Grain boundary liquation of zirconium-containing nickel aluminides

Scripta METALLURGICA et MATERIALIA

GRAIN

Vol. 27, pp. 121-126, 1992 Printed in the U.S.A.

BOUNDARY

LIQUATION

OF ZIRCONIUM-CONTAINING

Pergamon Press Ltd. All rights reserved

NICKEL

ALUMINIDES

G.H.Chen and C.Chen Institute

of Materials Science and Engineering National Taiwan University Taipei, Taiwan 10764, Republic of China (Received December 16, 1991) (Revised May 7, 1992) Introduction

Nickel aluminide, Ni3AI , is an intermetallic compound with the L12 crystal structure. The yield strength of the compound increases with increasing temperature [1,2], which together with good oxidation resistance [3], makes it suitable for structural applications at elevated temperatures. In the early developing stage of the compound, the extremely brittle nature of polycrystalline Ni3AI presented a major problem. However, the addition of boron in trace amounts significantly improves the ductility of Ni3AI [4,5], and the fracture mode changes from intergranular brittle to transgranular ductile fracture. Boron doped in Ni3AI exhibits a strong segregation behavior to grain boundaries, and the degree of segregation is strongly temperature dependent [6,7]. The m e c h a n i c a l s t r e n g t h of Ni3AI can be improved t h r o u g h ternary or quaternary alloying additions [8,9]. But, some alloying elements might cause grain boundary liquation at a temperature lower than 1395°C, the peritectic temperature of the compound [i0]. Grain boundary liquation could cause heataffected zone c r a c k i n g in the weld, hence lowering the w e l d a b i l i t y of the compound. The element additions, therefore, should be carefully controlled to prevent grain boundaries from melting at a temperature which is significantly lower than normal. The present investigation examined the influence of Zr additions on the grain boundary liquation of boron doped nickel aluminides.

Materials

and Experimental

procedures

The composition of the ternary compounds was Ni77A123_xZr x (X=0.5 or i), and all alloys were doped with 500 ppm (wt%) boron. Commercially pure nickel, aluminum t o g e t h e r w i t h Ni-B and Ni-Zr master alloys were v a c u u m - i n d u c t i o n melted to obtain ingots with the various compositions. Contents of oxygen and sulfur in the ingots were controlled at levels below 50 and i0 ppm, respectively. The ingots were sliced and cold rolled with intermediate annealing at I000-I050°C, i n t o strips of approximately 1.2 mm thick. After thermomechanical treatment (TMT), the microstructure of the samples was r' (LI 2 structure) single phase with equiaxed grains of about 20 ~m. Specimens used in the heat treatment e x p e r i m e n t s had d i m e n s i o n s of 15x15xl.2 mm. The heat t r e a t m e n t temperature was in the range of 1200°C to 1340°C, with temperature controlled within ±3°C. Samples were kept in furnace for 15 minutes at a predetermined temperature and then water quenched (W.Q.). The surface layer of specimens 121 0956-716X/92 $5.00 + .00 Copyright (c) 1992 Pergamon Press Ltd.

122

GRAIN BOUNDARY

was r e m o v e d at least 200 #m to assure problem in the e v a l u a t i o n process.

LIQUATION

that the o x i d a t i o n

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did

27, No.

not p r e s e n t

a

All h e a t - t r e a t e d specimens were subjected to m e t a l l o g r a p h i c observations. Fractographic studies were performed on impact-fractured specimens (U-notched) at room t e m p e r a t u r e and examined by a Philips 515 scanning electron microscope (SEM). S t r u c t u r a l i d e n t i f i c a t i o n was p e r f o r m e d on a P h i l i p PW 1710 X - r a y d i f f r a c t o m e t e r using CuK~ radiation. Compositional analyses were carried out by a J E O L J X A - 8 6 0 0 S X e l e c t r o n probe m i c r o a n a l y z e r (EPMA) on a l i g h t - e t c h e d m e t a l l o g r a p h i c specimen.

