Search Results for Grain boundary diffusion

An alloy IN-690 (N06690) incinerator liner approximately 0.8 mm (0.031 in.) thick failed after only 250 h of service burning solid waste. Investigation supported the conclusion that the root cause of the failure was overfiring during startup and sulfidation of the nickel-base alloy. No recommendations were made.

After only a short time in service, oil-fired orchard heaters made of galvanized low-carbon steel pipe, 0.5 mm (0.020 in.) in thickness, became sensitive to impact, particularly during handling and storage. Most failures occurred in an area of the heater shell that normally reached the highest temperature in service. A 400x etched micrograph showed a brittle and somewhat porous metallic layer about 0.025 mm (0.001 in.) thick on both surfaces of the sheet. Next to this was an apparently single-phase region nearly 0.05 mm (0.002 in.) in thickness. The examination supported the conclusion that prolonged heating of the galvanized steel heater shells caused the zinc-rich surface to become alloyed with iron and reduce the number of layers. Also, heating caused zinc to diffuse along grain boundaries toward the center of the sheet. Zinc in the grain boundaries reacted with iron to form the brittle intergranular phase, resulting in failure by brittle fracture at low impact loads during handling and storage. Recommendation included manufacture of the pipe with aluminized instead of galvanized steel sheet for the combustion chamber.

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.

Severe pitting corrosion of a carbon steel tube was observed in the air preheater of a power plant, which runs on rice straw firing. Approximately 1450 tubes were removed from Stage 3 of the preheater (air inlet and flue gas outlet) due to corrosion and local bursting. Samples from Stage 2 (where corrosion was low) and Stage 3 (severe corrosion) were taken and subjected to visual inspection, SEM, x-ray diffraction, microhardness measurement, and chemical and microstructural analysis. It was determined that extended non-operation of the plant resulted in the settlement of corrosive species on the tubes in Stage 3. The complete failure of the tube occurred due to diffusion of these elements into the base metal and precipitation of potassium and chlorine compounds along the grain boundaries, with subsequent dislodging of grains. The nonmetallic inclusions acted as nucleating sites for local pitting bursting. Nonuniform heat transfer in Stage 3 operation accelerated the selective corrosion of front-end tubes. The relatively high heat transfer in this stage resulted in condensation of some corrosive gases and consequent corrosion. Continuous operation of the plant with some precautions during assembly of the tubes reduced the corrosion problem.

Microstructural examinations on transverse cross sections of a steam reformer tube, showed the presence of large macrovoids elongated in the radial direction and emanating from the internal surface of the tube. The macrovoids were located at the interdendritic regions, and were partially filled by a Mn-Fe bearing chromium oxide film. The areas adjacent to the oxide film were chemically depleted in C, Cr and Mn and rich in Fe and Ni. Associated with this depletion were a large concentration of microvoids. It was suggested that the dissolution of carbides in areas surrounding the macrovoids and the concentration of stresses at their tips, caused extensive localized plastic deformation which led to the formation of microvoids and subsequently to the spalling of the oxide film. The non-protective character of the film induced a progressive deterioration of the grain boundaries properties. Grain boundary sliding and dislocation motion were enhanced, causing a local increase in the steady state strain rate and the premature failure of the tube.

During a planned shut-down in 1990 it appeared that the bottom manifold parts made of wrought Incoloy 800H had undergone diametrical expansion of up to 2% due to creep. Further, cracking at the outer diam was found. It was decided to replace these parts. Microscopical investigations showed that the cracking could not be caused by creep. It was found that the cracking was confined to a 4-mm deep coarse-grained zone (ASTM 0-1) at the outer diameter. The cracking appeared to be caused by strain-induced intergranular oxidation. When the cracks reached the fine-grained material, the oxidation-cracks stopped. To determine the residual creep life of the sound (non-cracked) bottom manifold material, iso-stress creep tests were performed. It was found that tertiary creep started at 7% strain. The time-to-rupture was greater than 100,000 h. It was concluded that the bottom manifold (and thus the furnace) could be used safely during the foreseen production period.

Pinhole defects were found in a main combustion chamber made from NARloy-Z after an unexpectedly short time in service. Analysis indicated that the throat section of the liner had been exposed to very severe environmental conditions of high temperature and high oxygen content, which caused ductility loss and grain-boundary separation. The excessive oxygen content in the liner was attributed to diffusion from an oxygen-rich environment that had resulted from nonuniform mixing of propellants. The internal oxygen embrittled the alloy and reduced its thermal conductivity, which resulted in a higher hot-gas wall temperature and associated degradation of mechanical properties.

Thermomechanical fatigue (TMF) is the general term given to the material damage accumulation process that occurs with simultaneous changes in temperature and mechanical loading. TMF may couple cyclic inelastic deformation accumulation, temperature-assisted diffusion within the material, temperature-assisted grain-boundary evolution, and temperature-driven surface oxidation, among other things. This article discusses some of the major aspects and challenges of dealing with TMF life prediction. It describes the damage mechanisms of TMF and covers various experimental techniques to promote TMF damage mechanisms and elucidate mechanism coupling interactions. In addition, life modeling in TMF conditions and a practical application of TMF life prediction are presented.

