Hydrogen Embrittlement Susceptibility of Gas Metal ArcWelded Joints from a High-Strength Low-Alloy Steel GradeS690QL
Martin Christ,* Xiaofei Guo, Rahul Sharma, Tianyi Li, Wolfgang Bleck,and Uwe Reisgen
1. Introduction
The use of high-strength low-alloy (HSLA) structural steels withfine-grained microstructure for applications such as steel con-struction, heavy commercial vehicle construction, or pipelineconstruction is becoming increasingly important for weightreduction.[1] In addition to the significant weight savings dueto wall thickness reductions, the use of HSLA steels in weldedstructures can significantly reduce the consumption of resources
and thus production costs. The importanceof conventional welding manufacturing forprocessing innovative materials is continu-ously increasing from an industrial point ofview to achieve the required technologicalproperties of welds. The fine-grainedhigh-strength structural steel S690QL(material number 1.8928) possesses goodlow-temperature toughness at �40 �Caccording to EN 10025-6. The fine-grainedmicrostructure is achieved by the additionof microalloying elements such as Nb andTi. As S690QL is widely used in weldedconstructions over a wide range of wallthicknesses, the weldability is a primaryrequirement criterion for a safe applicationof this material.[2] To ensure that themechanical properties of the weld metal(WM) meet the requirements of the base
metal, the heat input and the preheating or interpass tempera-ture during welding must be limited. This applies if suitablewelding consumables of the corresponding yield strength classare used. Gas metal arc welding (GMAW) is used for processinghigh-strength fine-grained structural steels both in pre- andonsite production. Although the welding of those steels usingthe GMAW process is widely and successfully used in industry,the risk of hydrogen-induced damage such as embrittlement orcold cracking must be considered. In general, the sensitivity tohydrogen damage increases with raised steel strength levels.[3]
Hydrogen-related damages to welded structural componentsoften occur unexpectedly without prior visible signs and clearlybelow the load limit of a construction, making them particularlycritical. For example, cold cracking or hydrogen embrittlement(HE) may occur in relation to welding due to a critical combina-tion of susceptible high-strength microstructure, stress state, andhydrogen absorption.[4,5] In particular, the high diffusivity ofhydrogen is regarded as favoring embrittlement.[6] Welding-related residual stresses and changes in microstructure as wellas external loading in addition to facilitate hydrogen diffusion,whereby a local critical accumulation in stress-concentratedregions and thus the risk of embrittlement increases.[7,8]
However, despite numerous investigations, the precisely defineddamage mechanisms caused by hydrogen have not yet been fullyexplained.[9] The understanding of the microstructure and dam-age mechanisms of HE sensitive regions in the welded structurebecome especially important.[9] Therefore, this research work
M. Christ, R. Sharma, Prof. U. ReisgenWelding and Joining Institute (ISF)RWTH Aachen UniversityAachen 52062, GermanyE-mail: [email protected]
Dr. X. Guo, T. Li, Prof. W. BleckSteel Institute (IEHK)RWTH Aachen UniversityAachen 52072, Germany
The ORCID identification number(s) for the author(s) of this articlecan be found under https://doi.org/10.1002/srin.202000131.
Herein, the hydrogen embrittlement (HE) susceptibility of base material and gasmetal arc welds from a high-strength low-alloy steel grade S690QL by slow strainrate test (SSRT) with in situ hydrogen charging is investigated. To investigate theinfluence of welding process parameters on the HE susceptibility of welded joints,gas metal arc welds in multilayer technique with two different heat inputs areconducted. The results reveal a considerable loss of fracture elongation for allhydrogen-charged specimens when compared with the uncharged referenceconditions. If the higher heat input is chosen, the required minimum yield strengthin the weld metal (WM) is not achieved. It is remarkable that this WM also revealedincreased HE susceptibility as measured by the relative elongation loss, despite itshigher ductility associated with the lower strength level. In combination withhydrogen measurements and fracture surface analysis, the crack propagationmechanisms and related HE mechanisms are intensively discussed.
applied GMAW with two different welding parameter sets tostudy their effects on microstructure and mechanical propertiesof the joint welds. Different locations from the welded jointswere further manufactured into tensile specimens to evaluatetheir HE sensitivity through the slow strain rate test (SSRT) within situ cathodic hydrogen charging.[10] With combined fracturesurface analysis, the damage mechanisms in the welded struc-tures were discussed.
