microstructurally mediated changes in fracture characteristics for electrochemically hydrogenated...

Materials Performance andCharacterization

Mobbassar Hassan Sk1 and Ruel A. Overfelt2

DOI: 10.1520/MPC20140026

MicrostructurallyMediated Changesin FractureCharacteristics forElectrochemicallyHydrogenated4340 Steel

VOL. 4 / NO. 2 / 2015

Mobbassar Hassan Sk1 and Ruel A. Overfelt2

Microstructurally Mediated Changesin Fracture Characteristics forElectrochemically Hydrogenated4340 Steel

Reference

Hassan Sk, Mobbassar and Overfelt, Ruel A., “Microstructurally Mediated Changes

in Fracture Characteristics for Electrochemically Hydrogenated 4340 Steel,” Materials

Performance and Characterization, Vol. 4, No. 2, 2015, pp. 68–83, doi:10.1520/

MPC20140026. ISSN 2165-3992

ABSTRACT

The effect of hydrogen on the fracture characteristics of low alloy 4340 steel

was studied using double-notched tensile samples electrochemically charged

in situ with hydrogen in 1 N H2SO4þ 5 mg/l As2O3 solution. Fracture behaviors

of samples with prior austenitic grain sizes of 10 and 40 lm and martensitic

hardness of 43–52 HRC (Rockwell hardness, C scale) were examined after

hydrogen charging times of 0–40 min. As expected, increases in hydrogen

charging time and hardness resulted in decreased failure strains and decreased

evidence of ductile fracture. Harder samples showed predominant

intergranular fracture close to the notch and a combination of ductile and flat

surfaces away from the notch. Softer samples showed mixed mode

intergranular and quasi-cleavage fracture close to the notch and predominant

ductile fracture as distance from the notch increased.

Keywords

hydrogen embrittlement, intergranular fracture, quasi-cleavage, fracture mode change,

prior austenitic grain size, 4340 steel, electrolytic hydrogen charging

Manuscript received February 18,

2014; accepted for publication

June 26, 2014; published online

August 28, 2014.

1

National Center of Excellence

for Research in the Intermodal

Transport Environment,

Materials Engineering, Auburn

Univ., Auburn, AL 36849,

United States of America,

e-mail: [email protected]

2

National Center of Excellence

for Research in the Intermodal

Transport Environment,

Materials Engineering, Auburn

Univ., Auburn, AL 36849, United

States of America,

e-mail: [email protected]

Copyright VC 2014 by ASTM International, 100 Barr Harbor Drive, P.O. Box C700, West Conshohocken, PA 19428-2959 68

Materials Performance and Characterization

doi:10.1520/MPC20140026 / Vol. 4 / No. 2 / 2015 / available online at www.astm.org

Introduction

Hydrogen-based fuel has been identified as a promising alternative to fossil fuel to

reduce the emission of greenhouse gases and to enable sustainable energy supplies

worldwide [1,2]. In spite of such promising applications of hydrogen as a prime

energy source, serious technical challenges exist because of metallurgical interactions

of hydrogen with fracture sensitive metals and alloys [2]. Hydrogen has caused

major failures [2] in the oil and gas industries. The uptake of hydrogen into steel

tanks, pressure vessels and pipelines has induced unexpected and premature fracture

of such structures.

Atomic hydrogen can be generated at the surface of metals and alloys by ther-

mal dissociation of gaseous diatomic hydrogen [3] or by electrochemical decomposi-

tion of hydrogen bearing aqueous solutions [4] and vapor condensates [5]. The

atomic hydrogen generated is first adsorbed onto the surface and then either diffuses

or is transported by dislocations into the interior of the metals [6]. Once inside

hydrogen-sensitive materials, e.g., high strength steel, hydrogen atoms can interact

with various microstructural features resulting in enhanced tendency for brittle

fracture.

Hydrogen is believed to primarily diffuse along the prior austenitic grain boun-

daries from the steel surface resulting in weakening those grain boundaries and

inducing their decohesion [7]. Hydrogen is also attracted to the triaxial stresses at

material discontinuities (e.g., notches, microcracks) reducing the strain energy for

crack initiation and propagation [7]. Additionally, hydrogen can be trapped within

the stress fields surrounding dislocations leading to either reduction in dislocation

mobility [8] or enhancements to dislocation mobility [9]. Hydrogen-enhanced local

plasticity (HELP) could induce grain boundary failures from amplified dislocation

pileups.

