Investigating the microstructure and mechanical properties of deposited metal with ENiCrFe-7 covered electrode

The effect of post weld heat treatment (PWHT) on microstructure and mechanical properties of deposited metal using ENiCrFe-7 covered electrode was investigated. The PWHT was conducted at 615 °C for 24 h. The results showed that three types of precipitates existed in both as-welded and PWHT deposited metals, i.e., NbC carbides, Al-Ti oxides in grains, and M23C6 carbides on grain boundaries. The M23C6 carbides coarsened after PWHT. The room temperature and elevated temperature (350 °C) tensile strength decreased by about 30 MPa, and elongation increased by about 3%–4% after PWHT. The tensile specimens presented ductile fracture feature and the dimples in PWHT condition were larger than that in as-welded condition. The hardness of deposited metal in as-welded condition was higher than that of PWHT condition. The impact toughness showed no significant change after PWHT. Both as-welded and PWHT impact specimens showed a mixed fracture mode with dominant ductile fracture and cleavage fracture.


Introduction
Nickel-based alloy with high temperature strength, excellent corrosion resistance, and good microstructure stability, is widely used in pressurized water reactors (PWR) in the third generation nuclear power plants [1][2][3][4].
As an important connection method for metal alloy, welding process has been extensively used in the construction of nuclear power construction. In the early years, nickel-base alloy 600, welding wire ERNiCr-3 and welding electrode ENiCrFe-3 had been used to manufacture nuclear reactors. However, it was found that those materials were susceptible to primary water stress corrosion cracking (PWSCC) due to their insufficient chromium content (15 wt% Cr). Therefore, it has been replaced by high chromium nickel-based alloy 690, welding wire ERNiCrFe-7, ERNiCrFe-7A and welding electrode ENiCrFe-7 (30 wt% Cr) nowadays [5,6]. In addition, the filler metal ERNiCrFe-7, ERNiCrFe-7A, and ENiCrFe-7 are typical welding materials for joining nickel-based alloy and dissimilar alloy [7,8].
Several studies on weld metal or deposited metal of ENiCrFe-7 cover electrode have been carried out. Qin et al [9] studied the carburization phenomenon of deposited metal in three different slag systems. They found the acidic (CaOTiO 2 SiO 2 type) and neutral (CaOCaF 2 TiO 2 type) slag systems exhibited relatively lower carburization than the basic (CaOCaF 2 type) slag, and the carbon came from the carbonates in the flux coating. Moreover, they investigated the microstructure and ductility-dip cracking (DDC) susceptibility of deposited metal [10]. They found discontinuous Cr-rich M 23 C 6 carbides on the grain boundaries tended to coarsen during reheating, but began to dissolve above approximately 1273 K (1000°C). The DDC susceptibility increased sharply as the carbides coarsened in the temperature range of 973 K to 1223 K (700°C to 950°C). Wang et al [11] investigated the metallurgical behavior and the slag detachability of the CaO-CaF 2 -SiO 2 type ENiCrFe-7 covered Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. electrode. The results showed that the separability of slag could be improved with the decrease of SiO 2 in flux coating, the SiO 2 in the flux coating was 10.9 pct, about 28.3 pct CaF 2 resulted in the best slag detachability. Comparing the slag compositions, the slag detachability was best at CaO:CaF 2 :SiO 2 = 1.7:1.8:1, and worst at CaO:CaF 2 :SiO 2 = 1.3:0.9:1. Jeng et al [12] investigated the effects of Mn and Nb on the microstructure and mechanical properties of ENiCrFe-7 weld metal. The base metal consisted of Alloy 690 and SUS 304 L. They found that an increase in Mn could induce finer grains and smaller Nb-rich phases in the weld, the elongation of weld metal increased and the hardness decreased at room temperature. High contents of Mn and Nb could improve the notch strength and ductility at 300°C, while they had little effects on the tensile strength at room temperature. In addition, Guo et al [13][14][15][16] studied the effects of post weld heat treatment (PWHT) on microstructure, mechanical properties, residual stress, and corrosion resistance of ENiCrFe-7 deposited metal at 615°C for 16 h and 48 h.
