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Some considerations on the occurrence of intergranular fracture during fatigue crack growth in steels
Materials Science and Engineering, 83 (1986) L11-L16
Lll
Letter
S o m e considerations on the occurrence of intergranular fracture during fatigue crack growth in steels
K. S. R A V I C H A N D R A N , and K I S H O R E
E. S. D W A R A K A D A S A
Department of Metallurgy, Indian Institute of Science, Bangalore 560012 (India) (Received March 25, 1986, in revised form June 3, 1986)
ABSTRACT
The occurrence o f a maximum in the percentage of intergranular fracture on the fracture surface during the transition from intermediate to low fatigue crack growth rates has been observed for a high strength steel. It is suggested that transgranular planar slip leading to slip localization is essential in promoting intergranular fracture when the cyclic plastic zone size becomes equal to the prior austenite grain size.
I. INTRODUCTION
It is very apparent that microstructure can affect fatigue crack growth behaviour at low values of the alternating stress intensity range AK, which is referred to as "microstructuresensitive" crack growth [1]. This is described as the condition in which the crack tip tends to follow the easy crack path, e.g. crystallographic slip planes or grain boundary regions. At this stage, a variety of fracture modes such as transgranular faceted fracture, intergranular fracture and transgranular quasicleavage fracture can occur depending on the deformation behaviour of the microstructure. A transition from this "microstructure-controlled" crack growth m o d e t o a "continuum-controlled" striation crack growth m o d e occurs as AK is increased. Generally, a consequent change in 0025-5416/86/$3.50
the fracture behaviour from a more crystallographic nature to a striation mode of crack advance accompanies this transition. This can be physically appreciated in terms of the relationship between the size of the cyclic plastic zone ahead of the crack tip and the grain size. An increasing number of investigations [1-4] on the dependence of fatigue crack growth behaviour on microstructure clearly demonstrate this phenomenon. A crystallographic crack path can be expected to occur if the crack tip deformation is localized along slip planes when the size of the cyclic plastic zone is comparable with the grain size [1, 2]. However, when the cyclic plastic zone size exceeds the grain size, more homogeneous deformation resulting in crack growth by striation formation occurs. In steels, this transition in macroscopic crack growth is reflected as the incidence of the maximum a m o u n t of intergranular fracture at a AK level (hereafter referred to as AKcPzs= ~s where CPZS means the cyclic plastic zone size and GS the grain size} at which the cyclic plastic zone becomes equal to the grain size [2]. While there is a significant a m o u n t of experimental evidence [2~ 3, 5, 6] available to show the incidence of the m a x i m u m percentage of intergranular fracture at AKcPzs=Gs, it appears that no complete explanation is available to describe this behaviour. Also, disagreement with this trend is reported in an investigation [4] in which intergranular fracture was n o t observed at AKcPzs=Gs. The present Letter is addressed to the discussion of some recent results obtained in our investigations with regard to the occurrence of intergranular fracture at AKcFzs=Gs. The discussion is intended (1) to provide additional experimental evidence for the occurrence of a maximum in the percentage of intergranular fracture on the fracture surface at AKc~zs=6 s and (2) to describe the deformation conditions ahead of the crack tip which will or will n o t lead to the incidence of intergranular fracture at AKcPzs=Gs. © Elsevier Sequoia/Printed in The Netherlands
L12 2. P R E V I O U S W O R K
The results on steels reported b y other investigators [2, 3, 5, 6] indicate that good correlation exists between the occurrence of a m a x i m u m in the a m o u n t of intergranular fracture and AKcPzs=Gs. Cooke et al. [2] observed a m a x i m u m in the percentage of intergranular fracture (approximately 60%) at AKcPzs=Gs in a quenched-and-tempered medium carbon steel, and it was attributed to be a consequence of reverse crack tip plasticity rather than a tensile tearing fracture m o d e . Irving and Kurzfeld [3], in a comparison of the effects of steel purity* and vacuum on near-threshold fatigue crack growth behaviour, reported a maximum in the incidence of intergranular fracture at AK = 12-16 MPa m 1/2 . Interestingly a commercially pure steel tested in laboratory air exhibited intergranular fracture of the order of 60% at AK = 12-16 MPa m 1/2 while tests in vacuum resulted in a reduction in the extent of intergranular fracture to a b o u t 30%. A high purity quenched-and-tempered steel had a m a x i m u m of a b o u t 30% intergranular fracture during tests in air, whereas tests in a vacuum on the same steel indicated that intergranular fracture was nearly completely absent. Frandsen and Marcus [6] also showed that fatigue fracture surfaces of specimens tested in a hydrogen environment exhibited intergranular fracture at AKcPzS=GS , while intergranular fracture was n o t observed in tests conducted in vacuum. In contrast with these observations, Cheruvu [4] did n o t observe intergranular fracture at ~KcPzs=Gs in a commercially pure steel (similar to that used in ref. 2) tested in laboratory air and confirmed that no correlation existed between the prior austenite grain size and the reverse plastic zone size at the maximum incidence of intergranular fracture. Contrary to these observations [2, 3, 5, 6] of intergranular fracture at AKcPzs=os, Clarke et al. [7] and Irving and Kurzfeld [3] argued that the occurrence of intergranular fracture during fatigue crack growth is instead related to crack velocity through the crack growth rate at which the crack tip environmental species m a y be subject to favourable conditions *Here, purity refers to a material greatly refined b y v a c u u m arc r e f i n i n g .