Hydrogen induced brittle fracture is microscopically manifested as (i) intergra-

nular fracture along the prior austenitic grain boundaries, or (ii) transgranular frac-

ture along the cleavage planes and/or martensitic lath boundaries as well as mixtures

of these two fracture modes. Another interesting fractographic feature often associ-

ated with hydrogen embrittlement is termed quasi-cleavage and displays characteris-

tics often associated with macroscopically brittle fracture with local plasticity. Recent

results [10,11] discuss some mechanistic aspects associated with quasi-cleavage but

the precise details and applicable models remain the subject of controversy [12].

Quantitative explanation of these oftentimes subtle fractographic differences has

remained difficult over the years.

The present paper shows how the fracture growth process quantitatively

changes with increasing distance from the notch for notched samples with a variety

of grain sizes, hardnesses, and hydrogen charging times. In this paper, we examined

the effects of hydrogen on embrittlement of 4340 steel using double-notched tensile

samples electrochemically charged in situ with hydrogen in 1N H2SO4þ 5mg/l

As2O3 solution. The fracture behaviors of samples were observed for prior austenitic

grain sizes of 10 and 40 lm, martensitic hardness of 43–52 HRC (Rockwell hardness,

C scale), and hydrogen charging time of 0–40min. Data quantifying different types

of observed fracture modes (i.e., intergranular, quasi-cleavage, flat-surfaces, or duc-

tile microvoid) were obtained using scanning electron microscopy.

SK AND OVERFELT ON FRACTURE CHARACTERISTICS FOR STEEL 69


Materials and Methods

4340 round bar tensile samples from the same heat (see Table 1) were procured from

Laboratory Devices Company (Placerville, CA). The diameter and the gage length

of the samples were 6mm and 58.5mm, respectively. As-received samples were

austenitized in argon atmosphere, oil quenched, and then double tempered at

Auburn University (Auburn, AL). The various austenitizing and tempering condi-

tions and the resultant prior austenitic grain sizes (PAGS) and hardnesses are shown

in Table 2.

After heat treatment, one sample from each heat-treated condition was

randomly chosen from the batch of 30–40 samples for microstructural analysis,

measurement of the prior austenitic grain sizes, and hardness measurements.

Specimens were sectioned transverse to their long axis for metallographic

analysis. Sectioned samples were mounted in bakelite, ground, and polished by

standard metallographic techniques. Observation of the prior austenitic grain

boundaries were revealed by etching with modified Winstead’s reagent at

65�C–70�C (100ml of 1.2 % saturated aqueous solution of picric acid, 2.5ml of

40 % solution of sodium do-decyl benzene sulfonate as a wetting agent, 5ml of

ethanol, and 10 drops of HCl). All etchants were applied by cotton swabs. Met-

allographic examination of the samples was performed with an Olympus PME3

optical microscope and grain size measurements used the intercept method per

ASTM E112-96 [13].

After heat treatment, the samples were returned to Laboratory Devices Com-

pany (Placerville, CA) where two semi-circular notches were ground according to

the dimensions shown in the Fig. 1. After grinding of the notches, all samples were

shipped back to Auburn University where they were stress-relieved at 190�C6 10�C

for 4.5 h.

The in situ hydrogen charging and tensile loading process is shown in Fig. 2.

The double-notched tensile samples were electrochemically charged in situ with

TABLE 1

Composition of AISI 4340 test specimens.

Elements C Cr Mn Ni P Si S Cu Al Mo N

Composition (wt. %) 0.4 0.79 0.74 1.69 0.007 0.29 0.014 0.18 0.036 0.24 0.009

TABLE 2

Prior austenitic grain sizes (PAGS) and martensitic hardnesses.

Tempering

350�C 500�C 350�C 500�C

Austenitizing(�C) (1 h)

Tempering(�C) (1 hþ 1 h)

Prior Austenitic Grain Size (PAGS)�Average Nominal Diameter in lm

HRC (Rockwell hardness,C scale)

870 350 500 106 3 106 3 52 45

1100 350 500 406 10 406 10 50 43

70 SK AND OVERFELT ON FRACTURE CHARACTERISTICS FOR STEEL


http://www.astm.org/Standards/E112

hydrogen in 1N H2SO4þ 5ml/g As2O3 solution for different time durations (0, 5,

10, 20, 30, and 40min) with a current density of 10mA/cm2 (characterized with

respect to the surface area of the sample exposed to electrolyte). 99.95 % pure plati-

num wire with nominal diameter of 0.5mm (from Alfa Acer) was used as the anode.