Results and Discussion Figure l(a) shows a metallograph of the Ni77AI22Zr I alloy after a 1200W (1200°C/15 m i n u t e s + W.Q.) treatment, where the structure is single r' phase with grain size of a p p r o x i m a t e l y 60 ~m. In the case of the alloy receiving a 1320W (1320°C/15 minutes + W.Q.) treatment, the grain boundaries displayed a zigzag a p p e a r a n c e (Fig. l(b)) due to grain boundary liquation. In fact, grain b o u n d a r y l i q u a t i o n of this alloy became n o t i c e a b l e in the 1300W (1300°C/15 minutes + W.Q.) specimen. Figure 2(a) is a fractograph of 1200W specimens r e v e a l i n g an almost entirely i n t e r g r a n u l a r fracture, while the f r a c t u r e s u r f a c e of the specimens just after TMT showed a completely transgranular fracture mode. For the case of a 1300W (1300°C/15 minutes + W.Q.) treatment, there was apparent evidence of grain b o u n d a r y liquati0n (marked As) on the fracture surface in addition to the i n t e r g r a n u l a r fracture as shown in Fig.2(b). For 1320W specimens, most of grain b o u n d a r i e s had m e l t e d except for some smooth areas (marked Bs in Fig.2 (c)) r e m a i n i n g on the impact-fractured surface. Compositional analysis of the liquated phase (marked As in Fig.2(b)) at the grain boundary by SEM/EDX suggested that it was a Ni-Zr compound as illustrated in Fig.3. EDX analyses of the 1 3 2 0 W s p e c i m e n r e v e a l e d that the m e l t e d r e g i o n s had s i m i l a r c h e m i c a l compositions as d i s p l a y e d in Fig.3, while the composition of the smooth intergranular surface (marked Bs in Fig.2(c)) was the same as the bulk material. Figure 4(a) is the impact-fractured surface of the N i 7 7 A I 2 2 . 5 Z r 0 . 5 specimen s u b j e c t e d to a 1 3 2 0 W t r e a t m e n t , w h i c h shows c o m p l e t e l y i n t e r g r a n u l a r fracture w i t h o u t grain boundary liquation. The fractograph of the same alloy after a 1320F (1320°C/15 minutes + furnace cooled) treatment, however, reveals a t r a n s g r a n u l a r fracture mode as shown in Fig.4(b). Similar o b s e r v a t i o n s have been f o u n d on the N i 7 7 A I 2 2 Z r l s p e c i m e n s s u b j e c t e d to a l o w e r t e m p e r a t u r e (<1300°C) treatment. Choudhury et al [6] reported that slow cooling promotes grain b o u n d a r y s e g r e g a t i o n of boron in Ni3AI, w h i l e rapid q u e n c h i n g retains the less boron level at high temperature. Therefore, the more ductile fracture of the f u r n a c e - c o o l e d specimen could be attributed to the s e g r e g a t i o n of boron back to the grain boundaries during a slow cooling process. F i g u r e s 5(a) t h r o u g h (e) are the results of EPMA on a cross s e c t i o n of the N i 7 7 A I 2 2 Z r I alloy after a 1320W treatment. Figures 5(b) through (e) indicate that the o x i d e layer (40-50 ~m thick) c o n s i s t e d of m a i n l y a l u m i n u m and z i r c o n i u m o x i d e s . The m e l t e d g r a i n b o u n d a r i e s as i n d i c a t e d by a r r o w s in Fig.5(a) were e n r i c h e d in Zr (Fig°5(e)) but depleted in A1 (Fig.5(c)). However, t h e r e w e r e no s i g n i f i c a n t d i f f e r e n c e in Ni and O c o n t e n t s b e t w e e n the melted region of the grain boundary and the grain interior. It is important to

