Microstructure and Stress Corrosion Cracking Behavior
of the Weld Metal in Alloy 52-A508 Dissimilar Welds
Wei-Chih Chung
1, Jiunn-Yuan Huang
2, Leu-Wen Tsay
3and Chun Chen
1;* 1Department of Materials Science and Engineering, National Taiwan University, Taipei 106, Taiwan, R. O. China
2Institute of Nuclear Energy Research, Lungtan 325, Taiwan, R. O. China
3Institute of Materials Engineering, National Taiwan Ocean University, Keelung 202, Taiwan, R. O. China
In the nuclear power industry, dissimilar metal welding is widely used for joining low alloy steel to austenite stainless steel components with nickel-base filler metals. In this study, attention was paid to the weld metal in multi-pass Alloy 52-A508 dissimilar welds. An approximately 2 mm wide transition zone was observed that consisted of a martensitic layer (1020um) along the weld interface and the austenite phase region with varying degrees of dilution. After post-weld heat treatment, the microstructures near the weld interface consisted of martensite, carbides and Type II boundaries. The presence of Type II boundaries significantly reduced the resistance to stress corrosion cracking (SCC) and formed intergranular cracking under simulated reactor coolant conditions. Constant extension rate tensile (CERT) tests were performed on the notched tensile specimens in 300C water at two extension rates,3104and1106mm/s. A fast CERT test can be
regarded to have no contribution of corrosion, and its results can be used as standards for comparison. In the slow CERT tests, the ductility losses of round-bar specimens with a circumferential notch at various regions in the weld metal were ranked accordingly. The relative susceptibility to SCC in terms of the ductility loss in increasing order of severity was as follows: the undiluted weld metal, the transition zone and the weld interface. SEM fractographic observations were consistent with the SCC results, i.e., an increased ductility loss or SCC susceptibility was associated with more brittle fractures. [doi:10.2320/matertrans.M2010294]
(Received August 31, 2010; Accepted October 12, 2010; Published November 24, 2010)
Keywords: A508 steel, Alloy 52 filler metal, dissimilar welds, microstructures, stress corrosion cracking, high-temperature water
1. Introduction
In nuclear power plants, dissimilar metal welding is widely used in joining low alloy steel components, e.g., reactor pressure vessels (RPVs) and steam generators, to the stainless steel piping.1,2)For such applications, Alloy 52 (a nickel-base
filler metal) is often applied to reduce the difference in thermal expansion coefficients between the ferritic and austenitic base metals.3,4)Additionally, the excellent
corro-sion resistance and the ability to retard carbon migration adjacent to the weld interface increase the usage of Alloy 52 filler metal in dissimilar welds,4,5)e.g., A508/Alloy 52/316L
material combinations. After welding, a post-weld heat treatment (PWHT) is required to lower residual stresses and temper the heat-affected zone (HAZ) of the welds. A transition zone is also formed within the weld metal, that exhibits complex microstructures and significant composi-tion gradients due to the effect of dilucomposi-tion,3,6,7)particularly at near the weld interface of the Alloy 52-A508 welds. As a result, a change in the susceptibility to stress corrosion cracking (SCC) can be expected in the transition zone of such welds.
It has been reported that the environment-assisted cracking (EAC) is one of the most common failure mechanisms of RPVs after long-term services.8)In some cases, the EAC in
boiling/pressurized water reactors (BWR/PWR) occurs in low-alloy steel piping components, such as nozzle safe ends, core shroud support welds and bottom head penetration housings, but rarely in the RPV itself.9)Consequently, crack-related problems, such as SCC and corrosion fatigue of dissimilar welds in high-temperature water, have become a major concern in the nuclear industry. It has also been
reported that interface cracking is often associated with a hardened interface region10,11)in the weld, implying that the weld interface plays an important role in determining SCC susceptibility. Although many investigations have conducted research on the SCC of dissimilar welds,11,12)little attention
has been paid to assess the SCC behavior at the weld interface and transition regions in the weld metal.
