Free Access
Issue
Metall. Res. Technol.
Volume 117, Number 4, 2020
Article Number 407
Number of page(s) 12
DOI https://doi.org/10.1051/metal/2020035
Published online 24 July 2020

© EDP Sciences, 2020

1 Introduction

Nickel base superalloy and Duplex Stainless Steel (DSS) are extensively used in several fields such as chemical, petrochemical, marine and nuclear power industries owing to their excellent combination of corrosion resistance in harsh environments and mechanical properties [15]. DSS is a mixed microstructure of austenite and ferrite which has improved mechanical properties over ferritic and austenitic stainless steel. Their good properties are related to the presence of two phases in approximately equal amounts of austenite γ and δ-ferrite [6]. On the other hand, Inconel 625 (INC 625) superalloy are Ni–Cr–Mo–Nb solid solution strengthening [7,8]. It can be hardened by the precipitation of several varieties of carbides such as MC (M : Nb, Ti), M6C (M : Si, Ni, Cr) and M23C6 (M : Cr) during an adequate thermal ageing [9,10]. However, the presence of Nb element in high concentration decreases the mechanical properties by forming hard and brittle intermetallic phases [11,12].

Dissimilar welding is used to joint two different alloys in order to enhance industrial product. It is well known that the Nickel base superalloys are expensive compared to DSS which provide economic advantages. The particularity of this kind of weld is to associate within the same welding joint a multitude of microstructures. This microstructural heterogeneity leads to direct impact on mechanical and corrosion properties. In this context, countless experimental investigations have been carried out in order to study the metallurgical, mechanical behavior and weldability between stainless steel and Inconel alloy welds [1318].

Devendranath Ramkumar et al. [19] investigated the weldability between Inconel 625 superalloy and UNS S32205 duplex stainless steel using the Electron Beam weld process. The results showed the absence of any secondary phase in the HAZs. Nevertheless, segregation of Mo rich phases at the interdendritic arms of weld zone was observed which caused a loss in toughness and tensile strength. Sridhar et al. [20] studied a dissimilar weld between Inconel 625 and SAF 2205 joined by manual GTA Welding process. Their study relates to the study of the influence of the filler metals used in this case ER2209 and ERNiCrMo-3. The results showed that all the fractures occurred in the weld zone. ERNiCrMo-3 filler metal exhibit high impact toughness and lower strength compared with ER2209 filler metal. A comparative study was carried on dissimilar joint of SAF 2507 and Incoloy 825 by Kangazian et al. [21]. They adopted two techniques during this study Pulsed Current (PCGTAW) and Continuous Current Gas Tungsten Arc Welding (CCGTAW) and used two filler wires ER2594 and ERNiFeCr-1. They reported that the weld zone contain Ti-rich phase when using ERNiFeCr-1 filler metal and Cr-rich phase in the case of ER2594 filler metal. Moteshakker et al. [22] investigated the effect of three filler wires AWS ER 347, AWS ER 316L and AWS ER 309L on the electrochemical behavior in NaCl solution of dissimilar weld between 2205 DSS and AISI 316L stainless steel obtained from GTA welding. They demonstrated that satisfying results were obtained using AWS ER 309L filler metal.

During the solidification in the Inconel alloys, secondary phases precipitated in the interdendritic regions. It is well known that Nickel based superalloy solidify firstly with primary solidification followed by two eutectic reaction [2326]:

Silva et al. [27] presented another solidification path in Inconel 625 alloy. They revealed the presence of another kind of precipitate with cuboidal morphology consisting of an inner core which contain a titanium nitride (TiN) surrounded by carbonitride of titanium and niobium (NbTi)C.

