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Fracture behaviour in secondary hardening embrittlement of isothermally peak-aged 6W3Ni steel
Materials Science and Engineering, A 147 ( 1991 ) 175-180
175
Fracture behaviour in secondary hardening embrittlement of isothermally peak-aged 6W-3Ni steel Kon Bae Lee, Chang Ki Choi and Hoon Kwon Department of Metallurgical Engineering, Kookmin University, Seou1136- 702 (Korea)
Kwang-Hoon Kim Department of Surface Finishing and Engineering, Tae-Hun Technical College, lncheon 401- 714 (Korea) (Received April 10, 1991 ; in revised form May 17, 1991 )
Abstract Secondary hardening embrittlement (SHE) in a 6W-3Ni steel isothermally peak aged in the range 550-650 °C was investigated by conducting impact testing in the range 20-250 °C. In the 600 and 650°C peak-aged conditions the impact toughness underwent a brittle-to-ductile transition with increasing test temperature, although the ductile-brittle transition temperature in the latter condition was higher by 200 °C than that in the former condition. In the 550 °C peak-aged condition, however, the impact toughness was not recovered and remained at the lower shelf energy level, about 5 J, even at 250 °C. The fracture mode was intergranular in the 550 °C peak-aged condition whereas it comprised mostly transgranular dimples in the 600 °C peak-aged condition. In the 650 °C peak-aged condition, where some intergranular area was observed, the area of transgranular cleavage changed to transgranular dimples at temperatures above 200 °C. The fracture behaviour has been discussed in terms of the impurity segregation associated with W2C precipitation, the presence of coarse cementite and the hardening behaviour. SHE is caused by: the easy occurrence of intergranular fracture due to impurity segregation in the 550 °C peak-aged condition; a decrease in ductile fracture energy due to the presence of coarse cementite and intrinsic hardening in the 600 °C peak-aged condition; and the combined effect of all three factors in the 650 °C peak-aged condition.
1. Introduction Secondary hardening steels have been used as heat-resistant alloys because of the high hardness retained at elevated temperatures. Recently, secondary hardening steels such as the AF1410 systern have been used for ultrahigh strength structures owing to their very high toughness [ 1 ] . After tempering in the secondary hardening ternperature range, however, embrittlement can occur. This embrittlement is referred to as secondary hardening embrittlement (SHE). According to its fracture mode, SHE can be classified as either intergranular or transgranular, It has been suggested that intergranular SHE is caused by impurity segregation [2, 3] and that transgranular SHE is associated with coarse boundary carbides [4-6]. Kwon [7] and Kim and Kwon [8] reported recently that intergranular 0921-5093/91/$3.50
SHE occurred in the overaged condition compared to transgranular SHE in the underaged condition for molybdenum steel and W-Ni steel. Hence the fracture behaviour in SHE can be greatly affected by the aging condition. The intrinsic secondary hardening due to precipitation of fine alloy carbides [9-11] also contributes to SHE. Since the precipitation of fine alloy carbides (i.e. Mo2C or W2C ) in molybdenum and tungsten steels leads to secondary hardening, the degree of secondary hardening is affected by the precipitation kinetics of alloy carbides [11]. Recently, Kwon [12] reported that a difference in degree of secondary hardening for molybdenum and tungsten steels has an important effect on SHE. In addition, Kim and Kwon [13] suggested that the different behaviour of SHE in the various iso© Elsevier Sequoia/Printed in The Netherlands
176
thermally peak-aged (i.e. aged to peak hardness) conditions for 6W-6Ni steel is associated with the differences in carbide morphology and impurity segregation. In 6W-6Ni and 6W-3Ni steels isochronally aged they [8] also investigated the differences in carbide morphology and fracture behaviour, particularly in the overaged condition where both steels had low impact toughness even at a high temperature of 250 °C. Hence there may be a variation in carbide morphology in isothermally peak-aged conditions for 6W-3Ni steel and the fracture behaviour in SHE may be different from that for 6W-6Ni steel. In this study SHE in a 6W-3Ni steel isothermally peak aged in the range 550-650 °C was analysed in terms of variations in impact toughness and fracture behaviour with carbide morphology and impurity segregation.
