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Effect of chemical composition on the mechanical properties and ordering behavior of a Ni-Mo alloy
SaiptaMetaUurSicaelMaterialia, Vol. 32, No. IOpp. 1525-1531,1995 Copyight 0 1995 Else&r Scie Ltd I’rinkdtitheUSAAllrights-ed 0956-716x95 $9.50+ .Ocl
Pergamon
0956716x(94)ooo26-3
EFFECT OF CHEMICAL COMPOSITION ON THE MECHANICAL PROPERTIES AND ORDERING BEHAVIOR OF A Ni-Mo ALLOY H. M. Tawancy Materials Characterization Laboratory Metrology, Standards, and Materials Division Research Institute, King Fahd University of Petroleum and Minerals P. 0. Box 1639, Dhahran 31261, Saudi Arabia (Received October 20, 1994) Introduction
One of the corrosion-resistant alloys which has many applications in the chemical process industry is the wrought, Ni-Mo HASTELLOY* alloy B2 (1,2). Its nominal chemical composition is shown in Table 1. When exposed to temperatures in the range of 600 “C to 800 “C, the alloy undergoes a series of long-range ordering reactions to Ni4Mo and other related phases, which can adversely aEect its mechanical strength (3-7) and corrosion resistance (3,5,8). For binary Ni-Mo alloys, only Ni4Mo is formed at MO concentrations of 5 30 weight %, and Ni3Mo also forms when the MO concentration exceeds about 30 weight % (4). It is the objective of this paper to demonstrate that the nature of the ordering reactions, which occur in ahoy B2 and the corresponding etfects on properties can considerably vary from one alloy heat to another depending upon the exact heat composition within the specified range of Table 1. ExDerimentd
Procedure
Two heats of alloy B2 in the form of sheets about 2.5 mm in thickness were included in this study Table 2 show their chemical compositions. As can be seen, both heats contained about the same concentrations of minor elements such as Fe, Cr and Mn, however, the MO content of heat A was about 2 weight % less than that of heat A. Metallographic samples as well as standard tensile samples (SO.8 mm gage length) fi-om the two heats were thermally aged for up to 1000 hours at 600°C 700°C and 800°C in argon atmosphere and
TABLE 1 Nominal Chemical Composition of Alloy B2 (Weight Percent) Ni Balance
MO 26-30
FL*
co 1.o*
Cr 1.o* *Maxirnum
Mn 1.0*
*HAYNES is a rei+sterd badema& of the Haynes Inhnational Company, Kokomo, Indiana.
1525
Si 1.o*
C 0.02*
1526
EFFECT OF CHEMICAL
Vol. 32, No. 10
COMPOSITION
TABLE 2 Chemical Compositions of the Heats of Alloy B2 Studied (Weight Percent) Heat
Ni
MO
A B
Balance Balance
26.92 28.78
co
k3 0:98
<%I
Cr 0.64 0.60
E4 0:25
Si <0.02 <0.02
C 0.002 0.002
then water-quenched to room temperature. All tensile tests of aged samples were conducted at room temperature. An analytical electron microscope operating at 200 keV was used to characterize the ordered microstructure. Thin-foil specimens were prepared by the jet polishing technique in a solution consisting of 30 % nitric acid in methanol by volume at -20°C. X-ray d&&ion (Cu-Ka radiation) and conventional scanning electron microscopy were respectively used to determine the overall structure and morphology of aged samples. Results and Discussion
