Met. Mater. Int., Vol. 20, No. 2 (2014), pp. 297~305 doi: 10.1007/s12540-014-1023-z
Modeling and Characterization of As-Welded Microstructure of Solid Solution Strengthened Ni-Cr-Fe Alloys Resistant to Ductility-Dip Cracking Part I: Numerical modeling Jimy Unfried-Silgado1,2,3,* and Antonio J. Ramirez1 1
Metals characterization and processing laboratory, Brazilian Nanotechnology National Laboratory, CNPEM/ABTLuS. Campinas, Brazil. 2 Universidade Estadual de Campinas - UNICAMP. Faculdade de Engenharia Mecânica - FEM, Campinas, Brazil. 3 Universidad Autónoma del Caribe. Programa de ingeniería mecánica. Grupo IMTEF, Barranquilla, Colombia. (received date: 25 August 2012 / accepted date: 25 July 2013) This work aims the numerical modeling and characterization of as-welded microstructure of Ni-Cr-Fe alloys with additions of Nb, Mo and Hf as a key to understand their proven resistance to ductility-dip cracking. Part I deals with as-welded structure modeling, using experimental alloying ranges and Calphad methodology. Model calculates kinetic phase transformations and partitioning of elements during weld solidification using a cooling rate of 100 K.s−1, considering their consequences on solidification mode for each alloy. Calculated structures were compared with experimental observations on as-welded structures, exhibiting good agreement. Numerical calculations estimate an increase by three times of mass fraction of primary carbides precipitation, a substantial reduction of mass fraction of M23C6 precipitates and topologically closed packed phases (TCP), a homogeneously intradendritic distribution, and a slight increase of interdendritic Molybdenum distribution in these alloys. Incidences of metallurgical characteristics of modeled as-welded structures on desirable characteristics of Ni-based alloys resistant to DDC are discussed here. Key words: metals, welding, ductility, fracture, computer simulation
1. INTRODUCTION Several types of solid solution strengthened Nickel alloys (SSS-Ni-based alloys), when subjected to tensile stress at high temperature conditions have tendency to undergo Ductility-Dip Cracking (DDC) [1,2]. DDC is a hot cracking phenomenon, in which occurs a marked drop of ductility and, consequently, intergranular fracture in solid state at homologous temperature range between 0.5 and 0.8 [3-6]. Nevertheless, despite of having several hypotheses about DDC, its fundamental mechanism is not understood yet. Relevant hypotheses about operational mechanism of DDC in as-solidified structures are strongly related to grain boundary sliding [7,8], presence of intergranular precipitates and morphology of grain boundaries [9,10], and they are weakly related to grain boundary embrittlement by impurity segregation [11,12]. DDC is a serious problem during fabrication of SSS-Ni-based alloys components, because it promotes hot microcracking, which *Corresponding author:
[email protected] ©KIM and Springer
could induce to catastrophic failure during service, subsequently [13]. Last decade it has had an important increase of DDC researches in SSS-Ni-based alloys due to potential applications of these materials in nuclear and petrochemical industries. Currently, Ni-Cr-Fe alloys used in above mentioned applications are the 690 alloy and its respective equivalent filler metal AWS ERNiCrFe-7. Nonetheless, these materials present a high susceptibility to undergo DDC [14]. However, despite it is not clear the DDC mechanism, today there are successful developments to counteract this. One of the most important is the chemical composition’s alteration of Ni-based alloys, which has increased their DDC resistance. In this direction, several works have studied the additions of carbide-forming elements, as Nb, Zr, B, V, and Ti, and solid solution hardening elements, as Mo in ERNiCrFe-7 [15-18]. Most of the aforementioned works have discussed the links among DDC, intermetallic precipitates, as-solidified microstructure characteristics, and general chemical composition. Although, it has been possible to identify evidence of beneficial effects provided by carbides primary precipitation on the DDC increase [19,20], there is not a clear explanation of