2. Experimental Section
2.1. Base Material
A quenched and tempered high-strength steel S690QL with aplate thickness of 15mm was chosen as the base material withinthe framework of this study. Table 1 shows the contents of themain alloying elements as analyzed using spark optical emissionspectrometry.
2.2. Welding Conditions
Welded joints of S690QL steel plates were produced usingGMAW. The base material has a plate thickness of 15mm, theworkpiece geometry in the welded joint is 400mm� 300mm(length�width), whereby the welding direction follows the roll-ing direction. A V-butt joint with a weld preparation angle of 60�
was chosen. The workpiece was prefixed by tack welding of twoequal-sized plates with a gap of 2mm. The welding of the rootwas conducted with a reduced energy parameter set and a slightlybackward pointing electrode with the support of a ceramicbacking. To ensure stable welding process conditions on theHSLA steel workpiece, from which specimens are taken subse-quently, run-on and run-off plates were used for process start andend. Adapted to the strength of the base material, a type GMn4Ni1,5CrMo after EN ISO 16834 solid wire electrode witha diameter of 1.2 mm was used as filler metal for all weldingbeads. A mixture of argon (82%) and carbon dioxide (18%)was used as shielding gas, with the volume flow rate set to15 Lmin�1. The welds were successively filled in several layers,whereby the number of beads per layer naturally increases in thedirection of the top layer.
To investigate the dependence of the welding processing con-ditions on the resulting hydrogen sensitivity of the welded joint,two joint welds with different energy inputs per unit length wereproduced by varying the welding speed. The measured meanvalues for welding current and welding voltage are 250 A and30 V, respectively, derived from a wire feed rate of 12mmin�1
and resulting in transition arc mode. For the first joint weld(weld A), a welding speed of 0.45mmin�1 was used, whichresulted in an energy input per unit length of �1.0 kJ mm�1.For the second joint weld (weld B) with lower welding speed of0.30mmin�1, a correspondingly higher energy input per unit
length of 1.5 kJ mm�1 was induced. Within a joint weld (weldA or B), all weld beads of the filler layers following the root werewelded with constant parameters. Preheating of the workpieceswas not conducted in either of the two tests and an interpasstemperature of 150 �C was not exceeded. As a characteristic valueto describe the cooling curve of the weld beads, the cooling timefrom 800 to 500 �C (t8/5) was measured. In case of lower energyinput per unit length (weld A), the cooling time t8/5 of a singlefiller layer bead is �7 s, whereas it is �15 s using higher energyinput (weld B).
2.3. Microstructure Characterization
The microstructures from the base material and the weldedjoints under two different welding conditions A and B were char-acterized under optical laser microscope VK-X1000 (Keyencecompany). The base material is prepared transversal to the roll-ing direction. The welded specimens were cut from the crosssections at the middle height of the welded joints. The positionfor microstructure investigation corresponds to the latersampling position of the SSRT transversal tensile specimens.The specimens were mechanically ground, polished to 1 μmfinishing, and subsequently etched in alcoholic 1% nitric acid.
2.4. Hardness Measurements
In addition to metallographic images, hardness will give furtherinformation of local microstructure and properties. The hardnessprofiles proceeding horizontally through the welding zone weremeasured by hardness tester HMV-2000 (Shimadzu Company)according to Vickers (HV1) on the embedded cross-section speci-mens. The single hardness indentations were positioned at a dis-tance of 0.25mm from each other over a total width of 30mm.
2.5. SSRTs with Optional In Situ Hydrogen Charging
SSRTs were conducted with the specimens cut from two differ-ent positions of the welded joints as well as from the unaffectedbase material. As shown in Figure 1, the round tensile specimenswere taken longitudinally from the multilayer WM as well astransversally to the weld seam, with the specimen always beingcentered at half the thickness of the workpiece. Two longitudinaland three transversal specimens were taken from each weldedworkpiece, with the transversal specimens being taken at thebeginning, middle, and end of the weld seam and the longitudi-nal specimens in between. The SSRT specimen has a gaugelength of 25mm with a diameter of 3.5mm. The specimensfrom the base material were manufactured transversal to therolling direction of the plates.
