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Metal-induced embrittlement of materials
Metal-Induced Embrittlement of Materials S. P. Lynch Aeronautical Research Laboratory, Defence Science & Technology Organisation, Department of Defence, Melbourne, Victoria 3207, Australia The phenomena of liquid-metal embrittlement (LME) and solid-metal-induced embrittlement (SMIE) of materials are reviewed, emphasizing metallographic, fractographic, and other characteristics of embrittlement that are relevant to failure analysis and prevention. Structural-metal/metal-environment combinations for which there are potential risks of failure, variables affecting susceptibility to embrittlement, and sources of embrittling metals are outlined. The mechanisms and kinetics of embrittlement are then discussed-including recent data for SMIE that show that rates of cracking are often the same for different materials and embrittling elements at the same homologous temperatures, T/Tm, where Tm is the melting temperature (K) of the embrittling element.
INTRODUCTION
cerned with identifying the minimum temperatures likely to produce SMIE, are then discussed. Before these aspects are considered, structural-metal/metal-environment combinations for which there are potential risks of failure, variables affecting susceptibility to embrittlement, and sources of embrittling metal environments in industry are outlined.
The resistance of materials to the initiation and growth of cracks is often considerably lower in certain liquid-metal, solid-metal, and metal-vapor environments than in inert environments, although bulk properties such as yield stresses and workhardening rates are not affected [1, 2]. Failure of components because of liquid-metal embrittlement (LME) and solid-metal induced embrittlement (SMIE) is less c o m mon than failures caused by other processes, such as fatigue, hydrogen-embrittlement (HE), and stress-corrosion cracking (SCC), but a significant number of industrial failures caused by metal-induced embrittlement (MIE) do occur. Failures arise partly because there is a lack of awareness of the phenomenon and partly because there is a lack of knowledge of the conditions under which it can occur, In the present article, metallographic, fractographic, and other techniques for identifying failures caused by LME/SMIE are described. Mechanisms and kinetics of embrittlement, including recent work con-
OCCURRENCE OF METAL-INDUCED EMBRITTLEMENT Embrittlement may occur during fabrication (e.g., hot-working processes) or during service, providing that there are tensile stresses above a threshold value and intimate contact, i.e., no intervening oxide or other films, between the structural and erabrittling metals. Embrittling environments for some common structural materials are listed in Table 1. Comprehensive surveys of embrittling and (possible) nonembritfling systems can be found in references [3, 4]. 279
©Commonwealth of Australia
MATERIALS CHARACTERIZATION 28:279-289 ~1992)
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Table 1 Examples of Embrittling Metal
Environments for Some Structural Materials Structural material
High-strength Martensitic steels "y-Stainless steels Titanium alloys Aluminium alloys Copper alloys Zirconium alloys Nickel alloys Magnesium alloys
Embrittling environments a Hg, In, Sn, Pb, Cd, Zn, Li, Cu Zn, Cu, Li Hg, Cd, Ag, Au Hg, Ga, In, Sn, Pb, Cd, Zn, Na Hg, Ga, Bi, Zn, Li, Sn, Pb, In Hg, Cd (Cd-Cs), Zn Hg, In, Li, Zn, Ag Na, K, Rb, Cs, Zn
a Alloys containing these elements are usually embrittling, but the presence of other elements can increase or decrease embrittlement. The degree of embrittlement varies widely and d e p e n d s on n u m e r o u s variables.
VARIABLES INFLUENCING SUSCEPTIBILITY TO EMBRITTLEMENT
ALLOY STRENGTH/COMPOSITION/ MICROSTRUCTURE The degree of embrittlement is generally greater in higher-strength materials, as is the case for embrittlement by hydrogen and by other environments. The effects of microstructure on susceptibility to LME often depend on the material and testing conditions, but some effects common to a number of materials include (1) increases in the degree of LME with increasing grain size, (2) greater susceptibility to LME for materials with planar-slip rather than wavy-slip characteristics, and (3) decreases in the degree of LME with increasing cold work [1, 2]. Slip bands in materials witl~ large grain sizes, and planar rather than wavy slip, probably cause greater stress concentrations at slip-band/grain-boundary intersections (for a given strain), so that lower strains are required for crack initiaLion at these sites. Increasing cold work produces a more fibrous grain structure that results in a more tortuous intercrystalline crack path and, hence, a greater resistance to intercrystalline cracking for ma-
terial stressed parallel to the fiber axis. For high-strength steels, susceptibility to LME is increased by the presence of segregated impurity elements, e.g., Sb and Sn, to grain boundaries [5]. However, the effects of impurity segregation to grain boundaries on intercrystalline LME, in general, will depend on the segregated and liquidmetal species involved. TEMPERATURE The degree of LME is often greatest at the embrittler melting-point, T,,, with embrittlement becoming less severe with increasing temperature, probably because stress relaxation at potential crack-initiation sites occurs more readily at higher temperaLures. The range of temperature o v e r which LME occurs varies widely (10100°C), and depends on the embrittlement couple and the testing conditions [1, 2]. LME can occur below the embrittler melting point if a eutectic forms between the liquid metal and a component of the structural metal. SMIE has been observed at temperatures, T, as low ~0.38 T/T,, (T and Tm K) for embrittlement of titanium alloys by silver and gold [6], and would generally be expected to occur at temperatures above 0.5 TIT,n, with higher rates of cracking at higher temperatures, as discussed in detail in a later section. SPECIMEN GEOMETRY, STRESS-MODE, STRAIN-RATE The effects of specimen geometry, stress state, and stress mode on susceptibility to MIE have not been investigated in detail, but limited data suggest that their effects are similar to those observed for HE. For example, embrittlement may sometimes be observed in notched specimens but not in unnotched specimens, plane-strain conditions are more severe than plane-stress conditions, mode-I loading is more severe than mode-III loading, and cyclic loading may be more severe than sustained loading [1, 2]. The degree of LME is also often greater at higher strain rates, especially in
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materials that exhibit increases in yield strength with increasing strain rate. For solid-metal environments, however, strain rates that produce rates of crack growth greater than rates of transport of embrittling atoms to crack tips do not produce embrittlement.
lead can produce subsequent cracking during stress relieving [9]. Fine particles of low-melting point metals have been added to lubricants to facilitate machining [10], and residues could cause subsequent embrittlement.
SOURCES OF EMBRITTLING METAL
METAL VACUUM SEALS
ENVIRONMENTS IN INDUSTRY
Very soft metals such as indium, plated on Inconel gaskets to ensure a good seal between components in vacuum systems, have been known to produce embrittlement [11].
Some examples of the diverse sources of embrittling metals that have led to failures are listed in the following. Very small amounts (0.1 g) of embrittling metals can produce extensive cracking (crack lengths 10-100 mm) by LME, although cracks can "run out of" liquid metal if the amount is very small. COATI NGS Cadmium, zinc, and other low-melting point metals are commonly applied to steels and other materials for protection against corrosion. Metals are also coated for other reasons, e.g., electrical contacts are sometimes silver plated to reduce electrical resistance. LME can occur during application of coatings by "hot-dipping" if residual stresses are present, or LME/SMIE can occur during service if temperatures and stresses are sufficiently high and coatings are in intimate contact with the substrate. Thin films of "inert" material are sometimes deposited between coatings and substrates to prevent embrittlement from occurring. SOLDERS, BRAZES, WELD-METALS LME can occur during joining of components by soldering, brazing, or welding if one of the constituents of the joining material embrittles the base material and if residual or assembly stresses are sufficiently high [7, 8]. METAL LUBRICANTS Lead is sometimes used as a lubricant during cold forming of steels, and a residue of
O V E R H E A T EBEARINGS D Steel locomotive axles sometimes fail because of LME by copper-containing bearing materials that have overheated because of loss of oil lubrication [12].
NUCLEAR FISSION PRODUCTS In the nuclear industry, the effects of alkali-metal coolants and liquid/solid metal fission products on the mechanical properties of structural materials have been extensively studied [1, 13]. Failures of zircalloy nuclear-fuel cladding caused by embrittlement by cadmium-caesium fission products have been observed.
