Assessing Hydrogen-Assisted Cracking Fracture Modes in High-Strength Steel Weldments

Test results substantiate and extend the Beachem theory on hydrogen embrittlement

BY S. A. GEDEON AND T. W. EAGAR

ABSTRACT. The stress intensity that causes causes crack propagation as a function of then builds to very high values, which crack propagation in high-strength steel stress intensity for a specific material and adds to the applied external stresses. By weldments was quantified as a function of temperature. This relationship is then com­ applying Sievert's law, it is estimated that the hydrogen content at the crack loca­ pared to existing cracking mechanism the­ a steel with 5 ppm hydrogen would have tion. This relationship was used to assess ories. over 17,000 atmospheres pressure in the previously proposed theoretical hydro­ Previous literature concerning cracking voids at 20 °C (68 °F). However, several gen-assisted cracking mechanisms. In­ theories, stress intensity, determination experimental observations conflict with deed, it was found that the microplastic- and hydrogen content determination is this mechanism. Hydrogen embrittlement ity theory of Beachem can best describe briefly reviewed. can be eliminated by degassing even after how the stress intensity factor and hydro­ exposure to room temperature. The low gen content affect the modes of inter­ Hydrogen-Assisted Cracking Mechanism temperature of the degassing would not granular, quasi-cleavage and microvoid Theories be high enough to dissociate the diatomic coalescence fracture. hydrogen into monatomic hydrogen The results of theoretical studies of hy­ Implant test results were analyzed with which could diffuse out of the steel. Also, drogen embrittlement mechanisms pro­ the aid of fracture mechanics to deter­ the observation of hydrogen-induced posed by physical metallurgists have rarely mine the stress intensity associated with cracks growing on a free surface pre­ been applied to the field of welding. various modes of fracture. Diffusible weld cludes an internal pressure gradient as the Sawhill's study (Ref. 1) of HY-130 steel hydrogen results were analyzed with the driving force for crack growth. weldments, however, provides a good aid of a hydrogen distribution model de­ background for the ensuing analysis of the The adsorption theory of Petch and veloped by Coe and Chano to determine most often proposed hydrogen embrit­ Stables (Ref. 3) and further modifications the amount of hydrogen present at the tlement mechanisms. Even though the (Ref. 4) propose a lowering of the surface crack location at the time of fracture. problem of hydrogen embrittlement has free energy by hydrogen so that a crack The stress intensity and hydrogen con­ been studied extensively, no one theory can grow under a lower applied stress. tent responsible for the microvoid coales­ has become generally accepted. This theory has been criticized on the ba­ cence fracture mode have been quanti­ sis that the small but finite plastic defor­ The planar pressure theory, proposed fied for the high-strength steel used in this mation observed on hydrogen-induced by Zapffee (Ref. 2), is based on the study. The resulting relationship agrees fracture surfaces requires more energy decrease in solubility of hydrogen as the with the results of Beachem but extend his than could be explained by the adsorption temperature is lowered. The atomic hy­ theory to a wider range of hydrogen con­ theory. In addition, fracture surfaces indi­ drogen is postulated to reassociate into tents. cate rapid void formation and coales­ diatomic hydrogen in pores and micro- cence at low temperatures where rate of voids. The pressure of diatomic hydrogen Introduction surface migration would be negligible. A theory proposed by Troiano (Ref. 5) It is known that hydrogen-assisted suggests that hydrogen interacts with dis­ cracking is a complex function of the location pileups in areas of triaxial stress to amount of hydrogen, the stress, the tem­ KEY WORDS lower the cohesive strength. It is known perature and the microstructure of the that hydrogen will diffuse toward regions steel. The purpose of this study is to Hydrogen Cracking of high triaxial stress such as those associ­ quantify the amount of hydrogen that High-Strength Steei ated with a stress riser. When the con­ Fracture Modes centration reaches a given level, the inter­ High-Strength Welds action of hydrogen with dislocation arrays Stress Intensity ahead of the stress riser is postulated to be 5. A. GEDEON is a Materials Engineering Consul­ Crack Propagation sufficient to cause fracture. Troiano sug­ tant at Temav Center for Research, Venice, It­ Implant Test Results gests that this interaction is due to the va­ aly. T. W. EAGAR is a Professor at Massachuetts Diffusible Hydrogen lence electrons from hydrogen atoms en­ Institute of Technology, Cambridge, Mass. Microplasticity tering the unfilled "d" shells of the iron MIL-A-46100 Steel Paper presented at the 69th Annual AWS and modifying the repulsive forces which Meeting, held April 17-22, 1988, in New Or­ determine the interatomic spacing in tran­ leans, La. sition metals.