The constant tensile strain testing machine manufactured byFritz Fackert Company was used at a controlled strain rate of5� 10�6 s�1. One set of tests was conducted in atmosphere,whereas another set of tests was conducted under in situ hydro-gen charging. As shown in Figure 2, a housing was installed withneutral solution of 3% NaCl and 0.3% NH4SCN at roomtemperature.
A platinum wire was used as the anode, whereas the speci-mens represented the cathode. In hydrogen charging condition,
Table 1. Chemical composition of S690QL steel, element contents givenin wt.%.
a constant current density of 1 mA cm�2 was applied to chargethe specimens artificially during the whole testing periods ofSSRT. Hydrogen atoms were generated continuously duringthe electrochemical process and diffused into the specimens.In case of hydrogen charging, the specimens were first surfaceactivated for 30 s with 20% HClþ 5 g L�1 C6H12N4 (hexamethy-lenetetramine, HMTA) and subsequently neutralized in 5 g L�1
NaOH. The selected boundary conditions of SSRT play a decisiverole in the final result.[11]
2.6. Hydrogen Measurement
After completion of SSRT, the gauge length region of one of thefracture halves was immediately cut off and stored in liquid
nitrogen before subjecting to hydrogen measurement.Hydrogen contents were measured with the analyzer G8Galileo (Bruker Company) by thermal desorption spectrometry(TDS) based on the carrier gas method. The specimens werecleaned in acetone and dried before the measurement. Duringthe measurement, the specimens were heated at a constant heat-ing rate of 20 �Cmin�1 from room temperature to 900 �C. Thehydrogen desorption rate curves were obtained within the mea-sured temperature range. The amounts of diffusible hydrogenare regarded as hydrogen contents below the first desorptionpeak which is usually ranging from room temperature to 450 �C.
2.7. Fracture Surface Analysis
The second half of the broken SSRT specimens was used to char-acterize the fracture surface. The investigations were conductedwith a field emission scanning electron microscope (FESEM) byZeiss Company at an operating voltage of 15 kV.
3. Results
3.1. Microstructure Examination
As a result of the chemical composition and the quenching andtempering process during production, the S690QL base metalhas a primarily martensitic–bainitic microstructure with finegrains. A light optical microscope image of the base metal micro-structure transverse to the rolling direction is shown in the leftpart of Figure 3.
The middle and right parts of Figure 3 show light opticalmicroscope images of both joint welds (weld A and B) fromthe area of the fusion line (FL) at the middle height of the multi-layer weld, i.e., the transition between WM and heat-affectedzone (HAZ). Since the welding heat input into the base materialdecreases with increasing distance from the FL, the HAZ doesnot have a uniform microstructure, but rather different charac-teristic subzones. The subzone immediately adjacent to the FLtherefore experiences the highest temperatures (well aboveAc3) within the HAZ and is called coarse grain heat-affected zone(CGHAZ) due to the grain coarsening tendency present here.
Figure 1. Schematic illustration of the positions of SSRT longitudinal and transversal specimens (Detail A) in the welded workpiece and the SSRTspecimen geometry.
Figure 2. Experimental setup for SSRT with in situ hydrogen charging.
The WM of both welds shows a needle-like structure typicalfor HSLA steels, consisting mainly of acicular ferrite. Becauseof its superior crack arrest capacity and the related high level oftoughness, it is generally desirable to achieve the formationof acicular ferrite in high-strength steel WM.[12–16] These acicularferritic structures in the WM are a little finer at lower heat input(weld A). However, the different heat inputs are particularly evi-dent in the microstructure of the CGHAZ. Here, the higherenergy input per unit length of weld B leads to extensive graingrowth with the prior austenite grain sizes above 100 μm. Theaccumulation of precipitates along the prior austenite grainboundaries was also observed.
3.2. Hardness Profiles in the Joint Welds
To identify local properties in the respective microstructurezones resulting from the selected welding parameters, hardnessprofiles proceeding horizontally through the welding zone aredetermined by means of hardness testing on cross sectionsaccording to Vickers (HV1), Figure 4.