PWHT is usually required for weldment to reduce residual stress after welding. The PWHT parameters commonly include temperature with a holding time that depends on the component's size, ranging from a few hours for small components up to tens of hours for large components. For example, the typical heat treatment parameter for the steam generator and reactor pressure with weld metal of ERNiCrFe-7 wire is usually 610°C-620°C with holding time 24 h [17]. It is known that PWHT not only reduces the welding residual stress, but also affects the microstructure and properties. The effects of PWHT on microstructure and mechanical properties of deposited metal using ENiCrFe-7 covered electrode at 615°C for 16 h and 48 h had been reported, it was found that the tensile strength decreased after holding time for 16 h, while increased after holding time for 48 h [13,14]. However, the interval between the two holding time was long, and researches on the changes of impact toughness had not been reported yet. For structural manufacturers, on the one hand, they hope to obtain welded metal with good comprehensive performance after PWHT, such as good tensile strength, plasticity, and toughness, on the other hand, they hope to save economic costs and reduce PWHT time in manufacturing. Therefore, it is significant to study the microstructure, tensile properties and toughness changes of deposited metal at an intermediate PWHT time, such as 615°C with 24 h. In this work, the effects of PWHT at 615°C for 24 h on microstructure, hardness, tensile properties and impact toughness of deposited metal of ENiCrFe-7 covered electrode were studied.

Materials and methods
The deposited metal of ENiCrFe-7 covered electrode was fabricated in the form of butt weld using a shielding metal arc welding (SMAW) process. The nominal chemical composition of deposited metal is given in table 1. The dimension of ENiCrFe-7 covered electrode was 3.2 mm, the base metal was 20 mm thick carbon steel (SAE 1020) plate and the backing strip was 10 mm thick carbon steel (SAE 1020). Both base metals were machined to the dimensions of 350 × 100 × 20 mm. All plates were prepared as single V grooved joints with a 10°bevel angle. In order to keep the deposited metal undiluted, the prepared edge and backing strip were buttered with ENiCrFe-7 covered electrode up to 6 mm thickness. Figure 1 is the schematic of the weldment. A total of 19 weld passes were carried. The welding parameters are summarized in table 2. The welding process was performed using a direct current of 90 ∼ 100 A and a voltage of 24 ∼ 28 V, a travel speed of 130 ∼ 160 mm min −1 . The heat input was 1.00-1.05 kJ mm −1 . The interpass temperature was below 100°C for all weld beads. Two groups with four weldments were welded under the same welding parameter. PWHT then was carried out to relieve the internal stress for one group with two weldments. The PWHT parameters were 615°C with holding time 24 h, the maximum heating rate was limited to 55°C/h in the temperature interval 350°C-615°C.