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GRAIN BOUNDARY LIQUATION

[23

note that the surface layer of the specimens was removed at least 200 ~m prior to the t e s t i n g . As a r e s u l t , e m b r i t t l e m e n t of the c o m p o u n d c a u s e d by the p e n e t r a t i o n of o x y g e n a l o n g g r a i n b o u n d a r i e s [II] was e l i m i n a t e d in this study. F i g u r e 6 s h o w s the X - r a y d i f f r a c t i o n peaks of the N i 7 7 A I 2 2 Z r I alloy subjected to a 1320W treatment. In addition to the strong r' peaks, there are several w e a k p e a k s ( i n d i c a t e d by the arrows) w h i c h h a v e b e e n i d e n t i f i e d as Ni5Zr. Ni5Zr is an A u B e 5 - t y p e compound having the F C C s t r u c t u r e and a lattice parameter of 0.67 nm [12,13]. This is consistent with the p r e v i o u s results of the SEM/EDX and EPMA analyses. Since only a small amount of Ni5Zr was observed at the grain boundaries, the relative intensities in the d i f f r a c t i o n spectrum are also weak. The X-ray analyses of the Ni77AI22.5Zr0. 5 alloy (1320W treatment) which did not reveal grain boundary liquation, showed only diffraction peaks of Ni3AI. Furthermore, a Du Pont 2000 d i f f e r e n t i a l t h e r m o a n a l y z e r (DTA) was a l s o u s e d to s t u d y the p o s s i b l e p h a s e c h a n g e of the N i 7 7 A I 2 2 Z r I alloy after a 1 3 2 0 W t r e a t m e n t . It was found that t h e r e was a small e n d o t h e r m i c reaction at about i180°C in the heating curve. A c c o r d i n g to the Ni-Zr binary p h a s e d i a g r a m [14], the e u t e c t i c t e m p e r a t u r e of N i 5 Z r is i172°C, w h i c h is close to the DTA result of i180°C. The results of DTA and X-ray diffraction are in good agreement, the grain boundary liquation a s s o c i a t e d with the formation of Ni5z ~ can be p o s i t i v e l y identified. In contrast, no e n d o t h e r m i c reaction at 1180 C in the DTA e x p e r i m e n t was o b s e r v e d for the N i 7 7 A I 2 2 Z r I alloy just after TMT or after heat treatments which had no grain b o u n d a r y liquation. On the other hand, x-ray and DTA results indicated that the weak peaks and the endothermic reaction at i 1 8 0 ° C w e r e n o t f o u n d f o r t h e 1 3 2 0 F s p e c i m e n (Ni77AI22Zrl) . The t r a n s f o r m a t i o n behavior of Hi5Zr a p p a r e n t l y d e p e n d e d on the c o o l i n g r a t e in t h e e x p e r i m e n t . M o r e d e t a i l a n a l y s i s of t h i s low m e l t i n g product is in progress. The f o r g o i n g r e s u l t s i n d i c a t e d t h a t the N i 7 7 A I 2 2 Z r I a l l o y d i s p l a y e d significant g r a i n b o u n d a r y liquation at a temperature of a p p r o x i m a t e l y i00 C lower than the m e l t i n g point of Ni3AI (1395°C). A l t h o u g h a severe grain boundary l i q u a t i o n o c c u r r e d in the N i 7 7 A I 2 2 Z r I alloy, no m e l t i n g p h e n o m e n o n has been found in the Ni77AI22.5Zr0. 5 alloy at temperature of up to 1320°C. It was suggested that the H i - r i c h Ni3AI compound might have r phase (Hi solid solution, d i s o r d e r e d FCC structure) at the grain boundary [15,16]. Since the solubility of Zr in r is lower than that in r', the rejected Zr atoms combine with enriched Ni atoms at grain boundaries [17-19], thereby p r o m o t i n g grain boundary l i q u a t i o n and f o r m a t i o n of N i 5 Z r c o m p o u n d s . A l o w e r Zr c o n t e n t in the compound, of course, w o u l d reduce the tendency of grain b o u n d a r y liquation as d e m o n s t r a t e d in the Ni77AI22.5Zr0. 5 alloy.

Conclusion In the d e v e l o p m e n t of order intermetallic alloys for s t r u c t u r a l applications, m a c r o a l l o y i n g is an i m p o r t a n t m e t h o d to s u b s t a n t i a l l y m o d i f y the p r o p e r t i e s of the compound. Zr addition has been found to improve the strength of nickel a l u m i n i d e s at high temperatures [8]. However, grain b o u n d a r y liquation a s s o c i a t e d w i t h the f o r m a t i o n of a low m e l t i n g e u t e c t i c phase, N i 5 Z r , o c c u r r e d in the N i 7 7 A I 2 2 Z r I a l l o y at a p p r o x i m a t e l y 1 3 0 0 ° C for s h o r t h o l d i n g times f o l l o w e d by r a p i d l y quenching. The existence of Ni5Zr in the N i 7 7 A I 2 2 Z r 1 a l l o y w a s d e p e n d e n t on t h e c o o l i n g r a t e f r o m e l e v a t e d t e m p e r a t u r e s . G r a i n b o u n d a r y l i q u a t i o n w a s not o b s e r v e d in the N i 7 7 A I 2 2 . 5 Z r 0 . 5 a l l o y e v e n at

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GRAIN BOUNDARY

LIQUATION

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O

1320 C. Thus, lowing the Zr content in the ternary compounds could reduce the tendency of grain boundary liquation. This is particularly important to the alloy design in the weld structure.

Acknowledqments The authors gratefully acknowledge the support Science Council (Contract No. NSC 79-0405-E002-25).

of the

R.O.C.