The etching of the macrosections reveals that with increasingheat input, as expected, the dimension of the HAZ becomesconsiderably larger overall. On the basis of the hardnessprofiles, it can be observed that the slower cooling (t8/5 �15 s)associated with the increased heat input of weld B results in soft-ening of the WM when compared with weld A (t8/5 �7 s). Withthe average values of 250–270 HV1, the hardness of WM B isaltogether lower than that of WM A, which has the averagehardness values of 270 to over 300 HV1. Associated withthe extensive grain growth in the CGHAZ of weld B, as shownin Figure 3, local hardness peaks of about 370 HV1 wereobserved near the FL of weld B in Figure 4. In the finegrain or tempering zones of the HAZ, typical softening effects
become apparent, as compared with the base metal hardness ofabout 300 HV1.
3.3. Slow Strain Rate Tests
SSRTs were conducted with the specimens cut from twodifferent positions of the welded joints as well as from the basematerial. One representative curve for each testing condition wasdisplayed on the SSRT diagrams. The average values of mech-anical properties from repeated tests are shown in Table 2and Figure 5.
Figure 6 shows the engineering stress–strain curves ofthe base material S690QL tested in air condition and withadditional in situ hydrogen charging. The measured diffusiblehydrogen contents are indicated on the corresponding stress–strain curves. The base material obtains an average yieldstrength of 841� 7 MPa, a tensile strength of 880� 23 MPa,and a fracture elongation of 12.5� 0.4%. No diffusiblehydrogen has been detected in the base material tested inatmosphere.
In the comparison, a loss of fracture elongation is observed inthe in situ hydrogen-charged condition. The fracture of thehydrogen-charged specimen occurs almost immediately uponreaching the maximum force (ultimate tensile strength) withthe elongation at fracture reduced to 4.8� 2.4%. After comple-tion of the SSRT under in situ hydrogen charging, the specimensobtained a diffusible hydrogen content of 3 ppm.
Figure 7 shows the SSRT stress–strain curves of the tensilespecimens taken longitudinally from the multilayer WM of weldA that is conducted with an energy input per unit length of�1.0 kJmm�1.
Both the yield strength and tensile strength from WM of weldA are lower than those of the base material but well above therequired values, which amount to 707� 37 and 823� 46MPa,
Figure 3. Light optical microscope images of S690QL base metal (left) and joint welds A (middle) and B (right) from the area of the FL (marked by dottedlines) separating the WM from the CGHAZ.
Figure 4. Hardness profiles (HV1) across the multilayer joint welds A (left) and B (right). The indentation line positioned horizontally in the middle can beseen in each case.
respectively. It is primarily noticeable that the fracture elongationreduces substantially from 14.4� 2.5% to 2.8� 1.1% due to thein situ hydrogen charging. The yield strength has been raised to
773� 10MPa in the in situ hydrogen-charged condition,whereas the tensile strength has been reduced by 21MPa aver-agely. While the uncharged tensile specimen shows a soft tran-sition from the linear elastic to the plastic range, the artificiallyintroduced hydrogen causes an increase in the yield strengthwith a kink-shaped transition to the highly limited section ofplastic deformation. The considerable difference in deformabilitycan also be seen in the different necking behavior visible in thescanning electronmicroscopy (SEM) overview images of the frac-ture heads (Figure 7, right). The fracture head of the hydrogen-charged specimen is hardly constricted in cross section, whereasthe uncharged specimen shows a distinct necking. This limitedarea reduction at fracture in the hydrogen-charged condition canbe interpreted as a typical sign of HE. Furthermore, circularcracks can be seen on the lateral surface of the hydrogen-chargedspecimen (Figure 7, bottom right), which indicate additionalhydrogen damage.
Figure 8 shows the SSRT stress–strain curves of the tensilespecimens taken longitudinally from the multilayer WM of weldB. The energy input per unit length is �1.5 kJ mm�1 and thushigher than that in weld A.
