Search Documents

By R. R. Judd, H. W. Paxton

The rate of dissolution of cementite was studied in three low-carbon materials: a zone-refined Fe-C alloy, an Fe-0.5pct Mn-C alloy, and a commercial low-carbon steel. The materials were spheroidized, ad then held isothermally at temperatures above the Al. The isothermal anneal was interrupted periodically by a water quench and the specimens were analyzed by quantitative metallography for the amount of aus-tenite formed during the anneal. The results of this study were compared with an analytical model for the process, which assumes that carbon diffusion in aus-tenite is the rate-controlling step for the cementite dissolution process. The correlation between the model and the experimental data is excellent for the zone-refined Fe-C alloys; however, the Fe-0.5 pct Mn-C alloys and the commercial steel deviate from the calculated model. This deviation is thought to be a result of manganese segregation between the carbide and the matrix. The rate of nucleation of austenite at carbide interfaces was reduced by the manganese addition and enhanced by the presence of ferrite-ferrite grain boundaries. PREVIOUS investigations of the nucleation and growth of austenite from ferrite-carbide aggregates are not entirely satisfying for at least one of several reasons. The most prevalent of these is a lack of quantitative data. Engineering studies have been run on many steels with little control over important parameters such as composition and initial aggregate structure. The data obtained are valid only for material with identical chemistry and thermal history. A more informative approach to the problem of aus-tenitization would be to determine the mechanism that controls the rate of solution of carbide in austenite and how it is modified by alloying elements. This information could then be used to calculate an austeniti-zation rate for any material, provided its composition and structure are known. The object of the present work is to establish the rate-controlling step for cementite dissolution in Fe-C austenite and to investigate the modification of this rate by small manganese additions. The composition and structure of the material used were carefully controlled and all measurements were designed to allow a quantitative analysis of the kinetic process that controls the austenitization rate. A MODEL FOR DISSOLUTION OF CEMENTITE Cementite dissolution has been analyzed mathematically by a model that approximates the material used in the experiments. This model postulates a regular ar-array of identical cementite spheroids with 4 C( diam, embedded in a grain boundary- free ferrite matrix. The analysis provides a detailed description of the dissolution of one carbide spheroid and a generalization of the solution by summation over all the carbides in the material. The carbides may be isolated by defining identical, space-filling cells of ferrite around them. If the cell dimensions are greater than the diameter of the austenite sphere resulting from complete dissolution of the carbide, and no interaction (through diffusion in ferrite) takes place between cells during the dissolution process, the model need concern only one cell, since the solution in each cell is identical. In the experimental material, the dimensions of the cell, the carbide, and the final austenite sphere are approximately 24, 4, and 8 p, respectively; use of the single cell is therefore justified. The experimental observations are made on the austenite nodules that form around each carbide during the dissolution process. The model concerns the growth of these austenite nodules. The attendant shrinking of the carbide can be obtained from the same analysis by an extension of the calculations. Several a priori assumptions are necessary to make the analysis of the growth problem tractable. They are: 1) carbon diffusion through the austenite nodule is the rate-controlling process; 2) local equilibrium exists at all interfaces, 3) the austenite nucleus that forms on each carbide instantaneously envelops the carbide; 4) during the austenite growth process, the diffusion flux of carbon in ferrite is insignificant; 5) a quasi-steady state exists in the austenite concentration field; that is, at any instant during the dissolution process, the austenite carbon concentration gradient closely approximates that for a steady-state solution; and 6) the effects of capillarity on the dissolution rate of the carbides can be neglected. Referring to Fig. 1, a mass balance at the y-a interface for an infinitesimal boundary movement gives: Where rb is the outer radius of the austenite shell, C1 and C are carbon concentrations at the interface in austenite and ferrite, respectively, see Fig. 2, is the diffusion coefficient of carbon in austenite for the concentration of carbon at the interface, and t is time. The fifth assumption permits the austenite carbon concentration to be approximated by the Laplace solution for the spherical case. Therefore, where C(Y) is the carbon concentration at r, and A and B are constants. Local interfacial equilibrium fixes the boundary conditions for the diffusion problem. They are:

Jan 1, 1969

For some years MINING ENGINEERING has published abstracts of the articles appearing in each issue and also of the Transactions papers published in the Quarterly. Beginning this month, a list of key words is being given with each abstract as an aid to indexing the articles for ease in retrieving them when they are needed. The value of a home or office system of indexing and retrieving technical articles is hard to over-estimate. How often is it that a technical problem comes up and someone says: "Oh yes, I've read something about that." Many hours can be expended tracing down that nebulous source. A modest indexing and retrieval system will save those tedious hours, and the person who is able to produce technical documents when they are needed will achieve an enviable status for being on the ball.