The aim of this study was to investigate the microstructure and SCC susceptibility of Alloy 52-A508 welds, with an emphasis on the fusion zone (weld metal), which includes the weld interface, transition zone and undiluted weld metal regions. To evaluate the SCC susceptibility in various regions of the welds, notched round-bar specimens were used to perform constant extension rate tensile (CERT) tests at two different rates under simulated coolant conditions in a nuclear reactor. For a given specimen, the result of a fast CERT test represented little or no environmental contribution to SCC, while that of a slow CERT test showed the effect of corrosion. The percentage loss in notched tensile strength (NTS) or ductility of various specimens in high-temperature water due to SCC could be calculated accordingly. The fracture surfaces of various specimens after CERT tests were also examined, particularly in the regions showing brittle fractures. Furthermore, the correlations among micro-structure, fracture characteristics and SCC susceptibility at various locations in the weld metal were highlighted.
2. Materials and Experimental Procedures
2.1 Materials and welding procedures
The materials used in this experiment were A508 Class 2 forged steel and Alloy 52 filler wire that was 2.4 mm in diameter. Table 1 lists the chemical composition of these materials. The steel was cut and solution-treated at 900C
*Corresponding author, E-mail: gchen@ntu.edu.tw
for 6 h and then water quenched, followed by tempering at 680C for 4 h. Figure 1 shows the schematic diagram of
the joint design and welding sequence for the welds. A508 specimens were buttered with Alloy 52 prior to filling a single 55 bevel joint using the multi-pass gas tungsten arc welding (GTAW) method. The detailed welding variables and procedures used to manufacture Alloy 52-A508 welds are listed in Table 2. It should be noted that all weld passes were processed with similar welding parameters, regardless of whether a higher current could be used to fill the butt joints. PWHT at 621C for 24 h was also carried out to
relieve the residual stress and soften the HAZ of all welds. Additionally, X-ray inspection was conducted to insure that there were no defects in the welded specimens prior to the SCC tests.
2.2 Metallographic and fractographic examinations
Metallographic examinations and microhardness measure-ments were performed on several dissimilar welds. A Vicker’s micro-hardness tester was used to determine the hardness in the vicinity of the weld interface regions using a
load of 50 g and a dwell time of 15 s. For microstructural observations, the Alloy 52 weld metal was electro-etched in 10% H2C2O4 + 90% H2O solution, while the A508
base metal and HAZ were etched using a 4% nital solution. The etched specimens were then observed with a scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS). Moreover, the microstruc-tural features near the weld interface were verified by electron backscatter diffraction (EBSD), as well as EDS semi-quantitative analysis.
[image:2.595.321.523.74.167.2]2.3 Constant extension rate tensile tests in high-temper-ature water
Figure 2 shows the configuration of the notched round-bar specimens used in this study. To evaluate the cracking susceptibility at distinct regions in the weld metal, circum-ferential V-notches (notch radius 0.025 mm) were introduced 0, 0.5 and 5 mm away from the vertical weld interface to represent the weld interface, transition zone and undiluted weld metal specimens, respectively. The CERT tests were conducted in pure water at 300C with a pressure of 10 MPa
to simulate coolant conditions in a nuclear reactor. Figure 3 is a schematic diagram showing the high-pressure water test loop in the CERT tests.13)Filtered compressed air was
injected into the water tank to maintain a saturated oxygen level (about 7 ppm dissolved oxygen) in the water. The test conditions of the water environment are summarized in Table 3. The CERT tests at two different rates, i.e.,
[image:2.595.46.292.93.255.2]3104mm/s (fast) and 1106mm/s (slow), were conducted to study the influence of environment on the
Table 1 The chemical composition (mass%) of A508 steel base metal and Alloy 52 filler wire.
Material
A508 Alloy 52
Element
C 0.17 0.04
Ni 0.67 Bal.
Cr 0.37 29.05
Fe Bal. 9.47
P 0.004 0.003
S 0.015 0.001
Mn 0.92 0.27
Mo 0.66 0.05
Si 0.19 0.18
V 0.03 —
Nb + Ta — 0.08
A508 55° Alloy 52
(3 buttering layers)
A508
Alloy 52
3 3 15
Alloy 52 (3 buttering layers)
3
Unit: mm
[image:2.595.306.548.216.365.2]Fig. 1 Schematic diagram showing the multi-pass welds for the SCC tests. Table 2 Welding parameters used in the experiment.
Welding current 140 A
Voltage 12 V
Travel speed 128 mm/min
Shielding gas Ar 15 l/min
Preheat temperature 150C
Interpass temperature 150C
ø12.2 ø10
ø4
60°
ø8
20
44 3 14
78
M12×1 C1
Unit: mm
R6
Fig. 2 Dimensions of the notched tensile specimen for the CERT test.