700 600
_~ ~ ~ 50o ~ 400
300
• ~ 1'0 TEMPERinGT~ME~h)
100
Fig. 1. Isothermal hardness variations at 550, 600 and 650 °C.
25 20
2. Experimental procedure
~ t5
An ingot of mass about 20 kg was prepared from high purity materials by vacuum induction melting. The ingot was hot rolled to 14 mm plate. Standard Charpy V-notch impact specimens were machined from the plate. The chemical composition of an experimental steel was 0.37C-6.18W3.18Ni-0.006P-0.004S (weight per cent). Specimens were austenitized under a flowing argon atmosphere at 1200°C for 1 h and oil quenched. They were then isothermally aged in the secondary hardening range 550-650 °C and water quenched, The micro-Vickers hardness was measured to determine the peak-aged conditions at 550, 600 and 650 °C. The peak-aging times were 20, 1 and 0.4 h at 550, 600 and 650 °C respectively. For the peak-aged conditions Charpy impact testing was performed in the range 20-250 °C. Microstructures were investigated using a Jeol 200CX transmission electron microscope operated at 160 kV. To investigate the fracture initiation area, which may be a primary factor in impact properties of relatively low toughness materials, all fractographs were taken from areas near the notch.
~
3. Experimental results Isothermal hardness variations in the range 550-650°C are shown in Fig. 1. The hardness initially decreased from the as-quenched value of 570 HV to values of 4 0 5 - 4 4 0 HV and then increased to the peak values of 4 3 5 - 4 9 0 H V
0. I
~
10
~
•
"
J
• 5 -
• '
100
'
200
300
TESTTEMPERATURE( OC) Fig. 2. Impact toughness variations in the 550, 600 and 650°Cpeak-ag edconditions.
before final softening. As the aging temperature increased, the time to peak hardness became progressively shorter and the peak hardness decreased. The peak hardness values were 490, 460 and 435 HV at 550, 600 and 650 °C respectively. This is very common for precipitation-hardening alloys. The impact toughness changes with test ternperature in the peak-aged condition are shown in Fig. 2. In the 600 and 650 °C peak-aged conditions the impact toughness underwent a brittleto-ductile transition with increasing temperature. If the ductile-brittle transition temperature (DBTT) is assumed to be a temperature at which the impact toughness has a mid-value (12.5 J) between the upper shelf energy (20 J) and the lower shelf energy (5 J), the DBTTs are about 0 and 210°C in the 600 and 650°C peak-aged conditions respectively. The DBTT in the 650 °C condition was shifted to a very high temperature compared to the 600°C condition, though the
177
hardness (435 HV) in the 650 °C condition was lower than that (460 HV) in the 600 °C condition. The 650 °C peak-aged condition results in more severe embrittlement in spite of its lower hardness. Hence SHE is not simply correlated with the hardness variation. On the other hand, the impact toughness in the 550°C peak-aged condition remained at a low level of 5 J even at 250°C. The fracture surfaces are shown in Figs. 3-5. The 600 °C peak-aged specimen exhibits ductile dimples containing a small amount of intergranular fracture at room temperature (Fig. 3). The 650 °C peak-aged specimen shows a mixture of intergranular fracture and transgranular quasicleavage at room temperature (Fig. 4(a)). The transgranular quasi-cleavage had a transition to dimples on approaching 250 °C near the upper shelf energy level, but some intergranular area remained (Fig. 4(b)). In the 550°C peak-aged condition, however, large amounts of intergranular fracture were observed at both room temperature and 250°C (Fig. 5). Thus no increase of impact toughness occurs even at 250 °C because of the easy occurrence of intergranular fracture. Transmission electron micrographs are shown in Figs. 6 and 7. During aging, the cementite initially precipitates and later dissolves into the matrix and M2C precipitates primarily at dislocations in the matrix. Alternatively, the cementite disintegrates by the in situ formation of M2C and finally M23C6 and/or M6C precipitate as the M2C dissolves [ 11 ]. In the 550 °C peak-aged condition
Fig. 3. Fractograph of the 600 °C peak-aged specimen tested at 20 °C.