Jan 1, 1964

By Klaus Schwerdtfeger

The rules of solution of oxygen from H2O-H2-He gas and of nitrogen from N2-H2 gas in shallow melts of liquid iron were measured at 1610o and 1600o C, respectiuely. Concentration profiles were detemined in the liquid iron. Tire rate data indicate that the solution process is controlled by diffusion in the iron melt. The diffusivities for oxygen and nitrogen in liquid iron, as calculated from the present data, are DFe-o = (12 ± 3) < 10-5 sq cm per sec and DFe-N = 11 ± 2) X 10-5 sq cm per sec at the temperatures employed. AN attempt was made by Shurygin and Kryukl to measure the diffusivity of oxygen in liquid iron. In their experiments a silica disc was rotated in liquid iron containing oxygen, and the rate of formation of liquid iron silicate was measured. Assuming that the rate of dissolution of silica is controlled by diffusion of oxygen in the iron, the oxygen diffusivity was computed from the rate data giving Dfe-0 = 6.1 X 5 sq cm per sec at 1600°C. Although this value seems to be of the right order of magnitude, there is no proof of the correctness of the assumptions involved in the interpretation of these rate data. The oxygen concentration in the iron at the iron-iron silicate interface was taken to be that in equilibrium with the silica-saturated silicate melt. That is, it was assumed that no concentration gradient existed in the liquid silicate. This is a questionable assumption, unless it is proved that the thickness of the silicate layer is very much smaller than that of the diffusion boundary layer in the iron. Furthermore, Shurygin et al.1 used the Levich equation2 to interpret their rate data. This equation was derived for mass transfer between a solid disc and a single-phase liquid. The hydrodynamic and diffusion boundary layers in the iron stirred by a disc, via coupling of the silicate melt, may be appreciably different from those predicted by Levich&apos;s derivations. In the present work the diffusivities of oxygen and nitrogen in liquid iron were measured at 1610" and 1600oC, respectively. EXPERIMENTAL METHOD Iron melts contained in high-purity gas-tight alumina crucibles were reacted with H2O-H2-He gas for the determination of the oxygen diffusivity and with N2-H2 gas for the determination of nitrogen diffusivity. At the end of the reaction period, the samples were quenched in a cold H2-He gas stream at the top of the furnace. Oxygen or nitrogen contents in the iron were determined by chemical analysis. Two different types of diffusion experiments were perforxed. To determine concentration profiles, a few rate measurements were made using 4-cm-deep melts. The solidified samples were sliced into discs and each disc was analyzed for oxygen or nitrogen. In another series of experiments, oxygen or nitrogen was diffused into shallow melts (about 0.5 to 1 cm in depth) and the total sample was analyzed to obtain an average concentration of the diffusate. In most experiments, 4- to 5-mm-ID alumina crucibles were used. Some experiments were also made in smaller (3 mm) and larger (7 mm) diam crucibles. This variation in diameter caused no difference in the reaction rate, within the limits of experimental uncertainty. To promote the establishment of a stable density profile in the melt, all the samples were suspended in the lower end of the hot zone so that the top of the melt was hotter by a few degrees. Molybdenum wire resistance heating was used. The reaction tube of the furnace was a gas-tight recrystal-lized alumina tube. In most experiments the furnace was heated by an ac power supply. To check the possibility of inductive stirring, some experiments were carried out in a dc operated furnace, with essentially the same results. The temperature of the furnace was controlled automatically in the usual manner. The temperature was measured with a Pt/Pt-10 pet Rh thermocouple and is estimated to be accurate within ±5°C. The iron used was prepared by melting and vacuum-carbon deoxidizing electrolytic "Plastiron" in a zir-conia crucible. The main impurities are: Si 0.004 pct P, S <0.002 pct Cr 0.005 pct N 0.001 pct Zr 0.002 pct O 0.003 pct Mn 0.004 pct C 0.002 pct The gas composition was controlled by constant pressure head capillary flowmeters. Oxygen was removed from the gas mixture by passing it through columns of platinized asbestos (450°C) and anhydrone. Selected H2O contents were obtained by passing the purified gas through oxalic acid dihydrate-anhydrous oxalic acid mixtures held at constant temperature in a water bath. Water vapor pressure data for the oxalic acid dihydrate-anhydrous oxalic acid equilibrium were taken from the 1iterature.3 The flow rate used was about 1.5 liters per min. The whole system was checked for tightness at regular intervals.

Jan 1, 1968

By R. L. Stephenson

The effect of fabrication variables on the creep-rupture properties of TZM (Mo-0.5 pct Ti-0.08 pct Zr-0.03 pct C), Cb-TZM [columbium-modified TZM (7140-1.5 pct Nb-0.5 pct Ti-0.8 pct Zr-0.05 pct C)], and TZC (Mo-1.2 pct Ti-0.25 pct Zr-0.15 pct C) has been studied at 982º. 1093º, 1204º, and 1315°C for rupture times up to 1000 hr. The creep tests were performed in vacuum creep furnaces, employing cold-trapped oil-diffusion pumps, operating at pressures 2 X 10-7 tory. After-test chemical analyses showed no significant changes in interstitial content. The creep-rupture properties of TZM, fabricated according to current mill practice, were determined in the stress-relieved and in the re crystallized conditions to provide control data. Specimens of TZM, Cb-TZM, and TZC fabricated according to each of two experimental fabrication schedules were then tested. The schedules differed in terms of in-process annealing temperatures and warm -working temperatures. The data make several conclusions evident. Fabrication variables do indeed have a profound effect on the long-time creep-rupture properties of molybdenum alloys. Strengths of these alloys compare very favorably with those of other refractory metal alloys in the temperature range studied. FOR several years we have been engaged in the evaluation of the creep-rupture properties of promising refractory metal alloys. We consider the determination of these properties to 1000 hr at several temperatures necessary for a valid comparison. Such an evaluation indicated that the creep strengths of Mo-0.5 pct Ti and TZM (Mo-0.5 pct Ti-0.08 pct Zr-0.03 pct C) compare favorably with those of other refractory metal alloys. Tensile and recrystallization data of other investigators1,2 show an effect of fabrication variables. More specifically, those fabrication schedules which produce high strength through work hardening exhibit the least thermal stability. A similar conclusion was reported by perkins3 for Cb-TZM. This would lead one to expect that, in terms of creep, those materials showing higher strength at short times and low temperatures would exhibit lower strength at long times and high temperatures. If these expectations were fact, a substantial improvement could be made in these alloys by optimizing their fabrication schedule in terms of their anticipated service temperature. We therefore undertook a study of the effect of fabrication variables on the long-time creep properties of TZM, Cb-TZM (Mo-1.5 pct Cb-0.5 pct Ti-0.08 pct Zr-0.05pct C) and TZC (Mo-1.2 pct Ti-0.25 pct Zr-0.15 pct C). MATERIALS All materials used were supplied by the Climax Molybdenum Co. of Mich. Our original intention was to test the same heat of each alloy in different conditions. Fabrication difficulties prevented this; however, heats with nearly identical analysis were selected. The compositions of the materials used are listed in Table I. All materials were supplied in the

