Mechanisms of intergranular fracture in alloy steels

Mechanisms of Intergranular Fracture in Alloy Stccls C. J. M c M a h o n , Jr. Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia, PA 19104-6272

aries in quenched and tempered steels that have been heated in, or cooled slowly through the 300-600°C temperature range. It is reversible; reheating an embrittled steel above 600°C and rapidly cooling removes most of the embrittlement. It is now known to be associated with the segregation of certain impurities to grain boundaries; many fracture studies coupled with analyses of fracture surfaces by Auger electron spectroscopy (AES) have shown that the degree of embittlement increases with the intergranular impurity concentration. This has been documented by means of the time-honored Charpy impact test, as well as by measurements of the local stress needed to start fracture in a notched specimen. Figure 1 gives an example of the effects of three of the most important impurities in alloy steels.

INTRODUCTION

Intergranular fractures are generally macroscopically brittle in that they usually occur with little overall plastic deformation. In some cases, the fractures are not brittle on the microscopic scale, but instead result from rupture along grain boundaries that contain a weak or weakly bonded second phase. Examples of this include some precipitation-hardened aluminum alloys [1], overheated steels [2], and steels with precipitated carbides or nitrides [3, 4]. Here, we are interested not in intergranular rupture by the displacement-controlled accumulation of damage, but rather in the stress~controlled failure of cohesion that can occur under a number of different conditions, ranging from the catastrophically rapid to the quasi-statically slow. Since the microscopic fracture appearance is often similar for diverse forms of intergranular fracture, one cannot deduce the fracture mechanism merely from the appearance. The purpose here, then, is to describe various mechanisms having to do with steels so that one can make better sense out of a particular instance of such fracture, once the probable conditions are known. These various mechanisms will be discussed in turn.

THE MAIN EMBRITTLING ELEMENTS: Sb, Sn, P, AND Mn

As shown in Fig. 1, antimony and tin are much more powerful embrittling elements than phosphorus. However, they are only effective in steels that contain several percent of nickel, which is believed to enhance the intergranular segregation of these elements by virtue of an attractive interaction between the nickel and the impurity in an iron matrix. This effect has been rationalized by the Guttmann regular-solution model [5, 6]. In the absence of nickel, e.g., in a 241 Cr-1 Mo steel, these elements do not cause temper embrittlement even at a level of several hundred parts per million [7].

TEMPER EMBRITTLEMENT

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Table I Comparison of Mn and P Effects, Separate and Combined Effect

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Intergranular Fracture in Alloy Steels

fore, should be treated as individual embrittling elements. The ultimate explanation of why manganese acts as an embrittling element in steels will pose a real challenge for theorists. One does not have too much difficulty accepting that segregated nonmetallic impurities can disrupt bonding along grain boundaries, but manganese is closely related to iron chemically, and there does not appear to be even the beginnings of a theoretical explanation of its embrittling effect. THE STRENGTHENING EFFECT OF CARBON

In the context of temper embrittlement, the profound effect of carbon in strengthening grain boundaries in iron has been generally overlooked until the last few years. It has been known for a long time that intergranular fracture of iron could be induced at low temperatures by decarburization [10] or even by solution treating and quenching [11]. In a study of F e - C r - P alloys, Erhart and Grabke [12] drew attention to the importance of the scavenging of carbon by chromium, which promoted phosphorus segregation and intergranular embrittlement in their alloys. They attributed this mainly to the removal of the site competition between phosphorus and carbon at the grain boundaries. However, carbon segregated in preference to phosphorus would also make its inherent contribution to intergranular strength. This effect can be seen clearly in Fig. 3; removing the residual carbon from the Fe-P alloy having segregated phosphorus at a level corresponding to a P12o/Fe705 Auger peak height ratio of about 30% reduces the critical fracture stress from 1,200 to about 800 MPa. The strengthening effect of carbon has not been treated theoretically by atomistic calculations, presumably because no one has yet found a way to describe the electronic state of a carbon atom in bcc iron. However, the effect is likely to involve an increase in the cohesive energy of the iron grain boundary. The other conceivable

