Tempering effect on the fusion boundary region of

Tempering effect on the fusion boundary region of alloy 625 weld overlay on 8630 steel Tao Dai & John C. Lippold

Welding in the World The International Journal of Materials Joining ISSN 0043-2288 Weld World DOI 10.1007/s40194-018-0560-3

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Author's personal copy Welding in the World https://doi.org/10.1007/s40194-018-0560-3

RESEARCH PAPER

Tempering effect on the fusion boundary region of alloy 625 weld overlay on 8630 steel Tao Dai 1

&

John C. Lippold 1

Received: 3 September 2017 / Accepted: 18 January 2018 # International Institute of Welding 2018

Abstract The NACE standard MR0175 requires that the HAZ hardness of AISI 8630 steel overlaid with Ni-base Alloy 625 cannot exceed 250 VHN, requiring a postweld heat treatment (PWHT) to decrease the as-welded HAZ hardness. However, PWHT results in the diffusion of carbon from the steel into the overlay and may lead to interface embrittlement. To understand the effect of PWHT on the microstructure and hardness of the regions near fusion boundary, a wide range of PWHT conditions were investigated. The regions of interest include the coarse-grain HAZ (CGHAZ), the partially-mixed zone, the Bswirl^ structure, the planar growth zone (PGZ), and the Alloy 625 weld metal. The microstructure and hardness were evaluated using optical metallography, SEM, Vickers hardness testing, and nanoindentation. With an increase of the Hollomon-Jaffe parameter (HJP), the CGHAZ hardness decreases, the PGZ hardness increases, and the weld metal hardness also increases but at a much lower rate. At HJP > 19,300, carbide precipitation was observed in the PGZ. Following PWHT, the partially-mixed zone was found to contain fresh martensite since the PWHT temperature was above the A3 temperature of this region. Reducing the cooling rate from the PWHT temperature did not eliminate the untempered martensite in this region. Finally, the behavior of the 8630/625 overlay is compared and contrasted to the F22/625 overlay behavior which was reported previously. This work builds a foundation for future studies on hydrogen-assisted cracking and sulfide stress cracking in Alloy 625 overlay on steels used in the oil and gas industry. Keywords Alloy 625 . Low alloy steel . Heat-affected zone . Fusion boundary . Postweld heat treatment . Carbon migration . Metallography . Hardness . Hardness testing

1 Introduction AISI 8630 steel is a quench and tempered, low alloy forging steel containing lower alloy content but higher carbon content than other steels used in the oil and gas industries such as F22 (2.25Cr-1Mo) [1]. Both of these steels are used for components of subsea architecture for oil extraction and transportation. To protect 8630 steel components from being exposed to the production fluid containing H2S, the corrosion-resistant alloy 625 (ERNiCrMo-3) is clad on the inner surface of the components.

Recommended for publication by Commission IX - Behaviour of Metals Subjected to Welding * Tao Dai [email protected] 1

Welding Engineering Program, The Ohio State University, 1248 Arthur E Adams Dr, Columbus, OH 43221, USA

For some of the components, only the connecting outlet is clad with alloy 625 for leak prevention. At the connecting outlet with alloy 625 overlay, the fusion boundary is exposed to the crude oil and risks the problem of sulfide stress cracking (SSC) [2]. The cladding procedure creates a heat-affected zone (HAZ) in 8630 which contains mainly fresh martensite with high hardness, particularly the coarse-grain HAZ (CGHAZ). In this scenario, there are four conditions for the occurrence of SSC: martensite with high hardness, tensile stress exerted by highpressure production fluid on the components, hydrogen from the production fluid, and ambient temperature. In order to reduce the possibility of SSC, the hardness of the HAZ was reduced below a certain critical level which is dependent on the stress level and hydrogen content. In order to insure resistance to hydrogen embrittlement, the National Association of Corrosion Engineers (NACE) Standard MR0175 and the International Standards Organization (ISO) 15156 requires that maximum hardness of the as-tempered steel does not exceed 22 HRC or 250 VHN [2].

Author's personal copy Weld World

Postweld heat treatment (PWHT) or tempering is commonly used to reduce the hardness of the HAZ. During the PWHT, the carbon in 8630 steel diffuses from the base metal across the fusion boundary to the weld metal. The diffusivity of carbon in austenitic weld metal is low, resulting in a Bpile up^ of carbon in a narrow region along the fusion boundary on weld metal side. High carbon concentration results in very high hardness levels at the fusion boundary [3]. In cases where extreme PWHT is required to soften the HAZ, a reduction in base metal strength may occur that compromises the performance of the component. Ideally, a well-balanced PWHT condition is desirable in order to decrease the HAZ hardness while avoiding embrittlement of the fusion boundary and subsequent resistance to SSC. The 8630/625 overlay has been studied by a number of researchers [3–7], including both the microstructure and chemical composition gradients across the fusion boundary. In general, these studies were limited to the as-welded condition and two to three PWHT conditions [4, 5]. These studies provided valuable insight, but were not sufficient to determine the optimal PWHT conditions at which the susceptibility of SSC is minimum. The evolution of HAZ microstructure and hardness of 8630/625 overlay has also not been thoroughly evaluated. This investigation represents a parallel study of the F22/ Alloy 625 work that has been previously reported [1]. As in the previous study, a wider range of PWHT (11 conditions) was studied to reveal the variation of microstructure and hardness, including some extreme PWHT conditions, i.e., 660 °C/ 100 h, which is not applied in practice but provides research insight. Optical and SEM microscopy was used to study the microstructure near the fusion boundary, especially the CGHAZ. Vickers hardness testing and nanoindentation were utilized to reveal the hardness variation with PWHT. The investigation methods of 8630/625 overlay are similar to those of F22/625 overlay, but the results contain some significant difference due to the difference in chemical composition of the two steels. This study serves as the groundwork of future hydrogen cracking tests and sulfide stress cracking tests on 8630/625 overlay samples subjected to various PWHTs.