WELDING RESEARCH SUPPLEMENT I 213-s the K,c of the final fractured area. Scanning electron microscopy (SEM) of H = 500 fractured implant specimens can be used 100 lo = 0.3 to determine the crack geometry. Speci­ mens with the same crack geometry as /© \ that studied by Daoud, etal., can then be used to determine the stress intensity fac­ 80 tor associated with that fracture.

Determination of Hydrogen Content in the /© \\ Cracking Zone c 60 "™ 1 / / \ \1 o The diffusible hydrogen test can be O used to determine the amount of hydro­ c gen initially solidified into the weld pool. a) 40 11 fo ^^ However, since fracture in the implant specimens will occur sometime after the weld has cooled down, and some hydro­ ll i r& gen will have been lost by diffusion, these 20 results must be analyzed to determine the amount of hydrogen remaining in the ©\©\j cracking zone at the instant of fracture. The amount and distribution of hydro­ fig. 1 —Hydrogen 0 V i i distribution as a gen remaining in an implant specimen as a function of distance 0.6 0.8 1.0 function of time after welding can be es­ in the weld for r = timated with the aid of a model initially 0.011, 0.044, 0.10, developed by Coe and Chano (Ref. 13). 0.20 and 0.50. Top Surface Fusion Line Bottom Surface They used an iterative procedure using small time-at-temperature increments to calculate the effect of time on the hydro­ gen distribution. The results are presented Others have modified the planar pres­ this means will not behave in the same as hydrogen as a function of the nondi­ sure theory and the adsorption theory by way as hydrogen introduced by an actual mensional parameter 9. This value is de­ assuming that hydrogen atoms are trans­ welding process. fined as: ported to the void or crack tip as Cottrell e = Dt/|2 (1) atmospheres. Bastein (Ref. 6) has pro­ Quantification of the Stress Intensity in a posed that hydrogen atoms are carried where D is the diffusivity of hydrogen in Weld along by the movement of dislocations solid iron, t is time, and lo is the weld bead during plastic deformation. Thus, he rea­ Among the various testing methods for depth. A sample of their distribution plots sons, dislocation pileups at structural de­ assessing hydrogen embrittlement, the is shown in Fig. 1, which shows the fects will produce an oversaturation of implant test has become one of the most hydrogen distribution as a function of dis­ hydrogen, which will result in an increase popular for scientific investigations of the tance in the weld for various values of 9. in pressure which in turn produces triaxial cracking phenomenon in welds. This is A better plot for the purposes of this re­ stresses and embrittlement. Research by due to the fact that the stress, hydrogen search is shown in Fig. 2, which shows the Graville (Refs. 7, 8) supports the hypoth­ level and microstructure can be indepen­ hydrogen concentration as a function of 9 esis that hydrogen transport by disloca­ dently varied and controlled. Crack sus­ at various weld locations. tions to the site of crack initiation is a nec­ ceptibility using this test is typically defined It has been postulated that dislocation essary part of the embrittlement process. as the lower critical stress (LCS). The LCS sweeping will increase the actual amount Beachem (Ref. 9) has proposed a theory is the maximum stress at which fracture of hydrogen at the crack tip. The in­ of hydrogen-assisted cracking, based on a does not occur for an arbitrarily long pe­ creased solubility of hydrogen under an microplasticity mechanism rather than em­ riod of time (usually 1 to 3 days). applied axial tensile stress has been esti­ brittlement. He suggests that the hydro­ Fracture mechanics can be used to de­ mated to be 5 times higher than the nom­ gen in the lattice ahead of the crack tip termine the stress intensity associated inal solubility by Louthan, et al. (Ref. 14). assists whatever microscopic deformation with fracture in the implant specimens. Andersson (Ref. 15) used a finite element processes the microstructure will allow. Since the helical notch used on the implant technique to estimate that the hydrogen Thus, intergranular, quasi-cleavage or mi­ specimens is too blunt to use linear elastic in front of the crack tip is about 1.2 times crovoid coalescence fracture modes will fracture mechanics (LEFM), the crack initi­ the nominal bulk hydrogen value. Schulte operate depending on the microstructure, ation process is difficult to quantify. How­ and Adler (Ref. 16), using nuclear reaction the crack tip stress intensity, and the con­ ever, once hydrogen embrittlement oc­ analysis of deuterium distribution, deter­ centration of hydrogen. The model unifies curs, the embrittled region itself will act as mined that the maximum hydrogen will be several theories, but shows that the planar a sharp crack tip, and one can use LEFM to about 1.4 times the nominal bulk hydro­ pressure and adsorption theories are un­ investigate the fracture of the remaining gen concentration. necessary. He proposes that the basic hy­ area, at least in high-strength welds. drogen-steel interaction appears to be an Daoud, etal. (Refs. 11, 12), have deter­ Experimental Procedure easing of dislocation motion or genera­ mined the stress concentration factor for tion, or both. an edge cracked circular bar in tension, Although no standard procedure exists In all of the above studies, the specimen and since modified this to include the ef­ for the implant method (Ref. 17), the IIW was charged with hydrogen in order to fect of the crack geometry. Even though published a document (Ref. 18) containing examine the effect on fracture. However, their analysis does not include the effect guidelines for performing this test. These Bonisewski and Moreton (Ref. 10) have of a restraining weld close to the crack, it procedures were followed using the I65 observed that hydrogen introduced by can be used to give an approximation of notch geometry and a helical notch.

