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By A. Eilers

THE telluride of gold and silver, the " Tellurgold-silver" of Hausmann, and " Petzite" of other mineralogists, has been found in so few localities, that a late discovery of the mineral in a new locality may not be uninteresbing. So far this mineral has been found in Europe only in a few mines in Transylvania, and in Hungary; and in the United States in the Stanislaus Mine, in Calaveras County, California, in the Golden Rule Mine, Tuolumne County, California, and in a mine in Montana, the precise locality of which I have forgot, ten. The new locality referred to is at the Red Cloud Mine, Gold Hill District, Boulder County, Colorado.

By A. Eilers

THE telluride of gold and silver, the " Tellurgold-silver" of Hausmann, and " Petzite" of other mineralogists, has been found in so few localities, that a late discovery of the mineral in a new locality may not be uninteresting. So far this mineral has been found in Europe only in a few mines in Transylvania, and in Hungary ; and in the United States in the Stanislaus Mine, in Calaveras County, California, in the Golden Rule Mine, Tuolumne County, California, and in a mine in Montana, the precise locality of which I have forgotten. The new locality referred to is at the Red Cloud Mine, Gold Hill District, Boulder County, Colorado.

Jan 1, 1873

By Anthony Anable

HOOVER DAM, rapidly nearing completion in the Black Canyon of the Colorado River at Boulder City, Nev., taxes the superlatives of the vocabulary to describe. For by all odds, it is the largest construction project of its type ever essayed by man and the plants providing the essential materials of construction are also in a class by themselves not only in size but in advanced design, mechanization, and accurate control. Picture, if you will, the great aggregate plant, as it is called, situated about 4 mi. upstream from the site of

Jan 1, 1934

By G. W. Imholz

The effect of new discoveries and development of other pools recently discovered is reflected in the increase in production from 586,000 bbl. during 1940 to 758,000 bbl. during 1941 in Coleman County. The States Oil Company's No. I Hudson, Coleman County, was the discovery well of the Silver field, with a total depth of 3540 ft.; initial production, 10,000,000 cu. ft. of dry gas and 60 bbl. of oil per day from the Gray sand, topped at 3482 ft. The first well drilled in the new Jim Ned field was the Merry Bros. and Perini No. I Greer, also in Coleman County. It had a total depth of 3937 ft. and an initial production of 152 bbl. of 42" gravity oil in 6 hr., through I-in. choke, from Gray sand encountered from 3909 to 3937 feet. J. E. Farrell's No. I Rose, Haskell County, was a new discovery, around which no additional wells have been drilled. Total depth was 3002 ft., plugged back to 2780 ft. to test an oil sand encountered from 2663 to 2685 ft., which when perforated with 76 shots at that depth was good for 85 bbl. of 40" gravity oil. In Fisher County, General Crude drilled its No. 12 Flanagan to a total depth of 6746 ft. After plugging back to 4578 ft. to test a sand encountered from 4515 to 4528 ft. it was abandoned as a dry hole. The top of the Ellenburger was reported as 6094 feet. Interest in Jones County continued unabated, largely through the development of the Wimberly pool in the southern part of the county, where five Cisco sands are producing. At present the pool has been defined on the south and the southwest. The structure has 170 ft. of closure; the lower limestone pays producing on top of the structure and the fist pay encountered, the Akard sand, producing on the flanks of the structure. Five deep wells were drilled in Stonewall County (H. and T.C. survey)—the Shell Oil Company's No. I-A Patterson, in block D, with a total depth of 6827 ft., dry and abandoned; the Shell Oil Company's No. I Rutherford, block I, with a total depth of 6239 ft., dry and abandoned; the Leader Oil Company's No. I Bilberry, block D, with a total depth of 5700 ft., dry and abandoned; Forest Development Company's No. 2 Boyd, block D, with a total depth of 4753 ft., dry and abandoned; and the Hunter et al. No. I Boyd, block D, total depth 4757 ft., dry and abandoned. J. E. Farrell drilled his No. I Tipton south of a well that produced from a Strawn sand for some time, to a total depth of 5202 ft., where it was abandoned. The well is in the southwest corner of sec. 43, block 19, T. and P. survey, Nolan County.