extensive formation of very fine W2C carbides occurred within the lath as the long-time aging at a relatively low temperature led to a gradual dissolution of a large amount of cementite, as shown in Fig. 6(a), In the 600 °C peak-aged condition fine W2C particles were observed within laths whereas the coarse plate-like cementite persisted predominantly at the lath boundaries (Fig. 6(b)). Even in the 650°C overaged condition some coarse cementite still remained, although relatively fine W2C carbides were observed (Fig. 7). Hence it is certain that much more coarse cementite is retained in the 650 °C peak-aged condition relative to the other two conditions. 4. Discussion Before we discuss the peak-aged conditions, the results for the isothermally underaged and
Fig. 4. Fractographs of the 650 oC peak-aged specimen tested at (a) 20 and (b) 250 °C.
178
Fig. 6. Transmission electron micrographs of (a) 550 and (b) 600 °C peak-aged specimens.
Fig. 5. Fractographs of the 550°C peak-aged specimen tested at (a) 20 and (b) 250 °C.
overaged conditions in the range 550-650 °C are summarized as follows [14]. In the underaged conditions the impact toughness had high values above those at the DBTT and approached the upper shelf energy levels, 35-42 J. In the overaged conditions, however, the impact toughness stayed at the lower shelf energy levels, 2-5 J, even at 250°C. The fracture surfaces showed mostly transgranular dimples in the underaged condition whereas intergranular fracture was most often observed in the overaged condition. Since tungsten has quite similar properties to molybdenum, we postulate that tungsten in the matrix can tie up phosphorus, presumably by lowering its activity or by the formation of tungsten phosphide, resulting in inhibition of phosphorus segregation to grain boundaries in a similar way as in molybdenum steel [7, 15].
I
Fig. 7. Transmission electron micrograph of the specimen overaged at 650°C for 1 h.
There may be no available tungsten in the matrix to tie up phosphorus in the 550°C peak-aged condition because the cementite was mostly replaced by W2C. Hence impurity segregation to
179
grain boundaries can occur in that condition. In the 600 °C peak-aged condition, however, a relatively large amount of coarse cementite plates remained, in particular at the lath boundaries, The fact that the lath boundary cementite plates persist means that there may still remain available tungsten to tie up phosphorus. In the 650°C peak-aged condition, since a relatively small amount of coarse cementite remained and the amount of tungsten in the matrix may be insufficient to tie up all the phosphorus, residual impurifles beyond a certain amount to be tied up by the available tungsten can segregate to the grain boundaries, lntergranular fracture can be caused by impurities and/or carbides at grain boundaries. The impurities directly affect the grain boundary strength whereas the carbides act as stress concentrators. Thus impurity segregation plays a primary role in intergranular fracture in high strength alloy steels, Denser precipitation of carbides leads to higher stress concentrations. In addition, the difference in hardness (i.e. difference in extent of secondary hardening) affects the stress concentration susceptibility at grain boundaries. Stress relaxation in a higher hardness microstructure is more difficult because of the greater difficulty of plastic deformation. Hence in the 550°C peakaged condition of higher hardness and denser precipitation of carbides, brittle intergranular fracture can occur. In that condition, however, the easy occurrence of intergranular fracture fundamentally implies that impurity segregation