The aforesaid secondary phases have a negative impact on the mechanical and electrochemical behavior. Li et al. [28] reported a significant drop in toughness, tensile strength and corrosion resistance of the weld zone in dissimilar joint between Inconel 625 and SUS 304 obtained from laser welding. The loss of mechanical and electrochemical properties was mainly attributed to the formation of Laves phase and segregation of Nb and Mo in the interdendritic regions. Some investigations [18,20,29] reported that the precipitation of Laves phase in the base nickel superalloy deteriorate mechanical and corrosion properties. Mainly, the previous investigations in dissimilar weld between stainless steel and Inconel superalloy was focalized on the study of microstructural evolution and their impact on the global mechanical and corrosion behavior. Hitherto, few studies have been interested on micromechanical properties in such kind of weld. More information relating mechanical gradient with microstructural evolution should be carried out. The present study attempt to extend the previous works by investigating local heterogeneities of microstructural and mechanical properties in dissimilar weld between Nickel base superalloy and DSS. A nanoindentation measurement is a real field of local characterization by indenting small area and can be used to evaluate local properties with more precision.

In this paper, INC 625 and UNS S32205 DSS are welded by Gas Tungsten Arc Welding (GTAW) process using ERNiCrMo-3 filler metal. Microstructural changes are investigated using Optical Microscopy and Scanning Electron Microscopy. Mechanical properties (hardness and elastic modulus) at micro-scales in each zone of the weld are obtained from the indentation load-depth curves. The purpose of this work is to correlate between microindentation results and metallurgical evolution in each zone of the weld: Base Metal (BM), Heat Affected Zone (HAZ) and Weld Metal (WM).

2 Materials and experimental procedure

The materials studied in this work are UNS S32205 DSS and INC 625 received in plate form of 3 mm thick. The filler wire used for welding is ERNiCrMo-3 as Nickel based filler metal. Its chemical composition is almost similar to INC 625 which prevents overly heterogeneities variation induced during welding process. The nominal chemical composition of base and filler materials is given in Table 1. These base materials were prepared and cleaned by acetone to prevent contaminations before welding which could generate undesirable defects. Two pass weld joint was performed using Gas Tungsten Arc Welding (GTAW) process with argon shielding gas. The details of welding parameters are given in Table 2. Transverse section of the weld was prepared for microstructural observation and indentation tests. The metallographic preparation was carried out using SiC papers in gradual grit size followed by polishing with 1 and 0.5 μm diamond pastes. Finally, fine polishing was performed with 0.05 μm colloidal silica. Then, samples were electropolished using a solution of 98% ethyl alcohol and 2% perchloric acid under 30 V for 60 s at room temperature. Finally, the samples were electrolytically etched with a solution containing 10% oxalic acid and 90% methyl alcohol under 6 V DC for 20 s to create contrast difference. The samples were examined by an optical microscope (OM, Olympus GX51) and a scanning electron microscope (SEM, ZEISS SUPRA 55 VP) equipped with an Energy Dispersion Spectroscopy (EDS) to investigate structural evolution and secondary phases in each region of the weld.

In this study, microindentation tests were performed in each zone of the weld. Mechanical properties such as: hardness H and elastic modulus E are determined using Loubet method [30]. The load was applied using dynamic method known as Continuous Stiffness Measurement (CSM) [31]. Two microindentations grid are carried out across the weld using MTS indenter machine in load control mode. An attentive preparation for sample was carried to obtain a flat area with vibratory polisher (Buhler VibroMet®2) in colloidal silica 50 nm for 10 h of lapping. The indents were performed with Berkovich diamond tip. The indenter was calibrated using a standard fused silica specimen. The grid spacing was taken to be 50 μm between indents to avoid plastic deformation caused by the neighbor indent. The maximum load of 200 mN was applied. Combining nanoindentation measurements with direct observation in SEM was carried out on some location to determine nanohardness. SEM in-situ nanoindentation indents were performed in both Laves phase and austenitic matrix. Cube corner tip was used to perform the indents and the maximum load of 3 mN was applied.

Table 1

Chemical composition of base and filler metals (wt.%).

Table 2

Welding parameter.

