The integer quantum Hall effect requires a conducting two-dimensional electron gas with well-resolved Landau levels. In the simple parabolic-band model their orbital spacing isLow temperature and sufficiently low disorder broadening are needed so that $k_{B}T$ and the relevant level width $\Gamma$ are small compared with this spacing. A high-mobility sample with ${\omega }_c\tau \gg 1$ is therefore useful.

Some disorder has a constructive role: localized states between extended Landau-level states allow the Fermi energy to move through a range of carrier populations while the transport response stays on a plateau. The longitudinal response is then very small and the transverse response is quantized. Whether odd or only even integer plateaux are resolved also depends on whether spin splitting is resolved.

In the classical single-carrier model,for a two-dimensional sheet. The curve is a straight line through the origin, with its sign determined by the carrier charge and the voltage convention. The sheet Hall coefficient has units ${\mathrm{m}}^2/C$, consistent with the numerical slope used later.

The classical transverse response is not quantized. With carrier density and scattering parameters fixed, reversing $B$ reverses the transverse signal. Several carrier populations can produce a nonlinear classical curve, so straight-line behaviour here uses the stated single-population approximation rather than being a property of every conductor.

Take a two-dimensional conducting sheet in the $xy$ plane with magnetic flux density $\mathbf{B}=B\hat{\mathbf{z}}$ and electric current density along $x$. Moving carriers feel the transverse Lorentz force $q{\mathbf{v}}_d\times \mathbf{B}$. They accumulate at opposite edges, producing a transverse electric field until its force balances the magnetic deflection:For the transverse-voltage convention $E_y/j_x$, this gives $E_y=v_{x}B$ andThe Hall effect therefore measures carrier sign as well as density in the single-carrier approximation: electrons give the opposite transverse-voltage sign from positive carriers. Some tensor conventions call $E_y/j_x$ ${\rho }_{yx}$ rather than ${\rho }_{xy}$; the physical sign is fixed by the stated lead and field directions.

The Hall voltage is the integral of this transverse electric field across the sample width. In a three-dimensional slab of thickness $t$, its magnitude is $|V_H|=IB/(n|q|t)$; for a sheet, use the sheet number density $n_s$ and omit $t$.

An atomic layer deposition recipe depends on the surface chemistry during the first cycles. A silicon wafer normally carries native oxide with reactive hydroxyl groups. A pristine graphene basal plane has no comparable dense population of dangling bonds or hydroxyl groups, so precursor adsorption and initial nucleation are much less uniform.

Consequently the nominal number of cycles calibrated to give $2\mathrm{nm}$ on silicon can instead give isolated alumina islands, preferential decoration of edges, grain boundaries and defects, and uncovered graphene between islands. Adsorbed transfer residues may provide additional, irregular nucleation sites. Once alumina has nucleated, later cycles grow on the alumina normally; the problem is delayed, spatially uneven nucleation. The resulting local island height can exceed the film's area-averaged equivalent thickness.^[ref]

Two suitable improvements are:

Ozone or a mild plasma is another possible activation route, but it must be controlled because creating covalent defects can degrade the electrical properties that motivate using graphene.

The gas impingement flux onto a surface follows from the normal component of the Maxwell-Boltzmann distribution:Here $P=0.1\mathrm{Pa}$ and $m=18(1.66\times 10^{-27})=2.99\times 10^{-26}\mathrm{kg}$, givingAssume a conventional monolayer density $N_s=10^{19}{\mathrm{m}}^{-2}$ and unit sticking coefficient. The incident-monolayer time isThus a few milliseconds is the usual contamination-time estimate at this pressure.

The assumed molecular footprint matters: neither $N_s$ nor the sticking coefficient is specified. More generally $t_{\mathrm{ML}}=N_s/(sJ)$ if the sticking probability is $s$ and desorption is negligible. Water can desorb rapidly from clean graphene at room temperature, so the incident-monolayer estimate does not establish that a stable, complete water layer actually forms in that time. Residues and defects can increase local adsorption.