Jan 1, 1970

By R. Lagneborg, R. Gullberg

An X-ray method for determinig volume fractions of phases in textured materials has been developed. The method involves measurements of the integrated intensities of reflections of each phase for one single orientation of the sample. The method was checked by evaluating the phase contents in samples from two stainless Cr-Ni steels containing austenite and ferrite in known amounts. In the present case measurements of about five reflections from each phase were required. The agreement between calculated and known volume fractions was found to be very close. AN X-ray diffraction method is often employed for determining the amounts of two or more phases in a material. The details of the method have been described at several instances.&apos;" In short, the method is based upon the fact that the integrated intensity of a reflection is directly proportional to the volume fraction of the phase considered. In principle, the method only requires the measurement of the intensity of one reflection for each phase. The most severe shortcoming with this method is that it disregards influences from preferred orientation in the material. In a diffractometer, where this type of analysis is most conveniently carried out, the geometry is such that the diffracting planes for a specific Bragg peak are all parallel to the surface of the specimen. This implies that it must be required that the occurrence of the diffracting planes parallel to the specimen surface must be as frequent as for any arbitrarily oriented physical plane of the sample. This requirement will generally only be fulfilled for a material free of texture. Since all materials in reality always exhibit a preferred orientation to some extent. the deficiency of the method is quite serious. In cold-deformed material the texture is appreciable, but also annealed material will in many cases display a considerable texture. The problem of quantitative phase analysis in textured material is one of obtaining the integrated intensity averaged over all orientations of the specimen with respect to the X-ray beam. Lopata and ~ula~ have shown that this can be accomplished by integration of intensities over the whole pole figure, or by random rotation of the specimen in the X-ray beam, or by a combination of both. Hence, in this method the average integrated intensity for a specific reflection is obtained by directly determining the intensity for a sufficient number of physical orientations of the sample. However, since an excess of diffracting planes in a certain surface plane necessarily must be accompanied by a deficiency of another type of diffracting plane, it appears to be possible to obtain the correct average intensity for a specific diffracting plane by performing an appropriate averaging of intensities belonging to a large number of diffraction planes of different hkl when all intensities have been produced from one single orientation of the specimen surface. In the present work such a method has been attempted and tested experimentally for two stainless Cr-Ni steels containing austenite and ferrite. 1) THEORY The expression for the integrated intensity of a Bragg reflection diffracted from a phase in a polycrys-talline sample in a diffractometer is: where C is a constant only dependent on the incident X-ray power and the diffractometer geometry, v the volume of the unit cell, F the structure factor, m multiplicity of the diffracting plane, L Lorentz factor, P polarization factor, e-2 temperature factor, p average linear absorption coefficient of the irradiated vO1umej and V volume fraction of the phase &apos;Onsidered. Hence, the factor R in the abbreviated form of Eq. [I] can be calculated and depends only on 0 and the diffracting plane {hkl}. For an alloy containing two phases, a and y, we can write

Jan 1, 1967

By T. M. Geffen, H. N. Hall

One-third to one-half of the original oil is left in most reservoirs even after thorough gas driving or waterflooding. In contrast, essentially all of the oil contacted con be recovered by flooding with a solvent (a fluid miscible with the reservoir oil). It is not economical to fill a reservoir with solvent since the value of the solvent would very likely exceed that of the oil. A practical method would be to inject a small bank of solvent to displace the oil. This solvent bank in turn could be driven through the rock by a less valuable scavenging fluid, such as natural gas which is also miscible with the solvent. Such a method of oil recovery was studied in the laboratory using long (up to 95 ft) core systems. It was found that, as the solvent moves through the reservoir, the front end of the solvent bank becomes mixed with the oil and forms a relatively short mixing zone between undiluted crude oil and pure solvent. An engineering correlation was obtained which can be used to predict the length of this zone. Displacement of the solvent by natural gas can be achieved with complete miscibility (above the critical pressure of the gas-solvent system at reservoir temperature) or with partial miscibility at lower pressures. Under conditions of complete miscibility, it was found that the length of the gas-solvent mixing zone could be predicted from the correlation developed for the solvent-oil mixing zone. With partial miscibility, two phases exist and the solvent is displaced by a modified gas drive in which vaporization and condensation occur. Equations are developed to predict the performance of this type of displacement. Two plans are suggested for application of miscible fluid displacement. one for miscible displacement of the solvent and the other applying to lower pressures where partial miscibility occurs during solvent displacement. Calculations based on the laboratory tests indicated that a bank of solvent with a volume of less than 10 per cent of the hydrocarbon volume under consideration will generally be adequate to obtain high oil recovery from the portion of the reservoir contacted in miscible fluid displacement. INTRODUCTION The conventional methods by which oil is displaced from reservoir rock—solution gas drive, conventional gas drive, and waterflooding—are all processes in which the oil is expelled from the rock by direct displacement with a fluid immiscible with the oil. Because of the immiscible type displacement, complete recovery of the reservoir oil is not achieved. In economically successful flooding operations, as much as 50 per cent of the original oil may be abandoned in the reservoir as unrecoverable. The use of a solvent (miscible with oil, such as LPG) as an injection fluid does not have the undesirable displacement features common to the use of gas or water (immiscible with oil). It can be readily appreciated, however, that it would not be economically feasible to displace the oil and entirely replace it with a solvent. A process was therefore devised for use in field operations in which only a small volume of solvent would be required to obtain essentially complete oil recovery in the portion of the reservoir contacted. In this process, the solvent acts as a buffer between the reservoir oil and injected gas, allowing the natural gas to displace indirectly the oil with much greater efficiency than if the gas displaced the oil directly. DESCRIPTION OF OIL RECOVERY PROCESS The solvent oil recovery process will be described in

