Juraj Blach and Ladislav Falat* The Influence of Thermal

DOI 10.1515/htmp-2013-0053 High Temp. Mater. Proc. 2014; 33(4): 329 – 337

Juraj Blach and Ladislav Falat*

The Influence of Thermal Exposure and Hydrogen Charging on the Notch Tensile Properties and Fracture Behaviour of Dissimilar T91/TP316H Weldments

Abstract: The effects of ageing and hydrogen charging on the notch tensile properties and fracture behaviour of individual heat-affected zones (HAZ) and Ni-based weld metal (Ni WM) of T91/TP316H weldments were investi-gated. After the post-weld heat treatment at 750 °C for 1 h the weldments were annealed at 600 °C for 1000 h and 5000 h, respectively. All heat-treated states were studied in condition without as well as with hydrogen charging. Thermal expositions led to additional precipitation and microstructure coarsening but their influence on tensile strength was insignificant. In contrast, remarkable plas-ticity decrease and the fracture mode transition from ductile dimple tearing to transgranular cleavage were observed. The combined effects of thermal exposure and hydrogen charging were more complex. Whereas the regions of Ni WM and TP316H HAZ did not show any sig-nificant change in strength, the hydrogen effect caused the strength increase in T91 HAZ. Although the hydrogen embrittling effects were clearly manifested by decreas-ing plasticity, their significance was getting smaller with increasing annealing duration. The fracture behaviour of thermally exposed and hydrogen charged regions ex-hibited mixed fracture modes including transgranular cleavage, intergranular dimple fracture and intergranular decohesion.

Keywords: dissimilar weldment, thermal exposure, hydro-gen embrittlement, tensile properties, fracture behaviour

PACS® (2010). 81.70.-q

*Corresponding author: Ladislav Falat: Institute of Materials Research, Slovak Academy of Sciences, Košice SK 04001, Slovak Republic. E-mail: [email protected] Blach: Institute of Materials Research, Slovak Academy of Sciences, Košice SK 04001, Slovak Republic

1 IntroductionCr-Mo-V ferritic-martensitic and Cr-Ni-Mo austenitic creep- resistant steels and their welded joints are frequently used in classical coal-fired as well as nuclear power plant in-dustry for construction of boiler equipments operating at supercritical steam conditions [1–3]. Dissimilar welded joints of these steels are commonly produced by gas tung-sten arc welding (GTAW) using Ni-based weld metals (Ni WM) of “Inconel-type” to suppress undesired “up-hill” diffusion of carbon throughout the weldments [4].

Heat-affected zones (HAZ) represent potentially weak localities of weldments because of their microstructural degradation induced by welding thermal cycle, mani-fested in general by coarsening of prior austenitic grains and secondary phase particles compared to unaffected base materials [5]. Inner parts of power plant boiler tubing and/or piping may operate in hydrogen containing environments since the used work media, i.e. supercriti-cally heated and pressurized steam can be a source of free atomic hydrogen. Thus the hydrogen absorption may contribute to the local toughness deterioration and final failure. Embrittlement of metallic materials including creep-resistant steels by the absorbed hydrogen is a well-known phenomenon, commonly denoted as hydrogen embrittlement (HE) [6–8]. In the case of cooling-down of the power plants below 150 °C after their shutdowns, the hydrogen dissolved in steels may lead to the low threshold stress intensity for crack propagation at ambient tem-perature under tensile stress conditions. Such effect often results in hydrogen-induced cracking (HIC) or hydrogen- assisted cracking (HAC) phenomena [9].