to grain boundaries took place. In contrast, in the 650 °C peak-aged condition of lower hardness with less dense precipitation than in the 550 °C peak-aged condition, the stress concentration susceptibility may be lower. Furthermore, because relatively small amounts of impurities may segregate to grain boundaries, the occurrence of intergranular fracture is resisted, Thus the amount of intergranular fracture is small and the impact toughness undergoes a DBTT. However, the fact that the impact toughness approaches the upper shelf energy level only at temperatures above 250 °C in spite of the lower hardness indicates that some impurity segregation contributes significantly to the shift of DBTT to a very high temperature. This may be compared to the DBTT of the 600 °C peak-aged condition where there may be sufficient tungsten available to tie up the phosphorus. In contrast, for
the 6W-6Ni steel isothermally peak aged at 650 °C, since the available tungsten was consumed as a result of the near absence of coarse cementite and the extensive precipitation ofW2C, the impurity segregation led to the easy occurrence of intergranular fracture [ 13]. In the 600°C peak-aged condition the fracture surface consisted mostly of dimples with some intergranular area even at room temperature. Such intergranular areas were also observed in the 200 °C tempered condition in which almost all segregation of impurities occurred during austenitizing rather than during tempering [16]. In the isothermally underaged conditions in the range 550-650 °C the upper shelf energy levels were 35-42 J [14]. Thus a decrease in ductile fracture energy, which is caused by the presence of coarse cementite and the intrinsic hardening due to fine W2C precipitation, results in a large drop in upper shelf energy. A more detailed discussion was presented in previous papers [8, 13]. Therefore SHE is caused by easy intergranular fracture due to impurity segregation in the 550 °C peak-aged condition, by the presence of coarse cementite and intrinsic hardening in the 600 °C peak-aged condition and by the combined effect of all three factors in the 650 °C peak-aged condition.
5. Conclusions The fracture behaviour in SHE of a 6W-3Ni steel isothermally peak aged in the range 550-650 °C was analysed in terms of the carbide morphology and the impurity segregation, which may be indirectly affected by the formation of tungsten carbides. (1) In the 600 and 650°C peak-aged conditions the impact toughness underwent a brittleto-ductile transition with increasing test temperature and the DBTT in the 650 °C condition was higher by 200 °C than that in the 600 °C condition. In the 550°C peak-aged condition, however, the impact toughess remained at the lower shelf energy level, about 5 J, even at 250 °C. (2) The fracture mode was intergranular in the 550 °C peak-aged condition whereas it comprised mostly transgranular dimples in the 600°C peak-aged condition. In the 650 °C peak-aged condition, where some intergranular area was observed, the transgranular cleavage area
180 c h a n g e d to t r a n s g r a n u l a r d i m p l e s at t e m p e r a t u r e s a b o v e 2 0 0 °C.
(3) SHE is caused by the easy occurrence of intergranular fracture due to impurity segregation in the 550 °C peak-aged condition, by the prese n c e of coarse ~ementite and intrinsic hardening in the 600°C peak-aged condition and by the combined action of all three factors in the 650 °C peak-aged condition.
Acknowledgment
This work was supported by the Sammi Foundation, 1990.
References 1 E.W. Lee, C.E. NeuandJ. Kozol, J. Met.,42(5)(1990) 16. 2 M. K. Tseng, D. Y. Lee and H. L. Marcus, Mater. Sci. Eng., 60 (1983) 73.