Jan 1, 1958

By D. N. Dietz

A method for determining average reservoir pressure is presented, which is simpler to apply than that devised by Matthews. Brons and Hazebroek. For bounded reservoirs, identical results are obtained if stabilized-flow conditions prevail. The present method yields inferior results in the transient state. The method can, with a slight modification, also be used for water-drive reservoirs. INTRODUCTION In the method proposed by Matthews, Brons and Haze-broek&apos; for determining average reservoir pressure in a multi-well reservoir, the cumulative production (the production time of each well) enters twice: once when the buildup is plotted against In (t + t)/At to arrive at p*, and a second time when the correction, p*-p is determined with one of the several formulas for differently shaped drainage areas. Once a steady state has been attained, however, the previous production history should be immaterial. The same pressure distribution could have been arrived at after different cumulative productions of the individual wells. In principle, therefore, it should be possible to determine average pressures without referring to cumulative productions. In the following paragraphs an expression is presented for the difference between the pressure in a producing well and the average pressure of its drainage area. Then the build-up time needed to overcome this difference is indicated. CASE OF A CIRCULAR DRAINAGE AREA AND A CENTRAL WELL PRESSURE DISTRIBUTION BEFORE SHUT-IN The general differential equation of radial flow (see Muskat, Eq. 10.2) may be written In the steady state, the rate of production of the well is equal to the rate of expansion of the fluid contained in the the drainage area; thus

Jan 1, 1966

By C. Richards Honeycutt, Frederick C. Langenberg, Guenter Pestel

Solidification in a moving electromagnetic field was successful in altering the as-cast grain structure of steel ingots. The equipment is described and exerimental results are presented for several different alloys. Conditions for optimum grain refinement are described and an empirical design constant presented. The effect of the grain refinement on hot workability is briefly discussed. T HE grain structure of a steel ingot teemed into a metal mold can be defined by macroscopic examination in semiquantitative terms. For example, the relative volume of obviously differing structures can be estimated with the classic definitions of chill, columnar, and equiaxed zones. Experimental evidence has proven that the grain structure of an ingot is affected by changes in such factors as teeming temperature, rate of teeming, degree of turbulence, size of ingot, mold characteristics (including preheat temperature, design and coating), and composition of the metal (especially as it affects the range of solidification). Significant advances have been made in fundamental solidification studies; but the principles which have been developed are difficult to apply tothe complex and changing phenomena that occur simultaneously during ingot solidification. In particular, many attempts have been made to refine the as-cast structure of ingots, with special emphasis on the columnar grains. The formation of large columnar grains is generally considered harmful to the further fabrication of an ingot. It may also result in undesired ani-sotropic properties. However, columnar grains are not undesirable in all grades, or product applications, and caution must be exercised to determine in which grades and applications columnar grains are not desired in the ingot. Although the literature on the refinement of as-cast ingot grain structures is too vast in general to review here, it is noteworthy that such methods are additions of artificial nuclei, mold vibration, mechanical stirring or ultrasonic agitation of the molten metal, and controlled cooling and reheating during solidification have been tried with varying degrees of success. However, it is doubtful that any method of refining the as-cast ingot grain structure has been completely successful on a commercial basis. The purpose of

Jan 1, 1962

By Raymond D. Sloan

During 1944, substantial gains were recorded in practically every phase of the petroleum industry in Oklahoma. With the spotlight of activity focused on other states during the more recent years, the production momentum that was achieved in Oklahoma through the discovery and development of such outstanding pools as Oklahoma City, Fitts. and major pools in the Seminole area during the late 1920&apos;s and early 1930&apos;s was practically spent by the end of 1943. However, with much of the industry&apos;s exploratory efforts greatly reduced in Illinois, Indiana and Kentucky, the strong demand for Mid-Continent crude and the optimism created by some discoveries in the state toward the end of 1943, Oklahoma has again become the recipient of much activity and interest, with its 1944 drilling activity, production and refinery runs to stills registering substantial gains over the preceding year. Development and Exploration Since the discovery of the West Edmond pool in April of 1943, with the substquent drilling and development boom, the scene of exploratory activity has shifted from the central. south and east portions of the state to the general trend of the . Granite Ridge, where leasillg and wild-catting were greatly activated within a wide radius of this high subsurface feature. Leasing activity, however, jumped the confines of the Granite Ridge trend and extended into all of the northwestern counties in the state. Seismograph activity at once began following and keeping pace with this shift in the leasing and exploration locale, by the concentration of a large number of seismograph crews in the area. At one time during the year,35 crews were concentrated in Oklahoma, the majority being west of the Granite Ridge and extending into the northwestern counties. To illustrate the increase in geophysical work done in the state during the year, records indicate that 377 crew months were worked in 1944, which is a 21 per cent increase over the 311.5 crew months activated in 1943. With the discovery in 1943 of several encouraging prospects and the finding of West Edmond as an incentive, drilling operations increased some 54.3 per cent, for a total of 1986 completions during the year. This is compared with a figure of 1287 for wells drilled in 1943, an increase of 699 wells. The industry, with a total of 1064 oil producers, enjoyed a success percentage of 53.6 per cent. Gas wells completed were 223, while 699 dry holes were drilled. In 1943, completions included 587 oil wells, 127 gas wells, and 573 dry holes. The number of oil producers completed during the Year was almost double the numher of wells in 1943. Of the total drilling wells completed in the state, 350, or 17.6 per cent, were wildcat operations—an increase of 5.1 per cent over the 194.3 total of 333. The wildcat success percentage was 13.4 per cent, 47 successful wells being completed for an average initial potential of 160 bbl. In the state, 23 gas wells were completed with an average initial production of

