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Correlation between microstructure and creep performanceof martensitic/austenitic transition weldment in dependenceof its post-weld heat treatment

Ladislav Falat a,⇑, Lucia C iripová a, Ján Kepic  a, Jir í Buršík b, Ivana Podstranská b

a Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, SK 040 01 Košice, Slovak Republic b Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Z ˇ iz ˇ kova 22, CZ 616 62 Brno, Czech Republic 

a r t i c l e i n f o

 Article history:

Received 13 December 2013Received in revised form 7 February 2014Accepted 25 February 2014Available online 3 March 2014

Keywords:

Transition weldment: T92/TP316HPost-weld heat treatmentMicrostructureCreep failure mode

a b s t r a c t

This paper deals with the influence of post-weld heat treatment (PWHT) of T92/TP316Hmartensitic/austenitic transition weldment on the resulting microstructure and creepcharacteristics. Experimental weldments were fabricated by gas tungsten arc weldingusing a nickel-based weld metal (Ni WM). After the welding, two individual series of produced weldments were heat-treated according to two different PWHT procedures.The first ‘‘conventional PWHT’’ was carried out via subcritical tempering (i.e. bellow Ac 1

temperature of T92 steel), whereas the other one, the so-called ‘‘full PWHT’’ consisted of a complete reaustenitization of the weldments followed by water-quenching and finaltempering. The use of ‘‘conventional PWHT’’ preserved microstructural gradient of T92steel heat-affected zone (HAZ), consisting of its typical coarse-grained and fine-grainedsubregions with tempered martensitic and recrystallized ferritic–carbidic microstructures

respectively. In contrast, the ‘‘full PWHT’’ led to the complete elimination of the originalHAZ via transformation processes involved, i.e. the reaustenitization and back on-coolingmartensite formation. The observed microstructural changes depending on the initialPWHT conditions were further manifested by corresponding differences in the weldments’creep performance and their failure mode. The weldments in ‘‘conventional PWHT’’ stateruptured after long-term creep tests by premature ‘‘type IV failure’’ within their recrystal-lized intercritical HAZs. On the contrary, the long-term creep behavior of the weldmentsprocessed by ‘‘full PWHT’’ was characterized by their remarkable creep life extension butalso by the occurrence of unfavorable ‘‘decohesion failure’’ along T92/Ni WM interface.

 2014 Elsevier Ltd. All rights reserved.

1. Introduction

The 9%Cr martensitic steels are used in power generation industry for thick-walled boiler components such as steamheaders and main steam piping because of the relatively low coefficient of thermal expansion and favorable cost comparedto the high-alloyed austenitic steels [1]. However, the austenitic steels with their excellent corrosion and creep resistance arefrequently used for construction of superheaters  [2]. This indicates that the joining of martensitic and austenitic steels israther necessary in supercritical boiler constructions.

http://dx.doi.org/10.1016/j.engfailanal.2014.02.018

1350-6307/ 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding author. Tel.: +421 55 7922447; fax: +421 55 7922408.

E-mail address: lfalat@imr.saske.sk (L. Falat).

Engineering Failure Analysis 40 (2014) 141–152

Contents lists available at  ScienceDirect

Engineering Failure Analysis

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During the first implementation of transition weldments between ferritic (e.g. tempered bainitic or tempered martens-itic) steels and austenitic steels, the use of austenitic steel type welding consumables was typical [3]. However, further expe-rience with these welds revealed a serious problem related to the carbon diffusion processes across the ferritic/austeniticweld interface, resulting in the formation of soft carbon-depleted zones, the so-called ‘‘white bands’’ at the ferritic side of the welded joints with deteriorated creep strength. To prevent this problem, alternative nickel-based welding filler materialshave been invented [4]. The main benefit of Ni-based weld metal (Ni WM) comes from its low carbon solubility, so it acts as acarbon diffusion barrier. Moreover, thermal properties such as thermal conductivity and expansivity of Ni-based alloys lie ina medium range between the corresponding properties of ferritic and austenitic steels. Thus the use of Ni WM is very suit-able in transition weldments for the lowering and/or moderation of the gradient of residual welding stresses  [5,6].

