P92 dissimilar weld joints

Journal Pre-proof Effect of post-weld heat treatment on the microstructure and hardness of P92 steel in IN740H/P92 dissimilar weld joints

Wi-Geol Seo, Jin-Yoo Suh, Jae-Hyeok Shim, Hansang Lee, Keunbong Yoo, Shi-Hoon Choi PII:

S1044-5803(19)33075-X

DOI:

https://doi.org/10.1016/j.matchar.2019.110083

Reference:

MTL 110083

To appear in:

Materials Characterization

Received date:

10 November 2019

Revised date:

17 December 2019

Accepted date:

17 December 2019

Please cite this article as: W.-G. Seo, J.-Y. Suh, J.-H. Shim, et al., Effect of post-weld heat treatment on the microstructure and hardness of P92 steel in IN740H/P92 dissimilar weld joints, Materials Characterization (2018), https://doi.org/10.1016/j.matchar.2019.110083

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© 2018 Published by Elsevier.

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Effect of Post-Weld Heat Treatment on the Microstructure and Hardness of P92 Steel in IN740H/P92 Dissimilar Weld Joints Wi-Geol Seo1, Jin-Yoo Suh2**, Jae-Hyeok Shim2, Hansang Lee3, Keunbong Yoo3 and Shi-Hoon Choi1* 1

Department of Printed Electronics Engineering, Sunchon National University, Sunchon 57922, Korea 2

Center for Energy Materials Research, Korea Institute Science and Technology, Seoul 02792, Korea

3

Power Generation Laboratory, Korea Electric Power Corporation Research Institute, Daejeon 34056, Korea

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* Corresponding author. Tel.: +82 61 750 3556; Fax: +82 61 750 5260. E-mail: [email protected] (Shi-Hoon Choi). ** Co-corresponding author. Tel.: +82 2 958 6805; Fax: +82 2 958 5449. E-mail: [email protected] (Jin-Yoo Suh).

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Abstract: The microstructural factors that contribute to hardening mechanisms were investigated

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to explain the effect of post-weld heat treatment (PWHT) on the hardness of P92 steel in IN740H/P92 dissimilar weld joints. This study presents experimental analysis of the distribution

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of microstructural factors such as precipitate size, precipitate fraction, grain size and dislocation density in the heat-affected zone (HAZ) and base metal (BM) of the martensitic heat-resistant

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steel. Although precipitates mainly decorated the prior-austenite grain boundaries (PAGBs), packet boundaries (PBs), and block boundaries (BBs), their spatial distribution strongly depended

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on the distance from fusion line and the PWHT conditions. The grain size and the densities of geometrically necessary dislocations (GNDs) also exhibited a non-uniform distribution. The

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individual contributions of the microstructural factors to hardness were explained by introducing

mechanisms.

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a simple hardening equation that considers the independent effect of different hardening

Keywords: PWHT, weld joints, precipitate, dislocation, hardness

1. Introduction High Cr martensitic heat-resistant steels have been developed due to the demand for enhanced thermal efficiency and reduced CO2 emissions from fossil-fuel power plants. Advanced Ultra-Super-Critical (A-USC) power generation systems operate at steam temperatures above 700 °C and at pressures of 35 MPa [1]. Operating a turbine under steam temperatures and pressures that are higher than those conventionally used requires the use of materials with higher levels of creep and oxidation resistance. This new A-USC system requires the use of Ni-based superalloys in both 1

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superheater tubing and header piping to replace Fe-based alloys. However, for the low-temperature components operating below 650 °C, 9Cr steels are still considered in order to reduce costs. Therefore, dissimilar metal welding between 9Cr steels and Ni-based superalloys is expected to become more common for the construction of A-USC systems. For 9Cr steels, two grades have been widely considered: 91 and 92 grades. P91 steel (9Cr-1Mo-VNb) strengthened by the addition of Nb and V is widely used at temperatures approaching 600 °C [2,3]. P91 steel was developed by adding a

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small amount of carbide- and nitride-forming elements such as N, V, and Nb to ordinary 9Cr-1Mo

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steel. P91 steel has excellent mechanical properties and creep resistance, which are achieved by the

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formation of a thermally stable fine V- and Nb-rich precipitates. P92 (9Cr-1.8W-0.5Mo-VNb) steel is strengthened by replacing part of Mo with W, and is currently being applied to the boiler components

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of the USC power generation systems operating at temperatures around 625 °C. P92 steel is a

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promising material for next-generation nuclear and thermal power plants due to its excellent

