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The effects of grain boundary carbide on the low temperature tensile properties of type 316 stainless steel
Materials Science and EngineeringA, 102 (1988) 193-199
193
The Effects of Grain Boundary Carbide on the Low Temperature Tensile Properties of Type 316 Stainless Steel S. S. WU* and D. GAN Institute of Materials Science and Engineering. National Sun Yat-Sen University, Kaohs'iung(Taiwan) (Received October 13, 1987; in revised form Februar~ 4. 19881
Abstract
The effects of grain boundary M:3C~ prectpitates on the low temperature ( - 196 °C) tensile properties are systematically studied. For specimens aged at 650 °C for 240 h and at 750 °C for 100 h, the low temperature tensile strength, elongation and reduction in area are reduced severely and the specimens fail by brittle intergranular fracture. This is probably due to the dense grain boundary carbide and also to the grain boundary martensite in the chromium-depleted zone. The amounts" of reduction in tensile strength, elongation and reduction of area increase with increasing grain size for specimens aged at 650 °C for 240 h. The fracture strength is proportional to the inverse of the square root of grain size, and a critical stress intensity factor is determined. 1.
Introduction
Type 316 stainless steel has found wide applications for high temperature and low temperature uses owing to its good mechanical properties and corrosion resistance. Recent studies [1, 2] showed that significant reduction in room temperature impact energy resulted from the grain boundary M23C~, precipitates. The grain boundary carbide was also shown to reduce the room temperature tensile elongation moderately but to have little effect on the yield stress and tensile strength. Fine dense grain boundary carbide was most damaging and the embrittling mechanisms were mainly grain boundary cracks, grain boundary separation and shearing, and grain boundary sliding [1, 2]. Although there was no martensitic transformation in the room temperature tensile test in type 316 stainless steel [3-5], the low temperature ( - 196 °C) tensile test [5] showed a secondary stage of work
*Present address: Materials Research Lab., Industrial Technology Research Institute, Kaohsiung, Taiwan. (J921-5093/88/$3.50
hardening with marked increase in hardening rate, and X-ray diffraction confirmed that it was due to the martensitic transformation induced by plastic deformation. Martensitic transformation along the grain boundary was found to be induced in type 304 stainless steel by a low temperature ( - 196°C) in the chromium-depleted zone if the specimens were aged previously at 600-750°C to cause the precipitation of grain boundary carbide [6]. The low temperature impact tests of type 316 stainless steel after aging at 650, 750 and 900°C to form grain boundary carbide demonstrated that the impact energy at - 196 °C was severely reduced by up to about 90% by grain boundary carbide and grain boundary martensite [7]. However, the grain boundary martensite showed no further effect on the room temperature tensile properties but caused further severe reduction in the room temperature impact energy, in addition to the effect of the grain boundary carbide [7]. The effects of grain boundary carbide on the - 1 9 6 ° C tensile properties of type 316 stainless steel have not been studied. In this experiment, specimens of different grain sizes were prepared and aged to form grain boundary carbides and then tensile tested at - 1 9 6 ° C . Specimens with grain boundary carbides of different sizes and densities were also prepared by exposing them at 650, 75(I and 900°C for various times and tensile tested at - 196 °C to study the change in tensile properties. 2.
Procedures
The composition of the type 316 stainless steel used in this experiment is 16.99 wt.% Cr, 10.46 wt.% Ni, 2.36 wt.% Mo, 1.45 wt.% Mn, 0.45 wt.% Si, 0.06 wt.% C, 0.03 wt.% P and the balance iron. A group of specimens, as shown in Table 1, were solution treated at from 1000 to 1250°C with increasing time to develop different grain sizes. A prestrain of 30%, applied to specimens solution treated at © Elsevier Sequoia/Printed in The Netherlands
194 TABLE 1 The heat treatment and the grain size of specimens
Heat treatment
Grain size
Cold working 30%,1000°C, 10 min 1050°C, 40 min 1100°C, 5 h ll50°C, 10h 1250°C, 20 h
43 75 98 135 195
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6.255
10
i Fig. 1. Configuration(in millimetres)of tensile specimen. 3. Results
3.1. Microstructure
1000°C for 10 min to reduce the final grain size, was of no effect on the final tensile properties except the grain size as recrystallization should remove all the effect of cold working. This group of specimens was then aged at 650°C for 240 h to induce fine dense grain boundary carbide. A second group of specimens, solution treated at 1050 °C for 40 min with a grain size of 75/~m, were aged at 650°C for 8, 30, 60, 100, 240 and 480 h, at 750°C for 8, 30, 60 and 100 h, and at 900°C for 8, 30 and 100 h. The different aging times were to induce the carbide to precipitate gradually along the grain boundary, and the different temperatures were to develop grain boundary carbides of different sizes and densities. The aging treatment at each temperature was stopped before the precipitation of intragranular carbide, as determined by time-temperature-precipitation curve [8] and by metallographic examination [2]. The grain boundary carbide is of the M23C 6 type, as determined by electron diffraction in a transmission electron microscope and by a Debye-Scherrer pattern of extracted residuals
[1]. The heat-treated specimens were machined into round tensile specimens of diameter 6.25 mm and gauge length 25 mm, as shown in Fig. 1. The oxide scales developed in the solution and aging treatments were all removed during the machining process. Tensile tests were performed on an Instron 1125 universal testing machine with a low temperature jig dipped in a Dewar of liquid nitrogen. Specimens were tensile tested at a cross-head speed of 0.5 mm min- ~. Marks of length 25 mm were made on the specimens before testing, and the 0.2% offset yield stress, ultimate tensile stress, elongation and reduction in area were calculated after testing. From the load-displacement curves the true stress vs. true plastic strain curves were calculated. The grain size, grain boundary carbide and fracture surfaces were analysed by scarming electron microscopy (using a JEOL JSM-35CF instrument). X-ray diffraction (Diano 8536; M o K a radiation) was conducted on the polished surface of tensile-tested specimens.