There are many kinds of hydrogen trapping sites (e.g. grain boundaries, dislocations, and precipitates, etc.) in microstructure which can influence the mechanical prop-erties and fracture behaviour of steels and alloys in tensile loading conditions. It is well-known that plastic deforma-tion processes are in general strongly affected by the pres-ence of hydrogen in steels since free dislocations (but also specific configurations of dislocations such as grain and

330 J. Blach and L. Falat, Influence of Thermal Exposure and Hydrogen Charging

subgrain boundaries) represent reversible hydrogen trap sites [10]. Another category of hydrogen trap sites such as precipitated particles (in particular the carbide/matrix interfaces) represent irreversible hydrogen traps which may suppress the hydrogen embrittling effect and thus positively affect the mechanical properties and fracture behaviour [6, 11]. In general, there are several fundamen-tal hypotheses regarding the hydrogen acting mechanisms in microstructures and their influence on the plasticity and fracture behaviour, e.g. hydrogen enhanced decohe-sion (HEDE), hydrogen enhanced localized plasticity (HELP), adsorption-induced dislocation emission (AIDE), and hydrogen-enhanced strain-induced vacancy forma-tion (HESIV) [6, 12–16] but the investigation of real struc-tural materials always requires an individual approach taking into account complex influence of microstructural state, loading and environmental conditions.

The influence of hydrogen charging on the notch tensile properties and fracture behaviour of HAZ regions and Ni WM of dissimilar welded joint T91/TP316H in condition after post-weld heat treatment (PWHT) was in-vestigated in our previous work [17]. Present study rep-resents a continuation of research work performed on the same welded joint in conditions after thermal exposures at 600 °C for up to 5000 hours. Newly obtained results about the influence of hydrogen charging on the room- temperature tensile properties and fracture behaviour of long-term annealed states are compared with those of the initial PWHT state.

2 Experimental procedureThe tubes (38 mm outer diameter and 5.6 mm wall thick-ness) of T91 and TP316H creep-resistant steels were welded by gas tungsten arc welding (GTAW) using a Ni-based filler metal Thermanit Nicro 82. Chemical composi-tions of the materials used for fabrication of the investi-gated welded joint are listed in Table 1.

After the welding a PWHT procedure was carried out at 750 °C for 1 h with subsequent cooling down in air.

All experimental work was performed using cross-weld samples. Schematic cross-weld (c-w) sampling of tubular welded joint T91/TP316H is shown in Fig. 1. The first series of prepared c-w samples was investigated in the PWHT state. The other c-w samples were isothermally annealed at 600 °C for 1000 and 5000 h, respectively. From the c-w samples in all material states cylindrical tensile speci-mens (4 mm body diameter, 40 mm gauge length and M6 head thread) were machined by conventional turning. The body surfaces of the specimens were etched in order to visualize their HAZ regions. The compositions of etching solutions used for individual dissimilar materials are listed in Table 2. After the etching, the tensile specimens were circumferentially “V”-notched. The notch was alter-nately located in different parts of the specimens i.e. either in Ni WM or in the individual HAZ regions i.e. T91 HAZ or TP316H HAZ. One half of the notched specimens in all heat-treated states (i.e. PWHT, 600 °C/1000 h, 600 °C/5000 h) was subjected to tensile testing in condition without hydrogen charging. The second half of the specimens was subjected to electrolytic hydrogen charging at room tem-

Table 1: Chemical composition [mass%] of the materials used for fabrication of T91/TP316H welded joint with Ni-based filler metal Thermanit Nicro 82.

C N Si Mn S P Cr Ni Mo V Ti Nb Cu Fe

T91 0.092 0.045 0.39 0.44 0.003 0.011 8.68 0.25 0.92 0.2 – 0.064 – BalanceTP316H 0.052 – 0.51 1.77 0.006 0.031 16.76 11.13 2.05 – – – – BalanceThermanit Nicro 82 0.011 – 0.07 3.21 0.001 0.004 20.71 Balance 0.004 – 0.368 2.6 0.01 0.31

Fig. 1: Schematic cross-weld (c-w) sampling of tubular welded joint T91/TP316H.

Table 2: Etchants used for the etching of dissimilar materials.

