Handles welded to the top cover plate of a chemical-plant downcomer broke at the welds when the handles were used to lift the cover. The handles were fabricated of low-carbon steel rod; the cover was of type 502 stainless steel plate. The attachment welds were made with type 347 stainless steel filler metal to form a fillet between the handle and the cover. The structure was found to contain a zone of brittle martensite in the portion of the weld adjacent to the low-carbon steel handle; fracture had occurred in this zone. The brittle martensite layer in the weld was the result of using too large a welding rod and too much heat input, melting of the low-carbon steel handle, which diluted the austenitic stainless steel filler metal and formed martensitic steel in the weld zone. Because it was impractical to preheat and postheat the type 502 stainless steel cover plate, the low-carbon steel handle was welded to low-carbon steel plate, using low-carbon steel electrodes. This plate was then welded to the type 502 stainless steel plate with type 310 stainless steel electrodes. This design produced a large weld section over which the load was distributed.

The welds joining the liner and shell of a fluid catalytic cracking unit failed. The shell was made of ASTM A515 carbon steel welded with E7018 filler metal. The liner was made of type 405 stainless steel and was plug welded to the shell using ER309 and ER310 stainless steel filler metal. Fine cracks starting inside the weld zone and spreading outward through the weld and toward the surface were observed during examination. Decarburization and graphitization of the carbon steel at the interface was noted. The high carbon level was found to allow martensite to form eventually. The structure was found to be austenitic in the area where the grain-boundary precipitates appeared heaviest. The composition of the precipitates was analyzed using an electron microprobe to reveal presence of sulfur. Microstructural changes in the weld alloy at the interface were interpreted to be caused by dilution of the alloy and the presence of sulfur caused hot shortness. The necessary internal stress to produce extensive cracking was produced by the differential thermal expansion of the carbon and stainless steels. Periodic careful gouging of the affected areas followed by repair welding was recommended.

Grease-wiper springs for cams formed from stampings of 0.25-mm thick carbon spring steel (0.65 to 0.80% C) fractured at the 0.025 mm radius on the stamped 135 deg corner at a 90 deg bend after 5,000,000 cycles. Tool marks 2 to 2.3 mm from the center of the stamped bend were disclosed by visual examination. Fatigue striations originating from cracks at the 0.025 mm radius inside corner at the bend were revealed by SEM of the fractured surface. The maximum stress at the bend, in stock of maximum thickness and as a function of the radius of the 135 deg corner, was indicated by stress calculations to be very close to the maximum allowable fluctuating stress for the material. The wiper springs were concluded to be fractured in fatigue and the cyclic loading resulted from cam rotation. The maximum applied stress approached the allowable limit due to high stress-concentration factor in the spring (caused by the very small inside radius). The corner radius was increased to 0.76 mm and the tools were re-polished to avoid tool marks.

The connecting end of two forged medium-carbon steel rods used in an application in which they were subjected to severe low-frequency loading failed in service. The fractures extended completely through the connecting end. The surface hardness of the rods was found to be lower than specifications. The fractures were revealed to be in areas of the transition regions that had been rough ground to remove flash along the parting line. The presence of beach marks, indicating fatigue failure, was revealed by examination. The fracture origin was confirmed by the location and curvature of beach marks to be the rough ground surface. An incipient crack 9.5 mm along with several other cracks on one of the fractured rods was revealed by liquid penetration examination. Metallographic examination of the fractured rods indicated a banded structure consisting of zones of ferrite and pearlite. It was established that the incipient cracks found in liquid-penetrant inspection had originated at the surface in the banded region, in areas of ferrite where this constituent had been visibly deformed by grinding. Closer control on the microstructure, hardness of the forgings and smooth finish in critical area was recommended.

A plant had manufactured and heat treated their product in house for years. As time went on, the special steel that they had been using became more expensive, and a switch was made to a more common and less highly alloyed material. However, no change in hardness specifications were made, because calculations of ideal critical diameter and analysis of available hardenability data indicated that the original hardness specification could be met. There was, however, less room for process variation. The parts ended up containing temper carbides, developed heavy decarburization, and experienced excessive distortion because they were left in the furnace for extended and varying periods with the temperature “turned down a couple hundred degrees.”

An Inconel-clad SA-212 Grade B carbon steel inlet cone with an anticipated 25-year service life failed in a localized area after only seven years of service. The failure was caused by an erosion/corrosion leak at the midsection. Erosion/corrosion was confined to a localized area directly facing the steam inlet nozzle. The Inconel cladding was intact elsewhere in the inlet cone with insignificant corrosion-related degradation. In the absence of the conditions that led to erosion/corrosion, the Inconel clad carbon steel was considered adequate for the intended service. As a corrective measure, a solid Inconel liner was recommended in the areas of direct steam impingement.