Jan 1, 1942

By G. W. Imholz

The effect of new discoveries and development of other pools recently discovered is reflected in the increase in production from 586,000 bbl. during 1940 to 758,000 bbl. during 1941 in Coleman County. The States Oil Company's No. I Hudson, Coleman County, was the discovery well of the Silver field, with a total depth of 3540 ft.; initial production, 10,000,000 cu. ft. of dry gas and 60 bbl. of oil per day from the Gray sand, topped at 3482 ft. The first well drilled in the new Jim Ned field was the Merry Bros. and Perini No. I Greer, also in Coleman County. It had a total depth of 3937 ft. and an initial production of 152 bbl. of 42" gravity oil in 6 hr., through I-in. choke, from Gray sand encountered from 3909 to 3937 feet. J. E. Farrell's No. I Rose, Haskell County, was a new discovery, around which no additional wells have been drilled. Total depth was 3002 ft., plugged back to 2780 ft. to test an oil sand encountered from 2663 to 2685 ft., which when perforated with 76 shots at that depth was good for 85 bbl. of 40" gravity oil. In Fisher County, General Crude drilled its No. 12 Flanagan to a total depth of 6746 ft. After plugging back to 4578 ft. to test a sand encountered from 4515 to 4528 ft. it was abandoned as a dry hole. The top of the Ellenburger was reported as 6094 feet. Interest in Jones County continued unabated, largely through the development of the Wimberly pool in the southern part of the county, where five Cisco sands are producing. At present the pool has been defined on the south and the southwest. The structure has 170 ft. of closure; the lower limestone pays producing on top of the structure and the fist pay encountered, the Akard sand, producing on the flanks of the structure. Five deep wells were drilled in Stonewall County (H. and T.C. survey)—the Shell Oil Company's No. I-A Patterson, in block D, with a total depth of 6827 ft., dry and abandoned; the Shell Oil Company's No. I Rutherford, block I, with a total depth of 6239 ft., dry and abandoned; the Leader Oil Company's No. I Bilberry, block D, with a total depth of 5700 ft., dry and abandoned; Forest Development Company's No. 2 Boyd, block D, with a total depth of 4753 ft., dry and abandoned; and the Hunter et al. No. I Boyd, block D, total depth 4757 ft., dry and abandoned. J. E. Farrell drilled his No. I Tipton south of a well that produced from a Strawn sand for some time, to a total depth of 5202 ft., where it was abandoned. The well is in the southwest corner of sec. 43, block 19, T. and P. survey, Nolan County.

Jan 1, 1942

MINING AND METALLURGICAL SOCIETY OF AMERICA San Francisco Section The San Francisco Section of the Mining and Metallurgical Society of America held a joint meeting with the local section of the American Institute of Mining Engineers at the Engineers' Club, San Francisco, on Tuesday, Apr. 16, 1918. The meeting was preceded by a dinner. Letters from the parent body in NewYork asking for the consideration of the proposed bill before Congress, and an expression of opinion relative to the appointment of a Federal Administrator of Mines justified the joint meeting of the local sections of the two mining societies, in that a more representative opinion of the engineers of California could he obtained. The meeting was called to order at 7:15 P.M. by Mr. Frank H. Probert acting as Chairman. About thirty engineers were present. Mr. Probert read a letter of March 14, from Secretary Huntoon, asking that the meeting be called. He read excerpts from Bulletin 117 of February 28, 1918, of the Mining and Metallurgical Society containing a copy of the proposed bill, so that all present might be advised of its essential contents. He briefly commented on the urgent need of such legislation and invited discussion of the subject. After long and thorough discussion on the part of the members present, Professor Lawson made the following motion: I move that a resolution expressing the sense of the joint meeting of the local sections of the American Institute of Mining Engineers and Mining and Metallurgical Society of America be forwarded to the secretaries of the two societies to the effect that we heartily approve of the purpose and intention of the bill designed to promote and stimulate the production of minerals necessary for the successful prosecution of the war; and that we suggest prices should be fixed for the period of the war and two years thereafter, and that these prices, guaranteed by the Government, should apply at the nearest shipping point to place of production of the minerals. This motion was carried. The meeting then adjourned, and a telegram was sent by Mr. Probert, the Chairman, to Mr. Huntoon, notifying him of the motion that had been carried at the meeting.

