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Materials Science and Engineering A 505 (2009) 105–110

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

icrostructural influences on hydrogen delayed fracture of high strength steels

i Soo Kima, You Hwan Leeb, Duk Lak Leeb, Kyung-Tae Parkc, Chong Soo Leea,∗

Department of Materials Science and Engineering, Pohang University of Science and Technology, Pohang 790-784, Republic of KoreaPOSCO Technical Research Laboratories, Pohang 790-100, Republic of KoreaDivision of Advanced Materials Science and Engineering, Hanbat National University, Daejeon 305-719, Republic of Korea

r t i c l e i n f o

rticle history:eceived 25 July 2008eceived in revised form 31 October 2008ccepted 24 November 2008

eywords:ydrogen delayed fractureempered martensiteull pearlite

a b s t r a c t

This study was aimed at investigating the effect of the microstructural constituents of high strengthsteels on their hydrogen delayed fracture properties. For this purpose, a series of constant loadingtests, slow strain rate tests, cyclic corrosion tests, and thermal desorption spectrometry analysis wasconducted on the hydrogen pre-charged specimens with tempered martensite or full pearlite show-ing a similar tensile strength level of 1600 MPa. Constant loading tests and slow strain rate testsrevealed that the tempered martensitic steel was more susceptible to hydrogen delayed fracture thanthe fully pearlitic steel. In slow strain rate tests, the maximum tensile strength decreased with increasingdiffusible hydrogen content in a power-law manner. The content of hydrogen inflowing from environ-

hermal desorption spectrometry ment was also simulated by cyclic corrosion tests. It was found that the fully pearlitic steel has thehigher equilibrium hydrogen content than the tempered martensitic steel. The primary trapping siteswere prior austenite grain boundaries for the tempered martensitic steel, and ferrite/cementite inter-faces and dislocations for the fully pearlitic steel. SEM fractographs revealed that the cracks inducedby hydrogen propagated along the prior austenite grain boundaries resulting in brittle intergranularfracture for the tempered martensitic steel while the fully pearlitic steel was fractured in a ductilemanner.

. Introduction

Up to the present, considerable efforts have been devoted tonvestigate the effects of hydrogen on the mechanical propertiesf steels. One of the common consensus of such studies is thathe higher the strength of steels, the more susceptible to hydrogenelayed fracture (HDF), especially over ∼1.2 GPa [1–3]. Nowadays,ith increasing demands for advanced high strength steels due to

conomical and environmental views, the concerns about HDF arelso increased [4]. There are several factors affecting the HDF prop-rties such as environment, stress and materials. Besides, HDF islso closely related to the microstructural factors which can affectydrogen diffusion and trapping [5,6]. Hydrogen is preferred to berapped at structural defects such as grain boundaries, dislocations,nd carbide/matrix interfaces as well as interstitial sites [7–9]. Inddition, by means of thermal desorption spectrometry, the con-

ent of hydrogen trapped at various structural defects has been

easured quantitatively [10]. As a result, it was found that the dif-usible hydrogen which can evolve out from specimens at roomemperature is responsible to HDF [11].

∗ Corresponding author. Tel.: +82 54 279 2141; fax: +82 54 279 2399.E-mail address: cslee@postech.ac.kr (C.S. Lee).

921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2008.11.040

© 2008 Elsevier B.V. All rights reserved.

Most of high strength steels for the use of fasteners in construc-tion and automotive industries consist of tempered martensite.Tempered martensite usually contains film-like carbides alongprior austenite grain boundaries, known to be one of the mostpreferential sites for excessive hydrogen trapping [12]. However,in the case of pearlitic steels, precipitation of carbides along theprior austenite grain boundaries usually does not occur for (hypo)eutectoid composition and can be suppressed even for hypereutec-toid composition by controlled cooling during austenite–pearlitetransformation. This fact indicates that the fully pearlitic (FP) steelsmay be more resistant to HDF compared to the tempered marten-sitic (TM) steels. However, limited researches have been performedto compare the hydrogen degradation properties of temperedmartensite vs. pearlite structures and, if any, their HDF proper-ties cannot be compared directly because of their different strengthlevel.