The samples were always located at the central position of the platinum coil with a

coil diameter of approximately 4 cm. To facilitate uniformity in this arrangement

and maintain a constant distance between the platinum coil and the sample, the

location of the platinum wire was permanently established by winding the wire

against a fixed support. To maintain electrical isolation between the sample and the

fixture of the Q-Test machine, the portions of the fixtures (bolt heads) inside the

electrolyte were sealed by using electrolyte-resistant and leak-resistant silicone

rubber sealant. After the samples were screwed into the fixture thread, junctions

between the sample and the bolt heads were also sealed using RTV silicone rubber.

This arrangement ensured the physical as well as electrical isolation of electrolyte

from the tensile machine and the entire electrochemical charging arrangement was

isolated from the tensile machine.

After electrochemical charging was completed, the electrolyte was quickly

removed from the charging cell and tensile testing started within 5 s from the end of

FIG. 1

Double notched round bar

tensile sample (measurements

are given in inches).

FIG. 2 Hydrogen charging apparatus was mounted on the tensile instrument (MTS Q-Test). Samples were charged with

hydrogen and immediately loaded under tension at cross head speed of 0.01 mm/min and the electrolyte was

simultaneously removed.



charging at a cross-head speed of 0.01mm/min using a Q-Test Elite 100 tensile test-

ing machine (MTS Systems Corporation). All tests were repeated four times at each

test condition.

The samples failed along either of the two notches (top or bottom). No practical

bias was seen for fracture preference of either of the two notches. Examination of

the fracture surfaces of the broken notch of the samples was performed using a

JEOLJSM7000F field emission scanning electron microscope operating at 20 kV with

energy dispersive x-ray spectroscopy (EDS) employing an ultrathin window detector

and Princeton Gamma-Tech analyzer.

Data quantifying the percentages of the various fractures modes (i.e., intergra-

nular, quasi-cleavage, flat, or dimpled) were quantified from the notch root along a

radial line to a depth of �1000 lm using intervals of 100lm and then to the center

of the sample using intervals of 500lm. Lines normal to the radial line were drawn

on the micrographs of the fracture surfaces and visual estimates of the percentages

of the fracture mode(s) exhibited characterized. A sampling distance of� 500lm,

normal to the crack growth direction, was employed for the quantification of the

fracture features.

Results and Discussion

MICROSTRUCTURE

The prior austenitic grain sizes (PAGS) are shown in Figs. 3(a) and 3(b) for the 10

and 40 lm PAGS, softer (45 and 43 HRC) samples and in Figs. 4(a) and 4(b) for the

10 and 40lm PAGS, harder (52 and 50 HRC) samples. The associated heat treat-

ment conditions are listed in Table 2. As expected, increasing the austenitizing tem-

peratures led to increases in the observed PAGS. Follow-on tempering of the

austenitized and quenched samples resulted in a range of martensitic hardnesses for

the samples produced. No significant differences in PAGS were observed because of

the various tempering procedures. However, the prior austenitic grain boundaries

for the tempering temperature of 350�C were revealed more prominently compared

to the tempering temperature of 500�C.

FIG. 3 Prior austenitic grain sizes for softer samples. Austenitized and tempered at (a) 870�C (1 h) and 500�C (2 h), and (b)

1100�C (1 h) and 500�C (2 h). Modified Winstead’s reagent.



OVERALL MECHANICAL RESPONSE

Figure 5 shows the average failure strain at 0.01mm/min cross-head speed for the

two PAGS and various hardnesses as a function of different hydrogen charging

times. The failure strain (ef) continually decreased with increasing hydrogen charg-

ing time. As expected, larger failure strains were observed with the softer samples

compared to the harder samples. Up to 10min hydrogen charging, the softer sam-

ples did not show any statistically significant difference (at 95 % confidence level) in

their failure strains compared to their uncharged counterparts. No statistically sig-

nificant (95 % CL) differences in failure strain values were observed between the 10

and 40 lm PAGS samples of comparable hardnesses. This result agrees with the

finding of Lessar and Gerbarich [14] who did not observe any significant difference

in fracture behavior of hydrogenated samples up to the prior austenitic grain size

range of �40lm.