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mechanical properties at high temperatures, good weldability, low thermal expansion, high thermal conductivity, and adequate resistance to corrosion and stress corrosion cracking [4]. Stabilized

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microstructure by tempered martensite lath, M23C6 type carbide, high dislocation density, precipitate

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hardening of MX carbonitrides and solid solution hardening of W is advantageous for high-temperature strength of this steel [5–9]. P92 steel has been extensively studied and is currently being applied to the high-temperature components of fossil-fuel power plants [1,4,6]. Large plants composed of high-temperature components require joining processes that generally are accomplished using fusion welding. Fusion-welded joints of P92 steel consist of three areas: a base metal (BM), a weld zone (WZ) (or weld metal), and a heat-affected zone (HAZ). The HAZ is a transition region between the weld metal and the heat-unaffected BM and can be subdivided into several zones depending on the temperature history each zone experiences during welding. The HAZ typically consists of a coarse-grained HAZ (CGHAZ) having a relatively coarse prior-austenite grains (PAG) located next to the weld metal, a fine-grained HAZ (FGHAZ) with a relatively fine PAG adjacent to the CGHAZ, and an inter-critical 2

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HAZ (ICHAZ) in which the experienced welding peak temperature lies between Ac1 and Ac3, and tempered BM without a significant thermal effect (which is also known as over-tempered region) [9– 11]. The heterogeneity of the microstructure formed across the weld joint, geometry of the weld joint, and loading conditions all affect creep deformation and fracture behavior of the joint by creating a complex stress state throughout the joint [12]. Cracks or fractures in weld joints are classified according to the location of the cracks that occur across a joint. In general, fractures can be classified

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into Type I (cracking in weld metal), Type II (cracking in weld metal and HAZ), Type III (cracking in

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CGHAZ), and Type IV (cracking in either FGHAZ or ICHAZ) [13,14]. Type IV cracks in weld joints

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have been examined by many researchers to understand the phenomenon from a metallurgical point of view [1,4,11,15–17]. The microstructure of weld is also known to be influenced by welding

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methods and welding conditions [18–21]. In order to optimize the microstructure of welds, it is

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common to perform post-weld heat treatment (PWHT) for the 9Cr steels [1,22–25]. The conditions of

weld joints [19].

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the PWHT may affect the mechanical properties as well as the microstructural development in the

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However, the effect of PWHT on the microstructure and mechanical properties at room temperature

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of weld joints consisting of P92 steel has not been fully understood. In the present study, dissimilar weld joints were fabricated using IN740H and P92 steel via conventional fusion welding technology. The changes in the microstructural factors such as precipitate size, precipitate fraction, grain size and dislocation density in P92 steel of IN740H/P92 dissimilar weld joints before and after PWHT were investigated experimentally. The contribution made by individual microstructural factors to hardness was evaluated using a simple hardening equation.

2. Experimental and Computational Procedures IN740H/P92 dissimilar weld joints were fabricated using TIG (Tungsten Inert Gas) welding with IN740H with a thickness of 12 mm and P92 steel of 11.8 mm. The IN740H was provided in the form of rolled plate. The P92 steel was initially extruded in the form of a pipe and processed into a plate for 3

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this study. ERNiCr-3 was used as filler metal. P92 steel extruded to an outer diameter of 168.3 mm with a thickness of 11.8 mm was normalized at 1,050 C for 20 minutes and tempered at 775 C for 30 minutes. Welding direction (WD) for IN740H and P92 steel was set perpendicular to the rolling direction (RD) and the extrusion direction (ED) of each specimen coordinate system. Ar and 25% He were used as protective gases during the TIG welding. The preheating and interval temperatures were

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220 and 300 C, respectively.

Fig. 1. IN740H/P92 dissimilar weld joints: a section perpendicular to the WD of IN740H/P92

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dissimilar weld joints.

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Table 1 shows the chemical composition of the alloys used in this study. Specimens of dissimilar metal joints were fabricated by TIG welding with a total of 10 passes. In order to analyze the microstructure developed in the P92 steel before and after the heat treatment of IN740H/P92 dissimilar weld joints, PWHT was performed in air at 740, 760, and 790 C for 2 hours.

Table 1. Chemical compositions of the alloys used in this study. Material

C

Si

Mn

Cr

Mo

W

P

S

Al

V

Ti

Nb

Ni

Co

Fe

P92

0.11

0.50

0.45

9.0

0.45

1.75

0.02

0.01

0.02

0.2

0.01

0.05

0.4

-

Bal.

