LM-05K074 May 9, 2005

Auger Spectroscopy Results from Ductility Dip

Cracks Opened Under Ultra-High Vacuum

T Capobianco, M Hanson

This report was prepared as an account of work sponsored by the United States Government. Neither the United States, nor the United States Department of Energy, nor any of their employees, nor any of their contractors, subcontractors, or their employees, makes any warranty, express or implied, or assumes any legal liability or responsibility for the accuracy, completeness or usefulness of any information, apparatus, product or process disclosed, or represents that its use would not infringe privately owned rights.

NOTICE

Auger Spectroscopy Results from Ductility Dip Cracks Opened

Under Ultra-High Vacuum

TE Capobianco and ME Hanson Lockheed Martin, Schenectady, NY

Abstract Ni Cr Fe Filler Metal 52 is susceptible to a form of intergranular solid state cracking known as ductility dip cracking (DDC). In Ni Cr Fe alloys, DDC cracking is usually associated with a loss of tensile ductility in the temperature range of 1500° F to 1700° F. While this tensile ductility loss in austenitic materials has been identified in the literature as early as 1912 (1), and many potential mechanisms have been suggested, the metallurgical basis for this phenomenon is poorly understood. The purpose of the present study is to characterize the grain boundary composition and microstructure of pristine ductility dip cracks in Filler Metal 52 (FM 52). Specifically, this report describes a technique to locate, prepare, and expose DDC fracture surfaces in ultra-high vacuum (<5 x 10-10 torr). These fracture surfaces were then characterized using Auger Electron Spectroscopy (AES) and Scanning Electron Microscopy (SEM). Using this methodology, two ductility dip cracks were located in specimens extracted from a heavy section multipass weld via x-ray computed tomography (CT). Once the DDCs were located by CT, the specimens were prepared for in-situ fracture in the vacuum chamber of the Auger instrument. Note that the DDCs were entirely contained within the specimens and were not exposed to the atmosphere prior to the in-situ fracture. The crack faces showed heavy coverage of chromium rich carbide particles (Cr23C6) as well as titanium and sulfur enrichment. While elevated grain boundary sulfur levels are well known to be detrimental to grain boundary strength in nickel alloys, the Auger analysis suggests that the high sulfur concentrations observed on the DDC fracture surfaces may be the result of segregation after the DDC formed. The post-fracture segregation is likely driven by the high energy of the newly created fracture surface and the thermal energy from subsequent weld passes.

Introduction FM 52 is the welding alloy for Ni Cr Fe Alloy 690. This alloy is resistant to stress corrosion cracking but is susceptible to a form of solid state cracking that occurs in the temperature range of 1500° F to 1700° F. This temperature range is typically produced in weld deposits as subsequent beads are made during multipass welds. Because this cracking is associated with a loss of tensile

ductility (% reduction in area) at these temperatures, it is referred to as ductility dip cracking. A consequence of the susceptibility of FM 52 to ductility dip cracking is that DDC-resistant Filler Metal 82H is used in some welds where the greater corrosion resistance of FM 52 is preferred. As part of an effort to improve the DDC resistance of FM 52, Lockheed Martin is conducting research on the metallurgical basis for ductility dip cracking. Weld simulation tests based on either the Varestraint or Gleeble have been proposed to assess the sensitivity of nickel base filler metals to DDC (2, 3). However, until now there have been few, if any, studies of DDC fracture surfaces extracted from actual multipass weld joints.

Specimen Preparation Weld wire from a commercial heat of FM 52 was deposited in a narrow groove test weld. A section of this weld was cut from the weldment and small cylindrical specimens (Auger pins) were machined from this section (Figure 1). The pins were visually inspected for surface breaking ductility dip cracks. Pins with surface breaking cracks were chosen to develop a crack imaging technique using CT. Figure 1: Weld section from which Auger pins were machined (rows of small holes) The surface breaking crack in the first pin chosen, specimen A4, was successfully imaged and an indication of an embedded DDC was also observed in the center of the pin (Figure 2). Flaw location data determined from the CT scan were used to set the position of a circumferential groove machined into the pin (Figure 3).

Welding direction

Figure 2: CT images of DDC in specimen A4. Both surface breaking and embedded cracks are visible in image (c). Figure 3: Pin machining dimensions were based on the CT scan. Axial location of groove was determined relative to surface-breaking crack. Groove diameter was based on the smallest dimension between the crack and the pin outside diameter (Figure 2b) then subtracting 0.010 in. The pin was cleaned, mounted in a holder, and inserted into the vacuum space (Figure 4). The unbroken specimen was left in the chamber under vacuum for 12 hours; the circumferential groove was then sputtered with an Ar+-ion beam (for additional cleaning) and fractured (Figure 5). Figure 6 shows a scanning electron micrograph of the end of the fractured pin with the DDC fracture surface outlined.