Jan 1, 1945

By Charles H. Jacoby, Stanley J. LeFond

Salt, or halite, has had a long and most varied history. While we know the Chinese were producing salt as early as 3000 B.C., the first written reference to salt appears in the book of Job recorded about 2250 B.C., which reads: "Can nothing which is unsavory be eaten without salt?" We do not know when man first used or realized that he needed salt. Perhaps one of our early Eolithic ancestors noted a piece of salt shining along the shore of a salt lake or a bay. He may have been attracted by its shining luster and tasted it. He found that it had a tart, zesty, distinctive flavor and perhaps chewing a large piece, became ill. However, by experimentation he soon learned that taken in the proper proportions salt was a very pleasant additive to his food. He may have noticed, too, that it increased his feeling of well-being. Thus, when the curtain of recorded history rises, we find man using salt for his health&apos;s sake, seasoning his food with it, feeding it to his animals, and using it in many superstitious and religious ceremonies. These uses, however, are a far cry from the uses for salt today. The first attempts of white man to make salt in continental United States were reported in the year 1614, while the first commercial production was in 1790. Practically all of this salt was for human consumption, preserving food, and curing hides. Prior to the coming of the "white man" Indians are known to have made salt from such salt springs as those that existed at Avery Island, La., Charleston, W.Va., Saltville, Va., and Onondaga, N.Y. In 1862 the first rock salt was mined in the U.S. from the Avery Island salt dome. Almost immediately thereafter, rock salt was mined at Saltville, Va., and then at Retsof, N.Y., in 1885. Of the more than 40 million tons produced annually in the United States, about 68% or 27 million tons are consumed by the chemical industry in the production of chlorine, caustic soda, hydrochloric acid, sodium metal, and some 30 other basic chemicals. These basic chemicals, in turn, are used in the preparation of approximately 14,000 chemicals which range from soap and detergents, to such chlorinated hydrocarbons as DDT, BHC, carbon tetrachloride, etc.

Jan 1, 1975

By A. A. Watts, R. I. Jaffee, F. C. Holden, H. R. Ogden

Hypoeutectoid Ti-Cu alloys are responsive to heat treatment, and considerable variation of mechanical properties may be produced by transformation of the ß phase. Control of cooling rate, isothermal transformation, or quench tempering determine the transformation structure and properties. The transformation products are similar to those of carbon steel. THE equilibrium diagram for the Ti-Cu system has been studied by various investigators, and the most recent diagram is that determined by Jou-kainen, Grant, and Floe.&apos; The titanium-rich portion of this diagram is reproduced in Fig. 1. The alloy compositions and annealing temperatures used in this work are shown. The Ti-Cu alloys containing up to 17 wt pct Cu are representative of an active eutectoid alloy system. Joukainen et a1.l have stated that ß phase cannot be retained on quenching and this has been confirmed in the present work. The similarity between the Ti-Cu system and the Fe-C system is evident and, as will be shown later, many of the transformation products appear to have similar microstructures. Heat treatments in this program were designed to produce the structures of primary interest for study and determination of mechanical properties. Experimental Procedures Preparation of Alloys: Five ½ lb ingots, ranging in composition from 0.8 to 6.83 pct Cu, were prepared by double arc melting in an atmosphere of high purity argon. Titanium melting stock was obtained from high purity iodide titanium, and copper additions were made from copper wire clippings. Segregation was reduced by inversion melting, in which the once-melted ingots were inverted in the crucible and remelted. The double-melted ingots were forged at 1600°F to ¾ in. diam rods, with both heating and forging operations being carried out in air. Surface scale was removed by grit blasting and grinding, after which the forgings were vacuum annealed at 900°C. The alloys then were swaged to ¾ in. diam rod at 750°C and test specimens were cut from this material. The vacuum annealing operation was to remove dissolved hydrogen, which is commonly present in the iodide titanium as-received. Although vacuum fusion analyses were not obtained for all these alloys, past experience indicates that the 6 hr treatment with a limiting pressure of 10- to 10-5 mm of Hg is sufficient to reduce the hydrogen level to 10 to 30 ppm. Vacuum fusion analysis of the 0.8 pct Cu alloy showed 19 ppm of hydrogen. In addition to the five hypoeutectoid alloys, three 15 g ingots were prepared with nominal copper additions of 9, 10, and 11 pct. Fabrication procedures were similar, except that these alloys were rolled to 0.040 in. sheet. Microstructure specimens only were obtained. Table I shows analyzed compositions and fabrication temperatures for these alloys. Heat Treatments: All heat treatments were made in potentiometer-controlled, resistance-wound tube

Jan 1, 1956

By G. Sachs, E. P. Klier, W. Beck

HYDROGEN embrittlement of steels has recently attracted much attention because it is asso- ciated with a variety of failures, especially those of aircraft structural components.&apos; For instance, one of the undesirable properties of the ultra-high strength steels used for landing gears is their pronounced susceptibility to hydrogen embrittlement during plating. The presence of hydrogen in these steels becomes particularly apparent through delayed or sustained load failures of parts which have been subjected to low loads for an extended time. A number of variables are involved in causing a reduction of strength and ductility through hydrogen

Jan 1, 1957

By Paul H. Lu, Fred D. Wright

Numerous laboratory tests and field measurements conducted by different persons have indicated that most mine rocks behave viscoelastically rather than elastically. Hence, for an accurate and realistic design of underground structure, the time factor should be taken into account, and the theory of viscoelasticity can be used. By considering the time factor, the time-safe-strength may be used for economical design instead of liberal safety factors with design based on the elasticity theory. The deflection and moment distributions on an elastic roof slab resting on elastic pillars have already been studied.l In this chapter, the stress- strain analysis for a viscoelastic mine roof resting on pillars of the same material or on supports of different material is first considered theoretically. Viscoelastic solutions for deflections, moments, and shears are de- rived from the elastic plate theory on the basis of the elastic-viscoelastic correspondence principle. The Laplace transformation technique is employed as the operational method in the mathematical treatment. The roof slab materials are represented by four different rheological models and four different loading histories are considered. The method of finite difference equations is employed for solving the partial differential equations, using FORTRAN programs to calculate roof deflections, moments, and shears. Numerical examples are given for one of the four models under two of the four loading histories with regular and staggered pillar patterns. The results are discussed in terms of the variance with time of deflections, moments, and shears as well as their variation in space due to the dimensions and configuration of the structure system. A viscoelastic slab resting on viscoelastic pillars or on elastic supports