271 way in which a segregated impurity could affect intergranular strength would be to change the tendency for dislocations to be emitted from the tip of an intergranular crack. Inhibiting dislocation emission would promote intergranular decohesion. Since carbon is very effective in pinning dislocations in bcc metals, it could hardly be expected to have the opposite effect on dislocation sources in grain boundaries. Therefore, one should look to the explanation of the carbon effect in some kind of strong bonding that acts across the grain boundary and that does not disrupt the bonding between iron atoms in the boundary and those further away. This might be achieved if the Fe-C bonding were mainly covalent, rather than ionic, for example. The theoretical understanding of how embrittling elements work is far from well developed. Two sets of calculations, one on unrelaxed atomic clusters [13] and the other of a model grain boundary [14], have both suggested that bonding between embrittling elements and the matrix atoms involves charge transfer that weakens the bonding between matrix atoms that do not interact directly with the impurity. The latter calculation suggested that the impurity atoms, in this case phosphorus in iron, can be spread over several atomic planes along a boundary. This has received support from recent work [15] on Fe-3%Si bicrystals with a symmetrical ~ = 5 (013) boundary that were doped with phosphorus. In this study it was found that the intergranular phosphorus distribution had enough width that the fracture could travel to one side of the center of it, resulting in a 2:1 ratio of phosphorus on the mating fracture surfaces, as shown in Fig. 4. The conclusion was that the phosphorus was most probably distributed over five atomic planes, which, in fact, is the number of planes needed to define the structural units which make up this boundary, as shown in Fig. 5. The concentration of phosphorus which resulted in the fracture of these boundaries was about 20 at. ~ , i.e., about one phosphorus atom per structural unit shown in Fig. 5.

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THE ROLE OF MOLYBDENUM

Since molybdenum interacts with carbon in iron even more strongly than does chromium, and since chromium is believed, as related above, to enhance the embrittling effect of phosphorus through scavenging of carbon, one might expect that molybdenum would be even more detrimental to a steel with phosphorus. However, just the reverse is true. Molybdenum has been shown to inhibit temper embrittlement, not only in steels containing phosphorus, but also in steels containing the other known embrittling elements, including manganese [16], A recent study of F e - M o - P alloys [17] has led to the conclusion that the beneficial effect of molybdenum results from an enhancement of carbon segregation. This can be understood by considering a McLeantype segregation isotherm: Q Xl, exp R-T Xi-

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value of Xb is reached. Therefore, one could postulate that the effect of chromium on carbon is mainly to reduce Xb without changing Q significantly but that, with molybdenum, the increase in Q dominates. The effect on Q would, of course, be rationalized in the framework of the regularsolution model [5]. This effect would only operate as long as there is enough molybdenum in solid solution to cosegregate to grain boundaries with carbon. Several long-time aging studies [16, 18, 19] have shown that the beneficial effect of molybdenum ultimately disappears when the molybdenum has been precipitated in alloy carbides. One strength of the molybdenum-carbon hypothesis is that it would be equally valid for any impurity, which is consistent with the experimental observations.

HYDROGEN-INDUCED

INTERGRANULAR

CRACKING

Hydrogen-induced fracture in heat-treated alloy steels is generally intergranular. This is true whatever the source of hydrogen, whether from a gaseous atmosphere or from the cathodic part of an electrochemical reaction. A long series of experiments [20-23] on several types of alloy steel has shown that the tendency for intergranular fracture in hydrogen is directly related to the presence of embrittling impurities in the grain boundaries. For a given yield strength and hydrogen activity, as the intergranular impurity concentration is reduced, the hydrogen-induced fracture shifts from a stress-controlled intergranular decohesion to a displacement-controlled transgranular fracture, often referred to as quasi-cleavage. The hydrogen effect can best be understood by considering experiments [23] done on notched specimens of an impurity-doped steel stressed in a n H2 atmosphere. Because of plastic flow at the tip of the notch, the maximum principal (tensile) stress and the maximum hydrostatic stress lie at a distance ahead of the notch

274

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ferent aging treatments so as to vary the intergranular concentration of antimony. The loading rate was slow enough to approximate static loading, as far as the diffusion rate of hydrogen is concerned. It is apparent that the stress necessary for intergranular cracking is reduced by about 40% by the hydrogen. It was found possible to detect the first hydrogen-induced intergranular microcracks by monitoring the acoustic emissions [23]. In this way the threshold stress CrTH(i.e., the maximum tensile stress ahead of the notch, obtained from a finite-element analysis) for the onset of hydrogeninduced cracking could be obtained as a function of the maximum antimony concentration that is likely to be found in the highly stressed region. The latter is obtained from a statistical analysis of the results of AES determinations of antimony concentrations on a large number of grain boundaries after the particular aging time and knowing the grain size and the volume of the highly stressed region. The results for three impurities in the Ni-Cr steel are shown in Fig. 7. Also shown here is the value of the stress at which fast fracture occurs in the H 2 , designated as (rH*. Once CrTH is reached, microcracking continues, and the patches of isolated intergranular decohesion become larger, until there is