3 Results

3.1 Microstructural characterization

The micrographs of the base metals INC 625 and DSS are shown in Figures 1a and 1b. The INC 625 microstructure is fully austenitic with twin boundaries and intermetallic precipitates in blocky form which can be Nb, Mo and Ti carbides. The DSS microstructure is a dual phase containing both austenite γ and δ-ferrite without any other secondary precipitates. The austenite (γ) is like equiaxed islands surrounded by the continuous (δ)-ferrite matrix in approximately equal amounts. The rate balance between ferrite and austenite is approximately 52/48%. The austenite grains are elongated in the rolling direction and easily differentiated from the δ-ferrite by the presence of twins boundaries. Figures 1c and 1d are the WM micrographs at high magnification. The microstructure reveal the presence of both columnar and equiaxed dendrites. Secondary phases in both equiaxed and columnar are founded and identified by SEM as Nb, Ti and Mo rich phase. Figures 1e and 1f illustrated respectively the HAZ of the DSS and HAZ of the INC 625. Columnar growth from the WM is observed with the precipitation of secondary phases in the interdendritic region. However, cellular growth characterizes the HAZ of the INC 625. No secondary phases have been observed in the HAZ of DSS. Interestingly, well defined fusion line between the WM and DSS has been clearly revealed at the interface which could be ascribed to less compatibility between ERNiCrMo-3 filler wire and DSS base metal. Nevertheless, progressive microstructural evolution has been observed in INC 625 with an unmixed zone. The extent of the HAZ of the DSS is much larger comparatively with other side in the HAZ of the INC 625 due to the difference in thermal conductivity (15 W/mK in DSS and 9.8 W/mK in INC 625). Various irregular morphology of austenite is founded in HAZ of the DSS: grain boundary austenite (GBA), Widmanstätten austenite (WA), inter-granular austenite (IGA) and partially transformed austenite (PTA). Near the fusion line an overheating zone can be delimited and characterized by δ-ferrite grain coarsening. However, the HAZ of the INC 625 has no fusion line well distinguished but an unmixed zone having about 400 μm in thickness resembling as small islet (cellular grain) with the formation of some secondary phases.

Figure 2a is the SEM interface between the WM and HAZ of the DSS. A fusion line was clearly distinguished in more details. Grain coarsening of δ-ferrite is well observed near to fusion line. EDS Line mapping analysis shown in Figure 2b was carried out at the interface between the HAZ of DSS and WM in order to underline the mobility of chemical element. The result shows a migration of Fe from the HAZ of DSS to the WM, however the Ni which is a gamagenic element has migrated from the WM to the HAZ of DSS. As seen, a progressive decline in content of Fe element has been observed as an indicator of Fe diffusion from the HAZ of DSS to the weld zone. As far as goes, diffusion of Ni has been occurred from the weld zone to the HAZ of DSS. Also, the result showed that the welded zone presents various peaks indicating a strong heterogeneity and non-uniformity of distribution in chemical elements such as: Fe, Ni, Mo, Nb, Si, Ti and Mn.

Since the WM exhibit heterogeneous microstructure mixing both cellular and columnar dendrite, the distribution of secondary phases is inhomogeneous. Software image J was used to quantify the precipitation of secondary phase based on the difference in contrast. After the image acquisition, the background noise was excluded and the image was transformed to the grayscale. The adjustment of brightness and contrast allow obtaining binary image where the secondary phase appears in the desired color and the matrix remains in white color. The volume fraction and the average size of secondary phase are determined from the SEM microstructure and the results are showed in the Figure 3. The results showed high amount of secondary phases in WM compared to the other regions of weld. Secondary phases precipitation in HAZ of INC 625 is meager in comparison with the WM. Moreover, the columnar dendrite contains more secondary phases than the cellular dendrite. The maximum average size is founded in HAZ of the INC 625. The size of these secondary phases is about 4 μm or less.