214-s|JUNE 1990 A loading time of 5 min was chosen based on research by Peng (Ref. 19), who showed that variations in loading time H = 500 from 2 to 7 min after welding did not af­ 100 lo = 0.3 fect the lower critical stress (LCS). A 24- hour loading time was also chosen so that the hydrogen distribution model of Coe and Chano (Ref. 13) could be used to de­ IS 80 termine the amount of hydrogen in the cracking zone. o The material studied in this investigation o is a high-strength steel conforming to MIL- A-46100C (Ref. 20). Its main use is for ar­ 8" 40 mor in military applications and it is the main structural steel used in the M1 tank. Chemical composition requirements in MIL-A-46100C are very broad as the main performance criteria are good hardenabil­ ity and ballistic integrity. Due to its ex­ tremely high hardness, this material is very 2.0 3.0 susceptible to hydrogen embrittlement. 0.005 0.01 0.02 0.05 0,1 0.2 0.3 03 The composition of the 46100 used Dimensionless Diffusion Time,T throughout this investigation is listed in Fig. 2 - Hydrogen concentration as a function of T for different locations in the weld. C0 is the amount Table 1. It is composed primarily of tem­ of hydrogen initially in the weld, and Q is the total amount in the weld region. pered martensite with some banding. Due to the high hardness (53 RC), this steel had to be normalized to 35 RC in order to be Table 1—Composition of the Steel Used in This Study machined into implant specimens. The C Si Mn Cu P Ni S Al Cr Mo specimens were then austenitized in vac­ uum, quenched in oil, and tempered in air MIL-A-46100 0.31 0.41 0.97 0.38 0.011 1.21 0.008 0.044 0.51 0.50 to their original condition. The implant specimens were machined longitudinal to research program. Each LCS value was the rolling direction. was performed to map the various failure zones across the failed surfaces of the im­ determined from a plot of time-to-fracture Diffusible hydrogen testing was per­ plant specimens, and the crack geome­ vs. applied stress. Two such plots show formed in accordance with AWS A4.3-86. tries were compared with those modeled the effect of preheat (Fig. 3) and hydrogen The gas chromatography method was by Daoud, et al. Other details of the in the shielding gas (Fig. 4). In the majority used with a Yanaco hydrogen analyzer experimental procedures, materials used, of specimens, fracture occurred in the fu­ model C-1006. weld parameters and results have been sion zone. By varying the amount of hydrogen and published elsewhere (Ref. 21). The ratio Q/Q0 in Fig. 2 is approxi­ oxygen in argon GMAW shielding gas, mately equal to the ratio of the hydrogen time to loading, and preheat temperature, content at 24 h divided by the initial Experimental Results amount of hydrogen in the weldment at hydrogen content. Using the data from the time of cracking can be varied. A ma­ Table 2 summarizes the experimental Table 2, this ratio averages 0.25 for welds trix of seven conditions was studied for results acquired in this portion of the made without preheat. From Fig. 2, this each shielding gas composition: 1) Diffusible weld hydrogen content (as per AWS A4.3-86). 100 2) Hydrogen remaining 24 h after weld­ o ing (the AWS specimen was allowed to A X cool for 24 h before being analyzed). 3) Hydrogen remaining 24 h after weld­ ing with preheat. 80 4) LCS when loaded 5 min after weld­ ing. 5) LCS when loaded 24 h after welding. 6) LCS when loaded 5 min after weld­ ing with preheat. S 60 7) LCS when loaded 24 h after welding with preheat. Seven different shielding gas composi­ x 250°F Preheat tions were studied although not every gas 40 - A 150OF Preheat was evaluated both with and without o No Preheat preheat. A scanning electron microscope (SEM) was used to examine the fractured sur­ 20 j i i i 111 |I j i i i i 11 ii j—i i i 11 ul faces of the implant specimens. The initial 1 2 J—i—J. fracture in the majority of specimens was 100 10 10 103 due to hydrogen embrittlement, with the Time to Failure (min) remaining area failing due to microvoid Fig. 3 — Implant tests results for CMA welds made with 0% H2/2% 02/Ar and loaded 5 min after coalescence. Quantitative fractography welding. Curves depict welds made with 250"F, 150°F and no preheat.

