A detailed fracture mechanics evaluation is the most accurate and reliable prediction of process equipment susceptibility to brittle fracture. This article provides an overview and discussion on brittle fracture. The discussion covers the reasons to evaluate brittle fracture, provides a brief summary of historical failures that were found to be a result of brittle fracture, and describes key components that drive susceptibility to a brittle fracture failure, namely stress, material toughness, and cracklike defect. It also presents industry codes and standards that assess susceptibility to brittle fracture. Additionally, a series of case study examples are presented that demonstrate assessment procedures used to mitigate the risk of brittle fracture in process equipment.

This article focuses on characterizing the fracture-surface appearance at the microscale and contains some discussion on both crack nucleation and propagation mechanisms that cause the fracture appearance. It begins with a discussion on microscale models and mechanisms for deformation and fracture. Next, the mechanisms of void nucleation and void coalescence are briefly described. Macroscale and microscale appearances of ductile and brittle fracture are then discussed for various specimen geometries (smooth cylindrical and prismatic) and loading conditions (e.g., tension compression, bending, torsion). Finally, the factors influencing the appearance of a fracture surface and various imperfections or stress raisers are described, followed by a root-cause failure analysis case history to illustrate some of these fractography concepts.

A roll assembly consisting of a forged AISI type 440A stainless steel sleeve shrink fitted over a 4340 steel shaft and further secured with tapered keys on opposite ends was crated and shipped by air. Upon arrival, the sleeve was found to have cracked longitudinally between the keyways. A roll manufacturer had successfully used the above procedure for many years to make them. Analysis (visual inspection; 150x micrograph of sections etched with a mixture of 2 parts HNO3, 2 parts acetic acid, and 3 parts HCI; electron microscopy; and stress testing) supported the conclusion that superficial working of the metal, probably insufficient hot working, produced a microstructure in which the carbide particles were not broken up and evenly distributed. As a result, the grains were totally surrounded with brittle carbide particles. This facilitated the formation of a crack at a fillet in the keyway. Crack growth was rapid once the crack had initiated, causing brittle fracture to occur.

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.

Chain link, a part of a mechanism for transferring hot or cold steel blooms into and out of a reheating furnace, broke after approximately four months of service. The link was cast from 2% Cr austenitic manganese steel and was subjected to repeated heating to temperatures of 455 to 595 deg C (850 to 1100 deg F). Examination included visual inspection, macrograph of a nital-etched specimen from an as-received chain link 1.85x, micrographs of a nital-etched specimen from an as-received chain link 100x/600x, normal microstructure of as-cast standard austenitic manganese steel 100x, micrograph of a nital-etched specimen that had been austenitized 20 min at 1095 deg C (2000 deg F) and air cooled 315x, and micrograph of the same specimen after annealing 68 h at 480 deg C (900 deg F) 1000x). Investigation supported the conclusions that the chain link failed in a brittle manner, because the austenitic manganese steel from which it was cast became embrittled after being reheated in the temperature range of 455 to 595 deg C (850 to 1100 deg F) for prolonged periods of time. The alloy was not suitable for this application, because of its metallurgical instability under service conditions.

The T-section cross member of the lifting sling failed in service while lifting a 966 kg (2130 lb) load. The L-section sling body and the cross member were made of aluminum alloy 5083 or 5086 and were joined by welding using aluminum alloy 4043 filler metal. The fracture was found by visual examination to have occurred at the weld joining the sling body and the cross member. Inadequate joint penetration and porosity was revealed by macrographic examination of the weld. Lower silicon content and a higher magnesium and manganese content than the normal for alloy 4043 filler metal were found during chemical analysis. It was revealed by examination of the ends of the failed cross member that a rotational force that had been applied on the cross member caused it to fracture near the sling body. It was concluded that brittle fracture at the weld was caused by overloading which was attributed to the misalignment of the sling during loading. Aluminum alloy 5183 or 5356 filler metal was recommended to be used to avoid brittle welds.

