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The effect of constant-load creep on fracture toughness and tensile behavior of precipitation-free zone aluminum alloy type 2618
Scripta METALLURGICA
Vol. 22, pp. 1503-1508, 1988 Printed in the U.S.A.
Pergamon Press plc All rights reserved
THE EFFECTOF CONSTANT-LOAD CREEPON FRACTURE TOUGHNESS AND TENSILE BEHAVIOR OF PRECIPITATION-FREE ZONEALUMINUMALLOYTYPE 2618 D. BobrowI, A. Arbel2 and D. EliezerI I. Department of Materials Engineering, Ben-Gurion University, Beer Sheva, 84153 Israel 2. Department of Metallurgical Engineering, The Ohio State University, Columbus, OH 43210 USA (Received June 8, 1988) (Revised July i, 1988) Introduction Modern design c r i t e r i a of energy generating and conversion systems which operate at elevated temperatures incorporate information related to tensile, creep, fatigue and fracture toughness properties of potential candidate materials to be used in such systems. Hightemperature mechanical loading can introduce several microstructural changes in crystalline materials such as: (I) nucleation, growth and migration of new phases; (2) dimensional and morphological changes, migration and disappearance of old phases; (3) grain boundary sliding and migration; (4) formation of subgrains; (5) changes of dislocation density, configuration and distribution; (6) nucleation, growth and coalescence of voids especially in the v i c i n i t y of grain boundaries; (7) migration of undesirable elements to grain boundaries. Someof these microstructural changes are interdependent, some w i l l happen by mere exposure to elevated temperatures, while others may be accelerated by high-temperature loading and some may not occur at all without m~chanical loading. Someof the aforementioned microstructural changes can cause reduction in the fracture toughness of the materials to dangerous values well below those which these materials had before been put to service. Thus, a situation may develop whereby the fracture toughness of materials, which serve at elevated temperatures, may reach low unacceptable values long before these materials have exhausted their lives as determined from creep or fatigue life-prediction methods. I t was suggested that the useful lives of materials which serve while loaded at high temperatures may have to be determined by their residual fracture toughness rather than by more conventional creep and fatigue considerations [ I ] . Two exploratory investigations revealed that formation and coalescence of grain-boundary voids caused sharp reductions in fracture toughness long before a solid solution and a precipitation strengthened nickel base superalloys approached the end of their creep lives [1,2]. Intergranular and intragranular carbides, which formed during the early stages of exposing 316 stainless steel to a high temperature soak and to creep, were considered as the main contributors to the observed reduction in uniform and t o t a l tensile elongations and to the decrease in impact energy to fracture of Charpy keyhole specimens [3]. Monotonic and cyclic creep reduced somewhat the fracture toughness of precipitation-hardened aluminum alloy type 2618-T61 primarily because of an increase of dislocation density due to the need to compensate for the reduction in cross section which accompanied longitudinal elongation [4]. These two modes of creep did not weaken the grain boundaries of the aluminum alloy and the percentage of intergranular separation observed on fracture surface of fracture-toughness specimen was the same for the virgin and for the crept alloy [4]. Improper heat treating of p r e c i p i t a t i o n strengthened aluminum alloys may result in a precipitation free zone (PFZ) especially along grain boundaries. The purpose of this note is to report the effect of constant-load creep on the resultant fracture toughness of a commercial aluminum alloy which was intentionally treated to produce PFZ along i t s grain boundaries. Experimental The commercial precipitation-strengthened aluminum alloy type 2618 was selected for this study. Pieces of the alloy having dimensions slightly greater than those of the creep specimens were solution-treated for 2 h at 525°C and rather than receiving a subsequent quench i n t o water at 70-80°C, the pieces were cooled to IOO°C at a constant rate of 13.7°C/min, followed by the standard aging for the alloy of 20 h at 2OO°C. Creep specimens of 16 mm diameter and 150 mmgage length having a total length of over 213 mm, were machined from the heat treated pieces. The location and orientation of all the creep specimens were the same with respect to the original 152 mmdiameter extruded stock-bar of the alloy from 1503 0036-9748/88 $3.00 + .00
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BEHAVIOR OF PRECIPITATION-FREE
ZONE ALUMINUM
Vol.
