Magnetic properties,microstructures and domain structures of arc-plasma sprayed Nd-Fe-B permanent magnet

Thin Nd-Fe-B films prepared by arc-plasma spraying at different substrate temperatures were investigated for their magnetic and structural properties. The isotropic magnets with the best magnetic properties (_M H _c=1.2 MA m^–1, _o M _r=0.6 T, (BH)_max=64 kJ m^–3), were obtained after plasma spraying the Nd-Fe-B powders on water-cooled copper substrates and subsequently annealing the films for 0.5 h at 750 °C. The optimum magnetic properties of the anisotropic Nd-Fe-B films, i.e. _M H _c=1.2 MA^–1, _o M _r=0.9 T and (BH)_max=180 kJ m^–3, were obtained in films sprayed on to heated to 600 °C substrates. The magnetic properties of the sprayed films were strongly influenced by the microstructure. The domain structure of these films is also presented.     

Stress-strain rate relations for high-temperature deformation of two-phase Al-Cu alloys

Al-Cu alloys containing 6, 11, 17, 24 and 33 wt% Cu, annealed for 0.5–100 h, were deformed by the differential strain-rate test technique over a strain-rate range of 4×10^–6 to 3×10^–2s^–1 at temperatures ranging from 460–540°C. Superplastic behaviour, with strain-rate sensitivity, m0.5, and activation energy, Q=171.5 kJ mol^–1, is shown by the Al-24Cu and Al-33Cu alloys at lower strain rates and higher temperatures. All the alloys show m0.20 at higher strain rates, but the average activation energy for deformation of the Al-6Cu, Al-11Cu, and Al-17Cu alloys is evaluated to be 480.7 kJ mol^–1, in contrast to a lower value of 211 kJ mol^–1 for the Al-24Cu and Al-33Cu alloys. Instead of grain size, the mean free path between particles is suggested to be a more appropriate microstructural parameter for the constitutive relationship for deformation of the Al-Cu alloys.     

Effects of lithium addition and wheel velocity on the microstructure of aluminium-lithium alloy ribbons

The effects of Li addition and wheel velocity on the microstructure of as-cast AI-Li ribbons were studied by optical metallography, scanning and transmission electron microscopy and by X-ray diffraction. Li addition has a marked effect on the ribbon solidification mechanism. Ribbons spun at 12.5 ms^–1 and which contained up to 1.79 wt% Li experienced a cooling rate of about 10⁴ Ks^–1 and solidified under the melt jet with a growth rate greater than 0.3 m s^–1. However, when Li content exceeded about 3 wt%, the ribbons were undercooled (cooling rate 10⁶ Ks^–1) and formed far from the melt puddle effect. Under these conditions, unusual secondary featureless zones occurred through the ribbon thickness and the detected phases corresponded to those expected to be formed under equilibrium conditions. Nevertheless, (AI₃Li) phase did not occur in ribbons containing up to 1.79 wt % Li or in the featureless zones of the most Li-rich ribbons (Li >about 3 wt%). Changing of the wheel velocity from 12.5 to 22.2 or 35.7 ms^–1 did not affect the ribbon formation mechanism, but it favoured the precipitation and decreased slightly the grain size.     

Microstructural evolution and calorimetric evaluation of non-equilibrium states in rapidly solidified Al-Mn-Co ternary alloys

Evolutionary structures were observed in as-melt spun ribbons of Al-3.5Mn-0.8Co, Al-3.5Mn-1.3Co, Al-3.5Mn-2Co, Al-5Mn-0.8Co and Al-5Mn-2Co alloys. The evolutionary structures contained two zones, zone A and zone B, at the chilled and unchilled sides, respectively. A banded structure was observed between the two zones as a transitional region in the Al-3.5Mn-1.3Co alloy. The enthalpy differences, H _O ^ne , between a non-equilibrium state in as-melt spun ribbons and the equilibrium state in fully annealed ribbons of the Al-Mn-Co alloys were evaluated using data obtained by differential scanning calorimetry (DSC) measurements. The values of H _O ^ne for Al-Mn-Co alloys are lower than those for Al-Mn binary alloys. H _O ^ne corresponds to a decrease in Gibbs free energy, G, which is stored in a solidified solid during rapid solidification. The formation of the banded structure is interpreted on the basis of an analysis of the distribution of H _O ^ne or the fluctuation in G, which should oscillate in order to correspond with a gradual reduction of cooling ability for the constrained solidification by the substrate.     