Many previous investigations, e.g. [7–9] were focused on the Ni-based transition joints between 2.25Cr–1Mo ferritic–bainitic steel and ‘‘AISI 300 series’’ austenitic steels. Recently published studies [2,10–15] include the results obtained forthe martensitic/austenitic weldments involving 9%Cr martensitic base materials. Nevertheless, literature concerning theeffects of the variation of post-weld heat treatment (PWHT) conditions on microstructure and creep behavior of martens-itic/austenitic weldments is very limited. In general, the weldments of martensitic steels require an application of PWHTin order to relieve residual stresses and stabilize the microstructure with strengthening precipitates  [16–18]. Conventionalway of PWHT of these weldments consists of subcritical tempering with respect to the steels’ Ac 1 temperature, typically inthe range from 720 to 760 C [16]. It is also well-known that the ‘‘conventional PWHT’’ improves the weldments’ toughnessbut the remaining problem is their premature ‘‘type IV failure’’ in the heat-affected zone (HAZ) during operation in creepconditions [1,19,20]. Albert et al. [19] concluded that this failure mode cannot be suppressed by any variation of subcriticalPWHT conditions. Recently, Abe et al. [21] found out that the welded joints of newly developed 9Cr–3W–3Co–0.2V–0.05Nbsteel with 160 ppm boron and 85 ppm nitrogen exhibited no ‘‘type IV failure’’ in relatively short to medium-term creepconditions as a result of specific modification of HAZ microstructure. However, long-term creep tests of this new promisingmartensitic steel with modified  B  and N  contents (from there  ‘‘MARBN’’ steel) are still in progress [21]. On the other hand,Tezuka and Sakurai [22] and Kimura et al. [23] suggested that a possible way to avoid ‘‘type IV failure’’ of the weldments istheir full renormalization. In a specific case of the weldments between martensitic and austenitic steels, PWHT conditionsare commonly specified according to the ‘‘conventional’’ (subcritical) procedure with regard to the martensitic base materialof welded joint   [24]. However, available information about the application of ‘‘full PWHT’’ for martensitic/austeniticweldments is rather scarce [25–27].

The present study represents an extended and continuing research work of the former study by Falat et al.  [25]. It sum-marizes the results regarding the effects of ‘‘conventional’’ as well as ‘‘full PWHT’’ on microstructure and creep performanceof T92/TP316H transition weldments.

2. Experimental procedure

Dissimilar steels T92 and TP316H in the form of tubes with outer diameter of 38 mm and wall thickness of 5.6 mm werewelded by gas tungsten arc welding (GTAW) using Ni-based filler alloy Nirod 600. The electrode diameter was 2.4 mm andthe applied welding parameters were: welding current 70–110 A, voltage 12–17 V and heat input 9–12 kJ/cm. Chemicalcompositions of the individual materials used for fabrication of T92/TP316H weldments are listed in  Table 1.

After the GTAW process, two different PWHT procedures were applied to the first and second series of the produced weld-ments respectively. The ‘‘conventional PWHT’’ consisted of subcritical tempering at 760 C for 1 h, followed by air cooling infurnace to room temperature. The second series of the weldments was subjected to the so-called ‘‘full PWHT’’ including acomplete reaustenitization at 1060 C for 15 min with subsequent water-quenching and subcritical tempering at 760 Cfor 1 h and air cooling down in furnace. Detailed PWHT diagrams of both the used procedures are shown in Fig. 1.

