Stress relief cracking in MnMoNi and MnMoNiCr pressure vessel steels

Materials Science and Engineering, 37 (1979) 179 - 186

179

© Elsevier Sequoia S.A., Lausanne - - P r i n t e d in the Netherlands

Stress Relief Cracking in MnMoNi and MnMoNiCr Pressure Vessel Steels C. J. M c M A H O N , Jr., R. J. DOBBS and D. H. G E N T N E R

Department of Metallurgy and Materials Science, University of Pennsylvania, Philadelphia, Pa. 191 74 (U.S.A.) (Received May 18, 1978; in revised form August 21, 1978)

SUMMARY

A constant displacement stress relaxation test at 615 °C on notched bars which were given a prior treatment to simulate a heataffected zone in a welded joint was used to characterize the susceptibility to stress relief cracking (SRC) of various heats of two types of commercial pressure vessel steels, SA 533-B and SA 508-2. The susceptibility of the former steel was quite variable, and that of the latter was uniformly high. It is shown that the SRC susceptibility of these steels can be rationalized in terms of their impurity and chromium concefftrations. The susceptibility increases linearly with a composition parameter which contains both impurity elements and chromium. The physical basis for this is discussed.

1. I N T R O D U C T I O N

Stress relief cracking, sometimes called " r e h e a t " cracking*, refers to the formation of intergranular cracks in the coarse-grained regions of the heat-affected zone (HAZ), or occasionally the weld metal, of a welded assembly when it is reheated to relieve residual stresses or when it is put in service at an elevated temperature [2]. The cracking accompanies the plastic relaxation of residual stresses by creep and is a manifestation of poor creep ductility. It tends to be found in alloy steels in which some measure of high temperature strength is imparted by the presence of strong carbide formers, such as chromium, molybdenum, vanadium, niobium, titanium or tantalum.

*Also called " u n d e r - c l a d " cracking w h e n it occurs under a weld overlay, usually o f stainless steel, used to clad a steel c o m p o n e n t [ 1 ].

In addition to the factors already mentioned, the tendency toward stress relief cracking (SRC) increases with the degree of mechanical constraint imposed on the weld region, and it has recently been shown that the SRC tendency is related to the purity of the steel [3]. In a study of MnMoNi pressure vessel steels of the SA 533-B type, Brear and King [3] showed that a high purity heat (prepared by vacuum induction melting) was n o t susceptible to SRC. They reported that the cracking tendency increased with the value of the following impurity parameter (compositions in weight per cent): 0.20[Cu] + 0.44[S] + 1.0[P] + 1.8[As] + + 1.9[Sn] + 2.7[Sb] This parameter was obtained from experiments on laboratory heats of fixed base composition in which the impurity contents were systematically varied. The susceptibility to SRC has been evaluated by a variety of tests [2], which can be placed in two general categories: (1) those involving stress relief of welded specimens, and (2) those carried out on either tensile or bend specimens in which the condition of the HAZ is simulated by a short-time treatment at a high temperature (above about 1300 °C) followed by rapid cooling. The former m e t h o d has the advantage of unquestioned authenticity, but with it one lacks knowledge of the stress levels involved, the type of microstructure in the most highly stressed regions and the exact thermal history of these regions. Hence, simulated-HAZ specimens are best suited for studies of SRC in which reproducibility is desired. The Charpy-type three-point bend specimen is an attractive candidate as a simulatedHAZ specimen because it offers a fairly high degree of constraint and is convenient to make and use. In fact, it was applied to some of the earliest studies of SRC by Christoffel

180 TABLE 1 Compositions of steels Element

SA 533-B

CAB C (wt.%) Mn (wt.%) Mo (wt.%) Ni (wt.%) Cr (wt.%) V (wt.%) W (wt.%) Si (wt.%) Al (wt.%) N (wt.%) Cu (wt.%)

P (ppm) Sn (ppm) As (ppm) Sb (ppm) S (ppm)