WELDING RESEARCH SUPPLEMENT 1215-s 100 corresponds to a 9 of 0.9, which is very a 0% Hydrogen close to the value of 1.0 found if 9 is cal­ o 0.5% Hydrogen culated directly from the cooling curve + 2% Hydrogen and diffusivity vs. temperature data. 80 Based on the experimentally deter­ mined 9 value of 0.9, the amount of hy­ drogen located at the fusion line (I = lo = 0.3 in Fig. 2) will be equal to 12.5% of the initial hydrogen content in 2 60 the weld. For welds made with 250°F (121 °C) preheat, Q/Q0 averages 0.13, 24 h after welding, which corresponds to a 9 of 1.9. In this way, the amount of hydrogen at 40 - the cracking zone during the final fracture can be found for each of the implant specimens. This amount of hydrogen is termed the bulk hydrogen in the cracking 20 i i i i 11111 i i i i 11111 H zone and does not include any increased 10° 102 103 amount which may be due to the in­ creased stress state at the crack tip. Time to Failure (min) In order to determine the stress inten­ Fig. 4 —Implant test results for GMA welds made with Ar/ 2% O2 and loaded after 5 min. Curvessit y that caused cracking, a fractographic depict 0%, 0.5% and 2% H added to the shielding gas. 2 analysis of the fractured implant speci­ mens was performed. The cracking zone in all of the implant fractures studied was Table 2—Summary of Experimental Results at the weld fusion line. %H Figure 5 shows a typical overall view of 2 %o2 Added to Added to LCS LCS H H the fractured surface of an implant spec­ Shield Gas Shield Cas Preheat 5 min (ksi) 24 h 3 s (ppm) 24 h imen. Figure 6 shows the location of each 0 2 none 48.5 58 2.14 0.74 of the magnified photos taken of this sur­ face. Figure 7 shows the fracture mor­ 0 2 150°F 79 0.29 phology typical of hydrogen embrittle­ 0 2 250°F 82.5 86 - 0.22 ment as evidenced by intergranular facet­ 0.01 2 none 56 52.5 4.52 1.09 ing. The fracture morphology associated 0.1 2 none 46 48.5 6.80 1.23 with microvoid coalescence is shown in Fig. 8 as evidenced by the ductile dimples. 0.42 0.1 2 250°F 47.5 52 - The transition region showing areas of in­ 0.5 2 none 28 39 8.28 1.96 tergranular faceting below or next to 0.5 1 none - - 8.42 1.80 areas of microvoid coalescence is shown in Fig. 9. The resulting quantitative fracture 0.5 1 250° F 45 53 1.58 - map developed for this specimen is de­ 2 2 none 25 - 14.0 picted in Fig. 10. 2 1 none 26.5 34 8.17 2.19 Not all of the specimens exhibited such a clear distinction between the different 2 1 250°F 45 1.31 fracture zones. For example, a number of

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ft ,

Fig. 5 —Overall SEM view of a fractured implant specimen surface (20X). Fig. 6—Schematic showing the regions from which the magnified pho­ tos were taken.

