21-Mat-B6 Ceramic Materials · May 2016
Nivaar worked solution (AI-drafted; not reviewed by a licensed engineer)
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.
| Grade | C | Cr | Ni | W | V | Cu | Mn | S | P |
|---|---|---|---|---|---|---|---|---|---|
| T1 | 0.65–0.80 | 3.75–4.00 | 0.3 | 17.25–18.75 | 0.9–1.3 | 0.25 | 0.1–0.4 | 0.03 | 0.03 |
T1 carries very large amounts of strong carbide-forming elements — tungsten and vanadium especially, with chromium contributing as well — which form coarse, extremely thermally stable alloy carbides (M6C-type W-rich carbide, MC-type VC) during solidification and subsequent processing. These carbides have very LOW solubility in austenite at the moderate austenitizing temperatures (800–900 °C) used for plain-carbon or low-alloy steels, and they do not dissolve appreciably at such temperatures. To get enough W, Cr, V and C into solid solution in the austenite — essential so that on quenching and tempering they can re-precipitate as the FINE secondary carbides responsible for this steel's characteristic secondary hardening and hot hardness (red hardness) — the steel must be heated close to its incipient-melting temperature, roughly 1250 °C for T1, where the coarse carbides finally dissolve substantially. Austenitizing too low leaves the alloy content locked up in undissolved primary carbides, starving the matrix of the alloying needed for both adequate hardenability and for the secondary-hardening response on tempering.
Once the heavy alloy content (W, Cr, V) is taken into solution by the high-temperature austenitizing soak, that SAME dissolved alloy content dramatically depresses the pearlite/bainite "nose" of the steel's TTT/CCT diagram to very long times, i.e. it gives the steel extremely high hardenability. With the nose pushed out that far, even the comparatively slow, uniform cooling rate available from still air (or a gentle forced-air fan, which mainly improves temperature UNIFORMITY through the section rather than raw cooling speed) is already faster than the now-very-long critical cooling time needed to avoid the nose — so the steel can be taken essentially straight to martensite without ever needing a severe oil or water quench. This is exactly why high-speed and other highly-alloyed tool steels are classed as "air-hardening": a slow, gentle air cool is not merely tolerated, it is the RECOMMENDED practice, because it also minimizes the thermal-gradient-driven distortion and quench-cracking risk that a fast liquid quench would otherwise impose on the complex, often thin-sectioned tool geometries these steels are used for.
As-quenched high-speed steel martensite contains a substantial fraction of RETAINED austenite (not fully martensite), because the heavy alloying that depresses the pearlite nose (needed for 5.2) also depresses $M_s$/$M_f$ well below room temperature, so the quench never fully completes the austenite-to-martensite transformation. Each temper cycle serves two compounding purposes at once: (1) it precipitates a dense dispersion of very fine, coherent secondary alloy carbides (the "secondary hardening" reaction, from the same dissolved W/V/Cr that was put into solution during austenitizing) from the primary martensite; and (2) the resulting compositional and volume change, together with the temper's own heating/cooling cycle, converts a substantial fraction of the previously RETAINED austenite into fresh, untempered (and therefore brittle) martensite on cooling back to room temperature after each cycle. A single temper therefore leaves behind a mixture of properly tempered martensite AND freshly formed, still-untempered martensite from the just-converted retained austenite. A SECOND temper is needed to temper THAT freshly formed martensite; and because even the second temper still converts a further (smaller) increment of any remaining retained austenite into new untempered martensite, a THIRD temper is the standard practice to ensure the retained-austenite-to-martensite conversion is driven essentially to completion and every generation of martensite formed along the way is properly tempered, leaving a stable, fully tempered, minimal-retained-austenite final microstructure.
The very high final hardness is the SUM of two mechanisms set up by parts 5.1–5.3, not a single effect. First, and dominant: secondary hardening — during tempering, the alloy content dissolved by the 1250 °C austenitizing soak precipitates as an extremely fine, dense dispersion of coherent (or semi-coherent) M2C/MC alloy carbides directly WITHIN the martensite laths. These particles are far too fine and closely spaced for a dislocation simply to bow between them at low stress (Orowan bypass requires a stress inversely proportional to the interparticle spacing), so the fine secondary-carbide dispersion is a strong, direct obstacle to dislocation motion — strong enough that hardness actually RISES with tempering temperature (up to a secondary-hardening peak, typically 500–570 °C) instead of monotonically falling the way a plain-carbon steel's tempered martensite does. Second, each of the three tempers converts a further increment of retained austenite (5.3) into FRESH, untempered, high-carbon martensite, which is itself very hard; by the third temper this conversion is essentially complete and the microstructure is a uniform matrix of secondary-carbide-strengthened tempered martensite with only a small residual retained-austenite fraction. The combination — fine coherent-carbide precipitation strengthening plus a matrix that is almost entirely tempered martensite rather than partly soft retained austenite — is what gives high-speed tool steel its characteristic very high hardness AND, unusually, its hot hardness (red hardness): the same fine alloy carbides that resist dislocation motion at room temperature also resist coarsening at the elevated temperatures a cutting edge reaches in service.