0.25 1.41 0.49 0.46 0.11 0.003 <0.01 0.26 0.016 0.008 0.12

80 80 120 13 140

SA 508-2

CBB 0.21 1.45 0.64 0.55 0.05 0.003 <0.01 0.23 0.029 0.014 0.13

60 150 350 15 90

CDB

CEB

0.22 1.37 0.60 0.59 0.18 0.004 <0.01 0.23 0.035 0.008 0.12

90 80 110 18 150

A

0.21 1.35 0.48 0.55 0.03 0.003 <0.01 0.25 0.021 0.009 0.12

90 10 100 14 210

[4] in type 347 stainless steel. Recently, a constant displacement bend test apparatus was introduced b y Miller and Batte [5] and applied to the problem of SRC in CrMo and CrMoV steels. In their test, grain-coarsened specimens are bent to various initial notchopening displacements (NODs) in the apparatus (the thermal expansion coefficient of which is approximately that of the specimen material) and the whole assembly is given a stress relief treatment by a slow heating and cooling cycle through the temperature range of interest. The specimens are broken open at 77 K, and the extent of cracking below the notch is measured. These values are plotted against the imposed NOD as shown schematically in Fig. 1 in which the susceptibility to SRC is seen to decrease in the order A, B, C. The purposes of the present research are to understand the factors which lead to SRC in the MnMoNi (SA 533-B) and MnMoNiCr (8A 508-2) pressure vessel steels, to develop methods to minimize its incidence and to develop a test method capable of discriminating among heats of varying degrees of susceptibility. After experimentation with other methods, the m e t h o d of Miller and Batte was. adopted and it proved to be well suited to the above purposes. This paper describes the results obtained from a comparison of five commercial heats of SA 533-B and three of SA 508-2. It is shown that the SRC susceptibility

BAB

0.17 1.5 0.56 0.58 0.15 <0.01 0.03 0.26 0.007 0.18 100 150 650 30 90

BBB

0.213 0.64 0.57 0.56 0.36 0.004 <0.012 0.26 0.001 0.002 0.04

0.223 0.64 0.58 0.63 0.34 0.022 <0.012 0.28 0.003 0.005 0.02

50 230 20 40 110

70 40 30 40 120

BCB 0.212 0.63 0.62 0.55 0.38 0.002 <0.012 0.26 0.001 0.002 0.04 60 110 40 40 120

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Fig. I. Schematic plot of the percentage brittle cracking vs. N O D for Charpy specimens given a constant displacement stress relaxation test,showing susceptibility to S R C decreasing in the order A, B, C.

can be accurately predicted by means of a comparison parameter which takes both impurity elements and chromium content into account.

2. E X P E R I M E N T A L

PROCEDURE

The compositions of the steels tested are listed in Table 1. The samples were all taken from production heats and all except heat A were removed from fracture toughness specimens used in previous work [6, 7]. Heat A was one tested earlier by Brear and King [ 3 ]. After the Charpy blanks were prepared, they were given a HAZ-simulation treatment consisting of heating in a muffle furnace to

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1300 °C, holding for 10 s at t h a t temperature and air cooling. The time elapsed from the start of heating to the start of cooling was about 4 min. This treatment produced a bainitic microstructure, as shown in Fig. 2, with a prior austenite grain size of ASTM no. 3 and hardness and yield strength values as shown in Table 2. The blanks were machined into Charpy specimens after this treatment; approximately 0.025 in was ground from each lateral surface in this step. The test equipment is illustrated in Fig. 3. The specimens were loaded in three-point bending to a chosen NOD as measured with a standard COD gauge attached to a set of knife edges clamped to the specimen. The gauge was then removed and the apparatus was sealed

(by welding) in a stainless steel container along with a small piece of titanium sponge which acted as a scavenger of active gases. The container was then placed in a circulating-air furnace and given the stress relief cycle shown in Fig. 4. After the stress relief treatment the specimen was removed and fractured at 77 K; the cracking which occurred during stress relief was readily apparent, as shown in Fig. 5. The fraction of the total area cracked was determined by photography of the fracture surface and measurement of the cracked area with the used of a planimeter. Scanning electron microscopy (SEM) was used for examination of the fracture surfaces,. and a specimen of heat A was fractured in ultrahigh vacuum (UHV) in a scanning Auger microprobe so t h a t the intergranular fracture facets could be analysed by Auger electron spectroscopy (AES).

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Fig. 5. Appearance of stress-relief cracking in an unbroken specimen and of the SRC portion of the fracture surface (discolored region).

182 TABLE 2 Properties of steels Heat

Average grain diameter ( m m )

Hardness R B

Yield strength (0.02%, lbf in -2 )

SA 533-B CAB CBB CDB CEB A

0.48 0.40 0.40 0.48 0.40

94 99 100 99 95

74300 80700 85000 74200 76000

SA 508-2 BAB BBB BCB

0.45 0.48 0.45

95 100 100

90000 75700 76000

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3. RESULTS

The plots of the degree of SRC as a function of NOD for the SA 508-2 and SA 533-B steels are shown in Fig. 6. It is apparent that there is a significant heat-to-heat variation in the SRC susceptibility in the SA 533-B steel, whereas the three heats of SA 508-2 show uniformly low resistance to SRC. In several cases a wide range of NOD was explored and it was found that the degree of

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SRC passed through a maximum. In two SA 508-2 heats the curve is composed of two branches. The reason for this behavior has not yet been determined; bend tests carried out with an Instron machine on similar specimens did not indicate any anomalies which might explain this behavior. It was decided that the first branch of the curve would be used to characterize the SRC resistance. Comparison of the behavior of the two types of steel indicates that the SA 508-2 is considerably more susceptible to SRC than the SA 533-B, although heat A of the latter group does approach the former group in susceptibility. This difference is not explicable on the basis of differences in grain size or strength, as indicated by Table 2 and Fig. 7, which shows the NOD for 15% cracking plotted against the yield strength.