Figure 4: Specimen in vacuum chamber before fracture. Figure 5: Specimen in vacuum chamber after fracture. Several distinct regions can be seen: (1) the DDC surface approximately centered in the pin, (2) the deformation (smearing) produced by the actuator that fractured the pin, (3) the overload fracture surrounding the exposed DDC, and (4) a raised section in the middle of the crack surface that may have been scraped when the pin was sheared. A higher magnification view of the DDC area is shown in Figure 7. Figure 6: SEM image of fractured pin with DDC surface

(3) Mechanical overload damage in shear direction

(2) Force applied here

Shear direction

(4) Possible damage to crack face during fracture

(1) DDC area

b

Dia. = 0.135 in.

a

Section shown in (b)

Section shown in (c)

c Surface breaking crack

Embedded crack

Figure 7: SEM image of vacuum-exposed DDC fracture surface (outlined areas). The dendritic morphology of the grain boundary is evident on the fracture surface.

Auger Spectroscopy Auger Electron Spectroscopy (AES) was performed on two vacuum-exposed DDC surfaces (specimens A4 and A5). Both crack surfaces showed extensive coverage of chromium carbides (Cr23C6), sulfur, and titanium. At a magnification of approximately 120x, the sulfur appears to be a nearly continuous layer over the majority of the crack face (Figure 8) with roughly corresponding titanium coverage at lower concentrations (Figure 9). At higher magnifications, secondary electron images of the crack surfaces showed extensive coverage by small particles (Figures 10 and 11). The chromium concentration map aligns with these particles (Figure 12). Spectra collected from these particles point to identification as Cr23C6 (e.g. area 1 in Figure 11). This is further reinforced by analytical electron microscopy results which indicate that Cr23C6 is the predominant carbide in FM 52 (4). The images in Figures 6 to 12 were produced on specimen A4; similar results were obtained on A5. The sulfur maps out with other elements on the DDC surfaces but none are known sulfide formers like manganese or magnesium. This indicates that sulfur probably exists in elemental form on the fracture surface. Additionally, areas of high sulfur concentration are also associated with elevated concentrations of titanium, but perfectly coincidental mapping does not exist. A higher magnification view of the sulfur map (Figure 12) shows that sulfur coverage is not uniform at the sub-micron scale. Areas of local high sulfur concentration are roughly coincident with the 0.5 to 1 µm Cr23C6 particles seen in Figure 12. Caution needs to be exercised during the interpretation of these results since the exact moment that sulfur appears on the fracture surface is not yet known. It is quite possible that sulfur may populate the DDC crack face after the crack has formed (sulfur “bloom”). The free surfaces of the DDC cracks are highly energetic since they have been formed internally and have not been exposed to adsorbing species

like water, oxygen, or organics as in the case of air-exposed surfaces. Figure 8: Sulfur map of vacuum-exposed DDC in specimen A4. The outlined areas in Figure 7 correspond to areas of high sulfur concentration (bright areas in map). Figure 9: Titanium map of vacuum-exposed DDC in specimen A4. Atomically pristine surfaces will adsorb impurities in order to lower their energy. Facile diffusion of impurities is likely in the range of temperatures between crack formation and cool down. The combination of favorable diffusion conditions and high-energy sites for the sulfur to populate tends to support sulfur bloom as the mechanism underlying the observed result. Figure 11 shows three areas that were selected for Auger spectral elemental analyses. All elements except He and H are detectable in the Auger electron kinetic energy ranges displayed for these areas. Empirical sensitivity factors applied to peak-to-peak intensities of each element’s unique Auger transition provide atomic concentrations within approximately ±20% of the values reported in Table 1. The Auger electron spectra (Figure 13) clearly show the presence of sulfur and elevated levels of Ti both on and off Cr23C6 particles. The only detectable metals were the major alloying elements of Ni, Cr, and Fe. P was detected in sample A4, but not in sample A5. Auger depth profiles were collected on specimen A5 at the points shown in Figure 14. The depth profiles are shown in Figure 15. The depth profiling technique is defined by