ACCIDENTAL CONTAMINATION Spillages of mercury from thermometers, and spillage of a consignment of gallium, have occurred in aircraft, necessitating replacement of c o n t a m i n a t e d sections. Splashes of zinc-based paints on stainlesssteel components have caused embrittlement during subsequent welding of the steel [14]. Cracking of stainless steel piping caused by contamination by liquid zinc (from galvanized steel during a fire) in a chemical plant at Flixborough, United Kingdom, resulted in leakage of hydrocarbons that may have exacerbated the disastrous explosion that occurred [15].
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INTERNAL SOURCES
Small amounts (-0.1%) of low-melting point metals, e.g. Pb and Bi, are sometimes deliberately added to materials to improve their machining behavior ("free-machining alloys"). These additions result in the presence of low-melting point inclusions, especially at grain boundaries, and these materials are brittle at elevated temperatures [16, 17]. Trace (ppm) impurities segregated to grain boundaries can also cause ofembrittlement" sbF°r example, segregation P, Sn, and to prior-austenite grain boundaries in high-strength steels causes temper embrittlement, which could also be termed "metal-induced embrittlement" for the embrittling elements that are metallic . Similarly, trace amounts of bismuth, lead and other elements can cause intergranular embrittlement in aluminum alloys and nickel-base superalloys [18]. Hot shortness, e.g., caused by the presence of copper or tellurium in steels, is also a form of metal-induced embrittlement. In high-strength A1-Li base alloys, which are currently being developed for use in aircraft structures, traces (several ppm) of sodium, potassium, and possibly other impurities, have deleterious effects on fracture properties [19, 20]. Fracture toughness is unacceptably low in some alloys, and exposure of stressed specimens at elevated temperatures (>60°C) in inert environments can produce subcritical cracking, with higher rates of cracking at higher temperatures (cracking rates - 10 mm/h have been observed at 180°C for several alloys [21]). These effects are probably largely caused by metal-induced embrittlement, although other effects are possibly also involved. METALLOGRAPHIC AND FRACTOGRAPHIC CHARACTERISTICS OF LME/SMIE CRACK PATH AND FRACTURE SURFACE APPEARANCE LME and SMIE produce brittle intercrystalline and transcrystalline (cleavagelike)
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h ~ FIG. 1. SEM fractographs of fracture surfaces of (a) Ni-200crackedin liquidlithiumat 210°C;(b) 7075-T651 aluminiumalloy crackedin solid indium at 20°C;and (c) a mill-annealedTi-6%A1-4%V alloy cracked in liquid mercuryat 20°C, showingintercrystalline(I) and cleavagelike(C) areas. Flutes (F) are also visiblein (c). fracture surfaces in a wide range of normally ductile materials (Fig. 1). Brittle intercrystalline fracture predominates in fcc and bcc materials, but substantial amounts of cleavagelike fracture, especially on basal planes, often occur in hcp materials. The fracture path and appearance produced by SMIE are generally similar to those produced by LME, although the em-
Metal-Induced Embrittlement
283
FIG. 3. TEM replica of brittle intercrystalline fracture surface of D6ac steel cracked in liquid mercury at 20°C, showing tear ridges (T) and small, shallow dimples (D), which were not clearly resolved by SEM.
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Fic. 2. SEM fractographs of fracture surface of D6ac steel tempered at various temperatures and cracked in solid indium at 140°C, showing (a) dimpled fractures along martensite lath boundaries for a specimen tempered at 290°C (~54HRC), (b) relatively smooth fractures along prior-austenite grain boundaries for a specimen tempered at 400°C (~51HRC), and (c) dimpled fracture along prior-austenite grain boundaries for a specimen tempered at 650°C (-41HRC). Overload fractures were transcrystalline and dimpled for all these conditions, and dimples were larger and deeper than those produced on MIE fractures.
b r i t t l e m e n t of t i t a n i u m alloys b y c a d m i u m is o n e e x c e p t i o n - - L M E p r o d u c e s p r e d o m inately intercrystalline fracture, b u t SMIE p r o d u c e s m i x t u r e s of intercrystalline a n d transcrystalline fractures. T h e fracture p a t h a n d a p p e a r a n c e in s o m e materials dep e n d o n the m i c r o s t r u c t u r e . For e x a m p l e , the fracture characteristics of a h i g h -
s t r e n g t h t e m p e r e d martensitic D6ac steel c r a c k e d in l i q u i d - m e r c u r y [22] or in solid i n d i u m [2l] e n v i r o n m e n t s d e p e n d e d on the t e m p e r i n g t e m p e r a t u r e of the steel (Fig. 2). S h a l l o w d i m p l e s , tear ridges, a n d slip lines are p r e s e n t o n "brittle" intercrystalline a n d cleavagelike fracture surfaces in m o s t materials. T h e s e features are s o m e times r e s o l v e d b y SEM [Fig. 2(a, c)] but are often o n l y a p p a r e n t w h e n fracture surfaces are e x a m i n e d b y h i g h - r e s o l u t i o n transmission electron m i c r o s c o p y f r a c t o g r a p h i c t e c h n i q u e s (Figs. 3, 4). LME a n d SMIE in h i g h - s t r e n g t h materials are u s u a l l y not associated w i t h a n y m a c r o s c o p i c signs of plasticity, b u t sensitive t e c h n i q u e s reveal small plastic z o n e s a r o u n d cracks. ]in low-
. . . . . . . . 23t F,c. 4. TEM replica of cleavagelike fracture surface of Ti-6%A1-4%V alloy cracked in liquid mercury at 20°c, showing tear ridges and small dimples that were not clearly resolved by SEM. Flutes (F) are also evident.
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b FIG. 5. Light micrograph of a polished and etched section showing the slip pattern around a cleavagelike crack in the interior of a solution treated A1-Zn-Mg alloy single crystal cracked in a liquid bismuth alloy environment at 60°C, showing extensive slip, particularly on planes intersecting cracks. The slip pattern on the specimen side surface for the same crack is shown in the inset . . . .
strength materials, on the other hand, large plastic zones, with slip occurring par . ticularly on planes intersecting crack fronts, are often associated with brittle intercrystalline and cleavagelike cracks (Fig.
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5). DETECTION OF EMBRITTLING METALS AFTER FRACTURE Globules or films of embrittling metal a r e usually evident on fracture surfaces produced by LME, and can sometimes be detected with the unaided eye or by using a binocular microscope. For example, examination of high-strength (-44 HRC) martensitic steel aircraft engine nuts that had fractured into two or three pieces (Fig. 6) showed that some fracture surfaces exhibited distinctive blue and gold colors that are characteristic of oxidized cadmium. Energy-dispersive x-ray analysis (EDXA) of fracture surfaces confirmed that cadmium was present in substantial amounts. This failure occurred because cadmium-plated, high-strength steel nuts were inadvertently used instead of silver-plated steel nuts in part of the engine that experienced temperatures above the melting point of cadmium! Metallographic sections through
Fnc. 6. (a) Macroscopic appearance of a fractured high-strength cadmium-plated 4140 steel nut that was inadvertently used on bolts joining ducts carrying hot (500°C) air from the compressor of a military jet engine, and (b) SEM fractograph of fracture surface showing brittle intercrystalline facets.