The specimens of all grain sizes after solution treatments are free from carbide. For specimens of all grain sizes aged at 650°C, there was no grain boundary carbide after 8 h. The carbide is observed to precipitate and increase in density along the grain boundary after 60-100 h. After 240 h the carbide is fine and dense along the grain boundary (Fig. 2(a)) which shows no further observable change on further aging to 480 h. For specimens aged at 750°C, the carbide precipitates more rapidly. After 30 h the carbide has already precipitated along the grain boundary and after 60 h the carbide is nearly continuous along the grain boundary and increases slightly in density on further aging to 100 h (Fig. 2(b)). The grain boundary carbide particles are larger but less dense than at 650 °C. Aging at 900 °C causes the grain boundary carbide to precipitate rapidly. After only 8 h the carbide has already precipitated along the grain boundary. After 30 h the grain boundary carbide is large and sparsely distributed (Fig. 2(c)) which remains so with further aging to 100 h. For all specimens, there is little intragranular carbide. A detailed analysis was reported previously [1, 2]. 3.2. Effect o f grain size
For the specimens of different grain size and aged at 650 °C for 240 h, the low temperature tensile test results are shown in Figs. 3-5. Figure 3 shows that the yield stresses of the aged specimens are slightly larger than the solution-treated specimens, but the difference is small, only about 6%. Both groups of specimens obey the Hall-Petch equation which, from the least-squares fit, are, for the aged specimens, a y = 4 0 1 + 1.17d-~/2 and, for the solution-treated specimens, ay = 376 + 1.08d- 1/2 where ay and d are in megapascals and metres respectively.
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Fig. 2. Microstructures of specimens aged at (a) 650°C for 240 h, (b) 750 °C for 100 h and (c) 900 °C for 30 h. The tensile fracture properties of the specimens with and without grain boundary carbide are quite different. For the solution-treated specimens the fracture occurs in a normal way, i.e. with considerable necking after maximum load. However, for the aged specimens the fracture occurs abruptly without any necking. The tensile strengths of the solution-treated specimens and the aged specimens are shown in Fig. 4 in which the tensile strengths of the aged specimens are significantly reduced with respect to those of the solution-treated specimens. The amount of reduction increases with increasing grain size: about 15% for specimens of grain size 43 # m and 45% for specimens of grain size 195 /~m.