To this end, it is worthwhile to compare the HDF properties ofTM and FP steels at the same tensile strength level. Therefore, theobjective of the present study is to analyze the effect of microstruc-

ture on hydrogen degradation properties of high strength steelswith tempered martensite and full-pearlite structures having thesame tensile strength level of 1600 MPa via constant loading tests,slow strain rate tests, cyclic corrosion tests and thermal desorptionspectrometry analyses.

1 and Engineering A 505 (2009) 105–110

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. Experimental procedures

The steel used in this study (Fe–0.82C–0.23Si–0.78Mn–0.01P–.005S–0.18Cr in wt. %) was supplied in the form of hot rolledars with a diameter of 13 mm. Some bars were austenitized at223 K for 3 min and oil-quenched. In order to introduce the tem-ered martensite structure, a subsequent tempering treatment wasonducted at 723 K for 1 h. In order to obtain the fully pearliticicrostructure having the same strength with the above temperedartensite structure, some hot rolled bars were austenitized at

253 K and then lead-patented at 813 K. The diameter of the lead-atented bars was reduced to 7.25 mm by cold drawing. Subsequenteat treatment at 573 K for 2 h was conducted to remove the effectf residual stress introduced during cold drawing.

To evaluate the hydrogen degradation properties, constant load-ng tests and slow strain rate tests were conducted on the hydrogenre-charged specimens. The specimens were pre-charged withydrogen in a 0.1N NaOH and 3% NaCl + 0.3%NH4SCN aqueous solu-ion at the current density of 0.1–20 A/m2 for 48 h. After hydrogenre-charging, the specimens were plated with cadmium to preventydrogen desorption during the slow strain rate test. Constant load-

ng tests were performed on the pre-charged and notched roundar specimens with a notch root radius of 0.1 mm, i.e., a stress con-entration factor of 4.9 up to 100 h. The applied stress was 0.9�TS,here �TS is the tensile strength of the smooth specimen. Slow

train rate tests were conducted using an Instron 5861 machineith a constant stoke rate of 0.005 mm/min which is equivalent tostrain rate of 1 × 10−6 s−1. The specimens for the slow strain rate

est had the same dimension with those of the constant loadingest. The fracture stress was defined as a nominal maximum ten-ile stress in the slow strain rate test. After constant loading testsnd slow strain rate tests, the specimens were immediately putnto liquid nitrogen to prevent hydrogen release until the hydrogen

easurement.To obtain the equilibrium hydrogen content to the environment,

yclic corrosion tests were carried out on the cylindrical specimensØ 6 mm × 25 mm). A cycle consists of a series of dry (303 K, humid-ty of 50% for 2.75 h)–wet (303 K, humidity of 98% for 0.75 h)–saltpray (303 K, 0.5% NaCl for 0.5 h) procedures to simulate the cor-osive seashore environment. To evaluate the hydrogen contentnflowing from the corrosive environment, the thermal desorptionpectrometry analysis was performed on the specimens in every0 cycles. Here, the equilibrium hydrogen content was defineds the saturated hydrogen content. The hydrogen content in thepecimens was measured by using a gas chromatography with aeating rate of 100 K/h up to 1073 K. The heating rates of 200 and00 K/h were also employed to obtain the activation energy forydrogen desorption. The activation energy was calculated accord-

ng to the method reported by Choo and Lee [9]. The amount ofydrogen gas was analyzed at 5 min intervals using He carrieras.

Transmission electron microscopy (TEM) was used to observehe sub-microstructures. Samples for TEM observation were

echanically ground and polished to a thickness of about 0.1 mm,nd then electrochemically thinned with a double jet polishingachine in a mixed solution of 8 vol.% perchloric acid and 92 vol.%

cetic acid. The fracture surface was observed by scanning electronicroscopy (SEM) operating at 15 kV.

