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MSE 323 Materials Characterization Laboratory Spring 2020 MSE 323 – Materials Characterization Laboratory Lab 2: SEM Fractorgraphy Lab Dates: 1/15, 1/22, 1/29 Report Due on 2/7 Aim of project: • To obtain an overall understanding and hands-on experience of the operation of FEI FEGSEM Sirion 200, FEI Quanta 200, or FEI Apreo. • To obtain images of the fracture surfaces of different materials and from which, identify the failure mechanisms. Experimental Procedures: • You will be given 5 fractured samples, they are: 1. 2. 3. 4. Mild/ Low carbon steel (A36) Aluminum YAG polycrystal Copper Each sample has been broken in some manner. • • • Observe and sketch the fracture surfaces of each sample with unaided eye and /or using an optical microscope. Obtain representative images (secondary and/or backscattered electron images) of the fracture surfaces from each sample. Determine from the observation the type of failure occurred in each sample. Report: Please refer to document on Blackboard Learn regarding • Style of report; and • Contents expected in the report (grading guidelines) References: You may find the following sources of information helpful: • W.D. Callister, Materials Science and Engineering: An Introduction, 10th ed., Wiley, New York, 2018. (MSE 201 book) See chapter 8 Failure (or other similar introductory materials text that discusses fracture). • ASM Handbook Volume 12 – Fractography, in the reference section in Owen Library. • http://www.ammrf.org.au/myscope/ -- this link provide a lot of detail information about SEM. You may also use the virtual SEM to practice your operation skills! Due to the large size of this file, it is broken into two parts for easier downloading. ASM Handbook, Volume 12: Fractography ASM Handbook Committee, p 12-71 DOI: 10.31399/asm.hb.v12.a0001831 Copyright © 1987 ASM International® All rights reserved. www.asminternational.org Modes of Fracture Victor Kerlins, McDonnell Douglas Astronautics Company Austin Phillips, Metallurgical Consultant METALS FAIL in many different ways and for different reasons. Determining the cause of failure is vital in preventing a recurrence. One of the most important sources of information relating to the cause of failure is the fracture surface itself. A fracture surface is a detailed record of the failure history of the part. It contains evidence of loading history, environmental effects, and material quality. The principal technique used to analyze this evidence is electron fractography. Fundamental to the application of this technique is an understanding of how metals fracture and how the environment affects the fracture process. This article is divided into three major sections. The section "Fracture Modes" describes the basic fracture modes as well as some of the mechanisms involved in the fracture process. The section "Effect of Environment" discusses how the environment affects metal behavior and fracture appearance. The final section, "Discontinuities Leading to Fracture," discusses material flaws where fracture can initiate. Fracture Modes Fracture in engineering alloys can occur by a transgranular (through the grains) or an intergranular (along the grain boundaries) fracture path. However, regardless of the fracture path, there are essentially only four principal fracture modes: dimple rupture, cleavage, fatigue, and decohesive rupture. Each of these modes has a characteristic fracture surface appearance and a mechanism or mechanisms by which the fracture propagates. In this section, the fracture surface characteristics and some of the mechanisms associated with the fracture modes will be presented and illustrated. Most of the mechanisms proposed to explain the various fracture modes are often based on dislocation interactions, involving complex slip and crystallographic relationships. The discussion of mechanisms in this section will not include detailed dislocation models or complex mathematical treatments, but will present the mechanisms in more general terms in order to impart a practical understanding as well as an ability to identify the basic fracture modes correctly. Dimple Rupture When overload is the principal cause of fracture, most common structural alloys fail by a process known as microvoid coalescence. The microvoids nucleate at regions of localized strain discontinuity, such as that associated with second-phase particles, inclusions, grain boundaries, and dislocation pile-ups. As the strain in the material increases, the microvoids grow, coalesce, and eventually form a continuous fracture surface (Fig. 1). This type of fracture exhibits numerous cuplike depressions that are the direct result of the microvoid coalescence. The cuplike depressions are referred to as dimples, and the fracture mode is known as dimple rupture. The size of the dimples on a fracture surface is governed by the number and distribution of microvoids that are nucleated. When the nucleation sites are few and widely spaced, the microvoids grow to a large size before coalescing and the result is a fracture surface that contains large dimples. Small dimples are formed when numerous nucleating sites are activated and adjacent microvoids join (coalesce) before they have an opportunity to grow to a larger size. Extremely small dimples are often found in oxide dispersion strengthened materials. The distribution of the microvoid nucleation sites can significantly influence the fracture surface appearance. In some alloys, the nonuniform distribution of nucleating particles and the nucleation and growth of isolated microvoids early in the loading cycle produce a fracture surface that exhibits various dimple sizes (Fig. 2). When microvoids nucleate at the grain boundaries (Fig. 3), intergranular dimple rupture results. Dimple shape is governed by the state of stress within the material as the microvoids form and coalesce. Fracture under conditions of uniaxial tensile load (Fig. la) results in the formation of essentially equiaxed dimples bounded by a lip or rim (Fig. 3 and 4a). Depending on the microstructure and plasticity of the material, the dimples can exhibit a very deep, conical shape (Fig. 4a) or can be quite shallow (Fig. 4b). The formation of shallow dimples may involve the joining of microvoids by shear along slip bands (Ref 1). Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user Fracture surfaces that result from tear (Mode I) or shear (Modes II and III) loading conditions (Fig. 5) exhibit elongated dimples (Ref 2, 3). The characteristics of an elongated dimple are that it is, as the name implies, elongated (one axis of the dimple is longer than the other) and that one end of the dimple is open; that is, the dimple is not completely surrounded by a rim. In the case of a tear fracture (Fig. 6a), the elongated dimples on both fracture faces are oriented in the same direction, and the closed ends point to the fracture origin. This characteristic of the tear dimples can be used to establish the fracture propagation direction (Ref 4) in thin sheet that ruptures by a full-slant fracture (by combined Modes I and III), which consists entirely of a shear lip and exhibits no macroscopic fracture direction indicators, such as chevron marks. A shear fracture, however, exhibits elongated dimples that point in opposite directions on mating fracture faces (Fig. 6b). Examples of typical elongated dimples are shown in Fig. 7. It should be noted that the illustrations representing equiaxed and elongated dimple formation and orientation were deliberately kept simple in order to convey the basic concepts of the effect on dimple shape and orientation of loading or plastic-flow directions in the immediate vicinity where the voids form, such as at the crack tip. In reality, matching dimples on mating fracture faces are seldom of the same size or seldom show equivalent angular correspondence. Because actual fractures rarely occur by pure tension or shear, the various combinations of loading Modes I, 1I, and III, as well as the constant change in orientation of the local plane of fracture as the crack propagates, result in asymmetrical straining of the mating fracture surfaces. Figure 8 shows the effect of such asymmetry on dimple size. The surface (B) that is strained after fracture exhibits longer dimples than its mating half (A). When fracture occurs by a combination of Modes I and II, examination of the dimples on mating fracture surfaces can reveal the local fracture direction (Ref 5). As illustrated in Fig. 8, the fracture plane containing the longer dimples faces the region from which the crack propagated, while the mating fracture plane containing the shorter dimples faces away from the region. With the different M o d e s of F r a c t u r e / 13 , ~ Upper s u r f a c e / ~ ~ (rmax o~ t~¢ surface Section A-A U (a) /Oval dimple Oval dimple Detail B ~ ' ~ C t " ~ - Uuppere Oma x \ Lower surface Ibl O'ma~ ~ma× Upper su rface ~ Lower surface (c) Fig. 1 Influenceof direction of maximumstress(Crmox)on the shape of dimples formed by microvoid coalescence. (a) In tension, equiaxed dimples are formed on both fracture surfaces. (b) In shear, elongated dimples point in opposite directions on matching fracture surfaces. (c) In tensile tearing, elongated dimples point toward fracture origin on matching fracture surfaces. combinations of Modes I, II, and III, there could be as many as 14 variations of dimple shape and orientation on mating fracture surfaces (Ref 5). Metals that undergo considerable plastic deformation and develop large dimples frequently contain deformation markings on the dimple walls. These markings Occur when slip-planes at the surface of the dimples are favorably oriented to the major stress direction. The continual straining of the free surfaces of the dimples as the microvoids enlarge produces slip-plane displacement at the surface of the dimple (Ref 6), as shown in Fig. 9. When first formed, the slip traces are sharp, well defined, and form an interwoven pattern that is generally referred to as serpentine glide (Fig. 10). As the slip process proceeds, the initial sharp slip traces become smooth, resulting in a surface structure that is sometimes referred to as rippies. Oval-shaped dimples are occasionally observed on the walls of large elongated dimples. An oval dimple is formed when a smaller subsurface void intersects the wall of a larger void (dimple). The formation of oval dimples is shown schematically in Fig. l(b) and 6(b). Cleavage Cleavage is a low-energy fracture that propagates along well-defined low-index crystallographic planes known as cleavage planes. The- Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user oretically, a cleavage fracture should have perfectly matching faces and should be completely flat and featureless. However, engineering alloys are polycrystalline and contain grain and subgrain boundaries, inclusions, dislocations, and other imperfections that affect a propagating cleavage fracture so that true, featureless cleavage is seldom observed. These imperfections and changes in crystal lattice orientation, such as possible mismatch of the low-index planes across grain or subgrain boundaries, produce distinct cleavage fracture surface features, such as cleavage steps, river patterns, feather markings, chevron (herringbone) patterns, and tongues (Ref 7). As shown schematically in Fig. 11, cleavage fractures frequently initiate on many parallel cleavage planes. As the fracture advances, however, the number of active planes decreases by a joining process that forms progressively higher cleavage steps. This network of cleavage steps is known as a river pattern. Because the branches of the river pattern join in the direction of crack propagation, these markings can be used to establish the local fracture direction. A tilt boundary exists when principal cleavage planes form a small angle with respect to one another as a result of a slight rotation about a common axis parallel to the intersection (Fig. 11 a). In the case of a tilt boundary, the cleavage fracture path is virtually uninterrupted, and the cleavage planes and steps propagate across the boundary. However, when the principal cleavage planes are rotated about an axis perpendicular to the boundary, a twist boundary results (Fig. 1 lb). Because of the significant misalignment of cleavage planes at the boundary, the propagating fracture reinitiates at the boundary as a series of parallel cleavage fractures connected by small (low) cleavage steps. As the fracture propagates away from the boundary, the numerous cleavage planes join, resulting in fewer individual cleavage planes and higher steps. Thus, when viewing a cleavage fracture that propagates across a twist boundary, the cleavage steps do not cross but initiate new steps at the boundary (Fig. 1 lb). Most boundaries, rather than being simple tilt or twist, are a combination of both types and are referred to as tilt-twist boundaries. Cleavage fractures exhibiting twist and tilt boundaries are shown in Fig. 12(a) and 13, respectively. Feather markings are a fan-shaped array of very fine cleavage steps on a large cleavage facet (Fig. 14a). The apex of the fan points back to the fracture origin. Large cleavage steps are shown in Fig. 14(b). Tongues are occasionally observed on cleavage fractures (Fig. 12b). They are formed when a cleavage fracture deviates from the cleavage plane and propagates a short distance along a twin orientation (Ref 8). Wallner lines (Fig. 15) constitute a distinct cleavage pattern that is sometimes observed on fractured surfaces of brittle nonmetallic materials or on brittle inclusions or intermetallic compounds. This structure consists of two sets 14 / Modes of Fracture (a) Fig. I 5 lira I I (bl 2 i~m I 2 Examples of the dimple rupture mode of fracture. (a) Large and small dimples on the fracture surface of a martempered type 234 tool steel saw disk. The extremely small dimples at top left are nucleated by numerous, closely spaced particles. (D.