3 M. S. Bhat, W. M. Garrison Jr. and V. F. Zackay, Mater. Sci. Eng., 41 (1979) 1. 4 G. R. Simcoe, A. E. Nerenberg, V. Biss and A. P. Coldren, Trans. ASM, 61 (1968)834. 5 R. D. Goolsby, Ph.D. Thesis (LBL Rep. 405), University of Californiaat Berkeley, 1975. 6 H. Kwon and C. H. Kim, Metall. Trans. A, 14 (1983) 1389. 7 H. Kwon,Scr. Metall., 23 (1989) 1001. 8 K. H. Kim and H. Kwon, Mater. Sci. Eng. A, 127 (1990) 51. 9 J. J. Irani and R. W. K. Honeycombe, J. Iron Steel Inst., 203 (1966) 826. 10 D. Raynor, J. A. Whiteman and R. W. K. Honeycombe, J. Iron Steel Inst., 204 (1966) 349. 11 A.T. Davenport and R. W. K. Honeycombe, Met. Sci., 9 (1975) 201. 12 H. Kwon, Metall. Trans. A, 22 (1991) 1119. 13 K. H. Kim and H. Kwon, Scr. Metall. Mater., 24 (1990) 515. 14 C. K. Choi, K. B. Lee, K. H. Kim and H. Kwon, J. Korean Inst. Met., 29 ( 1991 ) 740. 15 J. Y u a n d C . J. McMahonJr.,Metall. Trans. A, l l ( 1 9 8 0 ) 227. 16 J. S. Song and H. Kwon, Mater. Sci. Eng. A, 117(1989) 133.
175
Fracture behaviour in secondary hardening embrittlement of isothermally peak-aged 6W-3Ni steel Kon Bae Lee, Chang Ki Choi and Hoon Kwon Department of Metallurgical Engineering, Kookmin University, Seou1136- 702 (Korea)
Kwang-Hoon Kim Department of Surface Finishing and Engineering, Tae-Hun Technical College, lncheon 401- 714 (Korea) (Received April 10, 1991 ; in revised form May 17, 1991 )
Abstract Secondary hardening embrittlement (SHE) in a 6W-3Ni steel isothermally peak aged in the range 550-650 °C was investigated by conducting impact testing in the range 20-250 °C. In the 600 and 650°C peak-aged conditions the impact toughness underwent a brittle-to-ductile transition with increasing test temperature, although the ductile-brittle transition temperature in the latter condition was higher by 200 °C than that in the former condition. In the 550 °C peak-aged condition, however, the impact toughness was not recovered and remained at the lower shelf energy level, about 5 J, even at 250 °C. The fracture mode was intergranular in the 550 °C peak-aged condition whereas it comprised mostly transgranular dimples in the 600 °C peak-aged condition. In the 650 °C peak-aged condition, where some intergranular area was observed, the area of transgranular cleavage changed to transgranular dimples at temperatures above 200 °C. The fracture behaviour has been discussed in terms of the impurity segregation associated with W2C precipitation, the presence of coarse cementite and the hardening behaviour. SHE is caused by: the easy occurrence of intergranular fracture due to impurity segregation in the 550 °C peak-aged condition; a decrease in ductile fracture energy due to the presence of coarse cementite and intrinsic hardening in the 600 °C peak-aged condition; and the combined effect of all three factors in the 650 °C peak-aged condition.
1. Introduction Secondary hardening steels have been used as heat-resistant alloys because of the high hardness retained at elevated temperatures. Recently, secondary hardening steels such as the AF1410 systern have been used for ultrahigh strength structures owing to their very high toughness [ 1 ] . After tempering in the secondary hardening ternperature range, however, embrittlement can occur. This embrittlement is referred to as secondary hardening embrittlement (SHE). According to its fracture mode, SHE can be classified as either intergranular or transgranular, It has been suggested that intergranular SHE is caused by impurity segregation [2, 3] and that transgranular SHE is associated with coarse boundary carbides [4-6]. Kwon [7] and Kim and Kwon [8] reported recently that intergranular 0921-5093/91/$3.50
SHE occurred in the overaged condition compared to transgranular SHE in the underaged condition for molybdenum steel and W-Ni steel. Hence the fracture behaviour in SHE can be greatly affected by the aging condition. The intrinsic secondary hardening due to precipitation of fine alloy carbides [9-11] also contributes to SHE. Since the precipitation of fine alloy carbides (i.e. Mo2C or W2C ) in molybdenum and tungsten steels leads to secondary hardening, the degree of secondary hardening is affected by the precipitation kinetics of alloy carbides [11]. Recently, Kwon [12] reported that a difference in degree of secondary hardening for molybdenum and tungsten steels has an important effect on SHE. In addition, Kim and Kwon [13] suggested that the different behaviour of SHE in the various iso© Elsevier Sequoia/Printed in The Netherlands
176
thermally peak-aged (i.e. aged to peak hardness) conditions for 6W-6Ni steel is associated with the differences in carbide morphology and impurity segregation. In 6W-6Ni and 6W-3Ni steels isochronally aged they [8] also investigated the differences in carbide morphology and fracture behaviour, particularly in the overaged condition where both steels had low impact toughness even at a high temperature of 250 °C. Hence there may be a variation in carbide morphology in isothermally peak-aged conditions for 6W-3Ni steel and the fracture behaviour in SHE may be different from that for 6W-6Ni steel. In this study SHE in a 6W-3Ni steel isothermally peak aged in the range 550-650 °C was analysed in terms of variations in impact toughness and fracture behaviour with carbide morphology and impurity segregation.