. Results and discussion

.1. Initial characteristics of the present steels

The microstructures of the two steels are shown in Fig. 1. Duringuenching and tempering of high carbon steels, a mixed structuref lath martensite and plate martensite is usually formed [13]. The

Fig. 1. Microstructures of the TM and FP steels: (a) tempered martensite structure ofthe TM steel, (b) prior austenite grains of the TM steel and (c) full-pearlite structureof the FP steel.

prior austenite grains of the TM steel showed a homogeneous grainsize distribution of 14.6 �m in the average (Fig. 1(b)), indicating thatthe abnormal grain growth did not take place during austenitizingtreatment. As for the FP steel, the homogeneous lamella spacingwas obtained by lead patenting with an average lamella spacing of75 nm (Fig. 1(c)).

To observe the detailed microstructure, TEM observation wasperformed. As shown in Fig. 2(a), the film-like carbides residedalong prior austenite grain boundaries in the TM steel, which is

likely to have a high susceptibility to HDF [14]. Meanwhile, in the FPsteel (Fig. 2(b)), pearlitic cementite was plastically deformed by colddrawing. Besides, dislocations created during cold drawing wereobserved in the ferrite phase. These structural defects such as prioraustenite grain boundaries, film-like carbide/matrix interfaces, dis-

J.S. Kim et al. / Materials Science and Engineering A 505 (2009) 105–110 107

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below which HDF will not occur is often used [16]. According to thedefinition of the critical hydrogen content, the amount of the criti-cal hydrogen of the TM and FP steels was 0.20 and 0.41 mass ppm,respectively.

ig. 2. TEM micrographs showing (a) grain boundary carbides in the TM steel andb) ferrite/cementite lamellae structure in the FP steel.

ocations and ferrite/cementite interfaces may serve as the trappingites of hydrogen.

The normal stress–strain curves of both steels tested at ambientemperature and at a strain rate of 3 × 10−3 s−1 are plotted in Fig. 3.oth steels showed the similar tensile strength level of 1600 MPa1621 MPa for the TM steel and 1639 MPa for the FP steel), and thelongation of the TM steel (total elongation of 10.0%) was greaterhan that of the FP steel (total elongation of 6.8%). It is noticeablehat the nominal stress of the FP steel showed a relatively sharp dropith increasing strain resulted from the fragmentation of cementite

amellar [15].The notch tensile strengths of TM and FP steels without hydro-

en charging (HD = 0 ppm) were found to be 2183 and 2453 MPa,espectively. Even though the tensile strength values of both steelsere similar to each other, the strength of pearlite microstructureas slightly higher than that of tempered martensite structure at

small strain (Fig. 3). When the round-notched specimens were

nder tensile load, the resulting strains were very small. Accord-ngly, the fracture stress of the FP specimen without hydrogenharging was much higher (by 270 MPa) than that of the TM speci-

Fig. 3. Tensile curves of the TM and FP steels at ambient temperature.

men although both specimens had the same tensile strength levelof 1600 MPa.

3.2. HDF properties

3.2.1. Time-to-failureTime-to-failure of the hydrogen pre-charged and notched spec-

imens obtained from the constant loading test was plotted againstthe diffusible hydrogen content in Fig. 4. The arrows indicate thespecimens which were not fractured up to 100 h. It is apparent thatthe time-to-failure was prolonged with decreasing diffusible hydro-gen content. In addition, the FP steel was superior in the resistanceagainst HDF to the TM steel. As one of the criteria for estimating theresistance to hydrogen degradation, the critical hydrogen content

Fig. 4. A plot showing time-to-failure vs. the amount of diffusible hydrogen for theTM and FP steels. It is noted that the FP steel shows the lower susceptibility tohydrogen delayed fracture than the TM steel.

108 J.S. Kim et al. / Materials Science and En

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along with the [HC] and [HE] values inferred from the constant load-ing test and cyclic corrosion test, the susceptibility parameters forthe TM and FP steels were estimated as 0.525 and 0.707, respec-tively, implying that the susceptibility to HDF of the FP steel is lowerthan that of the TM steel.

ig. 5. A log–log plot showing the dependence of notch tensile strength on theiffusible hydrogen content for the TM and FP steels.

