Microstructure characteristics of LPBF&HIP fabricated graded composite transition joint between ferritic steel and austenitic stainless steel

ABSTRACTCylindrical shells made of stainless steels are widely used, e.g. in tanks and biogas plants [1]. Whereas austenitic stainless steels were commonly used in the past, austenitic‐ferritic and lean duplex stainless steels are more frequently used nowadays. By contrast with structural steels, stainless steels exhibit a significantly nonlinear stress‐strain behaviour which must be considered in the buckling analysis of these shells because the reduced material stiffness below the proof stress may cause premature buckling. For this reason, EN 1993‐1‐6 [2] indicates the use of either a reduced value of elastic modulus Ered or the secant modulus at the 0.2% proof stress when assessing the buckling resistance of shells with nonlinear stress‐strain curves. These provisions may produce over‐conservative low buckling strengths for stainless steel shells at both low and high slendernesses. Only stainless steel shells with medium slenderness are covered by this rule.For this reason, additional temperature‐ and slenderness‐dependent buckling strength correction factors for austenitic stainless steel axially loaded cylindrical shells were developed by Hautala and Schmidt (1999) [3,4] which were described in the ECCS Design Recommendations [5] for buckling of steel shells but are not yet implemented into EN 1993‐1‐6. Recent experimental, numerical and theoretical research, carried out within the European RFCS research project BiogaSS on austenitic‐ferritic and ferritic stainless steels has investigated the potential to apply these buckling strength correction factors to austenitic‐ferritic and ferritic stainless steels. This paper shows that shells made of austenitic‐ferritic and ferritic stainless steels have a less serious deviation from the reference mild steel than austenitic stainless steel, permitting less significant buckling strength correction factors.The results obtained for all families of stainless steels should be transformed into the generalized design verification format of EN 1993‐1‐6, in which a parameterized capacity curve [6,7] is used. This paper also shows how the correction factors devised for axially loaded shells made of austenitic stainless steel can be better represented by an appropriate characterisation using the capacity curve alone. The outcome should be relevant to austenitic, austenitic‐ferritic and ferritic stainless steels when sufficient data is available. The paper also illustrates how behaviours of considerable complexity in some shell buckling situations can still be represented by the unaltered universal capacity curve [6,8].

The influence of strain rate and grain size on the hot ductility of an austenitic and aferritic stainless steel has been examined. Samples were cooled from the austenitising temperature to temperatures in the range 1000 to 700°C and tensile tested at strain rates in the range 10−1 to 10−4 S−1. The austenitising temperature was varied to give two grain sizes, coarse, 600 μm and fine, ∼ 03 μm. For both steels, ductility was excellent at fine grain size throughout the temperature range and strain rates examined. The ferritic and austenitic stainless steels both gave ductility troughs at the coarse grain size, but the trough was favoured by higher strain rates in the ferritic steel and lower strain rates in the austenitic steel. The poor ductility was related to the presence of precipitation, mainly at the grain boundaries; this being FeTi phosphides in the case oftheferritic stainless steel and coarse chromium carbides in the austenitic steel. Grain boundary sliding was the major mode of intergranular failure in the austenitic steel while normal microvoid coalescent failure controlled ductility in the ferritic stainless steel.

The influence of strain rate and grain size on the hot ductility of an austenitic and aferritic stainless steel has been examined. Samples were cooled from the austenitising temperature to temperatures in the range 1000 to 700°C and tensile tested at strain rates in the range 10−1 to 10−4 S−1. The austenitising temperature was varied to give two grain sizes, coarse, 600 μm and fine, ∼ 03 μm. For both steels, ductility was excellent at fine grain size throughout the temperature range and strain rates examined. The ferritic and austenitic stainless steels both gave ductility troughs at the coarse grain size, but the trough was favoured by higher strain rates in the ferritic steel and lower strain rates in the austenitic steel. The poor ductility was related to the presence of precipitation, mainly at the grain boundaries; this being FeTi phosphides in the case oftheferritic stainless steel and coarse chromium carbides in the austenitic steel. Grain boundary sliding was the major mode of intergranular failure in the austenitic steel while normal microvoid coalescent failure controlled ductility in the ferritic stainless steel.

Colossal N supersaturation of ferritic as well as austenitic stainless steels during low temperature gaseous nitridation treatments has lately gained much technological significance. However, available thermodynamic models to calculate the N paraequilibrium solubility limits have failed to explain the levels of colossal N supersaturation observed in several cases of nitrided ferritic/austenitic stainless steels. In this work, we show that consideration of N dissolution induced spinodal decomposition is essential in calculating the N paraequilibrium solubility limit for both ferritic and austenitic stainless steels. This modification in the thermodynamic model has led to the successful explanation of the thermodynamic cause for the colossal N supersaturation in ferritic and austenitic stainless steels. Available experimental observations in literature support the occurrence of spinodal decomposition.