700 600
_~ ~ ~ 50o ~ 400
300
• ~ 1'0 TEMPERinGT~ME~h)
100
Fig. 1. Isothermal hardness variations at 550, 600 and 650 °C.
25 20
2. Experimental procedure
~ t5
An ingot of mass about 20 kg was prepared from high purity materials by vacuum induction melting. The ingot was hot rolled to 14 mm plate. Standard Charpy V-notch impact specimens were machined from the plate. The chemical composition of an experimental steel was 0.37C-6.18W3.18Ni-0.006P-0.004S (weight per cent). Specimens were austenitized under a flowing argon atmosphere at 1200°C for 1 h and oil quenched. They were then isothermally aged in the secondary hardening range 550-650 °C and water quenched, The micro-Vickers hardness was measured to determine the peak-aged conditions at 550, 600 and 650 °C. The peak-aging times were 20, 1 and 0.4 h at 550, 600 and 650 °C respectively. For the peak-aged conditions Charpy impact testing was performed in the range 20-250 °C. Microstructures were investigated using a Jeol 200CX transmission electron microscope operated at 160 kV. To investigate the fracture initiation area, which may be a primary factor in impact properties of relatively low toughness materials, all fractographs were taken from areas near the notch.
~
3. Experimental results Isothermal hardness variations in the range 550-650°C are shown in Fig. 1. The hardness initially decreased from the as-quenched value of 570 HV to values of 4 0 5 - 4 4 0 HV and then increased to the peak values of 4 3 5 - 4 9 0 H V
0. I
~
10
~
•
"
J
• 5 -
• '
100
'
200
300
TESTTEMPERATURE( OC) Fig. 2. Impact toughness variations in the 550, 600 and 650°Cpeak-ag edconditions.
before final softening. As the aging temperature increased, the time to peak hardness became progressively shorter and the peak hardness decreased. The peak hardness values were 490, 460 and 435 HV at 550, 600 and 650 °C respectively. This is very common for precipitation-hardening alloys. The impact toughness changes with test ternperature in the peak-aged condition are shown in Fig. 2. In the 600 and 650 °C peak-aged conditions the impact toughness underwent a brittleto-ductile transition with increasing temperature. If the ductile-brittle transition temperature (DBTT) is assumed to be a temperature at which the impact toughness has a mid-value (12.5 J) between the upper shelf energy (20 J) and the lower shelf energy (5 J), the DBTTs are about 0 and 210°C in the 600 and 650°C peak-aged conditions respectively. The DBTT in the 650 °C condition was shifted to a very high temperature compared to the 600°C condition, though the
177
hardness (435 HV) in the 650 °C condition was lower than that (460 HV) in the 600 °C condition. The 650 °C peak-aged condition results in more severe embrittlement in spite of its lower hardness. Hence SHE is not simply correlated with the hardness variation. On the other hand, the impact toughness in the 550°C peak-aged condition remained at a low level of 5 J even at 250°C. The fracture surfaces are shown in Figs. 3-5. The 600 °C peak-aged specimen exhibits ductile