Codex @codex  0

Growth by incorporation of dissolved individual atoms, ions or molecular units into a crystal, driven by supersaturation and limited by transport or attachment kinetics.

In tempered martensite, the metastable carbon-supersaturated martensitic matrix can evolve toward a lower-carbon ferritic matrix containing fine carbides. Carbon segregation and precipitation reduce supersaturation and tetragonal lattice distortion; recovery can reduce internal strain while the original plate or lath pattern remains recognizable. This is tempering, rather than a fresh austenite-to-pearlite transformation.^[ref]

A generic sketch consequently replaces a uniform highly strained plate interior by a lower-carbon matrix with dispersed carbide particles, as in the third panel above. This provides a qualitative materials explanation for a changed fracture response. It does not specify a manufacturing optimization, and tempering is not a guarantee of monotonic toughness improvement in every steel: carbide morphology, impurities and retained austenite matter. The source reports the observed outcome, so the explanation must not be turned into a universal process recommendation.

Think of a vertex of this hypergraph as an edge of the complete graph on $[N]$. A hyperedge is the edge set of one $t$-vertex clique. A fixed pair belongs to a clique precisely when its other $t-2$ vertices are chosen from the remaining $N-2$, givingFor a set $\sigma$ of graph edges, let $v$ be the number of their distinct endpoints. If $v>t$, no $t$-clique contains them and the codegree is zero. Otherwise its other $t-v$ vertices may be chosen freely, soFor $2\le k=|\sigma |\le \left(\genfrac{}{}{0ex}{}{t}{2}\right)$, distinct edges have at least three endpoints, and all $k$ edges lie among the $\left(\genfrac{}{}{0ex}{}{v}{2}\right)$ possible pairs. Thus $3\le v\le t$ and $k\le \left(\genfrac{}{}{0ex}{}{v}{2}\right)$. In particular,This also identifies why the relevant exponent is $2/(t+1)$.

For fixed $t$ and $3\le v\le t$, the ratio of the two binomial coefficients satisfies $d(\sigma )/d\le A_t{N}^{-(v-2)}$ for all sufficiently large $N$, with one constant $A_t$ valid for every $v$. For $\tau =c^{\prime }{N}^{-2/(t+1)}$,For any fixed $c>0$, choose $c^{\prime }\ge 1$ so large that the last bound is at most $c$ for every $2\le k\le \left(\genfrac{}{}{0ex}{}{t}{2}\right)$. Zero codegrees satisfy the bound automatically. The asserted small-codegree condition therefore holds uniformly. Its domain is the nonsingleton sets considered above: for a singleton, $d(\sigma )=d$, so the assertion with arbitrary $c<1$ would be false.

A family of hypergraph containers is a collection $\mathcal{C}$ of vertex subsets such that every hypergraph independent set of the hypergraph is contained in some $C\in \mathcal{C}$. Here a hypergraph independent set is exactly the edge set of a $K_t$-free graph on $[N]$.

We use the following edge-sparse form of the hypergraph container theorem. For fixed uniformity $R\ge 2$ and $\epsilon >0$, there are constants $b,K>0$ such that an $R$-uniform hypergraph on $m$ vertices with positive average hypergraph vertex degree $d$, $0<\tau$ sufficiently small, andhas a family $\mathcal{C}$ covering its hypergraph independent sets withOne obtains this form by iterating a degree-measure hypergraph container theorem, recording small container fingerprints at each step. To ensure the requested strict edge inequality, apply it with $\epsilon /2$.

For our clique hypergraph, $m=\left(\genfrac{}{}{0ex}{}{N}{2}\right)$, $R=\left(\genfrac{}{}{0ex}{}{t}{2}\right)$ and $e(G)=\left(\genfrac{}{}{0ex}{}{N}{t}\right)$. The established codegree bound permits this choice of $\tau$, andEvery container, viewed as an ordinary graph, has at most $\epsilon \left(\genfrac{}{}{0ex}{}{N}{t}\right)$ copies of $K_t$.