-W. Huang, Fuxin Mining Institute, and C.R. Brooks, University of Tennessee). (b) Large and small sulfide inclusions in steel that serve as void-nucleating sites. (R.D. Buchheit, Battelle Columbus Laboratories) I I 5 p.m Fig. 3 Intergronular dimple rupture in o steel specimen resulting from microvoid coalescence at grain boundaries. of parallel cleavage steps that often intersect to produce a crisscross pattern. Wallner lines result from the interaction of a simultaneously propagating crack front and an elastic shock wave in the material (Ref 9). Fatigue I (a) Fig. 4 I (b) I 2 iLm I Different types of dimples formed during microvoid coalescence. (a) Conical equioxed dimples in o spring steel specimen. (b) Shallow dimples in a maraging steel specimen Mode I Fig. 5 10 p.m Mode II Mode III Fracture loading modes. Arrows show loading direction and relative motion of mating fracture surfaces. Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user A fracture that is the result of repetitive or cyclic loading is known as a fatigue fracture. A fatigue fracture generally occurs in three stages: it initiates during Stage I, propagates for most of its length during Stage II, and proceeds to catastrophic fracture during Stage Ili. Fatigue crack initiation and growth during Stage I occurs principally by slip-plane cracking due to repetitive reversals of the active slip systems in the metal (Ref 10-14). Crack growth is strongly influenced by microstructure and mean stress (Ref 15), and as much as 90% of the fatigue life may be consumed in initiating a viable fatigue crack (Ref 16). The crack tends to follow crystallographic planes, but changes direction at discontinuities, such as grain boundaries. At large plastic-strain amplitudes, fatigue cracks may initiate at grain boundaries (Ref 14). A typical Stage I fatigue fracture is shown in Fig. 16. Stage I fatigue fracture surfaces are faceted, often resemble cleavage, and do not exhibit fatigue striations. Stage I fatigue is normally observed on high-cycle low-stress fractures and is frequently absent in low-cycle high-stress fatigue. The largest portion of a fatigue fracture consists of Stage II crack growth, which generally occurs by transgranular fracture and is more influenced by the magnitude of the alternating stress than by the mean stress or microstructure (Ref 15, 17, 18). Fatigue fractures generated during Stage II fatigue usually exhibit crack-arrest marks known as fatigue striations (Fig. 17 to 22), which are a visual record M o d e s of F r a c t u r e / 1 5 Principal loading , direction 1 ] ~ / / Lip of dim le -~ ~ Nucleating particle Fracture Fracture direction direction , I -c_ c c # c \ Microvoid * d I Top surface I I ~ C , "~- C ~- ~ c- I .c~ -- ic cC.cCc I? c.c -c ....... \ ° " [ ~ ' ~ %~C"~"Open"en f - k,._,.--,,,._ I ~ c ~-c- ~~ - Nucleating particle~" II '~..~-"- ' f ~ ~ i : ~ : ! \ " Bottom surface (a) Oval -----'------ C .c'-c C-'- C ~- C cl ,C%_ -r,._. ,~c,.. Top surface c..--~_ ~ _ ( t _ C ,,_. (-- ~ - ( " L'- C~ ., Principal loading direction ~ ..Q~ -~_~ ~ ~.~)*~.~)' ..t) ' ~ ~ ~) ~) -~'~) -~-~ Bottom surface (b) Fig. 6 Formation of elongated dimples under tear and shear loading conditions. (a) Tear fracture. (b) Shear fracture of the position of the fatigue crack front during crack propagation through the material. There are basically two models that have been proposed to explain Stage II striationforming fatigue propagation. One is based on plastic blunting at the crack tip (Ref 11). This model cannot account for the absence of striations when a metal is fatigue tested in vacuum and does not adequately predict the peak-topeak and valley-to-valley matching of corresponding features on mating halves of the fracture (Ref 8, 19-23). The other model, which is based on slip at the crack tip, accounts for conditions where slip may not occur precisely at the crack tip due to the presence of lattice or microstructural imperfections (Ref 19-21). This model (Fig. 23) is more successful in explaining the mechanism by which Stage II fatigue cracks propagate. The concentration of stress at a fatigue crack results in plastic deformation (slip) being confined to a small region at the tip of the crack while the remainder of the material is subjected to elastic strain. As shown in Fig. 23(a), the crack opens on the rising-tension portion of the load cycle by slip on alternating slip planes. As slip proceeds, the crack tip blunts, but is resharpened by partial slip reversal during the declining-load portion of the fatigue cycle. This results in a compressive stress at the crack tip due to the relaxation of the residual elastic tensile stresses induced in the uncracked portion of the material during the rising load cycle (Fig. 23b). The closing crack does not reweld, because the new slip surfaces created during the crack-opening displacement are instantly oxidized (Ref 24), which makes complete slip reversal unlikely. The essential absence of striations on fatigue fracture surfaces of metals tested in vacuum tends to support the assumption that oxidation reduces slip reversal during crack closure, which results in the formation of striations (Ref 19, 25, 26). The lack of oxidation in hard vacuum promotes a more complete slip reversal (Ref 27), which results in a smooth and relatively featureless fatigue fracture surface. Some fracture surfaces containing widely spaced fatigue striations exhibit slip traces on the leading edges of the striation and relatively smooth trailing edges, as predicted by the model (Fig. 23). Not all fatigue striations, however, exhibit distinct slip traces, as suggested by Fig. 23, which is a simplified representation of the fatigue process. As shown schematically in Fig. 24, the profile of the fatigue fracture can also vary, depending on the material and state of stress. Materials that exhibit fairly well-developed striations display a sawtooth~type profile (Fig. 24a) with valley-to-valley or groove-to-groove matching (Ref 23, 28). Low compressive stresses at the crack tip favor the sawtooth profile; however, high compressive stresses promote the groove-type fatigue profile, as shown in Fig. 24(c) (Ref 23, 28). Jagged, poorly formed, distorted, and unevenly spaced striations (Fig. 24b), sometimes termed quasistriations (Ref 23), show no symmetrical matching profiles. Even distinct sawtooth and groove-type fatigue surfaces may not show symmetrical matching. The local microscopic Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user plane of a fatigue crack often deviates from the normal to the principal stress. Consequently, one of the fracture surfaces will be deformed more by repetitive cyclic slip than its matching counterpart (Ref 29) (for an analogy, see Fig. 8). Thus, one fracture surface may show welldeveloped striations, while its counterpart exhibits shallow, poorly formed striations. Under normal conditions, each striation is the result of one load cycle and marks the position of the fatigue crack front at the time the striation was formed. However, when there is a sudden decrease in the applied load, the crack can temporarily stop propagating, and no striations are formed. The crack resumes propagation only after a certain number of cycles are applied at the lower stress (Ref 4, 23, 30). This phenomenon of crack arrest is believed to be due to the presence of a residual compressivestress field within the crack tip plastic zone produced after the last high-stress fatigue cycle (Ref 23, 30). Fatigue crack propagation and therefore striation spacing can be affected by a number of variables, such as loading conditions, strength of the material, microstructure, and the environment, for example, temperature and the presence of corrosive or embrittling gases and fluids. Considering only the loading condit i o n s - w h i c h would include the mean stress, the alternating stress, and the cyclic freq u e n c y - t h e magnitude of the alternating stress (Crmax -- ~min) has the greatest effect on striation spacing. Increasing the magnitude of the alternating stress produces an increase in the striation spacing (Fig. 25a). While rising, the mean stress can also increase the striation spacing; this increase is not as great as one for a numerically equivalent increase in the alternating stress. Within reasonable limits, the cyclic frequency has the least effect on striation spacing. In some cases, fatigue striation spacing can change significantly over a very short distance (Fig. 25b). This is due in part to changes in local stress conditions as the crack propagates on an inclined surface. For a Stage II fatigue crack propagating under conditions of reasonably constant cyclic loading frequency and advancing within the nominal range of 10 -5 to 10 -3 mm/cycle*, the crack growth rate, da/dN, can be expressed as a function of the stress intensity factor K (Ref 15, 31, 32): da i = C(AK)" dN (Eq 1) where a is the distance of fatigue crack advance, N is the number of cycles applied to advance the distance a, m and C are constants, and ZkK = Kmax - Kmi" is the difference between the maximum and minimum stress intensity factor for each fatigue load cycle. The *All fatigue crack growth rates in this article are given in millimeters per cycle (mm/cycle). To convert to inches per cycle (in./cycle), multiply by 0.03937. See also the Metric Conversion Guide in this Volume. 16 / Modes of Fracture o- 1 - di'° , es (a) I l mm I I (b) 20 pm T1 (T 1 Fig, 8 Effectof asymmetry on dimple size o" ....