Material Composition of etching solution

T91, TP316H 120 ml CH3COOH, 20 ml HCl, 3 g picric acid, 144 ml CH3OH

Thermanit Nicro 82 5 ml HF, 10 ml HNO3, 85 ml H2O

J. Blach and L. Falat, Influence of Thermal Exposure and Hydrogen Charging 331

perature prior to the tensile testing. The specimens were cathodically precharged with hydrogen for 24 h in a solu-tion of 1M HCl + 0.1N N2H6SO4 at a current density of 200 A.m−2. The used procedure was verified by our previous experience [17] which clearly indicated full saturation of the samples by hydrogen after 24 h of electrolytic charg-ing. Immediately after the hydrogen charging, room- temperature tensile tests were carried out at loading rate of 0.5 mm/min which allowed investigating of hydrogen effects for short term loads. From performed tensile tests the values of notch tensile strength (RmV) and reduction of area (RA) were determined. Microstructure charac-terisation and fracture analyses of broken tensile speci-mens were performed using scanning electron micros-copy (SEM) linked with energy dispersive X-ray (EDX) spectroscopy.

3 Results and discussion

3.1 Microstructure

Fig. 2 shows SEM microstructures of different regions of the dissimilar T91/TP316H welded joint in PWHT state. The microstructure of T91 steel HAZ (Fig. 2a) is formed of highly tempered martensite (i.e. ferritic laths with pre-cipitates) within significantly coarsened prior austenitic grains due to the heat input during welding. The tempered martensitic microstructure in the initial PWHT state typi-cally contains two types of precipitates, namely the elon-gated M23C6 (M = Cr, Fe) carbides at grain and subgrain boundaries and the globular MX (M = V, Nb, X = C, N) carbonitrides within the laths. This finding revealed by our SEM + EDX analysis fairly agrees with other literature data, e.g. [18]. In accordance with our previous study [17]

the microstructure of Ni WM Thermanit Nicro 82 (Fig. 2b) is formed of coarse migrated grains and the precipitates of two types, namely the primary blocky (Nb,Ti)C carbides and the secondary Cr23C6 carbides at the grain boundaries. The microstructure of TP316H steel HAZ (Fig. 2c) is formed of recrystallized polygonal grains with intergranular and intragranular M23C6 (M = Cr, Fe) carbides which are well-known to be the major type of precipitates in non- stabilised austenitic steels [19].

Fig. 3 shows the microstructures of different regions of the studied weldment after long-term annealing at 600 °C for 5000 h. In comparison to the microstructures in PWHT state (Fig. 2), the microstructures after thermal exposure are significantly coarsened with respect to the size of their precipitates (Fig. 3). The most significant change regarding the phase composition of T91 steel HAZ was the additional precipitation of Fe2Mo-based Laves phase during thermal exposure. This result agrees well with findings of other authors, e.g. [20]. Despite additional precipitation in the regions of Ni WM and TP316H HAZ (Fig. 3b and c), their qualitative phase compositions after the thermal exposure at 600 °C remained the same like in PWHT state.

3.2 Tensile properties

The results of room-temperature notch tensile tests of dif-ferent microstructural regions of the welded joint T91/TP316H in individual material states with respect to the heat treatment condition (PWHT, 600 °C/ 1000 h, and 600 °C/5000 h) and the application of hydrogen charging are graphically documented in Fig. 4. From Fig. 4a it is visible that the notch tensile strength (RmV) of T91 steel HAZ after both thermal exposures in condition without

Fig. 2: SEM microstructures of different regions of the dissimilar T91/TP316H welded joint in PWHT state: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).


hydrogen charging gradually decreases as a result of soft-ening processes related to the precipitate coarsening (Fig. 3a) and depletion of Mo atoms in solid solution via Laves phase precipitation. To the contrary, RmV values of Ni WM and TP316H HAZ regions without application of hydrogen charging gradually increase with thermal expo-sure. The observed strengthening effects are caused by additional fine precipitation in these regions, as shown in (Fig. 3b and c). In comparison to the base materials, the precipitation processes in HAZ regions are known to be enhanced by the welding-induced non-equilibrium thermal history resulting in high driving force for the ob-served microstructural changes [21]. Coarsening of sec-ondary phase particles in WM and HAZ regions during long-term thermal exposures seems to be a general reason for decreasing plasticity [22, 23] (see RA values in Fig. 4b).