Activation Energy of Bi2Sr2Ca1−x Er x Cu2O8+d Epitaxial Thin Films

The scaling behavior of the effective activation energy U(T, H) of epitaxial oriented Bi₂Sr₂Ca_1–x Er _x Cu₂O_8+d thin films with 0x0.10 has been studied as function of temperature and magnetic field. It has been found that the activation energy scales as U(H, T)=U ₀(1–t)^m H ⁿ with exponent m=1.20±0.03 and n=–1/2. The field and temperature dependence of the activation energy, as well as its overall magnitude, is consistent with a model in which U(T, H) arises from plastic deformations of viscous flux liquid above the vortex-glass transition temperature.     

Fabrication and mechanical properties of in situ formed carbide particulate reinforced aluminium composite

Stable carbide particles of TiC, ZrC and TaC were in situ synthesized in liquid aluminium by the reaction between Al-Ti, Al-Zr or Al-Ta systems liquid alloy and SiC or Al₄C₃ particles. It was possible to generate TiC particles of nearly 1 m diameter, even utilizing SiC of 14 m. However, the dispersion behaviour of TiC particles in the matrix depended on the size of the raw carbide. Finer SiC made the dispersion of TiC particles more uniform, resulting in a greater improvement of the mechanical properties. Furthermore, although Al-Ti-Si system intermetallic compound was detected in a TiC_p/Al-Si composite fabricated by the melt-stirring method, those compounds considerably decreased in the composite fabricated by the in situ method. The mechanical properties of in situ formed TiC_p/Al-5 wt% Mg and TiC_p/Al-5 wt% Cu composites were better than those fabricated by the melt-stirring method and by T6 heat treatment, the properties of in situ formed TiC_p/Al-5 wt% Cu composite were further improved. The experimental results were analysed by the reaction model based on the assumption that the overall reaction rate was controlled by both the reaction and the diffusion.     

Aerobic composting of tobacco industry solid waste—simulation of the process

Solid waste accumulated during the processing of tobacco for cigarette manufacture mostly contains tobacco particles and flavoring agents. Its main characteristics are a high content of nicotine (2,000 mg per kg of total solids), which is a toxic compound, and high value of total organic carbon of the aqueous extract (12,620.0 mg l^–1). Because of this fact tobacco waste cannot be disposed of with urban waste.The aim of this work was to stabilize tobacco solid waste by aerobic composting. The experiments were carried out in closed thermally insulated column reactors (1.0 l and 25 l) under adiabatic conditions and at an airflow rate of 0.9 l min^–1 kg^–1 of volatile solids for 16 days. During the process, temperature changes in the reactor, CO₂ production and the numbers of mesophilic and thermophilic organisms in the mixed microbial culture were closely monitored. Nicotine concentration in the samples was analyzed at the start and at the end of process. It was estimated that at the end of composting the volume and mass of total solids in the tobacco waste were reduced by about 50% and those of nicotine by 80%. A simple empirical model was used to simulate the biodegradation rate of the organic fraction of the solid waste. It was found that the selected model describes aerobic composting fairly well, although only two kinetic parameters (k₀ and n) were estimated.List of symbols c_pS specific heat capacity of the substrate, kJ kg^–1 K^–1 - c_pz specific heat capacity of air, kJ kg^–1 K^–1 - F_Ku and F_Ki molar airflow at the reactor inlet and outlet, mol h^–1 - H_r reaction enthalpy, kJ kg^–1 of dry substrate - k specific rate, Eqs. (5) and (9), h^–1 - k_o constant in Eq. (9), day^–1 - m_o initial mass of the substrate, kg - m_S mass of dry substrate, kg - n order of the reaction in Eq. (5) - n_K molar amount of oxygen, mol - Q_v airflow volume, m³ h^–1 - r_K oxygen depletion rate, mol kg^–1 h^–1 - r_S degradation rate, kg kg^–1 h^–1 - _z air density, kg m^–3 - SD mean square deviation - t time, h - T temperature in reactor, °C - T_o temperature of substrate at the beginning of reaction, °C - T_K temperature of compost at the end of reaction, °C - T_u temperature of air at the reactor inlet - space time, day - w_S mass fraction of compost, m_sm_o^–1, kg kg^–1     

A method of joint determination of the effective coefficients of horizontal mixing of large particles and of external heat exchange in an organized fluidized bed

