A group of control valves that regulate production in a field of sour gas wellheads performed satisfactorily for three years before pits and cracks were detected during an inspection. One of the valves was examined using chemical and microstructural analysis to determine the cause of failure and provide preventive measures. The valve body was made of A216-WCC cast carbon steel. Its inner surface was covered with cracks stemming from surface pits. Investigators concluded that the failure was caused by a combination of hydrogen-induced corrosion cracking and sulfide stress-corrosion cracking. Based on test data and cost, A217-WC9 cast Cr–Mo steel would be a better alloy for the application.

Nineteen out of 26 bolts in a multistage water pump corroded and cracked after a short time in a severe working environment containing saline water, CO 2 , and H 2 S. The failed bolts and intact nuts were to be made from a special type of stainless steel as per ASTM A 193 B8S and A 194. However, the investigation (which included visual, macroscopic, metallographic, SEM, and chemical analysis) showed that austenitic stainless steel and a nickel-base alloy were used instead. The unspecified materials are more prone to corrosion, particularly galvanic corrosion, which proved to be the primary failure mechanism in the areas of the bolts directly exposed to the working environment. Corrosion damage on surfaces facing away from the work environment was caused primarily by chloride stress-corrosion cracking, aided by loose fitting threads. Thread gaps constitute a crevice where an aggressive chemistry is allowed to develop and attack local surfaces.

An 18-MW gas turbine exploded unexpectedly after three hours of normal operation. The catastrophic failure caused extensive damage to the rotor, casing, and nearly all turbo-compressor components. Based on their initial review, investigators believed that the failure originated at the interface between two shaft sections held together by 24 marriage bolts. Visual and SEM examination of several bolts revealed extensive deterioration of the coating layer and the presence of deep corrosion pits. It was also learned that the bolts were nearing the end of their operating life, suggesting that the effects of fatigue-assisted corrosion had advanced to the point where one of the bolts fractured and broke free. The inertial unbalance produced excessive vibration, subjecting the remaining bolts to overload failure.

The failure of T12 reheater tubes that had been in service for only 3000 h was investigated. The thickness of the tubes was visibly reduced by heavy oxidation corrosion on the inner and outer walls. The original pearlite substrate completely decomposed. Uniform oxide scale observed on the inner wall showed obvious vapor oxidation corrosion characteristics. Corrosion originated in the grain boundary, and selective oxidation occurred due to ion diffusion in the substrate. The layered oxide scale on the inner wall is related to the different diffusion rates for different cations. Exposure to high temperature corrosive flux accelerated the corrosion on the outer wall. Microstructure degradation and the corrosion characteristics observed indicate that the tubes failed primarily because of overheating, which is confirmed by calculations.

This case study describes the failure analysis of a steel nozzle in which cracking was observed after a circumferential welding process. The nozzle assembly was made from low-carbon CrMoV alloy steel that was subsequently single-pass butt welded using gas tungsten arc welding. Although no cracks were found when the welds were visually inspected, X-ray radiography showed small discontinuous surface cracks adjacent to the weld bead in the heat affected zone. Further investigation, including optical microscopy, microhardness testing, and residual stress measurements, revealed that the cracks were caused primarily by the presence of coarse untempered martensite in the heat affected zone due to localized heating. The localized heating was caused by high welding heat input or low welding speed and resulted in high transformation stresses. These transformation stresses, working in combination with thermal stresses and constraint conditions, resulted in intergranular brittle fracture.

