The stress corrosion cracking (SCC) resistance of cold deformed thermally treated (TT) Alloy 690 has been questioned in recent years. As a step towards understanding its relevancy for weld deformed Alloy 690 in operating plants, Alloy 690 base metal and heat affected zone (HAZ) microstructures of three mockup components have been studied. All mockups were manufactured using commercial heats and welding procedures in order to attain results relevant to the materials in the field. Thermodynamic calculations were performed to add confidence in phase identification as well as understanding of the evolution of the microstructure with temperature. Ti(C,N) banding was found in all materials. Bands with few large Ti(C,N) precipitates had negligible effect on the microstructure, whereas bands consisting of numerous small precipitates were associated with locally finer grains and coarser
Many components in pressurized water reactors (PWRs) today are made of Ni-base alloys, for example, control rod drive mechanism (CRDM) nozzles, steam generator (SG) tubing, and SG divider plates. These components were previously made of Alloy 600 since it had exhibited good corrosion resistance in many aggressive environments in combination with having a coefficient of thermal expansion close to that of steel. Unfortunately, Alloy 600 soon turned out to be susceptible to stress corrosion cracking (SCC) in primary water, having a case reported in field in 1971 after only two years in service [
Alloy 690 was developed as a remedy to the SCC in Alloy 600, essentially being a high Cr version of Alloy 600 containing nominally 30% Cr instead of around 15-16%. Since the SCC resistance of Alloy 600 was found to significantly increase after being thermally treated (TT) [
Earlier experiments on cold worked TT Alloy 690 were part of an exploration of the effects of plastic strain on SCC in primary water. While Alloy 690 is not used in an intentionally cold worked state, the experiments were instead motivated by the concern for the heat affected zone (HAZ) of welded Alloy 690, which experiences high levels of plastic straining due to weld shrinkage, and can be further increased by repair welds. Estimates based on electron backscattered diffraction (EBSD) measurements have shown that the effective plastic strain in the HAZ typically increases towards a maximum around the fusion line [
Although cold working TT Alloy 690 base metal can generate a material with similar hardness and dislocation density as in the HAZ, the microstructural changes caused by the thermal cycles of a welding process would be absent. It has been shown that cold working Alloy 690 in a solution annealed (SA) state, where the M23C6 carbides have been dissolved, is not nearly as detrimental to SCC resistance as the same cold work performed in the TT state, where the grain boundaries would instead be densely covered in carbides [
How the differences in strain and thermal history affect the microstructure is an important step in understanding to which extent results obtained using cold worked materials can be extrapolated to describe HAZ behavior. The microstructure in the cold worked state has been described in literature and often exhibits damage such as cracked particles, particle-matrix interface debonding, grain boundary cavitation, or strain localization at particles and grain boundaries [
Segments of three mockup components from different manufacturers were provided for this work. All three mockups have been produced using commercial heats, realistic designs, and welding procedure specifications approved for actual plant components. Any observations made in this work are therefore expected to be found in materials in operating plants as well. Figure
The mockups denoted by CRDM1 and CRDM2 represented reactor pressure vessel head CRDM nozzle penetrations, and DP represented the lower part of the divider plate that separates the hot and cold sides of the SG.
All Alloy 690 base metals were welded in the as-received state, that is, hot worked followed by a SA and TT. The SA aimed at dissolving the carbides and is typically performed at around 1050°C for up to an hour followed by a water quench. The purpose of the TT is to precipitate carbides at the grain boundaries, which is achieved by several hours at around 700°C. These heat treatments also serve to recover the material and to relieve elastic stresses. The chemical compositions of the Alloy 690 base metals which used the mockups are given in Table
Chemical compositions in mass-% of the Alloy 690 base metals used in the mockups.
Base metal | Ni | Cr | Fe | C | Si | Mn | P | S | N | Ti | Al |
---|---|---|---|---|---|---|---|---|---|---|---|
CRDM1 tube | Bal. | 29.5 | 10.0 | 0.020 | 0.28 | 0.31 | 0.007 | 0.001 | 0.040 | 0.35 | 0.18 |
CRDM2 tube | Bal. | 30.0 | 8.9 | 0.020 | 0.32 | 0.30 | 0.007 | 0.001 | 0.028 | 0.21 | 0.38 |
DP bar | Bal. | 29.0 | 9.1 | 0.018 | 0.30 | 0.26 | 0.003 | 0.001 | 0.023 | 0.33 | 0.39 |
CRDM1 was manufactured using a hot extruded TT Alloy 690 tube inserted straight through a SA 533 grade B steel plate that was clad in grade 308 L stainless steel on the simulated internal surface, representing the reactor pressure vessel head. The tube had an outer diameter of 100 mm and a wall thickness of 16 mm. The tube was attached to the steel plate by gas tungsten arc welding (GTAW) with Alloy 52 M using a J-groove weld. A schematic sketch of the J-groove weld is shown in Figure
Schematic sketches of (a) CRDM1 and (b) DP.