As a result of the higher heat input and the associatedslower cooling rate (t8/5 �15 s) in comparison with weld A,the yield strength and tensile strength decline to 648� 14 and760� 13MPa, respectively, whereas the fracture elongationreaches 16.2� 2.1%. Due to the very slow cooling rate of weldB, the required minimum yield strength has not been achievedin the WM, so that this process parametrization does not satisfythe technical application. The raised elongation at fracture of theuncharged specimen could be caused by repeated tempering dueto the heat input of subsequent weld beads in the multilayerWM and is generally favored by the reduced strength level.The longitudinal WM specimen of weld B obtains no diffusiblehydrogen in the uncharged condition according to the measure-ment. In the in situ hydrogen-charged condition, the fractureelongation in the WM from weld B declines to 2.6� 0.7%, whichis very close to the longitudinal WM specimens in weld A.However, the specimens obtain a diffusible hydrogen contentof about 3 ppm, which is three times the value of the specimensin weld A. The charging duration for the two specimen sets isvery similar as concluded from their comparable fracture
Figure 5. Graphical representation of the characteristic mechanical prop-erties determined in the different SSRT series.
Table 2. Characteristic mechanical properties determined in the different SSRT series and the measured diffusible hydrogen contents.
Figure 6. SSRT stress–strain curves of S690QL base metal (transversal toRD), with and without hydrogen charging. The measured diffusible hydro-gen contents are indicated on the corresponding stress–strain curves.
elongations. Negligible area reduction of fracture surfaces andlateral surface cracks transversal to the specimen axis arealso observed in the hydrogen-charged specimens in Figure 8(bottom right).
The evaluation of transversal tensile tests on welded jointsis more complex compared with longitudinal tensile tests, astransversal specimens are metallurgically inhomogeneous andinclude several microstructure zones, such as the HAZs withtheir subzones and theWM. The unaffected base metal conditioncan be neglected for the transversal specimens used in this work,especially as there are only parts of the WM and HAZ included inthe gauge length. Figure 9 shows the SSRT stress–strain curvesof the transversal tensile specimens of weld A and the fracturefragments.
Compared with the longitudinal WM specimens from weldA, the transversal tensile specimens possess quite similar yieldstrength and tensile strength levels, however, a reduced fractureelongation of 10.4� 0.7%. The galvanostatic hydrogen charging
further reduces the elongation to 1.9� 0.7%. Despite thecomparatively short hydrogen charging duration due to itslow elongation, a mean diffusible hydrogen content of 3 ppmis measured immediately after fracture. In the chargedspecimen, circumferential surface cracks appear in the areanext to the fracture. The fracture position in each case is locatedin the WM.
The SSRT stress–strain curves and fracture fragments of thetransversal tensile specimens of weld B are shown in Figure 10.
The transversal tensile specimens of weld B reveal �93MPaless yield strength and �56MPa less tensile strength thanspecimens taken from the same position in weld A. The averagefracture elongation reaches 12.1� 1.6%, which is 1.7% highercompared with weld A and is close to the base material. Inthe hydrogen-charged condition, the specimen obtains almostidentical fracture elongation as the equivalent position ofweld A, however, with less detected diffusible hydrogen of2 ppm. In addition to the almost nonexistent necking of the
Figure 7. SSRT stress–strain curves of longitudinal WM specimens of weld A with and without hydrogen charging (left) and SEM overview images of onefracture half each (right). The measured diffusible hydrogen contents are indicated on the corresponding stress–strain curves.
Figure 8. SSRT stress–strain curves of longitudinal WM specimens of weld B with and without hydrogen charging (left) and SEM overview images of onefracture half each (right). The measured diffusible hydrogen contents are indicated on the corresponding stress–strain curves.
hydrogen-charged specimen, a distinct surface crack is visiblebelow the fracture surface (Figure 10, bottom right). The fractureposition of the uncharged reference specimen is located in theWM, whereas the fracture position of the hydrogen-chargedspecimens shifts to the region close to the FL.
As an overview, Table 2 shows the average characteristic val-ues of the different SSRT series and the measured diffusiblehydrogen contents. For better illustration, the SSRT resultsare additionally given as a bar chart in Figure 5, whereby the yieldstrength, the tensile strength and the elongation at fracture areconsidered. The error bars result from the deviation fromrepeated tests. In each case, the reference state is compared withthe hydrogen-charged state to highlight the effects of hydrogenon the mechanical properties.
