Effects of solution treatment on grain coarsening and hardness of laser welds in UNS N10003 alloy contained different carbon content

1. Introduction

Molten salt reactor (MSR), which possesses inherent safety, simplified fuel cycle and high power generation efficiency [1,2], is one of the generation IV nuclear reactors. The MSR operates at high temperature environment with 600-700 °C and uses molten salt as coolant which has severe corrosion. Thus, Hastelloy N alloy, designated as UNS N10003 alloy by American Society of Mechanical Engineers (ASME), was designed by Oak Ridge National Laboratory (ORNL) in 1950s and 1960s for Molten Salt Reactor Experiment (MSRE) because of its excellent mechanical performance and corrosion resistance [[3], [4], [5]]. In recent years, the MSR caught the attention of researchers again. Chinese have been developing the thorium molten salt reactor (TMSR) from 2011 [6]. GH3535 alloy (UNS N10003), made in china, was identified as the primary structure material for TMSR. However, the challenges of UNS N10003 alloy application still exist. The development of Hastelloy N alloy has been suspended several decades and the range of critical elements composition has not been optimized. In addition, the range of composition in ASME Boiler and Pressure Vessel Code is so broad that the properties of this alloy exist instability [7]. Therefore, the composition optimizing of alloy and evaluation of microstructure and mechanical properties of welds under different element composition condition are significant.

The composition of some elements is being modified to achieve much more excellent performance. The composition of element Si was optimized to about 0.32 wt% which could attain microstructural stability at high temperature [8,9]. The composition of element C affects the strength of alloy and weldability of alloy. Especially, the low melting phases contained carbon were always precipitated at the termination of solidification, which may induce hot cracking in nickel base alloy [[10], [11], [12]]. Thus, it is necessary to study the effects of different carbon content on microstructure of welds. The low carbon N10003 alloy (0.018 wt% C) (LC alloy) and high carbon N10003 alloy (0.054 wt% C) (HC alloy) have been prepared in previous work [13,14]. The laser beam welding was conducted to weld the two type alloys, and the microstructure and mechanical properties of laser welds were evaluated. The results showed that there are no secondary phases precipitated in fusion zone of LC laser welds. However, a great deal of eutectic carbides were precipitated at SGB and SSGB in fusion zone of HC laser welds. There are no any hot cracks in fusion zone of LC and HC laser welds, and the tensile strength of LC and HC laser joints could be up to 95% strength of base metal. In a word, the LC and HC laser welds exhibited excellent performance under as-welded condition.

In order to relieve residual stress for structure stabilization, solution treatment was often conducted. The element diffusion, grain coarsening and precipitates evolution are affected significantly by post weld heat treatment (PWHT) in nickel base alloy [[15], [16], [17], [18]]. However, there is no literature to study and compare the effects of solution treatment on microstructure of LC and HC laser welds in UNS N10003 alloy. In this work, the laser welds were treated at 1177 °C for 20 min, and then the microstructure and hardness evolution before and after solution treatment in LC and HC laser welds were characterized by a series of methods. The behavior of grain coarsening and precipitates evolution were discussed.

2. Experimental procedures

The UNS N10003 alloy (LC alloy and HC alloy) with 4 mm thick, which was provided in the form of solution treatment state, was used in this work. The plate of UNS N10003 alloy was sectioned into small plates with the size of 300 mm long and 100 mm wide.

The laser welding experiments were conducted using an IPG YLS.10000 fiber laser whose laser beam had a wavelength of 1.06 μm. Laser beam was transmitted by an optical fiber with a 200 μm core diameter to a laser welding head mounted on a 6-axis KUKA robot. The laser beam was focused onto the plate surface by a mirror with a 200 mm focal length and was focused as a spot of 0.26 μm in diameter.

The plates were grounded by a sand blasting to remove the surface oxide, then cleaned the surface with acetone prior to welding. Two plates were fixed in a butt configuration and welded without any filler metal. The welding direction was parallel to the rolling direction of the plate. The process parameters of laser beam welding for LC alloy and HC alloy are listed in Table 1. The welding process had been optimized according to nondestructive testing results.

After the laser welding process, half of LC and HC weldments were cut and carried out solution heat treatment at 1093 °C and 1177 °C for 5, 20, 40 min followed by air cooling.