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which all the specimens were machined. This procedure was employed in order to minimize variations in mechanical behavior due to possible anisotropy and inhomogeneity in the original stock bar. A constant nominal tensile load of 275 MPa was applied to the creep specimens while they were placed in a 3-zone s p l i t furnace where the temperature was maintained at 150±1°C along the specimen and its grips. Testing temperature of 150°C was selected because i t did not seem to affect the morphology and size of precipitates in this alloy nor its tensile properties and fracture toughness after the high temperature soak for times equal or somewhat longer than the durations of the creep tests [4]. The creep elongation of the specimens was continuously recorded for increasing loading periods of up to 180 hours. Upon termination of each creep t e s t , the s t i l l loaded specimens were cooled rapidly to room temperature with the aid of forced air. One standard tensile specimen and a multiple-notch three-point bend specimen were machined from the original large creep specimen. A sharp fatigue crack was introduced at the root of each of the original notches in the bend specimen. Both types of specimens were used to study the resultant tensile properties and fracture toughness as affected by the prior creep. The 3-point bending test of the multiple-notch and fatigue-precracked specimen is depicted in Fig. la. The horizontal arrows on both i t s sides indicate the direction of the loading during the prior creep process. The load F was applied at a constant mid-section deflection rate of I mm/min. A typical load deflection curve is shown in Fig. lb. The amount of specific energy spent to bend the precracked specimen beyond the peak of the load-deflection curve and down to a load equal to one tenth of the peak load was selected as the measure of the a l l o y resistance to extension of sharp cracks or i t s fracture toughness. This specific energy is equal to W/A where Wis the area under the load deflection curve and A is the area of the uncracked ligament before bending begins. The terms Wand A and the dimensions of the bend specimen are depicted in Fig. I. The reasons for selecting the aforementioned measure for fracture toughness will be discussed later. All the tensile tests in this work were carried out at a nominal strain rate of O.01/min. Optical, scanning and transmission electron microscopy were used to study the m i c r o s t r u c t u r e of the a l l o y before and a f t e r the various stages of mechanical testing. Results and Discussion Transmission electron microscopy revealed that the slow'cooling of the alloy from the solutionizing temperature indeed produced a precipitation free zone which was extended to an average distance of about 2 #m on both sides of every observed grain boundary. TEM of the same alloy receiving the standard T61 treatment revealed precipitates extending all the way to the grain boundaries [4]. The slow cooling also affected the tensile and the fracture toughness of the alloy. We observed in Table I that the treatment to obtain PFZ resulted in lower yield and ultimate stresses, nearly the same elongation to fracture, yet lower reduction in area of fracture but higher specific energy to fracture of precracked bend specimens (W/A). TABLE I:
Temper
T61 PFZ
Room-Temperature Tensile and Fracture Properties of the 2618-T61 and as Treated to Produce PFZ
2% Y.S. UTS
Elongation
[MPa]
[MPal
to fracture (%) in Area [%]
372 319
432 401
10.4 11.0
Reduction
20.6 15.9
W/A [KJ/m2l 14.7 19.2
fA IGF [%l 3.0 4.9
The term f~GF is the percentage of area over which intergranular fracture took place during the bending of the precracked specimens as shown in Fig. la. The fracture surface of the alloy, before the creep process, was primarily dimpled intragranular with intermetallic Al2FeNi particles observed at the bottom of the dimples. ThE reduction in y i e l d stress and ultimate tensile strength and the increase of fracture toughness of the alloy which received the PFZ treatment can be attributed to the coarser needles of the f i n e S'-AI2CuMg strengthening pre- cipitate as compared to the finer ones observed by TEM in this alloy after i t received the standard T61 treatment. I t was possible to machine one standard tensile specimen having a gage length of 25 mm and a fracture toughness specimen having 3 notches with 3 fatigue precracks from each of the large creep specimens. The room temperature
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22, No.