FRACTOGRAPHIC BEHAVIOR

Figures 6(a) and 6(b) show representative fracture behavior close to the notch and

sample centerline for uncharged 10lm and 40 lm PAGS, softer (45 and 43 HRC)

samples. Representative images are taken from 100 and 2000 lm away from the

ground notch. The uncharged softer samples for both grain sizes showed a fracture

FIG. 4 Prior austenitic grain sizes for harder samples. Austenitized and tempered at (a) 870�C (1 h) and 350�C (2 h), and (b)

1100�C (1 h) and 350�C (2 h). Modified Winstead’s reagent.

FIG. 5

Change in failure strain as

function of hydrogen charging

time (0–40 min) during

mechanical loading at 0.01 mm/

min. Error bars shows the 2r

limit of scattering in failure

strain value.



surface composed of dimples representative of microvoid coalescence at all distances

from the notch. Figures 7(a) and 7(b) show representative fracture behavior from

close to the notch and sample centerline for uncharged harder (52 and 50 HRC)

samples with 10lm and 40 lm PAGS. Extensive microvoid coalescence is seen close

to the notch (up to a depth of �750lm). However, from that depth to the sample

centerline, the fracture surface of the harder samples exhibited a mixture of dimples

mixed with flat features.

Figures 8(a) and 8(b) show representative fracture behavior from close to the

notch toward the sample centerline for softer samples hydrogenated for 40min:

10lm PAGS (45 HRC) and 40 lm PAGS (43 HRC), respectively. Fracture surfaces

dominated by dimples is exhibited at large distances from the notch for both sam-

ples (2000lm; Figs. 8(a) and 8(b)), whereas a quasi-cleavage fracture mode (also

known as brittle transcrystalline [15]) mixed with some evidence for intergranular

fracture is favored in the high stress region closer to the notch. A number of second-

ary cracks are also evident in the vicinity of the notch.

Figures 9(a) and 9(b) show some of the quasi-cleavage appearances found in

10lm PAGS (45 HRC) and 40lm PAGS (43 HRC) samples hydrogenated for

40min. The quasi-cleavage fracture features seen in this investigation are similar to

FIG. 6 (a) Fracture characteristics for uncharged 10 lm PAGS, 45 HRC samples at distances �100 lm and �2000 lm away lm

from the notch. Arrows indicate the crack growth direction. (b). Fracture characteristics for uncharged 40 lm PAGS, 43

HRC samples at distances �100 lm and �2000 lm away from the notch. Arrows indicate the crack growth direction.



those reported by Martin et al. [10], Nagao et al. [11], and Lynch [15]. Martin et al.

[10] asserted that quasi-cleavage can result from the growth and coalescence of

microvoids that initiate and grow within intense slip band intersections. Lynch [15]

argued that quasi-cleavage (referenced therein as transcrystalline cleavage) occurs

across martensitic laths and/or along the martensitic lath boundaries. The smooth

surface features, which did not show rock-like feature with prominent grain bounda-

ries (as observed in intergranular fracture), were considered in this study as portions

of a quasi-cleavage surface. The smooth surfaces associated with the quasi-cleavage

fraction presumably resulted from crack growth along the interface and/or bounda-

ries of martensitic plates or laths.

The fracture behavior as a function of distance from the notch to the sample

centerline for harder samples hydrogenated for 20min is shown in Fig. 10(a) (10 lm

PAGS, 52 HRC) and Fig. 10(b) (40lm PAGS, 50 HRC). The fracture surfaces up to

1000lm depth dominated by intergranular fracture along with some secondary

cracking and with evidence of dimples. The fracture surface from about 1000lm

depth to the sample centerline exhibited a mixed appearance of flat, cleavage-like

features along with dimples. Although some of the flat features exhibited river

FIG. 7 (a) Fracture characteristics for uncharged 10 lm PAGS, 52 HRC samples at distances �100 lm and �2000 lm away from

the notch. Arrows indicate the crack growth direction. (b) Fracture characteristics for uncharged 40 lm PAGS, 50 HRC

samples at distances �100 lm and �2000 lm away from the notch. Arrows indicate the crack growth direction.



patterns typical of classical cleavage fracture, the majority of the flat features did not.