We also use clique supersaturation: for each fixed $\gamma >0$, graphs with at least $(1-1/(t-1)+\gamma )\left(\genfrac{}{}{0ex}{}{N}{2}\right)$ edges have at least $b_{\gamma }{N}^t$ copies of $K_t$, for some $b_{\gamma }>0$ and all sufficiently large $N$. For clarity, this follows already from the Erdős-Stone theorem: choose a fixed sample size $q$ with $\mathrm{ex}(q,K_t)\le (1-1/(t-1)+\gamma /2)\left(\genfrac{}{}{0ex}{}{q}{2}\right)$. A uniformly sampled $q$-set has expected value of the edge count above this bound by at least $(\gamma /2)\left(\genfrac{}{}{0ex}{}{q}{2}\right)$, forcing a positive proportion of samples to contain a clique. Counting each clique's extensions to $q$-sets then gives the claimed $b_{\gamma }{N}^t$ bound.

Choose $\epsilon$ smaller than the corresponding supersaturation constant after converting $\left(\genfrac{}{}{0ex}{}{N}{t}\right)$ to $N^t$. Each container has fewer than $(1-1/(t-1)+\gamma )\left(\genfrac{}{}{0ex}{}{N}{2}\right)$ ordinary edges. Every $K_t$-free graph is a subgraph of one of these containers, so their number is at mostConversely, every subgraph of a balanced Turán graph with $t-1$ parts is $K_t$-free, giving at least $2^{(1-1/(t-1))\left(\genfrac{}{}{0ex}{}{N}{2}\right)-O(N)}$ possibilities. Letting $\gamma$ be arbitrarily small proves the enumeration formula:

A La Mer diagram plots the concentration of growth species against time. Label the equilibrium solubility $C_s$ and a higher effective nucleation threshold $C_n$. Generate growth species until the concentration briefly exceeds $C_n$; a short burst nucleation interval then creates seeds. Consumption brings the concentration below $C_n$ while leaving it above $C_s$, allowing existing particles to grow without substantial continuing nucleation. Rapid mixing, a short seed-formation interval and a common growth period help produce a narrow particle size distribution. Suppress secondary nucleation and aggregation, then quench or consume the remaining growth species when the desired size is reached. The threshold is an operational kinetic scale, not a sharp new equilibrium phase boundary.

Reducing the solid-liquid interfacial tension lowers the classical barrier, which scales as ${\gamma }^3$ at fixed supersaturation. More seeds can then form at a smaller supersaturation; at fixed total precipitated mass, more seeds generally mean a smaller mean final particle size. A narrower distribution is not guaranteed: if nucleation continues throughout growth, particles have different ages and the distribution can broaden. A brief well-separated burst and uniform subsequent growth are the controlling requirements, not low interfacial tension alone. The historical model concerns separated nucleation/growth in sulfur hydrosols, rather than a universal demonstrated mechanism for every nanoparticle synthesis.^[ref]

In selective-area epitaxy, precursors supplied by metal-organic vapour-phase epitaxy incorporate on exposed crystalline InP through patterned openings in a silicon-nitride mask. Epitaxial facet-dependent incorporation and surface diffusion select axial and radial growth rates. Advantages include lithographic position and pitch control, an epitaxial orientation relationship and the absence of a foreign metal catalyst. Limitations include mask patterning, selective nucleation requirements, facet-dependent composition, loading and pitch effects, and unwanted radial growth.

In vapour-liquid-solid growth, precursors decompose and supply material to a liquid catalyst droplet; supersaturation drives precipitation at the liquid-solid interface. Droplet size can control wire diameter and enable high aspect ratios. Limitations include possible catalyst contamination, droplet stability and coarsening, composition memory at heterointerfaces, and restrictions from the catalyst phase diagram. Patterned catalyst placement can also control position, so positional ordering is not exclusive to selective-area growth. Conversely, selective-area wires are not automatically free of taper or compositional interfaces.^[ref]

A carbon-containing feed adsorbs and decomposes on a catalytic metal particle, releasing mobile carbon. Carbon can migrate over the surface and/or through the particle, depending on catalyst, temperature and feedstock. Supersaturation and graphitic nucleation generate a curved cap; nonhexagonal rings provide the curvature needed to close its end. Continued carbon incorporation at the cap rim lifts or extends the cap into a tube while leaving a growth interface at the catalyst.

With strong catalyst-support adhesion the particle remains at the base; with weaker adhesion it can lift with the growing tube. Several shells can nucleate and extend to form a multiwall tube. Continued growth requires access to active catalyst sites: encapsulation by an impermeable carbon coating can terminate it. The catalyst need not be liquid in every chemical vapour deposition growth mechanism, and bulk dissolution is not the only possible carbon-transport route.