Two layers need not have the same electronic structure everywhere. Different rotational alignments create domains of twisted bilayer graphene; AB-stacked regions instead have the usual coupled low-energy bilayer bands. In a twisted domain, displaced Dirac cones hybridize where they meet, producing avoided crossings and saddle points in the two-dimensional dispersion relation. The resulting Van Hove singularities enhance optical transitions at energies that depend on the twist angle. Different domains therefore absorb different parts of the visible spectrum and can show different colours despite having the same layer count.^[ref]

On a silicon/silicon-dioxide substrate, reflected beams from the film stack also interfere. The wavelength-dependent optical response of the bilayer graphene changes this reflected contrast, making the domain differences visible. The observation is consistent with stacking domains; colour alone does not uniquely prove a twist angle. Variations in oxide thickness, trapped contamination, folds or local doping can also change contrast and should be checked independently, for example by Raman spectroscopy.

The important distinctions are changes of ring topology, missing or added atoms, and changes of local bonding:

Closed end caps require nonhexagonal rings, typically six pentagons per hemispherical cap; these are part of the cap topology rather than necessarily damage. Kinks, buckling, branches and incomplete closure are larger-scale structural imperfections.

A multi-walled carbon nanotube can contain any of these defects in individual shells, plus interrupted or terminated shells, extra incomplete graphitic layers, irregular interwall spacing, and covalent interwall bridges formed by displaced atoms. Such bridges change intershell coupling and sliding. Different shells having different wrapping indices is normal and is not, by itself, a defect.

Let ${\mathbf{a}}_1,{\mathbf{a}}_2$ be equal-length primitive translation vectors of the graphene lattice, separated by $60^{\circ }$. The carbon nanotube chiral vectoridentifies the two points joined when the sheet is rolled: $|{\mathbf{C}}_h|$ is the circumference, not the tube axis. The right-hand panel of the preceding original diagram shows ${\mathbf{C}}_h=6{\mathbf{a}}_1+6{\mathbf{a}}_2$ and ${\mathbf{C}}_h=7{\mathbf{a}}_1$.

The periodic boundary condition is $\mathbf{k}\cdot {\mathbf{C}}_h=2\pi \ell$, with $\ell$ an integer. Applying it at either Dirac valley gives the nanotube zone folding criterionfor a metallic tube. Thus $(6,6)$ is an armchair metallic example, whereas $(7,0)$ is a zigzag semiconducting example. Other choices with $n-m\not\equiv 0(\mathrm{mod}3)$ also give a semiconducting carbon nanotube.

This criterion is the ideal flat-sheet approximation. Curvature can open a small gap in nominally metallic non-armchair tubes; ideal armchair tubes remain gapless in that approximation. The property is determined by the wrapping vector, not simply by calling a tube “chiral” or by its diameter alone.

A single-walled carbon nanotube confines the electron around its circumference while allowing motion along its axis. Nanotube zone folding therefore turns the graphene bands into one-dimensional subbands. A semiconducting carbon nanotube has no allowed circumferential wave vector through a Dirac valley; a metallic tube has a band crossing at the Fermi energy in the ideal model. Each massive subband edge gives a one-dimensional van Hove singularity in the density of states, and corresponding peaks appear in the joint density of states.

For light polarized along the tube, the main electric-dipole selection rule conserves circumferential angular momentum: transitions connect matching valence and conduction subbands. These finite-energy transitions are conventionally $S_{11},S_{22},\dots$ in a semiconducting carbon nanotube and $M_{11},M_{22},\dots$ in a metallic tube. A metallic band crossing does not make its first strong interband peak have zero energy. Perpendicular polarization admits different angular-momentum channels, often much weaker because of depolarization. Excitons, dielectric screening and trigonal warping modify the simple band-edge transition energies.