2 Experimental procedures 2.1 Materials The 8630 steel was in the form of a ring forging with its mechanical properties listed in Table 1. The filler metal used Table 1 Mechanical properties of 8630 base metal 8630

for cladding was Ni-base Alloy 625 (ERNiCrMo-3). Overlay welding was conducted using the hot-wire gas tungsten arc welding (HW-GTAW) process with the welding parameters shown in Table 2. The chemical compositions of AISI 8630 and Alloy 625 are shown in Table 3. A cross section showing the Alloy 625 overlay on the 8630 ring forging is shown in Fig. 1. The thickness of the overlay is much greater than standard practice in order to facilitate the machining of samples for delayed hydrogen cracking tests (DHCT) and SCC tests. In both these tests, the fusion boundary is centered in the middle of the sample.

2.2 Postweld heat treatment on 8630/625 overlay samples In order to study the effect of tempering conditions on microstructure and hardness, 11 PWHT conditions were used, as shown in Table 4. Temperature (T) in kelvin and time (t) in hours are two parameters of each PWHT condition, which is quantified using the Hollomon-Jaffe parameter (HJP), as calculated in Eq. 1 [8]. The first three PWHTs with relatively small HJPs may be impractical since they cannot achieve the maximum HAZ hardness of 250 VHN. The last three PWHTs are so severe that they would compromise the yield strength of base metal by overtempering. These six are still useful from a research standpoint in order to better track the tempering and carbon diffusion effects. The samples were PWHTed in a horizontal Lindberg 59544-type furnace in an inert atmosphere with titanium sponge and continuous argon flow. Thermocouple monitoring was used to verify the actual temperature during the PWHT. The time of heating the sample to target temperature was usually 10–20 min and not considered in the calculation of the HJP. All PWHT samples were quenched in room-temperature water. Considering that air cooling is used in the PWHT in practice, the effect of the cooling rate on the microstructure and hardness needs also to be studied. Therefore, two more samples were prepared for PWHT at 670 °C for 10 h, and then cooled using air cooling (AC) and furnace cooling (FC) for comparison to the water quenched (WQ) condition. The cooling rates for the WQ, AC, and FC from 670 to 30 °C were measured at approximately 300, 1, and 0.013 °C/s (~ 0.8 °C/min), respectively, for the samples with 0.1 in. thickness [8]. HJP ¼ T *ð20 þ logðt ÞÞ

ð1Þ

Ultimate tensile strength

Actual yield strength

% EL

% RA

Hardness (VHN)

116 ksi or 800 MPa

98.3 ksi or 678 MPa

23.0

69.1

250

Author's personal copy Weld World Table 2

Welding parameters for Alloy 625 overlay on 8630 steel forging

Step increment

Voltage

Peak current

Background current

Hot wire voltage

Wire feed speed

0.150 in. or 3.81 mm

11.0 V

240.0 A

120.0A

3.3 V

Interpass temperature

Travel speed

Hot wire peak current

550 °F or 287.8 °C

7.9 in./min or 200.66 mm/min

179.0 A

Heat input Hot wire background current 55.0 A 10.0 kJ/in. or 0.394 kJ/mm

66.9–82.2 in./min 350 °F or 176.7 °C or 170–208.8 cm/min Deposition rate Shielding gas flow rate (argon)

2.3 Hardness measurements and microstructure analysis

30.0 cfh or 14.17 L/min

3 Results 3.1 As-welded hardness Figure 2 shows a microhardness map and indent pattern of the as-welded 8630/625 condition and illustrates the methodology of collecting and plotting the hardness data. In the as-welded condition, the base metal (not the HAZ) has similar hardness to the weld metal. The HAZ shows higher hardness but with a nonhomogeneous hardness distribution. The higher hardness regions (yellow and red regions) and the tempered regions (green regions) in the HAZ appear periodically, approximately every 4 mm, resulting from the tempering effect of the overlapping, multi-bead overlay process. The distinct transition between the higher hardness HAZ and the low hardness WM corresponds to the fusion boundary in Fig. 2a and c. The maximum hardness in the HAZ is 466 HV0.1, but an average hardness in the HAZ (or CGHAZ) cannot be obtained due to the large local variations. By plotting all the hardness data as a function of distance perpendicular to the fusion boundary, the collective hardness of the BM, FGHAZ,