.2.2. The maximum fracture stressThe slow strain rate test results are presented in Fig. 5; the

bscissa indicates the content of diffusible hydrogen diffused outrom specimens and the maximum fracture stress of notched spec-mens with a slow strain rate of 1 × 10−6 s−1 is manifested on therdinate. The maximum fracture stress of both steels decreasedith increasing diffusible hydrogen content which indicates theegradation of the mechanical properties due to hydrogen. Theaximum fracture stress of both steels drastically decreased up

o 0.5 mass ppm of the diffusible hydrogen content and then theecreasing rate was reduced. This relationship of the diffusibleydrogen content—maximum fracture stress can be expressed byqs. (1) and (2) for the TM and FP steels, respectively.

F = 877H−0.18D (1)

F = 1451H−0.11D (2)

here �F is the maximum fracture stress in MPa and HD is the dif-usible hydrogen content in mass ppm. Wang et al. [17] also reportedhe power-law relationship between the diffusible hydrogen con-ent and the maximum fracture stress for martensitic steels withensile strengths of 1300 and 1050 MPa. In their work, it was notedhat the maximum fracture stress of the steel with the lower ten-ile strength of 1050 MPa decreased slowly and linearly with thencrease of the diffusible hydrogen content up to 0.5 mass ppm dueo its relatively low tensile strength. On the contrary, in the presenttudy, the maximum fracture stress decreased with increasinghe diffusible hydrogen content in a power-law manner irrespec-ive of the microstructure, and the linear and slow drop of the

aximum fracture stress at the small content of diffusible hydro-en was not observed because of the higher tensile strength over600 MPa.

In Fig. 5, it is apparent that the FP steel exhibited the higherracture stresses for overall diffusible hydrogen contents than theM steel, indicating that the HDF resistance of the former is supe-ior to the latter. In the slow strain rate test, the critical hydrogenontent was defined as the hydrogen content equivalent to the max-

mum fracture stress of 0.9�TS where �TS is the nominal tensiletrength of the unnotched specimen in MPa. According to the def-nition, the critical contents of diffusible hydrogen were 0.07 and.43 mass ppm for the TM and FP steels, respectively. The criticalydrogen contents obtained from the slow strain rate test were in

gineering A 505 (2009) 105–110

good agreement with those estimated from the constant loadingtest.

3.2.3. Equilibrium hydrogen contentThe cyclic corrosion test was carried out to estimate the equi-

librium hydrogen content of both steels to the environment. Fig. 6illustrates the variation of the diffusible hydrogen content againstthe number of cycles. For the TM steel, the hydrogen content inflow-ing from environment reached the saturation point after 120 cycles.On the other hand, for the FP steel, it took 210 cycles to reach the sat-uration point. The saturated hydrogen content was 0.095 and 0.12mass ppm for the TM and FP steels, respectively. It is apparent thatthe longer time (90 cycles) was required for the FP steel to reach thesaturated (equilibrium) hydrogen content than the TM steel. More-over, the saturated hydrogen content of the FP steel was greater thanthat of the TM steel. The previous researches showed that the hydro-gen diffusivity of tempered martensite structure was greater bytwo orders of magnitude than that of full pearlite or ferrite/pearlitestructure since the ferrite/cementite interfaces served as the hydro-gen trapping sites [18,19]. The ferrite/cementite interfaces interferewith diffusion of hydrogen, resulting in more time required toreach the saturation point. Besides, if the ferrite/cementite inter-faces serve as the hydrogen trapping sites, the microstructure ofthe FP steel can provide a large area of trapping sites causing thegreater equilibrium hydrogen content compared to the TM steel.

From the above results, it is appropriate that the critical hydro-gen content and the hydrogen content from environment shouldbe considered simultaneously to evaluate the HDF resistance. Con-sidering these two factors for evaluation of the HDF resistance,Yamasaki and Takahashi [16] suggested a susceptibility parameteras Eq. (3).

Susceptibility parameter = ([HC] − [HE])[HC]

(3)

where [HC] is the critical hydrogen content and [HE] is the hydrogencontent from environment. According to Eq. (3), the larger the valueof parameter is the better the HDF resistance. In the present study,

Fig. 6. A plot showing the number of cycles vs. the desorbed hydrogen content forthe TM and FP steels.