Jan 7, 1918

By Harry K. Ihrig

The hot-working of iron and steel is an art dating back to antiquity, but until about 25 years ago, relatively few alloying elements were used, and these were present only in small percentages. With the exception of sulphur, the amounts of other elements in steels did not markedly affect their hot-workability. The references in the literature on the effect of other elements on the hot-workability of steel are usually reports of failure of these steels in forging or rolling.1-8 With the increased use of alloying elements, and their use in larger amounts, it was observed that many of the new steels were red short in certain critical ranges or at all hot-working temperatures. The stainless steels are particularly difficult to hot-work. The determining of the correct conditions for piercing, rolling and forging of these alloys by experimentation in the mill often results in large losses of these expensive metals. It also is found that the hot-workability limits of many types vary from heat to heat, and hence poor results are obtained on changing to a new heat after a previous one has rolled satisfactorily. To determine the effect on hot-workability of iron to which various elements have been added, it was necessary to devise a method that could be used on small quantities of metal. In such a method, the minimum number of variables should be present, and the procedure should be simple, quick and correlatable with large-scale hot-working operations. Sauveur,9 in the first Henry Marion Howe lecture in 1924, described the effect of carbon on the ductility of steels in torsion near their critical points. He found that carbon increases the plasticity of gamma iron. In the 1930 Campbell Memorial lecture,10 he gives the effect of carbon, chromium, and chromium and nickel on the ductility of steels under 2000°F. in torsion. Ellis11 has studied the malleability of steels at various elevated temperatures by means of a single-blow drop-hammer test. In this test the percentage decrease in height of standard cylinders is taken as a measure of their malleability. He has shown that in the low-alloy steels studied, carbon has a greater effect than nickel, chromium and vanadium. In 1938 the first tests were run on a new quantitative hot-workability test in the laboratories of the Globe Steel Tubes CO. Thum12 mentioned it in 1940. An illustration of the apparatus was published in January 1942.'' The effect of a few alloying elements on the hot-workability of metals was shown in a description of the test in 1944.14 Description of Test Method The method adopted after a trial of several others is a rapid twisting of bars to failure in a furnace. The standard-size bar was chosen as 9/16 in. It is prepared from larger pieces by forging and cold-

Jan 1, 1947

By Harry K. Ihrig

The hot-working of iron and steel is an art dating back to antiquity, but until about 25 years ago, relatively few alloying elements were used, and these were present only in small percentages. With the exception of sulphur, the amounts of other elements in steels did not markedly affect their hot-workability. The references in the literature on the effect of other elements on the hot-workability of steel are usually reports of failure of these steels in forging or rolling.1-8 With the increased use of alloying elements, and their use in larger amounts, it was observed that many of the new steels were red short in certain critical ranges or at all hot-working temperatures. The stainless steels are particularly difficult to hot-work. The determining of the correct conditions for piercing, rolling and forging of these alloys by experimentation in the mill often results in large losses of these expensive metals. It also is found that the hot-workability limits of many types vary from heat to heat, and hence poor results are obtained on changing to a new heat after a previous one has rolled satisfactorily. To determine the effect on hot-workability of iron to which various elements have been added, it was necessary to devise a method that could be used on small quantities of metal. In such a method, the minimum number of variables should be present, and the procedure should be simple, quick and correlatable with large-scale hot-working operations. Sauveur,9 in the first Henry Marion Howe lecture in 1924, described the effect of carbon on the ductility of steels in torsion near their critical points. He found that carbon increases the plasticity of gamma iron. In the 1930 Campbell Memorial lecture,10 he gives the effect of carbon, chromium, and chromium and nickel on the ductility of steels under 2000°F. in torsion. Ellis11 has studied the malleability of steels at various elevated temperatures by means of a single-blow drop-hammer test. In this test the percentage decrease in height of standard cylinders is taken as a measure of their malleability. He has shown that in the low-alloy steels studied, carbon has a greater effect than nickel, chromium and vanadium. In 1938 the first tests were run on a new quantitative hot-workability test in the laboratories of the Globe Steel Tubes CO. Thum12 mentioned it in 1940. An illustration of the apparatus was published in January 1942.'' The effect of a few alloying elements on the hot-workability of metals was shown in a description of the test in 1944.14 Description of Test Method The method adopted after a trial of several others is a rapid twisting of bars to failure in a furnace. The standard-size bar was chosen as 9/16 in. It is prepared from larger pieces by forging and cold-