After the solution treatment process, the welded joints were sectioned into specimens using an electrospark wire-electrode cutting. Cross-sections of welded joints for microstructure examination were mounted in epoxy resin, grinded with SiC papers from 400 grit to 2000 grit, and then polished with a 0.05 μm alumina paste for 2 min. The specimens for optical microscopy (OM) and scanning electron microscopy (SEM) examinations were etched with a mixture solution of (3 g CuSO₄ + 80 ml HCl) for 30 s. The specimen for electron backscatter diffraction (EBSD) examination was polished by vibratory finishing machine for 1 h without solution etched. The specimens for transmission electron microscopy (TEM) examination were punched into disks with a diameter of 3 mm and 0.1 mm in thickness.

The microstructure characterization was carried out by Zeiss M2m OM and Zeiss Merlin Compact SEM equipped with an Oxford EBSD system. The precipitates identification was analyzed by Tecnai G2 F20 S-TWIN TEM equipped with energy dispersive X-ray (EDX) analysis system.

Vickers hardness of welded joints was measured by Zwick/Roell ZHVμ-S machine under a load of 200 g sustaining 15 s. The specimen was measured with indentations at intervals of 0.1 mm, and three measurements were carried out for each indent.

3. Results and discussion

3.1. Microstructure of two kinds of base metals

The chemical composition of LC alloy and HC alloy in this work is listed in Table 2. It is shown that the carbon of LC alloy is 0.018 wt%, while that of HC alloy is 0.054 wt%. Fig. 1(a) and (b) shows the optical micrographs of LC alloy and HC alloy in as-received condition. Although the carbon of two kinds of alloys is different, the microstructure of alloys all consists of austenite γ phase and M₆C carbides distributed in the grain boundaries and also inside the grains [[19] [20],]. It is shown that the number of carbides in LC alloy is less than that of HC alloy.

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Fig. 1.   Microstructure of base metal: (a) low carbon alloy; (b) high carbon alloy.

3.2. Microstructure of fusion zone evolution

Fig. 2 (a) shows the microstructure of fusion zone in LC laser welds under as-welded condition. It can be seen that the microstructure of fusion zone consisted of dendrites, cells and equiaxed subgrains which were affected by the degree of constitutional supercoiling during solidification. Fig.2 (b) shows the microstructure of fusion zone in LC laser welds under solution treatment at 1177 °C for 20 min. It can be seen that the microstructure of fusion zone is coarse equiaxed grains contained twin crystal. Comparing with the microstructure of as-welded fusion zone, the subgrains consisted of dendrites and cell have been disappeared under solution treatment condition. Fig.2 (c) shows the microstructure of fusion zone in HC laser welds under as-welded condition. It can be seen that the microstructure of fusion zone consisted of dendrites, cells and equiaxed subgrains are same as that of LC laser welds. Fig.2 (d) shows the microstructure of fusion zone in HC laser welds under solution treatment at 1177 °C for 20 min. It can be seen that the solidification grain boundaries (SGBs) of fusion zone are clearly observed, and quantifies of black pots are found in the interior grain.

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Fig. 2.   Optical photographs of fusion zone: (a) as-welded fusion zone of LC alloy; (b) PWHT fusion zone of LC alloy; (c) as-welded fusion zone of HC alloy; (d) PWHT fusion zone of HC alloy.

In order to compare the grain evolution clearly, the EBSD analysis was carried out. Fig. 3 (a) shows the grain characterization of as-welded fusion zone in LC laser welds, it can be seen that the zigzag SGB of columnar grain was obviously observed. The grain boundaries misorientation are more than 15°, and ∑3 boundary was not found in the fusion zone. According to the statistics by Tango analysis software, the average grain diameter of fusion zone is 47.5 μm. Fig.3 (b) shows the grain characterization of fusion zone in LC laser welds under solution treatment, it can be seen that there are a great deal of ∑3 boundaries marked with red line in the SGB. The boundaries are flat rather than zigzag morphology as that of as-welded fusion zone. What is more, the grain coarsening of fusion zone is obvious when it is suffered from solution treatment. The average grain diameter of fusion zone suffered from solution treatment is 97.7 μm. Fig.3 (c) shows the grain characterization of as-welded fusion zone in HC laser welds, it can be seen that the grain boundary characterization is the same as that of LC fusion zone shown in Fig.3 (a). The average grain diameter of as-welded fusion zone is 48.1 μm. The microstructure of LC and HC fusion zone all exhibit inhomogeneity. After solution heat treatment, the columnar grains had still existed shown in Fig.3 (d). The zigzag grain boundaries of fusion zone exhibit a certain extent smoothly. The dimension of grains have a bit coarse than of as-welded grains. According to the EBSD analysis, the average grain diameter of fusion zone suffered from solution treatment is 57.1 μm. The ∑3 boundaries are not found other than “big misorientation” boundaries.