--Wall of dimple Original dimple surface (c) I 75 p.m I (d) I 20 I-tm ~ I w , slip-created surface Fig. 7 Elongated dimples formed on shear and torsion specimen fracture surfaces. (a) Shear fracture of a commercially pure titanium screw. Macrofractograph shows spiral-textured surface of shear-off screw. Typical deformation lines are fanning out on the thread. (b) Higher-magnification view of (a) shows uniformly distributed elongated shear dimples. (O.E.M. Pohler, Institut StraumannAG). (c) Elongated dimples on the surface of a fractured single-strandcopper wire that failed in torsion. (d) Higher-magnificationview of the elongated dimples shown in (c). (R.D. Lujan, Sandia National Laboratories) - - Preferably oriented slip planes (r stress intensity factor, K, describes the stress condition at a crack and is a function of the applied stress and a crack shape factor, generally expressed as a ratio of the crack depth to length. When a fatigue striation is produced on each loading cycle, da/dN represents the striation spacing. Equation 1 does not adequately describe Stage I or Stage III fatigue crack growth rates; it tends to overestimate Stage I and often underestimates Stage III growth rates (Ref 15). Stage III is the terminal propagation phase of a fatigue crack in which the striation-forming mode is progressively displaced by the static fracture modes, such as dimple rupture or cleavage. The rate of crack growth increases during Stage III until the fatigue crack becomes unstable and the part fails. Because the crack propagation is increasingly dominated by the static fracture modes, Stage III fatigue is sen- sitive to both microstructure and mean stress (Ref 17, 18). Characteristics of Fractures With Fatigue Striations. During Stage II fatigue, the crack often propagates on multiple plateaus that are at different elevations with respect to one another (Fig. 26). A plateau that has a concave surface curvature exhibits a convex contour on the mating fracture face (Ref 29). The plateaus are joined either by tear ridges or walls that contain fatigue striations (Fig. 19 and 20a). Fatigue striations often bow out in the direction of crack propagation and generally tend to align perpendicular to the principal (macroscopic) crack propagation direction. However, variations in local stresses and microstructure can change the orientation of the plane of fracture and alter the direction of striation alignment (Fig. 27). Large second-phase particles and inclusions in a metal can change the local crack growth Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user Fig. 9 Slip step formation resulting in serpentine glide and ripples on a dimple wall rate and resulting fatigue striation spacing. When a fatigue crack approaches such a particle, it is briefly retarded if the particle remains intact or is accelerated if the particle cleaves (Fig. 18). In both cases, however, the crack growth rate is changed only in the immediate vicinity of the particle and therefore does not significantly affect the total crack growth rate. However, for low-cycle (high-stress) fatigue, the relatively large plastic zone at the crack tip can cause cleavage and matrix separation at the particles at a significant distance ahead of the advancing fatigue crack. The cleaved or matrixseparated particles, in effect, behave as cracks or voids that promote a tear or shear fracture between themselves and the fatigue crack, thus significantly advancing the crack front (Ref 33, Modes of Fracture I \ Cleavage steps Cleavage planes ~ ~ / 17 Fracture direction subgrain boundary (a) /" \ Twist ~ ---L I Cleavage f e a t h e r s ~ ~ ~ River pattern I / ~ . . J Fracture d,rect,on I 5 i~m Fig. 10 Serpentine glide formation (arrow) in oxygen-free high-conductivity copper specimen Cleavage step subgrain boundary (b) Fig. 1 1 (b) (a) Fig. 1 2 Schematicof cleavage fracture formation showing the effect of subgrain and grain boundaries. (a) Tilt boundary. (b) Twist boundary I 20 p m I Examples of cleavage fractures. (a) Twist boundary, cleavage steps, and river patterns in an Fe-O.O1C-O.24Mn-O.02Sialloy that was fractured by impact. (b) Tongues (arrows) on the surface of a 30% Cr steel weld metal that fractured by cleavage Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user 1 8 / Modes Fig. of F r a c t u r e Cleavage fracture in Armco iron showing a tilt boundary, cleavage steps, and river patterns. TEM p-c replica 13 1 34). Relatively small, individual particles have no significant effect on striation spacing (Fig. 17b). The distinct, periodic markings sometimes observed on fatigue fracture surfaces are known as tire tracks, because they often resemble the tracks left by the tread pattern of a tire (Fig. 28). These rows of parallel markings are the result of a particle or a protrusion on one fatigue fracture surface being successively impressed into the surface of the mating half of the fracture during the closing portion of the fatigue cycle (Ref 23, 29, 34). Tire tracks are more common for the tension-compression than the tension-tension type of fatigue loading (Ref 23). The direction of the tire tracks and the change in spacing of the indentations within the track can indicate the type of displacement that occurred during the fracturing process, such as lateral movement from shear or torsional loading. The presence of tire tracks on a fracture surface that exhibits no fatigue striations may indicate that the fracture occurred by low-cycle (high-stress) fatigue (Ref 35). Decohesive Rupture A fracture is referred to as decohesive rupture when it exhibits little or no bulk plastic deformation and does not occur by dimple rupture, cleavage, or fatigue. This type of fracture is generally the result of a reactive environment or a unique microstructure and is associated almost exclusively with rupture along grain boundaries. Grain boundaries contain the lowest melting point constituents of an alloy system. They are also easy paths for diffusion and sites for the segregation of such elements as hydrogen, sulfur, phosphorus, antimony, arsenic, and carbon; the halide ions, such as chlorides; as well as the routes of penetration by the low melting point metals, such as gallium, mercury, cadmium, and tin. The presence of these constituents at the boundaries can significantly reduce the cohesive (a) J, 20 izm (b) Fig. 14 Examples of cleavage fractures. (a) Feather pattern on a single grain of a chromium steel weld metal that failed by cleavage. (b) Cleavage steps in a Cu-25 at.% Au alloy that failed by transgranular stress-corrosion cracking. (B.D. Lichter, Vanderbilt University) strength of the material at the boundaries and promote decohesive rupture (Fig. 29). Decohesive rupture is not the result of one unique fracture process, but can be caused by several different mechanisms. The decohesive processes involving the weakening of the atomic bonds (Ref 36), the reduction in surface energy required for localized deformation (Ref 37-39), molecular gas pressure (Ref 40), the rupture of protective films (Ref 41, 42), and anodic dissolution at active sites (Ref 43) are associated with hydrogen embrittlement and stress-corrosion cracking (SCC). Decohesive rupture resulting from creep fracture mechanisms is discussed at the end of this section. The fracture of weak grain-boundary films (such as those resulting from grain-boundary penetration by low melting point metals), the rupture of melted and resolidified grainboundary constituents (as in overheated aluminum alloys), or the separation of melted material in the boundaries (Ref 44) before it solidifies (as in the cracking at the heat-affected zones, HAZs, of welds, a condition known as hot cracking) can produce a decohesive rupture. Figures 30 to 32 show examples of decohesive rupture. A decohesive rupture resulting from hydrogen embrittlement is shown in Fig. 30. Figure 31 shows a decohesive rupture in a precipitation-hardenable stainless steel due to SCC. A fracture along a low-strength grainboundary film resulting from the diffusion of liquid mercury is shown in Fig. 32. More detailed information on hydrogen embrittlement, SCC, and liquid-metal embrittlement can be found later in this article in the section "Effect of Environment." When a decohesive rupture occurs along flattened, elongated grains that form nearly uninterrupted planes through Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user I I 1 ixm Fig. 1 5 Wallner lines (arrow) on the surface of a fractured WC-Co specimen. TEM formvar replica. Etched with 5% HCI. (S.B. Luyckx, University of the Witwatersrand) the material, as in severely extruded alloys and along the parting planes of some forgings, a relatively smooth, featureless fracture results (Fig. 33). Creep rupture is a time-dependent failure that results when a metal is subjected to stress for extended periods at elevated temperatures that are usually in the range of 40 to 70% of the absolute melting temperature of the metal. With few exceptions (Ref 45-49), creep ruptures Modes (a) I Fig. !6 Fig. ! 7 20 i~m I of Fracture/19 I Ib) 20 p.m I Stage I fatigue fracture appearance. (a) Cleovagelike, crystallographically oriented Stage I fatigue fracture in a cast Ni-14Cr-4.5Mo-] Ti-6AI-1 .SFe-2.0(Nb + To) alloy. (b) Stair-step fracture surface indicative of Stage I fatigue fracture in a cast ASTM F75 cobalt-base alloy. SEM. (R. Abrams, Howmedica, Div. Pfizer Hospital Products Group Inc.) Uniformly distributed fatigue striations in an aluminum 2024-T3 alloy. (a) Tear ridge and inclusion (outlined by rectangle). (b) Higher-magnification view of the region outlined by the rectangle in (a) showing the continuity of the fracture path through and around the inclusion. Compare with Fig. 18. Fig. exhibit an intergranular fracture surface. Transgranular creep ruptures, which generally result from high applied stresses (high strain rates), fail by a void-forming process similar to that of microvoid coalescence in dimple rupture (Ref 45-47). Because transgranular creep ruptures show no decohesive character, they will not be considered for further discussion. Intergranular creep ruptures, which occur when metal is subjected to low stresses (often well below the yield point) and to low strain rates, exhibit decohesive rupture and will be discussed in more detail. Creep can be divided into three general stages: primary, secondary, and tertiary creep. The fracture initiates during primary creep, propagates during secondary or steady-state creep, and becomes unstable, resulting in failure, during tertiary or terminal creep. From a practical standpoint of the service life of a structure, the initiation and steady-state propagation of creep ruptures are of primary importance, and most efforts have been directed toward understanding the fracture mechanisms involved in these two stages of creep. As shown schematically in Fig. 34, intergranular creep ruptures occur by either of two fracture processes: triple-point cracking or grain-boundary cavitation (Ref 50-63). The strain rate and temperature determine which fracture process dominates. Relatively high strain rates and intermediate temperatures promote the formation of wedge cracks (Fig. 34a). Grain-boundary sliding as a result of an applied Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user ! 8 Local variations in striation spacing in a Ni-O.O4C-21Cr-O.6Mn-2.5Ti-O.7AI alloy that was tested under rotating-bending conditions. Compare with Fig. 17(b). tensile stress can produce sufficient stress concentration at grain-boundary triple points to initiate and propagate wedge cracks (Ref 50-52, 55, 56, 58-61). Cracks can also nucleate in the grain boundary at locations other than the triple point by the interaction of primary and secondary slip steps with a sliding grain boundary (Ref 61). Any environment that lowers grainboundary cohesion also promotes cracking (Ref 59). As sliding proceeds, grain-boundary 20 / Modes Fig. 19 of Fracture Fatigue striations in a 2024-T3 aluminum alloy joined by tear ridges as deformation continues, the cavities join to form an intergranular fracture. Even though the fracture resulting from cavitation creep exhibits less sharply defined intergranular facets (Fig. 35c), it would be considered a decohesive rupture. Instead of propagating by a cracking or a cavity-forming process, a creep rupture could occur by a combination of both. There may be no clear distinction between wedge cracks and cavities (Ref 70-72). The wedge cracks could be the result of the linkage of cavities at triple points. The various models proposed to describe the creep process are mathematically complex and were not discussed in detail. Comprehensive reviews of the models are available in Ref 59, 63, 73, and 74. Unique cracks propagate and join to form intergranular decohesive fracture (Fig. 35a and b). At high temperatures and low strain rates, grain-boundary sliding favors cavity formation (Fig. 34b). The grain-boundary cavities resulting from creep should not be confused with microvoids formed in dimple rupture. The two are fundamentally different; the cavities are principally the result of a diffusion-controlled process, while microvoids are the result of complex slip. Even at low strain rates, a sliding grain boundary can nucleate cavities at irregularities, such as second-phase inclusion particles (Ref 54, 57, 63, 64). The nucleation is believed to be a strain-controlled process (Ref 63, 64), while the growth of the cavities can be described by a diffusion growth model (Ref 65-67) and by a power-law growth relationship (Ref 68, 69). Irrespective of the growth model, 76-83). In steels, the cleavage facets of quasicleavage fracture occur on the {100}, {110}, and