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tip which is proportional to the notch radius. Experiments have shown that, after some delay time, during which the hydrogen diffuses from the notch tip to the region of maximum stress, intergranular cracking begins at isolated locations, which are presumably the grain boundary facets with the greatest impurity concentration. The result of this process is illustrated by the load-displacement curves shown in Fig. 6 for an Sb-doped steel given two dif-

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Intergranular Fracture in Alloy Steels

275

enough local stress elevation to sustain unstable fracture. Presumably, this latter process is not dependent on the presence of hydrogen, because it occurs too rapidly for hydrogen diffusion to keep up. Therefore, the curves for the critical fracture stress in air (top curves in Fig. 7) can be interpreted as the stress needed for the initiation of microcracking under the influence of impurities alone, the propagation stress already having been surpassed. Figure 8 gives examples of (a) an isolated microcrack found (after opening the spec- (a) imen by fatigue cracking) ahead of the notch just after the first acoustic emission, and (b) and (c) patches of intergranular fracture found in a precracked specimen well below and just above the first detectible crack extension. Experiments showed that the apparent threshold stress intensity K;;, for the onset of cracking in precracked specimens in hydrogen represents only the accumulation of a detectible amount of damage in the form of patches of intergranular cracking. Thus, the values of CrTH (b) measured in notched specimens have a well-defined physical significance, but the K,, values measured in precracked specimens do not. The effect of impurity-induced embrittlement of the fracture mode in hydrogen can be seen most clearly in the results of experiments on HY130-type steel [21]. Here, the principal embrittling elements were manganese and silicon. (Because they were added to the steel intentionally, we cannot call them impurities.) As shown in Fig. 9, the heat with manganese and silicon is highly susceptible to both temper embrittlement and to hydrogen-induced cracking, whereas the heat containing neither of these elements is essentially immune to temper embrittlement and suffers much less loss of toughness in hydrogen. (c) The mode of cracking in hydrogen in a preFIc. 8. (a) Scanning electron micrograph of Sb-doped steel that was loaded in H2 just to CTH, then unloaded cracked specimen, shown in Fig. 10, is apand fatigued in air. One of two isolated facet-size inparently transgranular in the latter heat, tergranular microcracks ahead of notch is shown. (b, but largely intergranular in the former. Exc): Hydrogen-induced intergranular cracking adjacent amination of the crack-tip region after the to precrack tip in Sb-doped steel aged 10 h. Specimens onset of cracking showed that, while the were unloaded and fatigued in air after loading to (b) 75% of "pop-in" load and (c) just after "pop-in" load. embrittled specimen began cracking in the

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expected Mode I, the unembrittled specimen began with a bifurcated type of Mode II cracking, as shown schematically in Fig. 11. The morphology of initial hydrogen-induced cracking in a notched specimen of the unembrittled steel is shown in Fig. 12; this emphasizes that the cracking occurs along planes of high shear stress. The fracture surface is clearly transgranular with respect to the prior austenitic microstructure, and, as shown in Fig. 13, it is similar in some respects to low-temperature cleavage. This fracture mode has been interpreted [21, 24] as comprising a combination of actual cleavage along slip planes of martensitic laths and decohesion along lath boundaries. The slip-plane decohesion could result from the blockage of dis-

FIG. 10. Modes of hydrogen-induced cracking in (a) heat A and (b) heat F, aged for 1,000 h [21]. Reprinted with permission from Metallurgical TransactionsA, Vol. 12A, Y. Takeda and C.J. McMahon, 1981, ASM International.

locations that carry hydrogen in their cores [24]. A possible rationalization of this type of fracture is shown schematically in Fig. 14. Because it requires the transport of hydrogen by dislocations, it is necessarily a strain-controlled phenomenon.