The SEM micrographs, in Figure 4, shows the different forms of secondary phases and precipitates which nucleate and growth at the interdendritic regions of the WM. Various morphology are well noticed such as: lamellar, rod shaped and cuboidal. The lamellar precipitates display eutectic morphology. Solidification rate being faster in the weld core than the periphery, the resulting microstructure is cellular including a larger volume fraction of the rod-shaped Laves phases. Examining the two microstructures Figures 4a and 4b, the most widespread morphology in cellular microstructure (Fig. 4a) is the rod shaped form. However, the eutectic morphology is most widespread in columnar microstructure (Fig. 4b). The rod-shaped Laves phases have a small size compared to the eutectic Laves phases which is elongated along the interdendritic region like a chain particles. Laves phase in rod-shaped form is obtained with high cooling rate, however Laves phase in lamellar form is obtained with the slower cooling rate. Figure 4d show another kind of precipitates with cuboidal morphology which is observed in small size less than 1 μm. According to the literature, few studies have been reported their existence in the nickel based alloy. As shown, the cuboidal precipitates looks like a combination of two particles. The shape of inner core is circular with a rectangular planar facet. The EDS mapping (Fig. 5) show a strong segregation of alloying element such as Nb, Mo and Si. The formation of all morphologies of precipitates is caused by the segregation of alloying elements such as: Nb, Mo and Si. The interdendritic regions become enriched with these elements alloying, however the dendrites depleted with the abovementioned elements. Therefore, the distribution of chemical alloying elements will be destabilized and inhomogeneous. Figure 5 is the elemental chemical mapping by EDS technique indicates an enrichment in Nb, Mo, Si and also an impoverishment in Ni, Cr and Fe which lead to identify this precipitates as Laves phase as reported in previous studies [7,26]. As shown in Figure 6, the EDS chemical composition in the cuboidal precipitates shows a strong partition of Ti, Nb and N. The surrounding of this precipitates has a strong contrast comparing with the inner core. The difference in contrast indicates heterogeneity distribution of chemical element. High contrast is a revelator for the presence of heavy chemical element. SEM/EDS analysis was performed to determine chemical composition. The results show an enrichment of Ti, Nb and N. These particles are resulting from solidification process and can be identified as TiN and (NbTi)C.

thumbnail Fig. 1

Optical micrographs of the different zones of the weld: (a) DSS base metal; (b) INC 625 base metal; (c) equiaxed dendrite in WM; (d) columnar dendrite in WM; (e) HAZ of the DSS; (f) HAZ of the INC 625.

thumbnail Fig. 2

(a) Interface microstructure between HAZ of the DSS and WM; (b) EDS line mapping analysis across the HAZ of the DSS and WM.

thumbnail Fig. 3

The average size and volume fraction in all regions.

thumbnail Fig. 4

SEM showing precipitation of secondary phases in: (a) equiaxed weld; (b) columnar weld; (c) different forms of secondary phases; (d) cuboidal precipitate.

thumbnail Fig. 5

SEM/EDS elemental chemical map of eutectic secondary phase.

thumbnail Fig. 6

SEM/EDS chemical composition in cuboidal precipitates.

3.2 Indentation characterization

Figure 7 is the representative microindentation load-indentation depth curves obtained in the WM, HAZ of the INC 625 and DSS, INC 625 and DSS base metal. It’s clearly observed that the WM exhibit the maximum indentation depth with the value of 1782.6 nm, however the lower penetration depth is obtained in the HAZ of the DSS and founded to be 1431.2 nm. In both HAZ and BM of DSS, the δ-ferrite exhibit lower penetration depth than austenite. In consequence, the austenite has the lower nanohardness and elastic modulus than the ferrite. The average values of H and E in the δ-ferrite and austenite γ are presented in Table 3.

The evolution of mechanical properties (H, E) across the weld is presented in Figure 8. A total of six hundred indentations were carried out across the weld. The important number of indents brings a good statistical of mechanical properties. Table 4 summarizes the average values of H and E with standard deviation in all regions of weld. Significant variation in H and E is noticed across the weld joint. It is observed that the WM have the lower values of H and E, however the HAZ of the DSS has the highest values. The maximum value of standard deviation is founded in the HAZ of the INC 625. This dispersion of the results can be attributed to the heterogeneity of the microstructure and non-regular distribution of secondary phases. As seen in Figure 9c, the size of indent is much larger than the size of Laves phase. In this case both the mechanical properties of austenitic matrix and Laves phase are combined and lead to result dispersion.