possibly the {112} planes. The term quasicleavage can be used to describe the distinct fracture appearance if one is aware that quasicleavage does not represent a separate fracture mode. A quasi-cleavage fracture initiates at the central cleavage facets; as the crack radiates, the cleavage facets blend into areas of dimple rupture, and the cleavage steps become tear ridges. Quasi-cleavage has been observed in steels, including quench-and-temper hardenable, precipitation-hardenable, and austenitic stainless steels; titanium alloys; nickel alloys; and even aluminum alloys. Conditions that impede plastic deformation promote quasicleavage fracture--for example, the presence of a triaxial state of stress (as adjacent to the root of a notch), material embrittlement (as by hydrogen or stress corrosion), or when a steel is subjected to high strain rates (such as impact loading) within the ductile-to-brittle transition range. Flutes. Fractography has acquired a number of colorful and descriptive terms, such as dimple rupture, serpentine glide, ripples, tongues, tire tracks, and factory roof, which describes a ridge-to-valley fatigue fracture topography resulting from Mode III antiplane shear loading (Ref 84). The term flutes should also be included in this collection. Flutes exhibit elongated grooves or voids (Fig. 38 and 39) that connect widely spaced cleavage planes (Ref 85-90). The fracture process is known as fluting. The term flutes was apparently chosen because the fractures often resemble the long, parallel grooves on architectural columns or the pleats in drapes. Fractures Some fractures, such as quasi-cleavage and flutes, exhibit a unique appearance but cannot be readily placed within any of the principal fracture modes. Because they can occur in common engineering alloys under certain failure conditions, these fractures will be briefly discussed. Q u a s i - c l e a v a g e f r a c t u r e is a localized, often isolated feature on a fracture surface that exhibits characteristics of both cleavage and plastic deformation (Fig. 36 and 37). The term quasi-cleavage does not accurately describe the fracture, because it implies that the fracture resembles, but is not, cleavage. The term was coined because, although the central facets of a quasi-cleavage fracture strongly resembled cleavage (Ref 75), their identity as cleavage planes was not established until well after the term had gained widespread acceptance (Ref I I I 10 izm Fig. 20 Fatigue striations on adjoining walls on the fracture surface of a commercially pure titanium specimen. (O.E.M. Pohler, Institut Straumann AG) Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user I lO izm Fig. 21 Fatigue striations on the fracture surface of a tantalum heat-exchanger tube. The rough surface appearance is due to secondary cracking caused by high-cycle low-amplitude fatigue. (M.E. Blum, FMC Corporation) M o d e s of F r a c t u r e / 21 I I I B p~m Fig. 22 10 p.m High-magnlflcatlon views of fatigue striations, (a) Striations (arrow) on the fracture surface of an oustenitic stainless steel. (C.R. Brooks and A. Choudhury, University of Tennessee). (b) Fatigue striations on the facets of tantalum grains in the heat-affected zone of a weldment. (M.E. Blum, FMC Corporation) o- l Slip Orientation of active slip planes $ /" \\ '• ~// \\ \\ / \\ ~ \', c~ (Principal tensile tensile stress) (a) arrangement that resembles cleavage river patterns (Ref 89). Although fluting has been observed primarily in hexagonal close-packed (hcp) metal systems, such as titanium and zirconium alloys, evidence of fluting has also been reported on a hydrogen-embrittled type 316 austenitic stainless steel (Ref 90). Titanium alloys having a relatively high oxygen or aluminum content (o~-stabilizers) that are fractured at cryogenic temperatures or fail by SCC may exhibit fluting (Ref 89). Tearing Topography f (Compressive closure stress) Relatively smooth trailing edge of striation S l i p First Second Crack advance during one load cycle traces on leading edge of striation Third cycle (b) Mechanism of fatigue crack propagation by alternate sllp at the crack tip. Sketches are simplified to F l g o 2 3 clarify the basic concepts. (a) Crack opening and crack tip blunting by slip on alternate slip planes with increasing tensile stress. (b) Crack closure and crack tip resharpening by partial slip reversal on alternate slip planes with increasing compressive stress Although flutes are not elongated dimples, they are the result of a plastic deformation process. Flutes are the ruptured halves of tubular voids believed to be formed by a planar intersecting slip mechanism (Ref 85, 88, 89) and have matching tear ridges on opposite fracture faces. The tear ridges join in the direction of fracture propagation, forming an Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user Surface A tentative fracture mode called tearing topography surface (TTS) has been identified and described (Ref 91). The TTS fracture occurs in a variety of alloy systems, including steels, aluminum, titanium, and nickel alloys, and under a variety of fracture conditions, such as overload, hydrogen embrittlement (Ref 92), and fatigue. Examples of TTS fractures are shown in Fig. 40 to 42. Although the precise nucleation and propagation mechanism for TTS fracture has not been identified, the fracture appears to be the result of a rnicroplastic tearing process that operates on a very small (submicron) scale (Ref 91). The TTS fractures do not exhibit as much plastic deformation as dimple rupture, although they are often observed in combination with dimples (Fig. 40 and 41). The fractures are generally characterized by relatively smooth, often flat, areas or facets that usually contain thin tear ridges. Tearing topography surface fractures may be due to closely spaced microvoid nucleation and limited growth before coalescence, resulting in extemely shallow dimples. However, this hypothesis does not appear to be probable, because TTS is often observed along with well-developed dimples in alloys having relatively uniform carbide dispersions, such as HY-130 steel, and because TTS 22 / Modes of Fracture (a) (b) (c) Fig. 24 Sowtooth and groove-type fatigue fracture profiles. Arrows show crack propagation direction. (a) Distinct sawtooth profile (aluminum alloy). (b) Poorly formed sawtooth profile (steel). (c) Groove-type profile (aluminum alloy). Source: Ref 23 is observed under varying stress states. A detailed discussion of the TTS fracture mode is available in Ref 91. centered cubic (fcc) metals and alloys are generally considered to have good resistance to hydrogen embrittlement, it has been shown that the 300 series austenitic stainless steels (Ref 95-98) and certain 2000 and 7000 series highstrength aluminum alloys are also embrittled by hydrogen (Ref 99-107). Although the result of hydrogen embrittlement is generally perceived to be a catastrophic fracture that occurs well below the ultimate strength of the material and exhibits no ductility, the effects of hydrogen can be quite varied. They can range from a slight decrease in the percent reduction of area at fracture to premature rupture that exhibits no ductility (plastic deformation) and occurs at a relatively low applied stress. The source of hydrogen may be a processing operation, such as plating (Fig. 30) or acid cleaning, or the hydrogen may be acquired from the environment in which the part operates. If hydrogen absorption is suspected. prompt heating at an elevated temperature (usually about 200 °C, or 400 °F) will often restore the original properties of the material. The effect of hydrogen is strongly influenced by such variables as the strength level of the Effect of E n v i r o n m e n t The environment, which refers to all external conditions acting on the material before or during fracture, can significantly affect the fracture propagation rate and the fracture appearance. This section will present some of the principal effects of such environments as hydrogen, corrosive media, low-melting metals. state of stress, strain rate, and temperature. Where applicable, the effect of the environment on the fracture appearance will be illustrated. Effect of Environment on Dimple Rupture The Effect of H y d r o g e n . When certain body-centered cubic (bcc) and hcp metals or alloys of such elements as iron. nickel, titanium, vanadium, tantalum, niobium, zirconium. and hafnium are exposed to hydrogen, they are susceptible to a type of failure known as hydrogen embrittlement. Although the face- (a) Fig. J 25 2 i~m I (b) alloy, the microstructure, the amount of hydrogen absorbed (or adsorbed), the magnitude of the applied stress, the presence of a triaxial state of stress, the amount of prior cold work. and the degree of segregation of such contaminant elements as phosphorus, sulfur, nitrogen. tin. or antimony at the grain boundaries. In general, an increase in strength, higher absorption of hydrogen, an increase in the applied stress, the presence of a triaxial stress state. extensive prior cold working, and an increase in the concentration of contaminant elements at the grain boundaries all serve to intensify the embrittling effect of hydrogen. However. for an alloy exhibiting a specific strength level and microstructure, there is a stress intensity. K I. below which, for all practical purposes, hydrogen embrittlement cracking does not occur. This threshold crack tip stress intensity factor is determined experimentally and is designated as Kth. A number of theories have been advanced to explain the phenomenon of hydrogen embrittlement. These include the exertion of an internal gas pressure at inclusions, grain boundaries. surfaces of cracks, dislocations, or internal voids (Ref 40. 108. 109): the reduction in atomic and flee-surface cohesive strength (Ref 110-116): the attachment of hydrogen to dislocations, resulting in easier dislocation breakaway from the pinning effects of carbon and nitrogen (Ref 38. 112. 117-122): enhanced nucleation of dislocations (Ref 112. 123): enhanced nucleation and grow:h of microvoids (Ref 109. 110. 113. 116. 122. 124-126): enhanced shear and decrease of strain for the onset of shear instability (Ref 112. 127, 128): I 10 i~m I Variations in fatigue striation spacing. (a) Spectrum-loaded fatigue fracture in a 7475-T7651 aluminum alloy test coupon showing an increase in striation spacing due to higher alternating stress. (b) Local variation in fatigue striation spacing in a spectrum-loaded 7050-T7651 aluminum alloy extrusion. (D. Brown, Douglas Aircraft Company) Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user M o d e s of F r a c t u r e / 2 3 Convex Tear r i d g //~-~..]