Figure 1 illustrates the effect of up to 1000 hours of exposure at 6OO”C, 700°C and 800°C on the room-temperature tensile properties of heats A and B (Table 2). A considerable increase in room-temperature yield strength of each heat was observed after exposure at all tempemmres studied (FIG. 1). However, the behavior of tesile ductility varied substantially Tom one heat to another particularlyat 600°C and 800°C. Although both heats reached the same level of yield strength ai& 1000 hours of exposure at 600 “C heat A suffered a complete loss of ductlity while heat B maintained about 60% of its initial ductility (FIG. 1a). At 700°C both heats behaved in a similar fashion (FIG. 1b), however, heat B reached a higher level of yield of yield strength. When the temperature was raised to 8OO”C,the behavior observed at 600°C was reversed, i.e. heat A maintained about 60% of its initial tensile ductility while a complete loss of ductility occurred in heat B. As an example, the secondary electron SEM images of Figure 2 illustrates comparative tensile t?acture modes of heats A and B after 1000 hours of exposure at 600°C. It is observed that the fracture mode of heat A (26.92% MO)was predominantly intergranular. In contrast, the fracture mode of heat B (28.78% MO) was predominantly trausgranular. After 1000 hours of exposure at 7OO”C,the tensile fracture mode of both heats was found to be predominantly intergranular. However, after 1000 hours of exposure at 8OO”C,the tensile fracture mode of heat A (26.92% MO) became transgranular and that of heat B (28.78% MO) became intergranular in opposition to the case observed at 600 “C. In the as-quenched condition, each heat consisted of a disordered fee solid-solution (a-phase) containing short-range order as indicated by the appeamnce of { 1 l/2 0} reflections in electron d&&on patterns similar to the case of binary Ni-Mo alloys (7,9,10). X-ray d&&ion analysis showed that the only ordered phase in both heats after 1000 hours of exposure at 1000 oC was Ni4Mo (tetragonal; a = 0.5720 mn, c = 0.3564 run). Electron difI?action also contirmed that the only ordered phase fnmed in heat A was Ni4Mo as s . sd in Figure 3a. Characteristic reflection of the Dl a superlattice of Ni4Mo are observed at 1/4<42O>fcc and all equivalent positions in the [OOl]fcc electron diffraction pattern of Figure 3a. Da&field imagingrevealed that Ni4Mo in the matrix was in the form of a mosaic assembly of twin-related variants lined up along {lOO}fccplanes (Figure3a). However, there was an evidence for a discontinuous grain boundary reaction at grain boundaries resulting in a lame&r structure of Ni4Mo as shown in Figure 3a. In contrast, heat B, was found to consist of a mixture of Ni4Mo (D 1a superlattice) and the crystallographically related DC22 superlattice (basis for the Ni3Mo stmcture) as shown in the electron dif’lractionpattern of Figure 3~. Dark-field imaging with any of the superlattice reflections revealed dense arrays of discrete particles (Figure 3c), however, there was no evidence for a discontinuous grain boundary reaction. Figure 3d
Vol. 32, No. 10
EFFECT OF CHEMICAL
1100 1000 -
g
l llcn1A 0
Ihat
a 2 60 -
900 -
3m
f In 50c '; 40 .;
@ x 700 3 g
600
5 30 B 2 202 g IO-
* 500 :! 400 Lfff 300
1527
COMPOSITION
0 114112 I 2
4
8
24
100
300 IIIIII 1 1 0 l/4II2 1 2 4 8 24 100 Exposul-eTilne(Iluurs)
Figure 1. Effect of thermal exposure at various B2. (a) 600 “C! (b) 700 “C (c) 800 “C.