IN740H

0.03

0.15

1.0

24.5

0.1

-

0.03

0.03

1.35

1.35

1.5

Bal.

20

3.0

ErNiCr-3

0.02

0.11

3.09

20.3

-

-

-

-

-

-

-

Bal.

-

1.0

4

-

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The microstructure of the P92 steel was examined in a section perpendicular to the WD, as shown in Fig. 1. The specimens before and after the PWHT were subjected to mechanical polishing and final polishing using a diamond suspension and colloidal silica, respectively. To identify the interface between the WZ and the HAZ (WZ/HAZ interface) and the interface between the HAZ and the BM (HAZ/BM interface) developed in the P92 steel, a macro-etching was conducted using a solution

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(HCl 20mL, Ethanol 40mL, DI water 40mL), and the microstructure was observed using an optical microscope (Olympus GX-51). To observe the precipitates of P92 steel, chemical etching was

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performed using Vilella’s reagent (45 ml of Glycerol, 30 ml of HCl, 15 ml of HNO3). Then, the size

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and fraction of the precipitates existing in the HAZ and the BM were measured at the center in

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thickness direction using FE-SEM (JEOL JSM-7100F). For the precipitates not observable in

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FE-SEM due to the small size of nanometer scale, transmission electron microscopy (TEM) (Talos

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F200X, FEI, USA) operated at 200 kV equipped with a super-X EDS system was used to carry out

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energy dispersive spectrometer (EDS) elemental mapping on several different areas of interest to measure the number and size of the precipitates and selected area diffraction pattern (SADP) analyses to confirm the precipitate identification. The TEM samples were prepared using a focused ion beam (FIB, Helios Nanolab 600i, FEI, USA) by carefully defining the exact location of the FGHAZ for each specimen. Diffraction data were used to identify and confirm the precipitates. Details of the diffraction data, however, are not presented here. In order to observe the grain size and Kernel Average Misorientation (KAM) distribution, mechanical polishing was performed using a diamond abrasive solution and colloidal silica, and then an EBSD (electron backscatter diffraction) attached to FE-SEM (JEOL JSM-7100F) was used. EBSD analysis was performed by establishing a scan area of 5

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100100 m2 at a scan interval of 0.2 m. In order to understand the correlation between microstructure and mechanical behavior, hardness was measured using a Vickers hardness tester (Matsuzawa, MMT-X7 Hardness tester) under a load of 200 gf at a holding time of 10 seconds. The interval between indentations was 0.15 mm, and a mean value was taken after 3 measurements.

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Phase equilibria of the steel were calculated using the thermo-kinetic software, MatCalc (version

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6.01) [26] together with the mc_fe thermodynamic database (version 2.058). The phases considered

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in the calculation were liquid, BCC, FCC, cementite, M23C6, MX, M6C, M7C3, Laves-phase, -phase

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and Z-phase.

Fig. 2. Optical microscopy image showing the four different regions for microstructure analysis of P92 steel marked on a macro-etched WD section.

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3. Results Analysis of the optical microstructure of the specimens before and after the PWHT of P92 steel, which is not shown here, showed no significant difference along the thickness direction. Fig. 2 shows the region used for the microstructure analysis of P92 steel on the macro-etched WD section. The red dashed line corresponds to the WZ/HAZ interface (fusion line) and the blue dashed line corresponds

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to the HAZ/BM interface. As defined in Fig. 2, the microstructures of the four regions were analyzed.

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The regions were named differently according to their locations as follows. The nearest region to the

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WZ/HAZ interface was designated as CGHAZ, the middle region between the WZ/HAZ interface and the HAZ/BM interface as FGHAZ I, and the nearest region to the HAZ/BM interface as FGHAZ

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II. The FGHAZ II is expected to include ICHAZ and over-tempered region. Also, a region distanced

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about 3 mm from the HAZ/BM interface was defined as BM and used as a reference microstructure.

