21-Mat-A5 Phase Transformations and Thermal Treatment · December 2017
Nivaar worked solution (AI-drafted; not reviewed by a licensed engineer)
Paper format. National Exams, December 2017 — 10-Met-A5, Mechanical Behaviour and Fracture of Materials. Three hours, closed book, any Casio- or Sharp-approved calculator permitted. Eight questions of 20 marks each; the rubric states that five questions constitute a complete paper and that only the first five appearing in the answer book are marked. All eight are answered here, because this set is a study resource rather than an exam script. Several sub-parts explicitly call for an essay-format answer, and the rubric rewards clarity and organisation, so those answers are written as structured prose rather than as note form.
Nothing on the paper is a phase-transformation or heat-treatment question in the TTT/CCT, hardenability or tempering sense; the syllabus actually examined is dislocation theory, slip and twinning, strengthening mechanisms, creep, fatigue, toughness and fracture mechanics, and deformation processing.
Reference texts. The answers below are keyed to the works normally recommended for this syllabus code:
Question text not reproduced: the examination questions are © Engineers and Geoscientists BC. Open the official past paper (linked at the top of this page) to read the question, then follow the worked solution below.
All three candidate causes leave physically distinct evidence on the fracture surface, because each proceeds by a different micromechanism, and a fractographic investigator's task is to read that evidence under the SEM.
(i) Slow crack growth fatigue. Fatigue fracture surfaces show macroscopic beach marks (concentric arc markings) tracing the successive crack-front positions as growth rate and load history varied, and, at higher magnification, closely spaced parallel striations — each one the fingerprint of a single load cycle, with the local striation spacing tracking $da/dN$ at that point via the Paris law of Question 5. The origin is typically a single, small, well-defined point at a stress concentration (a fillet, fastener hole or surface scratch), and the final overload region shows a comparatively small area of dimpled or cleavage fast fracture where the crack outran the section's remaining capacity — small because fatigue is typically detected or causes failure only after most of the section is already cracked.
(ii) Slow crack growth by corrosion (stress-corrosion cracking). SCC surfaces are typically intergranular (following grain boundaries embrittled by the corrosive species) or, in some alloy/environment systems, transgranular with a characteristic branched, "river" crack-path morphology visible even macroscopically, since the crack repeatedly re-nucleates and searches for the locally most susceptible path. Secondary cracking branching off the main fracture, corrosion products or pitting at the origin, and the complete absence of fatigue striations are the diagnostic tells; unlike fatigue, no cyclic load is required, so SCC can occur under a static sustained stress in the presence of a specific corrodent (e.g. chloride SCC of austenitic stainless steel).
(iii) Impact with the ground. An overload fracture from a single high-rate impact shows dimpled rupture (microvoid coalescence, for a ductile material) or cleavage facets (for a brittle one or at high strain rate/low temperature) over the entire fracture surface, with no beach marks, no striations and no progressive crack-front evidence at all, because the entire crack propagated in a single, essentially instantaneous event. Shear lips at 45$^\circ$ to the loading axis around the fracture perimeter, gross plastic deformation of the surrounding structure, and multiple, randomly oriented secondary fractures throughout the airframe (rather than one isolated component) further support a single overload event rather than a progressive one.
The unifying diagnostic logic is that progressive mechanisms (fatigue, SCC) leave a record of many discrete increments on the surface itself, while a single-event mechanism (impact overload) does not; and among the progressive mechanisms, the striation spacing and crack-path character (transgranular fatigue vs. often-intergranular SCC) distinguish a mechanical from an environmental cause.
Creep rupture is intergranular because grain boundaries are simultaneously the fastest diffusion path and the weakest mechanical link at the elevated homologous temperatures where creep operates. Two effects combine. First, sustained load causes grain-boundary sliding as the diffusional/dislocation-climb creep mechanisms of Question 3 proceed; sliding must be accommodated at triple points and at any boundary irregularity (a second-phase particle, a ledge), and where accommodation fails, a cavity nucleates. Second, once nucleated, cavities grow preferentially by the same vacancy diffusion that drives Coble/Nabarro-Herring creep, because the boundary itself is the short diffusion path feeding vacancies to the cavity; boundaries lying roughly normal to the applied tensile stress are the ones held most open and therefore grow cavities fastest. As loading continues, cavities nucleate progressively along these transverse boundaries, link up into boundary-length cracks (often visible metallographically as "creep cavitation" or, in advanced stages, "wedge cracking" at triple points), and finally coalesce into the through-section intergranular fracture characteristic of creep rupture — in sharp contrast to the transgranular, often high-ductility fracture of a room-temperature tensile overload of the same alloy.
Intrusions and extrusions form at the free surface of a metal subjected to cyclic plastic strain, even at nominal stresses well below the monotonic yield strength, through the build-up of persistent slip bands (PSBs). Within favourably oriented surface grains, cyclic loading concentrates plastic strain into narrow bands of very high dislocation activity (PSBs), distinct from the surrounding matrix, because the dislocation substructure that develops (typically a ladder-like arrangement of dipole walls) offers an easy, repeatable glide path once established. Critically, glide on the forward and reverse halves of the load cycle is not perfectly reversible: cross-slip and different obstacles encountered on the way out versus the way back mean that a small net increment of slip offset accumulates at the free surface with every cycle, rather than the material returning exactly to its starting configuration. Over many thousands of cycles this irreversible ratcheting builds a visible surface topography — a thin ridge of extruded material (an extrusion) protruding from the surface where the net slip pushed material outward, paired with an adjacent narrow surface notch (an intrusion) where material was drawn in. The intrusion is geometrically a sharp, atomically fine notch with a stress concentration of its own, and it is precisely this self-generated notch — not any pre-existing surface flaw — that nucleates the fatigue crack in an otherwise defect-free, polished specimen, which is why PSB intrusion/extrusion formation is regarded as the crack-initiation mechanism for fatigue in ductile metals.