183 Examination of the fracture surfaces by SEM indicated that they were almost entirely intergranular and that the facets were quite smooth, as indicated in Fig. 8. This observation does not preclude the possibility that the cracks were formed by the linking-up of cavities along grain boundaries since the smoothness of the intergranular facets could have arisen from surface diffusion.

ent in excess amounts in the grain boundaries (owing to segregation). The elements found to have segregated in this specimen are phosphorus, sulphur, nitrogen, tin, antimony and copper. All except the last are well-known embrittling elements in steels. It is interesting to note that, although this heat contained a large amount of arsenic (about 650 ppm), no arsenic was found to have segregated to the grain boundaries. This m a y simply have been due to the absence of segregation sites after the other elements had segregated. 533-B HEAT A SA r M°

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4. DISCUSSION

(b) Fig. 8. Typical intergranular fracture facets exposed by SRC: (a) 100×; (b) 1000x. A grain-coarsened specimen of heat A was fractured at 100 K in UHV in the scanning Auger microprobe and the resulting fracture surface was partly intergranular. The results of Auger analysis of an intergranular area are given in Fig. 9; spectra taken before and after extensive ion sputtering are shown. Peaks which are present before sputtering and absent afterwards are due to elements which are pres-

The test method adapted from Miller and Batte [5] has been found suitable for the characterization of susceptibility to stress relief cracking in the MnMoNi and MnMoNiCr pressure vessel steels. The results are consistent with earlier reports [1, 8] that SA 508-2 is considerably more sensitive to SRC than is SA 533-B. The limited AES observations of segregation of embrittling elements to intergranular fracture facets are also consistent with the findings of Brear and King [3] that these impurities are involved in the phenomenon. However, a plot of our SRC parameter (NOD for 15% cracking) v e r s u s the CERL impurity parameter, as shown in Fig. 10(a), indicates t h a t the impurity content alone cannot rationalize the behavior of these steels.

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improvement. From this we conclude that the Nakamura and Ito parameters overemphasize the roles of the strong carbide formers molybdenum, vanadium etc. in this particular appli-

185 cation. Figures 11(d) and 11(e) show the effect of using the chromium content alone compared with the CERL parameter + [Cr]. The latter clearly shows a better correlation, b u t it is apparent that the chromium is an important factor. From consideration of the thermal histories of the present steels in the as-received conditions, involving h o t working and later heat treatment, we expect that the strong carbideforming elements had been tied up as carbides prior to the present HAZ-simulation treatment. Since the latter was so rapid, it is reasonable to postulate that only the least stable alloy carbides (those of iron and chromium) would dissolve and that the more stable carbides of molybdenum, vanadium etc. would remain undissolved in the short-time excursion to 1300 °C [11]. Hence, only chromium would be available in sufficient amounts to retard stress relaxation by interaction with carbon to impede dislocation creep. Thus it is n o t at all surprising to find the type of correlation shown in Figs. 10(b) and l l ( e ) . If the time at 1300 °C had been longer, say 10 min or more rather than 4 s, then the influences of m o l y b d e n u m , vanadium etc. would more likely be felt and the present parameter would have to be modified accordingly. Thus the appropriate parameter to use in any given situation depends on the details of the thermal history of the HAZ; if the final treatment puts m o l y b d e n u m , vanadium etc. in solution, then the appropriate parameter would be, perhaps, CERL plus Ito. Finally, we should note that the reason that the importance of the chromium content did n o t reveal itself in the CERL work [3] is that they used model heats of constant chromium content to obtain their parameter. It was necessary to do this so that the impurity parameter could be determined unambiguously. In support of the present analysis we note the following: (1) a statistical analysis of 34 heats of SA 508-2 has indicated [1] that heats with high SRC susceptibility tend to have high chromium and low manganese contents (within the specification range); (2) recent work by Kanazawa e t al. [12] has shown that a steel with 0.8% manganese and 0.6% chromium (with 0.07% vanadium) was considerably more susceptible to SRC than an otherwise similar steel with 1.5%