Hydrogen charging of the welded joint T91/TP316H in all heat-treated states (PWHT, 600 °C/ 1000 h, and 600 °C/5000 h) had mostly small effects on the notch tensile strength of the individual microstructural regions (Fig. 4a). However, it should be noted that especially in the case of T91 HAZ region remarkable strengthening effects have been observed after hydrogen charging of long-term thermally exposed states. The reason for this behaviour is not completely understood at this point of investigation. It can be assumed that interactions of the absorbed hydrogen with free dislocations and/or precipi-tates in tempered martensitic microstructure of T91 HAZ may play a significant role in this case. However, it is still unclear why the similar hardening effects did not appear after the hydrogen charging of T91 HAZ region in PWHT state. It can be assumed that pinning of dislocations by the additional precipitation in thermally exposed states causes that hydrogen atoms have less possibility to move

Fig. 3: SEM microstructures of different regions of the T91/TP316H weldment after long-term annealing at 600 °C for 5000 h: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).

Fig. 4: Notch tensile properties of different microstructural regions of the welded joint T91/TP316H in individual material states with respect to the heat treatment condition and the application of hydrogen charging: notch tensile strength (a), reduction of area (b).


dragging by dislocations [24] resulting in strength in-crease after the hydrogen charging. A more focused inves-tigation of this behaviour will be the subject of our future studies. The variations in RmV values of other microstruc-tural regions (i.e. Ni WM and TP316H HAZ) in dependence of thermal exposure and subsequent hydrogen charging were found to be rather insignificant (Fig. 4a).

On the other hand, hydrogen embrittling effects were clearly manifested by gradual decrease in plasticity (RA values) in all microstructural regions and all heat-treated states of the studied weldment (see Fig. 4b). The reason for plasticity reduction of the heat-treated states without application of hydrogen charging is generally known to be related to thermal degradation of microstructure such as coarsening of secondary phase precipitates and/or ad-ditional precipitation of detrimental brittle phases (e.g. intermetallic Laves phase). However, in the case of com-bined effects of temper and hydrogen embrittlement the situation is much more complex. There are many different sites in microstructure such as free dislocations, grain boundaries, inclusions, and precipitates which may inter-act with the absorbed hydrogen and thus influence the resulting mechanical properties and final fracture be-haviour. Fig. 4b indicates that in some cases the effect of carbide coarsening and/or their additional precipitation during thermal exposures may even have a positive influ-ence on the plasticity via suppression of hydrogen embrit-tlement. This effect can be related to hydrogen trapping at carbide/matrix interfaces [6, 11] which can bind up free atomic hydrogen within a crystal lattice. An occurrence of this effect can be supported by the fact that the difference in RA values between the hydrogen free and hydrogen charged thermally exposed states decreases with increas-ing duration of long-term annealing. Thus it seems that

the influence of hydrogen charging on the plasticity with increasing duration of thermal exposure decreases and the dominance of temper embrittlement over hydrogen embrittlement increases.

3.3 Fractography

Figs. 5 and 6 show the SEM-fractographs of different regions of the welded joint T91/TP316H in PWHT state after the room-temperature tensile testing in condition without and with hydrogen charging, respectively.

In PWHT state without hydrogen charging the fracture mechanism was found to be a ductile dimple tearing in all tested regions of the weldment (see Fig. 5). The size and morphology of dimples in the individual microstructural regions are mostly uniform, although some deviations from this rule occur occasionally in some locations. After subsequent hydrogen charging of the welded joint in PWHT state, the fracture mechanisms in individual re-gions (both HAZs and Ni WM) have changed by a presence of transgranular cleavage besides dimple fracture (see Fig. 6).