An investigation of a damaged crankshaft from a horizontal, six-cylinder, in-line diesel engine of a public bus was conducted after several failure cases were reported by the bus company. All crankshafts were made from forged and nitrided steel. Each crankshaft was sent for grinding, after a life of approximately 300,000 km of service, as requested by the engine manufacturer. After grinding and assembling in the engine, some crankshafts lasted barely 15,000 km before serious fractures took place. Few other crankshafts demonstrated higher lives. Several vital components were damaged as a result of crankshaft failures. It was then decided to send the crankshafts for laboratory investigation to determine the cause of failure. The depth of the nitrided layer near fracture locations in the crankshaft, particularly at the fillet region where cracks were initiated, was determined by scanning electron microscope (SEM) equipped with electron-dispersive X-ray analysis (EDAX). Microhardness gradient through the nitrided layer close to fracture, surface hardness, and macrohardness at the journals were all measured. Fractographic analysis indicated that fatigue was the dominant mechanism of failure of the crankshaft. The partial absence of the nitrided layer in the fillet region, due to over-grinding, caused a decrease in the fatigue strength which, in turn, led to crack initiation and propagation, and eventually premature fracture. Signs of crankshaft misalignment during installation were also suspected as a possible cause of failure. In order to prevent fillet fatigue failure, final grinding should be done carefully and the grinding amount must be controlled to avoid substantial removal of the nitrided layer. Crankshaft alignment during assembly and proper bearing selection should be done carefully.

A pressure vessel failed causing an external fire on a nine-story coke gasifier in a refinery power plant. An investigation revealed that the failure began as cracking in the gasifier internals, which led to bulging and stress rupture of the vessel shell, and the escape of hot syngas, setting off the fire. The failure mechanisms include stress relaxation cracking of a large diameter Incoloy 825 tube, stress rupture of a 4.65 in. thick chromium steel shell wall, and the oxidation of chromium steel exposed to hot syngas. The gasifier process and operating conditions that contributed to the high-temperature degradation were also analyzed and are discussed.

The main shaft in a locomotive turbocharger fractured along with an associated bearing sleeve. Visual and fractographic examination revealed that the shaft fractured at a sharp-edged groove between two journals of different cross-sectional area. The dominant failure mechanism was low-cycle rotation-bending fatigue. The bearing sleeve failed as a result of abrasive and adhesive wear. Detailed metallurgical analysis indicated that the sleeve and its respective journal had been subjected to abnormally high temperatures, increasing the amount of friction between the sleeve, bearing bush, and journal surface. The excessive heat also softened the induction-hardened case on the journal surface, decreasing its fatigue strength. Fatigue crack initiation occurred at the root fillet of the groove because of stress concentration.

An investigation was conducted to determine what caused a bearing sleeve in a locomotive turbocharger to fail. The sleeve, which is made of nitrided 38CrMoAl steel, fractured at the transition fillet between the cylinder and plate. Visual examination revealed significant wear on the external surface of the cylinder, with multiple origin fatigue fracture appearing to be the dominant fracture mechanism. Metallurgical examination indicated that the nitrided layer was not as deep as it was supposed to be and had worn away on the outer surface of the sleeve, exposing the soft matrix underneath. This led to further wear and an increase in friction between the sleeve and bearing bush. Fatigue crack initiation occurred at the root fillet because of stress concentration and large frictional forces. Insufficient nitriding depth facilitated the propagation of fatigue cracks.

A main steam pipe was found to be leaking due to a large circumferential crack in a pipe-to-fitting weld in one of two steam leads between the superheater outlet nozzles and the turbine stop valves (a line made of SA335-P22 material). The main crack surface was found to be rough, oriented about normal to the outside surface, and had a dark oxidized appearance. The cracking was found to be predominantly intergranular. Distinct shiny bands that etched slower than the remainder of the sample at the top of each individual weld bead were revealed by microscopic examination. These bands were found contain small cracks and microvoids. A mechanism of intergranular creep rupture at elevated temperature was identified as a result of a series of stress-rupture and tensile tests. It was revealed by the crack shape that cracking initiated on the pipe exterior, then propagated inward and in the circumferential direction in response to a bending moment load. It was concluded that the primary cause of failure was the occurrence of bending stresses that exceeded the stress levels predicted by design calculations and that were higher than the maximum allowable primary membrane stress.