for promoting intergranular fracture by diffusion to grain boundaries. Tests at different frequencies and load ratios delineated clearly the influence of the half-period at which the crack is open and Km~ in a fatigue cycle. At a particular crack velocity, obtained by calculating the p r o d u c t of the measured da/dN value and the imposed frequency of loading, a maximum in the amount of intergranular fracture was observed irrespective of the frequency and waveform of the fatigue cycle. The dependence of intergranular fracture on crack velocity strongly suggested an environmentally induced fracture mechanism involving hydrogen. Thus, the occurrence of intergranular fracture was mainly viewed [3, 7] as a p h e n o m e n o n influenced by crack velocity. However, tests in vacuum reported by other researchers [ 5] indicate that intergranular fracture up to a b o u t 30% can occur at AKcrzs=os in the absence of an environmental influence in a commercially pure steel. Despite the high purity levels of the steels used, tests [3] in air exhibited intergranular fracture of a b o u t 30% when the monotonic plastic zone size was close to the grain size. This indicates that, in the absence of an impurity effect, the environmental influence will contribute to intergranular fracture, and, in the absence of an environmental influence, the impurity effect will do so. From the tests conducted in vacuum, it m a y be suggested that the cyclic crack tip plasticity and its relation to microstructure can also be important. The susceptibility of a material to hydrogen embrittlement has been correlated with slip m o d e by some investigators [8, 9]. These correlations show that coplanar dislocation motion favours hydrogen embrittlement. Alloys which have a low stacking fault energy and short-range order are more prone to hydrogen embrittlement and dislocation transport o f hydrogen has been suggested as a vital step. The mechanism of hydrogen embrittlement by dislocation transport has been discussed elsewhere [8, 9]. In fatigue crack growth, the dislocations emitted from the crack tip can transport the hydrogen in the moist air at the crack tip and cause saturation of the hydrogen at the grain boundary. The efficiency of this process increases as the slip m o d e becomes more planar in character. When the transition from stage I to stage II
L13 occurs during fatigue crack growth, i.e. when the cyclic plastic zone size equals the grain size, localized fatigue damage along crystallographic planes could be an effective mechanism for the dislocation transport of hydrogen. It has also been shown [6, 10] that, in the presence of hydrogen, the amount of intergranular fracture at AKcPzs=Gs increases significantly. It has been suggested [10-12] that intergranular fracture during the transition in the crack growth behaviour is favoured by the segregation of impurities to the grain boundary [10, 13] as well as by an environmental influence through crack velocity [3, 7]. In an impurity-induced embrittlement effect, elements such as phosphorus, antimony and arsenic segregate to the grain boundaries and act as recombinant poisons for hydrogen at the grain boundary made available by a dislocation transport mechanism and lead to intergranular fracture. As discussed previously, results of tests on high purity and commercial purity steels in air and vacuum clearly demonstrate that intergranular fracture can occur when these t w o conditions, i.e. t h e impurity effect and the environmental influence, are n o t both present. However, in either case, if intergranular fracture is to occur at AKcPzs=Gs, a dislocation transport mechanism leading to the build-up of hydrogen at the grain boundary in tests in air or a dislocation pile-up at the grain boundary causing a stress concentration to result in intergranular fracture for tests in vacuum is evidently necessary.
3. EXPERIMENTAL PROCEDURE In the present investigation, fractographic observations of intergranular fracture on
fatigue fracture surfaces of a quenched-andtempered high strength steel were made. The steel had the composition 0.32 wt.% C, 0.96 wt.% Cr, 1.1 wt.% Mn, 1.2 wt.% Si, 0.22 wt.% Ti, 0.015 wt.% S, 0.02 wt.% P and balance iron. C o m p a c t tension samples, machined out of forged stock, were austenitized at 1100 °C for 30 min and oil quenched, followed by tempering at 300 °C, 4 0 0 ° C and 530 °C for 135 min, 35 min and 25 min respectively. Fatigue crack growth tests were performed in a servohydraulic machine at a frequency of 3 5 Hz in laboratory air with a load ratio R of 0.05. A conventional load-shedding technique [14] was used to approach low crack growth rates. The fracture surfaces were examined in the scanning electron microscope and a quantitative correlation of the intergranular fracture was made.