Figs. 7 and 8 show SEM-fractographs of different regions of the T91/TP316H weldment in thermally exposed state at 600 °C/1000 h after room-temperature tensile testing without and with hydrogen charging, respectively. In addition, Figs. 9 and 10 show SEM-fractographs of different regions of the studied weldment in thermally exposed state at 600 °C/5000 h after room-temperature tensile testing without and with hydrogen charging, re-spectively. In both thermally exposed states (600 °C/1000 h and 600 °C/5000 h) without hydrogen charging (Figs. 7 and 9) the fracture mechanisms of the individual regions

Fig. 5: SEM-fractographs of different regions of the welded joint T91/TP316H in PWHT state after the room-temperature tensile testing in condition without hydrogen charging: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).


Fig. 6: SEM-fractographs of different regions of the welded joint T91/TP316H in PWHT state after the room-temperature tensile testing in condition with hydrogen charging: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).

Fig. 7: SEM-fractographs of different regions of the T91/TP316H weldment in thermally exposed state at 600 °C/1000 h after room-temperature tensile testing without hydrogen charging: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).

Fig. 8: SEM-fractographs of different regions of the T91/TP316H weldment in thermally exposed state at 600 °C/1000 h after room-temperature tensile testing with hydrogen charging: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).


of the dissimilar weldment were affected by a growth of secondary phase particles (in all studied regions) and/or additional precipitation of new phases, such as the Laves phase in T91 steel HAZ. However, the fractographic mani-festations of the individual regions were different from each other. Moreover, the duration of thermal exposure had also a significant influence on the resulting fracture behaviour. After the thermal exposure at 600 °C for 1000 h in condition without hydrogen charging (Fig. 7), qualita-tively similar fractographic features have been observed on fracture surfaces of the individual microstructural regions like in the case of hydrogen exposed PWHT state without thermal exposure (Fig. 6), i.e. a mixed fracture mode consisting of ductile dimple tearing besides trans-granular cleavage. Further thermal exposition at 600 °C for 5000 h without application of hydrogen charging led to the pronounced microstructure coarsening which was

strongly manifested in fracture behaviour of T91 HAZ through its transition from the mixed fracture mode (Fig. 7a) to the completely brittle transgranular cleavage and/or quasicleavage fracture mechanism (Fig. 9a). In addition, secondary cracking has been clearly observed in this region. To the contrary, the region of Ni WM after the longest thermal exposure (Fig. 9b) did not exhibit any substantial changes in fracture behaviour. The region of TP316H HAZ showed in the same condition (Fig. 9c) typical signs of intergranular decohesion besides trans-granular cleavage and dimple fracture mode.

Hydrogen charging of both thermally exposed states (600 °C/1000 h and 600 °C/5000 h) has induced additional changes in fracture behaviour of the individual regions of the studied weldment (Figs. 8 and 10). Thermal exposure at 600 °C for 1000 h with subsequent hydrogen charging resulted in fractographic features differing significantly

Fig. 9: SEM-fractographs of different regions of the studied weldment in thermally exposed state at 600 °C/5000 h after room-temperature tensile testing without hydrogen charging: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).

Fig. 10: SEM-fractographs of different regions of the studied weldment in thermally exposed state at 600 °C/5000 h after room-temperature tensile testing with hydrogen charging: T91 steel HAZ (a), Ni WM Thermanit Nicro 82 (b), TP316H steel HAZ (c).