In the uncharged condition, the specimens of weld A, which isbased on a lower welding heat input, achieve higher strengthlevels in both longitudinal and transversal directions than the
specimens of weld B taken at equivalent sampling positions.According to Figure 5, the in situ charged hydrogen slightlyraises the yield strength and reduces the tensile strength forall the specimen conditions considered. A significant reductionin fracture elongation due to galvanostatic hydrogen charging hasbeen observed over the entire test series. The fracture elongationhas declined to below 40% of its original value in the chargedcondition of the base material. The transversal specimens fromweld B show the most critical reduction in fracture elongationdue to in situ charging, which amounts to 16% of its originalvalue in the uncharged condition.
3.4. Fracture Surface Analysis
Fractographic analysis is conducted under SEM to look for indi-cations of possible hydrogen-induced damage characteristics.The hydrogen-charged state is directly compared with the
Figure 10. SSRT stress–strain curves of joint weld transversal specimens of weld B with and without hydrogen charging (left) and SEM overview images ofone fracture half each (right). The measured diffusible hydrogen contents are indicated on the corresponding stress–strain curves.
Figure 9. SSRT stress–strain curves of joint weld transversal specimens of weld A with and without hydrogen charging (left) and SEM overview images ofone fracture half each (right). The measured diffusible hydrogen contents are indicated on the corresponding stress–strain curves.
uncharged reference specimen taken from the same position.Figure 11 shows the fracture surfaces of the SSRT specimenstaken longitudinally from the WM of weld A.
In the fracture pattern of the uncharged specimen, typical duc-tile fracture mode with fine dimples is observed in both thefibrous core area and the lateral shear lip zone (Figure 11a,c,e).In the center of the dimples, inclusions can be found (Figure 11c,e), the formation of which is desired in high-strength steel WMsto serve as intragranular nucleation sites for acicular ferrite.[17–21]
In the in situ hydrogen-charged condition, the longitudinal SSRTspecimen from the WM of weld A exhibits a brittle fracturemorphology (Figure 11b,d,f ). In the central area of the fracturesurface (Figure 11d), rosettes appear as signs of quasicleavagefracture, whereby the local crack surface runs almost concentri-cally starting from a center and the individual fracture paths areintercepted in dimple fields. In the edge area (Figure 11f ), aneven more extensive damage in the form of cleavage facetsand cracks with lengths over 200 μm can be recognized due tothe edge-intruding exogenous hydrogen uptake.
The fracture surfaces of the longitudinal tensile specimensfrom weld B are shown in Figure 12.
In accordance with the high strain values in the unchargedstate, the dimple structure around inclusions confirms a ductilefibrous fracture over the entire fracture cross section of the WM
specimen (Figure 12a,c,e). As a consequence of the in situ hydro-gen charging, a mixed mode fracture (Figure 12b) consisting ofdimples (Figure 12d) and cleavage fracture (Figure 12f ) isobserved. The latter mentioned brittle fracture area revealsnumerous secondary cleavage microcracks.
Figure 13 shows the fracture surfaces of those SSRT speci-mens that are taken transversally from joint weld A.
In the fracture cross section of the uncharged specimen(Figure 13a), the core and edge areas can already be distin-guished from one another in the SEM overview image as a typicalresult of the different local stress state and plasticizing directionsduring necking. The resulting ductile dimples tend to be a littlelarger and deeper in the core area, whereas in the surroundingperipheral shear lip zone, they are slightly finer and partiallystretched perpendicular to the specimen axis due to the multiax-ial stress state during necking. The fracture surface of the in situcharged transversal specimen of weld A shows a mixed fracturecharacter, whereby the limitation of the different areas isadditionally highlighted by the red dotted line in Figure 13d.From the magnifications of the respective areas, it is obvious thatthe zone in Figure 13(f1) shows a typical brittle fracture surfacewith embrittled flakes, whereas Figure 13(f2) shows the typicalductile fracture mode with the similar dimple size as theuncharged condition.
Figure 11. Fracture surfaces of longitudinal WM specimens of weld A; a,c,e) uncharged condition, b,d,f ) with in situ hydrogen charging; (c,e) is themagnifications of the marked regions in (a), and d, f ) is the magnifications of the marked regions in (b).