LME cracks also often show films of solidified embrittling metal within cracks. Alloying between solid and liquid metals may sometimes occur prior to, or after fracture, and intermetallic compounds or other reaction products may be present on fracture surfaces. Volatile elements, such as mercury, could have evaporated if ternperatures were sufficiently high. If components have been only partially cracked by LME, extending the crack at room temperature may result in fracture through the thin layer of solidified embrittling metal within cracks, producing a fracture-surface appearance characteristic of the (con-
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FIG. 7. SEM fractograph of the fracture surface of a 7475 aluminium alloy partially cracked in liquid indium at 170°Cand then subsequentlybroken open at 20°C. The dimple/veinpattern on the fracturesurface results from ductile cracking through the solidified indium layer in the crack.
strained) embrittling metal, e.g., a dimpled fracture surface for a ductile material (Fig. 7). The presence of the embrittling metal on fracture surfaces and within cracks produced by SMIE is often not readily detected because the layer of embrittling metal is probably in the range -0.01-0.1 ~m thick [6, 22], and the sampling depth of EDXA is 2-5 ~m. However, for SMIE of aluminium alloys by indium [23], and of steels and titanium alloys by cadmium [24, 25], the embrittling metals have been detected on fracture surfaces by EDXA after long analysis times. Small islands of a film can also sometimes be observed by secondary-electron mode (SEM) (Fig. 8). Embrittling elements may be more readily detected by EDXA on sections through secondary cracks than on fracture surfaces, because the layers would be approximately normal to the surface [24]. Thin layers of embrittling metals are readily detected on fracture surfaces by auger electron spectroscopy (AES), x-ray Photoelectron spectroscopy (XPS), and secondary-ion mass spectroscopy (SIMS), because the sampling depths of these techniques are 1-2 nm. However, removing films produced by atmospheric contamination may sometimes be necessary to detect underlying embrittling metal films [6].
FIG. 8. SEM fractograph of fracture surface of a Ti6% A1-4%V alloycrackedin solid cadmiumat 295°C, showingdark islands (which are probably films of cadmium)on fluted intercrystalline(I) and cleavagelike (C) areas. DIFFERENTIATING BETWEEN MIE AND HE/SCC
The appearance of fracture surfaces produced by LME and SMIE are often remarkably similar to those produced by HE and SCC [22, 26-28], and, hence, it is not possible to distinguish between these fracture modes solely on the basis of fracture surface appearance. Differentiation between failures by SMIE and HE/SCC in electroplated high-strength steels (where hydrogen is generated during the plating process) may require careful analysis for thin layers of plating material on fracture surfaces. A knowledge of the plating and baking procedures, and of the in-service operating conditions (temperature, environment, applied stresses) would also be helpful. For example, nonporous plating--which would inhibit removal of hydrogen during baking--would be a factor in favor of HE as a cause of failure. The presence of a corrosive environment during service, corrosion pits at crack-initiation sites, and corroded fracture surfaces would point to SCC as a probable cause of failure. A knowledge of the operating ternperature of the component would also help distinguish between HE and SMIE because susceptibility to HE is usually greatest near ambient temperature, whereas susceptibility to SMIE is greatest just below the embrittler's melting temperature.
286
The extent of brittle intercrystalline or transcrystalline cracking may also indicate whether cracking has occurred by HE or SMIE. The maximum crack lengths reported for SMIE of steels and titanium alloys by cadmium [24, 25], and 7075-T651 aluminium alloys by indium [23], are - 3 r a m at T/Tm -0.95 and - 1 . 5 mm at T/ Tm -0.7. Thus, brittle cracking to depths greater than these values would suggest that cracking has probably occurred by a process other than SMIE. However, crack growth could initially occur by SMIE and then continue by another process (fatigue, HE, SCC) that produced a similar fracture surface appearance to SMIE.
MECHANISMS OF METAL-INDUCED EMBRITTLEMENT EMBRITTLEMENT DUE TO ADSORPTION
Adsorption of environmental metal atoms at surface stress concentrations and at crack tips is the most common cause of embrittlement, particularly w h e n solid and liquid metals have low mutual solubilities and do not form intermetallic compounds [1, 2]. Removal or rupture of oxide or other films on surfaces is necessary to enable adsorption to occur, but diffusion of environmental atoms into solids is not involved, Thus, preexposure of unstressed specimens to liquid-metal environments, then removing the liquid metal prior to stressing, does not produce embrittlement in such systems even if some diffusion has occurred. Adsorption of environmental metal atoms on surfaces (crack tips) apparently results in changes of electron density and a consequent weakening of interatomic bonds between substrate atoms within a few atomic distances of the adsorbate. However, the details are not well understood and, hence, predicting which liquid metals embrittle particular solid metals is presently not possible from first principles. Workers studying LME in the 1960s and early 1970s concluded that adsorption low-
S. P.
Lynch
ered the stress required for "decohesion" at crack tips so that atomically brittle fractures were produced [29]. However, subsequent detailed metallographic and fractographic studies of LME, e.g., showing dimples on fracture surfaces and slip on planes intersecting cracks, suggested that adsorption facilitates the injection of dislocations at crack tips so that crack growth occurs by a more localized microvoid-coalescence process than occurs in inert environments, as discussed in detail elsewhere [22, 26-28]. Similarities between MIE and HE/SCC, and other observations, suggest that an adsorption-induced localized slip process is also applicable to HE and SCC for some materials and environments [22, 26-28]. EMBRITTLEMENT ASSOCIATED WITH DIFFUSION AND OTHER PROCESSES
Diffusion of environmental metal atoms into materials along grain boundaries ("grain boundary wetting")commonly occurs in systems in which the liquid metal is relatively insoluble in the solid, but the solid has an appreciable solubility in the liquid [30, 31]. For example, exposure of unstressed aluminium and zinc polycrystals to liquid gallium results in diffusion of gallium along grain boundaries and the formation of a thin, gallium-rich, liquid film along the grain boundaries; subsequent stressing causes the material to disintegrate into individual grains. This effect is useful for studying grain morphologies and grain-boundary cavitation caused by prior deformation [32, 33]. LME of stainless steels by liquid zinc is also sometimes associated with diffusion along grain boundaries; in this case, the formation of a NiZn compound produces local depletion of nickel at grain boundaries, thereby resulting in an austenite-to-ferrite transformation and an associated volume expansion that generates high stresses and cracking [30]. In liquid alkali-metal environments, which are commonly encountered in the nuclear industry, transfer of interstitial ele-
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ments, such as carbon, nitrogen and oxygen, between liquids and solids can occur, thereby producing changes in mechanical properties [1, 13]. Reactions between some liquid and solid metals produce easily cracked, brittle intermetallic compounds or selective dissolution of alloy phases/grain boundaries/slip bands, thereby facilitating crack initiation. In some systems, e.g., stainless steels-zinc, aluminium alloysgallium, different processes may be responsible for embrittlement under different conditions, e.g., adsorption may be responsible at one temperature/strain rate, and diffusion may be involved at another temperature/strain rate.
as -10% of the critical K for cracking in inert environments (Fig. 9). Transport of liquid metal to crack tips occurs by capillary flow, and a limiting rate of transport possibly controls cracking kinetics at high crack velocities. The different "plateau" velocities for different materials (and for aluminium alloys in different conditions [34] and observations [36] that threshold K values for LME of aluminium alloys depend on the experimental conditions, suggest that other factors also affect cracking kinetics. However, these effects are not well understood. SMIE
KINETICS OF LME AND SMIE LME Delayed failure of materials can occur in liquid-metal environments because of slow crack initiation processes (which are not well understood), but crack growth generally occurs rapidly. Crack velocities in the range of 10-100 mm/s have been reported for subcritical cracking of aluminium and titanium alloys in liquid mercury [34, 35]. The velocities are not affected by stress-intensity factor (K) values except near the threshold K which can be as low
Crack velocities during SMIE are considerably less than those during LME, and they decrease with decreasing temperature. Studies of SMIE of titanium alloys by cadmium [25], and 7075-T651 aluminium alloy by indium [23], showed that crack velocities were unaffected by K values above a threshold value, and decreased with increasing crack length (i.e., increasing distance from the source of embrittling atoms) (Fig. 10). At longer crack lengths, crackgrowth rates were greater when tests were conducted in vacuum rather than in air, presumably because oxidation of fracture surfaces and crack tips inhibited transport and adsorption processes when tests were done in air. For short crack lengths, the
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FIG. 10. Crack length versus time data for 7075 aluminium and titanium alloys cracked in solid-metal env i r o n m e n t s at different h o m o l o g o u s temperatures, T/ Tm w h e r e Tm is the melting point of the embrittler and T is the testing temperature, K.