Water Tank Chemical
Injection
Deionizer H2+N2
O2+N2
N2
Water Source
High Pressure
Pump
Deionizer Measuring Tank
Heat Exchanger
Preheater Cooler PressureBack
Regulator
Actuator Autoclave Measure the values
of pH, DO, DH, and conductivity
[image:2.595.51.285.302.496.2] [image:2.595.59.277.404.495.2]mechanical properties of various specimens. The cracking susceptibility in terms of the percentage loss in ductility (extension or elongation to failure) or NTS of a given specimen can be expressed as follows:
Ductility loss (%)
¼Extension (fast)Extension (slow)
Extension (fast) 100%
NTS loss (%)¼NTS (fast)NTS (slow)
NTS (fast) 100%
3. Results and Discussion
3.1 Microstructural and compositional analysis
The microstructure of the A508 base metal consisted primarily of tempered martensite and bainite (left half of Fig. 4). After depositing three buttering layers, the micro-structures in the HAZ of the A508 steel were refined and had a mixture of bainite and ferrite (right half of Fig. 4). In Alloy 52-A508 welds, a transition zone within the weld metal was observed due to compositional differences between the base and weld metals. Figure 5 shows the distribution of major elements in a cross-section of typical multi-pass welds, in which a transition zone with significant changes in compo-sition is about 2 mm wide. It was noted that the weld metal, including most of the transition region, consisted of an austenitic phase. However, a narrow layer approximately 10 to 20mmwide along the weld interface with a completely different microstructure was observed. EBSD phase identi-fication (Fig. 6) confirmed that such a layer was composed mainly of martensite after depositing a single layer of Alloy 52. The wider layer of martensite in the figure was due to a specimen that was cut at an inclined angle of about 70 degrees to the weld interface. Figure 7 illustrates the
hard-ness difference in various regions near the weld interface of a multi-pass weld, in which the hardness is highest in the martensitic layer and lowest in the weld metal.
The martensitic layer can also be predicted by the Schaeffler Diagram,14) which is Fe-rich with Ni and Cr
[image:3.595.323.527.73.225.2]contents of less than 10 and 5 mass%, respectively. The variation in composition leads to a change of the mar-tensite start temperature (Ms) and affects the width of the martensitic layer in the transition zone. The Ms temperature can be calculated according to the following formula in mass%:15)
Table 3 Test conditions of the specimens in pure water.
Temperature 300C
Pressure 10 MPa
Conductivity (inlet/outlet) 0.18/0.49mS/cm
Dissolved oxygen (inlet/outlet) 6.97/7.18 ppm
PH (inlet/outlet) 6.27/6.78
Autoclave exchange rate 1 h1
5
µ
m
Base metal HAZ
Fig. 4 SEM micrographs showing the A508 base metal and HAZ in the as-welded specimens.
-4 0 20 40 60 80 100
Distance, d/mm
Mass per
cent (%)
Transition zone
HAZ WM
Ni
Cr
Fe
Weld interface Undiluted WM
1 0 -1 -2 -3
Fig. 5 The composition profiles of major elements in the weld metal. The gray area is the transition zone.
20
µ
m
HAZ Weld interface
WM
Ferrite or Martensite Austenite
Fig. 6 Micrograph showing the EBSD phase identification of a thin layer of martensite along the weld interface.
10
µ
m
Martensitic layer Weld interface
HZA WM
188 318 226
[image:3.595.46.290.84.159.2] [image:3.595.70.268.129.325.2] [image:3.595.329.522.278.416.2] [image:3.595.332.522.468.603.2]MsðCÞ ¼540 ð497Cþ6:3Mnþ36:3Ni
þ10:8Crþ46:6MoÞ
Figure 8 shows the calculated Ms temperature according to the results of EDS analysis, in which the width of the martensitic layer was about 10 to 20mmwhen measured from the weld interface to a parallel line that intersects room temperature. The presence of martensite/austenite mixtures at the weld interface is often observed due to the non-uniform distribution of alloying elements within the transition zone. This is consistent with the metallographic and EBSD observations.