All experimental work was performed using cross-weld (c-w) samples. Prepared tubular weldments were cut into the c-wblocks, as schematically shown in Fig. 2. The creep tests were performed using cylindrical tensile samples with a gaugelength of 40 mm, body diameter of 4 mm, and M6 head thread. Metallographic analyses involved light microscopy (LM),

scanning electron microscopy (SEM) with energy dispersive X-ray (EDX) spectroscopy, and transmission electron micros-copy (TEM). Etched metallographic samples were used for LM and SEM analyses. The used etching solutions were specifiedin [25]. Thin foils for TEM observations were prepared using focused ion beam (FIB) technique, applied perpendicularly ontothe longitudinally sectioned tensile creep specimens at the locations immediately beneath their creep fractures. Thermody-namic calculations of phase equilibria were performed using the software Thermo-Calc [28] and thermodynamic databaseSTEEL16 formulated by Kroupa et al.  [29].

 Table 1

Chemical composition (wt.%) of individual materials of the dissimilar weldment.

C N Si Mn P S Cr Mo W B Ni Al V Nb Fe

T92 0.11 0.056 0.38 0.49 0.019 0.002 9.08 0.31 1.57 0.0023 0.33 0.014 0.2 0.069 BalanceNirod 600 0.05 – 0.3 3.0 0.03 0.015 20.0 – – – Balance – – 2.0 2.0

TP316H 0.052 – 0.51 1.77 0.031 0.006 16.76 2.05 – – 11.13 – – – Balance

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3. Results and discussion

 3.1. The effect of PWHT conditions on microstructure and hardness

Figs. 3–6 show the microstructures of individual regions of the studied T92/TP316H weldments after the application of different PWHT procedures. In Fig. 3a and b the microstructures of martensitic (T92) region close to the Ni WM are shownin conditions after the ‘‘conventional’’ and ‘‘full PWHT’’ respectively. The heat-affected zone (HAZ) of the T92 base material(T92 BM) consists of coarse-grained HAZ (CGHAZ), fine-grained HAZ (FGHAZ), intercritical HAZ (ICHAZ) and subcritical HAZ(SCHAZ) subregions as a result of phase transformations taking place in T92 BM in dependence of a local temperaturereached during the welding (see Fig. 3a). The subregions of ICHAZ and FGHAZ are both essentially fine-grained because of a very short duration of   a? c   matrix transformation processes hindering the growth of austenite grains   [16]. On thecontrary, microstructure in HAZ of T92 BM became coarse-grained and homogeneous after the ‘‘full PWHT’’ (see Fig. 3b)as a result of transformation processes involved during the complete reaustenitization and back on-cooling fresh martensiteformation. This retransformed martensitic microstructure shows significant coarsening with respect to its prior austeniticgrains, even in comparison to the CGHAZ microstructure after the ‘‘conventional PWHT’’ (see the details in Fig. 4). Althoughmicrostructure of T92 steel after different PWHT procedures differs significantly in the prior austenite grain size (see Fig. 4),

its phase composition in both PWHT states is basically the same, consisting of tempered martensite, i.e. ferritic–carbidic

Fig. 1.  Schematic time–temperature diagrams clarifying the conditions of individual PWHT procedures: ‘‘conventional PWHT’’ (a) and ‘‘full PWHT (b).

Fig. 2.  Schematic sampling of cross-weld (c-w) blocks from the tubular weldments.

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Fig. 3.  Microstructures of martensitic (T92) region of transition T92/TP316H weldment in conditions after the ‘‘conventional PWHT’’ (a) and ‘‘full PWHT’’(b).

Fig. 4.  Detailed microstructures of the martensitic region of the T92/TP316H weldment in conditions after the ‘‘conventional PWHT’’ (a) and ‘‘full PWHT’’(b).

Fig. 5.  Microstructures of Ni WM (Nirod 600) near the interface with T92 steel in conditions after the ‘‘conventional PWHT’’ (a) and ‘‘full PWHT’’ (b). After

the selective etching of Ni WM, marginal T92 regions appear dark due to their over-etching.

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mixture. In accordance with the literature  [16,30], the microstructures of Grade 92 steels typically contain two majortypes of precipitates, namely the intergranular M23C6   (M = Cr, Fe) carbides and intragranular MX (M = V, Nb; X = C, N)carbonitrides.