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Intergranular Fracture in Alloy Steels

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The pronounced effect of manganese and silicon was also found in 4340-type steels at yield-strength levels in the 150270 ksi range; an example is shown in Fig. 15 for precracked samples loaded in hydrogen gas. The drop in Kth with increasing yield stress is far more gradual for the steel with no manganese or silicon, and the fracture mode in this steel does not become mainly (stress-controlled) intergranular fracture until near the upper end of the yield stress range. The yield stress is important for two reasons: First, in an elasticplastic specimen the maximum hydrostatic tension ahead of a precrack or notch is pro-

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FIG. 13. Comparison of (a) plasticity-related HIC and (b) low-temperature cleavage on fracture surface shown in Fig. 12(b) [21]. Reprinted with permission from Metallury, ical Transactions A, Vol. 12A, Y. Takeda and C.J. McMahon, 1981, ASM International.

portional to the yield stress. It is about 2.4 times the yield stress in the elastic-plastic stress analysis of Rice and Johnson [25]. The hydrostatic tension, m,, governs the equilibrium concentration of hydrogen, CH, in the metal lattice through the relation O'hWm

CH = Coexp RT where Co is the equilibrium lattice concentration in the absence of applied stress (proportional to the square root of the partial pressure of hydrogen for a gaseous at-

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Intergranular Fracture in Alloy Steels

279

Table 2 Compositions of 4340-Type Steels Steel

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mosphere), and V,,, is the molar volume of hydrogen in the lattice. The reasons why this relation is so important for hydrogen are that the atomic volume of hydrogen in iron, as in most metals, is quite large (i.e., hydrogen has a large tendency to expand the crystal lattice), and hydrogen is also highly mobile, so that the equilibrium concentration is readily obtained. Since the reduction in cohesion is presumably a function of the hydrogen concentration, it is very sensitive to the yield strength. The second reason for the importance of the yield stress is that it also governs the maximum principal stress ahead of the precrack. Various stress analyses put the proportionality factor in the range 2.4-3. Obviously, the higher the maximum tensile stress, the more likely is a stress-controlled mode of fracture. The measured values of Kth can also be plotted against the calculated hydrogen concentration, as shown in Fig. 16; here, the yield stress and hydrogen pressure were varied independently to achieve the range of hydrogen concentrations shown. The difference wrought by the presence of manganese and silicon is again clear, but this plot emphasizes the role of yield stress in determining the hydrogen concentration ahead of the precrack. All these data can be fitted onto one curve, as shown in Fig. 17, when an impurity parameter is added to the hydrogen concentration. Here, the manganese and silicon concentrations are the only significant components of this parameter, as indicated by Table 2. The point to be made by this plot is that it is reasonable to treat hydrogen as one of a number of elements that can promote intergranular decohesion. It appears that the only fundamental difference between hydrogen and the other embrittling

elements in its extremely high mobility at room temperature.

TEMPERED MARTENSITE EMBRITTLEMENT

Tempered martensite embrittlement refers to the minimum in fracture energy as a function of tempering temperature in martensitic low-alloy steels that normally occurs in the vicinity of 350°C. In 4340-type steels, as well as many other commercial low-alloy steels, this is an impurity-related phenomenon [26-28] that is, again, associated with intergranular fracture [27,29]. It can be illustrated very well by the two steels discussed in the previous section, as shown in Fig. 18. This is actually a reflection of an impurity-induced increase in the ductile-brittle transition temperature, as shown in Fig. 19. The main difference between tempered martensite embrittlement and temper embrittlement is that the former occurs in

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Fro. 18. C h a r p y fracture e n e r g y as a function of tempering t e m p e r a t u r e for a v a c u u m m e l t e d h i g h - p u r i t y 4340-steel with (B6) a n d w i t h o u t (B7) m a n g a n e s e a n d silicon [281.

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hibit t e m p e r embrittlement are first tempered, before the embrittling treatment (i.e., aging or slow cooling) is applied. Because the yield stress is so high, the a m o u n t s of impurity segregation required for intergranular fracture are m u c h lower than for t e m p e r embrittlement. Also, the a m o u n t s of intergranular fracture necessary to p r o d u c e the t o u g h n e s s m i n i m u m are relatively small, as shown, e.g., in Fig. 20. At the high stresses involved here, e v e n isolated intergranular cracks can produce p r e m a t u r e fracture, because they can act as the nucleators of cleavage fracture. In the absence of these intergranular microcracks, the expected m o d e of fracture at r o o m t e m p e r a t u r e w o u l d have been rupture. The TME p h e n o m e n o n can also be studied by measuring the critical stress, or*, for

FIG. 20. Portion of fracture surface of steel with manganese and silicon (B6) tempered at 350°C and fractured at room temperature [28]. Reprinted with permission from Metallurgical Transactions A, Vol. 14A, N. Bandyopadhyay and C.J. McMahon, 1983, ASM International.