4. RESULTS AND DISCUSSION The mechanical properties after the different tempering treatments are presented in Table 1. The prior austenite grain size was 6.5 ~m. In Table 2 are presented the values of AKcPzs=6s and the percentage of intergranular fracture for microstructures tempered at different temperatures. Figure 1 and Fig. 2 illustrate the magnitudes of intergranular fracture occurring at different A K levels for the microstructures tempered at 400 °C and 530 °C respectively. A maximum in the percentage of intergranular fracture occurring at AKcrzs=Gs (cf. Table 2) is clearly evident. It should be noted that, for the microstructure tempered at 300 °C, no intergranular fracture was encountered at any AK level. The fracture mor-
TABLE 1 Mechanical property data for the steel investigated Tempering temperature
0.2% yield strength
Ultimate tensile strength
(°C)
(MPa)
(MPa)
300 400 530
1305 1249 1096
1467 1371 1164
Elongation (%)
Fracture strain ef
Strainhardening exponent n
11 11 12
0.55 0.66 0.8
0.55 0.4 0.19
L14 TABLE 2 Calculated and experimentally observed values of AKcPzs=Gs and the amount of intergranular fracture for the microstructures investigated
Tempering temperature (°C)
AKcpzs=~ s a (MPa ml/2)
300 400 530
20.4 19.6 17.2
b (MPa m 1/2)
AKIG
-
-
17-20 17-20
Maximum amount of intergranular fracture at AKcPzS=GS (%)
Amount of intergranular fracture at AKth
0 ~-63 ~ 42
0 ~40 ~ 15
(%)
a Calculated from equating the cyclic plastic zone size (1/3~)(AK/2Oys) 2, to the prior austenitic grain size, 6.5 pm. b AK at which a m ax i m u m in the percentage of intergranular fracture was observed experimentally.
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Fig. 1. Variation in the amount of faceted fracture with AK close to threshold for the microstructure tempered at 400 °C.
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Fig. 2. Variation in the amount of faceted fracture with AK close to threshold for the microstructure tempered at 530 °C.
phologies accompanying crack growth near threshold are illustrated in the fractographs presented in Fig. 3. In the present investigation, the microstructures tempered at 400 and 530 °C exhibited intergranular fracture at AKc~zs=~s while the microstructure tempered at 300 °C did not show such an occurrence. Tempering at 300 °C results in a microstructure contain-
ing highly dislocated lath martensite which exhibits a high strain-hardening behaviour {n = 0.55) and a high yield strength {1305 MPa), while tempering at 530 °C produces a microstructure with a low strain-hardening e x p o n e n t (n = 0.19) and a relatively low yield strength (1096 MPa). When the transition to microstructure-sensitive crack growth occurs at AKcPzs=Gs, during testing by load reduction the crack tip tends to propagate in a crystallographic manner. Crack propagation along slip planes will be difficult when the high strain-hardening matrix constitutes resistance to deformation along slip planes or slip localization. This difficulty results in a tensile stress normal to the crack plane and, when this stress exceeds the critical stress for cleavage, flat transgranular fracture occurs. The presence of an environmental effect {moist air in the laboratory environment) can also p r o m o t e such a fracture by a reduction in lattice cohesive strength [ 15 ]. Such a fracture m ode occurring in the microstructure tempered at 300 °C at AKcPzS=GS is illustrated in Fig. 3(a). In contrast, crack propagation along crystallographic planes is observed in structures tempered at 400 and 530 °C. In Fig. 3{b), for the microstructure tempered at 4 0 0 °C, a transition from a striation-controlled crack propagation m o d e to a predominantly faceted fracture m o d e is shown to occur at AK = 23 MPa m 1/2, when the size of the cyclic plastic zone approaches that of the prior austenitic grains. At a slightly lower AK level, i.e. at AKcpzs=GS = 20 MPa m 1/2, a maximum amount of intergranular fracture occurs, as shown in Fig. 3(c). The microstructure tempered at 530 °C exhibits a partially trans-
L15
Fig. 3. (a) Fractograph showing fiat transgranular fracture at ~/~CPZS----GSfor the microstructure tempered at 300 °C; (b) fractograph illustrating the transition from striation to transgranular faceted fracture at AK = 23 MPa ml/2; (c) at AKcPzS_--GS= 20 MPa m 1/2 the incidence of the maximum amount of intergranular fracture for the microstructure austenitized at 1100 °C and tempered at 400 °C; (d) mixed transgranular and intergranular fracture at AKcPzs=Gs = 20 MPa m1/2 for the microstructure austenitized at 1100 °C and tempered at 530 °C.
