The desulphurization of molten slags in the gaseous phase during top-blown with air has been studied from the kinetic standpoint, emphasis being made to study the effects of FeO, the j value (j=N_(Fe_2O_(3))/(N_(Fe_2O_3)+N_(FeO))) and the basicity of the slag upon the rate of removal of sulphur. It has been found that this gaseous desulphurization takes place according to the following reaction: (S~(2-))+3/2O_2=SO_2+(O~(2-)), (1)while the iron ions in the slag serve only as the oxygen-carriers:6(Fe~(3+))+(S~(2-))+2(O~(2-))=6(Fe~(2+))+SO_2, (2)6(Fe~(2+))+3/2O_2=6(Fe~(3+))+3(O~(2-)). (3)With slags containing very little or no FeO or slags with constant j value, as when air is blown upon the slag surface from a rather remote distance, the reaction has been found to be of the first order, the transfer of S~(2-) ions inside the slag being most probably ratecontrolling. An increase in the FeO content and the j value of the slag, as well as the temperature, accelerates the speed of removal of sulphur, while an increase in slag basicity retards it. It has been proposed that, owing to the oxidizing character of the blown-in air, SO_4~(2-) ions might have been formed inside the slag, but at the high temperatures prevailing, they decompose immediately into SO_2 and O_2. Any sulphate that may be present in the solidified slag is most probably being formed on its way of cooling in air. Attention has been called to resort to the correct method of analysis of total sulphur in the slag, the combustion method using CO_2 as the oxidizer being more accurate than the method using pure O_2. It has also been shown experimentally that the direct gaseous desulphurization of the metal is not so feasible as that of the slag, and that there exists some contradiction between the desulphurization of the metal with slag and the gaseous desulphurization of the slag. With a view to bringing about as a whole a more efficient desulphurization, the possibility of adapting the gaseous desulphurization to the actual steelmaking practice with the converter is briefly discussed.

The thermo-stability of K_2TaF_7 in air and argon atmosphere was investigated by means of thermo-differential, thermo-gravimetric and chemical analyses and also X-ray powder diffraction method. An obvious heat-absorbing effect was observed at 218°±3℃ with the appearance of break-up of larger crystals. The melting point of K_2TaF_7 was measured to be 730°±3℃. During the course of heating, by the action of moisture in air, K_2TaF_7 was hydrolyzed into K_2TaO_2F_3 which is soluble in molten K_2TaF_7 and thus affects its melting point. This might throw light on the disagreement of the melting points of K_2TaF_7 determined by the various authors. The end product of hydrolysis of K_2TaF_7 was identified to be KF and KTaO_3 (potassium meta-tantalate). Finally, the different stages of hydrolytic reaction of K_2TaF_7 and its relation with various temperatures were discussed.

The paper discusses the conditioning and stabilization of austenite during tempering, and the suggested new tempering process based upon the results of research on high speed steels. The effect of conditioning of austenite, expressed in term of the M' temperature at which the retained austenite transforms to martensite during cooling, depends upon the temperature and duration of tempering. On step tempering, i.e., holding at a higher temperature followed by a lower one (but above a certain temperature T_c), the effect of conditioning is additive and sometimes equals or exceeds that of tempering only at the higher temperature with identical duration. The amount of retained austenite transformed after optimum step tempering (e. g., 560℃ holding followed by 550℃ holding) is similar to that after multiple tempering process, but the tempering time is shortened. The M' temperature falls with the increasing amount of retained austenite and varies with the carbon content of austenite. Holding at temperatures below T_c and above M' causes a marked effect on the stabilization of austenite. The degree of stabilization and the initial rate of stabilization decrease with the increasing holding temperature. But the maximum amount of austenite stabilized increases with the holding temperature when the holding temperature is below M'. It is suggested that the stabilization of austenite during tempering is chiefly due to the reduction of stresses that promote the growth of martensite embryos. The conditioning may be regarded as a reversed process of stabilization and the following changes are suggested: the rearrangement of random dislocations, the formation of dislocation loops at α-γ phase boundaries and the removal of atmospheres formed from vacancies and solute atoms. A new tempering process consisting of step tempering followed by a normal tempering is suggested, e.g., holding at 590℃ for 25 minutes, then at 550°C for 25 minutes, and finally tempering at 560℃ for 60 minutes. The amount of austenite transformation, the mechanical properties and the cutting abilities of the high speed steel treated with the new process are comparable with those obtained by traditional multiple tempering. An optimum tempering process for reducing the deformation of tools can also be designed by the application of the stabilization of austenite. The total tempering durations are shortened in both cases.