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Fig. 3.   Grain morphology of fusion zone by EBSD analysis: (a) as-welded fusion zone of LC alloy; (b) PWHT fusion zone of LC alloy; (c) as-welded fusion zone of HC alloy; (d) PWHT fusion zone of HC alloy.

In order to further investigate the microstructure, SEM analysis was carried out. Fig. 4 (a) shows the microstructure of fusion zone in LC laser welds under as-welded condition. It can be seen that there are no any precipitates at interdendritic region or dendritic cores. Fig.4 (b) shows the microstructure of fusion zone in LC laser welds under solution treatment. It can be seen that the grain boundary is flat and no precipitates are found in the grain boundary. In addition, a small quantity of precipitates can be observed in the interior grain. Fig.4 (c) shows the microstructure of fusion zone in HC laser welds under as-welded condition. It can be seen that the eutectic precipitates were precipitated at interdendritic region. Fig.4 (d) shows the microstructure of fusion zone in HC laser welds under solution treatment. The eutectic precipitates have been transformed to spherical precipitates.

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Fig. 4.   Precipitates in the fusion zone: (a) as-welded fusion zone of LC alloy; (b) PWHT fusion zone of LC alloy; (c) as-welded fusion zone of HC alloy; (d) PWHT fusion zone of HC alloy.

In order to identify the crystal structure, TEM analysis was carried out. Fig. 5 shows the TEM analysis result of the precipitate in LC fusion zone after solution treatment. The results show that the precipitate is rich in Mo, Si and C. The selected area electron diffraction (SAED) pattern of the precipitate with [-110] zone axis in Fig. 5 (b) exhibits a Face centered cubic (FCC) crystal structure with a lattice parameter of about 1.103 nm. It can be identified as M₆C carbide. In our previous research work ^[14], the crystal structure of eutectic precipitates of HC fusion zone shown in Fig.4 (c) had been identified as eutectic M₆C carbides which are rich in Mo and Si. After solution treatment, the morphology of carbides have been changed obviously, while the crystal structure of carbides have not been changed. The selected area electron diffraction (SAED) pattern of the precipitate with [1-1-2] zone axis in Fig. 6 (b) exhibits a Face centered cubic (FCC) crystal structure with a lattice parameter of about 1.097 nm. The spherical carbide can be identified as M₆C carbide and is rich in Mo and Si shown in Fig.6 (c).

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Fig. 5.   TEM analysis of carbide in LC alloy after solution treatment: (a) image of carbide; (b) SAED pattern of the carbide; (c) spectrum of EDX point analysis at the carbide marked as 1 in.(a).

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Fig. 6.   TEM analysis of carbide in HC alloy after solution treatment: (a) image of carbide; (b) SAED pattern of the carbide; (c) spectrum of EDX point analysis at the carbide marked as 1 in.(a).

3.3. Microhardness of welded joints evolution

The microhardness of welded joints under different condition are tested and compared because the microhardness is an important index to evaluate the material properties shown in Fig. 7. The microhardness distribution of welded joints begin with weld center and end with BM because of symmetrical characteristic of the welded joints. The results show that the average microhardness of FZ and BM in LC as-welded joint are about 223 HV_0.2 and 220 HV_0.2, respectively. After solution treatment at 1177 °C for 20 min, the average microhardness of FZ and BM in LC alloy are 190 HV_0.2 and 187 HV_0.2, respectively. The average microhardness of FZ and BM in HC as-welded joint are about 260 HV_0.2 and 223 HV_0.2, respectively. When the welded joint was suffered from solution treatment, the average microhardness of FZ and BM are about 225 HV_0.2 and 190 HV_0.2, respectively.

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Fig. 7.   Microhardness distribution of welded joints under different treatment conditions.