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role of element additions, the matrix, and its synergic effects on DDC. The most recent hypotheses about DDC mechanism show that, besides the grain boundary morphology, the relationships between hot deformation behavior and element distribution (partitioning) in the crystal lattice are very important [21,22]. Therefore, it is necessary to develop a comprehensive study, involving numerical modeling and several experimental measurements of microstructure characteristics aiming to understand their possible role on DDC, contributing to reveal its fundamental mechanism. The Calphad methodology is a powerful tool based on numerical modeling and analytical calculations, widely used to design, model and calculate numerous alloys [23-26]. Recently, this methodology has been used to optimize the intermetallic precipitation in as-solidified structure of modified alloy ERNiCrFe-7 chemical composition with Ti and Nb additions, aiming to improve properties related to DDC resistance [27]. The present work aims to model and characterize the as-welded structure of a group of alloys based on alloy 690 (Ni-Cr-Fe alloy), which have been modified with Nb, Mo, and Hf additions. The objective is to establish relationships between their metallurgical characteristics and DDC resistance potential. In part I of this investigation are developed three tasks related to modeling procedures and fabrication of selected material. First task is to examine and select adequate alloying ranges of Nb, Mo and Hf, which will be add in experimental alloys. The priorities of this task are minimizing the mass fraction of TCP-phases and maximizing the mass fraction of primary carbides. Second task aims to model the phase transformation of selected alloys during welding. Third task aims to obtain the as-welded microstructure of the experimental alloys, comparing it with modeling results. Part II of this work will deal with detailed metallurgical characterization, emphasizing the discussion on relationships between microstructure characteristics and its role on DDC resistance of these alloys.
Fig. 1. Flowchart shows the procedures of alloying range selection.
2. EXPERIMENTAL PROCEDURES
thermodynamic calculations were: FCC-matrix gamma (γ), FCC-carbonitrides (MX), ordered-cubic carbide M23C6, orderedcubic gamma prime (γ'), Sigma (σ), Laves, alpha Cr-rich (α), orthorhombic Ni3Nb delta (δ), and ordered-cubic gamma double prime (γ''). Used criteria to determine if an alloy system has favorable conditions to DDC resistance were the mass fraction maximization of primary carbides precipitation and mass fraction minimization of TCP-phases. The feedback of calculated thermodynamic results was developed using literature data. Once the alloying ranges have been determined, it was selected a final chemical composition for each alloy system. This procedure was developed through refining the selected alloying ranges and comparison of kinetic calculations results with current as-welded microstructure of fabricated alloys.
2.1. Procedure of alloying ranges selection Figure 1 shows the flow chart, summarizing analytical and operational steps for selection of alloying ranges. Selection of alloying ranges based on Calphad methodology was divided in two procedures. First procedure consisted in calculation of phase stability fields at equilibrium conditions obtaining solidi® fication sequence. This procedure used Thermo-Calc software ® and the commercial database TTNi-8 aiming to simulate the additions of Nb, Mo, and Hf into Ni-Cr-Fe alloy. Firstly, there were established the alloying ranges, in which have been minimized TCP-phases at equilibrium conditions. After that phase stability fields were computed, it has been studied the sequence of solidification, using non-equilibrium solidification based on Scheil-Gulliver model [28]. Involved phases in
2.2. Kinetic numerical modeling Phase transformation during weld solidification of selected alloys group has been simulated using kinetic modeling based on Calphad methodology. The structure of kinetic model was conceived in accordance with the results of solidification sequence studied using Scheil-Gulliver model. The following conditions and tools have been used for kinetic model: A −1 ® ® cooling rate of 100 K.s , Dictra software, TTNi-8 thermodynamic database, MobNi, and Mob2 kinetic databases. Simplified geometric models based on current software conditions were used. Scheil-Gulliver calculations were developed allowing back-diffusion of C and N. Kinetic calculations were divided in two stages. First stage represents the solidification of weld. During