dimples containing a small amount of intergranular fracture at room temperature (Fig. 3). The 650 °C peak-aged specimen shows a mixture of intergranular fracture and transgranular quasicleavage at room temperature (Fig. 4(a)). The transgranular quasi-cleavage had a transition to dimples on approaching 250 °C near the upper shelf energy level, but some intergranular area remained (Fig. 4(b)). In the 550°C peak-aged condition, however, large amounts of intergranular fracture were observed at both room temperature and 250°C (Fig. 5). Thus no increase of impact toughness occurs even at 250 °C because of the easy occurrence of intergranular fracture. Transmission electron micrographs are shown in Figs. 6 and 7. During aging, the cementite initially precipitates and later dissolves into the matrix and M2C precipitates primarily at dislocations in the matrix. Alternatively, the cementite disintegrates by the in situ formation of M2C and finally M23C6 and/or M6C precipitate as the M2C dissolves [ 11 ]. In the 550 °C peak-aged condition
Fig. 3. Fractograph of the 600 °C peak-aged specimen tested at 20 °C.
extensive formation of very fine W2C carbides occurred within the lath as the long-time aging at a relatively low temperature led to a gradual dissolution of a large amount of cementite, as shown in Fig. 6(a), In the 600 °C peak-aged condition fine W2C particles were observed within laths whereas the coarse plate-like cementite persisted predominantly at the lath boundaries (Fig. 6(b)). Even in the 650°C overaged condition some coarse cementite still remained, although relatively fine W2C carbides were observed (Fig. 7). Hence it is certain that much more coarse cementite is retained in the 650 °C peak-aged condition relative to the other two conditions. 4. Discussion Before we discuss the peak-aged conditions, the results for the isothermally underaged and
Fig. 4. Fractographs of the 650 oC peak-aged specimen tested at (a) 20 and (b) 250 °C.
178
Fig. 6. Transmission electron micrographs of (a) 550 and (b) 600 °C peak-aged specimens.
Fig. 5. Fractographs of the 550°C peak-aged specimen tested at (a) 20 and (b) 250 °C.
overaged conditions in the range 550-650 °C are summarized as follows [14]. In the underaged conditions the impact toughness had high values above those at the DBTT and approached the upper shelf energy levels, 35-42 J. In the overaged conditions, however, the impact toughness stayed at the lower shelf energy levels, 2-5 J, even at 250°C. The fracture surfaces showed mostly transgranular dimples in the underaged condition whereas intergranular fracture was most often observed in the overaged condition. Since tungsten has quite similar properties to molybdenum, we postulate that tungsten in the matrix can tie up phosphorus, presumably by lowering its activity or by the formation of tungsten phosphide, resulting in inhibition of phosphorus segregation to grain boundaries in a similar way as in molybdenum steel [7, 15].
I
Fig. 7. Transmission electron micrograph of the specimen overaged at 650°C for 1 h.