While the tensile strengths of the solution-treated specimens decrease slightly with increasing grain size, the fracture stresses of the aged specimens decrease significantly with increasing grain size. Figure 5 shows that the elongation and reduction in area of the aged specimens, similar to the tensile strength, also decrease significantly with respect to those of the solution-treated specimens. The amounts of reduction in elongation and reduction of area also increase with increasing grain size, reaching 79% and 85% respectively with respect to those of the solution-treated specimens of grain size 195 #m. For the solution-treated specimens, both the elongation and the reduction in area first increase and then level off with increasing grain size while, for the aged specimens, both decrease significantly with increasing grain size. The true stress vs. true plastic strain curves of the specimens of grain sizes 43 and 195 ,urn are shown in Fig. 6, and the curves of other specimens of intermediate grain sizes lie in between. The occurrence
196 of the second stage of hardening is apparent, and the work-hardening rate increases with decreasing grain size for all specimens. X-ray diffraction, shown in Fig. 7, confirmed that the second-stage hardening was accompanied by martensitic transformation similar to the reported results [5]. For the aged specimens the stress and the strain at which fracture occurs decrease significantly with increasing grain size and for the specimen of grain size 195 /~m the secondary hardening does not even begin at fracture. The work-hardening rates of the aged specimens are somewhat larger than those of the solution-treated specimens of corresponding grain size. In the aged specimens, the strain at which secondary hardening appears is less than that of the solution-treated specimens, as shown by the curves of the specimens of grain size 43/~m. However, the differences in the strain at which secondary harden-
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ing appears owing to different grain sizes in the solution-treated specimens are small and cannot be distinguished. For solution-treated specimens of all grain sizes the fracture is ductile with dimpled fracture surface, as shown in Fig. 8. However, after aging at 650°C for 240 h the fracture becomes brittle and intergranular with secondary cracks, as shown representatively in Figs. 9 and 10 for specimens of grain sizes 75 and 195/~m respectively. Secondary cracks along the grain boundary are also observed on the fracture surface. 3.3. Effect of aging temperature and time The yield stresses and tensile strengths of the specimens aged at 650, 750 and 900 °C for various times are shown in Fig. 11. The yield stresses generally show a small increase (up to about 6%) with the precipitation of grain boundary carbide. Aging at 900 °C has little effect on the tensile strengths and the specimens fail normally after necking. However, for specimens aged at 650 and 750 °C a significant
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Fig. 8. Ductile fracture of solution-treated specimen (grain size, 75/am).
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Fig. 9. Intergranular fracture of specimen aged at 650 °C for 240 h (grain size, 75 ¢tm).
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strength (upper three curves) for specimens aged at 900°C (o), 750°C ( A • ) and 650°C (rn m) as a function of aging time: x, solution-treated specimen; • , m, brittle fracture without necking.
Fig. 10. Intergranular fracture of specimen aged at 650 °C for 240 h (grain size, 195/am).
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decrease in the tensile strength gradually shows up with aging, reaching 28% and 21% respectively for specimens aged at 650°C for 480 h and at 750°C for 100 h. The fracture also changes to a brittle nature without necking, as indicated by the full symbols in Fig. 11. The changes in the elongation and the reduction in area with aging time and temperature are shown in Fig. 12. The tensile elongations show a small reduction up to 17% for the specimens aged at 900 °C and a severe reduction of up to 65% and 63% for specimens aged at 650°C for 480 h and at 750 °C for 100 h respectively. The trend of the change in the reduction of area is similar. For solution-treated specimens the fracture is ductile with dimpled fracture surface (Fig. 8). For specimens aged at 650°C the fracture surface gradually changes from a ductile transgranular fracture (8 and 30 h, with no grain boundary carbide) to a mixed mode (60 and 100 h, with the grain boundary carbide precipitating), as shown in Fig. 13, and to an intergranular fracture with secondary cracks (240 and 480 h, with stable fine and dense grain boundary carbide), as shown in Fig. 9. For the
Fig. 12. The elongation (EL) ( - - - ) and reduction of area (RA) ( ) of specimens aged at 900°C (o), 750 °C ( A, • ) and 650°C (n m) as a function of aging time: ×, solutiontreated specimen; • , m, brittle fracture without necking.
specimens aged at 750°C the fracture surface changes similarly, from a ductile transgranular fracture (8 h, with no grain boundary carbide) to a mixed mode (30 h, with grain boundary carbide precipitating), as shown in Fig. 14, and to a nearly complete intergranular fracture with secondary cracks (60 and 100 h, with a nearly continuous grain boundary carbide), as shown in Fig. 15. For the specimens aged at 900 °C with the rapid precipitation of large and sparsely distributed grain boundary carbide, the fracture is of a mixed mode with secondary cracks (Fig. 16). The appearance and evolution of the fracture surfaces are similar to those of the low temperature impact specimens except for the specimens aged at 900 °C which fail by intergranular fracture in impact tests i7].
198
Fig. 13. Mixed fracture of a specimen aged at 650°C for 100 h (grain size, 75/am).
Fig. 15. Intergranular fracture of a specimen aged at 750 °C for 60 h (grain size, 75/am).
Fig. 14. Mixed fracture of a specimen aged at 750 °C for 30 h (grain size, 75/am).
Fig. 16. Mixed fracture of a specimen aged at 900 °C for 30 h (grain size, 75/am).