IIIUO
I
loon
tempemmres
ll.iIIII) 1 I 0 l/4l/2 1 2 4 R 24 100 1000 E~posureTiln~(Iluurs)
o
l/4l/2 1 2 4
R
24
0
l/41/21 2 4
8
24
on the room-temperature
100
tensile proper&
1000
of heats A and B of alloy
schematically illustrates the relative positions of the characteristic reflections of the Dl a and DO22 superlattices in [OOl]fcc orientation. Both the Dla and DO22 superlattices observed in Figure 3 could be directly derived from the parent disordered fee lattice by minor atoms rearrangements on the (42O)fcc planes (11). In the case of the Dl a superlattice (Ni4Mo), every fifth plane becomes occupied by only MO atoms and planes in-between contain only Ni atoms producing characteristic super-lattice reflections at 1/542O~fcc and all equivalent positions. For the DO22 superlattice, every fourth plane becomes occupied by only MO atoms causing the superlattice reflections to appear at 1/442O>fcc positions (11). According to a disorder-order theory based upon static concentration waves, it was shown that it is possible that a DO22 superlattice with characteristic reflections at (lOO)fcc, { 1 l/2 0)fcc and (1 lO}fcc positions as shown in Figure 3 could form as an intermediate transient
1528
EFFECT OF CHEMICAL
COMPOSITION
Vol. 32, No. 10
phase during long-range ordering to Ni4Mo and Ni3Mo ( 12). Experiment had confirmed this prediction for both Ni4Mo ordering (13,14) and Ni3Mo ordering (15- 17). During long-range ordering to Ni3Mo, it was shown that a Dla superlattice (Ni4Mo) could also form as an intermediate transient phase (17). Thermodynamic calculationsbased upon first- and second-nearest neighbor interactions led to the conclusion that the gmmd-state structures of D 1a, DO22 superlatticeshave similar energies ( 18,19) suggesting that these phases could co-exist as demonstrated in the example of Figure 3.
Figure 2. !kcor&y eleckm SEM knages illustmtkgthe tensile fracture modes of heat A and B alter 1000 hoours of exposure at 600°C. (a) Heat A; interg~anular (b) Heat B; transgranular.
(a)
(4
Figure 3. Onkxl mim of Heat A afkr 1000 hours of exposure at 6000 C. (a) X-ray di&ction pat&n, @) [OOl]fcc electron dil%ction pattern,, (c) Dark-field TEM image formed with the encircled l/%420> fee reflection in (b), (d) Light optical micrograpb illwtratingthegrossmi~, and (e) Bright-field TEM image of an Ni4Mo lamellae at grain boundary
1529
EFFECT OF CHEMICAL, COMPOSITION
Vol. 32, No. 10
OfCC 0
Dl,
l
41
(Ni4Mo)
200
04
(4
(4
Figure 4. Ordered microstructure of Heat B af&r 1000 how of exposure at 600” C. (a) [OOl]fcc electron diGaction pattem, (b) A shematic illustmting the interpretation of the pat&m in (a), (c), A repnxmWive dark-field TEM image corresponding to any of the superlattice reflection in (a), (d) Bright-field TEM image illustrating the grain boundary structure.
@I
(4
Comparative
I
MO I
I
MO
Chemical Comp4 (Weight %)
33.52
26.32
Figure5 WeofHeatBafta 1OOOhoumezqmmre at 800°C. (a) Secondary electron SEM image illwtratingthegross D @) [OOl]fcc electron dit&ction pattern &amte&ic of Ni3Mo, (c) corresponding bright-field TEM image illustrating platelets of Ni3Mo, (d) Energy dispersive x-ray specbum representative of Mi3Mo, and (e) Results of quantifyingthe speckal data in cd).