A method is proposed for the joint determination of the coefficients of horizontal particle diffusion and external heat exchange in a stagnant fluidized bed.Notation c_f, c_s, c_n specific heat capacities of gas, particles, and nozzle material, respectively, at constant pressure - D effective coefficient of particle diffusion horizontally (coefficient of horizontal thermal diffusivity of the bed) - d equivalent particle diameter - d_t tube diameter - H₀, H heights of bed at gas filtration velocities u₀ and u, respectively - H_a height of active section - l width of bed - L tube length - l _o width of heating chamber - N number of partition intervals - p=H/H₀ expansion of bed - s_n surface area of nozzle per unit volume of bed - S_h, S_v horizontal and vertical spacings between tubes - t_c, t₀, t_s, t_n, t_w initial temperature of heating chamber, entrance temperature of gas, particle temperature, nozzle temperature, and temperature of apparatus walls, respectively - u₀, u velocity of start of fluidization and gas filtration velocity - y horizontal coordinate - ^*, coefficient of external heat exchange between bed and walls of apparatus and nozzle - ₁, ₁, ₂, ... coefficients in (4) - thickness of tube wall - _b bubble concentration in bed - ₀ porosity of emulsion phase of bed - _n porosity of nozzle - =(t_s – t₀)/(t_c – t₀) dimensionless relative temperature of particles - _n coefficient of thermal conductivity of nozzle material - f, _s, _n densities of gas, particles, and nozzle material, respectively - _be=_s(1 – ₀) (1 – _b) average density of bed - time - _max time of onset of temperature maximum at a selected point of the bed - R =l _o/l Fourier number - Pe = ₁ l ²/D Péclet number - Bi = /_n Biot number Translated from Inzhenerno-Fizicheskii Zhurnal, Vol. 41, No. 3, pp. 457–464, September, 1981.     

The influence of Cu and Mg additions on the decomposition of Al-(40 to 50) wt% Zn alloys during isothermal and continuous heating

The influence of 1 wt% copper and 0.2wt% magnesium additions on the heat evolution and structural changes of Al-40 and Al-50wt% Zn alloys during isothermal ageing at 30° C and continuous heating up to 400° C was studied using microcalorimetry, differential scanning calorimetry and transmission electron microscopy. Isothermal studies indicate rapid formation of Guinier-Preston (GP) zones at initial stages of ageing followed by discontinuous precipitation of-phase competing with the formation of transition phases also containing third elements. DSC curves show that the amount of metastable phases formed during room temperature ageing increases with ternary additions, especially to Al-50% Zn alloys. Up to 75° C, GP zones undergo composition but no size changes: during further continuous heating,' (Mg₂Zn) appears, the dissolution of which proceeds parallel to the continuous precipitation of at medium temperatures. Addition of copper reduces the discontinuous-precipitation to a lesser extent than magnesium.     

Estimation of the unknown parameters in the melt-spinning process

The free-surface temperature history of the melt spinning of copper measured by Tenwick and Davies [3] is compared with those calculated using a thermokinetic model assuming different parameters. The heat-transfer coefficient, nucleation temperature and the crystal-growth kinetics were thus estimated for the melt spinning of copper at a wheel speed of 35 ms^–1 as follows: heat-transfer coefficient during liquid cooling stage HL=1.0 × 10⁷ W m^–2K^–1, heat-transfer coefficient after solidification finished HS=1.0 × 10⁵ W m^–2K^–1, heat-transfer coefficient during solidificationH= 1.0 x 10⁷- 1.2 x 10¹¹ (t-t _n) (W m^–2K^–1), the nucleation temperatureT _n 1233 K and the crystal-growth kinetic lawV=4.0 × 10^–3 T^1.1 (ms^–1).     

The influence of surface kinetics in modelling chemical vapour deposition processes in porous preforms