9
B E H A V I O R OF P R E C I P I T A T I O N - F R E E
TABLE I I : Specimen No.
[h] 57
58
59
60
94
73
184
113
1505
RoomTemperature Mechanical Properties After Creep
Creep Data Time
ZONE A L U M I N U M
Strain
Resultant R.T. Properties UTS
YS
[%]
[MPa]
[MPa]
1.8 .82 .83 .88
3gl
342
.64 .94 .95 .92
384
2.5 .77 .87 .80
399
2.1 .92 .87 .85
390
338
356
348
Elong. [%] 8.5
8.5
7.9
8.3
RA [%]
W/A
fAGF
[KJ/m2]
[%]
13.4 17.2 17.2 18.2
6.0
17.6 18.8 1g.6
5.0
14.8 14.7 14.4
10.5
16.8 16.5 16.0
8.1
19.0
15.9
17.6
resultant properties are summarized in Table I I . The creep deformation was not absolutely uniform and the permanent deformation was determined from diametrical measurements at the planes of the fatigue precracks in the bend specimens. The prior creep strain related to the tensile properties is the average value along the gage length of the secondary tensile specimen machined out of the original creep specimen. The resultant higher yield stress, lower elongation to fracture and lower reduction in area as compared to these properties of the as-heat-treated alloy can be attributed to the increase in dislocation density during the creep process. Such an increase was also observed by TEM in the crept alloy which had received the standard T6I treatment [4]. As seen in Table I I , the prior creep strain at the plane of fracture of the precracked bend-specimens was less than I%. No correlation was found between .the prior creep strain and the resultant fracture toughness. Figures 2-3 are based on the data from Table I f . Figure 2 depicts the time dependence of the resultant fracture toughness on prior creep time. There is only a very slight reduction in fracture energy after 73 hours of creep followed by a Faster reduction reaching a 24% drop after 184 hours of creep with the total prior creep strain, again, being less than I%. Figure 2 also depicts that a f t e r what seems to be an incubation period of about 70 hours, there is nearly a linear relation between prior creeptime and the percentage of area over which intergranular fracture took place in the precracked bend specimens. Figure 3 shows a nearly linear dependence between the reduction of Fracture toughness and the percentage of intergranular fracture in the bend precracked specimens. Examining the resultant fracture toughness as presented numerically in Table I I and graphically in Figs. 2-3, i t is clear that some kind of time-dependent grain-boundary damage is accumulated during creep. The amount of damage is apparently independent of prior deformation at least for creep strains of up to I%. This damage resulted in increased separation along grain boundaries during fracturing of the precracked bend specimens. Figure 4 depicts such grain boundary facets which are covered with small dimples while at the upper area and the lower l e f t portion of the SEM photograph we observe the much larger and deeper intragranular dimples with intermetallic Al2FeNi particles s t i l l situated at their bottom. I t is not clear whether the small dimples over the grain boundaries were formed during the i n i t i a l creep process or only during fracturing by bending. Our observations did not reveal any separation along grain boundaries during creep and before bending was carried out. The intergranular fracture with small dimples requires less energy per unit area than does the
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intragranular fracture with the much larger and deeper dimples which are associated with larger amount of localized plastic deformation. Thus, the larger the fraction of the area over which fracture takes place intergranularly, the lower is the energy required for fracture and the lower is the fracture toughness of the crept alloy. The bend specimens machined out of the original large creep specimens were too small even for a direct valid JIC determination. Yet, work by others showed a good correlation between KIC values determined from large valid specimens of aluminum alloys and fracture energy measurements from precracked bend specimens similar to those used by us [5]. We did not use the short rod