CRDM2 was manufactured in a similar manner as CRDM1, but by a different manufacturer, and using a TT Alloy 690 tube inserted at a 46.5° angle through an SA 508 grade 3 steel plate instead of straight through. The tube was also hot extruded and had similar dimensions as the CRDM1 tube. Furthermore, no stainless steel cladding was applied on the steel plate, and Alloy 52 was used to fill the J-groove weld of this mockup. The completed weld joint was inspected using PT and UT.
The DP mockup had a steel bar as a base, representing the bottom end of a SG, buttered with a roughly 2 cm thick layer of Alloy 52 M. A hot forged bar of TT Alloy 690 was used to represent a SG divider plate. The bar was 132 mm thick in the S-direction, 105 mm in the T-direction, and roughly 410 mm long. The TT Alloy 690 bar was oriented as shown in Figure
In this work, the longitudinal, circumferential, and radial directions have been denoted by L, C, and R, respectively, for the tube materials. For the TT Alloy 690 bar used in DP, the longitudinal, long-transverse, and short-transverse directions have been denoted by L, T, and S, respectively. In order to study the possibility of microstructural anisotropy in the base metal, metallographic specimens were prepared with planes normal to the three orthogonal directions mentioned above. For characterization of the unaffected base metal, material sampled far from the weld was used to ensure minimal influence from the weld process. Specimens for HAZ characterization were prepared using material sampled near the weld root, where the weld induced plastic strains typically are significant. All samples were cut out of the mockups with care using plenty of coolant and low cutting speed to avoid unwanted evolution of heat.
The sampled materials were mounted in conductive plastic resin, ground on SiC paper, and polished on a rotating cloth disc with diamond suspension. The samples were investigated both in the as-polished condition and after being electrolytically etched in a nital solution (5% nitric acid in ethanol) with 4 V applied anodically for 15–25 s. The samples were studied in a light optical microscope (LOM) with optional dark field filter and in a scanning electron microscope (SEM) equipped with a backscattered electron (BSE) detector and equipment for energy dispersive X-ray spectroscopy (EDX).
Apart from the matrix, M23C6 type Cr-rich carbides and TiN/Ti(C,N) precipitates were assumed to be the only phases of significance in the material. These were differentiated in LOM by their color since TiN/Ti(C,N) appear golden or orange, while the Cr-rich carbides appear grey. High enough magnification is important as both phases appear as dark spots at insufficient magnification. They were also differentiated in SEM by atomic number contrast. The average atomic number of Ti(C,N) is significantly lower than M23C6, which results in a much darker color in the SEM using a BSE detector. EDX was used as a supplement in SEM to confirm locally elevated Ti or Cr levels at the particles for additional confidence in the phase identification.
Uneven illumination caused by the light source, lens system, and camera, for example, vignetting and hue gradients across the image, can affect the perception of hues in the LOM micrographs. The illumination over the image was quantified by fitting a second-order polynomial surface to the pixel intensities for each of the red, green, and blue color channels, respectively, over the image. Since Alloy 690 has a single phase matrix, such a fit of an ideal image would result in a flat plane for each color channel. All three fits would also be of similar intensity since they represent the grey color of the matrix. For each color channel, the fitted surface was divided by its mean value over the image to obtain a normalized division filter, which was used to compensate for the uneven illumination. Note that this method does not introduce any artifacts to, nor remove any details from, the images.
Thermodynamic calculations were applied in the present work in order to investigate the equilibrium phase fractions and the composition of each individual phase, with focus on the M23C6 carbide and the Ti(C,N) nitride/carbonitride precipitates. The considered alloy compositions were the same as shown in Table
The thermodynamic calculations presented in this work were performed using the Thermo-Calc software package [
Calculations such as these are strongly limited by the quality of the database used, which is why the choice of database is critical. In the case of TCNI5, the database does not include Mn, P, and S. However, these elements are only present in small quantities in the investigated alloys and were assumed to have little or no effect on the stability of the phases of interest.
The TT Alloy 690 tube of CRDM1 was homogeneous and had an equiaxed grain structure with plenty of annealing twins. The grain boundaries were densely covered by a network of Cr-rich carbides. Intragranular carbides were observed but very rarely. It is commonly accepted that these Cr-rich carbides in Alloy 690 are of the M23C6 type [
Microstructure of the CRDM1 tube. Ti(C,N) precipitates, visible as orange or dark spots, were common and tended to gather in faint streaks along the L-direction.