The microstructures of Ni WM (Nirod 600) in the vicinity of their interface with T92 steel are shown in both PWHT statesin Fig. 5. The microstructures of Nirod 600 weld metal are markedly heterogeneous with respect to the size, morphology anddistribution of grain boundaries and precipitates. In the state after the ‘‘conventional PWHT’’ (Fig. 5a) the Ni WMmicrostruc-ture revealed well recognizable solidification grain boundaries and dendritic cell sub-boundaries. On the contrary, after the‘‘full PWHT’’ (Fig. 5b) the microstructure of Nirod 600 contains elongated grains with clearly visible straightened boundariesas a result of their diffusional migration during the weldment austenitization and carbide precipitation. The size of precip-itates in Ni WM after the ‘‘full PWHT’’ (Fig. 5b) is markedly smaller in comparison to that after the ‘‘conventional PWHT’’(Fig. 5a). This refinement of the particles is likely to originate from the dissolution of precipitates during reaustenitizationand reprecipitation during subsequent subcritical tempering. The microstructures in  Fig. 5 contain intragranular as wellas intergranular precipitates with different sizes and morphological appearance. According to the previous investigation[25], both intragranular and grain-boundary precipitates were found to be rich in Nb by means of EDX analyses. This result,with respect to the overall Ni WM composition, indicated these precipitates to be the NbC carbides. The finest precipitateagglomerations were analyzed using EDX and electron diffraction on carbon extraction replicas and they were also foundto be the Nb-rich MC type precipitates with face-centered cubic crystal structure  [25].

In Fig. 6 the microstructures of austenitic steel (TP316H) region adjacent to the Ni WM are compared for the ‘‘conven-tional’’ and ‘‘full PWHT’’ state. With respect to the grain size, the austenitic microstructures (Fig. 6) were generally lessaffected by PWHT conditions compared to the martensitic ones (Fig. 4). Etching revealed a narrow band of coarse grainsfragmented by dendritic cell sub-boundaries along the fusion line in conditions after conventional PWHT (see  Fig. 6a). Afterthe ‘‘full PWHT’’, the dendritic band within the fusion zone disappeared as a result of homogenization and the rest of austen-itic microstructure became slightly coarser (Fig. 6b). Both austenitic microstructures contain recrystallized polygonal grainswith mostly intergranular M23C6 (M = Cr, Fe) carbides which are known to be the major type of precipitates in non-stabilizedaustenitic steels [31].

In order to evaluate the influence of different PWHT conditions on the local mechanical properties of the studied weld-ment, cross-weld hardness measurements were carried out on polished metallographic samples (Fig. 7). The hardness profile

Fig. 6.  Microstructures of austenitic (TP316H) region of the T92/TP316H weldment in conditions after the ‘‘conventional PWHT’’ (a) and ‘‘full PWHT’’ (b).

Fig. 7.  Cross-weld hardness profiles of the T92/TP316H weldment after different PWHT procedures.

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of T92/TP316H weldment in the state after the ‘‘conventional PWHT’’ exhibits a steep hardness gradient within the region of T92 HAZ which reasonably reflects the local microstructural characteristics of the individual T92 HAZ subregions ( Fig. 3a).The highest hardness was measured close to the fusion zone (FZ) of the T92 steel and Ni WM. This hardness peak can berelated to the transformation hardening, i.e. ‘‘fresh’’ martensite formation in the FZ during cooling after the welding  [16]as well as to the additional alloying originated from a mixture of partially melted T92 steel and Nirod 600 weld metal. Trans-formation hardening effects were also strongly manifested within the region of T92 CGHAZ heated up to the temperaturesbelow the range of solidification temperatures but still well above the Ac3 transformation temperature (see Fig. 8). Furtherdecreasing of the reached peak temperature of T92 HAZ down to the Ac3 at increasing distance from the weld FZ resulted in agradual decrease in hardness. In this region a short heating period allowed the austenite formation without its saturationwith alloying elements from the secondary carbide precipitates. Moreover, the undissolved carbides caused efficient hinder-ing of grain growth within the T92 FGHAZ. The lowest hardness values were measured within the region of T92 ICHAZheated up to the temperatures between Ac1  and Ac3  as a result of severe microstructure degradation during the welding