Figure 14 shows the fracture surfaces of transversalspecimens from weld B. The fracture surface of the unchargedtransversal specimen from weld B can be clearly divided into thefibrous central area and the surrounding shear lip zone at theedge (Figure 14a). Both regions exhibit a ductile dimple struc-ture, with the dimples in the center being larger and deeper over-all (Figure 14c,e). Inclusions can be found within the dimples.As a result of the hydrogen charging, an inhomogeneous fracturesurface with a stepped breakout is formed (Figure 14b). A ductiledimple structure forms in the central area (Figure 14d), althoughthe dimples are very fine and relatively flat. In the embrittlededge area (Figure 14f ) of the hydrogen-charged specimen, inter-granular quasicleavage fracture with river line pattern and manysecondary cracks can be observed. The distinct kink bandsobserved on the fracture surface (Figure 14b) support theassumption that the fracture position shifts from the WM tothe FL due to hydrogen charging.
4. Discussion
Within the framework of this study, the HE susceptibility of gasmetal arc welds of a HSLA structural steel S690QL was investi-gated. Welded joints with different heat inputs were conducted
to investigate the influences of welding parameters on HE sus-ceptibility. Longitudinal WM specimens and specimens takentransversally from the welded joints were tested in slow strainrate tensile tests without further conditioning and with addi-tional in situ hydrogen charging. In the case of GMAW, thehydrogen pickup resulting from the welding process is signifi-cantly lower than that caused by artificial charging. In this study,no diffusible hydrogen has been detected in the unchargedspecimens taken from both the base metal and the welds.The variation of the heat input during gas metal arc weldingof S690QL influences the respective microstructure and leadsto different hardness levels, as shown in Figure 3 andFigure 4. The higher heat input of weld B (1.5 kJ mm�1) leadsto slower cooling rates and therefore to intensified grain coars-ening in the CGHAZ and carbon segregations at the prior aus-tenite grain boundaries. As a result, the local hardness at theCGHAZ reaches a higher peak value of about 370 HV1. Theassociated slower cooling rate in weld B simultaneously reducesthe hardness of the WM by increasing the tempering effect. As aresult, the hardness of WM B is averagely 20–30 HV1 lower thanthat of WM A and experiences an even greater hardness gradientin the area of the FL at the transition from the WM to the HAZ.Corresponding to the hardness results, the SSRT reveal that thesoftening of WM B leads to a lower yield strength and tensile
Figure 12. Fracture surfaces of longitudinal WM specimens of weld B; a,c,e) uncharged condition, b,d,f ) with in situ hydrogen charging; (c,e) is themagnifications of the marked regions in (a), and d,f ) is the magnifications of the marked regions in (b).
strength and a raise of fracture elongation compared with WMA. However, due to the low cooling rate of weld B, the requiredminimum yield strength in the WM has not been achieved.Nevertheless, under the same in situ hydrogen charging condi-tion, specimens from weld B exhibit a more pronounced reduc-tion in fracture elongation by decreasing to 2.6� 0.7% for thelongitudinal specimens and 1.9� 0.7% for the transversal speci-mens. This reduction means that the fracture elongation of weldB decreases to values that correspond to weld A or even fallbelow it in the hydrogen-charged condition. The steep hardnessgradient between WM and HAZ of weld B could be a trigger forthe early failure of the transversal specimens under in situhydrogen charging. According to fracture surface analysis, mosttransversal specimens failed in theWM, except the fracture posi-tion of transversal specimens from weld B shifted to the FL inthe hydrogen-charged condition. The stepped kink bands on thefracture surface of the transversal specimen of weld B(Figure 14b) supports this assumption. The microstructuralinhomogeneity and the respective differences in hardness leadto uneven yielding responses upon loading, subsequently, a highHE sensitivity is observed. As seen across the experiments, thein situ hydrogen charging during SSRT generally leads to asevere material embrittlement, which can primarily be identifiedby the loss of total elongation. The raised diffusible hydrogen
content as a result of the galvanostatic charging is regardedas the main reason for this strain loss. The contents of diffusiblehydrogen measured by TDS represent average values of therespective analyzed tensile specimen fragment, whereas thelocal hydrogen concentrations near the crack tip are supposedto be beyond those averaged values because of the stress concen-tration present there. In welded constructions, the CGHAZ isregarded as particularly prone to hydrogen-induced crackingdue to its critical properties.[22] Nevertheless, at the slow loadingspeed of the SSRT, the transversal specimens predominantlyfailed in the WM. Figure 5 compares the characteristic mechan-ical properties of the SSRT specimens from different locationsin the welded joints according to the two welding conditions.The results demonstrate that, in the uncharged condition, boththe longitudinal and transversal specimens reveal similar yieldstrength and tensile strength levels for the same welding condi-tion. As combined with the fracture surface observationsin Figure 13a,c,e and Figure 14a,c,e, the crack propagation inthe transversal specimens is assumed advancing preferablyin the WM zone, as demonstrated by numerous inclusions inthe ductile dimples as observed in the WM in Figure 11a,c,eand Figure 12a,c,e.