The Ti(C,N) precipitates were often found concentrated in bands extending in the L-direction of the tube. These bands were nonetheless faint in this material and easily missed. Many of the precipitates also appeared elongated in shape along the L-direction. Virtually no free oxide inclusions were observed in the base metal, although oxide slags were frequently found at the core of the larger Ti(C,N) precipitates.
The microstructure of the TT Alloy 690 tube of CRDM2 was similar to that of the tube of CRDM1, having an equiaxed grain structure with plenty of annealing twins and with grain boundaries densely decorated by carbides. Large Ti(C,N) precipitates of comparable size as those in CRDM1 were also present, but fewer in number. The differences between these tubes were mainly related to particle banding, which was significantly more distinct in this tube than in CRDM1. In particular, the bands contained plenty of small intragranular particles, up to around 1
Bands of small TiN precipitates extend along the L-direction in the CRDM2 tube.
The small intragranular particles in the bands were also believed to be Ti(C,N). They also had an orange color in LOM, and EDX confirmed clearly elevated Ti levels at these particles. Table
EDX analyses of matrix and second phase particles in mass-%. The scans of the particles were strongly affected by the matrix, but the elevated Cr and Ti levels for respective phase were still very clear.
Phase | Ni | Cr | Fe | Ti | Other |
---|---|---|---|---|---|
Matrix | 58.10 | 30.09 | 8.57 | ND | Balance |
M23C6 | 13.08 |
|
2.88 | ND | Balance |
Ti(C,N) | 24.85 | 18.75 | 4.28 |
|
Balance |
ND = not detected.
The grain size tended to be smaller within the Ti(C,N) banded regions. This resulted in a bimodal grain size distribution in the CRDM2 tube, although the average grain size was comparable to that of the CRDM1 tube. Grain boundaries could be seen lining up along bands fairly commonly, forming relatively straight lines ranging up to around 1 mm; see Figure
Grain boundaries in the CRDM2 tube occasionally line up to form fairly straight lines within bands or along the edge of a band.
Observing the grain boundaries at high magnification in LOM revealed that the grain boundary carbides within the Ti(C,N) bands were coarse, which can be seen in Figure
Grain boundary carbides located within regions banded with small Ti(C,N) in the CRDM2 tube were noticeably coarser than carbides located in the regions free of fine Ti(C,N) particles.
The microstructure of the DP bar is shown in Figure
The TT Alloy 690 DP bar had coarse grains and showed carbide coarsening associated with Ti(C,N) banding just like the CRDM2 tube. The magnification shows that the intragranular particle banding consisted of mainly orange Ti(C,N) precipitates and also of grey M23C6 carbides.
Unlike the CRDM2 tube, coarse M23C6 carbides of a significant amount were also present intragranularly in the banded regions. High magnification is required to differentiate M23C6 and Ti(C,N) in LOM. If the resolution is insufficient, both types of particles appear as dark spots which can lead to erroneous phase identification.
The HAZ of the Alloy 690 base metals from all three mockups showed a trend of M23C6 carbide dissolution close to the fusion line. Approaching the fusion line from the unaffected base metal, the carbide distribution changed from the initial semicontinuous network of fine M23C6 carbides into discrete and slightly coarser particles. Further approaching the fusion line, the M23C6 carbides were absent in LOM. Since the nital etchant is less effective in carbide-free grain boundaries in Alloy 690 [
These stitched LOM images show the HAZ of the CRDM1. The M23C6 carbides closest to the fusion line, marked by the dashed line, have been dissolved.
In the region where the grain boundary carbides appeared dissolved in LOM, only the occasional intergranular carbides could be seen in SEM, but only along some of the grain boundaries, and generally of very fine size. An example from the CRDM2 tube is shown in Figure
(a) Intergranular carbides were rare in the carbide-dissolution region. They were not observed on all grain boundaries and were very small. This can be compared with (b) the unaffected base metal.
The microstructure in the carbide-free region would have been similar to that of SA Alloy 690 during the weld induced straining. Since the plastic strains typically are the highest near the fusion line, it is possible that cold working Alloy 690 in the SA state yields a better representation of weld deformed HAZ than cold working it in the TT state. This may explain why SCC crack growth experiments on HAZ specimens have shown crack growth rates closer to specimens of Alloy 690 cold worked in the SA state than those cold worked in the TT state.