2.4 Nanoindentation An MTS Nanoindenter® XP system was used to conduct nanoindentation. There are two modes for nanoindentation: XP mode and CSM mode. Only load control is available for the XP mode, which was used for making nanohardness line scans. The constant maximum load used is 3 gf for all the line indentations. The spacing between indents was 6 μm, less

8630 625

2.0 lb./h or 0.907 kg/h

than three indent-diameters. Different from the Vickers hardness indentation, the spacing between nanoindents has less effect on the measured hardness values. Even if the indent spacing is as small as one indent size, reliable nanohardness and modulus can be obtained [10]. Prior to nanoindentation, the sample was polished using 0.5-μm colloidal silica. Nanohardness data in units of gigapascals was also correlated to Vickers hardness (VHN) as described in a previous paper [1].

Samples were mounted in conductive Bakelite, ground with abrasive papers through no. 800 grit, and polished through 1-μm diamond. Vickers hardness mapping was conducted in the as-polished condition using a Vickers hardness testing device (LECO LM100AT) with a 100-g load. Each hardness map grid was 40 × 78 with 100-μm spacing between indents. The hardness grids were located in a way that the 40 indents in the y-axis spanned the steel base metal, HAZ, fusion boundary region, and weld metal, and the 78 indents in the x-axis direction spanned two adjacent overlay weld beads, each of which is 0.15 in. or 3810 μm. Individual Vickers hardness indentations were also made at multiple locations along the fusion boundary to determine the hardness in the planar growth zone as described in a previous paper [9]. The samples were then repolished as described above for nanoindentation and then etching. After nanoindentation, each sample was first chemically etched using 5% nital (5 vol.% nitric acid + 95 vol.% ethyl alcohol) to reveal the steel microstructure. The sample was electrolytically etched in 10% chromic acid at 5.0 V for 5 s to reveal the weld metal microstructure.

Table 3

Preheat temperature

Chemical composition (wt%) of 8630 steel and Ni-base filler metal Alloy 625 C

Ni

Cr

Mn

Si

Mo

S

P

Al

Cu

Ti

Nb

V

Ta

Fe

0.32 < 0.01

0.86 64.0

0.97 22.7

0.90 < 0.01

0.32 0.04

0.41 9.0

0.009 0.001

0.008 < 0.01

0.026 0.12

0.17 < 0.01

0.002 0.23

– 3.59

0.032 –

– 0.004

bal. 0.3


periodically, around which there is usually a particularly wide Bfeatureless^ zone. Figure 4b shows a typical swirl structure. Between the white featureless zone and the swirl, there is a composition-partially-mixed zone with different etching characteristics from the overlay weld metal. Much of the partiallymixed zone transforms to martensite during cooling from PWHT temperature. Figure 4c shows the Bbottom^ of the cusp, where the featureless zone is narrowest. The featureless zone consists of a continuous band of planar solidification zone, or so-called planar growth zone (PGZ). The penetration of weld metal into the base metal HAZ, referred to as a Bfinger^ structure, was also observed along the 8630/625 fusion boundary (Fig. 4c). This results from the penetration of Alloy 625 weld metal along prior austenite grain boundaries of the base metal.

Fig. 1 8630/625 dissimilar metal weld coupon

3.3 Tempering behavior of the CGHAZ CGHAZ, and WM can be represented as shown in Fig. 2d, and the average and maximum hardness can be determined.

3.2 As-welded microstructure The microstructures of the 8630 steel BM, FGHAZ, and CGHAZ and the Alloy 625 WM in the as-welded condition are shown in Fig. 3. The BM consists of fine grain ferrite with carbides. The FGHAZ and CGHAZ contain mainly martensite with apparent lath structures. The WM consists of a cellular dendritic microstructure with interdendritic precipitates that presumably result from the NbC/γ and/or Laves/γ eutectic reaction and exhibit a large variation in size and shape. There is no evidence of solidification cracking or other defects. The multi-bead overlay process produced a fusion boundary with periodic cusps, or curvatures. A continuous Bfeatureless^ zone is present at the fusion boundary, the width of which varies periodically. The swirl structures also appear Table 4

Tempering treatments on 8630/625 overlay samples

Heat treatment

Temperature

Time

HJP

Cooling

HT-1 HT-2 HT-3

600 °C or 1112 °F 640 °C or 1184 °F 620 °C or 1148 °F

5h 2h 6h

18,070 18,535 18,555

WQ WQ WQ

HT-4 HT-5 HT-6 HT-7 HT-8 HT-9 HT-10 HT-11

660 °C or 1220 °F 640 °C or 1184 °F 650 °C or 1202 °F 660 °C or 1220 °F 670 °C or 1238 °F 660 °C or 1220 °F 660 °C or 1220 °F 660 °C or 1220 °F