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Fig. 3 shows the equilibrium phase fraction of the P92 steel as a function of temperature calculated by the MatCalc program [26]. As expected, BCC phase is stable at temperature lower than

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approximately 780 °C, while FCC phase becomes stable at temperature higher than approximately

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850 °C. Both BCC and FCC phases are stable between these temperatures (inter-critical region). The results predict the thermodynamically stable temperature ranges for a few precipitation phases: Nb(C,N) (< 940 °C and 1,140 ~ 1,220 °C), V(N,C) (710 ~ 1,170 °C), Laves phase (< 740 °C), and M23C6 (< 875 °C). Since Laves phase is known to form during a long-term aging at high temperature [27], M23C6, V(N,C) and Nb(C,N) are the major phases expected to form in the steel after PWHT. Nb(C,N) appears in the two different temperature ranges. According to MatCalc result, the Nb(C,N) stable at high temperature (1,130 ~ 1,220 C) has a slightly higher nitrogen content than the low temperature Nb(C,N), although they have the same crystal structure (NaCl (B1) structure) and are composed mainly of Nb and C. Around the temperature between 1,100 and 1,200 C, Nb(C,N) and V(N,C), both of which have the same crystal structure, appear to compete with each other. At

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temperature lower than 1,170C, the fraction of V(N,C) increases with the decrease in temperature, which is accompanied by the decrease of the fraction of Nb(C,N). In this process, nitrogen in Nb(C,N) tends to move to V(N,C), which seems to be responsible for the dissolution of Nb(C,N) below 1,140 °C. According to the variation of the phases with temperature as predicted by the MatCalc program (at equilibrium), Nb(C,N) appears again below 940 °C, as the solubility of Nb and C in

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ferrite decreases with a decrease in temperature.

Fig. 3. The equilibrium phase amount of the P92 steel as a function of temperature calculated by the MatCalc program.

Figs. 4 and 5 display microstructures of AW and PWHT-760 specimens observed at the regions defined in Fig. 2 under 4,000 and 10,000 magnifications, respectively. For the visibility of the micrographs, not all of the microstructures were included in main paper. The microstructures of PWHT-740 and 790 that did not show any remarkable difference with PWHT-760 were included in the Supplementary material. Historically, a three-level hierarchy has been used to explain the lath 8

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martensitic microstructure. From the top of this hierarchy, the original austenite grains are divided

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into several packets to form a set of blocks that consist of martensitic laths [28–30].

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Fig. 4. SEM micrographs of the regions defined in Fig. 2 under 4,000 magnification.

Analyses of as-welded (AW) specimens have shown that BM has a typical tempered martensitic

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structure and relatively coarse precipitates with white color decorating the block boundaries(BBs),

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packet boundaries(PBs) and prior austenite grain boundaries(PAGBs) [31]. In the present study, however, it was difficult to distinguish between BBs and PBs in the FE-SEM images. In the AW

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specimen, the CGHAZ showed a typical martensitic structure. The number density of precipitates is lower than the other regions. The AW specimen showed that the distribution of precipitates in the FGHAZ I and FGHAZ II was different from those of the BM and the CGHAZ. In the FGHAZ I and FGHAZ II, the precipitates of significant amount were dispersed in the matrix. The precipitates dispersed in the matrix seemed to correspond to the precipitates that used to exist on the BBs, PBs and PAGBs before welding and those precipitates were not completely dissolved into the matrix during welding. In the FGHAZ II, precipitates appeared mostly in the matrix of the tempered martensite. The precipitates in the FGHAZ II were relatively larger than those in the FGHAZ I probably because the precipitates present on the BBs, PBs and the PAGBs before welding were less dissolved in the matrix 9

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by the lower peak temperature during welding than the peak temperature FGHAZ I region

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Fig. 5. SEM micrographs of the regions defined in Fig. 2 under 10,000 magnification.

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For the specimens PWHTed at different temperatures, 740, 760 and 790 °C, no remarkable

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difference in microstructure was observed among the three specimens for individual regions, CGHAZ, FGHAZ I, FGHAZ II and BM. However, the effect of PWHT on microstructure compared

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to that of AW specimen appears obvious. For example, the CGHAZ region of three specimens, PWHT-740, 760 and 790, has precipitates decorating both PBs and PAGBs unlike the CGHAZ of the AW specimen. The precipitates observed here are believed to have formed from the supersaturated solute elements during PWHT. Also, the precipitate size of the three specimens, PWHT-740, 760 and 790, looks relatively larger than that of AW. The result could probably be explained by the growth of existing precipitates during PWHT. The FE-SEM results of the P92 steel before and after PWHT, as shown in Fig. 4, were used to quantitatively analyze the area fraction of precipitates. In the present study, four different images were analyzed in one area to obtain statistically reliable data, which yielded an average value. Fig. 7(a) shows the average area fraction of the precipitates (observed by SEM) using the ImageJ software [32].