The isothermal chemical vapour infiltration (ICVI) process is a well known technique for the production of composites and the surface modification of porous preforms. Mathematical modelling of the process can provide a better understanding of the influence of individual process parameters on the deposition characteristics such as final porosity or deposition profiles in the pore network. The influence of different rate expressions for several binary compounds on the ICVI process is discussed. Experimental work is used to validate the importance of correct kinetic expressions in a continuous ICVI model for cylindrical pores. The predicted infiltration characteristics are compared with experimental results. The final densification and Thiele modulus, i.e. a number which is a measure for the diffusion limitations in a pore, are used for the evaluation of the presented model, and conditions are given for an optimal densification of a porous preform by the ICVI process for several binary compounds. The deposition profiles as predicted by the model calculations are in agreement with the experimentally determined deposition profiles of TiN and TiC in small tubes. Moreover, it can be concluded that the shape of the deposition profiles is determined by the heterogeneous reaction kinetics. There is only a qualitative agreement between the predicted densification and measured densification for the synthesis of TiN and TiB₂ in sintered porous alumina. This mismatch can be explained in terms of a complexity of the pore network and differences in reaction kinetics. Model calculations reveal that there is a scattering for the predicted residual porosity as a function of the Thiele modulus for TiN. Moreover, this Thiele modulus can not fully account for the changes in densification at different temperatures. Given these uncertainties it is likely that a residual porosity of less than one percent can be obtained if the Thiele modulus is smaller than 1 × 10^–4. However, a CVI process with such a small Thiele modulus will not be practical, because of the concomitant long process times. Therefore, more precise conditions for the individual process parameters, i.e. concentration, reactor pressure, and temperature are deduced from the model calculations.Nomenclature a, b, c reaction order constants - C _i(x, t) concentration of species i at axial position x and time t (mole m^–3) - C _i ^o bulk concentration of species i (mole m^–3) - C _i ^* (x, t) dimensionless concentration of species i at axial position x and time t - D _e(x, t) effective diffusion coefficient at axial position x and time t (m²s^–1) - D _ij(x, t) binary diffusion coefficient (m²s^–1) - D _K(x, t) Knudsen diffusion coefficient at position x and time t (m²s^–1) - F correction factor for effective diffusion coefficient - k growth rate constant (ms^–1(m³mole^–1)^a+b-1) - K _i adsorption-desorption equilibrium constant (m³mole^–1) - L length of a pore (m) - M _i molecular weight of species i (g mole^–1) - M _ij harmonic mean of the molecular weights of species i andj (g mole^–1) - M _s molecular weight of deposit (g mole^–1) - m _t measured mass increase (g) - n _i stoichiometric number - P reactor pressure (Pa) - R(C _i) growth rate (mole(m^–2s^–1)) - r(x, t) pore radius at position x and time t (m) - r ^o initial pore radius (m) - r ^* dimensionless pore radius - S geometrical surface area (m²) - s ^t fraction of free titanium sites at the surface of TiN - s ⁿ fraction of free nitrogen sites at the surface of TiN - T temperature (K) - t time (s) - t _p process time (s) - U K _HCl/(K _{H 2} C _{H 2})^1/2 (m³ mole^–1) - V volume of alumina substrate (m³) - W K _TiCl3(m³ mole^–1) - X volume of infiltrated deposit relative to initial pore volume - x axial distance (m) - x ^* dimensionless axial distance - z number of time steps - dummy variable for integration - porosity of sintered porous alumina substrate - ratio of the volume over the surface area perpendicular to the flux (m) - density deposit (kg m^–3) - _ij a characteristic length (Å) - tortuosity factor of substrate - Thiele modulus - _D collision integral     

Thermocapillary stability during gravity flow of a liquid film

Thermocapillary rupture of a film under conditions of turbulent undulatory flow is associated with the buildup of wave motion on its surface. Here an approximate solution to the problem and criterial relations are obtained for determining the limits of stable film flow.Notation _min, kg/m·sec minimum irrigation intensity at which no film rupture occurs - ₁, kg/m· sec irrigation intensity at which the first dry spot appears - q, W/m² thermal flux density - _D, °C temperature at the rupture section - x, m space coordinate along the warm surface in the direction of flow - y, m coordinate in the direction normal to the warm surface - _o, m mean thickness of the film between large waves - _c, m thickness of the continuous layer - _cr, m critical film thickness - _o=/_o andl _o=l _o/_o dimensionless initial amplitude and length of a wave - , sec^–1 recurrence frequency of large waves - t_cr, sec time till thermocapillary rupture of a film - t_p, sec time of penetration of a thermal perturbation through the film thickness - u, m/sec velocity of thermocapillary flow of the liquid - , W/m·°C thermal conductivity - c_p, kJ/kg·°C specific heat - , kg/m linear density - , N·sec/m² dynamic viscosity - a, m²/sec thermal diffusivity - , N/m surface tension - , N/m² tangential stress at the film surface - L, m length of the warm pipe segment - L_o, m distance from the inlet to the section where wave motion at the film surface occurs - ¯w, m/sec mean velocity of downward flow of liquid in the film - , m mean thickness of the laminar layer - g, m²/sec free-fall acceleration due to gravity Translated from Inzhenerno-Fizicheskii Zhurnal, Vol. 39, No. 4, pp. 581–591, October, 1980.     

