The AISI 1080 steel crankshaft of a large-capacity double-action stamping press broke in service and was repair welded. Shortly after the crankshaft was returned to service, the repair weld fractured. The repair-weld fracture was examined ultrasonically which revealed many internal reflectors, indicating the presence of slag inclusions and porosity. A low-carbon steel flux-cored filler metal was used in repair welding the crankshaft, without any preweld or postweld heating. This resulted in the formation of martensite in the HAZ. The repair weld failed by brittle fracture, which was attributed to the combination of weld porosity, many slag inclusions and the formation of brittle martensite in the HAZ. A new repair weld was made using an E312 stainless steel electrode, which provides a weld deposit that contains considerable ferrite to prevent hot cracking. Before welding, the crankshaft was preheated to a temperature above which martensite would form. After completion, the weld was covered with an asbestos blanket, and heating was continued for 24 h. During the next 24 h, the temperature was slowly lowered. The result was a crack-free weld.

A large shackle used in operating a dragline bucket failed in service. The shackle was made of a cast low-alloy steel (similar to AISI 4320) heat treated to a hardness of 415 BN. The shackle failed by fracturing through the load-bearing region. Examination of the fracture surface revealed a fatigue crack through about one-third of the cross section. A secondary fatigue crack, perpendicular to the main fracture, was also observed. The composition of the weld deposit corresponded to a heat treatable flux-cored arc welding filler material that was known to have been used for repair welding of these products. This shackle failed because of fatigue initiating at hydrogen cracks that had occurred in the HAZ of a repair weld. The weld had been made with a heat-treatable filler material, and a full postweld heat treatment had been performed. However, a low-hydrogen filler material had not been used to make the weld. Repair welds in high-strength steel castings should always be made with low-hydrogen filler materials.

A Ti-6Al-4V alloy pressure vessel failed during a proof-pressure test, fracturing along the center girth weld. The girth joints were welded with the automatic gas tungsten arc process utilizing an auxiliary trailing shield attached to the welding torch to provide inert-gas shielding for the exterior surface of the weld. A segmented backup ring with a gas channel was used inside the vessel to shield the weld root. The pressure vessel failed due to contamination of the fusion zone by oxygen, which resulted when the gas shielding the root face of the weld was diluted by air that leaked into the gas channel. Thermal stresses cracked the embrittled weld, exposing the crack surfaces to oxidation before cooling. One of these cracks caused a stress concentration so severe that failure of the vessel wall during the proof test was inevitable. A sealing system at the split-line region of the segmented backup ring was provided, and a fine-mesh stainless steel screen diffuser was incorporated in the channel section of the backup ring to prevent air from leaking in. A titanium alloy color chart was furnished to permit correlation of weld-zone discoloration with the degree of atmospheric contamination.

During the final shop welding of a large armature for a direct-current motor (4475 kW, or 6000 hp), a loud bang was heard, and the welding operation stopped. When the weld was cold, nondestructive evaluation revealed a large crack adjacent to the root weld. Investigation showed the main crack had propagated parallel to the fusion boundary along the subcritical HAZ and was associated with long stringers of type II manganese sulfide (MnS) inclusions. This supported the conclusion that the weld failed by lamellar tearing as a result of the high rotational strain induced at the root of the weld caused by the weld design, weld sequence, and thermal effects. Recommendations included removing the old weldment to a depth beyond the crack and replacing this with a softer weld metal layer before making the main weld onto the softer layer.

Parts of 21Cr-6Ni-9Mn stainless steel that had been forged at about 815 deg C (1500 deg F) were gas tungsten arc welded. During postweld inspection, cracks were found in the HAZs of the welds. Welding had been done using a copper fixture that contacted the steel in the area of the HAZ on each side of the weld but did not extend under the tungsten arc. In SEM examination, the cracks appeared to be intergranular and extended to a depth of approximately 1.3 mm (0.05 in.). The crack appearance suggested that the surface temperature of the HAZ could have melted a film of copper on the fixture surface and that this could have penetrated the stainless steel in the presence of tensile thermal-contraction stresses. The cracks in the weldments were a form of liquid-metal embrittlement caused by contact with superficially melted copper from the fixture and subsequent grain-boundary attack of the stainless steel in an area under residual tensile stress. The copper for the fixtures was replaced by aluminum. No further cracking was encountered.