4. Discussion
at 750°C, the separated grain boundaries also become coarser (Fig. 15). For these two reasons the grain boundary carbide is clearly a very important factor in the observed embrittlement of aged type 316 stainless steel in the low temperature tensile test. For the specimen aged at a high temperature to induce the grain boundary carbide, a layer of martensite network forms along the grain boundary in the chromium-depleted zone when cooled to - 196 °C [6, 7]. The grain boundary martensite has been shown to cause intergranular fracture and a reduction in impact energy in a room temperature impact test but to have little effect on the room temperature tensile properties [7]. It is possible that it can also contribute to the embrittlement and intergranular fracture in a - 1 9 6 ° C tensile test. The intragranular martensite, however, does not cause intergranular fracture, as observed in the solutiontreated specimens which fail by transgranular fracture (Fig. 8)[3]. In Fig. 6 the increase in the work-hardening rates of the aged specimens over the solution-treated
From the experimental results, it is clear that the tensile strength, elongation and reduction in area change with the precipitation of grain boundary carbide. For specimens aged at 650 and 750 °C, as the grain boundary carbide starts to precipitate, the tensile properties also start to deteriorate accordingly and the fracture mode starts to change to a mixed nature. After aging for 60 h at 750 °C and for 240 h at 650 °C when the grain boundary carbide is dense and stable, the reductions in tensile properties also level off and the fracture becomes intergranular with secondary cracks. Aging at 900°C causes rapid precipitation of large sparse grain boundary carbides and the fracture surface is of a mixed mode with secondary cracks. Furthermore, the appearances of the separated grain boundaries also change with the morphology of the grain boundary carbide. As the grain boundary carbide is fine and dense at 650 °C, the separated grain boundaries are smooth (Figs. 9 and 10). However, as the carbides are larger and less dense
199
specimens of the same grain size can be accounted for by the presence of grain boundary carbide and grain boundary martensite in the aged specimens. The grain boundary carbide alone can increase the work-hardening rate, as has been observed in the room temperature tensile test specimens in which no grain boundary martensite was formed [2]. Secondary hardening, which occurs owing to the formation of intragranular martensite [3], starts to appear at smaller strain for the aged specimens than the solution-treated specimens of the same grain size (Fig. 6). This is probably due to the larger work-hardening rate in the aged specimens so that a larger stress is reached earlier. The decrease in grain size is shown to increase the yield stress according to the Hall-Petch equation and also to increase the work-hardening rate which could be accounted for by the increased grain surface and the increased multiple-slip deformation zone [9]. For the specimens aged at 650 °C for 240 h, fracture occurs after considerable elongation and at a stress larger than the yield stress (Fig. 6). However, it is a brittle intergranular fracture without necking. The stress and the strain at which brittle fracture occurs decrease with increasing grain size (Figs. 4 and 5). Examination of the profile of failed specimens near the fracture surfaces shows that there is no intergranular crack below the fracture surface or on the specimen surface. It appears that, as long as one crack is formed along a grain boundary, brittle frailure follows and there is no need for several intergranular cracks to join together to cause brittle fracture. The fracture does not have to start from the specimen surface. If the true stress a t at fracture is plotted vs. the inverse of the square root of grain size, as shown in Fig. 17, a straight line can be fitted. If we consider the crack to be penny shaped and its size to be equal to the grain size, a stress intensity factor of o1{2d/:~) 1/2 is obtained which is approximately a constant within the tested range of grain sizes. The value, which can be taken as the critical stress intensity factor, is 8.9 MPa m u2 and the deviation is _+0.6 MPa m 1/2. For specimens of grain size 75/~m aged at 650 °C for 480 h and at 750 °C for 100 h, the values are 8.7 and 9.5 respectively; these are still in agreement with the above value.
5. Conclusions (1) For specimens of grain size 75 # m aged at 650 °C for 240 h and at 750°C for 100 h, the tensile strengths are reduced by 27% and 21% respectively and the elongations are reduced by 65% and 63%
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respectively with respect to those of the solutiontreated specimens. The specimens fail by brittle intergranular fracture. (2) The grain boundary carbide and the grain boundary martensite are probably the main factors resulting in the observed embrittlement. (3) The reductions in low temperature tensile strength and elongation for specimens aged at 650 °C for 240 h increase with increasing grain size, reaching 45% and 79% respectively with respect to those of the solution-treated specimens of grain size 195/~m. (4) The low temperature tensile fracture stress of the specimens aged at 650°C for 240 h is proportional to the inverse of the square root of the grain size, and a critical stress intensity factor of 8.9 MPa m 1/z is determined.
Acknowledgment This work was financially supported by the National Science Council, Taiwan.