Extra high strength zinc-coated 1080 steel welded wire was wound into seven-wire cable strands for use in aerial cables and guy wires. The wires and cable strands failed tensile, elongation, and wrap tests, with wires fracturing near welds at 2.5 to 3.5% elongation and through the welded joints in wrap tests. The welded wire was annealed by resistance heating. The wire ends had a chisel shape, produced by the use of sidecutters. Tests of the heat treatment temperatures showed that the wire near the weld area exceeded 775 deg C (1425 deg F). Metallographic examination revealed martensite present in the weld area after the heat treatment. The test failures of the AISI 1080 steel wire butt-welded joints were due to martensite produced in cooling from the welding operation that was not tempered adequately in postweld heat treatment, and to poor wire-end preparation for welding that produced poorly formed weld burrs. The postweld heat treatment was standardized on the 760 deg C (1400 deg F) transformation treatment. The chisel shape of the wire ends was abandoned in favor of flat filed ends. The wrap test was improved by adopting a hand-cranked device. Under these conditions, the welded joints withstood the tensile and wrap tests.

Two leaks were discovered at a sulfur recovery unit in a refinery. The leaks were at pipe-to-elbow welds in a 152 mm (6 in.) (NPS 6) diam line, operating in lean amine service at 50 deg C (120 deg F) and 2.9 MPa (425 psig). Thickness measurements indicated negligible loss of metal, and the leaks were clamped. A year later, 15 additional leaks were discovered, again at pipe-to-elbow welds in lean amine lines. Further nondestructive testing located other cracks, giving a total of 35. These lines had been in service for approximately eight years. Investigation (visual inspection, hardness testing, and micrographic cross-sections) supported the conclusion that the failure was caused by lean amine SCC. It was considered unlikely that these pipe welds had received such a postweld heat treatment, although it is industry practice to postweld stress relieve piping and pressure vessels in lean amine service if the temperature is expected to be above 95 deg C (200 deg F). Recommendations included inspecting all welds using shear wave ultrasonic testing and postweld heat treating all welds in lean amine service.

A large pressure vessel that had been in service as a hydrogen sulfide (H2S) absorber developed cracks and began leaking at a nozzle. The vessel contained a 20% aqueous solution of potassium hydroxide (KOH), potassium carbonate (K2CO3), and arsenic. The vessel wall was manufactured of ASTM A516, grade 70, low-carbon steel plate. A steel angle had been formed into a ring was continuously welded to the inside wall of the vessel. The groove formed by the junction of the lower tray-support weld and the top part of the weld around the nozzle was found to have a crack. Pits and scale near the crack origin were revealed by microscopic examination and cracking was found to be transgranular. Periods of corrosion alternated with sudden instances of cleavage, under a tensile load, along preferred slip planes were interpreted during examination with a microscope. It was concluded that the combination of the residual plus operating stresses and the amount of KOH present would have caused stress corrosion as a result of caustic embrittlement. It was recommended that the tray support should be installed higher on the vessel wall to prevent coincidence of the lower tray-support weld with the nozzle weld.

A wall section of a carbon steel choke body in gas service at 4400 psig blew out three months after the use of a corrosion inhibitor was stopped. Corrosion damage occurred in ripples, leaving both smoothly polished and unattacked areas. The corrodent in condensate wells was principally carbon dioxide dissolved in water condensed from the gas stream, with organic acids possibly an aggravating factor. A gas analysis showed no other corrosive agents. No metallurgical or fabrication defects were found in the carbon steel part. The mode of attack was corrosion-erosion, caused by the corrosive, high velocity gas flow. The corrosion rate of either the inhibited or uninhibited gas stream was too high for equipment in high pressure gas service. Type 410 (12% Cr) stainless steel was recommended for the choke bodies because other equipment such as valves made of type 410 showed no evidence of corrosion damage after three years' exposure. This change was made five years ago and there have been no failures since.

The power-type counterbalance spring, formed from hardened-and-tempered carbon steel strip and subsequently subjected to phosphating treatment, fractured at the two locations during fatigue testing. A rust colored dark band at the inside edge of the fracture surface was disclosed during investigation. Etch pits were revealed by the cleaned surface which were never observed on properly phosphated coating. It was interpreted that the spring had been subjected to an abnormal acid attack in pickling or phosphating which had resulted in considerable absorption of hydrogen by the metal and hence embrittlement. The part was concluded to have cracked during phosphating or excessive acid pickling before phosphating.

The center portions of two adjacent low-carbon steel boiler tubes (made to ASME SA-192 specifications) ruptured during a start-up period after seven months in service. It was indicated by reports that there had been sufficient water in the boiler two hours before start-up. The microstructure near the rupture edge was revealed by metallographic examination to consist of ferrite and acicular martensite or bainite. The microstructure and the observed lack of cold work indicated a temperature above the transformation temperature of 727 deg C had been reached. Swelling of the tubes was disclosed by the wall thickness and OD of the tubing. The tubes were concluded to have failed due to rapid overheating.

A resistance-welded carbon steel superheater tube made to ASME SA-276 specifications failed by pitting corrosion and subsequent perforation, which caused the tube to leak. The perforation was found to have occurred at a low point in a bend near the superheater outlet header. It was found that the low points of the superheater tubes could not be completely drained during idle periods. Water-level marks were noticed on the inside surface above the area of pitting. It was revealed by microscopic examination that localized pitting had resulted from oxidation. It was concluded that water contained in the tube during shutdowns had accumulated and cumulative damage due to oxygen pitting resulted in perforation of one of the tubes. Filling the system with condensate or with treated boiler water was suggested as a corrective action. Alkalinity was suggested to be maintained at a pH of 9.0 and 200 ppm of sodium sulfite should be added to the water.