A weld that attached the head to the shell of a preheater containing steam at 1.4 MPa (200 psi) and was used in the manufacture of paper cracked in service. The original joint contained a 6.4 by 50 mm backing ring that had been tack welded to the inside surface of the shell in a position to project beyond the fully beveled top edge of the shell. The projecting edge of the ring fitted against a deep undercut on the inner corner of the rim of the head. The internal 90-deg angle in this undercut was sharp, with almost no fillet. A bevel from the lower edge of the undercut to the outside of the head completed the groove for the circumferential attachment weld. Investigation (visual inspection and actual size views etched in hot 50% hydrochloric acid) supported the conclusion that cracking occurred in the HAZ in the head of the original design, originating in the sharp corner of the undercut, which was an inherent stress raiser. Recommendations included revised joint design to ensure full root penetration.

Several of the welds in a hoist carriage tram-rail assembly fabricated by shielded metal arc welding the leg of a large T-section 1020 steel beam to the leg of a smaller T-section 1050 steel rail failed in one portion of the assembly. Four weld cracks and several indefinite indications were found by magnetic-particle inspection. The cracks were revealed by metallographic examination to have originated in the HAZs in the rail section. Cracks in welds and in HAZs resulting from arcing the electrode adjacent to the weld and weld spatter were also revealed. The tram-rail assembly was concluded to have failed by fatigue cracking in HAZs. The fatigue cracking was initiated and propagated by vibration of the tram rail by movement of the hoist carriage on the rail. As a corrective measure, welding procedures were improved and the replacement rail assemblies were preheated and postheated.

Arms bolted to powerline towers were falling off two weeks after installation. Metallurgical and chemical analysis performed on the base metal, weld zone, and heat-affected zone showed acceptable quality material. Residual stress appeared to be responsible for the high failure rate. The sources of residual stress included welding, environment, and assembly operation.

Several fractures occurred in flange studs used for remote handling of radioactive equipment. The studs, of quenched-and-tempered type 414 stainless steel, fractured in the HAZs produced in the studs during the circumferential welding that joined the studs to the flanges. The weld deposits were of type 347 stainless steel, and the flanges were type 304 stainless steel. Metallographic examination of the failed studs revealed that the HAZs contained regions of martensite and that intergranular cracks, which initiated at the stud surfaces during welding, propagated to complete separation under subsequent loading. The studs fractured under service loads as a result of intergranular crack propagation in the HAZ. Rapid heating and cooling during attachment welding produced a martensitic structure in the HAZ of the stud, which cracked circumferentially from the combination of thermal-gradient and phase-change stresses. Joining the studs to the flanges by welding should be discontinued. They should be attached by screw threads, using a key and keyway to prevent turning in service.

Corrosion in a closed-loop cooling water system constructed of austenitic stainless steel occurred during an extended lay up of the system with biologically contaminated water. The characteristics of the failure were those of microbiologically influenced corrosion (MIC). The corrosion occurred at welds and consisted of large subsurface void formations with pinhole penetrations of the surfaces. Corrosive attack initiated in the heat affected zones of the welds, usually immediately adjacent to fusion lines. Stepwise grinding, polishing, and etching through the affected areas revealed that voids generally grew in the wrought material by uniform general corrosion. Tunneling or worm-holing was also observed, whereby void extension occurred by initiating daughter voids probably at flaws or other inhomogeneities. Selective attack occurred within the fusion zone, i.e., within the cast two-phase structure of the weld filler itself. The result was a void wall which consisted of a rough and porous ferritic material, a consequence of preferential attack of the austenitic phase and slightly lower rate of corrosive attack of the ferrite phase. The three-dimensional spongy surface was studied optically and with the scanning electron microscope.