The microstructure, room temperature and elevated temperature tensile properties, impact toughness and hardness were tested on both as-welded and PWHT deposited metals. The location of the samples on the weldments is shown in figure 2(a). The metallographic samples were prepared by a mechanical grinding and polishing process, etched using a solution of 10 g CuSO 4, 50 ml HCl and 50 ml H 2 O for 10 s. The macrostructure was observed by a VHX-1000 ultra-depth of field three-dimensional microscope, and the microstructure was observed using a Zeiss Axio Imager A2m optical microscope (OM) and a Zeiss Merlin Compact scanning electron microscopy (SEM) equipped with an energy dispersive spectrometer (EDS). In addition, the number of precipitates was counted using a COXEM EM-30AX SEM equipped with a backscattered electron detector (BSE). The structure and morphology of the precipitates were analyzed by a JEM-2100F transmission electron    microscope (TEM). The TEM specimens were mechanically ground to a thickness of about 50 μm, then thin foils were prepared by double jet thinner using an electrolyte of 7% HClO 4 , 4% CH 3 COOH and 89% C 2 H 5 OH at 38 V and −30°C. The room temperature and elevated temperature tensile specimens of as-welded condition were two pieces respectively, the same to PWHT condition. The tensile specimens tested by a UTM5205X tensile testing machine with a diameter of 10 mm and a gauge length of 50 mm ( figure 2(b)). The tensile rate of room temperature tensile test was 1 mm min −1 . Elevated temperature tensile test was set at 350°C with a heating rate of 20°C min −1 , keeping for 10 min, and the tensile rate was also 1 mm min −1 . The Charpy V-notch impact specimens of as-welded and PWHT conditions were six pieces respectively, with a size of 55 mm × 10 mm ×10 mm (figure 2(c)), tested by a JB-300B pendulum impact testing machine at room temperature. The fracture surfaces of tensile and impact specimens were observed by a JSM-6510 SEM. The hardness of metal was measured by an HVS-50Z Vickers hardness tester, the load was 10 kgf and dwell time was 10 s.

Microstructure
The macrostructures of deposited metals in as-welded and PWHT conditions by OM are shown in figures 3(a) and (b). The buffer layers and weld passes were clearly visible. Both the macrostructures showed solidified dendrite structure feature. The solidified grains in the weld pool tended to grow along the maximum temperature gradient direction during the solidification process, which grew in a direction perpendicular to each interface of weld pass, and the growth orientations were different in each pass. The macrostructure of aswelded and PWHT deposited metals had no obvious difference. The center-zone microstructure of as-welded and PWHT deposited metal is shown in figures 3(c) and (d). Precipitates with different sizes were found in grains. The grain boundary morphology was straight with slight bending. The grains were typically columnar with a length stretching several millimeters in one direction and relatively smaller in width direction. It was found that the grain boundaries of PWHT deposited metal were more obvious than those of as-welded one under the same agent and etched time parameter, indicating the grain boundaries coarsened after PWHT. In the solidification process of pure metal, the solid-liquid interface was usually flat unless there was severe undercooling. However, during the solidification process of alloys, the solid-liquid interface and the solidification mode that occurred could be planar, cellular, columnar or dendritic, depending on the solidification condition and material composition. The crystal morphology mainly depended on the effect of crystallization rate R and temperature gradient G in the liquid phase in the alloy. In the center of weld metal, the temperature gradient G was relatively small, while the crystallization rate R was relatively high, resulting in a columnar crystal growth. Figure 4 shows the grain width in deposited metals of as-welded and PWHT conditions. Due to columnar structure and large grain length, the grain width was measured to compare the changes in grain size before and after heat treatment (figures 4(a) and (b)). 15 regions were selected and more than 40 grains widths were calculated in each deposited metal, and the measurement position was the parallel segment width of the length direction. The average grain widths are shown in figure 4(c). The average grain width in as-welded deposited metal was 111.1 μm, while 115.4 μm in PWHT deposited metal, which indicated the grains grew but not obviously after PWHT. Figures 5(a) and (b) are the BSE image of deposited metals in as-welded and PWHT conditions. Two kinds of precipitates were found in the matrix. One was black with round or roughly round shape and the other was white with irregular shape. Subsequent analysis showed that the black precipitates were Al-Ti oxides, and the white precipitates were NbC. The density of these precipitates was counted by Image-Pro Plus software in twenty regions and the results were shown in figure 5(c). It was found no significant difference in the precipitates number per unit area in the two deposited metals, which indicated that there was no precipitation of NbC during PWHT.