specimen for which measured KICSR values are s i m i l a r to KIC values obtained for aluminum alloys using valid size specimens [6,7] because the plane of fracture of these specimens is parallel rather than being perpendicular to the loading axis of the original creep specimen. More damage was expected in cross sections perpendicular to the loading axis during the creep process. The precracked bend-specimens in this work were oriented to detect such damage. Attempts have been made over the years to empirically predict creep l i f e [8,9]. Others have attempted to predict rupture time at the end of the creep process due to growth of grain boundary cavities [10-13]. The results of this work indicate that the fracture toughness of a precipitation hardened aluminum alloy, which was treated to form PFZ along i t s grain boundaries, can be lowered considerably after less than I% of elevated-temperature creep strain. The creep process introduced time dependent damage, which, a f t e r what seemed l i k e an incubation period, resulted in increased separation along grain boundaries during fracturing of precracked bend-specimens machined from the original creep specimens. The results of this work indicate that the resultant fracture toughness of materials loaded at high temperature can be reduced appreciably after very small creep deformations. Thus, changes in fracture toughness can be an important factor in determining e f f e c t i v e and safe l i v e s of materials which have to serve while loaded at high temperatures. Conclusion Constant-load elevated-temperature creep with strains of less than I% caused considerable reduction in fracture toughness of an aluminum alloy which was treated to form a precipitation free zone along its grain boundaries. The reduction in fracture toughness was due to time-dependent damage resulting in increasing amounts of intergranular fracture as compared to that observed in the alloy before i t crept. The results of this work indicate the importance of considering possible deleterious changes in fracture toughness as a c r i t e r i o n which may determine the useful l i f e of materials serving while loaded at high temperatures. References I. A. Arbel, Scripta Met. 13, 1109 (]979). 2. A. Arbel and C.S. Pande, J. Mat. Sci. in press. 3. D. Gan, Metall. Trans. A, 13A, 2155 (1982). 4. D. Bobrow, A. Arbel and D. Eliezer, to be published. 5. G. Succop, R.T. Bubsey, M.H. Jones and W.F. Brown, Jr., in "Developments in Fracture Mechanics Test Methods Standardization", ASTM, STP 632, 153 (1977). 6. L.M. Barker and F.I. Barata, J. Test. Eval., 8, 97 (1980). 7. L.M. Barker, Int. J. Frac., 15, 515 (1979). 8. F.C. Monkman and N.J. Grant, Proc. ASTM56, 593 (1956). 9. F. Dobes and K. Milicka, Met. Sci. 10, 382 (1976). 10. W.D. Nix, Scripta Met. 17, I, (1983). I I . A.S. Argon, Scripta Met. 17, 5, (1983). 12. W. Beere, Scripta Met. ]7, 13, (1983). 13. J.S. Wang, L. Martinez and W.D. Nix, Acta Metall. 31, 873 (1983).
Vol.
22, No.
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B E H A V I O R OF P R E C I P I T A T I O N - F R E E
IF
(a)
ZONE A L U M I N U M
1507
(b)
Load
F.~x I0
--T 40 45
0.1Frnox " / / / / / / / / / / , 1 1 ~ Mid-Point Deflection
FIG. 1. (a) Three point bending of a multiple-notched and precracked fracture toughness specimens with i t s dimensions in mm; (b) Typical load vs. mid-point d e f l e c t i o n curve of the specimen in (a).
n2 ~, ~
~--,~- ~
10
19 I
o I,..,
L..
~._
L,
OIL
~ 15 0
•---,v I 0 50
I I00
I 150
4 200
n
FIG. 2. The dependence of f r a c t u r e energy (W/A) and percentage of i n t e r g r a n u l a r f r a c t u r e (f~GF)~ on creep time.
Creep Time (h)
21
E
I
I
I
I
I
I
19 FIG. 3. The dependence of f r a c t u r e energy on percentage of intergranular fracture.
17 E
w 15 0
13 4
I 5 Percent
I 6
I 7
I 8
I 9
I I0
of I n t e r g r a n u l a r F r a c t u r e
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B E H A V I O R OF P R E C I P I T A T I O N - F R E E
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FIG. 4. Fracture surface of bend specimen in the v i c i n i t y of grain boundary facets.
Vol.