thermal cycle. Primarily, this deterioration was caused by the grain-refinement originating (as in the case of T92 FGHAZ)from the non-saturated austenite formation during short-term heating and subsequent on-cooling formation of fine polyg-onal ferrite instead of lath martensite. The second contribution to the microstructure degradation of T92 ICHAZ was relatedto the additional tempering of the remaining untransformed (originally tempered) martensite, i.e. the so-called ‘‘double-tempering’’ effect [16]. The hardness values of martensitic (T92) part of the studied weldment in the state after the ‘‘fullPWHT’’ are mostly lower, compared to those of the weldment in ‘‘conventional PWHT’’ state. In addition, the original hard-ness gradient of the former T92 HAZ was completely suppressed after the ‘‘full PWHT’’ and the hardness within the wholeT92 region was invariable (see Fig. 7). This behavior can apparently be related to the elimination of the former T92 HAZ dur-ing the reaustenitization period of the ‘‘full’’ PWHT (Fig. 3b). On the other hand, the increased hardness values of Ni WM andTP316H regions after the ‘‘full PWHT’’ can likely be associated with redistribution of alloying elements and subsequent rep-recipitation of fine particles during the tempering period of this heat treatment (Fig. 5b and Fig. 6b).

 3.2. The effect of PWHT conditions on creep life and failure mode

Fig. 9 shows the dependencies of creep-rupture life (time to fracture) of the studied weldment on the applied stress andused PWHT conditions for two creep testing temperatures. With increasing stress and/or temperature the fracture time of the weldments in both the PWHT states decreases. In addition,  Fig. 9 indicates that for all the applied stresses at a givencreep temperature, the fracture times of the weldments in the state after the ‘‘full PWHT’’ are always higher than thoseof the weldments after the ‘‘conventional PWHT’’. However, with decreasing applied stress the difference in creep rupturelife between the weldments in different PWHT states decreases. This effect is more significant at a higher creep temperature.

Fig. 10 clarifies the effects of creep conditions (applied stress and creep temperature) and the used PWHT procedure onthe weldments’ failure mode [25]. Metallographic samples of cross-weld sections of ruptured specimens indicating typicalfailures for ‘‘conventional’’ and ‘‘full PWHT’’ conditions are shown in  Figs. 11 and 12 respectively. In accordance with Fig. 10it can be stated that the long-term creep failure mode of the weldments processed by ‘‘conventional PWHT’’ was the ‘‘type IVfailure’’ within fine-grained T92 ICHAZ (see also  Fig. 11b). On the other hand, the long-term creep failure mode of theweldments processed by ‘‘full PWHT’’ was characterized by the ‘‘decohesion failure’’ along T92/Ni WM interface (see Figs. 10and 12b). Significant differences in local plasticity close to the fracture were observed for different specimens. Therefore

Fig. 8.  Schematic formation of individual T92 HAZ subregions in correlation to the equilibrium phase diagram calculated using software Thermo-Calc  [28]and thermodynamic database STEEL16 [29]. The T92 steel composition is indicated by the corresponding isopleth at 0.11 wt.% C (vertical dash-line in thephase diagram).

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Fig. 9.  Creep fracture times of the T92/TP316H weldments in dependence of the applied stress, initial PWHT procedure and creep testing temperature:625 C (a) and 650 C (b). Creep-rupture data according to  [25].

Fig. 10.  Occurrence of different failure mode types in T92/TP316H weldments in dependence of the used creep conditions and initial PWHT procedure:‘‘conventional PWHT’’ (a) and ‘‘full PWHT’’ (b).

Fig. 11. The creep failure types of the T92/TP316H weldment processed by the ‘‘conventional PWHT’’: abrupt ‘‘over-load failure’’ at T92/Ni WM interface(a)

and ‘‘type IV failure’’ in T92 ICHAZ (b).