Austenitic stainless steels such as type 304 and 316 are susceptible to stress corrosion cracking in high temperature water environments typical of boiling water reactors (BWR) and pressurized water reactors (PWR). The accumulation of irradiation dose on the austenitic materials increases their susceptibility to environmental cracking. Ferritic stainless steels are less susceptible to irradiation damage including void swelling. Ferritic stainless steels also offer desirable higher thermal conductivity and lower thermal expansion coefficient. Little is known however about the stress corrosion cracking behavior of ferritic steels in high temperature water. Crack propagation rate studies were conducted using four types of wrought ferritic steels (5 to 17% Cr) in high purity water at 288°C containing dissolved oxygen or dissolved hydrogen and also on welded 9Cr ferritic steel. Results show that the ferritic steels are notably more resistant to environmental cracking than the austenitic materials.

This paper reveals the fatigue crack growth behaviour of the shielded metal arc welded AISI 409M grade ferritic stainless steel joints fabricated using austenitic stainless steel (ASS), ferritic stainless steel (FSS) and duplex stainless steel (DSS) electrodes. Centre cracked tensile (CCT) specimens were prepared to evaluate fatigue crack growth behaviour. Servo hydraulic controlled fatigue testing machine with a capacity of 100 kN was used to evaluate the fatigue crack growth behaviour of the welded joints. It is found that the joints fabricated by DSS electrode showed superior fatigue crack growth resistance compared to the joints fabricated by ASS and FSS electrodes. Higher yield strength and relatively higher toughness of the weld metal may be the reasons for superior fatigue performance of the joints fabricated by DSS electrode.

Austenitic stainless steels are providing excellent trouble-free service in sea water for pumps, propellers, valves and other items of marine equipment. Once in a while, a failure occurs as the result of deep localized pitting in a crevice. Data is given showing that austenitic, ferritic and martensitic stainless steels suffer pitting in crevices and under deposits in quiescent seawater. Austenitic stainless steels remain free from attack in high velocity sea water. Low purity ferritic and the martensitic stainless steels frequently pit in high velocity sea water. Crevice corrosion can be effectively controlled with cathodic protection from iron, zinc, aluminum or magnesium galvanic anodes or impressed current cathodic protection by polarization to -0.6 volts versus Calomel. Austenitic stainless steel, in many situations, performs well because it is a component of a multialloy assembly utilizing iron or steel. Examples from field experience are given. INTRODUCTION In the past decade, there has been a growing usage of austenitic stainless steel in marine equipment. Most of the applications have been successful but on occasion an unexpected failure has been observed. It is the purpose of this paper to describe when and how to use austenitic stainless steel with success. The selection of stainless steels appears to result from the engineering requirements of new advanced high speed, high reliability commercial, pleasure and military craft. Ocean science and engineering, offshore oil production, fishing and ocean mining are also contributing to the selection of stainless steels for sea water applications. The increasing usage of stainless steel in the marine environment is found in work boat propellers, pump components, bow thrusters, valves, shafting and shaft components, through hull fittings, parts on data gathering buoys, fasteners and housings of oceanographic instruments. When austenitic stainless steel has given good corrosion free service, it is most often found to be utilized as a key component in a multi component, multi-alloy assembly or system therapy receiving the benefit of built-in cathodic protection. For example, in Figure 1, a cast 304 (Alloy Casting Institute CF-4) propeller is being used on a steel seagoing tugboat with zinc anodes attached to rudder. In Figure 2, a cast ACI CE-30 (1) (2) power plant sea water circulation pump impeller free of any corrosion after six years of service, has been used in combination with an austenitic cast iron suction bell and diffuser. BEHAVIOR OF THE STAINLESS STEELS IN SEA WATER QUIET SEA WATER: A brief review of the behavior of stainless steels in sea water is in order before discussing the criterion to successfully use stainless steels in marine applications. The stainless steels are passive active alloys. In the passive state, stainless steels are highly resistant to corrosion, showing little or no corrosion depending on the exposure. In the active state, when passivity is destroyed, these alloys display localized attack, with little total metal loss, in the form of serious pitting properly identified as crevice corrosion.