Experimentally, measure absorption versus photon energy; Raman spectroscopy becomes resonantly enhanced when the excitation matches an optical transition. A photoluminescence excitation spectrum of dispersed semiconducting tubes correlates excitation near $S_{22}$ with emission near $S_{11}$ after relaxation. Metallic tubes generally lack comparable bright emission because nonradiative relaxation is efficient. Combining optical energies with radial-breathing-mode Raman spectroscopy or electron diffraction assigns tube diameter and roll-up indices; a Kataura plot organizes those assignments.^[ref]

Starting from the clean wafer, dehydrate it and use an appropriate adhesion treatment if required. Spin coat the positive resist to obtain a uniform film, then soft bake to remove solvent and stabilize handling. Align the mask and expose the clear pad windows. Develop away the exposed resist, rinse and dry, leaving openings at future metal locations.

Inspect the openings and clear residual resist gently if needed. Perform a compatible surface preparation immediately before metallization so that an unintended oxide or contaminant film does not dominate the contact. Deposit metal with a resist profile and film thickness that leave an accessible discontinuity at the resist edge; an undercut or suitable lift-off stack helps prevent continuous sidewall coating.

Dissolve the sacrificial resist and lift off metal above it, leaving the pads on the exposed semiconductor. Rinse, dry and inspect for residues or bridges. For source/drain test contacts, apply the intended contact anneal; for gate test contacts, preserve the intended Schottky interface. An aggressive final resist hard bake or a fully conformal bridging metal film can defeat lift-off, so these are not automatic extra steps.

Use equal-width rectangular contact windows in a row with several different edge-to-edge spacings. For positive-resist lift-off, the future metal pads correspond to clear windows in an otherwise chrome-covered mask: exposure removes resist there, and the unexposed background retains resist beneath unwanted metal.

For ordinary ohmic transmission line contact measurement, a conducting strip or isolated mesa of width $W$ connects adjacent pads. Measure several spacings $d$ so thatThe intercept separates two contact contributions from the sheet-resistance slope. Large enough pad length relative to the transfer length is needed for the simplest contact-resistivity extraction.

Prepare separate arrays on the appropriate surfaces: annealed source/drain metal on the n+ cap, and gate metal on cap-removed AlGaAs. A Schottky gate is nonlinear, so its interface cannot be assigned an ohmic zero-bias contact resistivity merely by fitting the same straight line. Use bias-dependent current–voltage or differential-contact analysis, accounting for the two rectifying contacts.

Higher donor concentration moves the semiconductor Fermi level toward the conduction band, reducing the semiconductor work function ${\Phi }_s=\chi +(E_c-E_F)$ when these quantities are expressed as energies. Equilibration with the metal aligns electrochemical potentials and creates band bending.

For an ideal n-type Schottky interface the barrier energy ${\Phi }_{Bn}={\Phi }_m-\chi$ is not automatically reduced merely by raising doping. The decisive change is the narrower depletion region:Large $N_D$ allows electrons to tunnel through the thin barrier, reducing the contact resistance and promoting an approximately Ohmic contact. Contact annealing can further modify the local interface and doping. This allows low-resistance source/drain injection while preserving a more rectifying gate contact; real interface-state pinning can modify the ideal work-function picture.

Choose process B, the citric-acid/peroxide wet etch. Its GaAs-to-AlGaAs etch selectivity is $200/1=200$, giving strong protection of the underlying barrier and a controllable removal rate for the thin cap.

For comparison, A has selectivity one and therefore no useful etch stop. It is also extremely fast: a nominal $20\mathrm{nm}$ GaAs layer would be removed in only $0.048\mathrm{s}$. C has selectivity $1000/20=50$, useful but inferior to B, and clears the same cap in about $1.2\mathrm{s}$. D has selectivity four, clears the cap in about $0.060\mathrm{s}$ and additionally risks plasma-induced surface damage or altered interface properties.

B clears the nominal cap in about $6\mathrm{s}$. Its slower underlying etch permits a more forgiving endpoint or modest overetch, although actual rates must be checked for the relevant composition and conditions. Wet-etch lateral undercut and contact-metal compatibility still matter; selectivity alone does not make every mask geometry acceptable.