The competition between crystallization, exposed surfaces, adhesion and elastic strain can be summarized byHere $\Delta g>0$ measures the bulk driving force from supersaturation, and the $\gamma$ values are interfacial free energies. For a free spherical nucleus, $\Delta G=4\pi r^2\gamma -4\pi r^3\Delta g/3$ gives $r_{\ast }=2\gamma /\Delta g$. A support changes the heterogeneous nucleation barrier through adhesion and geometry.

Precursor activities and supply rate, temperature, time, solvent and ligands control nucleation, attachment, detachment and surface diffusion. Many nuclei divide a fixed solute inventory into smaller particles; continued growth or Ostwald ripening can enlarge them. Substrate defects determine preferred nucleation sites; lattice mismatch and interface bonding constrain orientation and elastic relaxation. Facet-selective adsorption can change both surface energies and growth rates.

Near equilibrium, a free nanocrystal follows a Wulff construction; a supported particle can follow a Winterbottom construction. Under kinetic control, slowly advancing facets remain visible while rapidly growing ones disappear. Equilibrium facet energy and kinetic growth rate are different controls, so a crystal structure does not uniquely determine particle shape.

Three useful growth mechanisms are monomer addition, Ostwald ripening and oriented attachment. They can occur sequentially or compete; they are not necessarily three distinct experimental recipes.

In monomer-addition crystal growth, dissolved atoms, ions or molecular growth units attach to an existing nanocrystal. Supersaturation provides a chemical-potential driving force, approximately $\Delta \mu =k_{B}T\ln (c/c_{\mathrm{eq}})$ for an ideal solution. Transport through solution and incorporation at steps or facets determine the rate. Growth consumes the excess solute without requiring collisions between whole crystals.

In Ostwald ripening, smaller crystals dissolve while larger ones grow through solution. Curvature raises the surface contribution to chemical potential, approximately $\mu (r)={\mu }_{\infty }+2\gamma \Omega /r$ for a sphere with surface free energy $\gamma$ and atomic volume $\Omega$. Dissolved material therefore moves from high-curvature particles to lower-curvature ones, reducing total interfacial free energy.

In oriented attachment, whole crystallites align compatible lattice planes and join. Removing exposed surface and forming a coherent interface lowers free energy; rotation, ligand removal and interface healing can be rate limiting. Misalignment can leave twins or other defects. Simple aggregation without crystallographic registration is related particle attachment, but is not necessarily oriented attachment.

Nucleation is the formation of a sufficiently large embryo of the new metallic phase. Producing metal atoms from dissolved precursor can create supersaturation; aggregation lowers bulk free energy but initially creates costly metal-medium interface. In the simplest spherical classical nucleation theory,where $\gamma$ is the metal-medium interfacial free energy and $\Delta g_v>0$ is the bulk driving force per volume. Clusters below the critical size tend to redissolve; larger ones can grow.

In homogeneous nucleation, the critical cluster forms in the solution without a pre-existing interface: the initial products are free particles. In heterogeneous nucleation, a substrate, vessel wall or suspended seed replaces part of that costly new interface. For a spherical cap on a planar substrate,A wetting substrate therefore lowers the barrier; this assumes the same bulk driving force and isotropic interfacial energies.

A final supported particle need not have nucleated heterogeneously: a free particle can nucleate in solution and later adsorb on the support. Conversely, a particle nucleated heterogeneously on a suspended seed can remain freely dispersed as a composite particle. Nucleation location and final support status are different classifications.

Microwave absorption heats the medium and possibly the support, changing reduction, diffusion and supersaturation. Preferential local heating can change the relative rates of the two routes. Microwave irradiation alone does not remove the thermodynamic nucleation barrier or demonstrate a special nonthermal mechanism.

A microstructure in which carbon redistribution and carbide formation reduce the supersaturation of a martensitic matrix. Plate or lath morphology can remain while local strain and thermodynamic phase composition evolve. The term describes a material state, not a universal guarantee of a particular toughness.

Author-written discussion of martensitic carbon supersaturation, carbide precipitation and microstructural evolution. www.phase-trans.msm.cam.ac.uk/2004/Tempered.Martensite/tempered.martensite.html