During the lifting of a piece of machinery by means of an overhead travelling crane the hook fractured suddenly. The load was attached to the hook by means of fiber rope slings and rupture occurred in a plane which appeared to coincide with the sling loop nearest to the back of the hook. The rated capacity of the crane was 15 tons. At the time of the mishap it was being used to lift one end of a hydraulic cylinder with a total weight of about 27 tons. Fracture was of the cleavage type throughout. There was no evidence of any prior deformation of the material in the vicinity, nor was there any indication of a pre-existing crack or major discontinuity at the point of origin. A sulfur print suggested the hook had been forged from a billet cogged down from an ingot of semi-killed steel. Failure of this hook was attributed to strain-age embrittlement of the material at the surface of the intrados.

Several thousands of new 16 mm diam alloy steel sling chains used for handling billets failed by chain-link fractures. No failures were found to have occurred before delivery of the new chains. It was observed that the links had broken at the weld. It was found that all failures had occurred in links having hardness values in the range of 375 to 444 HRB. It was revealed by the supplier that the previous hardness level of 302 to 375 HRB was increased to minimize wear which made the links were made notch sensitive and resulted in fractures that initiated at the butt-weld flash on the inside surfaces of the links. A further reduction in ductility was believed to have been caused by lower temperatures during winter months. Thus, the failure was concluded to have been caused in a brittle manner caused by the notch sensitivity of the high hardness material at lower temperatures. The chains were retempered to a hardness of 302 to 375 HRB as a corrective measure and subsequently ordered chains had this hardness as a requirement.

A hook, which was marked for a safe working load of 2 tons, failed while lifting a load of approximately 35 cwts. Fracture took place at the junction of the shank with the hook portion, at which no fillet radius existed. Except for an annular region round the periphery, which was of a smooth texture, the fracture was brightly crystalline indicative of a brittle failure. Microscopic examination showed the material was a low-carbon steel in the normalized condition; no abnormal features were observed. The basic cause of failure was the presence of a fatigue crack at the change of section where the shank joined the hook portion. To minimize the possibility of fatigue cracking, it was recommended that a generous radius be provided at the change of section.

A 58.4 cm (23 in.) diam heavy-duty brake drum component of a cable-wound winch broke into two pieces during a shutdown period. Average service life of these drums was two weeks; none had failed by wear. The drums were sand cast from ductile iron. During haul-out, the cable on the cable drum drove the brake drum, and resistance was provided by brake bands applied to the outside surface of the brake drum. Friction during heavy service was sufficient to heat the brake drum, clutch mount, and disk to a red color. Examination of the assembly indicated that the brake drum would cool faster than its mounts and would contract onto them. Brittle fracture of the brake drum occurred as a result of thermal contraction of the drum web against the clutch mount and the disk. The ID of the drum web was enlarged sufficiently to allow for clearance between the web and the clutch mount and disk at a temperature differential of up to 555 deg C (1000 deg F). With the adoption of this procedure, brake drums failed by wear only.

The specially designed sand-cast low-alloy steel jaws that were implemented to stretch the wire used in prestressed concrete beams fractured. The fractures were found to be macroscale brittle and exhibited very little evidence of deformation. The surface of the jaws was disclosed by metallographic examination to be case carburized. The case was found to be martensite with small spheroidal carbides while the core consisted of martensite plus some ferrite. The fracture was revealed to be related to shrinkage porosity. Tempering was revealed to be probably limited to about 150 deg C by the hardness values (close to the maximum hardness values attainable) for the core. It was interpreted that the low tempering temperature used may have contributed to the brittleness. The procedures used for casting the jaws were recommended to be revised to eliminate the internal shrinkage porosity. Tempering at a slightly higher temperature to reduce surface and core hardness was recommended.