granular faceted fracture and partially intergranular fracture at AKcPzs=6s. In Fig. 3(d), transgranular facets together with intergranular regions can be seen. These observations suggest a mechanism involving transgranular planar slip, leading to intergranular fracture in the n e x t grain in microstructures exhibiting a low strain-hardening behaviour and a relatively low strength. These observations in fact are also consistent with the suggestion [10] of the mechanism of dislocation pile-up along slip planes for hydrogen transport to the grain boundary for the observed fracture m o d e transitions in fatigue. Hence, it m a y be argued t h a t the process of nucleation of intergranular
fracture at AKc~zs=Gs necessitates transgranular crystallographic slip either to promote local hydrogen concentration or to cause stress concentration by dislocation pile-up at the grain boundary as a preceding event. This also explains the observations of the absence of intergranular fracture at AKcezs=Gs in the microstructure tempered at 300 °C, in that the difficulty in promoting transgranular crystallographic slip as a result of high strainhardening behaviour makes both dislocation transport of hydrogen and dislocation pile-up difficult. The results of this investigation, discussed above, clearly suggest that transgranular
L16
crystallographic slip should act if intergranular fracture is to occur at ~KcPzs=Gs. When the cyclic plastic zone size becomes equal to the prior austenite grain size, extensive slip band activity causes increased slip reversibility and slip band fatigue damage, both of which are effective mechanisms for both dislocation transport of hydrogen from the laboratory air and dislocation pile-up leading to stress concentration. Hence, it can be argued that a correlation between the maxim u m in intergranular fracture and ~KcPzs=Gs is meaningful in situations of deformation localization along slip planes ahead of the crack tip. However, the influence of crack velocity on intergranular fracture as demonstrated by some investigators [3, 7] cannot be discounted, and this could be prominent in situations where crystallographic slip is restricted because of a high yield strength and a lightly tempered microstructure. In fact, the observations [2, 5] of intergranular fracture at AKcPzs=oswere made on Ni-Cr-Mo steels after tempering at around 400-500 °C, where the yield strengths are relatively low (1275-1350 MPa). However, the absence [4] of a correlation between AKcPzs=Gs and intergranular fracture was noted especially at high strength levels (1580-1620 MPa) after tempering at around 200 °C. While these differences cannot be used to rationalize completely in favour of the arguments made in this Letter because of a lack of knowledge of the strain-hardening behaviour in the other studies, these still suggest the possibility of a yield strength effect. The interrelationship between the crack velocity and AKcPzs=Gs is not clear at the present time. However, this Letter demonstrates the importance
of transgranular crystallographic slip to prom o t e intergranular fracture at AKc~s=Gs. Further research is needed to isolate the effect of each of these factors on the occurrence of intergranular fracture during fatigue crack growth.
REFERENCES 1 R. J. Cooke and C. J. Beevers, Mater. Sci. Eng., 13 (1974) 201. 2 R. J. Cooke, P. E. Irving, G. S. Booth and C. J. Beevers, Eng. Fract. Mech., 7 (1975) 69. 3 P. E. Irving and A. Kurzfeld, Met. Sci., 12 (1978) 495. 4 N. S. Cheruvu, Fraetography of Ceramic and Metal Failures, ASTM Spec. Tech. Publ. 827, 1984, p. 309. 5 R . P . V . Evans, N. B. Owen and B. E. Hopkins, Eng. Fract. Mech., 3 (1971) 463. 6 J. D. Frandsen and H. L. Marcus, Scr. Metall., 9 (1975) 1089. 7 G. Clarke, A. C. Pickard and J. F. Knott, Eng. Fract. Mech., 8 (1976) 449. 8 M. R. Louthan, Jr., G. R. Caskey, Jr., J. A. Donovan and R. E. Rawl, Jr., Mater. Sci. Eng., 10 (1972) 357. 9 M. R. Louthan, Jr., Hydrogen in Metals, American Society for Metals, Metals Park, OH, 1974, p. 53. 10 W. L. Morris, J. D. Frandsen and H. L. Marcus, in C. D. Beachem and W. R. Warke (eds.), Fractography-Microscopic Cracking Processes, ASTM Spec. Tech. Publ. 600, 1977, p. 49. 11 R. O. Ritchie, Eng. Fract. Mech., 7 ( 1 9 7 5 ) 187. 12 R. O. Ritchie, Metall. Trans. A, 8 (1977)1131. 13 C. J. McMahon, C. L. Briant and S. K. Banerji, in D. M. R. Taplin (ed.), Fracture 1977, Proc. 4th Int. Conf. on Fracture, Waterloo, Ontario, June 19-24, 1977, Voh 1, University of Waterloo Press, Waterloo, Ontario, 1977, p. 363. 14 R. O° Ritchie, Met. Sci., 11 (1977) 368. 15 R. A. Oriani and P. H. Josephic, ActaMetall., 22 (1974) 1065.