Figure 9a shows the microindentation impressions in the interface between the WM and HAZ of the DSS. It was seen that each phase of the weld can be characterized and local properties were determined. Figure 9b illustrates the indent impression and nanohardness values in Laves phase and weld matrix. Three nanoindents were carried out in both Laves phases and weld matrix. The average nanohardness in Laves phase was found to be 13.0 GPa which is much higher than the nanohardness of weld matrix 5.1 GPa. The result proves that Laves phase is too much stiffer than the weld matrix. The elevate hardness in Laves phase could be attributed to alloying with Nb and Mo element.

thumbnail Fig. 7

Load versus displacement curves for WM, HAZ of the INC 625, HAZ of the DSS, INC 625 and DSS base metal.

Table 3

The average H and E of ferrite and austenite in DSS and HAZ of DSS.

thumbnail Fig. 8

Evolution of mechanical properties across the weld: (a) microhardness H; (b) elastic modulus E.

Table 4

The average value of H and E in different regions of the weldment.

thumbnail Fig. 9

(a) Optical microindentation showing the grid impression in HAZ of the DSS; (b) SEM nanoindentation impressions in Laves phase and weld matrix; (c) SEM microindentation impression in WM.

4 Discussion

All optical microstructural observations are shown in Figures 1a1f. The morphology of dendritic structure is controlled by the temperature gradient and the solidification rate [32]. The solidification is faster in the middle of WM which resulting the equiaxed structure. However the columnar dendrite is resulted from the decline of temperature gradient. The heat input has an important effect on the microstructural transformations. The lineic energy applied in the present case (1.08 kJ/mm) is relatively high and leads to the decline of temperature gradient. The austenite formation mechanism that occurs in the HAZ of DSS has been fully discussed in the literature. It has been reported that the GBA forms along the prior δ-ferrite grain boundaries in the temperature range between 800 and 1350 °C [33]. In the temperature range between 650 and 800 °C occurs the nucleation of WA at δ/γ interfaces [34]. As reported by Ramirez et al. [35], The IGA forms during reheating δ-ferrite grain in temperature range between 1000 and 1100 °C by heterogeneous nucleation of secondary austenite. PTA in plateau morphology is the remained austenite when the structure is reheated to solidus temperature [33]. The δ-ferrite grain coarsening is related to rapid thermal cooling which prevent the formation of austenite. Also, this phenomenon is due to the temperature increase during welding passes.

The resulted microstructure by the welding operation is, as we have just seen, very heterogeneous; particularly in the junction between the WM and the DSS steel (Figs. 1 and 2). As a result, the mechanical properties are also heterogeneous. Moreover, the discontinuity of the microstructure is accompanied by a discontinuity of the mechanical properties (H, E). The higher average values of H and E are observed in the HAZ of the DSS. This is due to the high amount of δ-ferrite in the HAZ of the DSS (Fig. 1). Indeed, resulting of the thermal cycle, the δ-ferrite content is in the range of 57–63% which induces an increasing of H and E values.

It is evident from the microstructural observation, that both the WM and HAZ of INC 625 contain secondary phases identified as Laves phase (Fig. 5). As reported by some researchers [23,26], Laves phase in lamellar morphology results from the solidification path ending with two eutectic reactions. Laves phases precipitate during the solidification by rejecting Mo, Si and Nb elements in the inter-dendritic liquid. In this case the formation of Laves phase is attributed to the thermal cycle developed during the welding process [36] and the high Nb content. The increase of Fe and Si amount is also favorable to the formation of Laves phase. The content of Fe and Si is respectively 0.82 and 0.14 in the WM. This can be another reason of Laves phase formation. Also, the formation of Laves phase is assigned to the micro-segregation and higher atomic radius of alloying element such as Nb, Ti and Mo which lead to loss of solubility [37]. The morphology of Laves phase is related to the solidification conditions and the microstructure. Formation of discrete Laves phase is related to small equiaxed dendrite spacing under high cooling rate and low temperature gradient. Nevertheless, continuously distributed coarse Laves phase result from large columnar dendrite spacing under low cooling rate and high temperature gradient [38].