A railway tank car developed a fracture in the region of the sill and shell attachment during operation at -34 deg C (-30 deg F). On either side of the sill-support member, cracking initiated at the weld between a 6.4 mm thick frontal cover plate and a 1.6 mm thick side support plate. The crack then propagated in a brittle manner upward through the side plate, through the welds attaching the side plate to a 25 mm (1 in.) thick shell plate (ASTM A212, grade B steel), and continued for several millimeters in the shell plate before terminating. Other plates involved were not positively identified but were generally classified as semi-killed carbon steels. Investigation (visual inspection, hardness testing, chemical analysis, Charpy V-notch testing, and drop-weight testing) supported the conclusions that the fracture was initiated by weld imperfections and propagated in a brittle manner as a result of service stresses acting on the plate having low toughness at the low service temperatures encountered. Recommendations included that the specifications for the steel plates be modified to include a toughness requirement and that improved welding and inspection practices be performed to reduce the incidence of weld imperfections.

Two examples involved brittle fracture promoted by small fatigue cracks owing to welding deficiencies. Other parts involving inadequate welding were a ski-wheel axle flange, ski fitting (brackets), and undercarriage shock strut stub assembly. In an attach fitting for an engine mount, weld cracks (severe stress concentrations) formed during repair welding. Cracks were severely oxidized. The main cause was incorrect repair and inadequate inspection of the fitting. In a cast CrNi alloy ski wheel axle, brittle fatigue failure emanated from welding cracks (notches). These welding cracks formed during the fabrication of the axle mounting plate. So-called all-purpose electrodes were used. Thus, the main cause for failure was a manufacturing deficiency-fatigue failure developed because of improper welding during fabrication of the axle. The proper electrode should have been used.

During installation, a clamp-strap assembly, specified to be type 410 stainless steel-austenitized at 955 to 1010 deg C (1750 to 1850 deg F), oil quenched, and tempered at 565 deg C (1050 deg F) for 2 h to achieve a hardness of 30 to 35 HRC, and used for securing the caging mechanism on a star-tracking telescope, fractured transversely across two rivet holes closest to one edge of the pin retainer in a completely brittle manner. Comparison with a non-failed strap using microscopic examination, spectrographic analysis, and slow-bend tests showed that both fit the 410 stainless steel specs, but hardness and grain size were different. Reheat treatment of full-width specimens showed that coarse grain size (ASTM 2 to 3) was responsible for the brittle fracture, and excessively high temperature during austenitizing caused the large grain size in the failed strap. The fact that the hardness of the strap that failed was lower than the specified hardness of 30 to 35 HRC had no effect on the failure because that of the non-failed strap was even lower. Recommendation was that the strap should be heat treated as specified to maintain the required ductility and grain size.

A 3 in. diam shaft was found to have suffered excessive wear on one of the journals and was built up by welding. While it was in the lathe prior to turning down the built-up region, a crack was discovered in the root of the oil-seal groove and subsequently the end of the shaft was broken off with hammer blows. The fracture surface was duplex in nature, there being an annular region surrounding a central zone, which suggests that the fracture developed in two stages. Microscopic examination confirmed that the fracture was of the brittle type. The shaft material showed a microstructure typical of a medium-carbon steel (carbon approximately 0.4%) in the normalized condition, a material not weldable by ordinary methods. It was concluded that the post-welding crack arose primarily from the thermal contraction which developed in the weld metal on cooling. It is probable that if the built-up zone had extended beyond the oil seal groove, failure in the manner would not have occurred. Experience indicated however, that failure from fatigue cracking would still have been likely to occur.

A ring clamp (8740 (AMS 6322), steel forged and cadmium plated) used for attaching ducts to an aircraft engine became loose after three hours of service. When the clamp was removed from the engine, the hinge tabs on one clamp half were found to be broken. Analysis (visual inspection and microscopic and metallographic examination) supported the conclusion that both hinge tabs on the clamp half fractured in a brittle manner as the result of gross overheating, or burning, during forging. The mechanical properties of the metal, especially toughness and ductility, were greatly reduced by burning. Evidence that burning was confined to the hinge end of the clamp indicated that the metal was overheated before or during the upset forging operation. Recommendations included notifying the supplier of the burned condition on the end of the clamp. The clamps should be macroetched before cadmium plating to detect overheating. The clamps in stock should be inspected to ensure that the metal had not been weakened by overheating during the upset forging operation.

