A thermodynamic analysis and calculation of the Fe-Ni-Cr phase diagram

Phase diagrams of Fe-Ni-Cr are calculated as a function of temperature from the liquid phase down to 500 K using thermodynamic values of the pertinent phases. These phases are the liquid, fcc, bcc, sigma, and γ^′-FeNi₃ phases. Not only the thermodynamic values calculated from the models for the various phases are in agreement with experimental data available in the literature, the calculated phase diagrams are also in agreement with experimental data over a wide range of temperature. Below 1123 K, no conventional phase diagram data are available in the literature due to kinetic difficulties. However, the calculated diagram is able to explain the results obtained when the alloys are subjected to electron irradiation. Deceased, was with the Department of Metallurgical and Mineral Engineering, University of Wisconsin.     

On the iron-carbon phase diagram

Experimental results that are obtained by electron microscopy, X-ray diffraction, and carbide analysis and indicate the precipitation of carbon atoms clusters in a hypereutectoid steel during its annealing above the eutectoid temperature are presented. These results are compared to the reported data in order to construct a new Fe-C phase diagram, where cementite forms below the eutectoid temperature due to the tendency of the Fe-C system toward ordering and carbon unbound to iron precipitates above this temperature in the form of clusters or graphite particles due to the tendency of this system toward phase separation.     

Low-temperature extension of the lehrer diagram and the iron-nitrogen phase diagram

Iron layers are nitrided in mixtures of ammonia and hydrogen at low temperatures, using a thin nickel caplayer as a catalyst. In the coordinate field of inverse temperature vs nitriding potential, we determined the boundaries between areas in which the α, γ′, or ε phases are in thermal equilibrium. Using these data, the Fe-N phase diagram is extended from 350 °C to 240 °C and extrapolated down to 200 °C. The α, γ′, and ε phases probably coexist in a triple point in the Lehrer diagram around 214 °C.     

On the Mn-MnS phase diagram

Using a DTA technique the melting point of pure MnS has been determined to 1655 ±5°C. The monotectic temperature in the Mn-MnS system was found to be 1570 ±5°C. With these new data a thermodynamic analysis of the Mn-MnS system was carried out applying a previously developed thermodynamic model.     

Calculation of the titanium-aluminum phase diagram

The Ti-Al phase diagram has been calculated by optimization of Gibbs energies with respect to phase diagram and thermochemical data.T ₀ curves, the locus of compositions and temperatures where the Gibbs energies of the liquid and one of the solid phases are equal, have been calculated over the entire composition range. In order to assure physically reasonable extrapolations of theT ₀ curves of the ordered phases far from their equilibrium stability ranges, the Bragg-Williams approximation was used to provide start values for the empirical optimizations. This approximation led to good convergence of the optimizations, and only small deviations from the Bragg-Williams Gibbs energies were needed to obtain excellent agreement with experimental data.     

双相高强钢表面机清条纹形貌和形成分析

某酸轧机组在生产双相钢时发现冷硬卷表面有条纹缺陷,为确认条纹的来源和构成,跟踪生产了两批双相钢。利用电子显微镜观察酸洗前后带钢表面条纹区域和非条纹区域的形貌和成分,结果发现,酸洗后带钢表面的条纹和热轧卷表面的条纹具有相似的成分构造和形貌。追溯分析,判定条纹产生于炼钢机械火焰清理板坯表面的过程中。当火焰清理发生漏清现象时,残留的氧化层随热轧轧制过程被压碎裹入带钢后形成条纹,条纹处严重的氧化层在酸洗过程中无法彻底消除,导致冷硬卷降级,还会损伤轧辊,对酸轧批量稳定生产双相钢造成非常严重的影响。     

Liquidus surface of the Mg-Sm-Tb phase diagram

Differential thermal, electron microprobe, and X-ray diffraction analyses and metallography are used to study Mg-Sm-Tb alloys containing up to 30% Sm or Tb. Polythermal sections and the solidification surface of the Mg-Sm-Tb phase diagram are constructed for the Mg-rich region. In the composition range under study, nonvariant transition-type equilibrium L + Mg₂₄Tb₅ = (Mg) + Mg₄₁Sm₅ is found to exist at a temperature of 539°C.     

Al-rich portion of the Al-Hf phase diagram

Electrical resistivity measurements, differential thermal analysis, optical microscopy, and the deposition of hafnium compound particles from a melt are used to study the Al-rich portion of the Al-Hf phase diagram. Prominence is given to the hafnium solubility in solid and liquid aluminum. The studies show a peritectic character of the invariant reaction during the solidification of alloys with sufficiently high hafnium contents and a slight difference between the peritectic and melting temperatures of pure aluminum. The hafnium solubility in solid and liquid aluminum is shown to increase with the temperature. The hafnium solubility in solid aluminum at the peritectic temperature (maximum solubility) is 1.00 wt % (0.153 at %); the hafnium solubility in liquid aluminum at this temperature is 0.43 wt % (0.065 at %).     

The phase diagram of the Ni-VC-NbC system

Conclusions The Ni-VC_0.87-NbC_0.9 system has a ternary eutectic in the solidification of which the equilibrium phases are an Ni-base solid solution, (V, Nb)C carbide containing 10% NbC_0.9 and 90% VC_0.87, and (V, Nb)C carbide containing 20% VC_0.87 and 80% NbC_0.9. The point of the four-phase nonvariant equilibrium is in the region of the alloy containing 3% NbC_0.9 and 6% VC_0.87 and the temperature of the equilibrium is 1300±15°C. The diagram of the phase equilibria (Fig. 3) of this system has the same form as for the Ni-TiC-ZrC and Ni-TiC-HfC systems [9, 10].Translated from Poroshkovaya Metallurgiya, No. 8(296), pp. 67–79, August, 1987.     

Nickel-rich portion of the Ni-Al-Nb phase diagram

The solubility of Nb + Al in the γ solid solution decreases markedly with decreasing temperature; thus alloys can be prepared that are γ at 1200°C and yet contain 50 pct γ’ precipitate after aging at 800°C. Thermal stability of the γ’ precipitate is related to the lattice mismatch between the γ and γ’ phases; the smaller the mismatch the lower is the interfacial elastic energy and the more stable is the γ’. Upon aging certain alloys at 800°C a γ’ growth interface other than the normal (100)_γll(100)_γ’ is observed. The maximum solubility of the niobium in γ’ is ∼7 at. pct; the width of the γ’ field increases with increasing niobium content but it is essentially independent of temperature. Replacing aluminum by niobium in γ’ gives hardnesses of up to 400 Dpn.