2h 6h 10 h 10 h 10 h 20 h 50 h 100 h

18,941 18,970 19,383 19,593 19,803 19,874 20,154 20,526

WQ WQ WQ WQ WQ, AC, FC WQ WQ WQ

The as-welded 8630/625 samples were tempered over a range of PWHT conditions in order to evaluate the effect on microstructure and hardness of the HAZ. The CGHAZ which exhibits higher hardness than the FGHAZ is the region of interest in terms of the susceptibility to HAC and/or SSC. The actual CGHAZ below the overlay is not homogeneous due to the partial tempering effect of the multi-bead overlay process. Six CGHAZ microstructures that were not affected by the multi-bead tempering are shown in Fig. 5. Ranging from low to high HJP values (Fig. 5a–f), the progression of decomposition of martensite can be observed. For the slightly tempered martensite at 620 °C/6 h and 640 °C/6 h, the lath structures are still clear (Fig. 5a, b). For the medium-tempered martensite at 660 °C/10 h and 670 °C/10 h, the lath structure sub-cells become more diffuse and there appears to be no evidence of particular more carbide precipitation (Fig. 5c, d). At the extreme tempering conditions, martensite is replaced by ferrite. For the sample tempered at 660 °C/50 h, there are still some carbides and sub-cells left in the ferrite matrix (Fig. 5e). Tempering at 660 °C/100 h even dissolved the carbides and produced a homogenous ferritic microstructure (Fig. 5f). The ferrite grain size also grew larger during tempering due to suppression of boundary pinning by the carbides on the grain boundaries.

3.4 Microstructure of the swirl structure and partially-mixed zone Another microstructure that is potentially susceptible to HAC/ SSC is within the composition-partially-mixed zone adjacent to the fusion boundary and/or between Bswirl^ structures and the fusion boundary. Characteristic swirl structures and partially-mixed zone adjacent to them can be seen in Fig. 4a and are shown at a higher magnification in Fig. 6. Literally, the partially-mixed zone composition results from the partial


Fig. 2 8630/625 sample in as-welded condition. a HV0.1 hardness map. b Hardness indent matrix. c Overlap of hardness map on the indent matrix. d Hardness distribution across the fusion boundary

mixture of weld metal Alloy 625 and base metal 8630 steel. In the as-welded condition, the partially-mixed zone is the light gray region adjacent to the swirls or fusion boundary, as shown in Fig. 4b. Three samples were tempered at 670 °C/ 10 h and cooled using water quenching (Fig. 6a), air cooling

(Fig. 6b), and furnace cooling (Fig. 6c, d), respectively. The optical microscopy revealed fine lath structures in the partially-mixed zone of all three samples, similar to the martensitic microstructure in the CGHAZ (Fig. 6a). Vickers hardness of the partially-mixed zone with fine lath morphology is

Fig. 3 Representative microstructure in the 8630/625 overlay in the as-welded condition. a Base metal. b FGHAZ. c CGHAZ. d Alloy 625 WM

(a)

(b)

(c)

(d)


(a)

(b)

(c)

Fig. 4 As-welded 8630/625 sample. a Fusion boundary overview. b Higher magnification of swirls. c Narrow planar growth zone and weld metal penetration into the base metal HAZ

as high as typical hardness of martensite, higher than the hardness of the adjacent weld metal and base metal (Fig. 6a–c). Thus, it is assumed that the partially-mixed zone in Fig. 6a–c contain fresh martensite. In the partially-mixed zone, the composition of iron, nickel, chromium, and molybdenum reduces Ac3 to below the PWHT temperature, e.g., 670 °C. Therefore, during PWHT, the partially-mixed zone transformed to austenite and after quenching from the PWHT temperature, the austenitic partially-mixed zone transforms to fresh martensite. As expected, as the cooling rate decreases from water quenching to air cooling to furnace cooling, the portion of the partially-mixed zone transforming to martensite decreases. Interestingly, in the furnace cooled sample (at ~ 0.8 °C/min), there is still fresh martensite formation in the partially-mixed zone (Fig. 6c, d) reflecting the high hardenability of the mixed composition. This indicates that the formation of the untempered martensite may be unavoidable even at slow cooling rates from the PWHT temperature. If the PWHT temperature is reduced, e.g., 500 °C, the partially-mixed zone at the swirl structures does not form martensite and the hardness of the partially-mixed zone is only 215 HV0.1 (Fig. 6e), since this PWHT temperature 500 °C is probably below the Ac3 of the partially-mixed zone and did not transform to austenite during

PWHT. Such a low PWHT would not be effective in tempering the 8630 HAZ.