Lll
Letter
S o m e considerations on the occurrence of intergranular fracture during fatigue crack growth in steels
K. S. R A V I C H A N D R A N , and K I S H O R E
E. S. D W A R A K A D A S A
Department of Metallurgy, Indian Institute of Science, Bangalore 560012 (India) (Received March 25, 1986, in revised form June 3, 1986)
ABSTRACT
The occurrence o f a maximum in the percentage of intergranular fracture on the fracture surface during the transition from intermediate to low fatigue crack growth rates has been observed for a high strength steel. It is suggested that transgranular planar slip leading to slip localization is essential in promoting intergranular fracture when the cyclic plastic zone size becomes equal to the prior austenite grain size.
I. INTRODUCTION
It is very apparent that microstructure can affect fatigue crack growth behaviour at low values of the alternating stress intensity range AK, which is referred to as "microstructuresensitive" crack growth [1]. This is described as the condition in which the crack tip tends to follow the easy crack path, e.g. crystallographic slip planes or grain boundary regions. At this stage, a variety of fracture modes such as transgranular faceted fracture, intergranular fracture and transgranular quasicleavage fracture can occur depending on the deformation behaviour of the microstructure. A transition from this "microstructure-controlled" crack growth m o d e t o a "continuum-controlled" striation crack growth m o d e occurs as AK is increased. Generally, a consequent change in 0025-5416/86/$3.50
the fracture behaviour from a more crystallographic nature to a striation mode of crack advance accompanies this transition. This can be physically appreciated in terms of the relationship between the size of the cyclic plastic zone ahead of the crack tip and the grain size. An increasing number of investigations [1-4] on the dependence of fatigue crack growth behaviour on microstructure clearly demonstrate this phenomenon. A crystallographic crack path can be expected to occur if the crack tip deformation is localized along slip planes when the size of the cyclic plastic zone is comparable with the grain size [1, 2]. However, when the cyclic plastic zone size exceeds the grain size, more homogeneous deformation resulting in crack growth by striation formation occurs. In steels, this transition in macroscopic crack growth is reflected as the incidence of the maximum a m o u n t of intergranular fracture at a AK level (hereafter referred to as AKcPzs= ~s where CPZS means the cyclic plastic zone size and GS the grain size} at which the cyclic plastic zone becomes equal to the grain size [2]. While there is a significant a m o u n t of experimental evidence [2~ 3, 5, 6] available to show the incidence of the m a x i m u m percentage of intergranular fracture at AKcPzs=Gs, it appears that no complete explanation is available to describe this behaviour. Also, disagreement with this trend is reported in an investigation [4] in which intergranular fracture was n o t observed at AKcPzs=Gs. The present Letter is addressed to the discussion of some recent results obtained in our investigations with regard to the occurrence of intergranular fracture at AKcFzs=Gs. The discussion is intended (1) to provide additional experimental evidence for the occurrence of a maximum in the percentage of intergranular fracture on the fracture surface at AKc~zs=6 s and (2) to describe the deformation conditions ahead of the crack tip which will or will n o t lead to the incidence of intergranular fracture at AKcPzs=Gs. © Elsevier Sequoia/Printed in The Netherlands
L12 2. P R E V I O U S W O R K
The results on steels reported b y other investigators [2, 3, 5, 6] indicate that good correlation exists between the occurrence of a m a x i m u m in the a m o u n t of intergranular fracture and AKcPzs=Gs. Cooke et al. [2] observed a m a x i m u m in the percentage of intergranular fracture (approximately 60%) at AKcPzs=Gs in a quenched-and-tempered medium carbon steel, and it was attributed to be a consequence of reverse crack tip plasticity rather than a tensile tearing fracture m o d e . Irving and Kurzfeld [3], in a comparison of the effects of steel purity* and vacuum on near-threshold fatigue crack growth behaviour, reported a maximum in the incidence of intergranular fracture at AK = 12-16 MPa m 1/2 . Interestingly a commercially pure steel tested in laboratory air exhibited intergranular fracture of the order of 60% at AK = 12-16 MPa m 1/2 while tests in vacuum resulted in a reduction in the extent of intergranular fracture to a b o u t 30%. A high purity quenched-and-tempered steel had a m a x i m u m of a b o u t 30% intergranular fracture during tests in air, whereas tests in a vacuum on the same steel indicated that intergranular fracture was nearly completely absent. Frandsen and Marcus [6] also showed that fatigue fracture surfaces of specimens tested in a hydrogen environment exhibited intergranular fracture at AKcPzS=GS , while intergranular fracture was n o t observed in tests conducted in vacuum. In contrast with these observations, Cheruvu [4] did n o t observe intergranular fracture at ~KcPzs=Gs in a commercially pure steel (similar to that used in ref. 2) tested in laboratory air and confirmed that no correlation existed between the prior austenite grain size and the reverse plastic zone size at the maximum incidence of intergranular fracture. Contrary to these observations [2, 3, 5, 6] of intergranular fracture at AKcPzs=os, Clarke et al. [7] and Irving and Kurzfeld [3] argued that the occurrence of intergranular fracture during fatigue crack growth is instead related to crack velocity through the crack growth rate at which the crack tip environmental species m a y be subject to favourable conditions *Here, purity refers to a material greatly refined b y v a c u u m arc r e f i n i n g .