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reduction area (RA) was evaluated in dependence of time to rupture in order to assess the susceptibility to local embrittle-ment of the weldments (Fig. 13). For both the PWHT procedures, the measured RA values of T92/TP316H weldments show adecreasing tendency with increasing creep fracture time. The only deviation from this rule was observed for the creep testwith the shortest time to rupture of weldment after ‘‘conventional PWHT’’ ( Fig. 13a), which revealed ‘‘over-load failure’’ atT92/Ni WM interface (Fig. 11a). However, a general comparison of RA values corresponding to the weldments in differentPWHT states indicates that the application of ‘‘full PWHT’’ results in much greater local embrittlement than that of the weld-ments processed by ‘‘conventional PWHT’’. In accordance with Fig. 11b it is obvious that the ‘‘type IV failure’’ occurring inT92 ICHAZ region was still characterized by considerable plasticity at fracture, whereas the ‘‘decohesion failure’’ at T92/NiWMinterface (Fig. 12b) was completely brittle. This finding discredits the benefits of creep life extension via the ‘‘full PWHT’’performed under the conditions of the present study.

 3.3. Correlation between microstructure degradation and creep failure

Fig. 14 shows detailed SEM-images of T92 ICHAZ microstructure near the fracture surface after the ‘‘type IV failure’’. It isgenerally known [32,33] that the microstructure degradation in this region is caused by the welding process and subsequent

Fig. 12.   The creep failure types of the T92/TP316H weldment processed by the ‘‘full PWHT’’: ‘‘ductile failure in T92 BM’’ (a) and brittle ‘‘decohesion failure’’at T92/Ni WM interface in long-term creep conditions (b).

Fig. 13.  Reduction area (RA) of the T92/TP316H weldment depending on the time to fracture, PWHT procedure and creep temperature: 625 C (a) and650 C (b).

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creep exposure. Detrimental effects during the creep are enhanced by gradual decrease in dislocation density and coarseningof secondary phase precipitates [34,35]. Microstructural damage accompanied by nucleation and growth of creep voids ispronounced by a multiaxial stress state [20,32] originating from the mechanical constraint in T92 HAZ region with lowercreep strength surrounded by the regions with higher creep strength (i.e. base material and weld metal). The microstructurein Fig. 14 clearly reveals a crucial influence of coarsened grain-boundary precipitates on the intergranular fracture charac-teristics of this failure type. The individual precipitates were differentiated using an ‘‘atomic-number contrast’’ of back-scattered electrons (see Fig. 14b).

Fig. 15 shows detailed SEM-images of the T92/Ni WM interface after the final ‘‘decohesion failure’’. It is obvious that theobserved interfacial decohesion is closely related to a progressive growth and interlinking of individual precipitates alongthe T92/Ni WM interface. With respect to their chemical composition, these interfacial precipitates are mostly M23C6

carbides and occasionally also Fe2(W,Mo) based Laves phase. An occurrence of interfacial precipitates in several types of transition weldments has also been reported in many other studies, e.g. [6,36,37] and it has been ascribed to the alloyingelements’ redistribution across the weld metal interfaces  [38,39].

With respect to the observed differences in microstructures and creep failure types of the studied T92/TP316H weld-ments, cross-weld hardness measurements were carried out to estimate their local mechanical properties in dependenceof initial PWHT procedure and subsequent creep exposure. Typical hardness profiles of the weldments after the individualcreep failures are shown in Fig. 16. The general hardness profiles’ tendencies of creep-exposed weldments (Fig. 16) aremostly very similar to those of the weldments in their corresponding PWHT states without creep exposure ( Fig. 7). Thus,clear differences between the hardness profiles corresponding to the individual PWHT states remain preserved after thecreep expositions. Irrespective of applied creep conditions, the hardness profiles of the weldments processed by ‘‘conven-tional PWHT’’ show a steep hardness gradient within their T92 HAZ regions (Fig. 16a). In comparison to the correspondinginitial PWHT state (Fig. 7), the hardness gradients within T92 HAZs are even more pronounced after the creep testing(Fig. 16a) as a result of continuing microstructure degradation involving the effects of precipitate coarsening as well asgradual decreasing in dislocation density of tempered martensite due to the dynamic recovery processes  [15,16].