3.5 Effect of PWHT on Vickers hardness Figure 7a–h shows hardness maps from eight of the PWHT conditions and their corresponding hardness distributions, which span the BM, FGHAZ, CGHAZ, and WM. As the HJP increases, the WM hardness increases and the HAZ hardness decreases. After the 600 °C/5 h PWHT, hardness in the FGHAZ and CGHAZ both decreased, and the hardness in the FGHAZ decreased more, comparing Figs. 7a and 2d. Carbon diffusion and carbon pileup in the PGZ are not significant at 600 °C/5 h, 620 °C/6 h, and 640 °C/6 h, and only a few hardness peaks were observed along the fusion boundary in their hardness maps. After 660 °C/10 h, the CGHAZ hardness almost equals to WM hardness, and there are multiple hardness peaks up to 500 HV0.1 along the fusion boundary. After 670 °C/10 h, the CGHAZ hardness becomes less than the WM hardness and the peak hardness at the fusion boundary reaches 600 HV0.1. However, the CGHAZ hardness after 660 and 670 °C/10 h is still above 250 HV0.1, the required maximum of the NACE standard [2] (Fig. 7e, f).

Author's personal copy Weld World Fig. 5 CGHAZ microstructure of 8630-625 samples after PWHT (HJP value): a 620 °C/6 h (HJP = 18,555), b 640 °C/6 h (HJP = 18,970), c 660 °C/10 h (HJP = 19,593), d 670 °C/10 h (HJP = 19,803), e 660 °C/50 h (HJP = 20,154), and f 660 ° C/100 h (HJP = 20,526)

At PWHT 660 °C/50 h and 660 °C/100 h (Fig. 7g, h), the hardness in the base metal and HAZ decreases to less than 200 HV0.1. In fact, the hardness of the CGHAZ is less than that of the base metal. The black regions in Fig. 7g, h represent hardness levels lower than 160 HV0.1. This is because of extreme carbon depletion in these regions due to carbon diffusion from the HAZ across the fusion boundary and pileup along the fusion boundary. The peak hardness points at the fusion boundary of sample at PWHT 660 °C/100 h is up to 700 HV0.1. The measured hardness peaks are discrete rather than continuous along the fusion boundary, because of the relatively low resolution of the Vickers hardness map with 100 μm as the spacing distance. The hardness peaks are contained within the PGZ, in the range 10–50 μm in width. Therefore, independent Vickers hardness indentations were made along the fusion boundary to reveal the peak hardness, and the nanoindentation was also applied to reveal the hardness at the fusion boundary. The average and maximum hardness of the CGHAZ, WM, and fusion boundary PGZ from Fig. 7 are summarized in Fig.

(a)

(b)

(c)

(d)

(e)

(f)

8 as a function of the HJP. The CGHAZ hardness decreases with the HJP and is not reduced below 250 VHN until HJP ≈ 19,900. Among the 11 PWHT conditions, only 660 °C/50 h and 660 °C/100 h reduced the CGHAZ hardness below 250 VHN, meeting the NACE Standard MR0175/ISO15156 [2], which is represented with the scarlet dashed line in Fig. 8. The PGZ hardness increases with the HJP rapidly when HJP > 19,000 due to carbon Bpileup.^ Below HJP ≈ 19,000, the PGZ hardness remains around 300 VHN, even less than the CGHAZ hardness. Thus, there are no hardness peaks appearing in the hardness maps when HJP < 19,000 as shown in Fig. 7a–c. When the CGHAZ hardness is tempered below 250 VHN at HJP ≈ 19,900, the average PGZ hardness reaches about 500 VHN, and the maximum hardness point is up to 650 VHN, which could be very susceptible to the HAC or SSC. The error bars show that the variation of the PGZ hardness is larger than the CGHAZ hardness and also WM hardness. The WM hardness remains essentially constant at 275 VHN when the HJP < 19,500, and it increases slightly when the HJP > 19,500. It seems that the precipitation hardening in the WM does not occur until HJP > 19,500.

Author's personal copy Weld World Fig. 6 Effect of cooling rate from the PWHT condition of 670 ° C/10 h, 5% nital etch only. a Water quenching at a rate ~ 300 °C/s. b Air cooling at the rate ~ 1 °C/s. c Furnace cooling at the rate 0.8 °C/min. d Expanded view of area in the yellow frame in c. e PWHT 500 °C/4 h, water quenching

(a)

(b)

(c)

(d)

(e)

3.6 Effect of HJP on nanohardness at the fusion boundary Figure 4 shows that there is a periodic variation of the PGZ width along the fusion boundary with some regions as narrow as 10 μm. The HV0.1 indent size is 20–30 μm which approximates the average width of the PGZ making it difficult to determine the actual hardness using Vickers hardness measurements. Nanoindentation can be applied to increase the Bfidelity^ of hardness measurement to a scale of 6 μm. Figure 9 shows the effect of the PWHT on the microstructure and the nanohardness distribution across the fusion boundary. Six samples were selected to represent a wide range of HJP for the 8630/625 overlay. The section of fusion boundary without a swirl structure or partially-mixed zone is designated as the continuous fusion boundary, and the section with a swirl

structure or partially-mixed zone is designated as the discontinuous fusion boundary [4, 5]. For each sample, regions of both continuous and discontinuous fusion boundaries were selected for the evaluation, considering the complexity of the microstructure along the fusion boundary. For low values of HJP such as 600 °C/5 h (Fig. 9a) and 620 °C/6 h (Fig. 9b), the nanohardness in the base metal is similar to the weld metal. There is a small hardness peak in the PGZ due to carbon diffusion and pileup. Interestingly, in Fig. 9a, the hardness traverse includes a finger structure, in which is the hardness peak. On the right edge of the swirl, it is the martensitic partially-mixed zone with higher hardness, indicated with the green arrow in Fig. 9a–f. The nanoindentation revealed a high hardness of the narrower partially-mixed zone than the Vickers hardness testing reported in Sect. 3.4.