for promoting intergranular fracture by diffusion to grain boundaries. Tests at different frequencies and load ratios delineated clearly the influence of the half-period at which the crack is open and Km~ in a fatigue cycle. At a particular crack velocity, obtained by calculating the p r o d u c t of the measured da/dN value and the imposed frequency of loading, a maximum in the amount of intergranular fracture was observed irrespective of the frequency and waveform of the fatigue cycle. The dependence of intergranular fracture on crack velocity strongly suggested an environmentally induced fracture mechanism involving hydrogen. Thus, the occurrence of intergranular fracture was mainly viewed [3, 7] as a p h e n o m e n o n influenced by crack velocity. However, tests in vacuum reported by other researchers [ 5] indicate that intergranular fracture up to a b o u t 30% can occur at AKcrzs=os in the absence of an environmental influence in a commercially pure steel. Despite the high purity levels of the steels used, tests [3] in air exhibited intergranular fracture of a b o u t 30% when the monotonic plastic zone size was close to the grain size. This indicates that, in the absence of an impurity effect, the environmental influence will contribute to intergranular fracture, and, in the absence of an environmental influence, the impurity effect will do so. From the tests conducted in vacuum, it m a y be suggested that the cyclic crack tip plasticity and its relation to microstructure can also be important. The susceptibility of a material to hydrogen embrittlement has been correlated with slip m o d e by some investigators [8, 9]. These correlations show that coplanar dislocation motion favours hydrogen embrittlement. Alloys which have a low stacking fault energy and short-range order are more prone to hydrogen embrittlement and dislocation transport o f hydrogen has been suggested as a vital step. The mechanism of hydrogen embrittlement by dislocation transport has been discussed elsewhere [8, 9]. In fatigue crack growth, the dislocations emitted from the crack tip can transport the hydrogen in the moist air at the crack tip and cause saturation of the hydrogen at the grain boundary. The efficiency of this process increases as the slip m o d e becomes more planar in character. When the transition from stage I to stage II
L13 occurs during fatigue crack growth, i.e. when the cyclic plastic zone size equals the grain size, localized fatigue damage along crystallographic planes could be an effective mechanism for the dislocation transport of hydrogen. It has also been shown [6, 10] that, in the presence of hydrogen, the amount of intergranular fracture at AKcPzs=Gs increases significantly. It has been suggested [10-12] that intergranular fracture during the transition in the crack growth behaviour is favoured by the segregation of impurities to the grain boundary [10, 13] as well as by an environmental influence through crack velocity [3, 7]. In an impurity-induced embrittlement effect, elements such as phosphorus, antimony and arsenic segregate to the grain boundaries and act as recombinant poisons for hydrogen at the grain boundary made available by a dislocation transport mechanism and lead to intergranular fracture. As discussed previously, results of tests on high purity and commercial purity steels in air and vacuum clearly demonstrate that intergranular fracture can occur when these t w o conditions, i.e. t h e impurity effect and the environmental influence, are n o t both present. However, in either case, if intergranular fracture is to occur at AKcPzs=Gs, a dislocation transport mechanism leading to the build-up of hydrogen at the grain boundary in tests in air or a dislocation pile-up at the grain boundary causing a stress concentration to result in intergranular fracture for tests in vacuum is evidently necessary.
3. EXPERIMENTAL PROCEDURE In the present investigation, fractographic observations of intergranular fracture on
fatigue fracture surfaces of a quenched-andtempered high strength steel were made. The steel had the composition 0.32 wt.% C, 0.96 wt.% Cr, 1.1 wt.% Mn, 1.2 wt.% Si, 0.22 wt.% Ti, 0.015 wt.% S, 0.02 wt.% P and balance iron. C o m p a c t tension samples, machined out of forged stock, were austenitized at 1100 °C for 30 min and oil quenched, followed by tempering at 300 °C, 4 0 0 ° C and 530 °C for 135 min, 35 min and 25 min respectively. Fatigue crack growth tests were performed in a servohydraulic machine at a frequency of 3 5 Hz in laboratory air with a load ratio R of 0.05. A conventional load-shedding technique [14] was used to approach low crack growth rates. The fracture surfaces were examined in the scanning electron microscope and a quantitative correlation of the intergranular fracture was made.