Intergranular Fracture in Alloy Steels

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unstable fracture in a notched bar tested at a temperature where the fracture is entirely stress controlled (i.e., either intergranular or cleavage), as depicted in Fig. 21. This test is obviously much more sensitive than the Charpy impact test. The dependence of ~r* on the impurity parameter

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used earlier is quite dramatic, as shown in Fig. 22. Even in the high-purity steel, the minimum in ¢* at 77K is associated with a peak in the amount of intergranular fracture, as shown in Fig. 23. Our study of these steels [28] confirmed what had been reported long before [30], that the toughness minimum is also associated with an increase in the amount of carbide platelets that form on the prior austenitic grain boundaries during tempering. Our interpretation of this phenomenon is that the segregated embrittling elements involved are inherited from the austenite phase and that their presence in the grain boundaries becomes more and more harmful as the unyielding carbide platelets are precipitated there. There have been reported other ways, not associated with intergranular fracture, in which a toughness minimum can occur [31-34]. These involve the transformation of retained austenite and the formation of carbide platelets between martensite laths. The importance of these other effects in

282

C. 1. McMahon, Jr.

commercial steels is unknown. The importance of the intergranular mode of TME is, however, well established.

STRESS-RELIEF CRACKING

The intergranular fracture modes discussed up to this point have all involved unstable, i.e., rapid, fracture. Even the hydrogen-induced microcrack formation was rapid, as evidenced by the acoustic emissions produced. There exists a stress-controlled, brittle mode of cracking in steels at elevated temperatures that is diffusion controlled and that proceeds at rates of the order of a few micrometers per second. There is evidence that this actually occurs in rapid bursts of crack extensions on the order of a few micrometers, so the difference between this and hydrogen-induced cracking may be one mainly of scale. However, the phenomenon appears to be quite different on the macroscale, because it permits quasi-static brittle crack growth over distances of the order of centimeters at stress levels far below those needed for unstable fracture. The brittle mode of high-temperature intergranular cracking in steels is actually usually mixed with the normal, displacement-controlled mode of high-temperature failure: intergranular cavitation. It is the result of a special kind of impurity segregation, and it occurs above a threshold stress and below a limiting strain rate; out-

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FIG. 25. Example of a microcrack produced by cavitation ahead of notch tip in specimen loaded at fixed displacement at 600°C [35]. Reprinted with permission from Acta Metallurgica, Vol. 32, J. Shin and C.J. McMahon, 1984, Pergamon Press PLC.

side these boundaries, cracking occurs by cavitation. The phenomenon has been studied recently in the context of stress-relief cracking (SRC), which is the damage that can occur in the heat-affected zone of weldments in alloy steels. In these studies a steel specimen is heated briefly to a high temperature, usually above 1,200°C, and cooled fairly rapidly to simulate the thermal cycle in a HAZ. This results not only in grain coarsening, but, more importantly, in the dissolution of sulfides and the segregation of sulfur to the austenitic grain boundaries. It was demonstrated by us [35] and later confirmed by others [36] that the brittle mode of cracking is caused by sulfur.

Intergranular Fracture in Alloy Steels

FIe. 26. (a) Example of a crack formed during 135min stress relaxation at 600°C. (b) Portion of one intergranular facet, showing that cracking mode is decohesion [35]. Reprinted with permission from Acta Metallurgica, Vol. 32, l- Shin and C.J. McMahon, 1984, Pergamon Press PLC.

The phenomenon of SRC can be illustrated by results on a MnMoNiCr pressure vessel steel (A508-2). When notched bars are loaded to a fixed displacement (to simulate a residual stress) at an elevated temperature, the stress relaxes with time, and cracking may occur, by a mode that depends on the initial stress. Figure 24 shows load-relaxation curves of a commercial heat of this steel, and Figs. 25 and 26 show the cracking mode of specimens whose load relaxation was interrupted at the points indicated in Fig. 24. In a specimen that started out at a low load, a small amount of microcracking by rupture occurred in the highly stressed region ahead of the notch tip (Fig. 25). The rupture cavities were nucleated at sulfide particles that formed on the austenite grain boundaries

283

FIG. 27. TEM replica from a fracture surface produced by decohesion at 600°C [35]. Reprinted with permission from Acta Metallurgica, Vol. 32, J. Shin and C.|. McMahon, 1984, Pergamon Press PLC.