In combination with the fracture surface analysis, the fractureprocess can be interpreted according to the following steps.
Figure 13. Fracture surfaces of joint weld transversal specimens of weld A; a,c,e) uncharged condition, b,d) (f1) (f2) with in situ hydrogen charging;c,e) is the magnifications of the marked regions in (a), and d) is the magnification of the marked region in (b), and (f1) and (f2) are the magnifications ofthe marked regions in (d).
At first, the diffusible hydrogen is continuously charged into thespecimens and the diffusivity rises with the raised externalstress.[23] Second, in the region of plastic deformation, the hydro-gen penetrating from the outer surface induces surface damageand leads to circumferential surface cracks. Venezuela et al.[24]
attribute such surface cracks to the high local stress in the neck-ing zone, which in turn leads to an increased hydrogen concen-tration in the near-fracture region. A promoted hydrogen entry isexpected, since the specimens were activated and the surface bar-riers are cracked by plastic deformation and, in addition, thenumber of hydrogen traps such as dislocations and vacanciesis rising.[24] It is also observed that hydrogen seems affectingthe proportion of plastic strain between the lower yield pointand the beginning of work hardening, represented by theLüders strain. Finally, as far as the material cohesive force underhydrogen degradation is overcome by the external load,[25,26] thehydrogen-induced brittle fracture initiates in the peripheral areaof the specimen due to higher local hydrogen contents andquickly propagates toward the center. The stress–strain curvesfrom the hydrogen-charged specimens usually show a suddenkink with a steep drop in stress. On the fracture surfaces, thetypical hydrogen-induced damage features such as intergranularquasicleavage fracture pattern and secondary (micro) cracks areidentified. Additional crack formation in the region of the
fracture head could be an indication that hydrogen-induced dam-ages, i.e., cracks, are competing with the actually desired ductilefracture mechanisms.
5. Conclusions
On the basis of the applied methodology, it is possible to estimatethe potential damage caused by hydrogen. The HSLA steel gradeS690QL shows a considerable HE susceptibility due to galvano-static in situ hydrogen charging both in the base metal and in thegas metal arc welded state manufactured with different weldingparameters using a similar strength class filler metal of type GMn4Ni1,5CrMo. The hydrogen charging during SSRT leads topremature failure in all cases due to a significant reduction ofthe elongation capacity, which can be regarded as an importantevaluation factor of HE susceptibility. Due to the higher weldingheat input and the associated slower cooling rate of weld B, lowerWM hardness and simultaneous severe hardening in theCGHAZ were induced, which resulted in a sharp hardnessgradient. As a result, although the WM of weld B achievedhigher ductility through deficient yield strength, the specimensrevealed higher HE susceptibility as demonstrated by higherelongation losses.
Figure 14. Fracture surfaces of joint weld transversal specimens of weld B; a,c,e) uncharged condition, b,d,f ) with in situ hydrogen charging; c,e) is themagnifications of the marked regions in (a), and d,f ) are the magnifications of the marked regions in (b).
AcknowledgementsThe research project IGF 19.540N/P1105 “Determination of the influenceof processing and environmental hydrogen sources on the tendency ofhigh-strength steels to hydrogen-induced cold crack formation” fromthe Research Association for steel Application (FOSTA), Düsseldorf,was supported by the Federal Ministry of Economic Affairs and Energythrough the German Federation of Industrial Research Associations(AiF) as part of the programme for promoting industrial cooperativeresearch (IGF) on the basis of a decision by the German Bundestag.The project was conducted at the RWTH Aachen University, Weldingand Joining Institute and Steel Institute. The authors are grateful for thissupport. The authors also thank the companies of the project-relatedworking group for their support.
Conflict of InterestThe authors declare no conflict of interest.
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