Effect of filler addition on solidification behaviour and hot tensile properties of GTA-welded tube joints of Super 304H austenitic stainless steel
© Vinoth Kumar et al. 2015
Received: 12 June 2015
Accepted: 9 November 2015
Published: 14 November 2015
Super 304H austenitic stainless steel tubes containing 2.3 to 3 (wt.%) of Cu are mainly used in superheaters and reheaters of ultra super critical boilers due their excellent corrosion and oxidation resistance. Cu addition causes precipitation strengthening effect by fine Cu-rich precipitates which evolve during creep conditions and results in increased creep strength. The microstructural evolution of stainless steels during welding significantly affects the material properties.
In this work, solidification mode microstructure and hot tensile properties of autogenous and filler-added Super 304H gas tungsten arc (GTA) welded tube joints were correlated.
Autogenous welds of Super 304H solidified as austenite with 1.57 % delta ferrite and intercellular (Fe,Cr)23(C,B)6 borocarbides. In filler-added welds, the higher Ni equivalent and addition of carbide stabilizing elements (Nb,Mo) eliminated δ ferrite and segregation of B as borocarbides.
The filler-added welds exhibited superior tensile strength than autogenous welds at both room and high temperature.
Efficiency of the power plant strongly depends on the operating steam temperature and pressure. The concern for reducing the CO2 emission and coal consumption leads to the efforts to build power plants operating at higher steam parameters (Brozda 2007; Shu Ping et al. 2010; David et al. 2013). Austenitic stainless steels are selected for use in sections of superheaters and reheaters, which require good oxidation and fire side corrosion resistance, in addition to the creep strength. The recent improvements in operating steam temperatures up to 620 °C were achieved through the development of creep-resistant stainless steels, containing Cr, Ni, W, Co, Cu and N. Future plans of increasing the efficiency of the power plants up to 49 % require creep-resistant steels operating at steam temperatures of 760 °C and pressures of 35 MPa (Viswanathan et al. 2005). Super 304H austenitic stainless steel is used in recently installed ultra super critical (USC) power units of Japan operating at steam temperature of about 613 °C (Thomas 2013). Further, Super 304H is listed as a candidate material for superheaters and reheaters in the US program for development of materials for USC with steam temperatures of 760 °C (Viswanathan et al. 2005). Super 304H with nominal composition of 0.1C-18Cr-9Ni-3Cu-Nb-N derives excellent creep strength from the distinct Cu addition which precipitates as fine Cu-rich particle under creep conditions, in addition to NbCrN, Nb(C,N) and M23C6 particles (Li et al. 2010).
The different modes of solidification mechanism (A-AF-FA-F) and the methods to predict the volume percentage of ferrite along with the solidification modes are described in detail elsewhere (Valiente Bermejo 2012). A ternary alloy with a composition of Super 304H is expected to solidify as austenite under equilibrium conditions (Hunter and Ferry 2002). In non-equilibrium cooling conditions such as welding, the solidification sequence of the stainless steel depends on the cooling rate and, in highly alloyed steels such as Super 304H, it may result in local compositional variation due to segregation, which alters the solidified phases at room temperature (Darja Steiner et al. 2011).
Recently, few investigations (Indrani et al. 2011; Yang et al. 2006; Li et al. 2010; Bai et al. 2013; Ha and Jung 2012) were carried out to understand the ageing behaviour of Super 304H alloy at high temperatures and evaluated the effect of varying Cu content on microstructure and mechanical properties of the alloy. Works carried out on weldability studies and high temperature properties of Super 304H weld joints reported the tendency of Super 304H to hot cracking and use of matching Super 304H filler metal as an alternate to Ni-based filler metal (Vekeman et al. 2014; Kim et al. 2014). However, the available information on solidification behaviour of gas tungsten arc (GTA) welds of Super 304H and its effects on high temperature properties are very scant. Hence in this work, an attempt has been made to compare the solidification modes, microstructure and hot tensile properties of autogenous and filler-added GTA-welded tube joints of Super 304H austenitic stainless steel. The amount of alloy segregation, composition of segregated elements in autogenous joints, and the role of alloying elements added in the filler metal to avoid detrimental segregation along with their effects on tensile strength of the joints are discussed in this paper.
Chemical composition (wt.%) of parent metal (PM) and filler metal (FM)
Tensile properties of parent metal in as-received condition
0.2 % yield strength (MPa)
Ultimate tensile strength (MPa)
Elongation in 25 mm gauge length (%)
Parameters for GTAW welding of Super 304H
Mode of welding
With filler addition
Power source type
Welding speed (mm/min)
Heat input (kJ/mm)
Percentage of delta ferrite in welds of Super 304H joint
δ ferrite (%)
Average δ ferrite (%)
Standard deviation (%)
1.5, 1.7, 1.6, 1.7, 1.5, 1.7, 1.4, 1.5, 1.6, 1.5
Tensile properties of autogenous and filler-added weld joints of Super 304H
Test temperature (°C)
0.2 % yield strength (MPa)
Ultimate tensile strength (MPa)
Elongation in 25 mm gauge length (%)
302.4 ± 5.5
564.3 ± 7.0
38.96 ± 2.0
201.6 ± 7.4
420.4 ± 10.0
32.09 ± 2.6
173.0 ± 7.4
388.2 ± 7.6
24.97 ± 3.0
206.2 ± 5.3
389.0 ± 8.5
34.70 ± 2.5
349.6 ± 10.2
614.6 ± 11.7
52.3 ± 3.0
247.3 ± 10.6
470.2 ± 10.7
39.2 ± 3.0
196.3 ± 6.7
413.6 ± 12.5
36.6 ± 2.3
240.8 ± 7.6
392.7 ± 13.7
27.3 ± 2.6
The SEM fractographs of filler-added joint tested at RT and 550 °C is shown in Fig. 8e, f, respectively, which reveals a similar behaviour with much lesser reduction in the cross-sectional area of the specimen tested at 550 °C than at RT. The fracture in the filler-added joints was located at the PM, which is evident from the absence of dendritic structure in the high magnification fractographs shown in Fig. 8g, h. The specimen tested at 550 °C reveals coarse and fewer dimples than the specimen tested at RT (refer to Fig. 8g). Flat featureless regions which are much prominent in specimen tested at 550 °C evidence the reduction in elongation of the joint at high temperatures. The Nb-rich precipitates (marked by arrow) associated with the voids in the fracture surface of RT specimen act as the crack initiation site.
