. A 33A (2002) 1437. [26] D. Kiener, C. Motz, M. Rester, M. Jenko, G. dehm, Mater. Sci. Eng. A 459 (2007) 262–272.