Jan 1, 1947

By D. A. Kraai, S. Floreen

The effect of 1 to 50 ppm S on the ductility of nickel at 800° to 1400°F was studied. Results at each temperature showed a decrease in the reduction of area from approximately 95 to 5 pet over the sulfur range studied. Ductility varied with grain size, but only to a minor extent relative to the sulfiw effect. The effects of sulfur were completely offset by the addition of small amounts of magnesium. The results indicate that the "hot-short" loss in ductility is not an inherent property of nickel. Some possible mechanisms which cause the loss in ductility are described. MANY metals or alloys that normally possess high ductility exhibit a ductility loss at intermediate temperatures. This loss in ductility is often called "hot-shortness". Numerous examples of this phenomenon have been reported in the literature. Much of this work has been reviewed by McLean1 and by Rhines and Wray.2 To date there is no fully satisfactory explanation of the cause of this intermediate-temperature hot-shortness. It is generally recognized that impurities, and particularly impurities that form low-melting phases, can cause embrittlement. Examples of hot-shortness have been reported, however, where there were no obvious impurities present which would lower the ductility. Thus there has been some basis for believing that hot-shortness is an inherent property, and that even the purest metal would display a hot-short loss in ductility. This latter hypothesis was recently put forward by Rhines and wray2 based on studies of nickel and nickel alloys. In the discussion of this paper, however, Guard noted that high-purity nickel showed no hot-shortness.3 Thus there is reason to doubt whether pure nickel, or by inference any other pure metal, will inherently exhibit hot-shortness. The present work was initiated to determine the extent to which hot ductility was sensitive to very small amounts of an impurity element. If it could be demonstrated that hot-shortness could be induced by only minor amounts of an impurity, then it might be argued that hot-shortness in general is an impurity effect, and not a fundamental property of pure metals. The particular impurity studied was sulfur in nickel. The deleterious effects of sulfur are well- known. It is also well-known, and will be shown below, that additions of magnesium will render sulfur innocuous. When no such refining agents are added, however, the Ni-S system is a very useful one for studying the influence of small amounts of impurities. EXPERIMENTAL PROCEDURE Two heats containing -24 ppm S were vacuum-melted and small amounts of magnesium were then added under an argon atmosphere. These alloys were used to show the effectiveness of the normal magnesium treatment in overcoming the influence of sulfur. A second series of alloys with a sulfur range of 1 to 50 ppm was then prepared by vacuum melting nickel in alumina crucibles. No elements, such as magnesium, which tend to combine with sulfur were added. The higher sulfur contents were attained by adding nickel sulfide. Lower sulfur contents were prepared using a method in which the melt was oxidized under vacuum to produce the reaction S + 2O = SO2 These heats were subsequently deoxidized with carbon. Ten- to twenty-pound ingots were cast of all of the alloys studied. The compositions are given in Table I. The ingots were forged and hot-rolled to 3/4-in. bar. They were then annealed at either 2000" or 1600°F to produce different grain sizes. One-quarter-in.-diam tensile specimens were machined from the bars. These were tested at 800°, 1000o, 1200°, and 1400°F. The specimens were held at temperature approximately 45 min before testing. The strain rates were 0.005 min-1 to yielding, and 0.05 min-' after yielding. No extensometers or gage marks were placed on the specimens because the higher sulfur heats tended to fracture at the knife-edge indentations or gage marks. The properties measured were ultimate tensile strength and reduction of area. The analytical technique for determining sulfur at low levels was that developed by Burke and Davis.4 They reported a standard deviation of 1 ppm at an average sulfur level of 4 ppm in NBS standards. A standard deviation of 3 ppm is probably more realistic for the alloys used in this investigation considering the possibility of some segregation in the ingots. RESULTS A summary of the tensile results is given in Table I. As shown in the table, both heats to which