The mode of solidification in stainless steels is strongly dependent on the chemical composition which is usually represented as Cr and Ni equivalents (Valiente Bermejo 2012; Brooks et al. 1983). The WRC-1992 constitution diagram Cr/Ni equivalents are commonly used to predict the solidification mode and ferrite content of the weld based on composition of the alloy (Srinivasan et al. 2012). The Cr and Ni equivalent of autogenous weld metal is determined as 18.2 and 14.7 %, respectively. Similarly, the Cr and Ni equivalent for filler-added weld metal is determined as 19.4 and 23.1 %. The welds of both autogenous and filler-added joints are predicted to solidify as primary austenite (A-mode) with no δ ferrite in the WM (as per WRC-1992 diagram). The WM of autogenous weld consists of 1.57 % δ ferrite (refer to Table 4), which is in contrary to the predicted weld metal microstructure.
The autogenous weld of Super 304H first solidifies as δ ferrite near the fusion line due to the peritectic solidification, followed by eutectic single austenite solidification phase evident from the cellular structure of austenite in the weld centre. The austenite continues to grow into the peritectic δ ferrite by solid state transformation, with austenite as the more stable phase. Metalographically hard to distinguish peritectic and eutectic ferrite was retained in the intercellular austenitic/austenitic and austenitic/δ-ferrite boundaries of the autogenous weld (refer to Fig. 2b), as the solid state transformation of δ ferrite to austenite is not complete due to segregation or cooling conditions (Brooks et al. 1983). The addition of B to stainless steel enhances the hardenability and creep strength, however the B in solid solution is beneficial and precipitated borides are detrimental. The segregation of Fe, Cr, B and C (refer to Fig. 4f) in the intercellular region of the austenite reveals the presence of (Fe,Cr)23(C,B)6 borocarbides (Hoffman and ASM 1989; Karlsson et al. 1982). The limited solubility of B in austenitic stainless steel of 30 PPM at 900 °C may cause the excess B to segregate along the grain boundaries and combine with Cr and Fe to form a low melting eutectic with the austenite (Carinci 1994). Such eutectic formation and borocarbide precipitation within the ferrite stringers alter the composition of the solidifying liquid ahead of the austenite/liquid interface and hinder the liquid to solidify as austenite in the intercellular boundaries.
In case of filler-added weld, the WM is fully austenitic with no δ ferrite as predicted (refer to Fig. 5a) and it is attributed to the increased addition of Ni, Nb, Mo and N to the weld metal. The higher Ni equivalent of filler-added WM resulted in A mode of solidification with primary austenite dendrites as single phase (refer to Fig. 5d) (Bonollo et al. 2004). The precipitation of (Fe,Cr)23(C,B)6 was to be suppressed by the Mo, Nb and N addition to the WM of filler-added joint. The formation of (Nb,Mo)C clusters or (Nb,Mo) carbonitrides (refer to Fig. 5e, f, g) prevents the supplement of C to the grain boundaries and retards the formation of detrimental carboborides in filler-added joint (Hara et al. 2004).
The failure of the autogenous weld joint was located at the weld centre regardless of the test temperature, recorded the lowest hardness in the weld joint (refer to Fig. 7). In single pass autogenous weld, the weld metal solidifies as coarse columnar grains by epitaxial growth towards the weld centre with preferential orientation, which may result in a weak centre line in the weld (Villafuerte and Kerr 1990). The higher tensile strength of filler-added joint than the autogenous joint at all test temperatures is attributed to the fine Nb- and Mo-rich carbides. In turn, B retained in the austenitic matrix without precipitation of detrimental carboborides in the WM of filler-added joint. The retention of elemental B in the WM increases the hardenability and thereby resulted in increased strength of the filler-added joint. The reduction in the strength values of weld joints with increase in test temperature is attributed to the accelerated recovery process (Choudhary and Rao Palaparti 2012).
Autogenously welded GTAW joints of Super 304H austenitic stainless steel resulted in weld metal with austenite and δ ferrite (1.57 %). Boron tends to segregate along the intercellular boundaries to form (Fe,Cr)23(C,B)6 borocarbides.
In the filler-added GTAW joints, the increased Mo, Nb and N content suppressed the precipitation of coarse borocarbides by formation of (Nb,Mo)C clusters or (Nb,Mo) carbonitrides and retained elemental B within the matrix to increase the hardenability of the joint.
Autogenous GTAW joints exhibited inferior room temperature and high temperature tensile strength compared to filler-added GTAW joints with increased austenite formers (Ni) and carbide formers (Nb and Mo).
The authors wish to express their sincere thanks to M/s Mailam India Ltd, Pondicherry, India for providing financial assistance to carry out this research work through the Mailam India Research (MIR) Fellowship, M/s Salzgitter Mannesmann Stainless Tubes Italia Srl of Italy for supplying the Super 304H tubes required to carry out this work and The Director, Naval Materials Research Laboratory, Ambernath, Mumbai for providing the facility to carry out hot tensile testing.
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