Modeling and Characterization of As-Welded Microstructure of Solid Solution Strengthened Ni-Cr-Fe Alloys
first stage, besides the liquid transformation in solid phases, it was studied the microsegregation of elements during solidification in the matrix phase. The second stage represents solid state phase transformation subsequent to solidification. By simplicity, the second phase precipitation was studied only on segregated region at matrix. The phases considered during kinetic calculations were FCC-matrix gamma (γ), FCC-carbonitrides (MX), ordered-cubic carbide M23C6, ordered cubic gamma prime (γ'), Sigma (σ), Laves and ordered cubic gamma double prime (γ'').Validation of kinetic modeling results was developed using measurements of global and local chemical composition, characterization of as-welded microstructure, and measurements of differential thermal analysis (DTA). As a rough approach, in this work is going suppose that the computed size of precipitates is equivalent to diameter of particle, if this has circular form or equivalent to width of precipitate, if this has elongated form. Solidification ranges were experimentally measured using a DTA/SG Netzsch® model STA409C. The heating and cooling rate used was 10 K.s−1, and maximum measured peak of temperature was 1773 K. There were used Argon atmosphere and three different samples for each measured alloy. 2.3. Fabrication and microstructure characterization of selected alloys Buttons (20 g) of selected alloys were melted in arc furnace, using metal pellets of ultra-high purity and a patternalloy ERNiCreFe-7, which chemical composition is shown in Table 1. Melted buttons were encapsulated, homogenized at 1473 K (1200 °C) during 12 h, and air cooled. Homogenized samples were cut and grinded, obtaining a piece with a plain surface, on which was applied autogenously an arc welding spot using the parameters shown on Table 2. Preliminary characterization of as-welded structure was performed ® using an optical microscope Olympus BX51M coupled to PaX-It image analysis software, and scanning electron micro® scope Jeol JLM 5900, which was operated to 25 kV and coupled with Energy Dispersive Spectrometry (EDS). Criteria for selection procedure during preliminary characterization were the mass fraction of primary carbonitrides and morphology of solidification grain boundaries. Table 1. Chemical composition experimentally measured of ERNiCrFe-7 alloy Ni Bal.
Cr 29.1
Fe 10.3
C 0.027
N 0.015
Al 0.45
Mn 0.28
Ti 0.48
Table 2. Welding parameters used for spot autogenously process application Process GTAW
I (A) 30-35
Gas Argon
Flux (L.min−1) Welding time (s) 5.0 2.0-3.0
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3. RESULTS AND DISCUSSION 3.1. Selection of alloying ranges In accordance with recent hypotheses, the improvement of DDC resistance for Ni-based alloys is related to grain boundary morphology and behavior of crystal lattice during hot deformation [17-22]. The main objective of this section is to select elements and alloying ranges, which will be added to ERNiCrFe-7 alloy with the aim to improve characteristics above mentioned. Several works have shown that additions of Nb and Mo in ERNiFeCr-7 alloy promote a high DDC resistance. These effects have been related with high fraction of carbides precipitates, wavy grain boundaries and expansion of crystal lattice, respectively [17-19,29]. On the other hand, Hf additions promote beneficial effects on hot ductility behavior in diverse wrought Ni-based alloys [30,31]. Both cases do not explain the methodology of alloying ranges selection or the consequences of these additions on the second phase precipitation. The increase of DDC resistance due to Nb additions in ERNiCrFe-7 is associated to Nb-rich carbides precipitation and its effects on pinning of grain boundaries, which increases GB waviness [17-18,27]. However, a previous work has demonstrated