216-s|JUNE 1990 •r r

:

I* win, Fig. 7 —Region of intergranular fracture characteristic of hydrogen em­ Fig. 8 —Region of microvoid coalescence showing ductile dimples brittlement (2000X). (2000X).

the low-hydrogen samples had areas of coalescence. Table 3 gives a presentation Table 3—Tabulation of Fracture Toughness fisheyes. A fisheye is an inclusion that is of the results for these 12 specimens. The Estimates locally surrounded by an area of hydrogen hydrogen values shown in Table 3 were embrittlement. The local area of hydrogen determined by estimating r from the cool­ Estimated embrittlement surrounding a fisheye is ing curve the time at which fracture H at presumed to be due to hydrogen trapping occurred, and finding the corresponding Stress Kic Crack Tip at the inclusion. Numerous investigators hydrogen concentration from Fig. 2. The Sample (ksi) a/D ksi(in.)1/2 (ppm) (Refs. 22-24) have found that hydrogen a/D ratio corresponds to the region of in­ 1.3 88.4 0.30 46.3 0.14 can be trapped at inclusions. tergranular fracture, which was assumed 1.4 82.2 0.42 65.0 0.15 Of 140 fractured implant specimens, to approximate a crack. 2.18 24.3 0.45 17.9 0.26 only 60 had cracks starting from one edge. The resulting plot of stress intensity vs. 2.20 113.0 0.13 32.0 0.245 Of the fracture maps developed for these amount of hydrogen in the cracking zone 3.6 24.0 0.41 14.4 0.275 60 specimens, the 12 that very closely ap­ is shown in Fig. 11. As can be seen, the 4.1 89.2 0.15 31.6 0.215 proximated the crack geometry studied stress intensity (an approximation of Kic at 4.6 82.8 0.35 57.8 0.19 55.0 0.50 by Daoud, et al. (i.e., circular-arc edge which microvoid coalescence occurs) de­ 17.2 70.5 0.135 2.2 82.2 0.41 68 0.81 crack), were used to determine the frac­ creases with increasing hydrogen in the 16.5 27.0 0.60 60.8 0.56 ture toughness (as estimated by K ) of 1C crack zone. However, at very high hydro­ 20.1 28.0 0.60 63.1 0.69 the area which fractured due to microvoid gen contents, intergranular fracture will

Fig. 9 — Transition region where both intergranular fracture and micro- void coalescence are evidenced (200X).

Fig. 10—Quantitative fracture map showing the regions of fracture types.

WELDING RESEARCH SUPPLEMENT | 217-s 70 _^< kips in.1/2), which is the same value found y" X by Herman and Campbell (Ref. 25) for the X fracture toughness of samples not ex­ <

8 posed to a corrosive environment. This

i n 1/2 ) \ * / datum point has been plotted on Fig. 11 - 50 \ MVC / for the Kic associated with hydrogen-free SMC yyJ specimens. X o \ / At medium levels of hydrogen (0.5% H2 added to the shielding gas), the KK varied UL. \ / with the applied load and time to failure. XX / - IG IG At low applied loads, the fracture tough­ #30 ness was almost as low as the K value. c scc At higher loads, the value increased to = 20 \ / approximately the nominal Kic value of 71 MPa m'/2 (65 kips in.'/2). In a few cases, t/1 xX — -- / very small amounts of hydrogen seemed £ io to increase the K^ of microvoid coales­ cence fracture above the K-ic of hydro­ 1 . . . , 0 1 i i , 1 I , , 1 i i 1 1 gen-free specimens. Hydrogen-induced 0,25 0.50 0.75 1.00 strengthening has been documented by Concentration of Hydrogen in Crack Zone (ppm) others (Ref. 14). White (Ref. 26) also noticed some slight strengthening of her Fig. 11 — Interrelationship developed in this study between the stress intensity factor, hydrogen con­implant specimens when welding with tent in the cracking zone, and mode of fracture. 0.05% hydrogen in the shielding gas. This phenomena may be due to hydrogen pinning the dislocations. be more energetically favorable than mi­ of Herman and Campbell (Ref. 25), who There were a number of specimens that crovoid coalescence until very high stress used fracture toughness samples of this had both high hydrogen contents and intensities are reached. identical type of steel and determined the high Kic values. These hydrogen values stress corrosion cracking toughness Kscc. were much higher than in specimens with Herman and Campbell found that the Kic values of 16 MPa m'/2 (15 kips in.,/2). Discussion fracture toughness of this material was 16 With the exception of the three points at The implant specimens that were MPa m'/2 (15 kips in.'/2) when exposed to very high hydrogen contents, the relation­ welded with a high-hydrogen content in distilled water. ship determined in this investigation quan­ the shielding gas (2% H2), had a K1C of At low hydrogen levels (0% H2 added to tifies the theoretical fracture mechanism about 16 MPa m/2 (15 kips in.'/2). This the shielding gas), the final fractures had a initially proposed by Beachem (Ref. 9). value agrees quite closely with the work toughness of approximately 71 MPa,/2 (65 One of the main features of the Beacham theory is the classification of fracture modes with respect to stress and hydrogen level. At relatively high stresses, hydrogen-assisted cracking can propa­ gate by microvoid coalescence, which is normally thought of as a ductile failure mechanism. Beachem proved that hydro­ gen can be responsible for microvoid co­ alescence by partially fracturing a sample in hydrogen, then freezing the sample in liquid nitrogen and sectioning the sample to find evidence of the processes occur­ ring ahead of Ithe crack tip. As the stress intensity decreases, crack propagation proceeds by the lower plastic deforma­ tion processes of quasi-cleavage, and fi­ nally, intergranular separation. Increasing hydrogen concentration at the crack tip has the effect of decreasing the stress in­ tensity at which these fracture processes occur. Beachem's model adequately explains the presence of plastic deformation pre­ ceding hydrogen cracking in the HAZ of welds (Ref. 27) and plastic deformation in other systems as well (Refs. 28-30). Also, H|G HQC HMVC the qualitative experimental results postu­ lated by Beachem in Fig. 12 bear a re­ Concentration of Hydrogen Dissolved markable resemblance to the quantitative results of the present investigation. in Crack Tip Material The major difference between the frac­ Fig. 12 — Suggested interrelationship by Beachem between stess intensity factor, dissolved hydro­ture map proposed by Beachem and the gen content, and HAC deformation mode of microscopically small volumes of crack-tip material one found in the present investigation, is