(a) 600˚C/5hr

(b) 620˚C/6hr

(c) 640˚C/6hr

(d) 650˚C/10hr

(e) 660˚C/10hr Fig. 7 Hardness maps and hardness distributions of PWHTed samples


(f) 670˚C/10hr

(g) 660˚C/50hr

(h) 660˚C/100hr Fig. 7 (continued)

The hardness peaks in the PGZ adjacent to the discontinuous fusion boundary are not clear even after severe PWHTs such as 650 or 670 °C/10 h, as indicated by the blue arrows in Fig. 9c, d. However, as the HJP increases, the hardness peak in the PGZ of the continuous fusion boundary becomes sharp and quite high. For PWHT 650 °C/10 h and more severe PWHT conditions, the PGZ becomes brownish in color, as indicated with the black arrow in Fig. 9d–h, rather than Bfeatureless.^ The brownish feature is due to carbide precipitation, which contributes to the sharp hardness increase.

As an example, Fig. 10 shows the correlation between nanohardness and Vickers hardness for the PWHT condition 670 °C/10 h. Because introducing a new sample and resetting the nanoindenter may create errors in testing, a correlation relationship was built independently for each sample shown in Table 5. Using these linear equations [1, 11–14], the nanohardness peak values of each sample can be converted to Vickers hardness values. The converted peak HV0.1 values are compared with actual measured HV0.1 values (from Fig. 8) in Table 5. The converted peak nanohardness values also increase with the HJP generally as the measured ones, but they exceed the measured ones in all cases, and the difference is

Author's personal copy Weld World Fig. 8 Summary of hardness response in the CGHAZ, WM, and PGZ areas as a function of the HJP

(a) 600˚C/5hr

(b) 620˚C/6hr

(c) 650˚C/10hr Fig. 9 Nanohardness traverses across the fusion boundary of the 8630/ 625 samples in six selected PWHT conditions. For each PWHT condition, the first figure shows a continuous fusion boundary and the second a discontinuous fusion boundary (containing a swirl pattern). The fusion

zone is on the left of the fusion boundary, and the base metal is on the right: a 600 °C/5 h, b 620 °C/6 h, c 650 °C/10 h, d 670 °C/10 h, e 660 °C/ 50 h, f 660 °C/100 h


(d) 670˚C/10hr

(e) 660˚C/50hr

(f) 660˚C/100hr Fig. 9 (continued)

Fig. 10 Correlation between nanohardness and Vickers hardness for the PWHT condition 670 °C/10 h


large under conditions where carbon diffusion is the greatest (high HJP).

4 Discussion 4.1 Effect of PWHT on microstructural features of 8630/625 overlay Eleven 8630/625 overlay samples over a wide range of PWHT conditions were investigated in this work, in an attempt to understand the evolution of microstructure and hardness at the fusion boundary, and determine HAC and SSC susceptibility as a function of HJP. This work is in preparation for actual HAC and SSC testing that will correlate susceptibility to PWHT conditions. Three distinct microstructural regions of interest near the fusion boundary have been observed. The first is the CGHAZ, which contains mainly martensite in the as-welded condition. As the HJP increases, the martensite decomposes as shown in Fig. 5. There is no obvious strong carbide precipitation in any of the PWHT conditions evaluated, nor a secondary hardening phenomenon as reported in F22 steels [1, 15]. In the samples PWHTed at 660 °C/50 h and 660 °C/100 h, the martensite was tempered to ferrite with clear grain boundaries (Fig. 5e, f). These microstructures are quite different from the CGHAZ of the F22/625 overlay in the same PWHT condition [1]. An explanation for this is that 8630 steel contains less chromium and molybdenum which are strong carbide formers within the PWHT temperature range. Thus, compared to the F22/625 sample, in the CGHAZ of the 8630/625 sample, more carbon can diffuse to the weld metal due to less carbide precipitation, and the dislocation and grain boundaries can migrate with less Bdrag force.^ During long-time tempering (660 °C/50 h and 660 °C/100 h), carbon diffusion out of the CGHAZ allows for ferrite formation and grain growth where the ferrite grains consume small grains, sub-cells, or lath structures left by martensite. Therefore, the CGHAZ microstructure of F22 steel after tempering is not as homogeneous or Bclean^ as that of 8630 steel. The residual carbides also dissolved as the tempering time increases, and the carbon continues to diffuse to the weld metal. This is why the ferrite grains are larger and Bcleaner^ with less sub-lath structures and less carbides in Fig. 5f (660 °C/100 h) than in Fig. 5e (660 °C/50 h). The second microstructure of interest is the so-called featureless zone, or the planar growth zone (PGZ), as reported by previous researchers [1, 3–6, 16]. The carbon diffuses from the CGHAZ to the featureless zone and piles up in this zone with an austenitic microstructure [3]. The PGZ has much higher content of chromium and molybdenum than the base metal, and it is hypothesized that carbide precipitation would occur in the PGZ after PWHT. Fenske and Dodge [4, 5] verified this hypothesis by identifying M7C3 carbides in the PGZ after PWHT