Jan 1, 1972

By L. C. Michels, O. G. Ricketts

The disorientations between s?nall grains, whose growth has been arrested at an early stage of recrys-tallization, and the deformed matrix in cold-rolled aluminum single crystals were determined using transmission Kikuchi line and electron diffraction patterns. The orientations of the recrystallized grains were found to be random, and the disorientations of these grains with the matrix weve found to be intermediate to large. This leads to the conclusion that the observed vecrystallization began in small areas of large disorientation present in the cold-worked structure. heavily cold-worked thin sections of aluminunz single crystals and of polycrystalline aluminum and nickel were produced directly by a mechanical technique. The specinlens thus prepared were heated with the electron beam to bring about vecrystallization during observation in the electron microscope. Motion pictures taken du.ring heating and the electvon, microg.raphs taken both before and aftev heating allowed the recrystallization process to be traced to its ovigin. Re cvystallized grains originated in very s,mall regions of the cold-worked structure and developed through rapid migration of high-angle boundaries. The boundaries either were present as such in the matrix or were formed out of dense dislocation networks. SIGNIFICANT advances have been made in recent years in the study of nucleation of recrystallization using the technique of transmission electron microscopy of thin metal foils. Bollman1 in a study of heavily rolled polycrystalline nickel found support for the Cahn-Cottrell2,3 theory of nucleation. According to this theory nuclei form by the initially slow growth of subgrains formed through polygonization. During this initial period of slow growth (the incubation period) the migrating boundary of the subgrain increases its disorientation with the cold-worked matrix and thereby increases its mobility to become a rapidly migrating high-angle boundary. Bailey4,5 investigated the annealing behavior of several metals deformed both in tension and by rolling and concluded that recrystallization took place through the migration of high-angle boundaries. With low deformations these boundaries were present in the metal before deformation. With high deformation it was not possible to tell whether the boundaries were pieces of the original grain boundaries or were produced either during deformation or by polygonization during ameal- ing. Direct observation during heating of metal foils indicated that subgrains form by polygonization and grow at an uneven rate. The grain size obtained decreased with decreasing foil thickness indicating that the foil surface resists boundary motion. Votava,6 in heating stage experiments on rolled copper, observed nuclei to appear suddenly and grow in jumps of differing magnitude. However, he found no special dislocation configurations where the nuclei appeared. Fujita,7 as a result of a study of subgrain growth in heavily worked aluminum, concluded that the boundary of a recrystallized grain initially forms from the boundary of a group of subgrains. This occurred by a process of deposition of vacancies and dislocations in the group boundary as the boundaries within the group disappear. HU8,9 directly observed a similar process in heating stage experiments on 70 pct rolled Si-Fe single crystals. The growth of subgrains appeared to proceed by a coalescence mechanism. The observed fading away of the boundary between two subgrains was explained by the moving out of dislocations from the disappearing boundary into the connecting or intersecting boundaries around the subgrains. The subgrain size and degree of disorientation with the surrounding structure were thus increased. With the increase in disorientation occurred a corresponding increase in boundary mobility, which eventually allowed the boundary to migrate rapidly. This process was observed to occur within "microbands" consisting of parallel narrow segments disoriented by a few degrees present in the as-rolled structure. The conclusion of Rzepski and Montuelle10 that growth is preceded by the coalescence of blocks through disappearance of their common boundaries supports this view. In contrast to Hu&apos;s coalescence model for nucleation were the conclusions of Walter and ~och.""~ Working with the same material as Hu, of the same orientation and rolled to the same reduction, they concluded that nucleation occurred by the Cahn-Cottrell mechanism. They observed, in agreement with Hu, that recrystallization began in the "microband" regions which they referred to as "transition" bands. Bartuska13 studied subgrain growth in heavily rolled nickel using a beam heating method in the electron microscope. He concluded that nuclei for recrystallization form from the largest most perfect subgrains present in the cold-worked structure by rapid intermittent migration of parts of subboundaries. In rare instances he observed subgrain growth by coalescence. EXPERIMENTAL PROCEDURE The materials used in this study were 99.999 pct A1 supplied by A.I.A.G. Metals, Inc., and 99.999 pct Ni supplied by Johnson and Matthey and Co., Ltd. The Hitachi HU-11 electron microscope, with uniaxial

Jan 1, 1968

By D. A. Martin, Laird Crocker, W. I. Nissen

The Federal Bureau of Mines has developed a process for removing SO2 from stack gases. The process comprises absorption of SO2 in an. aqueous solution of sodium citrate, citric acid, and sodium thiosulfate, followed by reaction of the absorbed SO2 with H2S to precipitate sulfur and regenerate the citrate solution for recycle. A pilot plant was operated at The Bunker Hill Co. lead smelter in Kellogg, Idaho, to assess the feasibility of using the process for flue gas desulfurization. Nominal capacity of the pilot plant, which treated .tail gas from the lead smelter sintering furnace, was 1699 standard meters cubed per hour (15 degrees celsius) of 0.3- to 0.7-percent-SO2 gas yielding about 272 net kilograms of sulfur per day. The pilot plant was operated for about 4500 hours and produced more than 45 net tonnes of sulfur. The operation demonstrated that (1) more than 95 percent of the SO2 could be removed from the lead smelter sintering furnace tail gases, (2) regeneration of the citrate solution with H2S and precipitation of sulfur in conventional stirred vessels was readily controlled and highly efficient, (3) the precipitated sulfur could be continuously recovered as a 99.5-percent-pure product by kerosine flotation and melting, and (4) the 77- to 79-percent-H2S gas used for sulfur precipitation could be readily produced from pilot plant product sulfur, natural gas, and steam. Projected operating costs, including capital charges, for a citrate plant, installed at a 118 000-tonne-per-year lead smelter and treating 49 271 standard meters cubed per hour of lead smelter sintering furnace tail gas, are estimated to be $384 per tonne of sulfur.