22, No.
9
Vol. 22, pp. 1503-1508, 1988 Printed in the U.S.A.
Pergamon Press plc All rights reserved
THE EFFECTOF CONSTANT-LOAD CREEPON FRACTURE TOUGHNESS AND TENSILE BEHAVIOR OF PRECIPITATION-FREE ZONEALUMINUMALLOYTYPE 2618 D. BobrowI, A. Arbel2 and D. EliezerI I. Department of Materials Engineering, Ben-Gurion University, Beer Sheva, 84153 Israel 2. Department of Metallurgical Engineering, The Ohio State University, Columbus, OH 43210 USA (Received June 8, 1988) (Revised July i, 1988) Introduction Modern design c r i t e r i a of energy generating and conversion systems which operate at elevated temperatures incorporate information related to tensile, creep, fatigue and fracture toughness properties of potential candidate materials to be used in such systems. Hightemperature mechanical loading can introduce several microstructural changes in crystalline materials such as: (I) nucleation, growth and migration of new phases; (2) dimensional and morphological changes, migration and disappearance of old phases; (3) grain boundary sliding and migration; (4) formation of subgrains; (5) changes of dislocation density, configuration and distribution; (6) nucleation, growth and coalescence of voids especially in the v i c i n i t y of grain boundaries; (7) migration of undesirable elements to grain boundaries. Someof these microstructural changes are interdependent, some w i l l happen by mere exposure to elevated temperatures, while others may be accelerated by high-temperature loading and some may not occur at all without m~chanical loading. Someof the aforementioned microstructural changes can cause reduction in the fracture toughness of the materials to dangerous values well below those which these materials had before been put to service. Thus, a situation may develop whereby the fracture toughness of materials, which serve at elevated temperatures, may reach low unacceptable values long before these materials have exhausted their lives as determined from creep or fatigue life-prediction methods. I t was suggested that the useful lives of materials which serve while loaded at high temperatures may have to be determined by their residual fracture toughness rather than by more conventional creep and fatigue considerations [ I ] . Two exploratory investigations revealed that formation and coalescence of grain-boundary voids caused sharp reductions in fracture toughness long before a solid solution and a precipitation strengthened nickel base superalloys approached the end of their creep lives [1,2]. Intergranular and intragranular carbides, which formed during the early stages of exposing 316 stainless steel to a high temperature soak and to creep, were considered as the main contributors to the observed reduction in uniform and t o t a l tensile elongations and to the decrease in impact energy to fracture of Charpy keyhole specimens [3]. Monotonic and cyclic creep reduced somewhat the fracture toughness of precipitation-hardened aluminum alloy type 2618-T61 primarily because of an increase of dislocation density due to the need to compensate for the reduction in cross section which accompanied longitudinal elongation [4]. These two modes of creep did not weaken the grain boundaries of the aluminum alloy and the percentage of intergranular separation observed on fracture surface of fracture-toughness specimen was the same for the virgin and for the crept alloy [4]. Improper heat treating of p r e c i p i t a t i o n strengthened aluminum alloys may result in a precipitation free zone (PFZ) especially along grain boundaries. The purpose of this note is to report the effect of constant-load creep on the resultant fracture toughness of a commercial aluminum alloy which was intentionally treated to produce PFZ along i t s grain boundaries. Experimental The commercial precipitation-strengthened aluminum alloy type 2618 was selected for this study. Pieces of the alloy having dimensions slightly greater than those of the creep specimens were solution-treated for 2 h at 525°C and rather than receiving a subsequent quench i n t o water at 70-80°C, the pieces were cooled to IOO°C at a constant rate of 13.7°C/min, followed by the standard aging for the alloy of 20 h at 2OO°C. Creep specimens of 16 mm diameter and 150 mmgage length having a total length of over 213 mm, were machined from the heat treated pieces. The location and orientation of all the creep specimens were the same with respect to the original 152 mmdiameter extruded stock-bar of the alloy from 1503 0036-9748/88 $3.00 + .00
1504
BEHAVIOR OF PRECIPITATION-FREE
ZONE ALUMINUM
Vol.