1530
EFFECT OF CHEMICAL
COMPOSITION
Vol. 32, No. 10
Intergranular embrittlement associated with Ni4Mo ordering in Ni-Mo alloys could be related to a discontiniuous grain boundary ordering reaction as shown in Figure 3a (4,20,21). Earlier studies had shown that the strain energy associated with intragranular ordering could be minimized by self-generated or strain-induced recrystallization resulting in a discontinuous ordering reaction at grain boundaries (22,23). It is possible that the co-existence of the D 1a and DO22 superlattices in heat B had reduced the strain energy associated with intragranular ordering suppressing the grain boundary reaction, which could explain at least partially the retained high ductility (Figure 1a).As the temperature was raised to 700 “C, both x-ray difh-action and electron difhaction showed that the only ordered phase formed in both heats a&r 1000 hours was Ni4Mo. Refarxlce to Figure lb shows that although most of the increase in yield strength occurred after about 1 hour exposure. However, at this both heats maintained most of their initial tensile ductilities. Considerable loss of ductility during the later stages of exposure was correlated with the discontinuous grain boundary reaction as described above. After 1000 hours of exposure at 800°C heat A remained to have Ni4Mo as the only ordered phase. However, in contrast with exposure at 600°C and 700 “, there was no evidence for a discontinuous reaction at grain boundaries, which could explain the high ductilitylevel maintained at& 1000 hours of exposure. This difference in behavior could be related to the effect of temperature on atomic mobility and nucleation site frequency (14). At relatively higher temperature where the rate of volume diffusion is rather high, nucleation site frequency favors matrix precipitation, however, at lower temperatures grain boundary precipitation becomes more pronounced because of the reduced rate of volume diffusion. Embrittlement of heat B during exposure at 800°C was correlated with precipitation of platelets of Ni3Mo as summarized in Figure 6. It is well known that platelets of Ni3Mo are extremely hard and brittle (7). If heat B were to behave as a binary Ni-Mo alloy of the same MO concentration, the only ordered phase formed would be Ni4Mo rather than Ni3Mo. However, as can be seen from the x-ray spectra of Figure 5, Ni3Mo contained marked concentrations of Fe and Cr suggesting that both elements even in relatively small concentration could have the effect of stabilizing Ni3Mo. Evidently, the stabilizing effects of Fe and Cr on Ni3Mo was a sensitive function of the MO content. In contrast, there was no evidence for segregation of either Fe or Cr in Ni4Mo formed in both heats as shown in the example of Figure 6. Conclusions
Variations in mechanical properties among heats of the Ni-Mo alloy B2 were correlated with the effect of small dif%erencesin chemical composition on the long-range ordering behavior of the alloy. It was concluded that for given concentrations of Fe and Cr, the ordering behavior of alloy B2 could considerably deviate Tom that of a binary Ni-Mo ahoy depending upon the MO concentration. Increasing the MO content of the alloy tends to change the ordered phase formed in the alloy from Ni4Mo into Ni3Mo due to the stabilizing effect of relatively small concentrations of Fe and Cr on Ni3Mo. Acknowledments
It is a pleasure to acknowledge the support of the Research Institute of King Fahd University of Petroleum and Minerals and its permission to publish this work. References 1. 2. 3.
F. Galen Hodge in “Con&m and Corrosion Protection Handbook” edited by P. ASchweitzer (Marcel Dekker, New York, 1983) p 55. W. Z. Friend, “Corrosion of Nickel and Nickel-Base alloys” (Wiley, New York, 1980) p 248. H. M. Tawancy, Scripta Metall. et Mater. 30 (a), 713 (1994).
Vol. 32, No. 10
4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 2 1. 22. 23.