The hardness profiles of creep-exposed weldments initially processed by the ‘‘full PWHT’’ ( Fig. 16b) differ vastly fromthose of the weldments initially processed by the ‘‘conventional PWHT’’ (Fig. 16a). The hardness values within T92 regionafter the ‘‘full PWHT’’ do not exhibit any strong variations among the individual creep-exposed states. In the case of 

Fig. 14.  SEM-images of typical T92 ICHAZ microstructure in the region of ‘‘type IV creep failure’’ of T92/TP316H weldment, initially processed by the‘‘conventional PWHT’’: secondary electrons (a) and back-scattered electrons (b) micrographs.

Fig. 15.   Detailed SEM-images of typical T92/Ni WM interfacial microstructure in the region of ‘‘decohesion failure’’ of T92/TP316H weldment, initiallyprocessed by the ‘‘full PWHT’’: secondary electrons (a) and back-scattered electrons (b) micrographs.

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long-term creep tests, the creep failure locations correspond reasonably with their most significant microstructural damageat the T92/Ni WM interface (Fig. 15).

Since the long-term creep failures of the investigated weldments always occurred in their martensitic parts, i.e. either inT92 ICHAZ or at T92/Ni WM interface, these regions were subsequently subjected to the detailed substructural characteriza-tion (see Fig. 17). The TEM observations were performed on creep-exposed samples with comparable creep-rupture lives forthe individual PWHT states. Sampling procedure for TEM was carried out using FIB technique in the locations immediatelybeneath individual creep fractures. In the case of creep-exposed weldment, initially processed by the ‘‘conventional PWHT’’,the T92 ICHAZ region was investigated (see Fig. 17a). On the other hand, in the case of creep-exposed weldment, initiallyprocessed by the ‘‘full PWHT’’, the TEM analysis was performed within T92 steel region next to the T92/Ni WM interface(Fig. 17b). The substructure of T92 ICHAZ region after the creep exposure for 4060 h at 625 C, 100 MPa shows completely

recrystallized polyhedric grains with significantly coarsened precipitates (Fig. 17a). On the contrary, the substructure of T92interfacial region after the creep exposure for 5143 h at 625 C, 120 MPa exhibits lath-like tempered martensite with muchfiner precipitates (Fig. 17b) in comparison to the previous case (Fig. 17a). Thus, the observed differences between individualsubstructures (Fig. 17) of different T92 regions of the studied weldments correspond well with the differences in their creepperformance and failure mode characteristics in dependence of the used PWHT procedure. Despite the observed beneficialeffects of the used ‘‘full PWHT’’ on the creep life increase of the studied T92/TP316H weldments via ‘‘type IV failure’’suppression, it is necessary to search for possibilities of creep failure mechanism modification from the completely brittleinterfacial decohesion towards less brittle failures.

4. Summary and conclusions

Variations in microstructures and creep performance of T92/TP316H martensitic/austenitic weldment were investigateddepending on the applied conditions of post-weld heat treatment (PWHT). The obtained results are summarized in thefollowing conclusions:

Fig. 16.  Cross-weld hardness profiles of the T92/TP316H weldments ruptured by individual creep failure modes in dependence of creep conditions andinitial PWHT procedure: ‘‘conventional PWHT’’ (a) and ‘‘full PWHT’’ (b). Interruptions in the hardness profiles indicate the locations of individual creepfailures.

Fig. 17.   Substructural characteristics of: T92 ICHAZ of the T92/TP316H weldment ruptured by type IV creep failure after 4060 h at 625 C, 100 MPa (a) and

T92 region near the T92/Ni WM interface of the weldment ruptured by ‘‘decohesion failure’’ at T92/Ni WM interface after the creep exposure at 625 C,120MPa for 5143 h (b).