manganese and no chromium (but with 0.02 0.03% niobium). The strengthening effect of chromium was illustrated nicely by Kanazawa e t al. [12]. The high chromium, low manganese steel relaxed more slowly at 600 °C than did the high manganese, low chromium steels. All other things being equal, the lower the rate of stress relaxation, the greater should be the tendency for intergranular cracking, since the rates of both nucleation and growth of grain boundary cavities increase with the level of the stress c o m p o n e n t normal to those grain boundaries [13]. The alloy elements in a steel to be welded can play other roles in addition to the creepstrengthening effect. They can alter the behavior of the embrittling elements and enhance or retard segregation [14, 15]. One effect of varying the Mn :Cr ratio would probably be to vary the ratio MnS:CrS. Since CrS dissolves at lower temperatures in steels than does MnS [16] and since sulphur is highly surface active, a reduction in the Mn:Cr ratio could contribute to sulphur-induced intergranular void growth. In this context, it should be noted that a tendency for hot cracking (intergranular liquation due to the low melting point of FeS) has been associated with the tendency for SRC in SA 508-2 steel [1]. The impurity content of the steel, which has been shown to be an essential part of the SRC phenomenon [3], influences the rates of both nucleation and growth of intergranular cavities. Impurities which segregate to carbide/ ferrite interfaces can lower their cohesive strength [17] or their surface energy, and this would greatly accelerate the rate of cavity nucleation on carbides [13]. Impurities which segregate to the free surfaces of cavities, lowering their surface free energy, cause the cavity shape to become less nearly spherical and more like a flat lens [18] ; this increases the rate of diffusion-controlled cavity growth. Finally, since the forging-grade steel SA 508-2 is so much more susceptible to SRC than is the plate-grade steel SA 533-B and since there is a forging-grade steel SA 508-3 with essentially the same composition (and low SRC susceptibility) as SA 533-B and the same mechanical properties as SA 508-2, one might ask w h y the latter is used in critical pressure vessels in the U.S.A. This is especially puzzling since European practice calls for crit-

186

ical pressure vessel forgings to be made from SA 508-3 because of its superior resistance to SRC. It would appear that a re-evaluation of U.S.A. practice in this regard is called for.

ACKNOWLEDGMENT

This research is supported by the Electric Power Research Institute under Contract RP559-1 and was carried o u t in the facilities of the LRSM, University o f Pennsylvania, supported by the National Science Foundation MRL program, Grant no. DMR-7C-80994.

REFERENCES 1 A. G. Vinckier and A. W. Pense, Weld. Res. Counc. Bull., 197 (Aug. 1974). 2 C. F. Meitzner, Weld. Res. Counc. Bull., 211 (Nov. 1975). 3 J. M. Brear and B. L. King, Proc. Conf. on Grain Boundaries, Metals Society, London, 1976. 4 R. J. Christoffel, Weld. J. (London), 41 (1962) 2515.

5 R. C. Miller and A. D. Batte, Met. Constr. Br. Weld. J., 7 (Nov. 1975) 550. 6 T. U. Marston, Combustion Engineering Inc., Rep. to EPRI on Contract No. RP-232-2-0, May 1, 1975. 7 Babcock and Wilcox Co., Rep. to EPRI on Contract No. RP 232-3. 8 K. Kussmaul and J. Ewald, 3rd Int. Conf. on Pressure Vessel Technology, ASME, New York, p. 627. 9 H. Nakamura, T. Naiki and H. Okabayashi, Proc. 1st Int. Conf. on Fracture, 1965, Vol. 2, p. 863. 10 Y. Ito and M. Nakanishi, IIW Document X-66872. 11 R. M. Brick and A. Phillips, Structure and Properties of Alloys, 2nd edn., McGraw-Hill, New York, 1949, p. 265. 12 S. Kanazawa, K. Yamoto, T. Takeda and K. Hashimoto, J. Jpn Weld. Soc., 44 (1975) 791. 13 R. Raj and M. F. Ashby, Acta Metall., 6 (1975) 653. 14 M. Guttmann, Surf. Sci., 53 (1975) 213. 15 C. J. McMahon, Jr., and L. Marchut, J. Vac. Sci. Technol., 15 (1978) 450. 16 B. J. Schulz and C. J. McMahon, Jr., Metall. Trans., 4 (1973) 2485. 17 J. R. Rellick and C. J. McMahon, Jr., Metall. Trans., 5 (1974) 2439. 18 C. J. McMahon, Jr., in J. L. Walter, J. H. Westbrook and D. A. Woodford (eds.), Grain Boundaries in Engineering Materials, 4th Bolton Landing Conf., Claitors Press, Baton Rouge, La., 1974, p. 252.