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The area fraction of precipitates in CGHAZ, FGHAZ I and FGHAZ II showed a relatively huge change by PWHT. On the other hand, the area fraction of precipitates in BM did not change significantly with increasing PWHT temperature, unlike other regions. Quantitative analysis of the precipitate size was conducted using the FE-SEM results measured at 10,000 magnification, as shown in Fig. 5. Fig. 7(b) shows the average size of the precipitates analyzed using the ImageJ software. For the BM, the size of precipitate was only slightly affected by PWHT, but, those in other

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regions were increased by PWHT, although they were not dependent on the PWHT temperature.

Fig. 6. TEM-EDS results measured for the FGHAZ I of AW and PWHT-760 specimens.

Fig. 6 shows the results of TEM-EDS for the FGHAZ I of AW and PWHT-760 specimens. The results of TEM-EDS for the FGHAZ I of PWHT-740 and PWHT-790 specimens were included in the Supplementary material. Analyses of the coarse precipitates revealed that carbides contained both Cr and W, and that the relatively fine precipitates were V and Nb containing carbonitrides (V(C,N)). 11

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TEM micrographs were used for the quantitative analysis of the average area fraction and the average

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size of the fine precipitates, V(C,N), which are shown in Figs. 7(c) and (d), respectively.

Fig. 7. Results from the analysis of the precipitates using the ImageJ software: (a) average area fraction of coarse precipitates, (b) average size of coarse precipitates, (c) average area fraction of fine precipitates, and (d) average size of fine precipitates.

The results, Figs. 7(c) and (d), showed relatively low levels for the area fractions of fine precipitates compared with those of coarse precipitates, and also showed that the size of the fine precipitates linearly increased with the increase in PWHT temperature. Fig. 8 shows the IQ (Image Quality) maps together with the grain boundaries obtained in the individual regions of AW and PWHT-760 specimens using EBSD. To aid the reader's understanding,

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not all of the IQ maps that did not show any remarkable difference under the different PWHT conditions were included in main text. Instead, only representative IQ maps (AW and PWHT-760)

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were included in the paper and the rest were included in the Supplementary material.

Fig. 8. An IQ (Image Quality) map showing the grain boundaries of the individual regions obtained via EBSD.

High-angle grain boundaries (HAGBs) were clearly distinguished in all specimens, and low-angle grain boundaries (LAGBs) were observed together in some specimens. There is a high possibility that either the BBs or the PBs are also composed of HAGBs. In the EBSD results, the BBs were easily identified because of the relatively straight lines of the HAGBs, but it was difficult to distinguish between the PBs and the PAGBs because of their similar shapes. The purpose of the EBSD analysis was to measure the grain size in each region. Fig. 9 compares the average grain size obtained from the EBSD results. Since the misorientation angle used in calculating the grain size was 5°, most of the 13

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PAG, packets, and blocks could be regarded as individual grains. Analysis of the AW specimen revealed that the finest grains were observed in the FGHAZ I, and relatively coarse grains were distributed in the remaining regions. An interesting result is that the grain size showed little change with PWHT at different temperatures. This result indicates that microstructural changes such as grain

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growth were not significant during PWHT.

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Fig. 9. The average grain size obtained from the EBSD.

In CGHAZ, the grain size of PWHT-760 specimen was finer than that of AW and other PWHT specimens. This result cannot be easily explained and, therefore, could be attributed to the scattering nature of the welding microstructure originated from the complex procedure of the welding. As described in Section 2 (Experimental and Computational Procedures), 5 layers are stacked with 10 passes during the TIG welding. During the manufacturing of several weld specimens, it is possible to change the heat input due to the difference of the welding conditions such as passing speed and layer thickness at every pass finely each time. As a result, the slight variation in CGHAZ microstructure such as grain size could be induced.

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Fig. 10 shows the density of the geometrically necessary dislocations (GNDs) calculated from the EBSD results. Only representative results (AW and PWHT-760) were included in the paper and the rest were included in the Supplementary material. When the density of GNDs was calculated, the orientation of the nearest neighboring pixels was considered, and the cases with a misorientation angle of 5° or more were excluded from the calculations. The density of GNDs in the AW specimen was relatively high regardless of the region. The microstructure of the BM of the AW specimen was

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tempered martensite. Also, in other areas, martensite was developed due to the phase transformation

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during welding. As a result, the density of GNDs was high in those regions. In the FGHAZ I and

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FGHAZ II of the PWHT-740 and PWHT-760 specimens, the density of GNDs was decreased due to the tempering effect. In the CGHAZ of the PWHT-740 and PWHT-760 specimens, the decrease in

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the density of GNDs was insignificant. In the PWHT-790 specimen, the density of GNDs was

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relatively high regardless of position. When PWHT was performed at 790 °C, the specimens were

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exposed to the two-phase region, as explained in Fig. 3. As a result, some regions could be involved in the phase transformation from austenite to martensite during cooling stage after PWHT resulting in

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the higher density of GNDs.