Structure,properties and response to heat treatment of melt-spun Al-Y and Al-La alloys

Al-Y and Al-La binary alloys containing 0.7–18 wt% (0.2–6.3 at%) Y and 0.9–18 wt% (0.2–4.2 at%) La, were rapidly solidified by chill-block melt-spinning to produce ribbons between 35 and 70 m thick. Microstructures were of the classical zone A/zone B type with a notable increase in Al lattice parameter for the Al-6.3 at% Y composition, which exhibited a Knoop hardness of 430±30 kg mm^–2 as-spun. Isochronal ageing for 2 h at 200–500 °C gave significant hardening at 200 and/or 300 °C for all of the more concentrated alloys, the largest responses being produced by Al-6.3 at % Y and Al-4.2 at % La at 200 °C. X-ray diffraction asspun indicated the presence of only Al and equilibrium Al₁₁La₃ in the Al-La alloy ribbons and Al and a non-equilibrium Al₄Y/Al₁₁Y₃ in the Al-Y ribbons. This non-equilibrium Al-Y phase was identified by X-ray diffraction as isomorphous with orthorhombic or tetragonal Al₁₁ La₃ with lattice parameters determined as a _o = 0.42 ± 0.02 nm (b _o = 1.26 ± 0.06 nm) and c _o = 0.97 ± 0.05 nm. TEM showed that it was present as an intercellular network with Energy dispersive spectroscopic analysis indicating an average composition Al-46 wt% Y consistent with the Al₄Y/Al₁₁Y₃ stoichiometry and diffraction patterns consistent with an orthorhombic or tetragonal cell with these lattice parameters. While no significant change in phase constitution of the Al-La ribbons was detected by X-ray diffraction as a result of heat treatment, the Al₁₁Y₃ in Al-Y ribbons was seen to be replaced by AI₃Y on heat treatment at 400 and 500 °C.     

Torsional Oscillator Study of Thin Helium Films on Hydrogen Substrates

We have made torsional oscillator measurements of thin helium films on a Mylar surface preplated with 2 and 5 layers of H ₂ . The minimum ⁴ He coverage needed for superfluidity, or inert layer, n _o on both surfaces is found to be 6 mole/m ² . This is equivalent to 1/2 of a monolayer at bulk liquid density. The superfluid transition in coverages above n _o is similar to that found on bare Mylar, exhibiting the standard Kosterlitz-Thouless characteristics. We find no anomaly in _s or the dissipation in the film that can be identified with a second transition reported in recent third sound experiments.     

Solidification and interfacial structure of in situ Al-4.5Cu/TiB2 composite

In situ particle reinforced Al-4.5Cu/TiB₂ composite was fabricated with TiO₂, H₃BO₃, Na₃AlF₆ powders and Al-4.5Cu alloy by reaction in melt. The composite can be directly casted into moulds to make composite parts. TiB₂ particles distribute uniformly in the matrix. The average size of TiB₂ particles is 0.93 m. At the atomic scale, TiB₂ is hexagonal, and exhibits hexagon or quadrilateral shape. The orientation relationships exist in the interfaces between TiB₂ particle and -Al, and between the reinforced small Al₂Cu phase and -Al in the composite. They are . TiB₂ particle is nucleation site for -Al matrix growth in the composite. The interface between TiB₂ particles and the matrix is clean and well bonded. No reaction product has been found through HREM observation. This is beneficial to the strength of the composite. The as-cast Al-4.5Cu/TiB₂ composite exhibits mechanical excellent properties: the tensile strength is 416.7 MPa, the yield strength is 316.9 MPa, and the elongation is 3.3 pct.     

Standard bases,critical tropisms and flatness

Using the standard basis notion, we adapt to the affine algebraic geometry the notion of the critical tropisms of Lejeune-Teissier and we give some effective criteria for testing geometrical properties for the generic projectionX=SpecK[t₁,...,t_m][x₁,...,x_n]/I y=SpecK[t₁,...,t_m].     