S

Ti

SEM

alternating cycles of Ar+-ion sputtering to remove a thin layer (5 to 10Å) of material, and characterization of the freshly exposed subsurface regions with Auger Electron Spectroscopy. Sulfur located between the carbide particles sputtered completely off during the first sputter cycle (10Å). A small sulfur signal persisted deeper into the carbide particle (point #1) and did not disappear entirely until a depth of approximately 150Å (these depths are not shown in Figure 15). The Ti enrichment was persistent in the sub-surface regions for points 1 and 3 (Cr23C6 particle and “Ti-rich” area) for the entire 125 Å, but the Ti enrichment is extremely thin in point #2 (off-carbide). Figure 10: Area of DDC fracture in specimen A4 showing Cr23C6 coverage as small particles. Smooth areas without apparent second phase particles (e.g. around void) were also found to have chromium carbides. These carbides did not have sufficient contrast to be visible in this image. Figure 11: Higher magnification view of Cr23C6 coverage showing areas used for spectral analysis. Specimen A5 was notable for small (5 to 10 µm) areas outside of the main DDC surface that showed DDC-like characteristics. These showed up initially as sulfur hot spots in the ductile overload region. Further investigation revealed voids with a DDC-like surface morphology, chromium carbides, high levels of sulfur, and titanium enrichment (Figure 16). This suggests that DDCs exhibit a range of sizes that spans at least two orders of magnitude. Re-examination of specimen A4 revealed a similar area which can be seen at the center-top of Figure 8.

Figure 12: Chromium and sulfur maps corresponding to SEM image in Figure 11. Maps show a general association between S and carbide particles. Figure 13: Auger spectra for areas 1, 2, and 3. Table 1: Atomic concentration (%)

Area 1 2 3 C 11.1 2.1 2.7 P 2.7 2.6 3.4 S 11.5 10.9 10.4 Ti 2.6 2.3 2.8 Cr 35.5 20.4 19.2 Fe 4.7 6.6 7.0 Ni 32.0 55.1 54.6

1 3

2

1 3

2 Cr

13

2 S

Figure 14: SEM images and Auger maps of sample A5 showing the analysis points used for depth profiles.

Figure 15: Auger depth profile for points shown in Figure 14.

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Figure 16: SEM image (top) of sulfur hot spot outside of the main fracture. Further investigation produced chromium and titanium maps characteristic of the main DDC fracture surface.

Conclusions Lockheed Martin developed a technique to identify ductility dip cracks embedded within weld metal specimens and open these cracks in ultra-high vacuum. Using this technique, the cracks are never exposed to atmospheric contamination prior to characterization. This technique has provided a unique view of the DDC fracture surface and enabled chemical analysis of a pristine fracture surface using Auger Electron Spectroscopy.

AES characterization of two DDC surfaces showed a significant amount of chromium carbide particles in a region of titanium enrichment overlaid with a thin (~ 10 angstrom) layer of sulfur. Sulfur concentrations increased in the vicinity of the carbides, but the sulfur distribution was not characteristic of particles such as sulfides. While it is recognized that sulfur segregates to grain boundaries in Ni-base alloys, it appears that sulfur also has a loose association with Cr23C6 at the grain boundary. The thickness of the sulfur layer on the fracture surface is indicative of a post-fracture bloom. Small pockets of sulfur concentration away from the main fracture surface also show high concentrations of Cr23C6 and Ti enrichment suggesting that ductility dip cracks exhibit a range of sizes that span at least two orders of magnitude. If the detected sulfur is the result of post-fracture bloom, the implication would be that sulfur is present in the near grain boundary region in lower concentrations than is measured on the fracture surface. This scenario suggests that while sulfur embrittles grain boundaries, it may be more of a contributing factor rather than a primary cause in the ductility dip cracking process. The presence of sulfur and Cr23C6 on UHV-exposed ductility dip crack faces provides a possible explanation at the atomic level for the cause of DDC in FM 52: (1) sulfur is present in the near grain boundary region as the result of solidification segregation and causes a loss of strength; (2) local stresses build up at the grain boundary from precipitation of Cr23C6 during cooling from the temperatures produced during subsequent welding passes; (3) global shrinkage stresses in the weld then provide the impetus for grain boundary decohesion when the total stress exceeds the grain to grain bonding strength or perhaps through grain boundary sliding.

References 1. G. D. Bengough, “A Study of the Properties of

Alloys at High Temperatures,” Institute of Metals, Vol. 7, 123-174 (1912).

2. J. M. Kikel, D. M. Parker, “Ductility Dip Cracking

Susceptibility of Filler Metal 52 and Alloy 690,” Proceedings of the 5th International Trends in Welding Research, pp. 757-762 (1998)

3. M. G. Collins, J. C. Lippold, “An Investigation of

Ductility Dip Cracking in Nickel-base Filler Materials – Part 1,” Welding Journal, vol. 82, no. 10, pp. 288s-295s, Oct 2003

4. M. J. Cola, D. F. Teter, “Optical and Analytical

Electron Microscopy of Ductility Dip Cracking in Ni-Base Filler Metal 52 – Initial Studies,” Proceedings of the 5th International Trends in Welding Research, pp. 781-786 (1998)

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