22, No.
9
which all the specimens were machined. This procedure was employed in order to minimize variations in mechanical behavior due to possible anisotropy and inhomogeneity in the original stock bar. A constant nominal tensile load of 275 MPa was applied to the creep specimens while they were placed in a 3-zone s p l i t furnace where the temperature was maintained at 150±1°C along the specimen and its grips. Testing temperature of 150°C was selected because i t did not seem to affect the morphology and size of precipitates in this alloy nor its tensile properties and fracture toughness after the high temperature soak for times equal or somewhat longer than the durations of the creep tests [4]. The creep elongation of the specimens was continuously recorded for increasing loading periods of up to 180 hours. Upon termination of each creep t e s t , the s t i l l loaded specimens were cooled rapidly to room temperature with the aid of forced air. One standard tensile specimen and a multiple-notch three-point bend specimen were machined from the original large creep specimen. A sharp fatigue crack was introduced at the root of each of the original notches in the bend specimen. Both types of specimens were used to study the resultant tensile properties and fracture toughness as affected by the prior creep. The 3-point bending test of the multiple-notch and fatigue-precracked specimen is depicted in Fig. la. The horizontal arrows on both i t s sides indicate the direction of the loading during the prior creep process. The load F was applied at a constant mid-section deflection rate of I mm/min. A typical load deflection curve is shown in Fig. lb. The amount of specific energy spent to bend the precracked specimen beyond the peak of the load-deflection curve and down to a load equal to one tenth of the peak load was selected as the measure of the a l l o y resistance to extension of sharp cracks or i t s fracture toughness. This specific energy is equal to W/A where Wis the area under the load deflection curve and A is the area of the uncracked ligament before bending begins. The terms Wand A and the dimensions of the bend specimen are depicted in Fig. I. The reasons for selecting the aforementioned measure for fracture toughness will be discussed later. All the tensile tests in this work were carried out at a nominal strain rate of O.01/min. Optical, scanning and transmission electron microscopy were used to study the m i c r o s t r u c t u r e of the a l l o y before and a f t e r the various stages of mechanical testing. Results and Discussion Transmission electron microscopy revealed that the slow'cooling of the alloy from the solutionizing temperature indeed produced a precipitation free zone which was extended to an average distance of about 2 #m on both sides of every observed grain boundary. TEM of the same alloy receiving the standard T61 treatment revealed precipitates extending all the way to the grain boundaries [4]. The slow cooling also affected the tensile and the fracture toughness of the alloy. We observed in Table I that the treatment to obtain PFZ resulted in lower yield and ultimate stresses, nearly the same elongation to fracture, yet lower reduction in area of fracture but higher specific energy to fracture of precracked bend specimens (W/A). TABLE I:
Temper
T61 PFZ
Room-Temperature Tensile and Fracture Properties of the 2618-T61 and as Treated to Produce PFZ
2% Y.S. UTS
Elongation
[MPa]
[MPal
to fracture (%) in Area [%]
372 319
432 401
10.4 11.0
Reduction
20.6 15.9
W/A [KJ/m2l 14.7 19.2
fA IGF [%l 3.0 4.9
The term f~GF is the percentage of area over which intergranular fracture took place during the bending of the precracked specimens as shown in Fig. la. The fracture surface of the alloy, before the creep process, was primarily dimpled intragranular with intermetallic Al2FeNi particles observed at the bottom of the dimples. ThE reduction in y i e l d stress and ultimate tensile strength and the increase of fracture toughness of the alloy which received the PFZ treatment can be attributed to the coarser needles of the f i n e S'-AI2CuMg strengthening pre- cipitate as compared to the finer ones observed by TEM in this alloy after i t received the standard T61 treatment. I t was possible to machine one standard tensile specimen having a gage length of 25 mm and a fracture toughness specimen having 3 notches with 3 fatigue precracks from each of the large creep specimens. The room temperature
Vol.
22, No.