References l D. Gan, Metall. Trans. A, 13 (1982) 2155. 2 S.Y. Chen and D. Gan, Mater. Sci. Eng., 84 (1986) 65. 3 S. L. Mannan, K. G. Samuel and P. Rodriguez, Ser. Metall., 16 (1982)255. 4 Y. Rosenthal, M. Mark-Markowistch, A. Stern and P. Eliezer, Mater. Sci. Eng., 67(1984) 91. 5 V. Seetharaman and R. Krishnan, J. Mater. Sei., 16 ( 1981 ) 523. 6 E.P. Bulter and M. G. Burke, Aeta Metall.. 34 (1986) 557. 7 T. Tseng, D. Gan and P. W. Kao, Mater. Sci. Eng., 95 (1987)73. 8 B. Weiss and R. Stickler, Metall. Trans. A, 3 (1972) 851. 9 M. F. Ashby, Philos. Mag., 2l (1969) 399.
193
The Effects of Grain Boundary Carbide on the Low Temperature Tensile Properties of Type 316 Stainless Steel S. S. WU* and D. GAN Institute of Materials Science and Engineering. National Sun Yat-Sen University, Kaohs'iung(Taiwan) (Received October 13, 1987; in revised form Februar~ 4. 19881
Abstract
The effects of grain boundary M:3C~ prectpitates on the low temperature ( - 196 °C) tensile properties are systematically studied. For specimens aged at 650 °C for 240 h and at 750 °C for 100 h, the low temperature tensile strength, elongation and reduction in area are reduced severely and the specimens fail by brittle intergranular fracture. This is probably due to the dense grain boundary carbide and also to the grain boundary martensite in the chromium-depleted zone. The amounts" of reduction in tensile strength, elongation and reduction of area increase with increasing grain size for specimens aged at 650 °C for 240 h. The fracture strength is proportional to the inverse of the square root of grain size, and a critical stress intensity factor is determined. 1.
Introduction
Type 316 stainless steel has found wide applications for high temperature and low temperature uses owing to its good mechanical properties and corrosion resistance. Recent studies [1, 2] showed that significant reduction in room temperature impact energy resulted from the grain boundary M23C~, precipitates. The grain boundary carbide was also shown to reduce the room temperature tensile elongation moderately but to have little effect on the yield stress and tensile strength. Fine dense grain boundary carbide was most damaging and the embrittling mechanisms were mainly grain boundary cracks, grain boundary separation and shearing, and grain boundary sliding [1, 2]. Although there was no martensitic transformation in the room temperature tensile test in type 316 stainless steel [3-5], the low temperature ( - 196 °C) tensile test [5] showed a secondary stage of work
*Present address: Materials Research Lab., Industrial Technology Research Institute, Kaohsiung, Taiwan. (J921-5093/88/$3.50
hardening with marked increase in hardening rate, and X-ray diffraction confirmed that it was due to the martensitic transformation induced by plastic deformation. Martensitic transformation along the grain boundary was found to be induced in type 304 stainless steel by a low temperature ( - 196°C) in the chromium-depleted zone if the specimens were aged previously at 600-750°C to cause the precipitation of grain boundary carbide [6]. The low temperature impact tests of type 316 stainless steel after aging at 650, 750 and 900°C to form grain boundary carbide demonstrated that the impact energy at - 196 °C was severely reduced by up to about 90% by grain boundary carbide and grain boundary martensite [7]. However, the grain boundary martensite showed no further effect on the room temperature tensile properties but caused further severe reduction in the room temperature impact energy, in addition to the effect of the grain boundary carbide [7]. The effects of grain boundary carbide on the - 1 9 6 ° C tensile properties of type 316 stainless steel have not been studied. In this experiment, specimens of different grain sizes were prepared and aged to form grain boundary carbides and then tensile tested at - 1 9 6 ° C . Specimens with grain boundary carbides of different sizes and densities were also prepared by exposing them at 650, 75(I and 900°C for various times and tensile tested at - 196 °C to study the change in tensile properties. 2.