National

References J.H.Westbrook, Trans. AIME 209, 898(1957). P.A.Flinn, Trans. AIME 218, 145(1960). S.Tanigushi and T.Shibata, Oxid. Met. 25, 201(1986). K.Aoki and O.Izumi, Nippon Kinzoku Gakkashi 43, 1190 (1979). C.T.Liu, C.L.White and J.A.Horton, Acta Metall. 33, 213 (1985). A.Choudhury, C.L.White and C.R.Brook, Scr. Metall. 20, 1061(1986). T.L.Lin, D.Chen and H.Lin, Acta Metall. 39, 523(1991). S.E.Hsu, N.N.Hsu, C.H.Tong and C.Y.Ma, Proc. Annual Conf. Chin. Soc. Mater. Sci., Taiwan R.O.C., p.424(1986). 9. C.T.Liu and C.L.White, High-temperature Ordered Intermetallic Alloys (edited by C.C.Koch, C.T.Liu and N.S.Stoloff), Mat. Res. Soc. Symp., Voi.39, p.365

(1985). i0. Metals Handbook Vol.8, 8th ed., American Society for Metals, Metals Park, OH, p.262 (1973). ii. M.Takeyama and C.T.Liu, Scr. Metall. 37, 2681(1989). 12. P.Nash and Y.Y.Pan, J. Phase Equilibria 12, 105(1991). 13. Powder Diffraction File, JCPDS International Center for Diffraction Data, Pennsylvania, U.S.A., File No. 10-229(1983). 14. Metals Handbook Vol.8, 8th ed., American Society for Metals, Metals Park, OH, p.327 (1973). 15. I.Baker and E.M.Schulson, Scr. Metall. 23, 1883(1989). 16. H.Kung, D.R.Rasmussen and S.L.Sass, Scr. Metall. Mater. 25, 1277(1991). 17. D.N.Sieloff, S.S.Brenner and H.Ming-Jian, High-Temperature Ordered Intermetallic Alloys III (edited by C.T.Liu, A.I.Taub, N.S.Stoloff and C.C.Koch), Mat. Res. Soc. Symp., Voi.133, p.155(1989). 18. J.E.Krzanowski, Scr. Metall. 23, 1219(1989). 19. I.Baker, E.M.Schulson and J.R.Michael, Phil. Mag. B 57, 379(1988).

FIG.1

Optical micrographs of Ni77AI22Zrl and (b) 1320W treatments.

specimens

after

(a) 1200W

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27, No. 1

GRAIN B O U N D A R Y L I Q U A T I O N

125

Ni

0

> ° ~

o

zr

-

!

O0

02

-e

04

i

-

06

- -

OB

10

Energy (KeY) Fig.3

EDX

spectrum

for

the

p h a s e ( m a r k e d As) a t boundary in Fig.3(b).

Fig.2

SEM f r a c t o g r a p h s of N i 7 7 A 1 2 2 Z r l specimens after heat treatment (a) 1 2 0 0 ° C , (b) 1 3 0 0 ° C , a n d (c) 1 3 2 0 ° C for 15 m i n u t e s a n d then w a t e r quenched.

FIG.4

liquated

the

grain

Fracture s u r f a c e of t h e N i 7 7 A122.5Zr0.5 alloy heat treated at 1 3 2 0 ° C a n d t h e n (a) w a t e r q u e n c h e d and (b) furnace cooled.

126

FIG.5

GRAIN BOUNDARY

LIQUATION

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27, No.

EPMA a n a l y s e s of the N i 7 7 A I 2 2 Z r l specimen after a 1320W treatment. (a) S e c o n d a r y e l e c t r o n image, (b) Ni, (b) AI, (c) O, and (d) Zr xray images.

)"(2 )) ['(111)

r'(311) ~" '(220)

~'(io o)

:

(/) c@ ¢m--o

tI/

Q;

>

Q;

¢r

• ,oo~j /(21o)(311),11 r'(11o) J/ I

(531) (440)

FIG. 6

I

I

I

100

90

80

"

,L~.~.-(~V;3 1 1 ( ~ 2 ) ' ~ ~

I

I

70

60

2e (degrees)

50

I

I

40

30

~

20

X-ray d i f f r a c t i o n peaks of the Ni77AI22Zr I alloy after a 1320W treatment. Note that the weak peaks as indicated by t h e a r r o w s were identified as Ni5Zr compound.

1