S. P. Lynch
288 embrittling metal itself probably p r e v e n t e d access of air into cracks, The previous data can be explained on the basis that crack-growth rates are controlled by transport of embrittling atoms from their source (surface or notch root) to the crack tip by a surface diffusion process, Values of surface diffusion coefficients, Ds, for SMIE at various t e m p e r a t u r e s have b e e n calculated from crack length (x) versus time (t) data using the standard random-walk diffusion equation (Ds = x2/2t), for various materials and environments, Plots of log D~ versus Tm/T (where Tm is the melting t e m p e r a t u r e of the embrittling elem e n t and T is the testing temperature) s h o w that a linear Arrhenius relationship can be fitted to most of the data (Fig. 11) [23]. C o m p a r i s o n of these data with "clean" surface self-diffusion data, obtained from the kinetics of sintering and other processes [37], s h o w e d that the surface selfdiffusion values were about one to two orders of m a g n i t u d e greater than D~ values obtained from SMIE data [23]. H o w e v e r , the values of D~ calculated from the SMIE data are u n d e r e s t i m a t e d , probably b y about an order of magnitude, because the diffusion distances were taken as the
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projectedcrack lengths. The actual diffusion distances will d e p e n d on the r o u g h n e s s of fracture surfaces on a microscopic and submicroscopic scale, and are probably three to four times the projected crack lengths w h e n crack g r o w t h produces dimpled intercrystalline fracture surfaces. Thus, corrected values of Ds for SMIE data would be within an order of m a g n i t u d e of values for surface self-diffusion obtained by other techniques. The data for the effect of t e m p e r a t u r e on surface diffusion/SMIE kinetics suggest that cracking caused by an external source of embrittling metal could occur at significant rates (>1 m m in 1 year) at homologous t e m p e r a t u r e s >0.5 T/Tm. SMIE of high-strength steels and titanium alloys by c a d m i u m has occasionally been observed at 38°C (T/Tm -0.52), but does not appear to be a c o m m o n problem, probably because electroplated c a d m i u m is not usually in intimate contact with the substrate. H o w e v e r , such contact could develop if intervening (oxide) films were r u p t u r e d by the applied stress or by local pressure, and if embrittling atoms were sufficiently mobile. Furthermore, small cracks p r o d u c e d by SMIE at low t e m p e r a t u r e s could initiate more extensive cracking by HE, SCC, or fatigue, leading to unstable fracture. SMIE of titanium alloys by silver and gold [6] has been observed at h o m o l o g o u s t e m p e r a t u r e s - 0 . 3 8 at rates of crack growth that are considerably larger than w o u l d be expected on the basis of the other SMIE data. Possible explanations for this a n o m o l o u s behavior are (1) higher surfacediffusion rates than " n o r m a l " because of impurities in the embrittling elements [37], and (2) a contribution to cracking from HE. L o w e r rates of crack g r o w t h than predicted from surface-diffusion data have been observed for SMIE of lower strength/higher t o u g h n e s s materials than those discussed earlier, p r e s u m a b l y because some other (slower) process controls the cracking kinetics [23]. SMIE at a given t e m p e r a t u r e could occur at m u c h faster overall rates for material containing closely spaced internal sources
Metal-Induced Embrittlement
of embrittling atoms than for material in contact with an external source, because diffusion distances to crack tips would never be greater than the spacing of sources for the former. Extrapolation of the SMIE surface-diffusion data to lower temperatures suggests that SMIE could possibly occur at significant rates at homologous temperatures as low as -0.3 T/Tm when the spacing of sources is - 1 p~m. The association of embrittling elements with other elements would also be expected to influence diffusion kinetics. For example, the addition of bismuth to A1-Mg alloys prevents embrittlement by sodium, possibly because a Bi-Na compound is formed, thereby preventing diffusion of sodium t o c r a c k tips [38].
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Vol. 11, Failure Analysis Metals Park, Ohio (1986),
and Prevention, ASM, pp. 715-727.
13. C. F. Old, I. Nucl. Mater. 92:2-25 (1980). 14. T. Shinohara and K. Matsumoto, Corr. Sci. 22:723-737 (1982). 15. D. Elliot, Process Eng. 67-71 (1976). 16. D. Bhattacharya, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 367-398. 17. M. C. Roth, T. A. Jones, G. C. Weatherly, and w . A . Miller, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME Warrendale PA (1984), pp. 399-413. 18. M. Guttmann, B. Quantin, and Ph. Dumoulin, Metal Sci. 17:123-140 (1983). 19. D. N. Fager, M. V. Hyatt, and H. T. Diep, Scripta Metall. 20:1159-1164 (1986). 20. D. Webster; Aluminium-Lithium Alloys 1II, Inst. of Metals., London, pp. 602-609 (1986). 21. s. p. Lynch, Mater. Sci. Eng. A136:45-57 (1991). 22. S. P. Lynch, Acta Metall. 32:79-90 (1984). 23. S. P. Lynch, Mater. Sci. Eng. A108:203-212 (1989).
References 1. M. H. Kamdar (ed.), Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984). (See Contents page and Index for subjects not referenced separately.) 2. N. S. Stoloff, Atomistics of Fracture, Plenum Press, New York, (1983), pp. 921-949. 3. F. A. Shunk and W. R. Warke, Scripta Metall. 8:519-526 (1974). 4. M. G. Nicholas, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 27-50. 5. S. Dinda and W. R. Warke, Mater. Sci. Eng. 24:199-208 (1976). 6. R. E. Stoltz and R. H. Stulen, Corrosion 35:165169 (1979). 7. M. G. Nicholas and C. F. Old, Weld. Met. Fabr. 47:527-531 (1979). 8. R. H. Edwards, N. I. Kilpatrick, E. G. Bradey, and J. Szczur, Mater. Forum 9;202-208, (1986). 9. M. H. Kamdar, Proc. Conf. on Environmental Degradation of Engineering Materials, Virginia Polytechnic Inst., Blacksburg, VA (1977), pp. 235-243. 10. J. R. Pickens and A. R. C. Westwood in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 51-64. 11. A. P. Druschitz and P. Gordon, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 285-316. 12. G. F. Vander Voort, Metals Handbook, 9th Ed.
24. D. N. Fager and W. F. Spurr, Corrosion 2:7:72-76 (1971). 25. D. N. Fager and W. F. Spurr, Corrosion 26:409419 (1970). 26. S. P. Lynch, Acta Metall. 29:325-340 (1981). 27. S. P. Lynch, J. Mater. Sci. 20:3329-3338 (1985). 28. S. P. Lynch, J. Mater. Sci. 21:692-704 (1986). 29. M. H. Kamdar, Prog. Mater. Sci. 15:289-374 (1973). 30. C. F. Old, Metal Sci., 14:433-440 (1980). 31. G. H. Bishop, Trans. Met. Soc. A1ME 242:13431351 (1968). 32. F. N. Rhines and A. B. Gokhale, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 265-276. 33. T. C. Reiley, Scripta Metall. 15:497-500 (1981). 34. M. O. Speidel, in The Theory of Stress Corrosion Cracking Alloys, NATO, Brussels, (1971), pp. 289344. 35. J. A. Feeney and M. J. Blackburn, in The Theory of Stress Corrosion Cracking Alloys, NATO, Brussels (1971), pp. 355-398. 36. D. A. Wheeler and R. G. Hoagland, Scripta Metall. 20:1433-1438 (1986). 37. H. P. Bonzel, in Surface Physics of Materials, Vol. II, Academic Press, New York, (1975), pp. 279338. 38. D. E. J. Talbot and C. E. Ransley, Metai'l. Trans. A, 8A:1149-1154 (1977).
AcceptedJanuary 1992.
INTRODUCTION
cerned with identifying the minimum temperatures likely to produce SMIE, are then discussed. Before these aspects are considered, structural-metal/metal-environment combinations for which there are potential risks of failure, variables affecting susceptibility to embrittlement, and sources of embrittling metal environments in industry are outlined.