that Nb additions higher than 3.0%-wt may produce solidification cracking and promotes the precipitation of undesirable phases, such as gamma-double-prime (γ'') and orthorhombic Ni3Nb (d) delta [27]. Hf additions in ERNiCrFe-7 were studied using numerical modeling. It has been demonstrated that these stimulate the highest fraction of primary carbides precipitates, increasing the wavy grain boundaries. Nevertheless, the additions higher than 1.0%-wt in ERNiCrFe-7 alloys may promote solidification cracking and harmful Hf-rich intermetallic precipitates [32]. Mo additions in Ni-Cr-Fe alloys are limited by Cr and Fe contents and their respective effects on fraction of TCP-phases precipitates, which may deteriorate toughness and formability of these alloys [33]. Ni-based alloys used in this work contain 30%-wt of Cr and 10%-wt of Fe allowing until 5.0%-wt of Mo without produce excessive precipitation of TCPphases and strong microsegregation [33-35]. Taking into account the considerations mentioned above, the preliminary alloying ranges selected in this work are: 0 to 3.0%-wt of Nb, 0 to 5.0%-wt of Mo and 0 to 1.0%-wt of Hf. Figure 2 shows the pseudobinary diagrams displaying the phase stability domains of ERNiCrFe-7 as function of Nb, Mo, and Hf contents at equilibrium conditions for temperatures between 873 and 2073 K. From above mentioned diagrams is observed that Liquidus temperature of ERNiCrFe-7 is increased with Nb, Mo and Hf additions, especially for last-mentioned. In accordance with equilibrium calculations results at 873 K the following phases are expected: Gamma (γ) FCC-matrix, MX primary carbonitrides, and secondary phases, such as cubic ordered-L12 M23C6 carbides, Gammaprime (γ'), sigma (σ), Cr-rich BCC, and delta (δ). It may
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Fig. 2. Pseudobinary diagrams show the phase stability fields as a function of temperature and additions of Molybdenum (a), Niobium (b), and Hafnium (c) for ERNiCrFe-7 alloy. Dashed lines represent specific composition of experimental alloys.
observe that when Hf content increases the M23C6 fraction could be totally suppressed. It is probably that high remaining Hf concentration at microsegregated regions and the presence of Carbon atoms available allow strongly tendency to form Hf-rich carbides [31]. A recent work has been established that Hf additions higher than 0.35%-wt in ERNiCrFe-7 may produce a remaining Hf concentration higher than 5.0%-wt, which may widen the solidification range greater than 673 K [32]. Therefore, in this work the Hf addition will be limited to maximum 0.35%-wt. After selection of alloying ranges, it was determined the solidification sequence, partitioning and phase mass fraction at welding conditions. A first reasonable approximation to solidification sequence is developed using Scheil-Gulliver model [28]. This calculation demands to specify the chemical composition for each alloy system; hence, there have been chosen four Ni-Cr-Fe alloys, with chemical composition
described in Table 3. The selection of these four Ni-Cr-Fe alloys is supported by experimental data about DDC behavior and its narrowed relationship with analysis above carried out [14,17,20-22]. The results of solidification sequence for each alloy are summarized in Table 4 and Scheil-Gulliver modeling results are shown in Fig. 3. Compared to measured solidification range, Scheil-Gulliver results exhibited reasonable agreement for alloy-1. Except for alloy-1, all other systems solidified with a primary eutectic reaction L → γ + MX followed by second phase precipitation, as function of alloy’s chemical composition. Cooling rate used during DTA experiments is lower than real cooling rate of welding; moreover, Scheil-Gulliver modeling supposes that solid phases are emerging from liquid until exhaust it. These mentioned reasons could explain differences between experimental measurements and calculations results. The results of solidification sequence for each alloy were used to develop the kinetic model-
Table 3. Chemical composition of selected alloys (wt%) Alloy Ni Cr Fe C(1) N(2) Ti 1 61.0 29.0 10.0 2 59.2 29.1 10.3 270 280 0.48 3 54.9 27.4 9.6 220 280 0.44 4 55.8 27.4 9.4 250 220 0.50 (s) Bal. 28 - 31.5 7 - 11