Comparing the microhardness of BM between LC alloy and HC alloy before solution treatment, although the carbon content between LC alloy and HC alloy is much different, it can be seen that the microhardness of BM in LC alloy is similar as that of BM in HC alloy. Comparing the microhardness of as-welded FZ between LC alloy and HC alloy, it is shown that the microhardness of FZ in LC alloy is lower than that of FZ in HC alloy. The reason is that there are no secondary phases precipitated in LC FZ, while quantifies of carbides precipitated in HC FZ which can strengthen the mechanical properties shown in Fig. 4. After post welded heat treatment, that is solution treatment, the microhardness of BM between LC alloy and HC alloy all decrease to equivalent value. The microhardness of FZ between LC alloy and HC alloy also decrease, while the microhardness of LC FZ is much lower than that of HC FZ. The reason is that the grains of FZ in LC alloy are coarsen seriously, while the grains of FZ in HC alloy are slight coarsen shown in Fig. 3. According to the Hall-Petch relationship [21], the reduction of hardness is expected as the grain structure of material coarsening.

3.4. Discussion of carbides evolution

Base on solidification theory, the segregation happen on interdenritic region in FZ. In our previous work [13], element Si, C and Mo with equilibrium distribution coefficient k<1 are rich in interdenritic region of FZ in LC alloy. However, there are no any carbides precipitated in FZ because there are no enough carbon segregation in LC alloy during the rapid cooling under LBW process shown in Fig. 4 (a). After solution treatment, the element diffusion exhibit easily because the atoms are activated under high temperature condition. The carbide forming elements tend to diffuse to interdenritic region where there are quantifies of atom defects that can catch the atoms. Therefore, there are some M₆C carbides precipitated in FZ shown in Fig. 4 (b). In this way, the free energy of system can be decreased and the microstructure exhibit stability.

Comparing with the microstructure of as-welded FZ in LC alloy, a great deal of eutectic M₆C carbides emerge at interdenritic region in HC alloy because the HC alloy has enough carbon. The Si, Mo and C, which exhibit k<1, segregated in interdendritic liquid region during the solidification procedure. The M₆C type eutectic carbides precipitate in the interdendritic region by eutectic type reaction: L→(γ+M₆C) at the terminal stages of solidification. After solution treatment, the eutectic carbides have transformed into spherical carbides shown in Fig.4 (c) and (d). The driving force for spheroidization of the carbides is the reduction of the Gibbs free energy under high temperature condition [22,23]. The atoms have been activated and driving force of microstructure transformation turn into greater under high temperature condition. The system always develop to stability from high Gibbs free energy to low free energy. The superficial area of eutectic carbide is larger than that of spherical carbide, that is the surface energy of eutectic carbide is higher than that of spherical carbide. Thus, the eutectic carbide can been transformed spontaneously into spherical carbide under solution treatment condition.

3.5. Discussion of grain coarsening

The change of free energy exhibits not only the evolution of carbide but also evolution of grain. The classical expression describing grain growth is given in Eq. (1).

$\bar{D}$(t)-$\bar{D}_0$=ktⁿ (1)

where $\bar{D}$(t) is the average gain size at time t, $\bar{D}_0$ is the initial grain size, n and k are exponent and rate constants respectively [24]. In order to treat grain growth as a thermally-activated process, the rate constant k can be defined by an Arrhenius equation in Eq. (2), where Q is the apparent activation energy for grain growth, k₂ is a constant depending on the physical kinetics, R is the ideal gas constant and T is the absolute temperature [25].

The grain growth kinetics of fusion zone in UNS N10003 alloy can be gained through determination of variables n, k₂ and Q. n can be found through the linear fitting of Eq.(3) to isothermal grain size plots, and k₂ and Q can be subsequently calculated according to Eq.(2).

$\bar{D}$(t)-$\bar{D}_0$=k₂exp-($\frac{-Q}{RT}$)tⁿ (2)

Ln($\bar{D}$(t)-$\bar{D_0}$=nlnt-lnk (3)