4. RESULTS AND DISCUSSION The mechanical properties after the different tempering treatments are presented in Table 1. The prior austenite grain size was 6.5 ~m. In Table 2 are presented the values of AKcPzs=6s and the percentage of intergranular fracture for microstructures tempered at different temperatures. Figure 1 and Fig. 2 illustrate the magnitudes of intergranular fracture occurring at different A K levels for the microstructures tempered at 400 °C and 530 °C respectively. A maximum in the percentage of intergranular fracture occurring at AKcrzs=Gs (cf. Table 2) is clearly evident. It should be noted that, for the microstructure tempered at 300 °C, no intergranular fracture was encountered at any AK level. The fracture mor-
TABLE 1 Mechanical property data for the steel investigated Tempering temperature
0.2% yield strength
Ultimate tensile strength
(°C)
(MPa)
(MPa)
300 400 530
1305 1249 1096
1467 1371 1164
Elongation (%)
Fracture strain ef
Strainhardening exponent n
11 11 12
0.55 0.66 0.8
0.55 0.4 0.19
L14 TABLE 2 Calculated and experimentally observed values of AKcPzs=Gs and the amount of intergranular fracture for the microstructures investigated
Tempering temperature (°C)
AKcpzs=~ s a (MPa ml/2)
300 400 530
20.4 19.6 17.2
b (MPa m 1/2)
AKIG
-
-
17-20 17-20
Maximum amount of intergranular fracture at AKcPzS=GS (%)
Amount of intergranular fracture at AKth
0 ~-63 ~ 42
0 ~40 ~ 15
(%)
a Calculated from equating the cyclic plastic zone size (1/3~)(AK/2Oys) 2, to the prior austenitic grain size, 6.5 pm. b AK at which a m ax i m u m in the percentage of intergranular fracture was observed experimentally.
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Fig. 1. Variation in the amount of faceted fracture with AK close to threshold for the microstructure tempered at 400 °C.
80
f
~ 60 ~ IJ
~r'~ . . . . . .
'Ly Tro.sgrQnul~r FOCetS
~
~
Inte~ranulor • Tror~granulor f0¢~1$
AUST'C TEMPE~'C .bo s3o
~2°oL ~< 6
to
l& &K
18
22
26
MPo
Fig. 2. Variation in the amount of faceted fracture with AK close to threshold for the microstructure tempered at 530 °C.
phologies accompanying crack growth near threshold are illustrated in the fractographs presented in Fig. 3. In the present investigation, the microstructures tempered at 400 and 530 °C exhibited intergranular fracture at AKc~zs=~s while the microstructure tempered at 300 °C did not show such an occurrence. Tempering at 300 °C results in a microstructure contain-
ing highly dislocated lath martensite which exhibits a high strain-hardening behaviour {n = 0.55) and a high yield strength {1305 MPa), while tempering at 530 °C produces a microstructure with a low strain-hardening e x p o n e n t (n = 0.19) and a relatively low yield strength (1096 MPa). When the transition to microstructure-sensitive crack growth occurs at AKcPzs=Gs, during testing by load reduction the crack tip tends to propagate in a crystallographic manner. Crack propagation along slip planes will be difficult when the high strain-hardening matrix constitutes resistance to deformation along slip planes or slip localization. This difficulty results in a tensile stress normal to the crack plane and, when this stress exceeds the critical stress for cleavage, flat transgranular fracture occurs. The presence of an environmental effect {moist air in the laboratory environment) can also p r o m o t e such a fracture by a reduction in lattice cohesive strength [ 15 ]. Such a fracture m ode occurring in the microstructure tempered at 300 °C at AKcPzS=GS is illustrated in Fig. 3(a). In contrast, crack propagation along crystallographic planes is observed in structures tempered at 400 and 530 °C. In Fig. 3{b), for the microstructure tempered at 4 0 0 °C, a transition from a striation-controlled crack propagation m o d e to a predominantly faceted fracture m o d e is shown to occur at AK = 23 MPa m 1/2, when the size of the cyclic plastic zone approaches that of the prior austenitic grains. At a slightly lower AK level, i.e. at AKcpzs=GS = 20 MPa m 1/2, a maximum amount of intergranular fracture occurs, as shown in Fig. 3(c). The microstructure tempered at 530 °C exhibits a partially trans-
L15
Fig. 3. (a) Fractograph showing fiat transgranular fracture at ~/~CPZS----GSfor the microstructure tempered at 300 °C; (b) fractograph illustrating the transition from striation to transgranular faceted fracture at AK = 23 MPa ml/2; (c) at AKcPzS_--GS= 20 MPa m 1/2 the incidence of the maximum amount of intergranular fracture for the microstructure austenitized at 1100 °C and tempered at 400 °C; (d) mixed transgranular and intergranular fracture at AKcPzs=Gs = 20 MPa m1/2 for the microstructure austenitized at 1100 °C and tempered at 530 °C.