Additional investigation of these flat fracture features is currently underway to deter-

mine their similarity with flat, quasi-cleavage fracture features reported by Martin

et al. [10], Nagao et al. [11], and Lynch [15].

FIG. 8

(a) Fracture characteristics for

40 min hydrogen charged

10 lm PAGS, 45 HRC samples

at distances �100 lm, 500 lm,

1000 lm, and 2000 lm away

from the notch. Arrows indicate

the crack growth direction. (b)

Fracture characteristics for






the crack growth direction.



FIG. 9 (a) Some of the characteristic quasi-cleavage fracture appearances found in the 10 lm PAGS, 45 HRC samples. (b) Some

of the characteristic quasi-cleavage fracture appearances found in the 40 lm PAGS, 43 HRC samples.



FIG. 10

(a) Fracture characteristics for


10 lm PAGS, 52 HRC samples at

distances �100 lm, 500 lm,



the crack growth direction. (b)

Fracture characteristics for






the crack growth direction.



The quantitative variation of percentage intergranular fracture, percentage

dimpled surface, and percentage flat surfaces with increasing distance from the

notch were quantified by scanning electron microscopy for the range of experimen-

tal conditions investigated and the results are shown in Figs. 11 and 12.

Figure 11(a) present the quantitative variation of (i) % intergranular fracture,

and (ii) quasi-cleavage with increasing distance from the notch for the 10 lm, 45

HRC samples hydrogen charged for 0, 20, 30, and 40min, whereas Fig. 11(b)

presents the same results for the 40lm, 43 HRC samples. The softer samples did

not show any evidence of intergranular failure until 30min of hydrogen charging

and then only in the vicinity of the notch. Evidence for quasi-cleavage fracture

close to the notch is noted after 20min of hydrogen charging. Increasing the

hydrogen charging time above 20min induced more intergranular fracture and

FIG. 11 (a) Quantitative variation of % fracture modes with increasing distance from the notch for softer samples for 0, 20, 30,

and 40 min hydrogen charging times: (i) intergranular (IG) fracture for 10 lm PAGS, 45 HRC samples; and (ii) quasi-

cleavage for 10 lm PAGS, 45 HRC samples. Error bars represent 2r scatter. (b) Quantitative variation of % fracture

modes with increasing distance from the notch for softer samples for 0, 20, 30, and 40 min hydrogen charging times: (i)

intergranular (IG) fracture for 40 lm PAGS, 43 HRC samples; and (ii) quasi-cleavage for 40 lm PAGS, 43 HRC samples.

Error bars represent 2r scatter.



quasi-cleavage. However, quasi-cleavage was predominant. The intensity and

depth of both these fracture modes away from the notch was generally observed

to increase with increasing charging time. The softer samples, unlike their harder

counterparts, did not show any featureless flat surfaces away from the notch and

rather showed fracture involving predominant microvoid coalescence. Additional

investigation of these fracture features is currently underway.

Figure 12(a) presents the quantitative variation of (i) % intergranular fracture,

and (ii) flat surfaces with increasing distance from the notch for the 10 lm, 52 HRC

samples hydrogen charged for 0, 5, 10, and 20min, whereas Fig. 12(b) presents the

same results for the 40lm, 50 HRC samples. The flat surfaces characterized here are

similar to the fracture feature shown in Figs. 7(a), 7(b), 10(a), and 10(b) at �2000 lm

FIG. 12 (a) Quantitative variation of percent fracture modes with increasing distance from the notch for softer samples for 0,

20, 30, and 40 min hydrogen charging times: (i) intergranular (IG) fracture for 10lm PAGS, 45 HRC samples; and (ii)

quasi-cleavage for 10 lm PAGS, 45 HRC samples. Error bars represent 2r scatter. (b) Quantitative variation of % fracture

modes with increasing distance from the notch for softer samples for 0, 20, 30, and 40 min hydrogen charging: (i)

intergranular (IG) fracture for 10 lm PAGS, 45 HRC samples; and (ii) quasi-cleavage for 10 lm PAGS, 45 HRC samples.