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C. R Brooks and Y. M. Wang, Mater&Characterization 25,185 (1990). H. M. Tawancy and N. M. Abbas, J. Mater. Sci. 24,1845 (1989). C. R Brooks and M. Saugeneria,Scripta Metall. 22,1683 (1988). C. R Brooks, J. E. Spruiell end E. E. Stansbuny, Inter. Met. Rev. 29 (3), 210 (1984). H. M. Tawancy andA I. Asphahaniin “High-Teqeratue ordered &rmetaUic Alloys” edited by C. C. Koch, C. T. Liu and N. S. Stoloff (Materials Research Society, Pittsburgh,Pennsylvania,vol. 39,1984) p 485. J. P. Chavalier end W. Stobbs, Acta MetalI. 27.1197 (1979). P. R Okamoto end G. Thomas, Acta Metall. 19,825 (1971). S. K. Das aud G. Thomas, Physica StatusSolidi, 21e, 177 (1974). A G. Kb&atuyan, “order-Disorder Traasitions in Alloys” edited by H. Warlimont, p 114, Springer-Verlag,New York, New York (1974). H. M. Tawancy and M. 0. Aboelfotoh, Physica StatusSolidi, 99a (2), 461(1987). H. M. Tawancy, Scripta Metall. 18,343 (1984). G. Van Tendekm, Mater. Sci. aad Eng. 26,209 (1976). E. Ruedl, Meter. Res. Bull. lo,1267 (1975). G. Van Ten&loo, R. De Ridder and S. Amelinckq Physica. StatusSolidi, 27a, 457 (1975). S. M. Allen and J. W. Calm, Acta Metall. 20,423 (1972). M. T. Richards and J. W. Calm, Acta Metall. 19,1263 (1971). H. M. Tawancy, J. Mater. Sci. 26,3955 (1991). B. Chakravarti,E. Strake and B. G. L.efevre,J. Mater. Sci. 5,394 (1970). L E. Tm and H. J. Leamy, “G&r-Disorder Traasitionsin Alloys” edited by H. Warlimont, p 180, Springer-Verlag,New York, New York (1974). W. B. Snyderand C. R Brooks, ‘Ordered Alloys: Sbu&ral Applications and Physical Metallurgy” edited byB. H. Kear, C. T. Sims, N. S. Stoloff and J. W. We&rook, p 275, Claitor’s Publishing Division, Baton Rouge, LA (1970).
Pergamon
0956716x(94)ooo26-3
EFFECT OF CHEMICAL COMPOSITION ON THE MECHANICAL PROPERTIES AND ORDERING BEHAVIOR OF A Ni-Mo ALLOY H. M. Tawancy Materials Characterization Laboratory Metrology, Standards, and Materials Division Research Institute, King Fahd University of Petroleum and Minerals P. 0. Box 1639, Dhahran 31261, Saudi Arabia (Received October 20, 1994) Introduction
One of the corrosion-resistant alloys which has many applications in the chemical process industry is the wrought, Ni-Mo HASTELLOY* alloy B2 (1,2). Its nominal chemical composition is shown in Table 1. When exposed to temperatures in the range of 600 “C to 800 “C, the alloy undergoes a series of long-range ordering reactions to Ni4Mo and other related phases, which can adversely aEect its mechanical strength (3-7) and corrosion resistance (3,5,8). For binary Ni-Mo alloys, only Ni4Mo is formed at MO concentrations of 5 30 weight %, and Ni3Mo also forms when the MO concentration exceeds about 30 weight % (4). It is the objective of this paper to demonstrate that the nature of the ordering reactions, which occur in ahoy B2 and the corresponding etfects on properties can considerably vary from one alloy heat to another depending upon the exact heat composition within the specified range of Table 1. ExDerimentd
Procedure
Two heats of alloy B2 in the form of sheets about 2.5 mm in thickness were included in this study Table 2 show their chemical compositions. As can be seen, both heats contained about the same concentrations of minor elements such as Fe, Cr and Mn, however, the MO content of heat A was about 2 weight % less than that of heat A. Metallographic samples as well as standard tensile samples (SO.8 mm gage length) fi-om the two heats were thermally aged for up to 1000 hours at 600°C 700°C and 800°C in argon atmosphere and
TABLE 1 Nominal Chemical Composition of Alloy B2 (Weight Percent) Ni Balance
MO 26-30
FL*
co 1.o*
Cr 1.o* *Maxirnum
Mn 1.0*
*HAYNES is a rei+sterd badema& of the Haynes Inhnational Company, Kokomo, Indiana.