J.S. Kim et al. / Materials Science and En

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an average value of the desorption from all of those trapping sites.

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ig. 7. Hydrogen desorption curves at a heating rate of 100 K/h for the TM and FPteels which were pre-charged with hydrogen at a current density of 10 A/m2 in a.1N NaOH aqueous solution.

.2.4. Hydrogen trappingThe thermal desorption spectrometry analysis was carried out

sing gas chromatography to analyze the hydrogen trapping phe-

omenon after slow strain rate tests and the results are shown inig. 7. At the heating rate of 100 K/h, the TM steel showed only oneydrogen desorption peak at 453 K. However, for the FP steel, thewo distinct hydrogen desorption peaks appeared at 410 and 565 K.

ig. 8. Macroscopic fracture surfaces of (a) the TM steel and (c) the FP steel after the slownd FP steels, respectively. The specimens were hydrogen pre-charged at a current densithe TM and FP steels corresponded to 0.67 and 0.56 mass ppm, respectively.

gineering A 505 (2009) 105–110 109

The corresponding hydrogen contents for the TM and FP steels were0.53 and 0.44 mass ppm, respectively. For the FP steel, the amount ofhydrogen corresponding to peak 1 was 0.12 mass ppm (27.3% of totaldesorbed hydrogen), and that to peak 2 was 0.32 mass ppm (72.7%of total desorbed hydrogen). The hydrogen trapping phenomenonwhich can be evidenced by the activation energy of hydrogen des-orption are strongly affected by the type of trapping sites. As a result,it can be inferred from Fig. 7 that the main trapping sites of the TMand FP steel are different.

Hydrogen in steels is prone to be trapped at the structural defectsdue to its low solubility in iron lattice. The hydrogen trapping sitescan be identified by calculating the trap activation energy for des-orption by Eq. (4) [20].

∂ ln(�/T2c )

∂(1/Tc)= −EaT

R(4)

where Tc is the peak temperature, � is the heating rate, EaT is thetrap activation energy for desorption and R is the gas constant. Thetrap activation energies of desorption for grain boundaries, disloca-tions and ferrite/cementite interfaces have been reported by manyresearchers to be 17.2, 26.8, and 65 kJ/mol, respectively [9,21]. In thepresent study, the activation energy was calculated from thermaldesorption spectrometry with various heating rates of 100, 200 and300 K/h. The activation energy for the hydrogen desorption peakof the TM steel was calculated to be 18.3 kJ/mol. Considering theearlier works [9], this corresponds to that of grain boundaries. How-ever, the martensitic steel has many kinds of trapping sites such asgrain boundaries, dislocations, interfaces and point defects. Thismeans that the calculated activation energy may be considered as

For the FP steel, the activation energies for hydrogen desorptionwere 24.8 and 66.3 kJ/mol for the desorption peak 1 and peak 2,respectively. As seen in Fig. 7, the thermal desorption spectrometryprofile of the FP steel consisted of the two peaks, i.e., high and low

strain rate test. Higher magnification images are shown in (b) and (d) for the TMy of 10 A/m2 in 0.1N NaOH aqueous solution and the diffusible hydrogen content of

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emperature peaks, and the portion of the high temperature peak isarger than that of the low temperature peak. It is considered thathe high temperature peak corresponds to the irreversible hydro-en trapping sites. This means that the large portion of hydrogennside of the FP steel is presumably trapped at irreversible trappingites, resulting in the superior hydrogen delayed fracture propertieso the TM steel. Takai and Watanuki [21] also reported that hydro-en occluded at high temperature about 673 K was innocuous dueo its high trap activation energy for desorption.