A section from a stop-block guide fell to the floor on a crane runway after it failed. A brittle crystalline-type break was disclosed by examination of the fracture surface. The point of initiation was in a hardened heat-affected layer that had developed during flame cutting and welding. The metal was identified to be 1020 steel. It was indicated by the coarse as-rolled structure (grain size of ASTM 00 to 4) of the base metal that the weldment (stop block and guide) had not been normalized. The brittle failure was evaluated to have been initiated at a metallurgical and mechanical notch produced by flame cutting and welding. As corrective measures, fully silicon-killed 1020 steel with a maximum grain size of ASTM 5 were used to make new stop-block weldments. The weldments were normalized at 900 deg C after flame cutting and welding to improve microstructure and impact strength. All flame-cut surfaces were ground to remove notches.

A large pressure vessel designed for use in an ammonia plant failed during hydrostatic testing. It was fabricated from ten Mn-Cr-Ni-Mo-V steel plates which were rolled and welded to form ten cylindrical shell sections and three forgings of similar composition. The fracture surfaces were metallographically examined to be typical for brittle steel fracture and associated with the circumferential weld that joined the flange forging to the first shell section. Featureless facets in the HAZ were observed and were revealed to be the fracture-initiation sites. Pronounced banding in the structure of the flange forging was revealed by examination. A greater susceptibility to cracking was interpreted from the higher hardenability found within the bands. Stress relief was concluded to have not been performed at the specified temperature level (by hardness and impact tests) which caused the formation of hard spots. The mode of crack propagation was established by microstructural examination to be transgranular cleavage. It was concluded that failure of the pressure vessel stemmed from the formation of transverse fabrication cracks in the HAZ fostered by the presence of hard spots. It was recommended that normalizing and tempering temperatures be modified and a revised forging practice explored.

The broken end of a shaft from a centrifugal pump had a smooth fracture surface characteristic of failure from fatigue. Failure occurred in the plane of the keyway end and followed a slightly helical path, indicating that combined bending and torsional stresses were responsible. The material was a Cr-Mo-Ni alloy steel of the En 19 type in the hardened and tempered condition and of satisfactory quality. The assembly also included a copper sleeve attached by a circumferential braze behind the plane of fracture. The cracks were examined for the presence of copper, thinking that penetration by molten copper may have played a role, but no evidence was seen. An absence of chromium plating at the region of the heat-affected zone was also observed but could not be explained. Unfortunately, the end portion of the shaft was not available for examination.

A roadarm for a tracked vehicle failed during preproduction vehicle testing. The arm was a weldment of two cored low-alloy steel sand castings specified to ASTM A 148, grade 120–95. A maximum carbon content of 0.32% was specified. The welding procedure called for degreasing and gas metal arc welding; neither preheating nor postheating was specified. The filler metal was E70S-6 continuous consumable wire with a copper coating to protect it from atmospheric oxidation while on the reel. Analysis of the two castings revealed that the carbon content was higher than specified, ranging from 0.40 to 0.44%. The fracture occurred in the HAZ , where quenching by the surrounding metal had produced a hardness of 55 HRC. Some roadarms of similar carbon content and welded by the same procedure had not failed because they had been tempered during a hot-straightening operation. Brittle fracture of the roadarm was caused by a combination of too high a carbon equivalent in the castings and the lack of preheating and postheating during the welding procedure. A pre-heat and tempering after welding were added to the welding procedure.

An austenitic stainless steel (type 316/316L stainless steel, schedule 40, 64 mm (2.5 in.) diam and larger) piping network used in the fire-sprinkler system in a large saltwater passenger and car ferry failed by rapid leaking. Operating conditions involved stagnant seawater at ambient temperatures. The pipe was in service for four weeks when three leaks appeared. Investigation (visual inspection and photographic images) supported the conclusion that the failure was caused by attack and corrosion damage of Cl ions in conditions that were ideal for three modes of highly accelerated pitting of austenitic stainless steel: the bottom surface, weld or HAZ pits, and crevices. Recommendations included proper material selection for piping, flanges, and weld rods with greater corrosion resistance. Proper filtering to prevent entrained abrasives and timely breakdown inspections were also advised.