Procedures
The composition of the type 316 stainless steel used in this experiment is 16.99 wt.% Cr, 10.46 wt.% Ni, 2.36 wt.% Mo, 1.45 wt.% Mn, 0.45 wt.% Si, 0.06 wt.% C, 0.03 wt.% P and the balance iron. A group of specimens, as shown in Table 1, were solution treated at from 1000 to 1250°C with increasing time to develop different grain sizes. A prestrain of 30%, applied to specimens solution treated at © Elsevier Sequoia/Printed in The Netherlands
194 TABLE 1 The heat treatment and the grain size of specimens
Heat treatment
Grain size
Cold working 30%,1000°C, 10 min 1050°C, 40 min 1100°C, 5 h ll50°C, 10h 1250°C, 20 h
43 75 98 135 195
14- 20 --~l <
32
~l ~-- 2 0 "-~l
6.255
10
i Fig. 1. Configuration(in millimetres)of tensile specimen. 3. Results
3.1. Microstructure
1000°C for 10 min to reduce the final grain size, was of no effect on the final tensile properties except the grain size as recrystallization should remove all the effect of cold working. This group of specimens was then aged at 650°C for 240 h to induce fine dense grain boundary carbide. A second group of specimens, solution treated at 1050 °C for 40 min with a grain size of 75/~m, were aged at 650°C for 8, 30, 60, 100, 240 and 480 h, at 750°C for 8, 30, 60 and 100 h, and at 900°C for 8, 30 and 100 h. The different aging times were to induce the carbide to precipitate gradually along the grain boundary, and the different temperatures were to develop grain boundary carbides of different sizes and densities. The aging treatment at each temperature was stopped before the precipitation of intragranular carbide, as determined by time-temperature-precipitation curve [8] and by metallographic examination [2]. The grain boundary carbide is of the M23C 6 type, as determined by electron diffraction in a transmission electron microscope and by a Debye-Scherrer pattern of extracted residuals
[1]. The heat-treated specimens were machined into round tensile specimens of diameter 6.25 mm and gauge length 25 mm, as shown in Fig. 1. The oxide scales developed in the solution and aging treatments were all removed during the machining process. Tensile tests were performed on an Instron 1125 universal testing machine with a low temperature jig dipped in a Dewar of liquid nitrogen. Specimens were tensile tested at a cross-head speed of 0.5 mm min- ~. Marks of length 25 mm were made on the specimens before testing, and the 0.2% offset yield stress, ultimate tensile stress, elongation and reduction in area were calculated after testing. From the load-displacement curves the true stress vs. true plastic strain curves were calculated. The grain size, grain boundary carbide and fracture surfaces were analysed by scarming electron microscopy (using a JEOL JSM-35CF instrument). X-ray diffraction (Diano 8536; M o K a radiation) was conducted on the polished surface of tensile-tested specimens.
The specimens of all grain sizes after solution treatments are free from carbide. For specimens of all grain sizes aged at 650°C, there was no grain boundary carbide after 8 h. The carbide is observed to precipitate and increase in density along the grain boundary after 60-100 h. After 240 h the carbide is fine and dense along the grain boundary (Fig. 2(a)) which shows no further observable change on further aging to 480 h. For specimens aged at 750°C, the carbide precipitates more rapidly. After 30 h the carbide has already precipitated along the grain boundary and after 60 h the carbide is nearly continuous along the grain boundary and increases slightly in density on further aging to 100 h (Fig. 2(b)). The grain boundary carbide particles are larger but less dense than at 650 °C. Aging at 900 °C causes the grain boundary carbide to precipitate rapidly. After only 8 h the carbide has already precipitated along the grain boundary. After 30 h the grain boundary carbide is large and sparsely distributed (Fig. 2(c)) which remains so with further aging to 100 h. For all specimens, there is little intragranular carbide. A detailed analysis was reported previously [1, 2]. 3.2. Effect o f grain size
For the specimens of different grain size and aged at 650 °C for 240 h, the low temperature tensile test results are shown in Figs. 3-5. Figure 3 shows that the yield stresses of the aged specimens are slightly larger than the solution-treated specimens, but the difference is small, only about 6%. Both groups of specimens obey the Hall-Petch equation which, from the least-squares fit, are, for the aged specimens, a y = 4 0 1 + 1.17d-~/2 and, for the solution-treated specimens, ay = 376 + 1.08d- 1/2 where ay and d are in megapascals and metres respectively.
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Fig. 4. The tensile strength of solution-treated specimens (A) and the fracture stress of the specimens aged at 650°C for 240 h (•) as a function of the grain size.
Fig. 2. Microstructures of specimens aged at (a) 650°C for 240 h, (b) 750 °C for 100 h and (c) 900 °C for 30 h. The tensile fracture properties of the specimens with and without grain boundary carbide are quite different. For the solution-treated specimens the fracture occurs in a normal way, i.e. with considerable necking after maximum load. However, for the aged specimens the fracture occurs abruptly without any necking. The tensile strengths of the solution-treated specimens and the aged specimens are shown in Fig. 4 in which the tensile strengths of the aged specimens are significantly reduced with respect to those of the solution-treated specimens. The amount of reduction increases with increasing grain size: about 15% for specimens of grain size 43 # m and 45% for specimens of grain size 195 /~m.