during cooling from the HAZ-simulation treatment. When the initial load on a specimen was relatively high, cracking occurred rapidly by the brittle mode, as shown in Fig. 26. The absence of cavitation on large areas of this kind of fracture was confirmed by TEM replicas, an example of which is shown in Fig. 27, and AES on brittle microcracks that formed ahead of the notch and were later exposed by fracture in UHV has shown that the main embrittling impurity on the surface was sulfur [35, 36]. There is not yet agreement on the precise route taken by the sulfur [37, 38], but Auger and acoustic-emission [39] studies support the following hypothesis [35, 37]: Brittle microcracking begins with the formation of a cavity, presumably at a weak inclusion, like a sulfide. This cavity would tend to grow in the usual way: by diffusion

284

C. J. McMahon, Jr.

of atoms from the surface of the cavity into the surrounding grain boundary, driven by the work done by the tensile stress normal to the boundary as the boundary region thickens, owing to the "plating out" of the atoms from the cavity. (This is the reverse of sintering, which is driven by the surface tension of the cavity.) The diffusive growth of the cavity can turn into growth by decohesion because of the presence of sulfur on the surface of the cavity. (It has been found by many investigators using AES that sulfur is an ubiquitous impurity on cavity surfaces, which

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FIC. 29. Examples of microcracks formed at 600°C by the combined effects of decohesion (smooth areas) and cavitation [35]. Reprinted with permission from Acta Metallurgica, Vol. 32, J. Shin and C.J. McMahon, 1984, Pergamon Press PLC.

is expected, because it is a highly surfaceactive element in steels and because the intergranular sulfides provide a ready source.) Since a significant fraction, probably on the order of one-half, of the surface atoms of the cavity are sulfur, the diffusive growth must result in a high sulfur concentration in the grain boundary surrounding the cavity. The concentration profile would depend on temperature, the local tensile stress, and, of course, time. It is well established [40] that sulfur reduces the strength of grain boundaries in iron, so if the sulfur concentration becomes high enough for the instantaneous level of the tensile stress to cause decohesion, then the rim of the cavity should advance unstably through the region of high sulfur content.

285

lntergranular Fracture in Alloy Steels

(This is analogous to the process that is ascribed to h y d r o g e n - i n d u c e d cracking.) The cavity could grow into a microcrack by this process, and microcracks could coalesce into a macrocrack of the kind s h o w n in Fig. 26. Acoustic emissions have been detected during this kind of crack g r o w t h that w o u l d be consistent with discontinuous brittle crack g r o w t h in steps of several micrometers; evidence for such growth steps, in the form of striations on the intergranular facets, was f o u n d in the same work [39]. This process is s u m m a r i z e d in Fig. 28. In order for this m e c h a n i s m to operate, the tensile stress must be high e n o u g h to

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cause local sulfur-induced decohesion, but not so high as to force cavity growth by plasticity. In m a n y cases decohesion and plastic rupture occur simultaneously, and the resulting fracture surface reflects this, as s h o w n in Fig. 29. Bika [41] has recently s h o w n that the d o m i n a n t cracking m o d e shifts gradually from decohesion to rupture as the crack advances and the plastic strain in the crack tip region intensifies. Again, this is also found in hydrogen-induced cracking.

(a)

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(b)

FIG. 30. (a) Stress-corrosion cracks in longitudinal section of P-doped NiCrMoV steel in step-cooled condition, loaded to 75'/~ of UTS in 9M NaOH at 98°C. (b) Fracture surface of steel in quenched and tempered condition [441. Reprinted with permission from Metallurgical Transactions A, Vol. 12A, N. Bandyopadhyay and C.L. Briant, 1981, ASM International.