The resistance of materials to the initiation and growth of cracks is often considerably lower in certain liquid-metal, solid-metal, and metal-vapor environments than in inert environments, although bulk properties such as yield stresses and workhardening rates are not affected [1, 2]. Failure of components because of liquid-metal embrittlement (LME) and solid-metal induced embrittlement (SMIE) is less c o m mon than failures caused by other processes, such as fatigue, hydrogen-embrittlement (HE), and stress-corrosion cracking (SCC), but a significant number of industrial failures caused by metal-induced embrittlement (MIE) do occur. Failures arise partly because there is a lack of awareness of the phenomenon and partly because there is a lack of knowledge of the conditions under which it can occur, In the present article, metallographic, fractographic, and other techniques for identifying failures caused by LME/SMIE are described. Mechanisms and kinetics of embrittlement, including recent work con-
OCCURRENCE OF METAL-INDUCED EMBRITTLEMENT Embrittlement may occur during fabrication (e.g., hot-working processes) or during service, providing that there are tensile stresses above a threshold value and intimate contact, i.e., no intervening oxide or other films, between the structural and erabrittling metals. Embrittling environments for some common structural materials are listed in Table 1. Comprehensive surveys of embrittling and (possible) nonembritfling systems can be found in references [3, 4]. 279
©Commonwealth of Australia
MATERIALS CHARACTERIZATION 28:279-289 ~1992)
280
S. P. Lynch
Table 1 Examples of Embrittling Metal
Environments for Some Structural Materials Structural material
High-strength Martensitic steels "y-Stainless steels Titanium alloys Aluminium alloys Copper alloys Zirconium alloys Nickel alloys Magnesium alloys
Embrittling environments a Hg, In, Sn, Pb, Cd, Zn, Li, Cu Zn, Cu, Li Hg, Cd, Ag, Au Hg, Ga, In, Sn, Pb, Cd, Zn, Na Hg, Ga, Bi, Zn, Li, Sn, Pb, In Hg, Cd (Cd-Cs), Zn Hg, In, Li, Zn, Ag Na, K, Rb, Cs, Zn
a Alloys containing these elements are usually embrittling, but the presence of other elements can increase or decrease embrittlement. The degree of embrittlement varies widely and d e p e n d s on n u m e r o u s variables.
VARIABLES INFLUENCING SUSCEPTIBILITY TO EMBRITTLEMENT
ALLOY STRENGTH/COMPOSITION/ MICROSTRUCTURE The degree of embrittlement is generally greater in higher-strength materials, as is the case for embrittlement by hydrogen and by other environments. The effects of microstructure on susceptibility to LME often depend on the material and testing conditions, but some effects common to a number of materials include (1) increases in the degree of LME with increasing grain size, (2) greater susceptibility to LME for materials with planar-slip rather than wavy-slip characteristics, and (3) decreases in the degree of LME with increasing cold work [1, 2]. Slip bands in materials witl~ large grain sizes, and planar rather than wavy slip, probably cause greater stress concentrations at slip-band/grain-boundary intersections (for a given strain), so that lower strains are required for crack initiaLion at these sites. Increasing cold work produces a more fibrous grain structure that results in a more tortuous intercrystalline crack path and, hence, a greater resistance to intercrystalline cracking for ma-
terial stressed parallel to the fiber axis. For high-strength steels, susceptibility to LME is increased by the presence of segregated impurity elements, e.g., Sb and Sn, to grain boundaries [5]. However, the effects of impurity segregation to grain boundaries on intercrystalline LME, in general, will depend on the segregated and liquidmetal species involved. TEMPERATURE The degree of LME is often greatest at the embrittler melting-point, T,,, with embrittlement becoming less severe with increasing temperature, probably because stress relaxation at potential crack-initiation sites occurs more readily at higher temperaLures. The range of temperature o v e r which LME occurs varies widely (10100°C), and depends on the embrittlement couple and the testing conditions [1, 2]. LME can occur below the embrittler melting point if a eutectic forms between the liquid metal and a component of the structural metal. SMIE has been observed at temperatures, T, as low ~0.38 T/T,, (T and Tm K) for embrittlement of titanium alloys by silver and gold [6], and would generally be expected to occur at temperatures above 0.5 TIT,n, with higher rates of cracking at higher temperatures, as discussed in detail in a later section. SPECIMEN GEOMETRY, STRESS-MODE, STRAIN-RATE The effects of specimen geometry, stress state, and stress mode on susceptibility to MIE have not been investigated in detail, but limited data suggest that their effects are similar to those observed for HE. For example, embrittlement may sometimes be observed in notched specimens but not in unnotched specimens, plane-strain conditions are more severe than plane-stress conditions, mode-I loading is more severe than mode-III loading, and cyclic loading may be more severe than sustained loading [1, 2]. The degree of LME is also often greater at higher strain rates, especially in
281
Metal-Induced Embrittlement
materials that exhibit increases in yield strength with increasing strain rate. For solid-metal environments, however, strain rates that produce rates of crack growth greater than rates of transport of embrittling atoms to crack tips do not produce embrittlement.
lead can produce subsequent cracking during stress relieving [9]. Fine particles of low-melting point metals have been added to lubricants to facilitate machining [10], and residues could cause subsequent embrittlement.
SOURCES OF EMBRITTLING METAL
METAL VACUUM SEALS
ENVIRONMENTS IN INDUSTRY
Very soft metals such as indium, plated on Inconel gaskets to ensure a good seal between components in vacuum systems, have been known to produce embrittlement [11].
Some examples of the diverse sources of embrittling metals that have led to failures are listed in the following. Very small amounts (0.1 g) of embrittling metals can produce extensive cracking (crack lengths 10-100 mm) by LME, although cracks can "run out of" liquid metal if the amount is very small. COATI NGS Cadmium, zinc, and other low-melting point metals are commonly applied to steels and other materials for protection against corrosion. Metals are also coated for other reasons, e.g., electrical contacts are sometimes silver plated to reduce electrical resistance. LME can occur during application of coatings by "hot-dipping" if residual stresses are present, or LME/SMIE can occur during service if temperatures and stresses are sufficiently high and coatings are in intimate contact with the substrate. Thin films of "inert" material are sometimes deposited between coatings and substrates to prevent embrittlement from occurring. SOLDERS, BRAZES, WELD-METALS LME can occur during joining of components by soldering, brazing, or welding if one of the constituents of the joining material embrittles the base material and if residual or assembly stresses are sufficiently high [7, 8]. METAL LUBRICANTS Lead is sometimes used as a lubricant during cold forming of steels, and a residue of
O V E R H E A T EBEARINGS D Steel locomotive axles sometimes fail because of LME by copper-containing bearing materials that have overheated because of loss of oil lubrication [12].
NUCLEAR FISSION PRODUCTS In the nuclear industry, the effects of alkali-metal coolants and liquid/solid metal fission products on the mechanical properties of structural materials have been extensively studied [1, 13]. Failures of zircalloy nuclear-fuel cladding caused by embrittlement by cadmium-caesium fission products have been observed.
ACCIDENTAL CONTAMINATION Spillages of mercury from thermometers, and spillage of a consignment of gallium, have occurred in aircraft, necessitating replacement of c o n t a m i n a t e d sections. Splashes of zinc-based paints on stainlesssteel components have caused embrittlement during subsequent welding of the steel [14]. Cracking of stainless steel piping caused by contamination by liquid zinc (from galvanized steel during a fire) in a chemical plant at Flixborough, United Kingdom, resulted in leakage of hydrocarbons that may have exacerbated the disastrous explosion that occurred [15].