Thermodynamic considerations of metal halide vapour complexes

Certain metal halide vapour complexes of the type ABX _m+n increase the equilibrium vapour concentration of one or both of the component salts in the binary system (AXm/BX_n). Such enhancement depends upon the free energy of complex formation, the saturated vapour pressure of the complexing salt and the melt activities, in a manner formulated here by means of a thermodynamic scheme. Certain trends in the periodic table are established.Nomenclature enhancement term=p ⁰ K _a - enhancement term=P ⁰ K _a - F enhancement - p partial vapour pressure (atm) - p ⁰ saturated vapour pressure of monomer (atm) - P ⁰ total saturated vapour pressure exerted by a salt (atm) - vapour concentration of a species - ⁰ vapour concentration of a species above the pure salt (mol m^–3) - x cation mole fraction - H _a enthalpy of complex formation from monomers - H _b vaporization enthalpy of AX _m - H _c vaporization enthalpy of BX _n - H _d vaporization enthalpy of ABX _m+n - H _e exchange enthalpy - H _m enthalpy of mixing in liquid pool - V change in coulombic potential energy - D _b bridge bond strength - D _t end bond strength     

A solution to the multivariable matrix factorization problem

Summary Given a para-Hermitian matrix,A(p ₁,p ₂, ...,p _n), whose elements are real, rational functions of the complex variablesp ₁,p ₂, ...,p _n, it is shown thatA (p ₁,p ₂, ...,p _n) can be factorized in the form,A (p ₁,p ₂, ...,p _n)=H ^t (–p ₁, –p ₂, ..., –p _n)H (p₁, p₂, ...,p _n), where the elements of the matrixH (p ₁,p ₂, ...,p _n) are also real rational functions of the specified variables, if and only ifA (j ₁, ...,j _n) is non-negative definite for all real _i, i=1, 2, ...,n. A rather simple computational method for the construction ofH (p ₁,p ₂, ...,p _n) is given, and examples are used to illustrate how, in many cases, the factorization can actually be carried to completion with little labour.     

Behaviour of hydrogen in pure aluminium,Al-4 mass% Cu and Al-1 mass% Mg2Si alloys studied by tritium electron microautoradiography

The effect of lattice defects and microstructures on hydrogen distribution in pure aluminium, Al-4 mass% Cu and Al-1 mass% Mg₂Si alloys has been studied by tritium electron microautoradiography. It was found that in pure aluminium and both the alloys, dislocations and grain boundaries act as short-circuiting diffusion paths and also as trapping sites for hydrogen. It was found that the Guinier-Preston (GP) zones in the Al-1 mass% Mg₂Si alloy do not act as trapping sites but as repellants for hydrogen, in contrast to the GP zones in the Al-4 mass% Cu alloy in which they act as trapping sites for hydrogen. In the Al-1 mass% Mg₂Si alloy, the interfaces between the matrix lattice and the metastable precipitate and between the matrix lattice and the equilibrium precipitate, were proved to be strong trapping sites for hydrogen. In the Al-4 mass% Cu alloy, the equilibrium precipitate itself has been found to be able to trap hydrogen.     

Coarsening of δ-Ni<Subscript>2</Subscript>Si precipitates in a Cu–Ni–Si alloy

The coarsening behavior of rod-shaped and spherical δ-Ni₂Si precipitates in a Cu–1.86 wt% Ni–0.45 wt% Si alloy during aging at 823–948 K has been investigated by measuring both precipitate size by transmission electron microscopy (TEM) and solute concentration in the Cu matrix by electrical resistivity. The rod-shaped δ precipitates have an elongated shape along 〈[`5] 5 8 \overline{5} 5 8 〉_m and a {110}_m habit-plane facet. The coarsening theory of a spherical precipitate in a ternary alloy developed by Kuehmann and Voorhees (KV) has been modified to a case of rod-shaped precipitates. The coarsening kinetics of average size of the rod-shaped and spherical δ precipitates with aging time t obey the t ^1/3 time law, as predicted by the modified KV theory. The kinetics of depletion of the supersaturation with t are coincident with the predicted t ^−1/3 time law. Application of the modified KV theory has enabled calculation of the energies of sphere, {110}_m and rod-end interfaces from the data on coarsening alone. The energy of the {110}_m interface having a high degree of coherency to the Cu matrix is estimated to be 0.4 J m⁻², the incoherent sphere-interface energy 0.6 J m⁻², and the rod-end interface energy 5.2 J m⁻².