at 691 °C/5 h (HJP = 19,954) and 675 °C/10 h (HJP = 19,908), respectively. This characterization work using TEM was not repeated in this study. With color etching and optical microscopy, it was found that as the PWHT increases, there are more and more brownish features appearing in the PGZ (as indicated by purple arrows in Fig. 9), particularly at PWHT 650 °C/10 h, 670 °C/10 h, 660 °C/50 h, and 660 °C/100 h. The brownish feature is thought to be clusters of M7C3 carbides, as reported by Fenske and Dodge [4, 5]. The brownish feature is not immediately adjacent to the fusion boundary, but is generally closer to the cellular solidification region as shown in Fig. 9c–f. This is consistent with the finding that there is a carbide-free zone immediately adjacent to the fusion boundary as revealed by Dodge using TEM [5]. The carbon can diffuse a short distance in the scale of several microns in the PGZ during PWHT. Nanoindenation reveals that the sharp hardness peak coincides with the brownish feature in the narrow PGZ, probably due to precipitation hardening resulting from the cluster of M7C3 carbides. The carbides also increase the HAC susceptibility because the interface between the carbides and the matrix is the trap for hydrogen [4, 5]. The third microstructure of interest is the partially-mixed zone at the swirl structures or adjacent to the fusion boundary (Fig. 6). Previous researchers Fenske [4] and Dodge [5] reported two types of fusion boundary: continuous fusion boundary and discontinuous fusion boundary. The discontinuous fusion boundary is the fusion boundary with partiallymixed zone and/or swirl (Fig. 4b), and the opposite is the continuous fusion boundary (Fig. 4c). Previous studies have indicated that the CGHAZ and the PGZ are the most likely microstructures susceptible to HAC or SSC. However, the partially-mixed zone with high hardness and fresh martensite formed during cooling after PWHT also needs to be considered. This microstructure results from both the physical mixture of alloying elements during the overlay process (increase in hardenability) and the cooling process after PWHT, rather than carbon migration. The partially-mixed zone is also a region with higher nickel, chromium, and molybdenum contents than the base metal. At some regions of the partially-mixed zone, their composition is similar to that of martensitic stainless steel and the Ac3 temperature is also lower than the PWHT temperature, meaning that this region is austenitic at the PWHT temperature. After cooling from the PWHT temperature to room temperature (even furnace cooling at a rate ~ 0.8 °C/min), the partially-mixed zone transformed to fresh martensite. The fresh martensite in the partially-mixed zone has high hardness and is potentially susceptible to HAC and SSC. In practice, air cooling is conducted on the overlay components after PWHT. Because the air cooling has a higher cooling rate than the furnace cooling in this study, the high hardness fresh martensite probably forms in the partially-mixed zone after PWHT in actual practice.

Author's personal copy Weld World Table 5 Comparison of peak values of converted and measured Vickers hardness

PWHT

HJP

Correlation relationship

600 °C/5 h

18,070

y = 103.86x − 154.92

620 °C/6 h 650 °C/10 h

18,555 19,383

y = 128.65x − 213.23 y = 43.391x + 117.71

670 °C/10 h

19,803

y = 61.071x + 12.765

660 °C/50 h 660 °C/100 h

20,154 20,526

y = 62.765x y = 77.2x − 66.577

4.2 Comparison of the hardness between 8630/625 overlay and F22/625 overlay Figure 8 does not show a range of HJP within which the CGHAZ hardness is tempered below or close to 250 VHN, and the PGZ hardness is in an acceptable range, e.g., around 300 VHN. It is difficult to determine the optimal HJP range where the HAC susceptibility is potentially low only according to Fig. 8. Instead, the HAC susceptibility is possibly generally high in the whole HJP range. Actual HAC testing, such as DHCT, is necessary to confirm or disprove this hypothesis. The hardness variation of the CGHAZ, PGZ, and WM of 8630/625 overlay with the HJP is quite different from that of F22/625 overlay [1]. The average hardness of CGHAZ of the as-welded 8630/625 is 384 VHN, only slightly higher than that of as-welded F22/625, 376 VHN [1]. The martensitic CGHAZ of 8630 steel has a less sensitive tempering response than the CGHAZ of F22 steel. The average hardness data of the CGHAZ and PGZ of both F22/625 and 8630/625 overlays at seven shared PWHT conditions are listed in Table 6 and also plotted in Fig. 11 [1]. In the range of HJP below 20,000, identical PWHT conditions reduce the CGHAZ hardness of the F22/625 overlay below that of the 8630/625 overlay. Above 20,000, the CGHAZ hardness of F22/625 is higher than that of 8630/625, as shown in Fig. 11. This occurs because the untempered CGHAZ is martensite and its hardness is mainly determined by carbon content. 8630 steel contains 0.32 wt% carbon as compared to 0.15 wt% for F22 steel. Above HJP = Table 6