218-s|JUNE 1990 that this investigation shows that inter­ 70 — X granular failure can still occur at much X higher values of hydrogen. It makes sense /< ~x~ that intergranular failure will occur at 60 4 x high-hydrogen contents, but Beachem ' \ / suggests that microvoid coalescence will 50 MVC occur faster, and thus predominate. The \ ^ / three points at high hydrogen concentra­ tions are beyond the range investigated 40 I \ / by Beachem. Thus, Fig. 11 shows a mod­ o ification to the original work by Beachem, / IG namely, that high hydrogen concentra­ 30 tn X. * / tions can suppress the microvoid coales­ c cence fracture mode, and that intergran­ CD 20 / ular fracture will still be operative. The . ,, - NO HAC ^'NX Beachem model appears to be the most tn tn 10 comprehensive model to date, and ac­ a> counts for most experimental observa­ CO tions of hydrogen cracking. 0 " , , , i 1 i i i i 1 i i i 1 i i i i 0.25 0.50 0.75 1.00 The present investigation did not quan­ tify the hydrogen concentration that Concentration of Hydrogen in Crack Zone (ppm) caused intergranular or quasi-cleavage Fig. 13 — Interrelationship between the stress intensity factor, hydrogen content, and mode of frac­ fracture. There were not enough speci­ ture, including a hypothesized no cracking region. mens that exhibited the proper amount of quasi-cleavage fracture along with a crack geometry which approximated the frac­ ture mechanical analysis of Daoud,ef al. Conclusions drogen on the mechanical behavior of steel. Special Report No. 73, The Iron and Steel Insti­ An attempt was made to quantify the tute. relationship between stress intensity and The fracture modes of high-strength 6. Bastien, P., and Azou, P. 1951. Effect of hydrogen content for which no hydro­ steel welds have been characterized as a hydrogen on the deformation and fracture of gen-assisted cracks will propagate. In an function of the stress intensity and hydro­ iron and steel in simple tension. Proc. First unfractured specimen, it is assumed that a gen content at the cracking zone in im­ World Metallurgical Congress. ASM, Materials very small crack exists for which a/D is less plant tested welds. The results indicate Park, Ohio, pp. 535-552. than 0.2. From Daoud, ef al., the stress in­ that the hydrogen embrittlement theory 7. Graville, B. A. 1968. Effect of hydrogen tensity factor will be approximately unity. originally proposed by Beachem can be concentration on hydrogen embrittlement. Here, the maximum hydrogen present in used to explain the effect of hydrogen on British Welding journal 15(4):191-195. the lower critical stress specimens can be cracking of high-strength steels. The re­ 8. Graville B. A., Baker, R. G., and Watkin- used along with the applied stress (the LCS sults of the present study increase the son, F. 1967. Effect of temperature and strain value) and the assumed a/D ratio to range of hydrogen above that used in the rate on hydrogen embrittlement of steel. British Welding lournal 14(6):337-342. develop the "no hydrogen-assisted crack­ original Beachem study to show that large 9. Beachem, C. D. 1972. A new model for amounts of hydrogen will increase the ing" region in Fig. 13. The numbers for this hydrogen-assisted cracking (hydrogen embrit­ region should be considered to be tenta­ propensity for intergranular fracture rather tlement). Metallurgical Transactions 3(2):437- tive now since the assumptions are not than microvoid coalescence. 451. necessarily justified. 10. Boniszewski, T., and Moreton, J. 1969. Hydrogen entrapment in mild steel weld metal The current research has attempted to A ckno wledgments quantify both the stress intensity factor with micropores. Metal Construction 1(6):269- 276. and the amount of hydrogen responsible The authors wish to express their grat­ 11. Daoud, O. E. K., and Cartwright D. ). for causing microvoid coalescence. This is itude to the U.S. Army Materials Technol­ ogy Laboratory for financial support of 1985. Strain energy release rate for a circular- the first time that this has been attempted. arc edge crack in a bar under tension or bend­ this project. In addition, portions of this Discrepancies may arise due to the fact ing, lournal of Strain Analysis 20(1):53-58. work performed at MIT were supported that the implant specimens were not well 12. Daoud, O. E. K., Cartwright, D. )., and suited to Kic measurements. Another by the Office of Naval Research under Carney, M. 1978. Strain energy release rate for shortcoming may be that the bulk hydro­ contract N00014-80-C-0384. Also, the a single-edge-cracked circular bar in tension. gen in the cracking zone was determined competent assistance of James Catalano lournal of Strain Analysis 13 (2):83-89. rather than the hydrogen due to disloca­ and Atillio Santoro are gratefully acknowl­ 13. Coe, F. R., and Chano Z. 1975. Hydro­ tion sweeping or stress concentrations at edged. gen distribution and removal for a single bead the crack tip. However, the relationships weld during cooling. Welding Research Institute 5:33-90. developed in this investigation may be ac­ References 14. Louthan, M. R., McNitt, R. P., and Sisson curate since the increased amount of hy­ 1. Sawhill,). M., lr. 1973. A Modified Implant R. D. 1982. Importance of stress state on drogen due to stress concentrations may Test for Studying Delayed Cracking. Ph.D. dis­ hydrogen embrittlement. Advanced Tech­ be only a factor of about 1.4. sertation. Rensselaer Poly. Institute, Troy, N.Y. niques for Characterizing Hydrogen in Metals. Hopefully, future research will enable 2. Zapffee, C. A., and Sims, C. E. 1941. Hy­ Eds. N. F. Fiore and B. ). Berkowitz, The Metal­ anticipated hydrogen levels to be used to drogen embrittlement, internal stress, and de­ lurgical Society of AIME, Warrendale, Pa., pp. 25-42. quantify the allowable defect size, which fects in steel. Trans. AIME 145:225-259. 3. Petch, N. O., and Stables, P. 1952. Delay­ 15. Andersson, B. 1981. Hydrogen Cracking will result in a stress intensity factor lower ed fracture of metals under static load. Nature in Weldments. Ph.D. dissertation. Gothenburg, than that which causes hydrogen-assisted 169:842-843. Sweden, Chalmers University of Technology. cracking. Thus, very low hydrogen welds 4. Petch, N. O. 1956. Lowering of the frac­ 16. Schulte, R. L, and Adler, P. N. 1982. can be designed to either allow higher ture stress due to surface adsorption. Philo­ Evauation of stress-induced hydrogen and deu­ stresses or larger defects than high-hy­ sophical Magazine, Series 8, Vol. 1: 331-335. terium redistribution in titanium alloys using nu­ drogen welds. 5. Troiano, A. R. 1962. The infuence of hy­ clear reaction analysis. Advanced Techniques