Jan 1, 1976

By George H. Gilman

THE campaign now being waged to improve the quality of the rock-drill bit is the natural outcome f the scientific development of the drilling machine during the past twenty years. In this development there has been a great increase of power output per unit of weight of machine. But the development of the rock drilling engine has been seriously hampered by the need of a drill bit that would withstand, without undue wear or breakage, the heavy rapid blows of the improved pneumatic hammer. This problem was difficult to solve because an increase in the weight or section of the drill-steel bar resulted in decreased cutting speed and higher operating costs of both the bit and its actuating engine. The solution, therefore, is to make the bit of steel that, for a given weight or section, will be more shock and wear resisting. In the attempt to meet these requirements, many of the drill-steel manufacturers alloyed the steel with vanadium, chromium, nickel, etc., on the assumption that such elements would improve the rock-drill bit but due to the sensitiveness of such steel to forging and heat treatment, higher cost, and lack f continuity of effort on the part of the steel manufacturer to perfect this product, but little progress has been made in the adoption of these steels and the so-called straight carbon steel is today recognized as the standard material for all rock-drill bits. The first step in improving drill-bit efficiency is the determination f the cause of breakage and wear. For this reason, in 1917, certain rock-drill manufacturers, in conjunction with some of the larger mining interests, conducted exhaustive tests in both the field and laboratory. It was then found that, in the main, the cause of drill-steel troubles is not attributable to the duty imposed upon it by the normal operation of the drilling engine, but to improper forging and heat treating of the bit and to the subjecting of the bit to improper use. The fact that it is possible to minimize the troubles attributable to improper heat treatment of the drill steel has been demonstrated and it is recommended that such knowledge be made available by teaching those interested in the production, use, and up-keep of drill bits the principles on which success in this direction may be achieved.

Jan 6, 1921

By W. A. Forbes

by the combustion of this gas as it reached the air was a familiar sight in the days when open-top furnaces were in vogue. As blast-furnace practice progressed, however, involving the use of hot blast and the development of power from the gas, blast-furnace construction was modified to conserve this gas for heating hot-blast stoves, raising steam under boilers, and various miscellaneous purposes. When the use of blast-furnace gas was extended to the driving of gas engines, the necessity for thoroughly cleaning the gas became apparent., and soon the advantages of also properly cleaning the gas for use in stoves and udder boilers became more fully recognized.

Jan 10, 1913

By W. J. Slatosky

As a means of effecting better control of endpoint temperatirres at the Jones & Laughlin basic oxygen furnace plant, a set of mathematical equations has been developed. The eqlutions are the product of a themlochemical anaysis of the process and aye designed to calculate the required scrap, lime, and hot metal additions in terms of a number of independent variables. Results of test heats have warranted adoption of this technique by the Prodrrction Department. BECAUSE of the autogeneous nature of the basic oxygen steel-making process, bath temperature can be controlled without an external fuel supply by charging the furnace with additions that are thermally balanced. The thermal requirements of the charge materials are such that, during the refining process, they throttle the heat generated by the metallurgical reactions in a manner designed to result in a speci-fied temperature at the completion of the heat. In the past, operating personnel at the basic oxygen furnace plant of Jones & Laughlin&apos;s Aliquippa Works relied on their experience and technical knowledge of the process to determine the quantities of charge additions needed to result in a finishing temperature in the range 2880"to 2920" F. (The charge consists primarily of 93 tons of scrap and hot metal plus an amount of lime sufficient to maintain a basicity ratio of 2.8 to 3.2). Estimates of these materials are based on a consideration of the effects on finishing temperature of 1) iron silicon content, having a variation of 0.8 to 1.8 pct; 2) iron temperature, ranging from 2250°to 2600°F; and 3)any excessive cooling of the furnace due to a production delay. The end temperature of the preceding heat also serves as a guide in that, if a heat was within the specified temperature range, the succeeding heat could be charged with materials of nearly the same proportions, provided the hot metal used in each of the two charges was of approximately the same temperature and composition. On the other hand, if a heat was outside the specified tapping range, or if the hot metal used in that heat was of different analysis and temperature from that of the iron to be charged, an adjustment in the proportion of additions is in order for the following heat. Due to the complex thermochemical behavior of the process and to the inexact and subjective nature of the described method of determining charge additions, consistently accurate temperature control was not to be expected. Therefore, those heats out- side the specified tapping range necessitated subsequent adjustments by either reblowing the cold heats for a suitable length of time so as to elevate the bath temperature to the desired level, or cooling hot heats with a proper amount of scrap. Because extra time is required to make these adjustments, production is delayed. In an attempt to devise a method for improving temperature control, an analysis of the thermochemistry of the process was undertaken. This, in turn, led to the development of a set of mathematical equations which enable the calculation of the quantities of scrap, lime, and hot metal needed to result in any specified tapping temperature range. The analysis was not intended to be a repetition of work done by others such as McMulkinl or ~hilbrook.&apos; It was meant to be an extension of their work so that charge additions could be calculated not in terms of silicon alone but, rather, as a function of all independent variables. This paper presents the derivation of these relationships, their effectiveness in controlling bath temperatures, and a method of utilizing them on an operational basis. The Heat Balance—The first step undertaken in the analysis of the problem was the enumeration of the pertinent variables. A list is presented in Table I where it is noticed that these quantities have been separated into the following three categories: important variables, variables considered as constants, and variables to be neglected. The breakdown was an arbitrary one designed to facilitate the analysis; otherwise, the mathematical treatment would have been exceedingly cumbersome and complex. Fortunately, experience has shown that these simplifying assumptions do not seriously impair the accuracy of the calculations. These variables along with the limiting assumptions listed in Table n were then used to write a heat balance of the process by applying the equation of continuity, Rate of Rate of Rate of Increase = Income - Outgo PI ] of Heat of Heat of Heat.