While the tensile strengths of the solution-treated specimens decrease slightly with increasing grain size, the fracture stresses of the aged specimens decrease significantly with increasing grain size. Figure 5 shows that the elongation and reduction in area of the aged specimens, similar to the tensile strength, also decrease significantly with respect to those of the solution-treated specimens. The amounts of reduction in elongation and reduction of area also increase with increasing grain size, reaching 79% and 85% respectively with respect to those of the solution-treated specimens of grain size 195 #m. For the solution-treated specimens, both the elongation and the reduction in area first increase and then level off with increasing grain size while, for the aged specimens, both decrease significantly with increasing grain size. The true stress vs. true plastic strain curves of the specimens of grain sizes 43 and 195 ,urn are shown in Fig. 6, and the curves of other specimens of intermediate grain sizes lie in between. The occurrence
196 of the second stage of hardening is apparent, and the work-hardening rate increases with decreasing grain size for all specimens. X-ray diffraction, shown in Fig. 7, confirmed that the second-stage hardening was accompanied by martensitic transformation similar to the reported results [5]. For the aged specimens the stress and the strain at which fracture occurs decrease significantly with increasing grain size and for the specimen of grain size 195 /~m the secondary hardening does not even begin at fracture. The work-hardening rates of the aged specimens are somewhat larger than those of the solution-treated specimens of corresponding grain size. In the aged specimens, the strain at which secondary hardening appears is less than that of the solution-treated specimens, as shown by the curves of the specimens of grain size 43/~m. However, the differences in the strain at which secondary harden-
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Fig. 8. Ductile fracture of solution-treated specimen (grain size, 75/am).
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decrease in the tensile strength gradually shows up with aging, reaching 28% and 21% respectively for specimens aged at 650°C for 480 h and at 750°C for 100 h. The fracture also changes to a brittle nature without necking, as indicated by the full symbols in Fig. 11. The changes in the elongation and the reduction in area with aging time and temperature are shown in Fig. 12. The tensile elongations show a small reduction up to 17% for the specimens aged at 900 °C and a severe reduction of up to 65% and 63% for specimens aged at 650°C for 480 h and at 750 °C for 100 h respectively. The trend of the change in the reduction of area is similar. For solution-treated specimens the fracture is ductile with dimpled fracture surface (Fig. 8). For specimens aged at 650°C the fracture surface gradually changes from a ductile transgranular fracture (8 and 30 h, with no grain boundary carbide) to a mixed mode (60 and 100 h, with the grain boundary carbide precipitating), as shown in Fig. 13, and to an intergranular fracture with secondary cracks (240 and 480 h, with stable fine and dense grain boundary carbide), as shown in Fig. 9. For the
Fig. 12. The elongation (EL) ( - - - ) and reduction of area (RA) ( ) of specimens aged at 900°C (o), 750 °C ( A, • ) and 650°C (n m) as a function of aging time: ×, solutiontreated specimen; • , m, brittle fracture without necking.
specimens aged at 750°C the fracture surface changes similarly, from a ductile transgranular fracture (8 h, with no grain boundary carbide) to a mixed mode (30 h, with grain boundary carbide precipitating), as shown in Fig. 14, and to a nearly complete intergranular fracture with secondary cracks (60 and 100 h, with a nearly continuous grain boundary carbide), as shown in Fig. 15. For the specimens aged at 900 °C with the rapid precipitation of large and sparsely distributed grain boundary carbide, the fracture is of a mixed mode with secondary cracks (Fig. 16). The appearance and evolution of the fracture surfaces are similar to those of the low temperature impact specimens except for the specimens aged at 900 °C which fail by intergranular fracture in impact tests i7].
198
Fig. 13. Mixed fracture of a specimen aged at 650°C for 100 h (grain size, 75/am).
Fig. 15. Intergranular fracture of a specimen aged at 750 °C for 60 h (grain size, 75/am).
Fig. 14. Mixed fracture of a specimen aged at 750 °C for 30 h (grain size, 75/am).
Fig. 16. Mixed fracture of a specimen aged at 900 °C for 30 h (grain size, 75/am).