Fig. 10. The density of geometrically necessary dislocations (GNDs), as calculated from the EBSD results.

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Fig. 11 quantitatively compares the variations in the density of GNDs in each region under the different PWHT conditions. A relatively high density of GNDs was observed in the AW and PWHT-790 specimens. Although it is not clear as to why the CGHAZ kept high level of GND even after PWHT, the thermal residual stress developed by the difference in thermal expansion coefficient of WZ and CGHAZ during cooling after welding and PWHT could have contributed to the unusual

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behavior of the area.

Fig. 11. Variation in the density of GNDs for each region under the different PWHT conditions.

Fig. 12 compares the hardness values measured in the regions defined in Fig. 2 under the different PWHT conditions. In AW specimen, the hardness ranged in the order of CGHAZ > FGHAZ I > BM > FGHAZ II. The reason for the relatively high level of hardness in the CGHAZ of the AW specimen seemed related to the effect of solid-solution hardening that is induced by solute atoms partially dissolved in the matrix without precipitation of carbides and carbonitrides in the region after welding. As shown in Figs. 4 and 5, the density of precipitates developed in this region of the AW specimen was relatively low with the exception of PAGBs. This tendency appears in Fig. 7(a), which is the area fraction of precipitates. The FGHAZ I of the AW specimen had the second-highest level of hardness. 16

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Precipitate hardening induced by small precipitates, as shown in Figs. 4 and 5 and solid-solution hardening seemed related to the relatively high level of hardness. Figs. 7(a) and (b) also indicate that the area fraction of precipitates was relatively large, whereas the size of precipitates was relatively small in this region. The PWHT-740 specimen showed a significant decrease in hardness in all regions compared to the AW specimen. In particular, the hardness of CGHAZ and FGHAZ I were significantly decreased to have hardness distribution in the order of CGHAZ > BM > FGHAZ I >

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FGHAZ II. In the AW specimen, the difference in hardness between different regions reached a

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maximum of ~175 Hv, whereas only ~50 Hv for PWHT-740. The reason why hardness was

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significantly decreased in the CGHAZ and FGHAZ I seemed related to precipitation in the form of carbides and carbonitrides induced by solute atoms such as Cr, W, C and N that were supersaturated

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in the AW specimen, as explained in Figs. 4 and 5. As shown in Figs. 7(a) and (b), the area fraction of

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the precipitates was increased in the PWHT-740 specimen, but since the size of precipitates

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insignificant.

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simultaneously increased, the increase in strength due to the precipitates was expected to be

Fig. 12. The hardness values measured for the regions shown in Fig. 2 under the different PWHT conditions.

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As shown in Fig. 12, the PWHT-760 specimen had a slightly higher hardness than that of the PWHT-740 specimen regardless of position. This tendency was caused by a grain size for the PWHT-760 specimen that was relatively small compared with that of the PWHT-740 specimen, as shown in Fig. 9. As shown in Fig. 12, the PWHT-790 specimen showed slight increase in hardness regardless of its position relative to the PWHT-760 specimen. As shown in Fig. 7(a) and (b), the PWHT-790 specimen had a distribution of coarse precipitates similar to that of the PWHT-760

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specimen, but the PWHT-790 specimen had a lower area fraction (Fig. 7(c)) and a bigger size (Fig.

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7(d)) of fine precipitates. This result indicates that the precipitate contribution to hardness is

hardening, as shown in Figs. 10 and 11.

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insignificant. Increases in the density of GNDs, however, seemed to exert a substantial effect on

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Fig. 13 schematically shows the precipitation behavior in each region before and after PWHT. The

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schematic drawing is based on the quantitative viewpoints acquired by SEM, TEM-EDS and EBSD.