Figure 9 shows a cross-sectional view across the weld interface of an Alloy 52-A508 weld. It clearly reveals the presence of Type II boundaries adjacent to the martensitic layer, in addition to Type I boundaries. Type II boundaries are grain boundaries that run about parallel to the weld interface at a very short distance into the weld metal.3)In
contrast, Type I boundaries are normal boundaries roughly perpendicular to the weld interface as a result of epitaxial growth.3)Both Type I and Type II boundaries are austenitic
grain boundaries, but the later has generally lower Ni and Cr contents than the former due to the dilution effect as illustrated in Fig. 5. Obviously, the composition of Type I boundaries depends on the distance from the weld interface, whereas the composition of Type II boundaries remains more or less the same. It should be noted that Type II boundaries
are formed only in dissimilar welds when the base and weld metals have different crystal structures, e.g. ferritic (bcc) steels welded with austenitic (fcc) filler metals.16–18)Nelson
et al.also reported that cracking or disbonding of dissimilar welds in service has been associated with Type II bounda-ries.18) The exact mechanism of this type of failure is not
clear but it may be due to impurity segregation and the orientation of Type II boundaries normal to the direction of principal stress.3) Nevertheless, the presence of such boun-daries might affect the SCC performance of dissimilar welds in high-temperature water.
Carbon migration from the A508 base metal (lower Cr content) to the Alloy 52 weld metal (higher Cr content) also occurred during the PWHT. Figure 10 shows the micro-structures near the weld interface of the weld after PWHT at 621C for 24 h. Carbon migration resulted in the formation
of a carbon-denuded zone (30mmwide) with an entirely ferritic structure in the HAZ. Further, the migrated carbon atoms were concentrated at the weld interface due to a much lower diffusion of carbon in the nickel alloy,5)resulting in the
precipitation of carbides (upper right corner in Fig. 10) at the interface. As a result, the microstructures near the weld interface were quite complicated and consisted of carbides, martensite and Type II boundaries.
3.2 SCC behavior of the dissimilar welds
Figure 11 shows the CERT results of the welds with notches at various locations in the weld metal. The stress-extension curves of all specimens exhibited serrations at different degrees. This could be attributed to the effect of dynamic strain aging, i.e., the interaction between disloca-tions and interstitial/substitutional atoms during plastic deformation at high temperatures.19,20) The serrations were
less obvious in the curves for the weld interface specimen than in the undiluted weld metal and transition zone specimens. The mechanism leading to the reduced amplitude of serrations of the weld interface specimen is not understood and requires further study. Figure 11(a) demonstrates that the undiluted weld metal (Alloy 52) had an excellent resistance to SCC; the values of NTS and extension to failure were independent of the extension rates of the CERT tests. The ductility and NTS losses in high-temperature water were 0.3 and 2.9%, respectively. For the transition zone specimen -0.10
-100 0 100 200 300 400 500 600
Ms temperatur
e,
T
/
°
C
Distance, d/mm Room temperature
Austenite Martensite
Weld interface
HAZ WM
0.02 0.00 -0.02 -0.04 -0.06 -0.08
Fig. 8 Calculated Ms temperature across the weld interface of the Alloy 52-A508 weld.
5
µ
m
Type I boundary Type II
boundary
WM
HAZ
Fig. 9 Micrograph showing Type II boundaries adjacent to the weld interface of an Alloy 52-A508 weld.
10
µ
m
Carbon-denuded zone
WM
[image:4.595.68.270.70.224.2]2 µm
Fig. 10 Micrograph showing the microstructures near the weld interface of the specimen after PWHT (621C/24 h). The carbon-denuded zone is
[image:4.595.333.521.73.208.2] [image:4.595.74.264.277.415.2](Fig. 11(b)), the NTS loss (4.1%) increased slightly, but the ductility loss (13.5%) was significantly increased relative to the undiluted weld metal. The transition zone specimen is essentially an austenitic stainless steel that is expected to have less SCC resistance than the nickel-base alloy, i.e., the undiluted weld metal or Alloy 52. In fact, the SCC suscepti-bility in the transition zone varied continuously with distance from the weld interface due to the dilution effect. In other words, the region with higher Ni and Cr contents would have more resistance to SCC. The least resistance to SCC was associated with the weld interface specimen (Fig. 11(c)), which had ductility and NTS losses of 20.6 and 4.3%, respectively. The above results suggested that the ductility loss is a better index with which to rank the relative susceptibility to SCC of weld metal specimens. In contrast, NTS loss is not an appropriate index to reflect the relative susceptibility to SCC among the specimens in this study.