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S. P. Lynch
INTERNAL SOURCES
Small amounts (-0.1%) of low-melting point metals, e.g. Pb and Bi, are sometimes deliberately added to materials to improve their machining behavior ("free-machining alloys"). These additions result in the presence of low-melting point inclusions, especially at grain boundaries, and these materials are brittle at elevated temperatures [16, 17]. Trace (ppm) impurities segregated to grain boundaries can also cause ofembrittlement" sbF°r example, segregation P, Sn, and to prior-austenite grain boundaries in high-strength steels causes temper embrittlement, which could also be termed "metal-induced embrittlement" for the embrittling elements that are metallic . Similarly, trace amounts of bismuth, lead and other elements can cause intergranular embrittlement in aluminum alloys and nickel-base superalloys [18]. Hot shortness, e.g., caused by the presence of copper or tellurium in steels, is also a form of metal-induced embrittlement. In high-strength A1-Li base alloys, which are currently being developed for use in aircraft structures, traces (several ppm) of sodium, potassium, and possibly other impurities, have deleterious effects on fracture properties [19, 20]. Fracture toughness is unacceptably low in some alloys, and exposure of stressed specimens at elevated temperatures (>60°C) in inert environments can produce subcritical cracking, with higher rates of cracking at higher temperatures (cracking rates - 10 mm/h have been observed at 180°C for several alloys [21]). These effects are probably largely caused by metal-induced embrittlement, although other effects are possibly also involved. METALLOGRAPHIC AND FRACTOGRAPHIC CHARACTERISTICS OF LME/SMIE CRACK PATH AND FRACTURE SURFACE APPEARANCE LME and SMIE produce brittle intercrystalline and transcrystalline (cleavagelike)
~
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.
i
.
h ~ FIG. 1. SEM fractographs of fracture surfaces of (a) Ni-200crackedin liquidlithiumat 210°C;(b) 7075-T651 aluminiumalloy crackedin solid indium at 20°C;and (c) a mill-annealedTi-6%A1-4%V alloy cracked in liquid mercuryat 20°C, showingintercrystalline(I) and cleavagelike(C) areas. Flutes (F) are also visiblein (c). fracture surfaces in a wide range of normally ductile materials (Fig. 1). Brittle intercrystalline fracture predominates in fcc and bcc materials, but substantial amounts of cleavagelike fracture, especially on basal planes, often occur in hcp materials. The fracture path and appearance produced by SMIE are generally similar to those produced by LME, although the em-
Metal-Induced Embrittlement
283
FIG. 3. TEM replica of brittle intercrystalline fracture surface of D6ac steel cracked in liquid mercury at 20°C, showing tear ridges (T) and small, shallow dimples (D), which were not clearly resolved by SEM.
~~,w~~'~ -:v':. ~ °" ~ ° i ~: ~ ~, ;;i~ ~
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Fic. 2. SEM fractographs of fracture surface of D6ac steel tempered at various temperatures and cracked in solid indium at 140°C, showing (a) dimpled fractures along martensite lath boundaries for a specimen tempered at 290°C (~54HRC), (b) relatively smooth fractures along prior-austenite grain boundaries for a specimen tempered at 400°C (~51HRC), and (c) dimpled fracture along prior-austenite grain boundaries for a specimen tempered at 650°C (-41HRC). Overload fractures were transcrystalline and dimpled for all these conditions, and dimples were larger and deeper than those produced on MIE fractures.
b r i t t l e m e n t of t i t a n i u m alloys b y c a d m i u m is o n e e x c e p t i o n - - L M E p r o d u c e s p r e d o m inately intercrystalline fracture, b u t SMIE p r o d u c e s m i x t u r e s of intercrystalline a n d transcrystalline fractures. T h e fracture p a t h a n d a p p e a r a n c e in s o m e materials dep e n d o n the m i c r o s t r u c t u r e . For e x a m p l e , the fracture characteristics of a h i g h -
s t r e n g t h t e m p e r e d martensitic D6ac steel c r a c k e d in l i q u i d - m e r c u r y [22] or in solid i n d i u m [2l] e n v i r o n m e n t s d e p e n d e d on the t e m p e r i n g t e m p e r a t u r e of the steel (Fig. 2). S h a l l o w d i m p l e s , tear ridges, a n d slip lines are p r e s e n t o n "brittle" intercrystalline a n d cleavagelike fracture surfaces in m o s t materials. T h e s e features are s o m e times r e s o l v e d b y SEM [Fig. 2(a, c)] but are often o n l y a p p a r e n t w h e n fracture surfaces are e x a m i n e d b y h i g h - r e s o l u t i o n transmission electron m i c r o s c o p y f r a c t o g r a p h i c t e c h n i q u e s (Figs. 3, 4). LME a n d SMIE in h i g h - s t r e n g t h materials are u s u a l l y not associated w i t h a n y m a c r o s c o p i c signs of plasticity, b u t sensitive t e c h n i q u e s reveal small plastic z o n e s a r o u n d cracks. ]in low-
. . . . . . . . 23t F,c. 4. TEM replica of cleavagelike fracture surface of Ti-6%A1-4%V alloy cracked in liquid mercury at 20°c, showing tear ridges and small dimples that were not clearly resolved by SEM. Flutes (F) are also evident.
284
S. P.
Lynch
b FIG. 5. Light micrograph of a polished and etched section showing the slip pattern around a cleavagelike crack in the interior of a solution treated A1-Zn-Mg alloy single crystal cracked in a liquid bismuth alloy environment at 60°C, showing extensive slip, particularly on planes intersecting cracks. The slip pattern on the specimen side surface for the same crack is shown in the inset . . . .
strength materials, on the other hand, large plastic zones, with slip occurring par . ticularly on planes intersecting crack fronts, are often associated with brittle intercrystalline and cleavagelike cracks (Fig.
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5). DETECTION OF EMBRITTLING METALS AFTER FRACTURE Globules or films of embrittling metal a r e usually evident on fracture surfaces produced by LME, and can sometimes be detected with the unaided eye or by using a binocular microscope. For example, examination of high-strength (-44 HRC) martensitic steel aircraft engine nuts that had fractured into two or three pieces (Fig. 6) showed that some fracture surfaces exhibited distinctive blue and gold colors that are characteristic of oxidized cadmium. Energy-dispersive x-ray analysis (EDXA) of fracture surfaces confirmed that cadmium was present in substantial amounts. This failure occurred because cadmium-plated, high-strength steel nuts were inadvertently used instead of silver-plated steel nuts in part of the engine that experienced temperatures above the melting point of cadmium! Metallographic sections through
Fnc. 6. (a) Macroscopic appearance of a fractured high-strength cadmium-plated 4140 steel nut that was inadvertently used on bolts joining ducts carrying hot (500°C) air from the compressor of a military jet engine, and (b) SEM fractograph of fracture surface showing brittle intercrystalline facets.
LME cracks also often show films of solidified embrittling metal within cracks. Alloying between solid and liquid metals may sometimes occur prior to, or after fracture, and intermetallic compounds or other reaction products may be present on fracture surfaces. Volatile elements, such as mercury, could have evaporated if ternperatures were sufficiently high. If components have been only partially cracked by LME, extending the crack at room temperature may result in fracture through the thin layer of solidified embrittling metal within cracks, producing a fracture-surface appearance characteristic of the (con-
285
Metal-Induced Embrittlement
FIG. 7. SEM fractograph of the fracture surface of a 7475 aluminium alloy partially cracked in liquid indium at 170°Cand then subsequentlybroken open at 20°C. The dimple/veinpattern on the fracturesurface results from ductile cracking through the solidified indium layer in the crack.
strained) embrittling metal, e.g., a dimpled fracture surface for a ductile material (Fig. 7). The presence of the embrittling metal on fracture surfaces and within cracks produced by SMIE is often not readily detected because the layer of embrittling metal is probably in the range -0.01-0.1 ~m thick [6, 22], and the sampling depth of EDXA is 2-5 ~m. However, for SMIE of aluminium alloys by indium [23], and of steels and titanium alloys by cadmium [24, 25], the embrittling metals have been detected on fracture surfaces by EDXA after long analysis times. Small islands of a film can also sometimes be observed by secondary-electron mode (SEM) (Fig. 8). Embrittling elements may be more readily detected by EDXA on sections through secondary cracks than on fracture surfaces, because the layers would be approximately normal to the surface [24]. Thin layers of embrittling metals are readily detected on fracture surfaces by auger electron spectroscopy (AES), x-ray Photoelectron spectroscopy (XPS), and secondary-ion mass spectroscopy (SIMS), because the sampling depths of these techniques are 1-2 nm. However, removing films produced by atmospheric contamination may sometimes be necessary to detect underlying embrittling metal films [6].