The SEM microstructure images of deposited metals in as-welded and PWHT conditions are shown in figures 6(a) and (b). Precipitates with two main types of shapes dispersed in the grains, i.e. spherical and irregular shape. The sizes of spherical precipitates were less than 1 um (indicated by arrows 1, 4). EDS results listed in  3 showed these spherical particles were rich in O, Al, Ti, which were inferred as Al-Ti oxides. Figure 7 is the bright-field TEM image of the spherical particle, the result of EDS attached on TEM further confirmed it was Al-Ti oxide. There were two main sources of oxygen in weld metal: one came from air, as oxygen was the second main component of air, resulting in absorbing O easily from air during welding; The other came from welding materials, which was decomposed by H 2 O absorbed by welding materials. The cover electrode was composed of coating and core. The coating was composed of minerals such as CaO, CaF 2 and SiO 2 , etc while those minerals easily absorbed H 2 O. It was known that Al and Ti elements were excellent deoxidizers, so a certain amount of Al and Ti elements were usually added to the electrode. The deoxidizers Al and Ti reacted with oxygen forming Al-Ti oxides and entered to slag, a few oxides might remain in the melted pool and enrich in the interdendritic region as the dendritic arms grew during solidification [18,19]. The results indicated that the circular black phases observed in the BSE images were Al-Ti oxides ( figure 5).
The irregular particles were bar-like or block-like shape with different sizes (bar-like indicated by arrows 2, 5, block-like indicated by arrows 3, 6). The EDS results showed they contained high Nb and C content, which were suggested as NbC carbides. Figures 8 and 9 are the bright-field TEM images of bar-like and block-like particle, respectively. the corresponding selected area electron diffraction (SAED) patterns further confirmed they were NbC with an FCC structure. Nb is prone to segregate in the dendrites during the solidification crystallization process [20]. Nb would form NbC carbides when enriched to a certain extent in the remained liquid at the later stage of solidification. The shape of NbC may be related to the degree of Nb enrichment. As a amount of Nb segregation in local area, it is prone to grow and transform from bar shape to block. Figures 6(c) and (d) are the magnified images of grain boundaries. The intergranular precipitates were clearly visible, which were blocky and chain-like distributing along the grain boundaries. It was obvious that the intergranular phases coarsened after PWHT, which transformed from small blocky shape to large blocky shape. EDS results (table 3) showed the intergranular phases (indicated by arrow7, 8) contained high Cr content, which was identified as Cr-rich M 23 C 6 carbides. M 23 C 6 has a cubic coherent relationship with one side of matrix, and its lattice parameters are about three times that of matrix [21][22][23][24]. The large misfit between M 23 C 6 and matrix interface easily causes stress concentration and crack nucleation [25].
It is known that the matrix microstructure of nickel-base alloy is γ austenite. The solidification reaction sequence of deposited metal of ENiCrFe-7 electrode can be described as follows [26]: Firstly, L → γ solidification transformation, dendrites form and Nb is enriched in interdendritic area. Secondly, a eutectic-type reaction L → γ + NbC occurs as Nb is enriched to a certain extent, which consumes the majority of the available carbon. Thirdly, M 23 C 6 precipitates from matrix on the grain boundary at a lower temperature than NbC. The precipitation temperature of M 23 C 6 is related to the content of Nb [27,28]. M 23 C 6 would grow and coarsen during PWHT, while a relatively stable phase NbC would not change owing to M 23 C 6 forming faster than NbC [29].

Hardness
The center Vickers hardness results of deposited metals in as-welded and PWHT conditions are shown in figure 10. The distance of the two test points was 2 mm. It was obvious that the hardness of deposited metal in asweld state was higher than PWHT one. The average hardness value of deposited metal in as-welded condition was 216-228 HV, while the PWHT one was 208-213 HV. The average hardness of deposited metal in as-welded condition was about 10 HV higher than that of PWHT condition.  Table 3. Chemical composition of phases detected by EDS in figure 6 (wt%).