A circumferential notch in a round-bar tensile specimen radically changes the stress distribution, causing the crack to grow inward. During the slow extension rate tests in high-temperature water, the area near the root of the notched tensile specimen would suffer most from SCC, and a change in fracture modes would be expected at the peripheral region of the fracture surface. Figure 12 shows typical fractographs of the weld interface specimens after the CERT tests at extension rates of 1106mm/s (slow rate) and 3
104mm/s (fast rate). The specimen tested at a slow rate
generally exhibited more brittle features on the fracture surface than that at a fast rate, as shown in Figs. 12(a) and 12(b). It should be mentioned that Type II boundaries are
very close to the weld interface and can be considered to be associated with the weld interface instead of the transition zone in the classification of the specimens. Figure 12(c) also reveals the intergranular cracking associated with Type II boundaries and the secondary cracking along Type I boun-daries after the SCC tests. The composition of Type II boundaries, as determined by EDS analysis, was about Fe-20 mass% Ni-12 mass% Cr, i.e., essentially an austenitic stainless steel. On the other hand, evidence of SCC forming a feather-like transgranular fracture propagation in the mar-tensitic layer of a weld interface specimen tested at a slow rate is shown in Fig. 12(d). Such a layer was confirmed by EDS analysis to have a composition (Fe-4.8 mass% Ni-2.0 mass% Cr) similar to a martensitic stainless steel. Conversely, the notched tensile fracture at the martensitic layer was not observed in the similar specimen tested at a fast rate, possibly due to the higher strength of martensite relative to its surroundings and the absence of a corrosion contribution.
In the case of the transition zone specimen tested at a slow rate, the fracture surface (Fig. 13(a)) exhibited more flat trangraunlar regions than that tested at a fast rate. Deforma-tion markings (left half of Fig. 13(b)), i.e., slip bands, on these areas could be clearly observed in the transition zone specimen tested at different rates. In addition, intergranular fractures (right half of Fig. 13(b)) were observed for the specimen tested at a slow rate, but were not found in the specimen tested at a fast rate. Under the influence of SCC, intergranular fractures could be related to the nearby Type II boundaries, due to the tendency of cracks to grow towards
0.0 0 200 400 600 800 1000 1200 1400 1600
Notched tensile str
ength,
σ
/MP
a
Extension, ∆l/mm
Fast Slow
(c)
890 MPa 930 MPa
0.85 mm 1.07 mm
NTS loss = 4.3% Ductility loss = 20.6%
0.0 0 200 400 600 800 1000 1200 1400 1600
Notched tensile str
ength,
σ
/MP
a
Extension, ∆l/mm
Fast Slow
(b)
879 MPa
917 MPa
0.96 mm 1.11 mm
NTS loss = 4.1% Ductility loss = 13.5%
0.0 0 200 400 600 800 1000 1200 1400 1600
Notched tensile str
ength,
σ
/MP
a
Extension, ∆l/mm
Fast Slow
(a)
904 MPa 931 MPa 1.43 mm 1.43 mmNTS loss = 2.9% Ductility loss = 0.3%
1.6 1.4 1.2 1.0 0.8 0.6 0.4
0.2 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6
[image:5.595.101.496.73.365.2]1.6 1.4 1.2 1.0 0.8 0.6 0.4 0.2
the weakest regions in the specimen. The amount of intergranular fractures in the transition zone specimen was considerably less than that of the weld interface specimen as a result of the notch was about 0.5 mm away from the weld
interface. For the undiluted weld metal specimen, necking was quite extensive in both the slow and fast extension rate tests. As a result, a circular cross-section of the macro-fracture appearance could not be obtained due to severe
1 mm
(a)
IG
IG
(c)
50
µ
m
Type II boundary
10
µ
m
Weld Interface Type I
boundary
Crack propagation
20
µ
m
(d)
(b)
1 mm
(d)
IGIG
IG
IG
(c)
IG IG
M IG
IG
[image:6.595.103.494.71.412.2]IG M
Fig. 12 Fractographs of the weld interface specimens after the CERT tests: (a) macro-fracture view of the specimen tested at a fast rate; (b) macro-fracture view of the specimen tested at a slow rate; (c) intergranular (marked IG in (b)) fracture associated with Type II boundaries; and (d) transgranular SCC fracture in the martensitic layer (marked M in (b)). Note that the unmarked areas on the macro-fracture surfaces were mainly composed of dimples.