~/ Concave~ Fig. 26 e ~ Crack propagation direction Schematic illustrating fatigue striations on plateaus the formation of methane gas bubbles at grain boundaries (Ref 129, 130); and, especially for titanium alloys, the repeated formation and rupture of the brittle hydride phase at the crack tip (Ref 131-137). Probably no one mechanism is applicable to all metals, and several mechanisms may operate simultaneously to embrittle a material. Whatever the mechanism, the end result is an adverse effect on the mechanical properties of the material. If the effect of hydrogen is subtle, such as when there is a slight decrease in the reduction of area at fracture as a result of a tensile test, there is no perceivable change in the dimple rupture fracture appearance. However, the dimples become more numerous but are more shallow at a greater loss in ductility (Fig. 43). Hydrogen Embrittlement of Steels. At low strain rates or when embrittlement is more severe, the fracture mode in steels can change from dimple rupture to quasi-cleavage, cleavage, or intergranular decohesion. These changes in fracture mode or appearance may not occur over the entire fracture surface and are usually more evident in the region of the fracture origin. Figure 44 shows an example of a hydrogen-embrittled AISI 4340 steel that exhibits quasi-cleavage. When an annealed type 301 austenitic stainless steel is embrittled by hydrogen, the fracture Fig. 27 Striations on two joining, independent fatigue crack fronts on a fracture surface of aluminum alloy 6061-T6. The two arrows indicate direction of local crack propagation. TEM p-c replica occurs by cleavage (Fig. 45a). An example in which the mode of fracture changed to intergranular decohesion in a hydrogen-embrittled AISI 4130 steel is shown in Fig. 45b. When a hydrogen embrittlement fracture propagates along grain boundaries, the presence of such contaminant elements as sulfur, phosphorus, nickel, tin, and antimony at the boundaries can greatly enhance the effect of hydrogen (Ref 111, 139). For example, the segregation of contaminant elements at the grain boundaries enhances the hydrogen embrittlement of high-strength low-alloy steels tempered above 500 °C (930 °F) (Ref 92). The presence of sulfur at grain boundaries promotes hydrogen embrittlement of nickel, and for equivalent concentrations, the effect of sulfur is nearly 15 times greater than that of phosphorus (Ref 140). I4ydrogen Embrittlement of Titanium. Although titanium and its alloys have a far greater tolerance for hydrogen than high-strength steels, titanium alloys are embrittled by hydrogen. The degree and the nature of the embritdement is strongly influenced by the alloy, the microstructure, and whether the hydrogen is :::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: .... iii: i (a) (b) Fig. 29 (c) Schematic illustrating decohesive rupture along grain boundaries. (a) Decohesion along grain boundaries of equiaxed grains. (b) Decohesionthrough a weak grain-boundary phase. (c) Decohesion along grain boundaries of elongated grains Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user I I 5 I*m Fig. 28 Tire tracks on the fatigue fracture surface of a quenched-and-tempered AISI 4140 steel. TEM replica. (I. Le May, Metallurgical Consulting Services Ltd.) present in the lattice before testing or is introduced during the test. For example, a Ti-8A11Mo-IV alloy that was annealed at 1050 °C (1920 °F), cooled to 850 °C (1560 °F). and water quenched to produce a coarse Widmanstatten structure exhibited cracking along the et-[3 interfaces when tested in 1-arm hydrogen gas at room temperature (Ref 137). The fracture surface, which exhibited crack-arrest markings. is shown in Fig. 46(a). The arrest markings are believed to be due to the discontinuous crack propagation as a result of the repeated rupture of titanium hydride phase at the crack tip (Ref 137). Also, Fig. 46(b) shows a hydrogen embrittlement fracture in a Ti-5AI-2.5Sn alloy containing 90 ppm H that was 13 processed at 1065 °C (1950 °F) and aged for 8 h at 950 °C (1740 °F). The fracture occurred by cleavage. Cleavage was also the mode of fracture for a Ti-6A1-4V alloy having a microstructure consisting of a continuous, equiaxed et phase with a fine, dispersed 13 phase at the (x grain boundaries embrittled by exposure to hydrogen gas at a pressure of 1 atm (Fig. 47a). However. when the same Ti-6A1-4V alloy having a microstructure consisting of a medium, equiaxed et phase with a continuous 13 network was embrittled by 1-atm hydrogen gas, the fracture occurred by intergranular decohesion along the ct-[3 boundaries (Fig. 47b and c). Hydrogen Embrittlement of Aluminum. There is conclusive evidence (Ref 99-107) that some aluminum alloys, such as 2124, 7050, 7075, and even 5083 (Ref 143), are embrittled by hydrogen and that the embrittlement is apparently due to some of the mechanisms already discussed, namely enhanced slip and trapping of hydrogen at precipitates within grain boundaries. The embrittlement in alumi- 24 / Modes of Fracture (a) I 1 pm I I I 2.5 i~m (b) Fig. 30 Decohesive rupture in an AISI 8740 steel nut due to hydrogen embrittlement. Failure was due to inadequate baking following cadmium plating; thus, hydrogen, which was picked up during the plating process, was not released. (a) Macrograph of fracture surface. (b) Higher-magnification view of the boxed area in (a) showing typical intergranular fracture. (W.L. Jensen, Lockheed Georgia Company) num alloys depends on such variables as the microstructure, strain rate, and temperature. In general, underaged microstructures are more susceptible to hydrogen embrittlement than the peak or averaged structures. For the 7050 aluminum alloy, a low (0.01%) copper content renders all microstructures more susceptible to embrittlement than those of normal (2.1%) copper content (Ref 106). Also, hydrogen embrittlement in aluminum alloys is more likely to occur at lower strain rates and at lower temperatures. The effect of hydrogen on the fracture appearance in aluminum alloys can vary from no significant change in an embrittled 2124 alloy (Ref 99) to a dramatic change from the normal dimple rupture to a combination of cleavagelike transgranular fracture and intergranular decohesion in the high-strength 7050 (Ref 106) and 7075 (Ref 105) aluminum alloys. Figure 48 shows an example of a fracture in a hydrogenembrittled (as measured by a 21% decrease in the reduction of area at fracture) 2124-UT (underaged temper: aged 4 h at 190 °C, or 375 °F) aluminum alloy. It can be seen that there is little difference in fracture appearance between the nonembrittled and embrittled specimens. However, when a low-copper (0.01%) 7050 in the peak-aged condition (aged 24 h at 120 °C, or 245 °F) is hydrogen embrittled, a cleavagelike transgranular fracture results (Fig. 49a). This same alloy in the underaged condition (aged 10 h at 100 °C, or 212 °F) fails by a combination of inter- lal I 1 mm I granular decohesion and cleavagelike fracture (Fig. 49b). The Effect of a Corrosive E n v i r o n m e n t . When a metal is exposed to a corrosive environment while under stress, SCC, which is a form of delayed failure, can occur. Corrosive environments include moist air; distilled and tap water; seawater; gaseous ammonia and ammonia in solutions; solutions containing chlorides or nitrides; basic, acidic, and organic solutions; and molten salts. The susceptibility of a material to SCC depends on such variables as strength, microstructure, magnitude of the applied stress, grain orientation (longitudinal or short transverse) with respect to the principal applied stress, and the nature of the corrosive environment. Similar to the Kth in hydrogen (b) I 100 p.m I F l g o 3 1 17-4 PH stainlesssteel main landing-gear deflection yoke that failed because of intergranular SCC. (a) Macrograph of fracture surface. (b) Higher-magnification view of the boxed area in (a) showing area of intergranular attack. (W.L. Jensen, Lockheed Georgia Company) Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user Modes embrittlement, there is also a threshold crack tip stress intensity factor, K=scc, below which a normally susceptible material at a certain strength, microstructure, and testing environment does not initiate or propagate stresscorrosion cracks. Stress-corrosion cracks normally initiate and propagate by tensile stress; however, compressive-stress SCC has been observed in a 7075-T6 aluminum alloy and a type 304 austenitic stainless steel (Ref 144). Stress-corrosion cracking is a complex phenomenon, and the basic fracture mechanisms are still not completely understood. Although such processes as dealloying (Ref 145-148) in brass and anodic dissolution (Ref 149, 150, 151) in other alloy systems are important SCC mechanisms, it is apparent that the principal SCC mechanism in steels, titanium, and aluminum alloys is hydrogen embrittlement (Ref 38, 100, 107, 137, 143, 152-166). In these alloys, SCC occurs when the hydrogen generated as a result of corrosion diffuses into and embrittles the material. In these cases, SCC is used to describe the test or failure environment, rather than a unique fracture mechanism. Mechanisms of SCC. The basic processes that lead to SCC, especially in environments containing water, involve a series of events that begin with the rupture of a passive surface film (usually an oxide), followed by metal dissolution, which results in the formation of a pit or crevice where a crack eventually initiates and propagates. When the passive film formed during exposure to the environment is ruptured by chemical attack or mechanical action (creepstrain), a clean, unoxidized metal surface is exposed. As a result of an electrochemical potential difference between the new exposed metal surface and the passive film, a small electrical current is generated between the anodic metal and the cathodic film. The relatively small area of the new metal surface compared to the large surface area of the surrounding passive film results in an unfavorable anode-tocathode ratio. This causes a high local current density and induces high metal dissolution (anodic dissolution) at the anode as the new metal protects the adjacent film from corrosion; that is, the metal surface acts as a sacrificial anode in a galvanic couple. If the exposed metal surface can form a new passive film (repassivate) faster than the new metal surface is created by film rupture, the corrosion attack will stop. However, if the repassivation process is suppressed, as in the presence of chlorides, or if the repassivated film is continuously ruptured by strain, as when the material creeps under stress, the localized corrosion attack proceeds (Ref 167-172). The result is the formation and progressive enlargement of a pit or crevice and an increase in the concentration of hydrogen ions and an accompanying decrease in the pH of the solution within the pit. The hydrogen ions result from a chemical reaction between the exposed metal and the water within the cavity. The subsequent I Fig. 50 p.m of F r a c t u r e / 25 I 32 Fracture surface of a Monel specimen that failed in liquid mercury. The fracture is predominantly intergranular with some transgranular contribution. (C.E. Price, Oklahoma State University) Stress-corrosion fracture that occurred by decohesion along the parting plane of an aluminum alloy forging reduction of the hydrogen ions by the acquisition of electrons from the environment results in the formation of hydrogen gas and the diffusion of hydrogen into the metal. This absorption of hydrogen produces localized cracking due to a hydrogen embrittlement mechanism (Ref 173, 174). Because the metal exposed at the crack tip as the crack propagates by virtue of hydrogen embrittlement and the applied stress is anodic to the oxidized sides of the crack and the adjacent surface of the material, the electrochemical attack continues, as does the evolution and absorption of hydrogen. The triaxial state of stress and the stress concentration at the crack tip enhance hydrogen embrittlement and provide a driving force for crack propagation. In materials that are insensitive to hydrogen embrittlement, SCC can proceed by the anodic dissolution process with no assistance from hydrogen (Ref 149, 155, 161). Alloys are not homogeneous, and when differences in chemical composition or variations in internal strain occur, electrochemical potential differences arise between various areas within the microstructure. For example, the grain boundaries are usually anodic to the material within the grains and are therefore subject to preferential Fig. 33 o- (T T / (T (a) Fig. 34 Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user = \ ff (b) Triple-point cracking (a) and cavitation (b) in intergranular creep rupture. Small arrows indicate grain-boundary sliding. 26 / Modes (a) Fig. 35 of Fracture I I 18 I~m I ! 10 ~m (c) I 1 iLm I Examples of intergranular creep fractures. (a) Wedge cracking in Inconel 625. (b) Wedge cracking in Incoloy 800. (c) Intergranular creep fracture resulting from grain-boundary cavitation in PE-16. Source: Ref 59 anodic dissolution when exposed to a corrosive environment. Inclusions and precipitates can exhibit potential differences with respect to the surrounding matrix, as can plastically deformed (strained) and undeformed regions within a material. These anode-cathode couplings can initiate and propagate dissolution cracks or fissures without regard to hydrogen. (al (bl Although other mechanisms may operate (Ref 175-178), including the adsorption of unspecified damaging species (Ref 177) and the occurrence of a strain-induced martensitic transformation (Ref 178), dezincification or dealloying (Ref 145-148) appears to be the principal SCC mechanism in brass (copper-zinc and copper-zinc-tin alloys). Dezincification is the preferential dissolution or loss of zinc at the fracture interface during SCC, which can result in the corrosion products having a higher concentration of zinc than the adjacent alloy. This dynamic loss of zinc near the crack aids in propagating the stress-corrosion fracture. Some controversy remains regarding the precise mechanics of dezincification. One mechanism assumed that both zinc and copper are dissolved and that the copper is subsequently redeposited, while the other process involves the diffusion of zinc from the alloy, resulting in a higher concentration of copper in the depleted zone (Ref 179). However, there is evidence that both processes may operate (Ref 180). Like hydrogen embrittlement, SCC can change the mode of fracture from dimple rupture to intergranular decohesion or cleavage, (b) Fig. 36 Examples of quasi-cleavage. (a) Fracture surface of an austenitized Fe-O.3C-O.6Mn-5.0Mo specimen exhibiting large quasi-cleavage facets, such as at A; elsewhere, the surface contains rather large dimples. (b) Charpy impact fracture in an Fe-O.18C-3.85Mo steel. Many quasi-cleavage facets are visible. The rectangle outlines a tear ridge. Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user Fig. 37 Small and poorly defined quasi-cleavage facets connected by shallow dimples on the surface of a type 234 tool steel. (D.-E. Huang, Fuxin Mining Institute, and C.R. Brooks, University of Tennessee) Modes (a) I I (b) of Fracture I 60 l~m to) } 17 i~m / 27 I 15 izm I (dl I 10 i~m I Examples of fluting. (a) Flutes and cleavage resulting from a mechanical overload of a Ti-0.350 alloy. (b) Flutes and cleavage resulting from SCC at 13-annealed F i g . 3 8 Ti-8AI-1Mo-IV alloy in methanol. (c) Flutes and cleavage in 13-annealed Ti-8AI-1Mo-IV resulting from sustained-load cracking in vacuum. (d) Flutes occurring near the notch on the fracture surface of mill-annealed Ti-SAI-1Mo-IV resulting from corrosion fatigue in saltwater. Source: Ref 89 although quasi-cleavage has also been observed. The change in fracture mode is generally confined to that portion of the fracture that propagated by SCC, but it may extend to portions of the rapid fracture if a hydrogen embrittlement mechanism is involved. Stress-corrosion fractures that result from hydrogen embrittlement closely resemble those fractures; however, stress-corrosion cracks usually exhibit more secondary cracking, pitting, and corrosion products. Of course, pitting and corrosion products could be present on a clean hydrogen embrittlement fracture exposed to a corrosive environment. SCC of Steels. Examples of known stresscorrosion fractures are shown in Fig. 50 to 56. Steels, including the stainless grades, stress corrode in such environments as water, seawater, chloride- and nitrate-containing solu- tions, and acidic as well as basic solutions, such as those containing sodium hydroxide or hydrogen sulfide. Stress-corrosion fractures in high-strength quench-and-temper hardenable or precipitation-hardenable steels occur primarily by intergranular decohesion, although some transgranular fracture may also be present. Figure 50 shows a stress-corrosion fracture in an HY- 180 quench-and-temper hardenable steel tested in aqueous 3.5% sodium chloride. The stress-corrosion fracture was believed to have occurred predominantly by hydrogen embrittlement (Ref 154). Increasing the stress intensity coefficient, K 1, resulted in a decreased tendency for intergranular decohesion; however, the oppposite was true for a cold-worked type 316 austenitic stainless steel tested in boiling aqueous magnesium chloride (Ref 181). It was Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user shown that increasing K~ or increasing the negative electrochemical potential resulted in an increased tendency toward intergranular decohesion (Fig. 51). When the 300 type stainless steels are sensitized----a condition that results in the precipitation of chromium carbides at the grain boundaries, causing depletion of chromium in the adjacent material in the grains-the steel becomes susceptible to SCC, which occurs principally along grain boundaries. Figure 52 shows the effect of the electrochemical potential, E, on the fracture path in a cold-worked AISI C-1018 low-carbon steel that stress corroded in a hot sodium hydroxide solution. At an electrochemical potential of E = - 0 . 7 6 VsH E, the fracture path is predominantly intergranular; at a freely corroding potential o f E = - 1.00 VsH E, the fracture path is transgranular (Ref 182). 2 8 / M o d e s of F r a c t u r e I Fig. 20 ~m I 39 Flutes and cleavage resulting from SCC of 13-annealed Ti-8AI-1Mo-IV in methanol. Source: Ref 89 SCC of Aluminum. Aluminum alloys, especially the 2000 and 7000 series, that have been aged to the high-strength T6 temper or are in an underaged condition are susceptible to SCC in such environments as moist air, water, and solutions containing chlorides. The sensitivity to SCC depends strongly on the grain orientation with respect to the principal stress, the shorttransverse direction being the most susceptible to cracking. Figure 53 shows examples of stresscorrosion fractures in a 7075-T6 (maximum tensile strength: 586 MPa, or 85 ksi) aluminum alloy that was tested in water. The fracture occurred primarily by intergranular decohesion. SCC of brass in the presence of ammonia and moist air has long been recognized. The term season cracking was used to describe the SCC of brass that appeared to coincide with the moist weather in the spring and fall. Environments containing nitrates, sulfates, chlorides, ammonia gas and solutions, and alkaline solutions are known to stress corrode brass. Even distilled water and water containing as little as 5 x 10 3% sulfur dioxide have been shown to attack brass (Ref 176, 178). Depending on the arsenic content of the Cu-30Zn brass, SCC in distilled water occurs either by intergranular decohesion or by a combination of cleavage and intergranular decohesion (Fig. 54). When brass containing 0.032% As is stress corroded in water containing minute amounts of sulfur dioxide, it exhibits a unique transgranular fracture containing relatively uniformly spaced, parallel markings (Fig. 55). These distinct periodic marks apparently represent the stepwise propagation of the stress-corrosion fracture. SCC of titanium alloys has been observed in such environments as distilled water, seawater, aqueous 3.5% sodium chloride, chlorinated organic solvents, methanol, red fuming nitric Fig. 40 Appearance of TTS fracture in bainitic HY-130 steel. (a) Areas of complex tearing (T) and dimple rupture (DR). (b) Detail from upper left corner of (a) showing particle-nucleated dimples (DR) and region of TTS. SEM fractographs in (c) and (d) show additional examples of TTS fractures. Source: Ref 91, 93 Fig. 41 Appearance of TTS fracture. (a) An essentially 100% pearlitic eutectoid steel (similar to AISI 1080) where fracture propagates across pearlite colonies. (b) Fractograph showing dimple rupture (DR) and TTS fracture in a quenched-and-tempered (martensitic) HY-130 steel. Source: Ref 91, 93 acid, and molten salts. Susceptiblity depends on such variables as the microstructure (Ref 183-185), the amount of internal hydrogen (Ref 186), the state of stress (Ref 187, 188), Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user and strength level (Ref 188). In general, microstructures consisting of large-grain et phase or containing substantial amounts of phase in relation to [3 phase, high levels of Modes Fig. 42 Examples of TTS fracture in Ti-6AI-4V a-J3 alloys. (a) Solution treated and aged microstructure consisting of about lO-~xm diam primary a particles in a matrix of about 70 vol% of fine Widmanst~itten a and 13. The microstructural constituents are not evident on the fracture surface as verified by the plateau-etching technique (Ref 91, 94). (b) Fractograph of a 13-quenched Ti-6AI-4V alloy consisting of a fine Widmanst6tten martensitic microstructure. The tearing portions of the fracture surface exhibit TTS. internal hydrogen, the presence of a triaxial state of stress, and high yield strengths all promote the susceptibility of an alloy to SCC. If hydrogen is present in the corrosive environment, SCC will probably occur by a hydrogen embrittlement mechanism. Depending on the environment, alloy, and heat treatment (microstructure), mild stresscorrosion attack can exhibit a fracture that cannot be readily distinguished from normal overload, while more severe attack results in cleavage or quasi-cleavage fracture. Figure 56 shows a stress-corrosion fracture in an annealed Ti-8AI-IMo-IV alloy that was tested in aqueous 3.5% sodium chloride. The stress-corrosion fractures in titanium alloys exhibit both cleavage (along with fluting) and quasi-cleavage. Corrosion products are a natural by-product of corrosion, particularly on most steels and aluminum alloys. They not only obscure fracture detail but also cause permanent damage, because a portion of the fracture surface is chemically attacked in forming the corrosion products. Therefore, removing the corrosion products will not restore a fracture to its original condition. However, if the corrosion damage is moderate, enough surface detail remains to identify the mode of fracture. Depending on the alloy and the environment, corrosion products can appear as powdery residue, amorphous films, or crystalline deposits. Corrosion products may exhibit cleavage fracture and secondary cracking. Care must be exercised in determining whether these fractures are part of the corrosion product or the base alloy. Some of the corrosion products observed on an austenitic stainless steel and a niobium alloy are shown in Fig. 57 and 58, respectively. Detailed information on the cleaning of fracture surfaces is available in the article "Preparation and Preservation of Fracture Specimens" in this Volume. Effect of E x p o s u r e to L o w - M e l t i n g M e t als. When metals such as certain steels, tita- nium alloys, nickel-copper alloys, and aluminum alloys are stressed while in contact with low-melting metals, including lead, tin, cadmium, lithium, indium, gallium, and mercury, they may be embrittled and fracture at a stress below the yield strength of the alloy. If the embrittling metal is in a liquid state during exposure, the failure is referred to as liquidmetal embrittlement (LME); when the metal is solid, it is known as solid-metal embrittlement (SME). Both failure processes are sometimes called stress alloying. Temperature has a significant effect on the rate of embrittlement. For a specific embrittling metal species, the higher the temperature, the more rapid the attack. In addition, LME is a faster process than SME. In fact, under certain conditions, LME can occur with dramatic speed. For liquid indium embrittlement of steel, the time to failure appears to be limited primarily by the diffusion-controlled period required to form a small propagating crack (Ref 189). Once the crack begins to propagate, failure can occur in a fraction of a second. For example, when an AIS14140 steel that was heat treated to an ultimate tensile strength of 1500 MPa (218 ksi) was tested at an applied stress of 1109 MPa (161 ksi) (the approximate proportional limit of the material) while in contact with liquid indium at a temperature of 158 °C (316 °F) (indium melts at 156 °C, or 313 °F), crack formation required about 511 s. The crack then propagated and fractured the 5.84mm (0.23-in.) diam electropolished round bar specimen in only 0.1 s (Ref 189). In contrast, at 154 °C (309 °F), when the steel was in contact with solid indium, crack nucleation required 4.07 x 103 s (1.13 h), and failure required an additional 2.41 x 103 S (0.67 h) (Ref 189). Although gallium and mercury rapidly erabrittle aluminum alloys, all cases of LME and, Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user of Fracture / 29 especially, SME do not occur in such short time spans. The embrittlement of steels and titanium alloys by solid cadmium can occur over months of exposure; however, when long time spans are involved, the generation of hydrogen by the anodic dissolution of cadmium in a service environment can result in a hydrogen embrittlement assisted fracture. The magnitude of the applied stress, the strain rate, the amount of prior cold work, the grain size, and the grainboundary composition can also influence the rate of embrittlement. In general, higher applied stresses and lower strain rates promote embrittlement (Ref 112), while an increase in the amount of cold work reduces embrittlement (Ref 189). The reduction in embrittlement from cold work is believed to be due to the increase in the dislocation density within