1525
Si 1.o*
C 0.02*
1526
EFFECT OF CHEMICAL
Vol. 32, No. 10
COMPOSITION
TABLE 2 Chemical Compositions of the Heats of Alloy B2 Studied (Weight Percent) Heat
Ni
MO
A B
Balance Balance
26.92 28.78
co
k3 0:98
<%I
Cr 0.64 0.60
E4 0:25
Si <0.02 <0.02
C 0.002 0.002
then water-quenched to room temperature. All tensile tests of aged samples were conducted at room temperature. An analytical electron microscope operating at 200 keV was used to characterize the ordered microstructure. Thin-foil specimens were prepared by the jet polishing technique in a solution consisting of 30 % nitric acid in methanol by volume at -20°C. X-ray d&&ion (Cu-Ka radiation) and conventional scanning electron microscopy were respectively used to determine the overall structure and morphology of aged samples. Results and Discussion
Figure 1 illustrates the effect of up to 1000 hours of exposure at 6OO”C, 700°C and 800°C on the room-temperature tensile properties of heats A and B (Table 2). A considerable increase in room-temperature yield strength of each heat was observed after exposure at all tempemmres studied (FIG. 1). However, the behavior of tesile ductility varied substantially Tom one heat to another particularlyat 600°C and 800°C. Although both heats reached the same level of yield strength ai& 1000 hours of exposure at 600 “C heat A suffered a complete loss of ductlity while heat B maintained about 60% of its initial ductility (FIG. 1a). At 700°C both heats behaved in a similar fashion (FIG. 1b), however, heat B reached a higher level of yield of yield strength. When the temperature was raised to 8OO”C,the behavior observed at 600°C was reversed, i.e. heat A maintained about 60% of its initial tensile ductility while a complete loss of ductility occurred in heat B. As an example, the secondary electron SEM images of Figure 2 illustrates comparative tensile t?acture modes of heats A and B after 1000 hours of exposure at 600°C. It is observed that the fracture mode of heat A (26.92% MO)was predominantly intergranular. In contrast, the fracture mode of heat B (28.78% MO) was predominantly trausgranular. After 1000 hours of exposure at 7OO”C,the tensile fracture mode of both heats was found to be predominantly intergranular. However, after 1000 hours of exposure at 8OO”C,the tensile fracture mode of heat A (26.92% MO) became transgranular and that of heat B (28.78% MO) became intergranular in opposition to the case observed at 600 “C. In the as-quenched condition, each heat consisted of a disordered fee solid-solution (a-phase) containing short-range order as indicated by the appeamnce of { 1 l/2 0} reflections in electron d&&on patterns similar to the case of binary Ni-Mo alloys (7,9,10). X-ray d&&ion analysis showed that the only ordered phase in both heats after 1000 hours of exposure at 1000 oC was Ni4Mo (tetragonal; a = 0.5720 mn, c = 0.3564 run). Electron difI?action also contirmed that the only ordered phase fnmed in heat A was Ni4Mo as s . sd in Figure 3a. Characteristic reflection of the Dl a superlattice of Ni4Mo are observed at 1/4<42O>fcc and all equivalent positions in the [OOl]fcc electron diffraction pattern of Figure 3a. Da&field imagingrevealed that Ni4Mo in the matrix was in the form of a mosaic assembly of twin-related variants lined up along {lOO}fccplanes (Figure3a). However, there was an evidence for a discontinuous grain boundary reaction at grain boundaries resulting in a lame&r structure of Ni4Mo as shown in Figure 3a. In contrast, heat B, was found to consist of a mixture of Ni4Mo (D 1a superlattice) and the crystallographically related DC22 superlattice (basis for the Ni3Mo stmcture) as shown in the electron dif’lractionpattern of Figure 3~. Dark-field imaging with any of the superlattice reflections revealed dense arrays of discrete particles (Figure 3c), however, there was no evidence for a discontinuous grain boundary reaction. Figure 3d
Vol. 32, No. 10
EFFECT OF CHEMICAL
1100 1000 -
g
l llcn1A 0
Ihat
a 2 60 -
900 -
3m
f In 50c '; 40 .;
@ x 700 3 g
600
5 30 B 2 202 g IO-
* 500 :! 400 Lfff 300
1527
COMPOSITION
0 114112 I 2
4
8
24
100
300 IIIIII 1 1 0 l/4II2 1 2 4 8 24 100 Exposul-eTilne(Iluurs)
Figure 1. Effect of thermal exposure at various B2. (a) 600 “C! (b) 700 “C (c) 800 “C.