FIG. 8. SEM fractograph of fracture surface of a Ti6% A1-4%V alloycrackedin solid cadmiumat 295°C, showingdark islands (which are probably films of cadmium)on fluted intercrystalline(I) and cleavagelike (C) areas. DIFFERENTIATING BETWEEN MIE AND HE/SCC
The appearance of fracture surfaces produced by LME and SMIE are often remarkably similar to those produced by HE and SCC [22, 26-28], and, hence, it is not possible to distinguish between these fracture modes solely on the basis of fracture surface appearance. Differentiation between failures by SMIE and HE/SCC in electroplated high-strength steels (where hydrogen is generated during the plating process) may require careful analysis for thin layers of plating material on fracture surfaces. A knowledge of the plating and baking procedures, and of the in-service operating conditions (temperature, environment, applied stresses) would also be helpful. For example, nonporous plating--which would inhibit removal of hydrogen during baking--would be a factor in favor of HE as a cause of failure. The presence of a corrosive environment during service, corrosion pits at crack-initiation sites, and corroded fracture surfaces would point to SCC as a probable cause of failure. A knowledge of the operating ternperature of the component would also help distinguish between HE and SMIE because susceptibility to HE is usually greatest near ambient temperature, whereas susceptibility to SMIE is greatest just below the embrittler's melting temperature.
286
The extent of brittle intercrystalline or transcrystalline cracking may also indicate whether cracking has occurred by HE or SMIE. The maximum crack lengths reported for SMIE of steels and titanium alloys by cadmium [24, 25], and 7075-T651 aluminium alloys by indium [23], are - 3 r a m at T/Tm -0.95 and - 1 . 5 mm at T/ Tm -0.7. Thus, brittle cracking to depths greater than these values would suggest that cracking has probably occurred by a process other than SMIE. However, crack growth could initially occur by SMIE and then continue by another process (fatigue, HE, SCC) that produced a similar fracture surface appearance to SMIE.
MECHANISMS OF METAL-INDUCED EMBRITTLEMENT EMBRITTLEMENT DUE TO ADSORPTION
Adsorption of environmental metal atoms at surface stress concentrations and at crack tips is the most common cause of embrittlement, particularly w h e n solid and liquid metals have low mutual solubilities and do not form intermetallic compounds [1, 2]. Removal or rupture of oxide or other films on surfaces is necessary to enable adsorption to occur, but diffusion of environmental atoms into solids is not involved, Thus, preexposure of unstressed specimens to liquid-metal environments, then removing the liquid metal prior to stressing, does not produce embrittlement in such systems even if some diffusion has occurred. Adsorption of environmental metal atoms on surfaces (crack tips) apparently results in changes of electron density and a consequent weakening of interatomic bonds between substrate atoms within a few atomic distances of the adsorbate. However, the details are not well understood and, hence, predicting which liquid metals embrittle particular solid metals is presently not possible from first principles. Workers studying LME in the 1960s and early 1970s concluded that adsorption low-
S. P.
Lynch
ered the stress required for "decohesion" at crack tips so that atomically brittle fractures were produced [29]. However, subsequent detailed metallographic and fractographic studies of LME, e.g., showing dimples on fracture surfaces and slip on planes intersecting cracks, suggested that adsorption facilitates the injection of dislocations at crack tips so that crack growth occurs by a more localized microvoid-coalescence process than occurs in inert environments, as discussed in detail elsewhere [22, 26-28]. Similarities between MIE and HE/SCC, and other observations, suggest that an adsorption-induced localized slip process is also applicable to HE and SCC for some materials and environments [22, 26-28]. EMBRITTLEMENT ASSOCIATED WITH DIFFUSION AND OTHER PROCESSES
Diffusion of environmental metal atoms into materials along grain boundaries ("grain boundary wetting")commonly occurs in systems in which the liquid metal is relatively insoluble in the solid, but the solid has an appreciable solubility in the liquid [30, 31]. For example, exposure of unstressed aluminium and zinc polycrystals to liquid gallium results in diffusion of gallium along grain boundaries and the formation of a thin, gallium-rich, liquid film along the grain boundaries; subsequent stressing causes the material to disintegrate into individual grains. This effect is useful for studying grain morphologies and grain-boundary cavitation caused by prior deformation [32, 33]. LME of stainless steels by liquid zinc is also sometimes associated with diffusion along grain boundaries; in this case, the formation of a NiZn compound produces local depletion of nickel at grain boundaries, thereby resulting in an austenite-to-ferrite transformation and an associated volume expansion that generates high stresses and cracking [30]. In liquid alkali-metal environments, which are commonly encountered in the nuclear industry, transfer of interstitial ele-
287
Metal-Induced Embrittlement
ments, such as carbon, nitrogen and oxygen, between liquids and solids can occur, thereby producing changes in mechanical properties [1, 13]. Reactions between some liquid and solid metals produce easily cracked, brittle intermetallic compounds or selective dissolution of alloy phases/grain boundaries/slip bands, thereby facilitating crack initiation. In some systems, e.g., stainless steels-zinc, aluminium alloysgallium, different processes may be responsible for embrittlement under different conditions, e.g., adsorption may be responsible at one temperature/strain rate, and diffusion may be involved at another temperature/strain rate.
as -10% of the critical K for cracking in inert environments (Fig. 9). Transport of liquid metal to crack tips occurs by capillary flow, and a limiting rate of transport possibly controls cracking kinetics at high crack velocities. The different "plateau" velocities for different materials (and for aluminium alloys in different conditions [34] and observations [36] that threshold K values for LME of aluminium alloys depend on the experimental conditions, suggest that other factors also affect cracking kinetics. However, these effects are not well understood. SMIE
KINETICS OF LME AND SMIE LME Delayed failure of materials can occur in liquid-metal environments because of slow crack initiation processes (which are not well understood), but crack growth generally occurs rapidly. Crack velocities in the range of 10-100 mm/s have been reported for subcritical cracking of aluminium and titanium alloys in liquid mercury [34, 35]. The velocities are not affected by stress-intensity factor (K) values except near the threshold K which can be as low
Crack velocities during SMIE are considerably less than those during LME, and they decrease with decreasing temperature. Studies of SMIE of titanium alloys by cadmium [25], and 7075-T651 aluminium alloy by indium [23], showed that crack velocities were unaffected by K values above a threshold value, and decreased with increasing crack length (i.e., increasing distance from the source of embrittling atoms) (Fig. 10). At longer crack lengths, crackgrowth rates were greater when tests were conducted in vacuum rather than in air, presumably because oxidation of fracture surfaces and crack tips inhibited transport and adsorption processes when tests were done in air. For short crack lengths, the
1 At 7075 10-~
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40
60
80
100
120
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Time (h)
FIG. 10. Crack length versus time data for 7075 aluminium and titanium alloys cracked in solid-metal env i r o n m e n t s at different h o m o l o g o u s temperatures, T/ Tm w h e r e Tm is the melting point of the embrittler and T is the testing temperature, K.