Positions
Suggested  Figure 11 shows the room temperature and 350°C elevated temperature tensile properties of the deposited metals in as-welded and PWHT conditions. The tensile strength, yield strength and elongation of as-welded specimens at room temperature were 686 MPa, 452 MPa and 38%, respectively. After PWHT, the tensile strength and yield strength decreased to 657 MPa and 427 MPa, while the elongation increased to 42%. When testing at 350°C, the elevated temperature tensile strength, yield strength and elongation of as-welded specimens were 565 MPa, 414 MPa and 36%, respectively. The tensile strength and yield strength decreased to 539 MPa and 363 MPa, while the elongation increased to 39% after PWHT. The PWHT deposited metals were a slight reduction in tensile strength and a small increment in elongation compared to the as-welded deposited metal. The tensile strength reduced by about 30 MPa, and elongation increased by about 3 ∼ 4% after PWHT. Guo et al [13] investigated the tensile strengths of ENiCrFe-7 deposited metals in as-welded condition, PWHT at     [30,31]. C and Cr elements play an important role in the strength of Ni-Cr-Fe alloys, more M 23 C 6 carbides precipitated from matrix on the grain boundary during PWHT weakened the solid solution strengthening effect [32][33][34]. The stress concentration formed by intersections of grain boundaries and slip bands leaded to a low bonding strength between the coarsened M 23 C 6 carbide and the matrix in the tensile test [23,32,35]. The release of residual stress leaded to a decrease in tensile strength during PWHT [13]. In addition, the small grain growth also played a role in the reduction of tensile strength after PWHT. Figure 12 shows the SEM images of the fracture surfaces of tensile specimens at room temperature. Both the as-welded and PWHT tensile specimens exhibited ductile fracture characteristic. A large number of dimples and  microvoids appeared on the fracture surface, while the dimples in PWHT sample were larger than as-welded one. Some second phase particles existed at the bottom of dimples. The deformation of second phase particles and matrix was not simultaneous during the tensile deformation, resulted in dimples formation around these particles. Figure 13 shows the SEM images of the fracture surfaces of tensile specimens at 350°C. Similar to tensile fracture characteristics at room temperature, both the fracture surfaces of as-welded and PWHT tensile specimens presented ductile characteristic. Numerous dimples and tear ridges were distributed on the fracture surfaces. It could be seen that the dimples in PWHT condition were larger and deeper than that in as-welded condition, which indicated larger plastic deformation happened during the tensile fracture of PWHT specimen. The larger plastic deformation corresponded to the higher elongation in figure 11. Figure 14 shows the Charpy impact energy of the deposited metals at room temperature. The impact energy values for deposited metals in as-welded and PWHT conditions were 104 J and 103 J, respectively. Both were close and showed good toughness. Figure 15 shows the SEM fracture surface morphologies of the impact specimens. Both the as-welded and PWHT specimens fracture surfaces were uneven, containing secondary cracks (figures 15(a) and (b)). Figure 15(c) shows the fracture feature of as-welded specimen was a mixed fracture mode with lots of shallow dimples and a few cleavage steps. The dimples and cleavage steps were typical characteristics of ductile fracture and brittle fracture, respectively. Ductile fracture was caused by the formation, growth, and interconnection of hollow nuclei under impact force, which was a high-energy absorption process of ductile fracture. Inclusions or second phase particles such as Al-Ti oxides and NbC commonly existed in the center of the dimples. Cleavage fracture was a transgranular brittle fracture caused by crack propagation along a specific crystallography plane, impact load and stress concentration often promoted the occurrence of cleavage  fracture. The results indicated the ductile fracture was dominant in the impact specimens. The fracture feature of PWHT specimen showed similar characteristic to as-welded specimen ( figure 15(d)).

Impact toughness
Although M 23 C 6 carbides along the grain boundaries coarsened after PWHT, no obvious intergranular fracture characteristics were found on the impact fracture surface, the impact toughness of deposited metal did not decrease or affect by M 23 C 6 precipitation. The reason was believed to due to a decrease in lattice distortion and dislocation density after residual stress release during PWHT, resulting in an improvement in impact toughness [36,37]. The precipitation of intergranular M 23 C 6 carbides after PWHT weakened the toughness,