Converted Vickers hardness (HV0.1)

Measured average Vickers hardness (HV0.1)

4.911

355.1

307

4.747 8.328

397.5 479.1

326 375

Average nanohardness peaks (GPa)

9.003

562.6

463

10.935 9.292

686.3 650.8

592 606

20,000, the strong PWHT (660 °C/50 h and 660 °C/100 h) causes more pronounced carbon diffusion to the weld metal and carbon depletion in the CGHAZ of 8630/625 overlay relative to that of F22/625. In these two PWHT conditions, the CGHAZ of the 8630/625 overlay transformed to a homogeneous ferrite phase as shown in Fig. 5e, f. In F22, higher content of carbides forming elements chromium and molybdenum reduces the carbon diffusion coefficient [1] and the structure is maintained as tempered martensite, even after severe tempering. The PGZ hardness of 8630/625 overlay is higher than that of F22/625 overlays, as shown in Fig. 11 [1]. This is consistent with the more pronounced carbon depletion in the CGHAZ 8630/625 discussed above. Thus, there is more carbon diffusion across the fusion boundary and pileup in the PGZ of 8630/625 overlay relative to the F22/625, because of reduced chromium and molybdenum and higher carbon content in the 8630 steel.

5 Conclusions 1. The fusion boundary of the 8630/625 overlay consists of continuous fusion boundary and discontinuous fusion boundary, the latter consisting of adjacent partiallymixed zone and/or swirl structures. The swirl structures appear periodically along the fusion boundary, and most of them contain a partially-mixed zone on their edge.

Comparison of Vickers hardness (VHN) variation with the PHWT between 8630/625 and F22/625 overlays [1]

Overlay

Region

As-welded

640 °C 2h HJP = 18,535

640 °C 6h 18,970

650 °C 10 h 19,383

660 °C 10 h 19,593

670 °C 10 h 19,803

660 °C 50 h 20,154

660 °C 100 h 20,526

F22/625 8630/625 F22/625 8630/625

CGHAZ CGHAZ PGZ PGZ

376 384 – –

284 348 280 282

266 330 274 318

275 300 284 375

246 287 307 421

219 280 349 463

220 180 362 592

210 175 439 606

Author's personal copy Weld World Fig. 11 Comparison of Vickers hardness variation of the CGHAZ and the PGZ with the PWHT of 8630/625 and F22/625 overlays [1]

2. The fusion boundary of the 8630/625 overlay exhibits a distinct planar growth region (PGZ), or featureless zone. The width of this zone varies periodically from approximately 10 to 50 μm along the fusion boundary as a result of the multi-bead overlay process. After PWHT, a Bbrownish^ feature appears in the PGZ due to carbide precipitation after PWHT with HJP > 19,300. 3. In the CGHAZ, martensite decomposes increasingly as the PWHT increases, resulting in a reduction in hardness without secondary hardening effect. At high HJP values (> 20,100), the CGHAZ becomes almost fully ferritic. 4. Within the HJP range 19,300–19,900, the average hardness of the CGHAZ decreases from 300 to 250 VHN, but the average hardness of the PGZ increases from 370 to 500 VHN. When HJP > 19,900, the average hardness of the CGHAZ is below 250 VHN, but the average hardness of the PGZ hardness is above 500 VHN.

due to the content of alloying elements chromium and molybdenum, and secondly due to carbon content. 8. The PWHT response of the 8630/625 dissimilar weld overlay can be used to determine optimum resistance to hydrogen-assisted cracking and sulfide stress cracking. Acknowledgements This work was supported by Cameron International (now Schlumberger) through the NSF I/UCRC, Manufacturing and Materials Joining Innovation Center (MA2JIC) at the Ohio State University. Thanks are due to Acute Technological Services for providing Alloy 625 filler metal and also producing the overlays on the 8630 forgings. In addition, special thanks are due to Mr. Dean Hannam and Mr. Nash Ubale from Schlumberger who arranged for the procurement of the 8630 steel and also for providing technical support throughout this investigation.

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5. The PWHT promoted the carbon diffusion from the CGHAZ into the PGZ and pile up in the PGZ. When HJP > 19,300, the carbide precipitation occurs in the PGZ as revealed through etching as Bbrownish^ features. The sharp and high nanohardness peaks were associated with these features. The nanohardness levels are higher than those determined using standard Vickers measurements. 6. Fresh martensite is present in the partially-mixed zone of samples with PWHT above 600 °C due to transformation to austenite at the PWHT temperature. This martensite was found to form even at very low cooling rates from the PWHT temperature and exhibited hardness levels from 330 to 450 VHN. 7. The difference of tempering behavior of microstructure and hardness between 8630/625 and F22/625 is mainly

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