A suitable sequence isFirst pattern the source/drain resist, then deposit and lift off the corresponding metal. Anneal these contacts to produce the low-resistance source/drain interfaces. Remove the exposed cap layer before gate lithography, leaving the cap beneath the contact metal; this uses the contact pads as an etch mask, with suitable protection where required. Finally define and deposit the gate onto exposed AlGaAs and lift off.

The key dependencies are that source/drain metal lands on the n+ cap, the gate lands on AlGaAs after cap removal, and the contact anneal precedes formation of the sensitive Schottky gate. A gate-recess variant can put gate lithography before the local cap etch if that resist is deliberately used as its mask, but it must additionally specify how unwanted conducting cap elsewhere is removed or isolated. The sequence above directly matches the cap remaining only beneath the source/drain contacts in the diagram.

Molecular beam epitaxy is a suitable research-scale choice. Independently controlled elemental beams and shutters allow abrupt composition changes, controlled silicon donor incorporation and thin spacer layers. Growth can proceed sequentially from the GaAs substrate through the undoped GaAs, undoped AlGaAs spacer, doped AlGaAs supply layer and highly doped GaAs cap.

A clean crystalline surface and controlled substrate temperature support epitaxial interfaces; separating the donor layer from the channel is modulation doping. Metalorganic vapor-phase epitaxy can also grow such structures, but MBE's controllable research-scale layer sequence is a strong match to the stated task. Merely depositing the right total elemental amounts would not guarantee the required crystallinity and interface abruptness.

For a face-centered cubic lattice, allowed reflections have indices all odd or all even. The first two are $\{111\}$ and $\{200\}$, withFor first-order Bragg diffraction, $2d\sin \theta =\lambda$. The detector intercepts the beam at the full scattering angle $2\theta$, givingThe inner detector radius supplies the minimum length for each ring:Thus both rings reach the annulus when $L\gtrsim 90.2\mathrm{mm}$. At this minimum the 200 ring is at roughly $3.12\mathrm{mm}$, well inside the $10\mathrm{mm}$ outer radius. Both remain within the detector for approximatelyEquality at the inner edge gives vanishing acceptance for an ideal infinitely narrow ring, so practical operation requires a small margin. The calculation concerns ring positions; individual nanoparticle orientation and the finite probe convergence also affect observed diffraction contrast. Using the Bragg angle instead of the full scattering angle would introduce an erroneous factor of about two.

An incident SEM probe undergoes elastic deflections by atomic potentials and inelastic interactions that excite or ionize the specimen. Repeated scattering generates an interaction volume wider and deeper than the incident probe. The measured resolution depends both on where a signal is generated and on whether it can escape.

Transmitted electrons from a sufficiently thin specimen can instead support scanning transmission electron microscopy. The values above are orders of magnitude, not intrinsic constants of the signal: beam energy, atomic number, specimen thickness, probe size and detector geometry all matter. The spatial origin explains why a fine topographic SEM image does not guarantee equally fine X-ray mapping.

The material figure of merit isLarge thermoelectric power factor supplies electrical output, while low thermal conductivity limits heat leaking without useful conversion. For a constant-property ideal generator, a common approximation isIncreasing the appropriately averaged figure of merit raises efficiency toward, but not beyond, the Carnot limit. Actual devices also incur contact, geometric and heat-transfer losses, and a single peak material ZT does not specify their efficiency.

The supplied trends show the trade-off with carrier density: conductivity rises, the magnitude of the Seebeck coefficient tends to fall, and electronic thermal conductivity rises approximately with conductivity through the Wiedemann-Franz law. Thus neither arbitrarily high doping nor maximization of conductivity alone gives the best ZT; the power factor and thermal terms produce an intermediate optimum.

Two complementary strategies are:

Both strategies must be checked for mobility penalties and electronic heat transport. Simply suppressing every kind of transport would also suppress electrical output.