Joining of austenitic stainless steel and ferritic stainless steel to sialon

Pitting Corrosion behaviour of similar and dissimilar metal welds of three classes of stainless steels, namely, austenitic stainless steel (AISI 304), ferritic stainless steel (AISI 430) and duplex stainless steel (AISI 2205), has been studied. Three regions of the weldment, i.e. fusion zone, heat affected zone and unaffected parent metal, were subjected to corrosion studies. Electron beam and friction welds have been compared. Optical, scanning electron microscopy and electron probe analysis were carried out to determine the mechanism of corrosion behaviour. Dissimilar metal electron beam welds of austenitic–ferritic (A–F), ferritic–duplex (F–D) and austenitic–duplex stainless steel (A–D) welds contained coarse grains which are predominantly equiaxed on austenitic and duplex stainless steel side while they were columnar on the ferritic stainless steel side. Microstructural features in the central region of dissimilar stainless steel friction welds exhibit fine equiaxed grains due to dynamic recrystallisation as a result of thermomechanical working during welding and is confined to ferritic stainless steel side in the case of A–F, D–F welds and duplex stainless steel side in the case of D–A welds. Beside this region bent and elongated grains were observed on ferritic stainless steel side in the case of A–F, D–F welds and duplex stainless steel side in the case of D–A welds. Interdiffusion of elements was significant in electron beam welding and insignificant in friction welds. Pitting corrosion has been observed to be predominantly confined to heat affected zone (HAZ) close to fusion boundary of ferritic stainless steel interface of A–F electron beam and D–F electron beam and friction weldments. The pitting resistance of stainless steel electron beam weldments was found to be lower than that of parent metal as a result of segregation and partitioning of alloying elements. In general, friction weldments exhibited better pitting corrosion resistance due to lower incidence of carbides in the microstructure.

Dislocations in austenitic and ferritic stainless steels (SSs) under cyclic loading were quantitatively evaluated via X-ray diffraction line-profile analysis to determine the relationship between the dislocation density and low-cycle fatigue (LCF) life in both SSs. The dislocation density of the austenitic and ferritic SSs varied linearly with respect to the LCF life in a double-logarithmic graph, with different slopes of the line. The dislocation density normalized by the maximum work hardening for both SSs exhibited a log–log linear relationship with the LCF life. The fraction of screw dislocations in the ferritic SS decreased with decreasing LCF life owing to the easy cross-slip of dislocations. Because of the difficulty of the cross-slip of dislocations in the austenitic SS, the fraction of screw dislocations remained almost constant throughout the LCF life. Analysis of the crystallite size and the dislocation arrangement with respect to the dislocation density under tensile and cyclic loading revealed that the dislocation arrangement for cyclic loading was smaller than that for tensile loading. Thus, the dislocation arrangement was related to the cyclic loading. In the plot of the dislocation evolution versus the number of cycles, two stages were observed in the variation of the dislocation characteristics for both SSs. In the first stage, the dislocation density increased, and the crystallite size decreased. The dislocation arrangement parameter of the ferritic and austenitic SS decreased and remained the same, respectively, in the first stage. In the second stage, the dislocation density, dislocation arrangement parameter, and crystallite size remained constant.

Practical Investigation of FSS (AISI 430) Weldments Welded by Pulse MIG Welding Process

Transition joints between ferritic steel and austenitic stainless steel are commonly encountered in high-temperature components of power plants. Service failures in these are known to occur as a result, mainly, of thermal stresses due to expansion coefficient differentials. In order to mitigate the problem, a trimetallic configuration involving an intermediate piece of a material such as Alloy 800 between the ferritic and austenitic steels has been suggested. In our work, modified 9Cr-1Mo steel and 316LN stainless steel are used as the ferritic and austenitic components and the thermal behavior of the joints between modified 9Cr-1Mo steel and Alloy 800 is described in this article. The joints, made using the nickel-base filler material INCONEL 82/182 (INCONEL 82 for the root pass by gas-tungsten arc welding and INCONEL 182 for the filler passes by shielded-metal arc welding), were aged at 625 °C for periods up to 5000 hours. The microstructural changes occurring in the weld metal as well as at the interfaces with the two parent materials are characterized in detail. Results of across-the-weld hardness surveys and cross-weld tension tests and weld metal Charpy impact tests are correlated with the structural changes observed. Principally, the results show that (1) the tendency for carbon to diffuse from the ferritic steel into the weld metal is much less pronounced than when 2.25Cr-1Mo steel is used as the ferritic part; and (2) intermetallic precipitation occurs in the weld metal for aging durations longer than 2000 hours, but the weld metal toughness still remains adequate in terms of the relevant specification.

Drawability of stainless steel sheets

Austenitic stainless steels have been used for boiler tubes in power plants. Since austenitic stainless steels are superior to ferritic steels in high temperature strength and steam oxidation resistance, austenitic stainless steel tubes are used in high temperature parts in boilers. Dissimilar welded joints of austenitic steel and ferritic steel are found in the transition regions between high and low temperature parts. In dissimilar welded parts, there is a large difference in the coefficient of thermal expansion between austenitic and ferritic steel, and thus, thermal stress and strain will occur when the temperature changes. Therefore, the dissimilar welded parts require high durability against the repetition of the thermal stresses. SUPER304H (18Cr-9Ni-3Cu-Nb-N) is an austenitic stainless steel that recently has been used for boiler tubes in power plants. In this study, thermal fatigue properties of a dissimilar welded part of SUPER304H were investigated by conducting thermal fatigue tests and finite element analyses. The test sample was a dissimilar welded tube of SUPER304H and T91 (9Cr-1Mo-V-Nb), which is a typical ferritic heat resistant boiler steel.

Effect of weld metal properties on fatigue crack growth behaviour of gas tungsten arc welded AISI 409M grade ferritic stainless steel joints