Jan 1, 1962

By R. E. Ogilvie, A. R. Kaufmann, S. H. Gelles

The solid-solubility limits of iron in beryllium were determined between 850o and 1200oC by analysis of differential type multiphase diffusion couples, using an X-ray absorption technique. The maximum value of the solubility limit was found to be 0.92 ± 0.02 at. pct (5.46 wt pet) at the eutectic temperature 1225°C. The solubilities of nickel and beryllium were determined between 900°and 1200°C by the same technique and the maximum solubility was found to be 4.93 + 0.01 at. pct (25.2 wt pet) at the eutectoid temperature, 1065°C. A previously unreported high-temperature phase which decomposes eutectoidally at 1065 °C was found to exist in the beryllium-nickel system at a composition of approximately 8 at. pct Ni (36 wt pet) by diffision-couple analysis. The presence of this phase was confirmed by thermal analysis and metallo-graphic analysis of the structure resulting from the eutectoid decomposition. G. V. Raynor1 has treated the solid solubilities of some of the elements in beryllium on the basis of the "Hume-Rothery" rules2 which have been modified to include ionic size and ionic distortion effects. It was predicted that the solubility of iron and nickel in beryllium should be slightly less than that of copper. The lowering of the solubility, according to Raynor, is due to a more unfavorable relative valency effect and an ionic size effect. Kaufmann and corzine3 have compiled data on the solubilities of elements in beryllium and have discussed them in the light of the Raynor paper. These authors feel that, because the elements having the greatest solubility in beryllium systematically fall in the Group VIII and IB Columns of the periodic table, the electronic structure greatly influences the maximum solid solubility of elements in beryllium. The solubility of iron in beryllium was determined by Teitel and cohen4 as part of the study of the beryllium-iron phase diagram. The determination was carried out by X-ray and thermal analysis and according to the phase diagram presented, the maximum solubility of iron in beryllium is 0.41 at. pct (2.5 wt pct) at 1225oC. However, it is estimated that the uncertainty in the position of the a-beryllium primary solid-solution boundary is about 0.5 at. pct (3wtpct). Losana and Goria3 in studying the beryllium-nickel phase diagram, determined the solid solubility of nickel in beryllium by thermal analysis. They found the maximum solubility to be between 1.65 and 2.65 at. pct (10 to 15 wt pct) at 1240°C. This value decreased rapidly with decreasing temperature. In determining approximate ranges of solubilities for different elements in beryllium, Kaufmann, et al,8 reported a value of between 1.3 and 1.7 at. pct (7.9 to 10.1 wt pct) for the solubility of nickel in beryllium. The value was obtained by metallographic examination of quenched alloys and lattice-parameter measurements. However, the authors also noted a single-phase structure for a 1.7 at. pct Ni alloy (10 wt pct) on cooling from the liquid. This would indicate a higher solubility range than was reported. ~isch,&apos; in his X-ray studies of beryllium-copper, beryllium-nickel, and beryllium-iron intermetallic compounds, reports the disappearance of a second phase (Ni,Be2) in the beryllium primary solid solution at approximately 4 at. pct (20 wt pct). THEORY The analysis of concentration gradients in diffusion couples has proven to be a useful tool in determining phase equilibria.8-14 In this particular study the diffusion couples were chosen to straddle the expected composition range of the phase boundary, then heat treated at a given temperature and the concentration gradient evaluated. The composition of the phase boundary for a given temperature appears at a point of discontinuity of the composition gradient. Examples of typical phase diagrams and the concentration gradients which should be found in such systems are shown in Fig. 1. In the present work, gradients of the form of Fig. l(c) were obtained in diffusion couples made of pure beryllium and two-phase alloys of beryllium with either iron or nickel. The composition at the point where the gradient becomes discontinuous, Cs, corresponds to the solubility limit of either iron or nickel in beryllium. The analysis of the concentration gradients was carried out by an X-ra absorption method developed and applied by Ogilvie and later used by Moll13 and Hilliard.l4 It depends on the fact that the absorption of X-rays by matter is determined by the concentration and type of the various atomic species present. The relationship for the intensity, I, of a monochro-

Jan 1, 1960

By Richard S. McCaffery

A Research on iron blast-furnace slags has been in progress since 1922 in the metallurgical department of the University of Wisconsin, supplemented by observations, tests and critical examination of operating records at 10 furnaces. From the data thus collected information has been put into forms suitable for use by furnace operators and is given in this report and the three accompanying papers.&apos; In the viscosity results reported, it is estimated that temperatures were measured within 5" C. and that the viscosity determination of the particular slag investigated is accurate within ± 4 per cent. In the viscosity diagrams in the various papers, the equal viscosity lines were drawn in such a manner that in all cases they are located correctly with respect to the actual experimental results, but when determining a viscosity for a particular composition from the diagrams, the viscosity may. be slightly less accurate than the experimentally determined viscosity for the same slag, for the reason that the iso-vis lines interpolated between the experimental results may be shifted slightly without violation of the experimental data. The viscosity diagrams represent our judgment of their position under the above conditions, and we believe the relative accuracy of viscosity data taken from the diagrams is more than sufficient for any use to which the data may be put. Hundreds of samples of slags synthetically produced in the laboratory and of actual furnace slags have been studied in various ways. The ranges of compositions of slags used for viscosity measurements are: silica 25 to 55 per cent.; alumina 0 to 35 per cent.; lime 25 to 55 per cent.; magnesia 0 to 20 per cent. These measurements were made at temperatures of 1250°, 1300°, 1350°, 1400°, 1450°, 1500°, 1550°, 1600" and 1650" Centigrade. Descriptions of apparatus, methods used and statements of the problems investigated have been published in the papers listed in the bibliography beginning on page 120.

Jan 1, 1932