Error bars represent 2r scatter.



away from the notch. Each condition represents four different samples that were

tested. The unhydrogenated samples did not show any evidence of intergranular fail-

ure, whereas increasing the hydrogenation times showed an increasing level of inter-

granular failure, which was observed to be most intense closer to the notch. The

maximum % intergranular fracture close to the notch is presumably the result of

stress-induced diffusion of hydrogen [16], which would increase the hydrogen con-

centration close to the notch where the triaxial stresses are a maximum. The

observed similarities in mechanical and fracture behavior for the 10 and 40lm

PAGS samples indicates independency of grain size within this PAGS range, in

agreement with the results of Lessar and Gerbarich [14].

The harder samples exhibited intergranular fracture at much lower hydrogen

charging times compared to the softer samples because of the commonly observed

tempered martensitic embrittlement in commercial low-alloy, high-strength steel

after tempering in the temperature range of 250�C–400�C [17,18]. Samples tem-

pered in this temperature range typically show grain boundary embrittlement

because of the singular or combined effects of alloying elements (P, S, N, Sn, Sb) seg-

regation and/or post-tempering (>250�C) carbide precipitation in the form of finely

dispersed cementite particles. However, the tempering temperature used for the

softer samples in this work (500�C) was too high to cause tempered martensitic

embrittlement. Above the temperature range of 400�C, the carbides would tend to

form spheroids and occupy a much smaller fraction of the grain boundary area [18].

Grain boundary segregation of the alloying elements is also not favorable above

400�C [17]. Thus, the loss of ductility observed for samples tempered at 500�C was

presumed to be caused predominantly by hydrogen.

Conclusions

1. The observed failure strain (ef) generally decreased with increasing hydro-gen charging time. Failure strains of the softer samples were noticeablylarger than the failure strains of the harder samples. Softer samples didnot show statistically significant difference in failure strains compared totheir uncharged counterparts for hydrogen charging times less than10min. Additionally, no statistically significant difference in failure strainswere observed between the 10 and 40 lm PAGS samples of comparablehardnesses.

2. Harder samples showed increases in % intergranular fracture with increase inhydrogen charging time. Intergranular fracture was predominant close to thenotch, whereas a combination of ductile and flat surfaces were exhibited inthe fracture surfaces away from the notch.

3. Softer samples showed a noticeable difference in fracture behavior withrespect to their uncharged counterparts after 20min of hydrogen charging.Softer samples charged above 20min showed mainly quasi-cleavage mixedwith intergranular fracture close to the notch and predominately ductile frac-ture as the distance from notch increased.

4. 10 lm and 40lm samples showed similar fractographic behavior, which is inagreement with their failure strain results. Similarities in mechanical behavioras well as fracture behavior between 10 and 40 lm PAGS samples indicatesindependency of grain size within this range.



References

[1] Kafle, J. and Qu, D., “Enhancement of Hydrogen Insertion into Carbon Inter-

layers by Surface Catalytic Poisoning,” J. Phys. Chem. C, Vol. 114, No. 44, 2010,

pp. 19108–19115.

[2] Louthan, M. R., Jr., “The Effect of Hydrogen on Metals,” Corrosion Mechanism,

F. Mansfield, Ed., Marcel Dekker, New York, 1987.

[3] Novak, P., Yuan, R., Somerday, B. P., Sofronis, P., Ritchie, R. O., “A

Statistical, Physical- Based, Micro-Mechanical Model of Hydrogen-Induced

Intergranular Fracture in Steel,” J. Mech. Phys. Solids, Vol. 58, No. 2, 2010, pp.

206–226.

[4] Shim, I. O. and Byrne, J. G., “A Study of Hydrogen Embrittlement in

4340 Steel I: Mechanical Aspects,” Mater. Sci. Eng. A, Vol. 123, No. 2, 1990, pp.

169–180.

[5] Sk, M. H., Overfelt, R. A., Haney, R. L., and Fergus, J. W., “Hydrogen

Embrittlement of 4340 Steel due to Condensation During Vaporized

Hydrogen Peroxide Treatment,” Mater. Sci. Eng. A, Vol. 528, Nos. 10–11, 2011,

pp. 3639–3645.

[6] Thompson, A. W. and Bernstein, I. M., Advances in Corrosion Science and Tech-

nology, Vol. 7, R. W. Staehle and M. Fontana, Eds., Plenum, New York, 1980.

[7] McMahon, C. J., Jr., “Hydrogen-Induced Intergranular Fracture of Steels,” Eng.