grains providing a large number of additional diffusion paths to dilute the concentration of embrittling atoms at grain boundaries. Smaller grain size should reduce embrittlement because of reduced stress concentration at grain-boundary dislocation pile-ups (Ref 189); however, in the embrittlement of Monel 400 by mercury, maximum embrittlement is observed at an approximate grain size of 250 Ixm (average grain diameter), and the embrittlement decreases for both the smaller and the larger grain sizes (Ref 112). The decrease in embrittlement at the smaller grain sizes was attributed to a difficulty in crack initiation, and for the larger grain sizes, the effect was due to enhanced plasticity (Ref 112). An example of a Monel specimen embrittled by liquid mercury is shown in Fig. 32. When fracture occurs by intergranular decohesion, the presence of such elements as lead, tin, phosphorus, and arsenic at grain boundaries can affect the embrittlement mechanism. The segregation of tin and lead at grain boundaries of steel can make it more susceptible to embrittlement by liquid lead, while a similar grainboundary enrichment by phosphorus and arsenic reduces it (Ref 190). Grain-boundary segregation of phosphorus has also been shown to reduce the embrittlement of nickel-copper alloys, such as Monel 400, by mercury (Ref 191, 192). It has been suggested that the beneficial effects of phosphorus are due to a modification in the grain-boundary composition that results in improved atomic packing at the boundary (Ref 192). The mechanisms proposed to explain the low-melting metal embrittlement process are often similar to those suggested for hydrogen embrittlement. Some of the mechanisms assume a reduction in the cohesive strength and enhancement of shear as a result of adsorption of the embrittling metal atoms (Ref 112, 114, 189, 193). It has also been suggested that the diffusion of a low melting point metal into the alloy results in enhanced dislocation nucleation at the crack tip (Ref 123,127, 194). A modified theory for crack initiation is based on stress and dislocation-assisted diffusion of the embrittling metal along dislocation networks and grain 30 / Modes Fig. 43 of Fracture Effect of hydrogen on fracture appearance in 13-8 PH stainlesssteel with a tensile strength of 1634 MPo (237 ksi). Top row: SEM fractographs of a specimennot embrittled by hydrogen. Bottom row: SEM fractographs of a specimen charged with hydrogen by plating without subsequentbaking. boundaries (Ref 189). The diffused atoms lower the crack resistance and make slip more difficult; when a critical concentration of the embrittling species has accumulated in the penetration zone, a crack initiates. The mechanism for the extremely rapid crack propagation for LME is not well understood. Diffusion processes are far too slow to transport the embrittling liquid metal to the rapidly advancing crack front. For embrittlement by liquid indium, it has been proposed that the transport occurs by a bulk liquid flow mechanism (Ref 189, 195); for the SME mode, the crack propagation is sustained by a much slower surface self-diffusion of the embrittling metal to the crack tip (Ref 189). Examples of low-melting metal embrittlement fractures are shown in Fig. 32 and 59 to 61. Figure 59 shows fractures in AISI 4140 steel resulting from testing in argon and in liquid lead. Figure 60 shows the embrittlement of a 7075-T6 aluminum alloy by mercury, and Fig. 61 shows the embrittlement of AISI 4140 steel by liquid cadmium. The articles "LiquidMetal Embrittlement" and "Embrittlement by Solid-Metal Environments" in Volume 11 of the 9th edition of Metals Handbook provide additional information on the effect of exposure to low melting point metals. Effect of State of Stress. This section will briefly discuss some effects of the direction of the principal stress as well as the state of stress, that is, uniaxial or triaxial, on the fracture modes of various metal systems. This section will not, however, present any mathematical fracture mechanics relationships describing the state of stress or strain in a material. The effects of stress will be discussed in general terms only. Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user The effect of the direction of the applied stress has been presented in the section "Dimple Rupture" in this article. Briefly, the direction of the principal stress affects the dimple shape. Stresses acting parallel to the plane of fracture (shear stresses) result in elongated dimples, while a principal stress acting normal to the plane of fracture results in primarily equiaxed dimples. Because the local fracture planes often deviate from the macroscopic plane and because the fracture is usually the result of the combined effects of tensile and shear stresses, it generally exhibits a variety of dimple shapes and orientations. The state of stress affects the ability of a material to deform. A change from a uniaxial to biaxial to triaxial state of stress decreases the ability of a material to deform in response to the Modes I 5 ~m I ,,~=:", 4 4 Quasi-cleavage fracture in a hydrogenembrittled AISI 4340 steel heat treated to an ultimate tensile strength of 2082 MPa (302 ksi). Source: Ref 138 applied stresses. As a result, metals sensitive to such changes in the state of stress exhibit a decrease in elongation or reduction of area at fracture and in extreme cases may exhibit a change in the fracture mode. The fcc metals, such as the aluminum alloys and austenitic stainless steels, and the hcp metals, such as the titanium and zirconium alloys, are generally unaffected by the state of stress. Although there can be a change in the nature of the dimples under biaxial or triaxial stresses, namely a reduction in dimple size and depth (Ref 196, 197), fcc and hcp metal systems usually do not exhibit a change in the mode of fracture. However, the bcc metals, such as most iron-base alloys and refractory metals, can exhibit not only smaller and shallower dimples but also a change in the fracture mode in response to the restriction on plastic deformation. This response depends on such variables as the strength level, microstructure, and the intensity of the triaxial stress. When a change in the fracture mode does occur as a result of a triaxial state of stress, such as that present near the root of a sharp notch, the mode of rupture can change from the normal dimple rupture to quasi-cleavage or intergranular decohesion (Ref 198). These changes in fracture mode are most evident in the general region of the fracture origin and may not be present over the entire fracture surface. Figure 62 shows the effect of a biaxial state of stress on dimples in a basal-textured Ti-6AI4V alloy. Under a biaxial state of stress, the size and the depth of the dimples decreased. For a pearlitic AISI 4130 steel (Ref 198) and a PH 13-8 precipitation-hardenable stainless steel, a triaxial state of stress resulting from the presence of a notch with a stress concentration factor of at least K t = 2.5 can change the fracture mode from dimple rupture to quasicleavage (Fig. 63). When a high-strength AISI 4340 steel is subjected to a triaxial stress, the mode of fracture can change from dimple rupture to intergranular decohesion. Effect o f S t r a i n R a t e . The strain rate is a variable that can range from the very low rates observed in creep to the extremely high strain rates recorded during impact or shock loading by explosive or electromagnetic impulse. Very low strain rates (about 10 - 9 to 10 - 7 s - l ) can result in creep rupture, with the (a) Fig° I 45 10 ~m I (b) of Fracture / 31 accompanying changes in fracture mode that have been presented in the section "Creep Rupture" in this article. At moderately high strain rates (about 102 s - l ) , such as experienced during Charpy impact testing, the effect of strain rate is generally similar to the effect of the state of stress, namely that the bcc metals are more affected by the strain rate than the fcc or the hcp metals. Because essentially all strain rate tests at these moderate strain rates are Charpy impact tests that use a notched specimen, the effect of strain rate is enhanced by the presence of the notch, especially in steels when they are tested below the transition temperature. A moderately high strain rate either alters the size and depth of the dimples or changes the mode of fracture from dimple rupture to quasicleavage or intergranular decohesion. For example, when an AISI 5140 H steel that was tempered at 500 °C (930 °F) was tested at Charpy impact rates, it exhibited a decrease in the width of the stretched zone adjacent to the precrack and an increase in the amount of intergranular decohesion facets (Fig. 64). The same steel tempered at 600 °C (1110 °F) showed no significant effect of the Charpy impact test (Ref 199). At very high strain rates, such as those observed during certain metal-shearing operations, high-velocity (100 to 3600 m/s, or 330 to 11 800 ft/s) projectile impacts or explosive rupture, materials exhibit a highly localized deformation known as adiabatic* shear (Ref *Adiabatic process is a thermodynamic concept where no heat is gained or lost to the environment. I 10 itm I Examples of hydrogen-embrittled steels. (a) Cleavage fracture in a hydrogen-embrittled annealed type 301 austenitic stainless steel. Source: Ref 98. (b) Intergranular decohesive fracture in an AISI 4130 steel heat treated to an ultimate tensile strength of 1281 MPa (186 ksi) and stressed at 980 MPa (142 ksi) while being charged with hydrogen. Source: Ref 111 Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user 32 / Modes of Fracture la) F |g./I.6 (al I 5 ixm I (bl Examples of hydrogen-embrittled titanium alloys. (a) Hydrogen embrittlement fracture in a Ti-8AI-1Mo-IV alloy in gaseous hydrogen. Note crack-arrest marks. Source: Ref 137. (b) Cleavage fracture in hydrogen-embrittled Ti-5AI-2.5Sn alloy containing 90 ppm H. Source: Ref 141 (b) I 10 ~m I (c) Fig. 47 Influence of heat treatment and resulting microstructure on the fracture appearance of a hydrogenembrittled Ti-6A-4V alloy. Specimens tested in gaseous hydrogen at a pressure of 1 atm. (a) Transgranular fracture in a specimen heat treated at 705 °C (1300 °F) for 2 h, then air cooled. (b) Intergranular decohesion along ~-I~ boundaries in a specimen heat treated at 955 °C (1750 °F) for 40 mln, then stabilized. (c) Coarse acicular structure resulting from heating specimen at 1040 °C (1900 °F) for 40 min, followed by stabilizing. The relatively flat areas of the terraced structure are the prior-~ grain boundaries. See text for a discussion of the microstructures of these specimens. Source: Ref 142 200-208). In adiabatic shear, the bulk of the plastic deformation of the material is concentrated in narrow bands within the relatively undeformed matrix (Fig. 65 to 67). Adiabatic shear has been observed in a variety of materials, including steels, aluminum and titanium alloys, and brass. These shear bands are believed to occur along slip planes (Ref 201, 202), and it has been estimated that under certain conditions, such as from the explosive-driven projectile impact of a steel target, the local strain rate within the adiabatic shear bands in the steel can reach 9 × 105 s -~ and the total strain in the band can be as high as 532% (Ref 204). An estimated 3 × 106-s - j strain rate has been reported for shear bands in a 2014-T6 aluminum alloy block impacted by a gun-fired (up to 900 m/s, or 2950 ft/s) steel projectile (Ref 205). The extremely high strain rates within the adiabatic shear bands result in a rapid increase in temperature as a large portion of the energy of deformation is converted to heat. It has been estimated that the temperature can go high enough to melt the material within the bands Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user (Ref 205,206). The heated material also cools very rapidly by being quenched by the large mass of the cool, surrounding matrix material; therefore, in quench-and-temper hardenable steels, the material within the bands can contain transformed untempered martensite. This transformed zone is shown schematically in Fig. 65(b). The hardness in the transformed bands is sometimes higher than can be obtained by conventional heat treating of the steel. This increase in hardness has been attributed to the additive effects of lattice hardening due to supersaturation