WELDING RESEARCH SUPPLEMENT I 219-s of Characterizing Hydrogen in Metals. Eds. 22. Caskey, G. R., Jr. 1982. Tritium autorad- 26. White, D. R. 1986. Process Measure­ Fiore and Berkowitz. The Metallurgical Society iography. Advanced Techniques for Character­ ment of Hydrogen in Welding. Ph.D. disserta­ of AIME, Warrendale, Pa., pp. 233-244. izing Hydrogen in Metals Eds. Fiore and Ber­ tion. University of Illinois, Champaign, III. 17. Bryhan, A. J. 1981. The effect of testing kowitz, The Metallurgical Society of AIME, 27. Homma, H. 1972. A Study of Delayed procedure on implant test results. Welding Warrendale, Pa., pp. 61-76. Cracking in Hy-80 Weldments. Ph.D. disserta­ lournal 59(9): 169-s to 176-s. 23. Kumnick, A. J., and Johnson, H. H. 1980. tion. Rensselaer Poly. Institute, Troy, N. Y. 18. Cold cracking test methods using im­ Deep trapping states for hydrogen in deformed 28. Bernstein, I. M., and Pressouyre, G. M. plants. Weldinginthe World. 23 (Vfe):1985. pp. iron. Acta Metallurgica 28:33-39. 1985. The role of traps in the microstructural 8-12. IIS/IIW-802-84 (exdoc. IX-1240-82). 24. Tuyen, D. L, and Wilde, B. 1982. An au- control of hydrogen embrittlement of steels. 19. Peng, |. 1981. Weldability Studies of toradiographic technique for studying the seg­ Hydrogen Degradation of Ferrous Alloys. Eds. High-Strength Steels Using the Implant Test regation of hydrogen absorbed into carbon Oriani, Hirth and Smialowski: Noyes Publica­ Method. Master's thesis. The Ohio State Uni­ and low-alloy steels. Current Solutions to Hy­ tions. 641-685. versity, Columbus, Ohio. drogen Problems in Steel. Eds. Interrante and 29. Oriani, R. A. 1972. A mechanistic theory 20. Military specification. 1983. Armor Plate, Pressouyre, ASM, Materials Park, Ohio, pp. of hydrogen embrittlement of steels. Ber. Bun- Steel, Wrought, High-Hardness. Mil-A-46100, 413-422. senges Phys. Chem. 76:848-857. revision C. 25. Herman, W. A., and Campbell, G. M. 30. Hirth, J. P. 1985. Theories of hydrogen- 21. Gedeon, S. A. 1987. Hydrogen-Assisted 1985. Environmental-assisted cracking in high- induced cracking in steels. Hydrogen Degrada­ Cracking of High-Strength Steel Welds. Ph.D. hardness armor steel. U. S. Army Materials and tion of Ferrous Alloys. Eds. Oriani, Hirth and dissertation. MIT, Cambridge, Mass. Mechanics Research Center TR, pp. 85-28. Smialowski. Noyes Publications.