S. P. Lynch
288 embrittling metal itself probably p r e v e n t e d access of air into cracks, The previous data can be explained on the basis that crack-growth rates are controlled by transport of embrittling atoms from their source (surface or notch root) to the crack tip by a surface diffusion process, Values of surface diffusion coefficients, Ds, for SMIE at various t e m p e r a t u r e s have b e e n calculated from crack length (x) versus time (t) data using the standard random-walk diffusion equation (Ds = x2/2t), for various materials and environments, Plots of log D~ versus Tm/T (where Tm is the melting t e m p e r a t u r e of the embrittling elem e n t and T is the testing temperature) s h o w that a linear Arrhenius relationship can be fitted to most of the data (Fig. 11) [23]. C o m p a r i s o n of these data with "clean" surface self-diffusion data, obtained from the kinetics of sintering and other processes [37], s h o w e d that the surface selfdiffusion values were about one to two orders of m a g n i t u d e greater than D~ values obtained from SMIE data [23]. H o w e v e r , the values of D~ calculated from the SMIE data are u n d e r e s t i m a t e d , probably b y about an order of magnitude, because the diffusion distances were taken as the
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10-,2 10-13 10-14
TiSAI4VandTiSAIIMolV :Cd
an~4340Stee~: C~| UTS • 4140 Steel : In o 4140Steer:Pb I _>1370 A 4140Steel:Zn MPa + TisAIsV2S,: Cd , Ti6AI6V2Sn:Au .
Ti6AI6V2Sn : Ag
\ ,b ,, "*~ N "-, ~o ~, ,, ", ,~ "O~
10-,, 1'0 1'2 1'4 tls
,, N "- ,, . , ,,, 1'8 a'o z2 24 2'6 z8
Trn/T
Fic. 11. Diffusion coefficients obtained from SMIE data for a range of materials and environments plotted (on a log scale) against TIn/T, showing that most of the data fits a linear Arrhenius relationship,
projectedcrack lengths. The actual diffusion distances will d e p e n d on the r o u g h n e s s of fracture surfaces on a microscopic and submicroscopic scale, and are probably three to four times the projected crack lengths w h e n crack g r o w t h produces dimpled intercrystalline fracture surfaces. Thus, corrected values of Ds for SMIE data would be within an order of m a g n i t u d e of values for surface self-diffusion obtained by other techniques. The data for the effect of t e m p e r a t u r e on surface diffusion/SMIE kinetics suggest that cracking caused by an external source of embrittling metal could occur at significant rates (>1 m m in 1 year) at homologous t e m p e r a t u r e s >0.5 T/Tm. SMIE of high-strength steels and titanium alloys by c a d m i u m has occasionally been observed at 38°C (T/Tm -0.52), but does not appear to be a c o m m o n problem, probably because electroplated c a d m i u m is not usually in intimate contact with the substrate. H o w e v e r , such contact could develop if intervening (oxide) films were r u p t u r e d by the applied stress or by local pressure, and if embrittling atoms were sufficiently mobile. Furthermore, small cracks p r o d u c e d by SMIE at low t e m p e r a t u r e s could initiate more extensive cracking by HE, SCC, or fatigue, leading to unstable fracture. SMIE of titanium alloys by silver and gold [6] has been observed at h o m o l o g o u s t e m p e r a t u r e s - 0 . 3 8 at rates of crack growth that are considerably larger than w o u l d be expected on the basis of the other SMIE data. Possible explanations for this a n o m o l o u s behavior are (1) higher surfacediffusion rates than " n o r m a l " because of impurities in the embrittling elements [37], and (2) a contribution to cracking from HE. L o w e r rates of crack g r o w t h than predicted from surface-diffusion data have been observed for SMIE of lower strength/higher t o u g h n e s s materials than those discussed earlier, p r e s u m a b l y because some other (slower) process controls the cracking kinetics [23]. SMIE at a given t e m p e r a t u r e could occur at m u c h faster overall rates for material containing closely spaced internal sources
Metal-Induced Embrittlement
of embrittling atoms than for material in contact with an external source, because diffusion distances to crack tips would never be greater than the spacing of sources for the former. Extrapolation of the SMIE surface-diffusion data to lower temperatures suggests that SMIE could possibly occur at significant rates at homologous temperatures as low as -0.3 T/Tm when the spacing of sources is - 1 p~m. The association of embrittling elements with other elements would also be expected to influence diffusion kinetics. For example, the addition of bismuth to A1-Mg alloys prevents embrittlement by sodium, possibly because a Bi-Na compound is formed, thereby preventing diffusion of sodium t o c r a c k tips [38].
289
Vol. 11, Failure Analysis Metals Park, Ohio (1986),
and Prevention, ASM, pp. 715-727.
13. C. F. Old, I. Nucl. Mater. 92:2-25 (1980). 14. T. Shinohara and K. Matsumoto, Corr. Sci. 22:723-737 (1982). 15. D. Elliot, Process Eng. 67-71 (1976). 16. D. Bhattacharya, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 367-398. 17. M. C. Roth, T. A. Jones, G. C. Weatherly, and w . A . Miller, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME Warrendale PA (1984), pp. 399-413. 18. M. Guttmann, B. Quantin, and Ph. Dumoulin, Metal Sci. 17:123-140 (1983). 19. D. N. Fager, M. V. Hyatt, and H. T. Diep, Scripta Metall. 20:1159-1164 (1986). 20. D. Webster; Aluminium-Lithium Alloys 1II, Inst. of Metals., London, pp. 602-609 (1986). 21. s. p. Lynch, Mater. Sci. Eng. A136:45-57 (1991). 22. S. P. Lynch, Acta Metall. 32:79-90 (1984). 23. S. P. Lynch, Mater. Sci. Eng. A108:203-212 (1989).
References 1. M. H. Kamdar (ed.), Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984). (See Contents page and Index for subjects not referenced separately.) 2. N. S. Stoloff, Atomistics of Fracture, Plenum Press, New York, (1983), pp. 921-949. 3. F. A. Shunk and W. R. Warke, Scripta Metall. 8:519-526 (1974). 4. M. G. Nicholas, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 27-50. 5. S. Dinda and W. R. Warke, Mater. Sci. Eng. 24:199-208 (1976). 6. R. E. Stoltz and R. H. Stulen, Corrosion 35:165169 (1979). 7. M. G. Nicholas and C. F. Old, Weld. Met. Fabr. 47:527-531 (1979). 8. R. H. Edwards, N. I. Kilpatrick, E. G. Bradey, and J. Szczur, Mater. Forum 9;202-208, (1986). 9. M. H. Kamdar, Proc. Conf. on Environmental Degradation of Engineering Materials, Virginia Polytechnic Inst., Blacksburg, VA (1977), pp. 235-243. 10. J. R. Pickens and A. R. C. Westwood in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 51-64. 11. A. P. Druschitz and P. Gordon, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 285-316. 12. G. F. Vander Voort, Metals Handbook, 9th Ed.
24. D. N. Fager and W. F. Spurr, Corrosion 2:7:72-76 (1971). 25. D. N. Fager and W. F. Spurr, Corrosion 26:409419 (1970). 26. S. P. Lynch, Acta Metall. 29:325-340 (1981). 27. S. P. Lynch, J. Mater. Sci. 20:3329-3338 (1985). 28. S. P. Lynch, J. Mater. Sci. 21:692-704 (1986). 29. M. H. Kamdar, Prog. Mater. Sci. 15:289-374 (1973). 30. C. F. Old, Metal Sci., 14:433-440 (1980). 31. G. H. Bishop, Trans. Met. Soc. A1ME 242:13431351 (1968). 32. F. N. Rhines and A. B. Gokhale, in Embrittlement by Liquid and Solid Metals, Met. Soc. AIME, Warrendale, PA (1984), pp. 265-276. 33. T. C. Reiley, Scripta Metall. 15:497-500 (1981). 34. M. O. Speidel, in The Theory of Stress Corrosion Cracking Alloys, NATO, Brussels, (1971), pp. 289344. 35. J. A. Feeney and M. J. Blackburn, in The Theory of Stress Corrosion Cracking Alloys, NATO, Brussels (1971), pp. 355-398. 36. D. A. Wheeler and R. G. Hoagland, Scripta Metall. 20:1433-1438 (1986). 37. H. P. Bonzel, in Surface Physics of Materials, Vol. II, Academic Press, New York, (1975), pp. 279338. 38. D. E. J. Talbot and C. E. Ransley, Metai'l. Trans. A, 8A:1149-1154 (1977).
AcceptedJanuary 1992.