IIIUO
I
loon
tempemmres
ll.iIIII) 1 I 0 l/4l/2 1 2 4 R 24 100 1000 E~posureTiln~(Iluurs)
o
l/4l/2 1 2 4
R
24
0
l/41/21 2 4
8
24
on the room-temperature
100
tensile proper&
1000
of heats A and B of alloy
schematically illustrates the relative positions of the characteristic reflections of the Dl a and DO22 superlattices in [OOl]fcc orientation. Both the Dla and DO22 superlattices observed in Figure 3 could be directly derived from the parent disordered fee lattice by minor atoms rearrangements on the (42O)fcc planes (11). In the case of the Dl a superlattice (Ni4Mo), every fifth plane becomes occupied by only MO atoms and planes in-between contain only Ni atoms producing characteristic super-lattice reflections at 1/542O~fcc and all equivalent positions. For the DO22 superlattice, every fourth plane becomes occupied by only MO atoms causing the superlattice reflections to appear at 1/442O>fcc positions (11). According to a disorder-order theory based upon static concentration waves, it was shown that it is possible that a DO22 superlattice with characteristic reflections at (lOO)fcc, { 1 l/2 0)fcc and (1 lO}fcc positions as shown in Figure 3 could form as an intermediate transient
1528
EFFECT OF CHEMICAL
COMPOSITION
Vol. 32, No. 10
phase during long-range ordering to Ni4Mo and Ni3Mo ( 12). Experiment had confirmed this prediction for both Ni4Mo ordering (13,14) and Ni3Mo ordering (15- 17). During long-range ordering to Ni3Mo, it was shown that a Dla superlattice (Ni4Mo) could also form as an intermediate transient phase (17). Thermodynamic calculationsbased upon first- and second-nearest neighbor interactions led to the conclusion that the gmmd-state structures of D 1a, DO22 superlatticeshave similar energies ( 18,19) suggesting that these phases could co-exist as demonstrated in the example of Figure 3.
Figure 2. !kcor&y eleckm SEM knages illustmtkgthe tensile fracture modes of heat A and B alter 1000 hoours of exposure at 600°C. (a) Heat A; interg~anular (b) Heat B; transgranular.
(a)
(4
Figure 3. Onkxl mim of Heat A afkr 1000 hours of exposure at 6000 C. (a) X-ray di&ction pat&n, @) [OOl]fcc electron dil%ction pattern,, (c) Dark-field TEM image formed with the encircled l/%420> fee reflection in (b), (d) Light optical micrograpb illwtratingthegrossmi~, and (e) Bright-field TEM image of an Ni4Mo lamellae at grain boundary
1529
EFFECT OF CHEMICAL, COMPOSITION
Vol. 32, No. 10
OfCC 0
Dl,
l
41
(Ni4Mo)
200
04
(4
(4
Figure 4. Ordered microstructure of Heat B af&r 1000 how of exposure at 600” C. (a) [OOl]fcc electron diGaction pattem, (b) A shematic illustmting the interpretation of the pat&m in (a), (c), A repnxmWive dark-field TEM image corresponding to any of the superlattice reflection in (a), (d) Bright-field TEM image illustrating the grain boundary structure.
@I
(4
Comparative
I
MO I
I
MO
Chemical Comp4 (Weight %)
33.52
26.32
Figure5 WeofHeatBafta 1OOOhoumezqmmre at 800°C. (a) Secondary electron SEM image illwtratingthegross D @) [OOl]fcc electron dit&ction pattern &amte&ic of Ni3Mo, (c) corresponding bright-field TEM image illustrating platelets of Ni3Mo, (d) Energy dispersive x-ray specbum representative of Mi3Mo, and (e) Results of quantifyingthe speckal data in cd).