9
B E H A V I O R OF P R E C I P I T A T I O N - F R E E
TABLE I I : Specimen No.
[h] 57
58
59
60
94
73
184
113
1505
RoomTemperature Mechanical Properties After Creep
Creep Data Time
ZONE A L U M I N U M
Strain
Resultant R.T. Properties UTS
YS
[%]
[MPa]
[MPa]
1.8 .82 .83 .88
3gl
342
.64 .94 .95 .92
384
2.5 .77 .87 .80
399
2.1 .92 .87 .85
390
338
356
348
Elong. [%] 8.5
8.5
7.9
8.3
RA [%]
W/A
fAGF
[KJ/m2]
[%]
13.4 17.2 17.2 18.2
6.0
17.6 18.8 1g.6
5.0
14.8 14.7 14.4
10.5
16.8 16.5 16.0
8.1
19.0
15.9
17.6
resultant properties are summarized in Table I I . The creep deformation was not absolutely uniform and the permanent deformation was determined from diametrical measurements at the planes of the fatigue precracks in the bend specimens. The prior creep strain related to the tensile properties is the average value along the gage length of the secondary tensile specimen machined out of the original creep specimen. The resultant higher yield stress, lower elongation to fracture and lower reduction in area as compared to these properties of the as-heat-treated alloy can be attributed to the increase in dislocation density during the creep process. Such an increase was also observed by TEM in the crept alloy which had received the standard T6I treatment [4]. As seen in Table I I , the prior creep strain at the plane of fracture of the precracked bend-specimens was less than I%. No correlation was found between .the prior creep strain and the resultant fracture toughness. Figures 2-3 are based on the data from Table I f . Figure 2 depicts the time dependence of the resultant fracture toughness on prior creep time. There is only a very slight reduction in fracture energy after 73 hours of creep followed by a Faster reduction reaching a 24% drop after 184 hours of creep with the total prior creep strain, again, being less than I%. Figure 2 also depicts that a f t e r what seems to be an incubation period of about 70 hours, there is nearly a linear relation between prior creeptime and the percentage of area over which intergranular fracture took place in the precracked bend specimens. Figure 3 shows a nearly linear dependence between the reduction of Fracture toughness and the percentage of intergranular fracture in the bend precracked specimens. Examining the resultant fracture toughness as presented numerically in Table I I and graphically in Figs. 2-3, i t is clear that some kind of time-dependent grain-boundary damage is accumulated during creep. The amount of damage is apparently independent of prior deformation at least for creep strains of up to I%. This damage resulted in increased separation along grain boundaries during fracturing of the precracked bend specimens. Figure 4 depicts such grain boundary facets which are covered with small dimples while at the upper area and the lower l e f t portion of the SEM photograph we observe the much larger and deeper intragranular dimples with intermetallic Al2FeNi particles s t i l l situated at their bottom. I t is not clear whether the small dimples over the grain boundaries were formed during the i n i t i a l creep process or only during fracturing by bending. Our observations did not reveal any separation along grain boundaries during creep and before bending was carried out. The intergranular fracture with small dimples requires less energy per unit area than does the
1506
B E H A V I O R OF P R E C I P I T A T I O N - F R E E
ZONE A L U M I N U M
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intragranular fracture with the much larger and deeper dimples which are associated with larger amount of localized plastic deformation. Thus, the larger the fraction of the area over which fracture takes place intergranularly, the lower is the energy required for fracture and the lower is the fracture toughness of the crept alloy. The bend specimens machined out of the original large creep specimens were too small even for a direct valid JIC determination. Yet, work by others showed a good correlation between KIC values determined from large valid specimens of aluminum alloys and fracture energy measurements from precracked bend specimens