A portion of two large spur tooth bull gears made from 4147H Cr-Mo alloy steel that had spalling teeth was submitted for evaluation. The gears were taken from a final drive wheel reduction unit of a very large open-pit mining truck. The parts had met the material and initial heat treat hardening specifications. The mode of failure was tooth profile spalling. By definition, spalling originates at a case/core interface or at the juncture of a hardened/nonhardened area. The cause of this failure was either insufficient or no induction-hardened case along the active profile. The cause was activated by a nonfunctioning induction hardening coil that did not or was not allowed to harden the midprofile of several teeth.

The tubes of a stationary industrial boiler, 64 mm in diam and made of 1.25Cr-0.5Mo steel (ASME SA-213, grade T-11) failed by two different types of rupture. Noticeable swelling of the tubes in the area of rupture was revealed by visual examination. The tubes with slight longitudinal splits were interpreted to have failed by stress rupture resulting from prolonged overheating at 540 to 650 deg C as the microstructure exhibited extensive spheroidization and coalescence of carbides. The larger ruptures were tensile failures that resulted from rapid overheating to 815 to 870 deg C as a completely martensitic structure was revealed at the edges of the ruptures in these tubes because of rapid quenching by escaping fluid. The prolonged-overheating failures were concluded to have been the primary ruptures and that local loss of circulation had caused rapid overheating in adjacent tubes. Poor boiler circulation and high furnace temperatures were believed to have caused the prolonged overheating.

Important clues about the probable cause of a gun tube explosion were obtained from a fractographic and metallographic examination of the fragments. The size, distribution, and surface markings of fragments may be used to localize the explosion and deduce its intensity. Microstructural features such as voids, adiabatic shear, and structural surface alterations also indicate the explosion intensity and further allow a comparison of the tube structure near and away from the explosion zone. These, and other metallurgical characteristics, are illustrated and discussed for cases of accidental and deliberately caused explosions of large caliber gun tubes.

Plunger shafts machined from 4150 steel bar stock were involved in a series of fatigue failures. The fractures consistently occurred at two locations on the shafts: the shaft fillet and either side of a machined notch. The material specification for the shafts required 41xx series steel with a carbon content of 0.38 to 0.53%, a hardness of 35 to 40 HRC for the shaft, and a hardness of 50 to 55 HRC for the notch (which was case hardened). Analysis (visual inspection, chemical analysis, hardness testing, and magnetic particle inspection) supported the conclusions that all the fractures were fatigue-induced failures due to sharp radii in the fillets. The stress-concentrating effects of the fillets caused fatigue cracks to initiate and grow under cyclic loading until the crack depth was critical, causing the shaft to fail and rendering the assembly inoperative. Recommendations included increasing the radii of the notch and shaft fillets. If fatigue cracking had continued to be a problem with this component, shot peening of the subject radii would be appropriate. This process produces residual compressive stresses in the surface of the part, thereby retarding initiation of fatigue cracks.

Hydraulic cylinder housings were being fabricated from 4140 grade seamless steel tubing. During production, magnetic-particle inspection indicated the presence of circumferential and longitudinal cracks in a large number of cylinders. Analysis (visual inspection, dye penetrant inspection, 50x/90x/400x SEM micrographs, and metallographic analysis) supports the conclusion that the cracking problem in these components was identified as quench cracks due to their brittle, intergranular nature and the characteristic temper oxide on the fracture surfaces. Although the steel met the compositional requirements of SAE 4140, the sulfur level was 0.022% and would account for the formation of the sulfide stringers observed. Apparently, the combination of the clustered, stringer-type inclusions and the quenching conditions were too severe for this component geometry. The result was a high incidence of quench cracks that rendered the parts useless. Recommendations included changing the specification, requiring the steel to have lower sulfur concentrations. Magnetic-particle cleanliness standards should be imposed that will exclude material with harmful clusters of sulfide stringers, for example, modified AMS 2301.