Many studies attempt to control the precipitation of Laves phase in Nickel base superalloy [39,40]. It reported that the precipitation of Laves phase is difficult to control. It is pertinent to mention that effective efforts are conducted to limit or suppress alloying elements segregation and Laves phase formation. Different methods are proposed in this context including the use of: (a) optimal weld parameters such as low heat input and fast cooling rates, (b) low Nb content in filler metal, (c) heat extraction techniques, (d) pulsing techniques, (e) post weld heat treatments at proper temperatures, (f) electron beam oscillation techniques [37].

As shown in Figure 4d, the cuboidal precipitates looks like a combination of two particles. The shape of inner core is circular with a rectangular planar facet. The size of these precipitates varies between 0.6 up to 1 µm. The EDS chemical composition in the cuboidal precipitates (Fig. 6) shows a strong partition of Ti, Nb and N. These cuboidal precipitates germinate with a complex mechanism of formation established by Silva et al. [27]. As reported, these precipitates are not simply a carbonitrides. It composed with a Titanium Nitride (TiN) in core surrounded by Niobium Titanium carbide (NbTi)C. SEM/EDS analysis reveal that these cuboidal precipitates contains notably higher Ti content (30.63 wt.%) than the matrix phase (0.27 wt.%). Also, it contains higher Nb content (19.81 wt.%) than the matrix phase (3.00 wt.%).

As seen from the optical observation shown in Figures 1c and 1d, the WM contains both columnar and cellular grains. These microstructural heterogeneities results from the solidification and multi-pass welding. The decline of mechanical properties is attributed specially to the precipitation of Laves phase. As reported by in the previous studies [41,42] Laves phase is brittle intermetallic which has an important effect on mechanical properties. The microstructure evolution and segregation of Nb and Mo element is controlled by the cooling rate [43]. Since the slow solidification and relatively high heat input during the GTA welding, the amount of secondary phases and segregation of Nb and Mo are relatively important which lead to degradation of mechanical properties. These segregations are assigned to the dilution of the ERNiCrMo-3 filler metal by Fe element. It is reported in the literature that the Nickel based superalloy with Mo and Nb elements seriously undergo segregation, owing to the increase in Fe content [44].

In Figure 8 microindentation studies shows that the higher average values of hardness H and elastic modulus E are observed in the HAZ of the DSS. It is well known that the δ-ferrite enhance mechanical properties [45]. The optical micrograph (Fig. 1e) reveal high amount of δ-ferrite in the HAZ of the DSS. The δ-ferrite amount is in the range of 57–63% which induces an increasing of hardness H and elastic modulus E values. However, graphic plots clearly visualized that the microindentation properties at the WM has been plummeted comparatively with the other weld zones and therefore it is considered to be the most vulnerable zone. Lower hardness H can be assigned to the coarse grains and the grains forms such as columnar and cellular dendritic. Moreover, it is also owed to the precipitation of secondary phases with segregation of alloying element such as Mo and Nb as investigated in microstructural characterization. These segregations are ascribed to the dilution of the filler wire by Fe element [44]. Another investigation conducted by Wang et al. confirmed this phenomenon [46].

In fact, Nb element enhances mechanical properties and stabilizes the austenite phase but unfortunately the segregation of Nb results in the formation of Nb-rich phase which has harmful effect on weld mechanical properties. The addition of Nb element increases the fragility of the weld and leads to inter-dendritic fracture. The precipitation of Laves phase was clearly observed in the interdendritic region of WM and consumed most of Nb, thus creating less Nb element available for solution strength.

Sridar et al. [20] performed a series of mechanical test such as tensile, hardness and impact test of this kind of dissimilar weld. Tensile tests have shown that the fracture occurs exactly in the weld joint in brittle mode. Also, they reported that the hardness has been dropped in the weld zone. In another study, similar results were found by Ramkumar et al. [13] which attributed to Mo segregation in interdendritic region caused by less solubility. Improved mechanical properties can be obtained using Pulsed Current GTAW technique which has interesting advantages like grain refinement and reducing segregation.