The results of tension, compression and plastic bending tests made on a Hadfield steel have been compared, from which a quantitative relationship between them has been found. It is suggested that plastic bending test may either be used as an independent test or as an estimation for the direct tension and compression tests. The correlation between them may be used to explain the technical specifications for cast products of Hadfield steel.

The influence of various factors on the aging embrittlement of Cr-Mn-N austenitic heat-resisting steels strengthened with Mo, V and B have been studied. It is shown that there are three factors controlling the aging embrittlement. (1) Loss of ductility is accompanied by fine precipitation of carbide and nitride along the grain boundaries. With increasing C and N content, the ductility of the aged steel is dropped sharply. The time-temperature-iso-impact values curves are in form of double C shape. The noses of the curves may possibly be related to the precipitation of carbide and X phase. (2) Strengthening elements such as Mo and B tend to segregate at grain boundaries. The effect of B in steel is to prevent the formation of the continuous fine precipitation film along grain boundaries, thus the aging embrittlement is minimized. Segregation of Mo along grain boundaries promotes brittlement. (3) After aging at temperature interval between 600—850℃ for a long time, austenite will become less stable due to the precipitation of carbide and nitride as M_(23)C_6 and (Cr, Fe)_2N of the solid solution. Thus the diffusionless transformation such as γ→s→α' or γ→α' can easily occur when shock load is applied, and therefore the aged specimens will show low ductility at temperature of testing.

The mechanism of age-hardening of magnesium-zinc alloy has been investigated by means of hardness measurement, tensile test, metallographic examination and X-ray analysis. It is shown that when the magnesium-5%zinc alloy is aged at 200℃ for not too long a period, a hexagonal transition phase β' is precipitated, which is rod-like, orientated perpendicularly to the basal plane of the magnesium matrix and coherent to it. The existence of such transition β' precipitates raises the yield strength of the alloy but does not change appreciably its rate of strain hardening in tension test at 77°K. It is presumed that when this age-hardened alloy was plastically deformed, the gliding dislocations were unable to cut through the β' rods but piled-up against them. The coherent stress field accompanying the β' phase is the principal reason to account for the rise of yield strength of the alloy after aging. Additional external stress must be applied in order to overcome such coherent stress field and to force the constricted gliding dislocations to cross-slip into the prismatic planes of the matrix.

The strain ageing of a-titanium under the influence of an applied stress has been studied. The kinetics of strain ageing obeys the following equation where n=2,; is the increase in flow stress σ_f due to strain ageing; n_0 is the original concentration of solute; n_∞ is the maximum concentration of solute segregated in the dislocations; D is the diffusion coefficient of solute atoms; U_m is the binding energy between a solute atom and a dislocation; b is the Burgers vector; t is the ageing time, k Boltzmann constant and T the absolute temperature. The activation energy of strain ageing obtained is about 1.30±0.15 eV. The strain ageing of α-titanium is very probably due to the presence of both iron as the substitutional atoms and oxygen as the interstitial atoms.

The deformation mechanism of α-titanium in the temperature range -196—400℃ has been investigated by differential tensile tests. The activation energies, activation volumes and frequency factor are thus obtained. After comparing with various dislocation mechanisms, it is shown that the deformation-controlling process is probably the overcoming of the Peierls-Nabarro force.

The development of the chromium-depleted zone with time of sensitization has been studied by the potentiostatic-coulometric method reported previously. The distribution curves of chromium in the depleted zone at different time of sensitization have been obtained. From these curves, the diffusion coefficient of Cr in the depleted zone at 650℃ is estimated to be 5×10~(-16)cm~2/sec. It is shown that the amount of chromium carbide precipitated at the grain boundaries increases nearly logarithmically with time of sensitization, and the precipitation of 1—2% of carbon contained in the steel at the grain boundaries is sufficient to cause severe intergranular corrosion.

Ni, Cr, W, Mo, Al, Ti, B and misch metal were added separately or in combination into pure iron containing 0.166% of oxygen. Optical and chemical properties as well as crystal structure of various types of nonmetallic inclusions present in the alloys were studied. The inclusions identified as in these alloys are CeO_2, FeO (with Ni), FeO (with Mo), Cr_2O_3, Cr_3O_4, FeO·WO_3, α-Al_2O_3, α-Al_2O_3 (with Ni), α-Al_2O_3 (with Ti, Ni, Cr), TiN, α-Ti_2O_3 and a certain kind of boride. The present results confirmed that the composition and properties of a deoxidation product are mainly determined by the affinity between the alloying element used and oxygen.

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