This article discusses the failure of cylinder clamping rods in single cylinder diesel engines. The AISI 4140 hardened and tempered steel clamping rods were failing after 200 to 250 h of operation. The fatigue failures initiated at the root of the last thread on the clamping rod that was engaged in a blind hole in the cylinder block. The failures were caused by loose tolerances on the threads that resulted in a non-uniform distribution of load. The load was concentrated on the last threads to engage, thus causing fatigue crack nucleation at the thread root and propagation until the rod broke by overload. Changing the tolerance on the threads virtually eliminated the fatigue problem.

A cast countershaft pinion on a car puller for a blast furnace broke after one month of service; expected life was 12 months. The pinion was specified to be made of 1045 steel heat treated to a hardness of 245 HRB. The pinion steel was analyzed and was a satisfactory alternative to 1045 steel. The pinion was annealed before flame or induction hardening of the teeth to a surface hardness of 363 HRB and a core hardness of 197 HRB. The broken pinion had a tooth which had failed by fatigue fracture through the tooth root because of the low strength from incomplete surface hardening of the tooth surfaces. Contributing factors included uneven loading because of misalignment and stress concentrations in the tooth roots caused by tool marks. Greater strength was provided by oil quenching and tempering the replacement pinions to a hardness of 255 to 302 HRB. Machining of the tooth roots was revised to eliminate all tool marks. Surface hardening was applied to all tooth surfaces, including the root. Proper alignment of the pinion was ensured by carefully checking the meshing of the teeth at startup.

A forged 4130 steel cylindrical permanent mold, used for centrifugal casting of gray- and ductile-iron pipe, was examined after pulling of the pipe became increasingly difficult. In operation, the mold rotated at a predetermined speed in a centrifugal casting machine while the molten metal, flowing through a trough, was poured into the mold beginning at the bell end and ending with the spigot end being poured last. After the pipe had cooled, it was pulled out from the bell end of the mold, and the procedure was repeated. Investigation supported the conclusion that failure of the mold surface was the result of localized overheating caused by splashing of molten metal on the bore surface near the spigot end. In addition, the mold-wash compound (a bentonite mixture) near the spigot end was too thin to provide the proper degree of insulation and to prevent molten metal from sticking to the bore surface. Recommendations included reducing the pouring temperatures of the molten metal and spraying a thicker insulating coating onto the mold surface.

Hydrogen-assisted stress-corrosion cracking failure occurred in four AISI 4137 chromium molybdenum steel bolts having a hardness of 42 HRC. The normal service temperature (400 deg C, or 750 deg F) was too high for hydrogen embrittlement but, the bolts were subjected also to extended shutdown periods at ambient temperatures. The corrosive environment contained trace hydrogen chloride and acetic acid vapors as well as calcium chloride if leaks occurred. The exact service life was unknown. The bolt surfaces showed extensive corrosion deposits. Cracks had initiated at both the thread roots and the fillet under the bolt head. Multiple, branched cracking was present in a longitudinal section through the failed end of one bolt, typical of hydrogen-assisted SCC in hardened steels. Chlorides were detected within the cracks and on the fracture surface. The failed bolts were replaced with 17-4 PH stainless steel bolts (Condition H 1150M) having a hardness of 22 HRC.

Deformation, surface cracking, and spalling on the raceway of the outer ring (made of 4140 steel) of a large bearing caused it to be replaced from a radar antenna. The raceway surfaces were to be flame hardened to 55 HRC minimum and 50 HRC 3.2 mm below the surface, according to specifications. Samples from both the inner and outer rings were examined. A much lower hardness (25.2 to 18.9 HRC) was indicated during a vertical traverse 4.1 cm from the outer surface of the outer ring while slightly lower hardness values (46.8 to 54.8 HRC) were seen on the hardness traverse on the inner ring raceway. The lower hardness values were attributed to improper flame hardening. It was confirmed by metallographic examination of a 3% nital etched sample that the inner ring (tempered martensite and ferrite) and the outer ring (ferrite, scattered patches of pearlite, and martensite) were not properly austenitized. Displacement of metal on the outer raceway was revealed by elongation of grain structure. It was concluded that the failure of the raceway surface was due to incomplete austenitization caused by the improper heat treatment during flame hardening process.