The previously mentioned coarser carbides in the bands of CRDM2 and DP appear to have been less affected by the weld. Compared to the carbides outside the bands, the coarser carbides within the bands appear to stay undissolved closer to the fusion line, as can be seen in Figure
The HAZ of CRDM2 with the fusion line marked by a dashed line. Coarser carbides seem to have been less affected by the weld than finer carbides. Ti(C,N) precipitates appear unaffected.
Unlike the M23C6 carbides, the Ti(C,N) bands of CRDM2 and DP were not noticeably affected by the thermal cycles of the welding process, as can be seen in Figure
The thermodynamic calculations verify that the intragranular precipitates, usually referred to as TiN in the literature, in fact correspond to a carbonitride while in equilibrium with the matrix phase in solid state. The composition of the phase Ti(C,N) at 1000°C based on the alloy composition of the CRDM1 tube is shown in Table
Calculated chemical composition of the Ti(C,N) phase at 1000°C for the CRDM1 overall material composition.
Element (mass-%) | Ti | N | C | Cr |
---|---|---|---|---|
Ti(C,N) | Bal. | 16.5 | 5.2 | 0.6 |
However, the C/N ratio varied strongly with temperature, as shown in Figure
The N (solid line) and C (dashed line) content in Ti(C,N) given in mole fraction as a function of temperature. The temperature for incipient melting of the matrix phase is marked by the number 1 and for full melting of the matrix by the number 2. The two vertical lines indicate the melting interval of the matrix phase.
Mole fraction of Ti(C,N) at equilibrium as a function of temperature for the three Alloy 690 compositions. The incipient melting of the matrix phase is indicated by the number 1. In all cases, the Ti(C,N) phase is stable up to and above the temperature of complete melting of the matrix, which is indicated by the number 2.
The main differences between the mockup alloys were the total phase fractions of Ti(C,N) and the temperature at which Ti(C,N) was fully dissolved in the liquid. This means that a complete solid solution of all elements into the matrix is not possible for any of the investigated alloys, if the calculations are to be trusted. Thus any heat treatment performed at solid state temperatures, or at insufficient superheating above the melting temperature of the matrix, will leave a certain fraction of undissolved TiN/Ti(C,N) in the microstructure. Figure
The temperature for full dissolution of the Cr-rich M23C6 phase was calculated to be 781°C for CRDM1. Only Fe and Ni were, to a small extent (~1-2 mass-%), dissolved in the carbide and the composition of this phase was more or less unaffected by the temperature. For the CRDM2 alloy, the dissolution temperature was calculated to be 983°C and for the DP alloy 889°C. Furthermore, the M23C6 carbide in CRDM2 also dissolved up to 4 mass-% Ni and 2.4 mass-% Fe and up to 9 mass-% Ni and 3.5 mass-% Fe in the DP alloy at higher temperatures.
A Cr-rich (95-96 mass-%) phase with body centered cubic (BCC) crystal structure was calculated to be stable below ~770–780°C for all three alloy compositions but was not observed in the samples. However, the M23C6 carbide was also stable in this temperature range and could be identified in the samples presented in Section
Particle banding was observed in the TT Alloy 690 base metal of all three mockups investigated in this work. In samples from CRDM2 and DP, the bands consisted mainly of fine Ti(C,N) precipitates. In these samples the Ti(C,N) bands were associated with locally finer grain sizes and coarser intergranular carbides. Furthermore, the grain boundaries tended to line up along the bands, which reduced the intergranular path tortuosity in the materials’ respective longitudinal direction. The TT Alloy 690 samples from CRDM1 had the highest volume fraction of the Ti(C,N) phase but these particles were distributed as fewer large precipitates instead of many fine ones. In this material, the particle banding had negligible effects on the base metal microstructure.
The thermodynamic calculations showed that the equilibrium composition of the Ti(C,N) phase corresponds to TiN at temperatures where the matrix is fully molten. According to the calculations, the Ti(C,N) in the HAZ would remain essentially unaffected by the heat from welding, which is verified by the observed microstructure near the fusion line. The calculations also indicated stability of a Cr-rich BCC phase below roughly 770–780°C, which so far has not been verified experimentally and could therefore be an artifact.
The main welding induced microstructural change was the full dissolution of intergranular M23C6 carbides near the fusion line. The width of the carbide-free region varied greatly between mockups, ranging roughly between 100 and 500
Since the weld induced plastic straining takes place during cooling after the carbide dissolution, cold worked SA Alloy 690 would be a better representation of this region than cold worked TT Alloy 690. This fits well with the SCC crack growth rates reported in the literature.
The authors declare that there is no conflict of interests regarding the publication of this paper.
This work has been performed under the financial support of the energy companies Vattenfall, Fortum, and Eon. The authors would also like to acknowledge EPRI and Ringhals AB for providing the mockups.