by carbon on quenching and the extremely fine grain size within the band (Ref 203). However, for an A1SI 1060 carbon steel, the hardness of the untempered martensite bands was no higher than that which could be obtained by conventional heat treating (Ref 206). In both cases, the hardness of the adiabatic shear bands was independent of the initial hardness of the steel. For a 7039 aluminum alloy, however, the hardness of the shear bands was dependent on the hardness of the base material. The adiabatic shear bands in an 80HV material exhibited an average peak hardness of about 100 HV, while those in a 150-HV material had an average peak hardness of about 215 HV (Ref 208). For the Ti-6AI-4V STA alloy shown in Fig. 66, there was no significant difference in hardness between the shear bands and the matrix. In materials that do not exhibit a phase transformation, or if the temperature generated during deformation is not high enough for the transformation to occur, the final hardness of Modes Fig. 48 Hydrogen-embrittled 2124-UT aluminum alloy that shows no significant change in the fracture appearance. (a) Not embrittled. (b) Hydrogen embrittled. Source: Ref 99 lal Fig. (b) 49 Effect of heat treatment on the fracture appearance of a hydrogen-embrittled low-copper 7050 aluminum alloy. (a) Transgranular cleavagelike fracture in a peak-aged specimen. (b) Combined intergranular decohesion and transgranular cleavagelike fracture in an underaged specimen. Source: Ref 106 the adiabatic shear band is the net result of the competing effects of the increase in hardness due to the large deformation and the softening due to the increase in temperature. The width of the adiabatic shear bands depends on the hardness (strength) of the material (Ref 206, 208). Generally, the harder the material, the narrower the shear bands. In a 7039 aluminum alloy aged to a hardness of 80 HV, the average band width resulting from projectile impact was 90 Ixm, while in a 150-HV material, the band width was only 20 Ixm (Ref 208). The average width of the shear bands observed in a Ti-6A1-4V STA alloy (average hardness, 375 HV~kg) was 3 to 6 txm. When an adiabatic shear band cracks or separates during deformation, the fractured surfaces often exhibit a distinct topography referred to as knobbly structure (Ref 205-208). The name is derived from the surface appearance, which resembles a mass of knoblike structures. The knobbly structure, which has been observed in 2014-T6 and 7039-T6 aluminum alloys, as well as in an AISI 4340 steel (Fig. 67) and AISI 1060 carbon steel, is believed to be the result of melting within the shear bands (Ref 205, 206). Although the cracked surfaces of adiabatic shear bands can exhibit a unique appearance, adiabatic shear failure is easiest to identify by metallographic, rather than fractographic, examination. Effect of T e m p e r a t u r e . Depending on the material, the test temperature can have a significant effect on the fracture appearance and in Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user of Fracture / 33 many cases can result in a change in the fracture mode. However, for materials that exhibit a phase change or are subject to a precipitation reaction at a specific temperature, it is often difficult to separate the effect on the fracture due to the change in temperature from that due to the solid-state reactions. In general, slip, and thus plastic deformation, is more difficult at low temperatures, and materials show reduced ductility and an increased tendency for more brittle behavior than at high temperatures. A convenient means of displaying the fracture behavior of a specific material is a fracture map. When sufficient fracture mode data are available for an alloy, areas of known fracture mode can be outlined on a phase diagram or can be plotted as a function of such variables as the test temperature and strain rate (Fig. 68). Similar maps can also be constructed for lowtemperature fracture behavior. Effect of Low Temperature. Similar to the effect of the state of stress, low temperatures affect the bcc metals far more than the fcc or hcp metal systems (see the section "Effect of the State of Stress" in this article). Although lower temperatures can result in a decrease in the size and depth of dimples in fcc and hcp metals, bcc metals often exhibit a change in the fracture mode, which generally occurs as a change from dimple rupture or intergranular fracture to cleavage. For example, a fully pearlitic AISI 1080 carbon steel tested at 125 °C (255 °F) showed a fracture that consisted entirely of dimple rupture; at room temperature, only 30% of the fracture was dimple rupture, with 70% exhibiting cleavage. At - 1 2 5 °C ( - 1 9 5 °F), the amount of cleavage fracture increased to 99% (Ref 210). This transition in fracture mode is illustrated in Fig. 69. Charpy impact testing of an AISI 1042 carbon steel whose microstructure consisted of slightly tempered martensite (660 HV) as well as one containing a tempered martensite (335 HV) microstructure at 100 °C (212 °F) and at - 196 °C ( - 3 2 0 °F) produced results essentially identical to those observed for the AISI 1080 steel. In both conditions, the fracture mode changed from dimple rupture at 100 °C (212 °F) to cleavage at - 1 9 6 °C ( - 3 2 0 °F), as shown in Fig. 70. Similar changes in the fracture mode, including a change to quasi-cleavage, can be observed for other quench-and-temper and precipitation-hardenable steels. A unique effect of temperature was observed in a 0.39C-2.05Si-0.005P-0.005S low-carbon steel that was tempered 1 h at 550 °C (1020 °F) to a hardness of 30 HRC and Charpy impact tested at room temperature and at - 8 5 °C ( - 1 2 0 °F) (Fig. 71). In this case, the fracture exhibited intergranular decohesion at room temperature and changed to a combination of intergranular decohesion and cleavage at - 8 5 °C ( - 1 2 0 °F). This behavior was attributed to the intrinsic reduction in matrix toughness by the silicon in the alloy, because when nickel is substituted for the silicon the matrix toughness 34 / Modes of Fracture (a) I 25 ~m I (bl Fig. 50 Stress-corrosion fractures in HY-180 steel with an ultimate strength of 1450 MPa (210 ksi). The steel was tested in aqueous 3.5% sodium chloride at an electrochemical potential of E = - 0 . 3 6 to - 0 . 8 2 VsHE (SHE, standard hyd,r.ggen electrode). Intergranular decohesion is more pronounced at lower values of stress intensity, K~ = 57 MPaVm (52 ksi~n~n.) (a), than at higher values, KI = 66 MPaN/'-mm(60 ksiN/~n.) (b). Source: Ref 154 room temperature to 600 °C (1110 °F) (Ref 213). Figure 73 shows the effect of temperature on the fracture mode of an ultralow-carbon steel. The steel, which normally fractures by dimple rupture at room temperature, fractured by intergranular decohesion when tensile tested at a strain rate of 2.3 × 10 -2 s -L at 9 5 0 ° C (1740 °F). The change in fracture mode was due to the precipitation of critical submicronsize MnS precipitates at the grain boundaries. This embrittlement can be eliminated by aging at 1200 °C (2190 °F), which coarsens the MnS precipitates (Ref 209). A similar effect was observed for Inconel X-750 nickel-base alloy that was heat treated by a standard double-aging process and tested at a nominal strain rate of 3 × 10 5 s ~ at room temperature and at 816 °C (1500 °F). The fracture path was intergranular at room temperature and at 816 °C (1500 °F), except that the room-temperature fracture exhibited dimples on the intergranular facets and those resulting from fracture at 816 °C (1500 °F) did not (Fig. 74). The fracture at room temperature exhibited intergranular dimple rupture because the material adjacent to the grain boundaries is weaker due to the depletion of coarse 3" precipitates. The absence of dimples at 816 °C (1500 °F) was the result of intense dislocation activity along the grain boundaries, producing decohesion at M 2 3 C 6 carbide/matrix interfaces within the boundaries (Ref 214). A distinct change in fracture appearance was also noted during elevated-temperature tensile testing of Haynes 556, which had the following composition: Element (al (b) I 50 ~m I Fig. 51 Stress-corrosion fractures in a 25% cold-worked type 316 austenitic stainless steel tested in a boiling (154 °C, or 309 °F) aqueous 44.7% magnesium chloride solution. At low (14 MPa~'-mm, or 12.5 ksi ~ - ~ ) KI values, t~:~ fracture exhibits a combination of cleavage and intergranular decohesion (a). At higher (33 MPaV m, or 30 ksiVin.) values of KI the principal mode of fracture is intergranular decohesion (b). Source: Ref 181 is increased and no cleavage is observed (Ref 211). The temperature at which a sudden decrease in the Charpy impact energy occurs is known as the ductile-to-brittle transition temperature for that specific alloy and strength level. Charpy impact is a severe test because the stress concentration effect of the notch, the triaxial state of stress adjacent to the notch, and the high strain rate due to the impact loading combine to add to the reduction in ductility resulting from the decrease in the testing temperature. Although temperature has a strong effect on the fracture process, a Charpy impact test actually measures the response of a material to the combined effect of temperature and strain rate. The effects of high temperature on fracture are more complex because solid-state reactions, such as phase changes and precipitation, are more likely to occur, and these changes affect bcc as well as fcc and hcp alloys. As shown in Fig. 72, the size of the dimples generally increases with temperature (Ref 209, 212, 213). The dimples on transgranular fractures and those on intergranular facets in a 0.3C- 1Cr- 1.25Mo-0.25V-0.7Mn-0.04P steel that was heat treated to an ultimate strength of 880 MPa (128 ksi) show an increase in size when tested at temperatures ranging from Downloaded from https://dl.asminternational.org/handbooks/chapter-pdf/268655/a0001831.pdf by Washington State University user Composition, % Iron ................................. 28.2 Chromium ............................ Nickel ............................... Cobalt ............................... Tungsten ............................. Molybdenum .......................... 21.5 22.2 19.0 2.9 2.9 Tantalum ............................. Manganese ............................ 0.8 1.4 Silicon ............................... Copper .............................. Nitrogen ............................. 0.5 0.1 0.1 Three specimens were tested at a strain rate of approximately 1 S - 1 at increasin...
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Materials characterization laboratory
Lab 2: SEM Fractography.
Abstract

In this laboratory, we worked with the concept of fractography. This concept is related to
the study and analysis of fractures on the surface of the materials. Through this type of
research, useful information about the operation, structure, and stability of certain
materials obtained. Besides, the importance of imaging fractures of surfaces in different
elements causes of the fissures and their mechanisms evidenced. Three different samples
of materials used: 1) Low carbon steel (A36), 2) Mild carbon steel (A36), and High carbon
steel (A36). In this way the different fractures were observed and sketched in the samples
with the naked eye and using an optical microscope, from this, representative images of the
fractured surfaces obtained, finally, the type of failure that occurs in the observation
determined each sample.
Introduction

Fractography, the study of fractures materials is related to the methods by which analyzes
of the characteristics of a split in the material performed to know its causes and
mechanisms to understand the reasons that cause the failure or situation. This analysis is in

a macroscopic and microscopic manner; the macroscopic, as the name implies, the eye is
used to observe characteristics. [1]
On the contrary, the analysis of tiny fracture surfaces performed with high magnifications
and different microscopes used. Typical details of the macroscopic characterization are the
radial marks, the fish-scratch marks, and the beach lines. Concerning the microscopic study,
there are gaps, facets of cleavage, intergranular factors, and stretch marks. The
identification of the origin of the
A fracture will also be an essential as...


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