among the individual microstructural regions (Fig. 8). The fracture surface of thermally exposed (600 °C/1000 h) and hydrogen charged region of T91 HAZ (Fig. 8a) exhibits a mixture of transgranular cleavage facets and fine dimples located presumably in intergranular areas. The region of Ni WM in the same experimental condition (Fig. 8b) shows also a mixture of the both above mentioned fracture modes but the size and morphology of dimples and cleav-age facets is different in comparison to the previous case. In contrast, the fracture surface of TP316H HAZ (Fig. 8c) is formed of transgranular cleavage facets, intergranular decohesion areas and a little portion of intergranular dimples. The fracture surfaces of individual microstruc-tural regions after the thermal exposure at 600 °C for 5000 h with subsequent hydrogen charging (Fig. 10) exhibit qualitatively the same fracture characteristics like in the previous case (Fig. 8). The differences related to the effect of thermal exposure duration are manifested by only small changes in size and/or portion of the indi-vidual fractographic features. After hydrogen charging of metals and/or alloys, the hydrogen tends to be distributed heterogeneously in their crystal lattice. At the stress- concentrator sites, such as grain boundaries, dislocations, precipitates, inclusions, micro-voids, and crack tip, the hydrogen concentrations are locally increased. Moreover, under external stress conditions the hydrogen migrates via diffusion paths represented by reversible hydrogen traps (mainly dislocations and grain boundaries) thus af-fecting the mechanisms of plastic deformation and final fracture behaviour [10]. On the other hand, irreversible hydrogen traps, such as particle/matrix interfaces (in particular carbide/matrix interfaces) may suppress the hy-drogen transport in the alloy with a positive effect on the plasticity (Fig. 4b) and fracture behaviour (Figs. 8 and 10).

4 Summary and conclusionsThis study was focused on the effects of long-term ageing and subsequent hydrogen charging on the room- temperature tensile properties and fracture behaviour of individual microstructural regions of dissimilar welded joint T91/TP316H with Ni-based weld metal. These are the main conclusions:

– Long-term annealing of the studied weldment led to the additional precipitation and significant coarsen-ing of secondary phase particles in all investigated regions but its influence on tensile strength was only small. To the contrary, microstructural changes induced by the annealing were manifested by remark-able decrease in plasticity.

– With increasing annealing duration the fracture mode transition occurred from pure ductile dimple tearing in the initial PWHT state to the mixed fracture modes containing also transgranular cleavage and/or inter-granular decohesion in thermally exposed states.

– The manifestation of combined effects of thermal ex-posure and hydrogen charging on the tensile proper-ties and fracture behaviour was much more complex than in the case of simple annealing. Whereas the hydrogen charging of thermally exposed regions of Ni WM and TP316H HAZ did not show any significant influence on their tensile strength, the effect of hydro-gen was clearly manifested by the strength increase in the region of T91 HAZ.

– The influence of hydrogen charging on the plasticity of thermally exposed regions varied in dependence of their individual microstructural characteristics and thermal exposure duration. Although the hydrogen embrittling effects were clearly manifested by gradual decrease in plasticity, the difference in RA values between the hydrogen free and hydrogen charged regions was decreasing with increasing thermal expo-sure duration.

– Fracture modes of the individual regions in PWHT state without hydrogen charging were characterised by ductile dimple tearing which initiated on the sec-ondary phase precipitates and/or inclusions. After subsequent hydrogen charging, all tested regions ex-hibited ductile dimple fracture with some portion of transgranular cleavage.

– The effects of long-term annealing in combination with hydrogen charging caused further manifestations in fracture behaviour. The region of T91 HAZ exhib-ited a mixture of transgranular cleavage and inter-granular dimple fracture, whereas the Ni WM did not show any significant changes. The region of TP316H HAZ was characterised by mixed fracture mode in-cluding intergranular decohesion besides transgran-ular cleavage and small portion of dimple fracture.

– In spite of certain small susceptibility of the weldment T91/TP316H to hydrogen embrittlement during long-term thermal exposure, it can be concluded that this type of dissimilar weldment is quite suitable for high temperature application in hydrogen containing steam environments of fossil-fired power plants.

This work has been financially supported by the Slovak Scientific Grant Agency (VEGA) under the Grant No. 2/0116/13.

Received: May 29, 2013. Accepted: September 21, 2013.


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Juraj Blach and Ladislav Falat* The Influence of Thermal