granular faceted fracture and partially intergranular fracture at AKcPzs=6s. In Fig. 3(d), transgranular facets together with intergranular regions can be seen. These observations suggest a mechanism involving transgranular planar slip, leading to intergranular fracture in the n e x t grain in microstructures exhibiting a low strain-hardening behaviour and a relatively low strength. These observations in fact are also consistent with the suggestion [10] of the mechanism of dislocation pile-up along slip planes for hydrogen transport to the grain boundary for the observed fracture m o d e transitions in fatigue. Hence, it m a y be argued t h a t the process of nucleation of intergranular
fracture at AKc~zs=Gs necessitates transgranular crystallographic slip either to promote local hydrogen concentration or to cause stress concentration by dislocation pile-up at the grain boundary as a preceding event. This also explains the observations of the absence of intergranular fracture at AKcezs=Gs in the microstructure tempered at 300 °C, in that the difficulty in promoting transgranular crystallographic slip as a result of high strainhardening behaviour makes both dislocation transport of hydrogen and dislocation pile-up difficult. The results of this investigation, discussed above, clearly suggest that transgranular
L16
crystallographic slip should act if intergranular fracture is to occur at ~KcPzs=Gs. When the cyclic plastic zone size becomes equal to the prior austenite grain size, extensive slip band activity causes increased slip reversibility and slip band fatigue damage, both of which are effective mechanisms for both dislocation transport of hydrogen from the laboratory air and dislocation pile-up leading to stress concentration. Hence, it can be argued that a correlation between the maxim u m in intergranular fracture and ~KcPzs=Gs is meaningful in situations of deformation localization along slip planes ahead of the crack tip. However, the influence of crack velocity on intergranular fracture as demonstrated by some investigators [3, 7] cannot be discounted, and this could be prominent in situations where crystallographic slip is restricted because of a high yield strength and a lightly tempered microstructure. In fact, the observations [2, 5] of intergranular fracture at AKcPzs=oswere made on Ni-Cr-Mo steels after tempering at around 400-500 °C, where the yield strengths are relatively low (1275-1350 MPa). However, the absence [4] of a correlation between AKcPzs=Gs and intergranular fracture was noted especially at high strength levels (1580-1620 MPa) after tempering at around 200 °C. While these differences cannot be used to rationalize completely in favour of the arguments made in this Letter because of a lack of knowledge of the strain-hardening behaviour in the other studies, these still suggest the possibility of a yield strength effect. The interrelationship between the crack velocity and AKcPzs=Gs is not clear at the present time. However, this Letter demonstrates the importance
of transgranular crystallographic slip to prom o t e intergranular fracture at AKc~s=Gs. Further research is needed to isolate the effect of each of these factors on the occurrence of intergranular fracture during fatigue crack growth.
REFERENCES 1 R. J. Cooke and C. J. Beevers, Mater. Sci. Eng., 13 (1974) 201. 2 R. J. Cooke, P. E. Irving, G. S. Booth and C. J. Beevers, Eng. Fract. Mech., 7 (1975) 69. 3 P. E. Irving and A. Kurzfeld, Met. Sci., 12 (1978) 495. 4 N. S. Cheruvu, Fraetography of Ceramic and Metal Failures, ASTM Spec. Tech. Publ. 827, 1984, p. 309. 5 R . P . V . Evans, N. B. Owen and B. E. Hopkins, Eng. Fract. Mech., 3 (1971) 463. 6 J. D. Frandsen and H. L. Marcus, Scr. Metall., 9 (1975) 1089. 7 G. Clarke, A. C. Pickard and J. F. Knott, Eng. Fract. Mech., 8 (1976) 449. 8 M. R. Louthan, Jr., G. R. Caskey, Jr., J. A. Donovan and R. E. Rawl, Jr., Mater. Sci. Eng., 10 (1972) 357. 9 M. R. Louthan, Jr., Hydrogen in Metals, American Society for Metals, Metals Park, OH, 1974, p. 53. 10 W. L. Morris, J. D. Frandsen and H. L. Marcus, in C. D. Beachem and W. R. Warke (eds.), Fractography-Microscopic Cracking Processes, ASTM Spec. Tech. Publ. 600, 1977, p. 49. 11 R. O. Ritchie, Eng. Fract. Mech., 7 ( 1 9 7 5 ) 187. 12 R. O. Ritchie, Metall. Trans. A, 8 (1977)1131. 13 C. J. McMahon, C. L. Briant and S. K. Banerji, in D. M. R. Taplin (ed.), Fracture 1977, Proc. 4th Int. Conf. on Fracture, Waterloo, Ontario, June 19-24, 1977, Voh 1, University of Waterloo Press, Waterloo, Ontario, 1977, p. 363. 14 R. O° Ritchie, Met. Sci., 11 (1977) 368. 15 R. A. Oriani and P. H. Josephic, ActaMetall., 22 (1974) 1065.