For the TM steel, considering both TEM micrograph in Fig. 2(a)howing the presence of film-like carbides along prior austeniterain boundaries and the trap activation energy for desorption of8.3 kJ/mol, it is inferred that pre-charged hydrogen was trappedxcessively at prior austenite grain boundaries and the film-like car-ides served as crack initiation sites. These cracks would propagatelong prior austenite grain boundaries, resulting in intergranularracture which is one of the typical characteristics of HDF [22].imilarly, TEM micrograph in Fig. 2(b) and the trap activationnergy for desorption of 66.3 kJ/mol for the FP steel inform that therapping site corresponding to the high temperature peak is fer-ite/cementite interfaces. Moreover, in the present FP steel, manyislocations induced by cold drawing existed in the ferrite phase aseen in Fig. 2(b). These dislocations are thought to be the trappingites for the hydrogen evolved at low temperature peak (desorptioneak 1) in Fig. 7 in view of its activation energy of 24.8 kJ/mol. These

acts are consistent with the results of the cyclic corrosion test.

.2.5. FractographyFig. 8 shows the fractographs for the steels after the slow strain

ate test. The specimens were pre-charged with hydrogen at aurrent density of 10 A/m2 in 0.1N NaOH aqueous solution. The dif-usible hydrogen content of the TM and FP steels in Fig. 8 was 0.67nd 0.56 mass ppm, respectively. The TM steel showed a mixedode fracture of intergranular and quasi-cleavage mode (Fig. 8(a)

nd (b)). The hydrogen-induced crack was initiated and propa-ated along prior austenite grain boundaries. This finding againonfirms that hydrogen is likely to be trapped at prior austeniterain boundaries and the film-like carbides along prior austeniterain boundaries are prone to crack initiation causing intergranularrittle fracture in the TM steel. By contrast, the FP steel exhibitedhe cup-and-corn type ductile fracture (Fig. 8(c)) and a dimpledtructure was well developed (Fig. 8(d)). For the FP steel, brittlentergranular fracture was constrained since the activity of priorustenite grain boundaries was suppressed due to eutectic trans-ormation. Besides, film-like carbides along prior austenite graintructure was not developed during the process of high strengthearlitic steel. As a result, intergranular fracture was not observedn the fracture surfaces of both specimens with and without dif-usible hydrogen. Only ductile fracture mode was observed in theP steel corresponding to the enhancement of resistance to HDF.

. Summary

1. The properties of hydrogen delayed fracture of the steels of tem-pered martensite and full-pearlite structure with the similartensile strength level of 1600 MPa were investigated by a seriesof constant loading tests, slow strain rate tests, cyclic corrosiontests and thermal desorption spectrometry analyses.

[[[

gineering A 505 (2009) 105–110

2. The mechanical tests, i.e. constant loading tests and slow strainrate tests, performed on the steels hydrogen pre-charged underthe same conditions revealed that the time-to-failure and max-imum fracture stress of the fully pearlitic steel were superior tothose of tempered martensitic steels.

3. The equilibrium saturation diffusible hydrogen content of bothsteels was also estimated by cyclic corrosion tests. It was foundthat the equilibrium saturation diffusible hydrogen content ofthe fully pearlitic steel was also higher than that of temperedmartensitic steels. Accordingly, the susceptibility to hydrogendelayed fracture of the former was lower than that of thelatter.

4. The primary hydrogen trapping sites of the tempered martensiticsteel were prior austenite grain boundaries, and those of the fullypearlitic steel were ferrite/cementite interfaces and dislocationsin pearlitic ferrite. As a result, the tempered martensitic steelwas fractured in an intergranular brittle manner while the fullypearlitic steel exhibited the dimpled and cup-and-cone shapedductile fracture mode.

5. From the above all results, it is obvious that, at the samestrength level, a fully pearlitic structure is more beneficial tosuppress hydrogen delayed fracture compared to a temperedmartensitic structure. In addition, the present results imply thatthe hydrogen delayed fracture behavior of high strength steelsprimarily depends on the microstructural constituents at thesimilar strength level.

Acknowledgements

This work was supported by POSCO and the grant from the Fun-damental R&D Program for Core Technology of Materials funded bythe Ministry of Knowledge and Economy of Korea. The authors arethankful to Dr. K. Tsuzaki and E. Akiyama of NIMS for the experi-mental help on hydrogen charging and analysis..

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