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 After the ‘‘conventional PWHT’’ the microstructure of the T92 base material adjacent to the interface with Ni-based weldmetal (Ni WM) shows a typical microstructural gradient. On the contrary, after the ‘‘full PWHT’’ the part of weldmentmade of T92 steel does not contain such a structured HAZ as a result of complete reaustenitization. In addition, the micro-structure is significantly coarser with respect to prior austenite grain size in comparison to the original HAZ.

 Solidification grains with intergranular and intragranular NbC precipitates occur in microstructure of Ni WM (Nirod 600)in conditions after the ‘‘conventional PWHT’’. After the ‘‘full PWHT’’ the microstructure of Ni WM is significantly affectedby diffusive migration of grain boundaries and the processes of NbC carbides’ dissolution and their back reprecipitationduring the individual stages of performed heat treatment. This resulted in a significant increase in density of precipitates.

 The austenitic microstructures of TP316H steel base material are generally less affected by PWHT conditions in compar-ison to the martensitic microstructures of T92 base material. However, after the ‘‘full PWHT’’ the microstructure ishomogenized and slightly coarsened in grain size.

 In long-term creep conditions, the weldments with the initial application of ‘‘conventional PWHT’’ procedure rupturedprematurely by the ‘‘type IV failure’’ mode within their T92 ICHAZ regions. On the contrary, the weldments initially pro-cessed by the ‘‘full PWHT’’ ruptured after their long-term creep tests by the ‘‘decohesion failure’’ within the fusion zonealong T92/Ni WM interface.

 Regardless of the applied creep conditions, the hardness profiles of the weldments initially processed by the ‘‘conven-tional PWHT’’ always showed steep hardness gradients within their T92 HAZ regions. Moreover, these gradients of creep-exposed weldments were even more pronounced, compared to the hardness profile corresponding to the initialPWHT state. In contrast, the hardness values within T92 regions of the weldments initially processed by the ‘‘full PWHT’’did not exhibit any strong variations among the individual creep-exposed states.

 In the case of long-term creep tests, the creep failure locations of the weldments initially processed by the ‘‘conventionalPWHT’’ were associated with the most remarkable microstructural deterioration of T92 ICHAZ regions with the lowestmeasured hardness values within the whole T92 regions of the investigated weldments. On the other hand, the creep fail-ures of the long-term creep-exposed weldments initially processed by the ‘‘full PWHT’’ were related to their significantmicrostructural damage at the T92/Ni WM interfaces, but their eventual correlations with performed cross-weld hardnessmeasurements were not revealed by the used experimental approach.

 Comparison of substructural characteristics from the individual creep failure locations indicated their good correlationswith creep life of the weldments processed by different initial PWHT procedures. Specifically, the substructure of T92ICHAZ region after the long-term creep exposure exhibits fully recrystallized polygonal grains with markedly coarsenedprecipitates. In contrast, after the comparable long-term creep exposure, the substructure of T92 interfacial region withthe Ni WM still contains highly tempered martensitic laths with relatively finer precipitates.

 The creep life of weldment processed by the ‘‘full PWHT’’ was significantly longer than that of the weldment processed bythe ‘‘conventional PWHT’’ procedure. The observed creep life extension of the weldments was directly related to their‘‘type IV failure’’ elimination, due to the complete suppression of their former T92 HAZ microstructural gradient duringthe ‘‘full PWHT’’. However, the occurrence of unfavorable ‘‘decohesion failure’’ along T92/Ni WM interface in long-termcreep conditions apparently discredited the achievements in creep performance of the studied martensitic/austeniticweldments obtained by the ‘‘full PWHT’’.

 Acknowledgement

Financial support by the Slovak Scientific Grant Agency (VEGA) under the Grant No. 2/0116/13 is gratefully acknowledged.

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