Carbon steel tubes from a boiler feedwater heater feeding a deaerator were treated to control scale formation, but the treatment instead produced more iron oxide. The additional iron oxide reduced the tubing to a totally corroded condition. Investigation showed that the chelate injected to control the scaling was added ahead of the preheater, where the boiler water still contained oxygen. As the chelate removed iron oxide, the O2 in the water continued to form more. Recommendations included moving the chelate addition to a point after the deaerator to stop the corrosion.

A carbon steel furnace tube which should have given good service for ten years ruptured after one year. The tube showed obvious swelling at the point of rupture, and the bulged surface of the tube was oxidized at a temperature far above the design temperature. There was little or no loss in wall thickness due to corrosion or scaling, and the tube wall was thinned to a knife edge at the rupture. Metallographic examination showed the condition of the material was satisfactory. The failure was mechanical in nature, typical of short time creep rupture. The localized oxidation indicated improper furnace operation or blockage of the tube. The furnace was checked and found to have a burner tip out of order. After the tip was repaired, localized overheating was minimized and further premature failures did not occur.

Welded steel storage vessels used to hold mildly alkaline solution were produced in exactly the same manner from deep-drawn aluminum-killed SAE 1006 low-carbon steel sheet. After the cylindrical shell was drawn, a top low-carbon steel closure was welded to the inside diameter. The containers were then filled with the slightly alkaline solution, pressurized, and allowed to stand under ambient conditions. A small number, less than 1%, were returned because they began to leak in service. Inspection revealed general corrosion and pitting on the inner surfaces. However, other tanks that experienced the same service conditions developed no corrosion. Corrosion was linked to forming defects that provided sites for localized corrosion, and to lack of steam drying after cleaning, which increased susceptibility to general corrosion.

A root cause failure analysis was performed on a vaporizer coil removed from a horizontal forced circulation vaporizer. The carbon steel coil was wound in a right-hand helix with a coil centerline diameter of about 2 m. The vaporizer was gas fired and used Dowtherm A as the heat transfer fluid. Design conditions are based on annular fluid flow to cool the coil wall. NDE, metallographic and fractographic examinations were performed. Numerous, circumferentially oriented, OD initiating cracks were found near the crown for two coils near the non-fired end of the vaporizer. The cracking was confined to the inner diameter of the vaporizer coil at positions from 4:00 to 7:00. The cracking was characterized as transgranular and the fracture surface had beach marks. The failure mechanism was thermal fatigue. The heat transfer calculation predicted that dryout of the coil would occur for coils at the non-fired end of the vaporizer during low flow transients. Dryout results in rapid increase in the tube wall temperature. Thermal cycling of the coil is completed by liquid quenching resulting from resumption of normal flow rates and the return to annular flow. The probable root cause of failure was low flow transient operation.

The pawl spring which was part of a selector switch used in telephone equipment failed. The springs were blanked from 0.4 mm (0.014 in.) thick tempered 1095 steel and then nickel plated. Numerous pits around the rivet holes were revealed by microscopic examination of longitudinal specimens. Delaminations that were formed at inclusion sites during punching of the rivet holes and that were filled with nickel during the plating operation were revealed by microscopic examination of the rivet hole. These delaminations were interpreted to have acted as stress raisers and initiated the fracture. Long, narrow sulfide stringers which were the probably the cause of delamination in this spring material were revealed in the raw material used to make the springs. It was concluded that fracture of the springs was caused by fatigue that had originated at delaminations around the rivet holes.

Welded low-carbon steel bomb fins were rejected because of poor weld practice. Visual and metallographic examination revealed that the resistance plug welds that attach the outer skin to the inner spar displayed inadequate weld penetration. Recommended changes to the resistance welding parameters resulted in acceptable welds.

A steel galvanizing vat measuring 3 x 1.2 x 1.2 m (10 x 4 x 4 ft) and made of 19 mm thick carbon steel plate (ASTM A285, grade B)) at a shipbuilding and ship-repair facility failed after only three months of service. To verify suspected failure cause, two T joints were made in 12.5 mm thick ASTM A285, grade B, steel plate. One joint was welded using the semiautomatic submerged arc process with one pass on each side. A second joint was welded manually by the shielded metal arc process using E6010 welding rod and four passes on each side. The silicon content of the shielded metal arc weld was 0.54%, whereas that of the submerged arc weld was 0.86%. After being weighed, the specimens were submerged in molten zinc for 850 h. Analysis (visual inspection, chemical analysis, 100x 2% nital-etched micrographs) supported the conclusions that the vat failed due to molten-zinc corrosion along elongated ferrite bands, possibly because silicon was dissolved in the ferrite and thus made it more susceptible to attack by the molten zinc. Recommendations included rewelding the vat using the manual shielded metal arc process with at least four passes on each side.