21-Mat-B7 Structure and Properties of Polymers · December 2016
Nivaar worked solution (AI-drafted; not reviewed by a licensed engineer)
Paper format. National Exams, December 2016 — 10-Met-B7, Physical Metallurgy of Non-Ferrous Metals and Alloys. Three hours, closed book, approved Casio/Sharp calculator only. Eight questions of 20 marks each; the rubric states that any five questions constitute a complete paper (100 marks total) 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. The rubric explicitly notes that most questions require an essay-format answer and that clarity and organization are marked, so the answers below are written as structured prose rather than as note form.
Nothing on this paper is a polymer question; the syllabus actually examined is the physical metallurgy, strengthening and heat treatment of non-ferrous engineering alloys — aluminum, magnesium, copper-base alloys (brasses and bronzes), nickel- and cobalt-base superalloys, titanium, and the refractory/noble metals and intermetallic compounds.
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.
"Hardenable" aluminum alloys are the age-hardening families — 2xxx (Al–Cu), 6xxx (Al–Mg–Si) and 7xxx (Al–Zn–Mg–Cu) — whose strength comes from a fine precipitate dispersion rather than from cold work alone. All three follow the same three-stage sequence.
Stage 1 — solution heat treatment. The alloy is soaked at a temperature above the solvus (but below the eutectic/incipient-melting temperature) for the relevant alloying elements — for Al–4%Cu, roughly 515–540 °C — long enough for the equilibrium second-phase particles present after casting or prior working (e.g. $\theta$, Al2Cu, in a 2xxx alloy; Mg2Si in a 6xxx alloy; MgZn2 in a 7xxx alloy) to dissolve completely. Resulting microstructure: a single-phase, solute-saturated $\alpha$ solid solution, with the coarse as-cast/as-worked second-phase particles gone.
Stage 2 — quenching. The alloy is cooled rapidly (typically a cold-water quench) to room temperature, fast enough that solute atoms have no time to diffuse to nucleation sites and precipitate on the way down. Resulting microstructure: a supersaturated solid solution (SSSS) — the same single $\alpha$ phase as after Stage 1, but now thermodynamically unstable at room temperature — together with a high concentration of quenched-in vacancies that will accelerate solute diffusion during the next stage.
Stage 3 — aging (precipitation heat treatment). The supersaturated solution is held at room temperature (natural aging) or at an elevated temperature, typically 120–190 °C depending on the alloy (artificial aging), allowing the excess solute to precipitate through a sequence of metastable, progressively less coherent transition phases before reaching the stable equilibrium phase (e.g., for Al–Cu: GP zones → $\theta''$ → $\theta'$ → equilibrium $\theta$). Resulting microstructure: peak strength is reached while the precipitates are still small, coherent or semicoherent, and closely spaced — the microstructure a designer actually wants in service — because dislocations must then either shear the particles or bow between them (Orowan mechanism), both of which require a much higher stress than moving through the solute-saturated but otherwise unobstructed lattice left after Stage 2. Continuing to age past this point (overaging) coarsens the particles into the widely spaced, incoherent equilibrium phase and strength falls again.
Factor 1 — second-phase particle (Zener) pinning. Dispersoid particles left in the microstructure — Al-Mn/Al-Fe-Mn constituents and dispersoids typical of structural sheet alloys — exert a retarding (Zener drag) pressure on a migrating grain boundary because detaching the boundary from a particle it intersects costs interfacial energy. Growth stops once the driving pressure from boundary curvature falls to the pinning pressure, giving a limiting grain size $D_{max}=4r/3f$, where $r$ is the mean particle radius and $f$ the volume fraction. A finer, more closely spaced dispersoid population (achieved by the homogenization/pre-heat practice before hot rolling) therefore locks in a finer ultimate grain size.
Factor 2 — annealing temperature and time (thermal schedule). For a fixed particle population, the anneal soak temperature and duration set both how many new grains nucleate during recrystallization (a higher temperature and faster heat-up favour a higher nucleation rate relative to growth rate, giving a finer initial recrystallized grain size) and how far normal grain growth is allowed to proceed afterward, since boundary mobility is itself thermally activated. A schedule held just long enough to complete recrystallization, then cooled promptly, yields a finer ultimate grain size than one that overshoots into the grain-growth regime; this is the second lever (alongside particle pinning) a mill uses to hit a target grain size for a structural sheet's formability and surface (orange-peel) requirements.