WRC Bulletin 341 February 1989 A Preliminary Evaluation of the Elevated Temperature Behavior of a Bolted Flanged Connection By J. H. Bickford, K. Hayashi, A. T. Chang and J. R. Winter

This Bulletin consists of four Sections that present a preliminary evaluation of the current knowledge of the elevated temperature behavior of a bolted flanged connection. Section I—Introduction and Overview, by J. H. Bickford; Section II—Historical Review of a Problem Heat Exchanger, by J. R. Winter; Section III—Development of a Simple Finite Element Model of an Elevated Temperature Bolted Flanged Joint, by K. Hayashi and A. T. Chang; and Section IV—Discussion of the ABACUS Finite Element Analysis Results Relative to In-the-Field Observations and Classical Analysis, by J. R. Winter.

Publication of this report was sponsored by the Subcommittee on Bolted Flanged Connections of the Pressure Vessel Research Committee of the Welding Research Committee. The price of WRC Bulletin 341 is $20.00 per copy, plus $5.00 for postage and handling. Orders should be sent with payment to the Welding Research Council, Suite 1301, 345 E. 47th St., New York, NY 10017.

WRC Bulletin 332 April 1988 This Bulletin contains two reports that characterize the mechanical properties of two different structural shapes of constructional steels used in the pressure vessel industry. (1) Characteristics of Heavyweight Wide-Flange Structural Shapes By J. M. Barsom and B. G. Reisdorf

This report presents information concerning the chemical, microstructural and mechanical (including fracture toughness) properties for heavyweight wide-flange structural shapes of A36, A572 Grade 50 and A588 Grade A steels. (2) Data Survey on Mechanical Property Characterization of A588 Steel Plates and Weldments By A. W. Pense This survey report summarizes, for the most part, unpublished data on the strength toughness and weldability of A588 Grade A and Grade B steels as influenced by heat treatment and processing. Publication of this Bulletin was sponsored by the Subcommittee on Thermal and Mechanical Effects on Materials of the Pressure Vessel Research Committee of the Welding Research Council. The price of WRC Bulletin 332 is $20.00 per copy, plus $5.00 for postage and handling. Orders should be sent with payment to the Welding Research Council, Suite 1301, 345 E. 47th St., New York, NY 10017.

220-s I JUNE 1990