A brazing-furnace muffle 34 cm (13 in.) wide, 26 cm (10 in.) high, and 198 cm (78 in.) long, was fabricated from nickel-base high-temperature alloy sheet and installed in a gas-fired furnace used for copper brazing of various assemblies. The operating temperature of the muffle was reported to have been closely controlled at the normal temperature of 1175 deg C (2150 deg F); a hydrogen atmosphere was used during brazing. After about five months of continuous operation, four or five holes developed on the floor of the muffle, and the muffle was removed from service. Analysis (visual inspection, x-ray spectrometry, and metallographic examination) supported the conclusion that the muffle failed by localized overheating in some areas to temperatures exceeding 1260 deg C (2300 deg F). The copper found near the holes had dripped to the floor from assemblies during brazing. The copper diffused into the nickel-base alloy and formed a grain-boundary phase that was molten at the operating temperature. The presence of this phase caused localized liquefaction and weakened the alloy sufficiently to allow formation of the holes. No recommendations were made.

Radiant tubes that failed prematurely in an ethylene cracking furnace were analyzed to determine the cause of their early demise. The tubes were made from austenitic heat-resistant steel and cracked along their longitudinal axis. New and used tubes were compared using scanning electron microscopy, energy dispersive x-ray spectrometry, and mechanical property testing. This provided critical information and revealed that improper coking and decoking had removed the protective oxide layer (Cr 2 O 3 ) that normally prevents coke deposits from forming on exposed surfaces. Without this layer, coke readily accumulates on the surface of the tubes, fueling carbon diffusion into the metal and a corresponding degradation in microstructure and loss of ductility at high temperatures.

An Incoloy 800H (UNS N08810) transfer line on the outlet of an ethane-cracking furnace failed during decoking of the furnace tubes after nine years in service. A metallographic examination using optical and scanning electron microscopy as well as energy-dispersive x-ray spectroscopy revealed that the failure was due to sulfidation. The source of the sulfur in the furnace effluent was either dimethyl disulfide, injected into the furnace feed to prevent coke formation and carburization of the furnace tubes, or contamination of the feed with sulfur bearing oil.

An investigation was conducted to determine why 16 out of 139 pipe bends cracked during hot induction bending. The pipe conformed to API 5L X65 PSL2 line pipe standards and measured 1016 mm (40 in.) in diam with a wall thickness of 18.5 mm. A metallurgical cross section was removed along a crack on the extrados to document the crack morphology using optical microscopy. In addition to cracking, golden-yellow streaks were visible at the extrados, and the composition was examined using scanning electron microscopy with energy dispersive spectroscopy. Based on the results, investigators concluded the pipe was contaminated with copper at the mill were it was produced.

Failure of a pilot scale test melter resulted from severe overheating of an Inconel 690 (690) jacketed molybdenum electrode. Extreme temperatures were required to melt the glass during this campaign because the feed material contained a very high waste loading. Metallurgical evaluation revealed the presence of an alloy containing nickel and molybdenum in several ingots found on the bottom of the melter and on a drip which had solidified on the electrode sheath. This indicates that a major portion of the electrode assembly was exposed to a temperature of at least 1317 deg C, the nickel/molybdenum eutectic temperature. Small regions on the end of the 690 sheath showed evidence of melting, indicating that this localized region exceeded 1345 deg C, the melting point of 690. In addition to nickel, antimony was found on the grain boundaries of the molybdenum electrode. This also contributed to the failure of the electrode. The source of the antimony was not identified but is believed to have originated from the feed material. Metallurgical evaluation also revealed that nickel had attacked the grain boundaries of the molybdenum/tungsten drain valve. This component did not fail in service; however, intergranular attack led to degradation of the mechanical properties, resulting in the fracture of the drain valve tip during disassembly. Antimony was not observed on this component.

This paper describes the superplastic characteristics of shipbuilding steel deformed at 800 °C and a strain rate less than 0.001/s. After the superplastic deformation, the steel presents mixed fractures: by decohesion of the hard (pearlite and carbides) and ductile (ferrite) phases and by intergranular sliding of ferrite/ferrite and ferrite/pearlite, just as it occurs in stage III creep behavior. The behavior is confirmed through the Ashby-Verrall model, according to which the dislocation creep (power-law creep) and diffusion creep (linear-viscous creep) occur simultaneously.

Hydrogen embrittlement is the brittleness affecting copper and copper alloys containing oxygen which develops during heat treatment at temperatures of about 400 deg C (752 deg F) and above in an atmosphere containing hydrogen. The phenomenon of hydrogen embrittlement of copper and its alloys is illustrated by examples from practice and reference is made to data from recent publications on the subject. Embrittlement due to this cause can only be identified by microscopic examination because other modes of failure in copper; e.g., from heat cracking, mechanical overload, the formation of low melting point eutectics or corrosion; show a similar appearance when investigated on a macroscopic scale.

The rear wall tube section of a boiler that had been in service for approximately 38 years was removed and examined. Visual examination of the tube revealed a small bulge with a through-wall crack. Metallography showed that the microstructure of the bulged area consisted of a few partially decarburized pearlite colonies in a ferrite matrix. The microstructure remote from the bulged area consisted of pearlite in a ferrite matrix. EDS analysis of internal deposits on the tube detected a major amount of iron, plus trace amounts of other elements. The evidence indicated that the bulge and crack in the tube resulted from hydrogen damage. Examination of the remaining water circuit boiler tubing using nondestructive techniques and elimination of any heavy deposit buildup was recommended.