In order to calculate activation energy (Q) for grain growth, k and n should be attained first. Consequently, the different solution treatment processes had been conducted using 5, 20 and 40 min at 1177 °C and 1093 °C, respectively. The grain sizes of FZ were analyzed by EBSD and calculated by Tango analysis software. The grain sizes for different PWHT conditions at 1177 °C and 1093 °C were shown in Fig. 8 and Fig. 9. The grain sizes of LC welds under as-welded, 1177 °C for 5 min, 1177 °C for 20 min and 1177 °C for 40 min conditions were 47.5, 81.7, 97.7, 128.1 μm respectively. While the grain sizes of HC welds under as-welded, 1177 °C for 5 min, 1177 °C for 10 min, 1177 °C for 20 min and 1177 °C for 40 min conditions were 48.1, 55.4, 56.8, 57.1, 65.6 μm respectively. The grain sizes of LC welds under as-welded, 1093 °C for 5 min, 1093 °C for 20 min and 1093 °C for 40 min conditions were 47.5, 60.6, 69.8, 75.1 μm respectively. The grain sizes of HC welds under as-welded, 1093 °C for 5 min, 1093 °C for 20 min and 1093 °C for 40 min conditions were 48.1, 53.0, 59.7, 65.4 μm respectively. It can be seen that the grain coarsening increases with the increase of holding time. The isothermal grain growth trends of FZ in LC and HC alloy is shown in Fig.10. The k and n of LC and HC welds can be given in Table 3. Further, the activation energy Q and rate constant k₂ can be calculated and listed in Table 4.

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Fig. 8.   Grain size of FZ for PWHT at 1177 °C.

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Fig. 9.   Grain size of FZ for PWHT at 1093 °C.

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Fig. 10.   Isothermal grain growth trends for welds.

According to the EBSD analysis, the grain coarsening value (i.e. $\bar{D}$(t)-$\bar{D_0}$) of FZ in LC alloy is 50.2 μm, while that of FZ in HC alloy is only 9.0 μm for 20 min at 1177 °C. It is suggested that Q of LC laser welds much lower than that of HC laser welds as shown in Table 4. The activation energy of FZ in LC alloy is 144.3 kJ mol^-1, while that of FZ in HC alloy is 309.5 kJ mol^-1. The energy of thermal activation in LC alloy is same as that of HC alloy under the same solution treatment condition. However, their grain coarsening exhibit different behavior due to effect of M₆C carbides in FZ. The system develops from high free energy to low free energy which exhibits the grain coarsening. The grain coarsening can lead to the surface energy decrease. Thus, the behavior of grain coarsening is spontaneous. The zigzag grain boundaries of as-welded FZ in LC alloy were migrated and transformed into straight grain boundaries under thermal activation shown in Fig.3 (a) and (b). This boundary migration was not impeded because of no carbides precipitated. Although the zigzag grain boundaries of as-welded FZ in HC alloy had been transformed into straight grain boundaries to a certain extent after solution treatment, the grain coarsening did not exhibit obviously shown in Fig.3 (c) and (d). The reason is that the boundary migration is impeded by the M₆C carbides distributed in the boundary. The migration of grain boundary needs more great energy, which is that Q is higher. The carbide evolution will come to be easy when the grain coarsening is difficult under thermal activation shown in Fig.4 (c) and (d). In this way, the free energy of system decreases and the microstructure comes to be stable.

It can be seen that the weld of HC alloy exhibits much more excellent thermal stabilities than the weld of LC alloy. The stable microstructure is good for MSR.

4. Conclusions

(1) There are no carbides precipitated in as-welded FZ of LC alloy while quantifies of eutectic M₆C carbides precipitate in as-welded FZ of HC alloy due to difference of carbon content.

(2) A few of spherical M₆C carbides precipitate in FZ of LC alloy and a great deal of eutectic M₆C carbides transform into spherical M₆C carbides after solution treatment due to thermal activation of atoms and diffusion of elements.

(3) The columnar grains of FZ in LC alloy transform into equiaxed grains, while the columnar grains of FZ in HC alloy exhibit little changes after solution treatment due to the grain coarsening is affected by M₆C carbides impeding. The activation energy of grain coarsening for FZ in LC and HC alloy is 144.3 kJ mol^-1 and 309.5 kJ mol^-1 respectively.

(4) The microhardness of FZ in LC alloy is lower than that of FZ in HC alloy because of grain coarsening in LC alloy.

(5) The weld of HC alloy exhibits excellent thermal stabilities and is good for molten salt reactor.

Acknowledgments

This work was supported by the National Key Research and development program of China (2016YFB0700404) and the National Natural Science (Grant Nos. 51371188, 51671122, 51671154, 51501216 and 51601213) and the Strategic Priority Research Program of the Chinese Academy of Sciences (Grant No. XDA02004210).

The authors have declared that no competing interests exist.