Fract. Mech., Vol. 68, No. 6, 2001, pp. 773–788.

[8] Gourmelon, A., “Effect of Hydrogen on the Plastic Deformation and Fracture

of Iron,” Memoires Scientifiques de la Revue de Metallurgie, Vol. 72, 1975, pp.

475–489.

[9] Birnbaum, H. K. and Sofronis, P., “Hydrogen Enhanced Localized Plasticity—A

Mechanism for Hydrogen Related Fracture,” Mater. Sci. Eng A, Vol. 176, Nos.

1–2, 1994, pp. 191–202.

[10] Martin, M., Fenske, A., Liu, G., Sofronis, P., and Robertson, I., “On the Forma-

tion and Nature of Quasi-Cleavage Fracture Surfaces in Hydrogen Embrittled

Steels,” Acta Mater., Vol. 59, No. 4, 2011, pp. 1601–1606.

[11] Nagao, A., Smith, C. D., Dadfarnia, M., Sofronis, P., and Robertson, I. M., “The

Role of Hydrogen in Hydrogen Embrittlement Fracture of Lath Martensitic

Steel,” Acta Mater., Vol. 60, Nos. 13–14, 2012, pp. 5182–5189.

[12] Lynch, S. P., “Interpreting Hydrogen-Induced Fracture Surfaces in Terms of

Deformation Processes: A New Approach,” Scripta Mater., Vol. 65, No. 10,

2011, pp. 851–854.

[13] ASTM E112-96: Standard Test Methods for Determining Average Grain Size,

Annual Book of ASTM Standards, ASTM International, West Conshohocken,

PA, 1996.

[14] Lessar, J. F. and Gerbarich, W. W., “Grain Size Effects in Hydrogen-Assisted

Cracking,” Metall. Trans., Vol. 7, No. 7, 1976, pp. 953–960.

[15] Lynch, S. P., “A Fractographic Study of Gaseous Hydrogen Embrittlement and

Liquid-Metal Embrittlement in a Tempered-Martensitic Steel,” Acta Metall.,

Vol. 32, No. 1, 1984, pp. 79–90.

[16] Li, J. C. M., Oriani, R. A., and Darken, L. S., “The Thermodynamics of Stressed

Solids,” Z. Physik. Chem. Neue Folge, Vol. 49, 1966, pp. 271–290.



http://dx.doi.org/10.1021/jp1064006

http://dx.doi.org/10.1016/j.jmps.2009.10.005

http://dx.doi.org/10.1016/0921-5093(90)90281-7

http://dx.doi.org/10.1016/j.msea.2011.01.100

http://dx.doi.org/10.1016/S0013-7944(00)00124-7

http://dx.doi.org/10.1016/S0013-7944(00)00124-7

http://dx.doi.org/10.1016/0921-5093(94)90975-X

http://dx.doi.org/10.1016/j.actamat.2010.11.024

http://dx.doi.org/10.1016/j.actamat.2012.06.040

http://dx.doi.org/10.1016/j.scriptamat.2011.06.016

http://www.astm.org/Standards/E112-96

http://dx.doi.org/10.1007/BF02644060

http://dx.doi.org/10.1016/0001-6160(84)90204-9

http://dx.doi.org/10.1524/zpch.1966.49.3_5.271

[17] Banerji, S. K., McMahon, C. J., Jr., and Feng, H. C., “Intergranular Fracture in

4340-Type Steels: Effects of Impurities and Hydrogen,” Metall. Trans A, Vol.

9A, 1978, pp. 237–247.

[18] Eliaz, N., Shachar, A., Tal, B., and Eliezer, D., “Characteristics of

Hydrogen Embrittlement, Stress Corrosion Cracking and Tempered Marten-

sitic Embrittlement in High Strength Steels,” Eng. Failure Anal., Vol. 9, No. 2,

2002, pp. 167–184.



Copyright by ASTM Int’l (all rights reserved); Tue Nov 18 18:38:26 EDT 2014Downloaded/printed byMobbassar Sk (Auburn University, Materials Engineering, Auburn, Alabama, United States)Pursuant to License Agreement. No further reproduction authorized.

http://dx.doi.org/10.1007/BF02646706

http://dx.doi.org/10.1016/S1350-6307(01)00009-7

microstructurally mediated changes in fracture characteristics for electrochemically hydrogenated...

Documents