5
µ
m
20
µ
m
(b)
FT
FT
IG
1 mm
(a)
FT
FT
FT FT
[image:6.595.102.493.487.676.2]deformation of the ductile material (Fig. 14(a)). The fracture surface of the undiluted nickel-base weld metal exhibited mainly dimples (Fig. 14(b)) with scattered flat transgranular fracture areas (Fig. 14(a)) on the fracture surface. At high magnification, such areas also revealed extensive slip bands on the flat surface.
For austenitic alloys such as stainless steels or nickel-base alloys, the SCC mechanism in high-temperature water is based mainly on the slip-dissolution/film-rupture model.21,22) The repeated formation and destruction of oxide layers are thought to cause deteriorated SCC resistance of the material. As the passive oxide film continuously ruptures in the CERT tests, localized corrosion attack in front of a circumferential notch tip proceeds. Consequently, an electrochemical cell forms between the fresh metal surface and the surrounding oxide film, which results in metal dissolution and the generation of hydrogen.23)Although hydrogen embrittlement
was not considered to cause SCC in this study, it might be related to SCC in the martensitic regions in the weld interface specimen. After SCC tests in high-temperature water, all specimens contained oxide particles with various composi-tions and morphologies on the fracture surfaces, as shown in Fig. 15. It is apparent that the composition and micro-structure of the specimens were different, which reflected in the SCC performance and oxide formation. Oxides on the fracture surface of the undiluted weld metal specimen were
covered with fine chromium oxide particles (Fig. 15(a)), whereas those of the transition zone specimen contained somewhat larger Ni-Cr-Fe oxides (Fig. 15(b)) in the auste-nitic stainless steel region, as determined by SEM/EDS analyses. In the case of the weld interface specimen, the oxides on Type II boundaries and martensitic regions were different in composition and contained some large particles (Fig. 15(c)). Ni-Fe oxides were observed on the fracture surface of Type II boundaries, and Fe-Ni oxides were found to be associated with the martensitic layer. It should be noted that the fracture surface might become obscure if oxide particles are formed extensively. SEM fractographic obser-vations were in good agreement with the ductility loss of the specimens after the CERT tests, i.e., the more brittle fracture is generally associated with more ductility loss or a higher SCC susceptibility. The SCC results of Alloy 52-A508 welds in high-temperature water indicated that the weld interface was the most susceptible region in the weld metal, and should be mentioned in the field applications.
1 mm
(a)
FT
FT
FT
(b)
10
µ
m
[image:7.595.74.263.70.409.2]1 µm
Fig. 14 Fractographs of the specimen with a notch located in the undiluted weld metal after a slow extension tensile test: (a) macro-fracture view; and (b) ductile fracture containing different dimple sizes.
2
µ
m
(a)
0.5
µ
m
Undiluted weld metal
2
µ
m
(b)
Transition zone2
µ
m
(c)
Type IIboundary
Martensitic layer
[image:7.595.330.520.71.497.2]2
µ
m
4. Conclusions
(1) The notched tensile tests of Alloy 52-A508 dissimilar welds at a slow rate were used to differentiate the SCC susceptibility of various regions in the weld metal under a simulated reactor coolant environment. The ductility loss seemed to be a better index than the NTS loss in ranking the relative SCC susceptibility among the specimens with comparable strength levels.
(2) The microstructures near the weld interface were complicated, consisting of martensite, carbides and Types II boundaries. Apparently, the presence of Type II boundaries caused intergranular cracking and significantly reduced the SCC resistance of the weld in 300C water. Additionally, the structural discontinuity at the interface also increased the SCC susceptibility of the weld interface specimen.
(3) The alloying elements in the transition zone varied with distance from the weld interface, as a result, the SCC susceptibility changed accordingly. The more Ni and Cr contents, the less the specimen would be susceptible to SCC. For instance, the transition zone specimen exhibited a much higher ductility loss (13.5%) than the undiluted nickel-base weld metal (0.3%) in high-temperature water.
(4) The relative susceptibility to SCC at various locations in the weld metal followed an increasing order of severity: the undiluted weld metal, the transition zone and the weld interface. SEM fractographic observations were in agreement with the SCC results, i.e., a higher susceptibility or ductility loss was associated with more brittle features on the fracture surface of the specimen.
Acknowledgments
The authors gratefully acknowledge financial support of this investigation by the ROC National Science Council (Contract Numbers NSC99-NU-E-002-001 and NSC98-NU-E-002-001) and technical assistance from the Institute of Nuclear Energy Research.
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