1530
EFFECT OF CHEMICAL
COMPOSITION
Vol. 32, No. 10
Intergranular embrittlement associated with Ni4Mo ordering in Ni-Mo alloys could be related to a discontiniuous grain boundary ordering reaction as shown in Figure 3a (4,20,21). Earlier studies had shown that the strain energy associated with intragranular ordering could be minimized by self-generated or strain-induced recrystallization resulting in a discontinuous ordering reaction at grain boundaries (22,23). It is possible that the co-existence of the D 1a and DO22 superlattices in heat B had reduced the strain energy associated with intragranular ordering suppressing the grain boundary reaction, which could explain at least partially the retained high ductility (Figure 1a).As the temperature was raised to 700 “C, both x-ray difh-action and electron difhaction showed that the only ordered phase formed in both heats a&r 1000 hours was Ni4Mo. Refarxlce to Figure lb shows that although most of the increase in yield strength occurred after about 1 hour exposure. However, at this both heats maintained most of their initial tensile ductilities. Considerable loss of ductility during the later stages of exposure was correlated with the discontinuous grain boundary reaction as described above. After 1000 hours of exposure at 800°C heat A remained to have Ni4Mo as the only ordered phase. However, in contrast with exposure at 600°C and 700 “, there was no evidence for a discontinuous reaction at grain boundaries, which could explain the high ductilitylevel maintained at& 1000 hours of exposure. This difference in behavior could be related to the effect of temperature on atomic mobility and nucleation site frequency (14). At relatively higher temperature where the rate of volume diffusion is rather high, nucleation site frequency favors matrix precipitation, however, at lower temperatures grain boundary precipitation becomes more pronounced because of the reduced rate of volume diffusion. Embrittlement of heat B during exposure at 800°C was correlated with precipitation of platelets of Ni3Mo as summarized in Figure 6. It is well known that platelets of Ni3Mo are extremely hard and brittle (7). If heat B were to behave as a binary Ni-Mo alloy of the same MO concentration, the only ordered phase formed would be Ni4Mo rather than Ni3Mo. However, as can be seen from the x-ray spectra of Figure 5, Ni3Mo contained marked concentrations of Fe and Cr suggesting that both elements even in relatively small concentration could have the effect of stabilizing Ni3Mo. Evidently, the stabilizing effects of Fe and Cr on Ni3Mo was a sensitive function of the MO content. In contrast, there was no evidence for segregation of either Fe or Cr in Ni4Mo formed in both heats as shown in the example of Figure 6. Conclusions
Variations in mechanical properties among heats of the Ni-Mo alloy B2 were correlated with the effect of small dif%erencesin chemical composition on the long-range ordering behavior of the alloy. It was concluded that for given concentrations of Fe and Cr, the ordering behavior of alloy B2 could considerably deviate Tom that of a binary Ni-Mo ahoy depending upon the MO concentration. Increasing the MO content of the alloy tends to change the ordered phase formed in the alloy from Ni4Mo into Ni3Mo due to the stabilizing effect of relatively small concentrations of Fe and Cr on Ni3Mo. Acknowledments
It is a pleasure to acknowledge the support of the Research Institute of King Fahd University of Petroleum and Minerals and its permission to publish this work. References 1. 2. 3.
F. Galen Hodge in “Con&m and Corrosion Protection Handbook” edited by P. ASchweitzer (Marcel Dekker, New York, 1983) p 55. W. Z. Friend, “Corrosion of Nickel and Nickel-Base alloys” (Wiley, New York, 1980) p 248. H. M. Tawancy, Scripta Metall. et Mater. 30 (a), 713 (1994).
Vol. 32, No. 10
4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 2 1. 22. 23.
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