The 8620 steel latch tip, carburized and then induction hardened to a minimum surface hardness of 62 HRC, on the main-clutch stop arm on a business machine fractured during normal operation when the latch tip was subjected to intermittent impact loading. Fractographic examination 9x showed a brittle appearance at the fractures. Micrograph examination of an etched section disclosed several small cracks. Fracture of the parts may have occurred through similar cracks. Also observed was a burned layer approximately 0.075 mm (0.003 in.) deep on the latch surface, and hardness at a depth of 0.025 mm (0.001 in.) in this layer was 52 HRC (a minimum of 55 HRC was specified). Thus, the failure was caused by brittle fracture in the hardness-transition zone as the result of excessive impact loading. The burned layer indicated that the cracks had been caused by improper grinding after hardening. Redesign was recommended to include reinforcing the backing web of the tip, increasing the radius at the relief step to 1.5 x 0.5 mm (0.06 x 0.02 in.), the use of proper grinding techniques, and a requirement that the hardened zone extend a minimum of 1.5 mm (0.06 in.) beyond the step.

A rocket-motor case made of consumable-electrode vacuum arc remelted D-6ac alloy steel failed during hydrostatic proof-pressure testing. Close visual examination, magnetic-particle inspection, and hardness tests showed cracks that appeared to have occurred after austenitizing but before tempering. Microscopic examinations of ethereal picral etched sections indicated that the cracks appeared before or during the final tempering phase of the heat treatment and that cracking had occurred while the steel was in the as-quenched condition, before its 315 deg C (600 deg F) snap temper. Chemical analysis of the cracked metal showed a slightly higher level of carbon than in the component that did not crack. X-ray diffraction studies of material from the fractured dome showed a very low level of retained austenite, and chemical analysis showed a slightly higher content of carbon in the metal of the three cracked components. Bend tests verified the conclusion that the most likely mechanism of delayed quench cracking was isothermal transformation of retained austenite to martensite under the influence of residual quenching stresses. Recommendations included modifying the quenching portion of the heat-treating cycle and tempering in the salt pot used for quenching, immediately after quenching.

The clapper in a 250 mm diam disk valve (made from ASTM A36 steel, stress relieved and cadmium plated) fractured at the welded joint between the clapper and a 20 mm diam support rod (also made of same material). The valve contained a stream of gas consisting of 55% H2S, 39% CO2, 5% H2, and 1% hydrocarbons at 40 deg C and 55 kPa during operation. Voids on the fracture surface and evidence of incomplete weld penetration were revealed by examination. Brittle fracture was indicated by the overall appearance through some fatigue beach marks were observed. Very narrow bands of high hardness were revealed at the edges of the weld metal. It was revealed by chemical analysis of this band that a stainless steel filler metal had been used which produced mixed composition at the weld boundaries. The plating material was revealed to be nickel by chemical analysis. It was concluded that clapper failed by fatigue and brittle fracture because it was welded with an incorrect filler metal. A clapper assembly was welded with a low-carbon steel filler metal, then cadmium plated.

Splined rotor shafts (constructed from 1151 steel) used on small electric motors were found to miss one spline each from several shafts before the motors were put into service. Apparent peeling of splines on the induction-hardened end of each rotor shaft was revealed by visual and stereo-microscopic examination. One tooth on each shaft was found to be broken off. It was revealed by metallographic examination of an unetched section through the fractured tooth that the fracture surface was concave and had an appearance characteristic of a seam. Partial decarburization of the surface was revealed after etching with 1% nital. The presence of a crack, with typical oxides found in seams at its root, was disclosed by an unetched section through the shaft in an area unaffected by induction heating. The etched samples revealed similar decarburization as was noted on the fracture surface of the tooth. It was concluded that the seam had been present before the shaft was heat treated and these seams acted as stress raisers during induction hardening to cause the shaft failure. It was recommended that the specifications should specify that the shaft material should be free of seams and other surface imperfections.