similar to those used by us [5]. We did not use the short rod specimen for which measured KICSR values are s i m i l a r to KIC values obtained for aluminum alloys using valid size specimens [6,7] because the plane of fracture of these specimens is parallel rather than being perpendicular to the loading axis of the original creep specimen. More damage was expected in cross sections perpendicular to the loading axis during the creep process. The precracked bend-specimens in this work were oriented to detect such damage. Attempts have been made over the years to empirically predict creep l i f e [8,9]. Others have attempted to predict rupture time at the end of the creep process due to growth of grain boundary cavities [10-13]. The results of this work indicate that the fracture toughness of a precipitation hardened aluminum alloy, which was treated to form PFZ along i t s grain boundaries, can be lowered considerably after less than I% of elevated-temperature creep strain. The creep process introduced time dependent damage, which, a f t e r what seemed l i k e an incubation period, resulted in increased separation along grain boundaries during fracturing of precracked bend-specimens machined from the original creep specimens. The results of this work indicate that the resultant fracture toughness of materials loaded at high temperature can be reduced appreciably after very small creep deformations. Thus, changes in fracture toughness can be an important factor in determining e f f e c t i v e and safe l i v e s of materials which have to serve while loaded at high temperatures. Conclusion Constant-load elevated-temperature creep with strains of less than I% caused considerable reduction in fracture toughness of an aluminum alloy which was treated to form a precipitation free zone along its grain boundaries. The reduction in fracture toughness was due to time-dependent damage resulting in increasing amounts of intergranular fracture as compared to that observed in the alloy before i t crept. The results of this work indicate the importance of considering possible deleterious changes in fracture toughness as a c r i t e r i o n which may determine the useful l i f e of materials serving while loaded at high temperatures. References I. A. Arbel, Scripta Met. 13, 1109 (]979). 2. A. Arbel and C.S. Pande, J. Mat. Sci. in press. 3. D. Gan, Metall. Trans. A, 13A, 2155 (1982). 4. D. Bobrow, A. Arbel and D. Eliezer, to be published. 5. G. Succop, R.T. Bubsey, M.H. Jones and W.F. Brown, Jr., in "Developments in Fracture Mechanics Test Methods Standardization", ASTM, STP 632, 153 (1977). 6. L.M. Barker and F.I. Barata, J. Test. Eval., 8, 97 (1980). 7. L.M. Barker, Int. J. Frac., 15, 515 (1979). 8. F.C. Monkman and N.J. Grant, Proc. ASTM56, 593 (1956). 9. F. Dobes and K. Milicka, Met. Sci. 10, 382 (1976). 10. W.D. Nix, Scripta Met. 17, I, (1983). I I . A.S. Argon, Scripta Met. 17, 5, (1983). 12. W. Beere, Scripta Met. ]7, 13, (1983). 13. J.S. Wang, L. Martinez and W.D. Nix, Acta Metall. 31, 873 (1983).
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B E H A V I O R OF P R E C I P I T A T I O N - F R E E
IF
(a)
ZONE A L U M I N U M
1507
(b)
Load
F.~x I0
--T 40 45
0.1Frnox " / / / / / / / / / / , 1 1 ~ Mid-Point Deflection
FIG. 1. (a) Three point bending of a multiple-notched and precracked fracture toughness specimens with i t s dimensions in mm; (b) Typical load vs. mid-point d e f l e c t i o n curve of the specimen in (a).
n2 ~, ~
~--,~- ~
10
19 I
o I,..,
L..
~._
L,
OIL
~ 15 0
•---,v I 0 50
I I00
I 150
4 200
n
FIG. 2. The dependence of f r a c t u r e energy (W/A) and percentage of i n t e r g r a n u l a r f r a c t u r e (f~GF)~ on creep time.
Creep Time (h)
21
E
I
I
I
I
I
I
19 FIG. 3. The dependence of f r a c t u r e energy on percentage of intergranular fracture.
17 E
w 15 0
13 4
I 5 Percent
I 6
I 7
I 8
I 9
I I0
of I n t e r g r a n u l a r F r a c t u r e
1508
B E H A V I O R OF P R E C I P I T A T I O N - F R E E
ZONE A L U M I N U M
FIG. 4. Fracture surface of bend specimen in the v i c i n i t y of grain boundary facets.
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