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During welding, the whole part of CGHAZ is transformed into austenite phase due to exposure above the AC3 temperature, and coarse austenite grains are formed due to the grain growth under the high

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peak temperature. Most of the M23C6 and MX precipitates inherited in the BM are dissolved to

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facilitate grain growth, and upon cooling, the coarse austenite phase is transformed into martensite to form PBs and BBs. The supersaturated solute atoms precipitate as M23C6 and MX during cooling and PWHT. M23C6 mostly precipitates in PAGBs and MX tends to form in PBs and BBs. The grain size change after PWHT was insignificant, and the precipitate tended to grow slightly. During welding, the whole region of FGHAZ I, is also transformed into austenite phases due to exposure to slightly higher temperatures than AC3. The fine precipitates are completely dissolved, but some of the relatively coarse precipitates are not completely dissolved but remain. The relatively low heat input and undissolved precipitates inhibits grain growth of austenite phase. The austenite phase is transformed into martensite phase during cooling after welding. The supersaturated solute atoms newly precipitate as M23C6 and MX in PAGBS and matrix during cooling and PWHT. The grain size after PWHT did not show any distinct change but slight growth of precipitate occurs. The whole 18

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region of FGHAZ II is exposed between the AC1 and AC3 temperatures during welding, resulting in the transformation of the austenite phase in only a partial region. Similar to FGHAZ I, the precipitate does not dissolve completely. After welding, the part transformed into austenite phase is transformed into martensite phase, and fine M23C6 is newly precipitated in PAGBs during cooling and PWHT.

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After PWHT, some growth of undissolved and newly precipitated precipitates occurs.

Fig. 13. Schematic diagrams showing the distribution behavior of precipitates in each region before and after PWHT.

4. Discussion The relationship between microstructure and strength of materials can be understood by classifying the influence of each hardening mechanism. The hardening mechanisms for these materials include solid-solution hardening (SSH), grain-boundary hardening (GBH), precipitation 19

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hardening (PH) and strain hardening (SH). The increase in yield strength brought about by all these hardening mechanisms was assumed to follow a simple additive rule, as shown by Eq. (1). (1) The yield strength,

, of the ferrous alloy can be simplified by accounting for the increases in

strength due to each of the hardening mechanisms. (2) is the friction stress of the ferrous alloy.

of

where

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Eq. (2) can be expressed more specifically as a function of the metallurgical factors as follows

-p

[33,34]. ∑

re

( (

(3)

is the concentration of the i-th solute

lP

where ki is the strengthening coefficient for the i-th solute,

))

atom, ky is the strengthening factor for the grain size, d is the grain size, f is the area fraction of the

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precipitate, X is the diameter of the precipitate,  is a constant with a value that ranges from 0.3 to 0.6,

ur

G is the ferrite shear modulus (83103 MPa), b is the ferrite Burger’s vector length (0.24910-3 m), and  is the dislocation density. The constant, ky, for the grain boundary strengthening was adopted

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from a known value for low-carbon steels, 600 MPam1/2. Because dislocations were generated only by transformation,  was assigned with the lowest value of 0.3. Since it is difficult to measure the amount of solute atoms in Eq. (3), we have combined the terms of friction stress and solute atoms into one term, ((

)(

(

, as shown in Eq. (4). ))

(

)(

(

)))

(4) where ff and Xf are the area fraction and diameter of the relatively fine precipitates, respectively, and fc and Xc are for the relatively coarse precipitates. Relatively coarse precipitates, as illustrated in Fig. 6, correspond to Cr-rich M23C6 and relatively fine precipitates to V(C,N).

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Fig. 14 compares the hardening terms, which contribute to the yield strength shown in Eq. (4), for each of the regions defined in Fig. 2 under the different PWHT conditions. The fraction corresponding to the contribution of each hardening term to the obtained by the following procedure in the present study. The

is added to Fig. 14.

of the BM was obtained by uniaxial

tensile tests using a miniature specimen developed in the laboratory [35]. determined from the difference between the

is the value

of the BM can be

and the sum of the terms of the whole hardening

of

mechanisms. However, it is difficult to perform the uniaxial tensile test of other regions except for the

ro

BM because the size of the regions is too small. Because of these experimental difficulties, the

and hardness have a linear relationship.

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other regions was determined under the assumption that

in

re

According to the calculation based on Eq. (4), as shown in Fig. 14, the contribution of GBH to the

lP

was relatively large and was estimated to have no significant change induced by PWHT. In general, it is known that the yield strength increases as the size of packets and blocks in martensitic alloys was 15.2 ~ 26.9% for BM specimens

na

decreases [36,37]. The percentage contribution of GBH to the

ur

and 26.1 ~ 46.5% for PWHT specimens. Because the finest grain size was observed in the FGHAZ I, regardless of the PWHT conditions, the contribution by GBH in this region represented the largest

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portion. The contribution of PH to the

The percentage contribution of PH to the

turned out to be relatively low compared to that of GBH. was 6.3 ~ 14.4% for BM specimens and 13.0 ~ 28% for

PWHT specimens. Also, the effect of PWHT on PH was not remarkable. In the case of the coarse precipitates (CP), the area fraction and size simultaneously increased by PWHT as shown in Figs. 7(a) and (b), which was not very effective for increasing the number density of strengthening particles. The fine precipitates (FP) showed low percentage contribution to the

because their area fraction

was at a very low level, as shown in Fig. 7(c). The contribution of SH to

was similar to that of PH.