There may be no available tungsten in the matrix to tie up phosphorus in the 550°C peak-aged condition because the cementite was mostly replaced by W2C. Hence impurity segregation to
179
grain boundaries can occur in that condition. In the 600 °C peak-aged condition, however, a relatively large amount of coarse cementite plates remained, in particular at the lath boundaries, The fact that the lath boundary cementite plates persist means that there may still remain available tungsten to tie up phosphorus. In the 650°C peak-aged condition, since a relatively small amount of coarse cementite remained and the amount of tungsten in the matrix may be insufficient to tie up all the phosphorus, residual impurifles beyond a certain amount to be tied up by the available tungsten can segregate to the grain boundaries, lntergranular fracture can be caused by impurities and/or carbides at grain boundaries. The impurities directly affect the grain boundary strength whereas the carbides act as stress concentrators. Thus impurity segregation plays a primary role in intergranular fracture in high strength alloy steels, Denser precipitation of carbides leads to higher stress concentrations. In addition, the difference in hardness (i.e. difference in extent of secondary hardening) affects the stress concentration susceptibility at grain boundaries. Stress relaxation in a higher hardness microstructure is more difficult because of the greater difficulty of plastic deformation. Hence in the 550°C peakaged condition of higher hardness and denser precipitation of carbides, brittle intergranular fracture can occur. In that condition, however, the easy occurrence of intergranular fracture fundamentally implies that impurity segregation to grain boundaries took place. In contrast, in the 650 °C peak-aged condition of lower hardness with less dense precipitation than in the 550 °C peak-aged condition, the stress concentration susceptibility may be lower. Furthermore, because relatively small amounts of impurities may segregate to grain boundaries, the occurrence of intergranular fracture is resisted, Thus the amount of intergranular fracture is small and the impact toughness undergoes a DBTT. However, the fact that the impact toughness approaches the upper shelf energy level only at temperatures above 250 °C in spite of the lower hardness indicates that some impurity segregation contributes significantly to the shift of DBTT to a very high temperature. This may be compared to the DBTT of the 600 °C peak-aged condition where there may be sufficient tungsten available to tie up the phosphorus. In contrast, for
the 6W-6Ni steel isothermally peak aged at 650 °C, since the available tungsten was consumed as a result of the near absence of coarse cementite and the extensive precipitation ofW2C, the impurity segregation led to the easy occurrence of intergranular fracture [ 13]. In the 600°C peak-aged condition the fracture surface consisted mostly of dimples with some intergranular area even at room temperature. Such intergranular areas were also observed in the 200 °C tempered condition in which almost all segregation of impurities occurred during austenitizing rather than during tempering [16]. In the isothermally underaged conditions in the range 550-650 °C the upper shelf energy levels were 35-42 J [14]. Thus a decrease in ductile fracture energy, which is caused by the presence of coarse cementite and the intrinsic hardening due to fine W2C precipitation, results in a large drop in upper shelf energy. A more detailed discussion was presented in previous papers [8, 13]. Therefore SHE is caused by easy intergranular fracture due to impurity segregation in the 550 °C peak-aged condition, by the presence of coarse cementite and intrinsic hardening in the 600 °C peak-aged condition and by the combined effect of all three factors in the 650 °C peak-aged condition.
5. Conclusions The fracture behaviour in SHE of a 6W-3Ni steel isothermally peak aged in the range 550-650 °C was analysed in terms of the carbide morphology and the impurity segregation, which may be indirectly affected by the formation of tungsten carbides. (1) In the 600 and 650°C peak-aged conditions the impact toughness underwent a brittleto-ductile transition with increasing test temperature and the DBTT in the 650 °C condition was higher by 200 °C than that in the 600 °C condition. In the 550°C peak-aged condition, however, the impact toughess remained at the lower shelf energy level, about 5 J, even at 250 °C. (2) The fracture mode was intergranular in the 550 °C peak-aged condition whereas it comprised mostly transgranular dimples in the 600°C peak-aged condition. In the 650 °C peak-aged condition, where some intergranular area was observed, the transgranular cleavage area
180 c h a n g e d to t r a n s g r a n u l a r d i m p l e s at t e m p e r a t u r e s a b o v e 2 0 0 °C.
(3) SHE is caused by the easy occurrence of intergranular fracture due to impurity segregation in the 550 °C peak-aged condition, by the prese n c e of coarse ~ementite and intrinsic hardening in the 600°C peak-aged condition and by the combined action of all three factors in the 650 °C peak-aged condition.
Acknowledgment
This work was supported by the Sammi Foundation, 1990.
References 1 E.W. Lee, C.E. NeuandJ. Kozol, J. Met.,42(5)(1990) 16. 2 M. K. Tseng, D. Y. Lee and H. L. Marcus, Mater. Sci. Eng., 60 (1983) 73.
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