In a nutshell, the present work articulated the microstructural gradient-micromechanical properties relationships of Inconel 625/UNS S32205 dissimilar welds. Although an extensive study of microstructure and micromechanical properties in all regions of weld was conducted, further investigations are required to bring supplement information of this dissimilar weld. For instance, the relation between macroscopic mechanical properties and micromechanical properties at the microstructure scale (cuboidal particles scale) could be investigated in a deeper way using refined microstructural characterization methods such as Transmission Electron Microscopy.

5 Conclusion

The microstructural evolution and the indentation behavior of dissimilar weld between Inconel 625 superalloy and UNS S32205 duplex stainless steel have been studied. The following important conclusions can be drawn:

  • Microstructural observations showed grain coarsening in HAZ of DSS and precipitation of secondary phases in both WM and HAZ of INC 625 at cellular and interdendritic region in lamellar, rod shaped and cuboidal form. These secondary phases were identified as Laves phase and carbonitrides of Nb and Ti;

  • The rod-shaped Laves phase is widespread in cellular dendrite and takes a small size compared to the eutectic phases Laves which is elongated at the interdendritic region and widespread in columnar dendrite;

  • Cuboidal precipitate is a combination of two particles included an inner core in circular shape and rectangular planar facet;

  • Microindentation tests showed a clear difference between mechanical properties (H, E) for all regions of the weld. The WM had the lowest average values of mechanical properties (H, E), however the HAZ of DSS exhibit the higher values;

  • The evolution of mechanical properties corroborated with microstructural examination, the decline of H and E in WM was attributed to the precipitation of secondary phases;

  • As a result of inhomogeneous distribution and size of secondary phases in the weld zone, a significant standard deviation of hardness was found in WM and INC 625 HAZ generating a wide point cloud;

  • Nanohardness in Laves phase was precisely measured using SEM in-situ nanoindentation technique. It was two and half times greater than the weld matrix due to the high concentration of alloying element such as: Nb and Mo.

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Cite this article as: Ammar Chabbi, Mabrouk Bouabdallah, Sergio Sao-Joao, Achraf Boudiaf, Guillaume Kermouche, Correlation between microstructural gradient and microindentation properties of dissimilar weld between INCONEL 625 and Duplex Stainless Steel, Metall. Res. Technol. 117, 407 (2020)

All Tables

Table 1

Chemical composition of base and filler metals (wt.%).

Table 2

Welding parameter.

Table 3

The average H and E of ferrite and austenite in DSS and HAZ of DSS.

Table 4

The average value of H and E in different regions of the weldment.

All Figures

thumbnail Fig. 1

Optical micrographs of the different zones of the weld: (a) DSS base metal; (b) INC 625 base metal; (c) equiaxed dendrite in WM; (d) columnar dendrite in WM; (e) HAZ of the DSS; (f) HAZ of the INC 625.

In the text
thumbnail Fig. 2

(a) Interface microstructure between HAZ of the DSS and WM; (b) EDS line mapping analysis across the HAZ of the DSS and WM.

In the text
thumbnail Fig. 3

The average size and volume fraction in all regions.

In the text
thumbnail Fig. 4

SEM showing precipitation of secondary phases in: (a) equiaxed weld; (b) columnar weld; (c) different forms of secondary phases; (d) cuboidal precipitate.

In the text
thumbnail Fig. 5

SEM/EDS elemental chemical map of eutectic secondary phase.

In the text
thumbnail Fig. 6

SEM/EDS chemical composition in cuboidal precipitates.

In the text
thumbnail Fig. 7

Load versus displacement curves for WM, HAZ of the INC 625, HAZ of the DSS, INC 625 and DSS base metal.

In the text
thumbnail Fig. 8

Evolution of mechanical properties across the weld: (a) microhardness H; (b) elastic modulus E.

In the text
thumbnail Fig. 9

(a) Optical microindentation showing the grid impression in HAZ of the DSS; (b) SEM nanoindentation impressions in Laves phase and weld matrix; (c) SEM microindentation impression in WM.

In the text

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