Intergranular cracking by localized anodic dissolution can occur in heat-treated, lowalloy steels and p r o d u c e a fracture surface that can closely resemble an intergranular fracture caused by decohesion. The impurity mainly responsible for this behavior appears to be p h o s p h o r u s [42], which is also the impurity responsible for the preferrential etching of grain boundaries in steels that are t e m p e r embrittled [43]. This type of cracking is fundamentally different from the types described heretofore; it has nothing to do with decohesion, and, as far as we know, it occurs in a continuous, stable manner. To e m p h a s i z e this distinction from the fracture types involving decohesion, it m a y be pointed out that the prob-

C. J. McMahon, Jr.

286

l e m is e x a c e r b a t e d b y the p r e s e n c e of m o l y b d e n u m in t h e steel [44]. T h e w o r k of B a n d y o p a d h y a y a n d Briant [44] h a s clarified t h e b o t h the role of steel c o m p o s i t i o n a n d t h e m e c h a n i s m of I G S C C of N i C r - b a s e d steels in a h o t caustic solution. Figure 30 s h o w s the i n t e r g r a n u l a r s t r e s s - c o r r o s i o n cracks a n d the r e s u l t i n g f r a c t u r e s u r f a c e in a s t e p - c o o l e d , P - d o p e d N i C r M o V steel, a n d Fig. 31 gives a c o m p a r i s o n of t h e stress v e r s u s failure t i m e c u r v e s for statically l o a d e d s m o o t h tensile s p e c i m e n s . T h e w o r s t case is the steel w i t h both molybdenum and segregated phosp h o r u s . T h e i r c o n c l u s i o n , b a s e d on their a n a l y s i s of the m o r p h o l o g y of the stressc o r r o s i o n cracks, w a s t h a t the p h o s p h o r u s a n d m o l y b d e n u m b o t h act b y f o r m i n g a p a s s i v a t i n g l a y e r o n the sides of a n interg r a n u l a r crack. T h a t is, t h e s e e l e m e n t s exa c e r b a t e a n existing t e n d e n c y for s o m e preferential intergranular dissolution (e.g., as in m e t a l l o g r a p h i c etching) a n d the t e n d e n c y for stress to c a u s e crack g r o w t h b y p l a s t i c i t y - i n d u c e d film r u p t u r e at the crack tip. O f c o u r s e , t h e r e are o t h e r m e c h a n i s m for I G S C C in steels, the m o s t f a m o u s b e i n g t h a t w h i c h o c c u r s in s e n s i t i z e d austenitic stainless steels. T h e r e a s o n for d i s c u s s i n g the caustic c r a c k i n g of r o t o r steels in the p r e s e n t c o n t e x t is the a p p a r e n t , albeit s p u rious, r e l a t i o n s h i p w i t h o t h e r k i n d s of intergranular embrittlement.

SUMMARY I n t e r g r a n u l a r f r a c t u r e of alloy steels c a n occur b y a v a r i e t y of m e c h a n i s m s , m o s t of them involving intergranular segregation of i m p u r i t i e s or alloy e l e m e n t s . T h e a p p e a r a n c e of t h e f r a c t u r e s u n d e r the micros c o p e is o f t e n i n d i s t i n g u i s h a b l e , o n e f r o m a n o t h e r . In o r d e r to e v a l u a t e the p r o b a b l e c a u s e of s u c h a fracture, it is n e c e s s a r y to h a v e s o m e a p p r e c i a t i o n of the c o n d i t i o n s t h a t lead to o n e t y p e or a n o t h e r . T h e disc u s s i o n s of the v a r i o u s t y p e s in the p r e s e n t article are i n t e n d e d to facilitate s u c h a n evaluation.

This article was prepared with the support of the U.S. D e p a r t m e n t of Energy under contract DE-FGO2-87ER45290.

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287

lntergranular Fracture in Alloy Steels

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35. J. Shin and C. J. McMahon, Jr., Acta. Metall. 32:1535 (1984). 36. C. A. Hippsley, H. Rauh and R. Bullough, Acta Metall. 32:1381 (1984). 37. C. J. McMahon, Jr., Z. fur Metallk. 75:496 (1984). 38. H. Rauh, C. A. Hippsley, and R. Bullough, Acta Metall. 37:269 (1989). 39. C. A. Hippsley, D. Buttle, and C. B. Scruby, Acta Metall. 36:441 (1988). 40. K. S. Shin and M. Meshii, Scripta Metall. 17:1121 (1983). 41. D. Bika, University of Pennsylvania 1990, unpublished research. 42. K. L. Moloznik, C. L. Briant, and C. J. McMahon, Jr., Corrosion 35:331 (1979). 43. A. H. Ucisik, C. J~ McMahon, Jr., and H. C. Feng, Metall. Trans. A 8A:321 (1978). 44. N. Bandyopadhyay and C. L. Briant, Metall. Trans. A 14A:2005 (1983). 45. Jun Kameda and C. J. McMahon, Jr., Metall. Trans. A 12A:31 Received October 1990; accepted November 1990.