A weld in a fuel-line tube broke after 159 h of engine testing. The 6.4-mm (0.25-in.) OD x 0.7-mm (0.028-in.) wall thickness tube and the end adapters were all of type 347 stainless steel. The butt joints between tube and end adapters were made by automated gas tungsten arc (orbital arc) welding. It was found that the tube had failed in the HAZ. Examination of a plastic replica of the fracture surface in a transmission electron microscope established that the crack origin was at the outer surface of the tube. The crack growth was by fatigue; closely spaced fatigue striations were found near the origin, and more widely spaced striations near the inner surface. The quality of the weld and the chemical composition of the tube both conformed to the specifications. However, the fuel-line assembly had vibrated excessively in service. The fuel-line fracture was caused by fatigue induced by severe vibration in service. Additional tube clamps were provided to damp the critical vibrational stresses. No further fuel-line fractures were encountered.

Two aircraft-engine tailpipes of 19-9 DL stainless steel (AISI type 651) developed cracks along longitudinal gas tungsten arc butt welds after being in service for more than 1000 h. Binocular-microscope examination of the cracks in both tailpipes revealed granular, brittle-appearing surfaces confined to the HAZs of the welds. Microscopic examination of sections transverse to the weld cracks showed severe intergranular corrosion in the HAZ. The fractures appeared to be caused by loss of corrosion resistance due to sensitization, that could have been induced by the temperatures attained during gas tungsten arc welding. Tests demonstrated the presence of sensitization in the HAZ of the gas tungsten arc weld. The aircraft engine tailpipe failures were due to intergranular corrosion in service of the sensitized structure of the HAZs produced during gas tungsten arc welding. All gas tungsten arc welded tailpipes should be postweld annealed by re-solution treatment to redissolve all particles of carbide in the HAZ. Also, it was suggested that resistance seam welding be used, because there would be no corrosion problem with the faster cooling rate characteristic of this technique.

A series of resistance spot welds joining Z-shape and C-shape members of an aircraft drop-tank structure failed during ejection testing. The members were fabricated of alclad aluminum alloy 2024-T62. The back surface of the C-shape members showed severe electrode-indentation marks off to one side of the spot weld, suggesting improper electrode contact. Visual examination of the weld fractures showed that the weld nuggets varied considerably in size, some being very small and three exhibiting an HAZ but no weld. Of 28 welds, only nine had acceptable nugget diameters and fusion-zone widths. The weld deficiencies were traced to problems in forming and fit-up of the C-shape members and to difficulties in alignment and positioning of the weld tooling. The failure of the resistance spot welds was attributed to poor weld quality caused by unfavorable fit-up and lack of proper weld-tool positioning. The problem could be solved by better forming procedures to provide an accurate fit-up that would not interfere with electrode alignment.

A throttle arm of an aircraft engine fractured and caused loss of engine control. The broken part consisted of a 6.4-mm (1/4-in.) diam medium-carbon steel rod with a thread to fit a knurled brass nut that was inserted into the throttle knob. The threaded rod had been welded to the throttle-linkage bar by an assembly-weld deposit made on the rod adjacent to the threaded portion. The fracture surface exhibited a coarse-grain brittle texture with an initiating crack at a thread root. The throttle-arm failed by brittle fracture because of the presence of cracks at the thread roots that were within the HAZ of the adjacent weld deposit. The heat of welding had generated a coarse-grain structure with a weak grain-boundary network of ferrite that had not been corrected by postweld heat treatment. The combination of the cracks and this unfavorable microstructure provided a weakened condition that resulted in catastrophic, brittle fracture under normal applied loads. The design was altered to eliminate the weld adjacent to the threaded portion of the rod.

An aircraft fuel-nozzle-support assembly exhibited cracks along the periphery of a fusion weld that attached a support arm to a fairing in a joint that approximated a T-shape in cross section. The base metal was type 321 stainless steel. Examination showed a good-quality weld penetrating to the support arm beneath, but revealed notch configurations at the inner mating surfaces at each edge of the fairing, the result of welding a poor fit-up of the support arm to the fairing. Fractures that originated at the cracks were examined by stereomicroscope and were found to contain fatigue marks that indicated crack propagation from multiple origins at the inner surface of the weld edge. Fatigue cracking was initiated at stress concentrations created by the notches at the inner surfaces between the support arm and the fairing, enhanced by poor fit-up in preparation for welding.