4. Discussion
at 750°C, the separated grain boundaries also become coarser (Fig. 15). For these two reasons the grain boundary carbide is clearly a very important factor in the observed embrittlement of aged type 316 stainless steel in the low temperature tensile test. For the specimen aged at a high temperature to induce the grain boundary carbide, a layer of martensite network forms along the grain boundary in the chromium-depleted zone when cooled to - 196 °C [6, 7]. The grain boundary martensite has been shown to cause intergranular fracture and a reduction in impact energy in a room temperature impact test but to have little effect on the room temperature tensile properties [7]. It is possible that it can also contribute to the embrittlement and intergranular fracture in a - 1 9 6 ° C tensile test. The intragranular martensite, however, does not cause intergranular fracture, as observed in the solutiontreated specimens which fail by transgranular fracture (Fig. 8)[3]. In Fig. 6 the increase in the work-hardening rates of the aged specimens over the solution-treated
From the experimental results, it is clear that the tensile strength, elongation and reduction in area change with the precipitation of grain boundary carbide. For specimens aged at 650 and 750 °C, as the grain boundary carbide starts to precipitate, the tensile properties also start to deteriorate accordingly and the fracture mode starts to change to a mixed nature. After aging for 60 h at 750 °C and for 240 h at 650 °C when the grain boundary carbide is dense and stable, the reductions in tensile properties also level off and the fracture becomes intergranular with secondary cracks. Aging at 900°C causes rapid precipitation of large sparse grain boundary carbides and the fracture surface is of a mixed mode with secondary cracks. Furthermore, the appearances of the separated grain boundaries also change with the morphology of the grain boundary carbide. As the grain boundary carbide is fine and dense at 650 °C, the separated grain boundaries are smooth (Figs. 9 and 10). However, as the carbides are larger and less dense
199
specimens of the same grain size can be accounted for by the presence of grain boundary carbide and grain boundary martensite in the aged specimens. The grain boundary carbide alone can increase the work-hardening rate, as has been observed in the room temperature tensile test specimens in which no grain boundary martensite was formed [2]. Secondary hardening, which occurs owing to the formation of intragranular martensite [3], starts to appear at smaller strain for the aged specimens than the solution-treated specimens of the same grain size (Fig. 6). This is probably due to the larger work-hardening rate in the aged specimens so that a larger stress is reached earlier. The decrease in grain size is shown to increase the yield stress according to the Hall-Petch equation and also to increase the work-hardening rate which could be accounted for by the increased grain surface and the increased multiple-slip deformation zone [9]. For the specimens aged at 650 °C for 240 h, fracture occurs after considerable elongation and at a stress larger than the yield stress (Fig. 6). However, it is a brittle intergranular fracture without necking. The stress and the strain at which brittle fracture occurs decrease with increasing grain size (Figs. 4 and 5). Examination of the profile of failed specimens near the fracture surfaces shows that there is no intergranular crack below the fracture surface or on the specimen surface. It appears that, as long as one crack is formed along a grain boundary, brittle frailure follows and there is no need for several intergranular cracks to join together to cause brittle fracture. The fracture does not have to start from the specimen surface. If the true stress a t at fracture is plotted vs. the inverse of the square root of grain size, as shown in Fig. 17, a straight line can be fitted. If we consider the crack to be penny shaped and its size to be equal to the grain size, a stress intensity factor of o1{2d/:~) 1/2 is obtained which is approximately a constant within the tested range of grain sizes. The value, which can be taken as the critical stress intensity factor, is 8.9 MPa m u2 and the deviation is _+0.6 MPa m 1/2. For specimens of grain size 75/~m aged at 650 °C for 480 h and at 750 °C for 100 h, the values are 8.7 and 9.5 respectively; these are still in agreement with the above value.
5. Conclusions (1) For specimens of grain size 75 # m aged at 650 °C for 240 h and at 750°C for 100 h, the tensile strengths are reduced by 27% and 21% respectively and the elongations are reduced by 65% and 63%
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respectively with respect to those of the solutiontreated specimens. The specimens fail by brittle intergranular fracture. (2) The grain boundary carbide and the grain boundary martensite are probably the main factors resulting in the observed embrittlement. (3) The reductions in low temperature tensile strength and elongation for specimens aged at 650 °C for 240 h increase with increasing grain size, reaching 45% and 79% respectively with respect to those of the solution-treated specimens of grain size 195/~m. (4) The low temperature tensile fracture stress of the specimens aged at 650°C for 240 h is proportional to the inverse of the square root of the grain size, and a critical stress intensity factor of 8.9 MPa m 1/z is determined.
Acknowledgment This work was financially supported by the National Science Council, Taiwan.
References l D. Gan, Metall. Trans. A, 13 (1982) 2155. 2 S.Y. Chen and D. Gan, Mater. Sci. Eng., 84 (1986) 65. 3 S. L. Mannan, K. G. Samuel and P. Rodriguez, Ser. Metall., 16 (1982)255. 4 Y. Rosenthal, M. Mark-Markowistch, A. Stern and P. Eliezer, Mater. Sci. Eng., 67(1984) 91. 5 V. Seetharaman and R. Krishnan, J. Mater. Sei., 16 ( 1981 ) 523. 6 E.P. Bulter and M. G. Burke, Aeta Metall.. 34 (1986) 557. 7 T. Tseng, D. Gan and P. W. Kao, Mater. Sci. Eng., 95 (1987)73. 8 B. Weiss and R. Stickler, Metall. Trans. A, 3 (1972) 851. 9 M. F. Ashby, Philos. Mag., 2l (1969) 399.