As shown in Fig. 11, the density of the GNDs was relatively high in the AW and PWHT-790 specimens, indicating that the contribution of SH to the

21

was relatively large in these specimens.

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Fig. 14. The hardening terms, which contribute to the yield strength shown in Eq. (4), for each region

Finally, the

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defined in Fig. 2 under the different PWHT conditions.

, which is the sum of the

and SSH, represented the largest contribution to the

of

the AW specimen. In particular, the CGHAZ showed the largest value. As shown by the hardness in Fig. 12, supersaturated solute atoms in the CGHAZ after welding induced relatively high portion of SSH. This effect tended to decrease as the region of analysis moved from the FGHAZ I to the BM. The higher peak temperature during welding tended to favor a more effective dissolution of either carbides or carbonitrides existing before welding to result in the better supersaturation of solute atoms at the completion of welding process (see the graph for CGHAZ in Fig. 14). The

decreased after

all conditions of PWHT, but the degree of decrease varied depending on the analysis region. In particular, both the CGHAZ and the FGHAZ I showed significant decreases because the

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supersaturated solute atoms were consumed to the growth of the existing precipitates and to the formation of new precipitates. In the FGHAZ II, as well as in the BM, the decrease of

was

relatively small. The FGHAZ II appears to be mainly composed of over-tempered region as supported by the lowest hardness shown in Fig. 12. Therefore, the only minor decrease in

of FGHAZ II and

BM is considered to be the result of a relatively small amount of supersaturated solute atoms in these

of

two regions.

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5. Conclusions

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In order to understand the effect that PWHT exerts on the microstructure and hardness of P92 steel

re

in IN740H/P92 dissimilar weld joints, the microstructural factors such as precipitate size, precipitate fraction, grain size, and dislocation density were experimentally analyzed as possible contributors to

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the hardening mechanisms. The distribution of microstructural factors in the four representative

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regions (CGHAZ, FGHAZ I, FGHAZ II, and BM) from the WZ/HAZ interface to the BM along the center of P92 steel was analyzed. The hardness values in the AW specimens showed distribution in

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the following order: CGHAZ > FGHAZ I > BM > FGHAZ II. The PWHT decreased the hardness of

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all regions especially with significant decrease for both the CGHAZ and the FGHAZ I. The contribution of GBH to

was relatively large, but there was no significant change induced by

PWHT. The precipitates mainly decorated the BBs, PBs and PAGBs in the AW and the PWHT specimens. The PH contributed relatively little to

compared with that of GBH. The density of

GNDs was relatively high in the AW and PWHT-790 specimens. The and SSH, represented the largest contribution to

, which is the sum of the

of AW specimens but not for the specimens after

PWHT. Supersaturated solute atoms in the CGHAZ and FGHAZ I after welding induced relatively high portion of SSH. In the CGHAZ and FGHAZ I,

decreased by a relatively large amount after

PWHT because the supersaturated solute atoms were consumed to the growth of the existing precipitates and to the formation of new precipitates.

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Acknowledgements This research was conducted by a research project titled “Evaluation of microstructure/mechanical properties in dissimilar weld zone of boiler tube” under the financial support of KEPCO. Also, the work done at KIST was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (Ministry of Science and ICT) (No. NRF-2016M3C1B5906957).

K.-Y. Shin, J.-W. Lee, J.-M. Han, K.-W. Lee, B.-O. Kong, H.-U. Hong, Transition of creep

ro

[1]

of

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Conflict of Interest

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There is no conflict of interest related to this paper.

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HIGHLIGHTS 

Effect of PWHT on the hardness of P92 steel in IN740H/P92 weld joints was investigated.



The PWHT significantly decreased the hardness of both the CGHAZ and the FGHAZ I.



The contribution of GBH to

was relatively large, but there was no significant change

induced by PWHT.. The density of GNDs was relatively high in the AW and PWHT-790 specimens.



Supersaturated solute atoms in the CGHAZ and FGHAZ I after welding induced relatively

of



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high portion of SSH.

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