Jan 1, 1964

By G. G. Brown

Enthalpy-entropy diagrams are presented for natural gases of 0.6, 0.7, 0.8, 0.9, and 1.0 gravity over the pressure range of 5 to 10,000 Ib. per sq. in. and temperature range of 32º to 700°F. The charts indicate directly the work requirement and temperature rise for adiabatic compression or temperature change for free emansion of natural eases. Computation and Uses of Charts The Mollier diagram (Fig. 4), in which the enthalpy (heat content) is plotted against the entropy with lines of constant temperature, pressure,, and in some cases volume, has been found most convenient when dealing with the compression, expansion, and flow of fluids. In dealing with the flow of fluids, the sum of the increase in the enthalpy, AH, plus the increase in kinetic energy, AM —, plus the increase in potential energy, ?MZ, representing the total increase in energy of the fluid in flow, is equal to the sum of the heat q, and work added, —w, to the fluid while flowing between the entrance and exit of the flow system. AH + AM— + ?MZ = q-w [1] In cases where there is no significant change in potential energy or in kinetic energy (velocity), it follows that the increase in enthalpy is equal to the total energy supplied to the fluid. Under such conditions' the changes in the property of the fluid as it flows through a throttling valve, choke, or any other similar arrangement, may be read directly from the enthalpy-entropy diagram by following a horizolltal line between the known pressures' When compressing Or expanding a gas by means of a Or or engine in which no heat is added to or subtracted from the gas, but only work done, the changes in the properties of the gas may be determined along a vertical line of constant entropy between the entering and exit pressures. The power required for the compression of the gas may be readily determined by converting the increase in enthalpy into the desired units. The enthalpy-entropy diagram for natural gases is to the gas engineer what the steam diagram is to the steam-power engineer. For this reason it would be extremely convenient if a reasonably satisfactory enthalpy-entropy diagram could be prepared as a function of the gravity of the gas. A careful study of the known properties of natural gas indicate that this is possible. The effect of temperature upon the enthalpy at constant pressure is expressed as the "heat capacity" or "specific heat" of the gas. The best available data for natural and petroleum refinery gases indicates the relationship shown in Fig. I. From this it is clear that the specific heat of natural gases is a function only of the gas gravity and the temperature at atmospheric pressure. The effect of pressure on the enthalpy of natural gases is dependent upon the pressure-volume-temperature relationships,

Jan 1, 1945

By W. H., Bassett

IN January, 1926, the writer began a study of the commercial value of beryllium in its relation to copper. The purpose of the investigation was not to make a mere laboratory study of the characteristics of beryllium-copper alloys, but to determine under mill conditions the behavior of copper when alloyed with beryllium and when other metals were added to such alloys; in other words, it was considered important to know whether the alloy bars could be hot-rolled or cold-worked, and what limits in composition would have to be established for these two methods of working. Consequently, more stress has been placed on the practical application of the alloys in engineering and commercial use than on a discussion of the metallography of the specimens examined, although this part of the matter has not been entirely neglected. Credit should be first given to G. Oesterheld,1 who, in 1916, published the copper-beryllium equilibrium diagram. Fig. 1 has been copied from the original article with the simple substitution of English words. It will be noted from Fig. 1 that alpha solid solutions are formed, up to about 1.5 per cent. Be (10.5 atom.). The freezing-point curve, after falling from 4.2 per cent. Be (23.5 atom.), enters into a very complicated form over the entire beta region. There is a possible compound, Cu2Be, at about 6.5 per cent. Be (32.3 atom.). The pure gamma range extends from 11.2 to 12.2 per cent. Be. (46.5 to 49 atom.), after which the liquidus rises rapidly to 30 per cent. Be. (75 atom.). The delta solution extends from 25 per cent. to an unknown percentage of beryllium. There is a eutectoid point at 575" C., extending from 1.5 to 11.2 per cent. Be., the maximum time interval of which occurs at 6 per cent. Be. Below this line the beta solution is described as breaking up into alpha and gamma. Microscopically, the eutectiferous alloy thus formed closely resembles pearlite, and requires a high magnification for its resolution.

Jan 1, 1927

By W. F. Carew, F. A. Crossley

IT has been reported that the Ti-8 pct A1 alloy is ductile as water quenched from 800°C but brittle as annealed at 650 °C." The present, somewhat limited, investigation was undertaken to discover the cause of the embrittlement. In order to evaluate to some extent the continuity of the embrittling phenomenon, alloys of 6, 8, and 10 pct A1 were studied. Subsequent to the completion of most of this work, Ence and Margolin reported that a Ti,A1 phase exists and that two or more unidentified compounds may exist in the Ti-A1 system.' Materials, Equipment, and Techniques The alloys were prepared using Kroll process sponge of 124 Bhn hardness. In addition, some 8 pct A1 material was made with iodide titanium. The sponge-base alloys were prepared as melts of 4 lb or more by double arc melting. The iodide Ti-8 pct A1 alloy was prepared as 200-g buttons. The alloys were forged at temperatures of 980" to 1150°C. Results of chemical analyses are given in Table I. A Baldwin-Southwark 60,000-lb capacity tensile machine, equipped with an autographic stress-strain recorder, was used for the tensile tests. Standard tensile test specimens, 0.252 in. diam, were used. Standard Charwy V-notch impact specimens were tested using a -kiehle machine. specimens were sealed in evacuated Vycor bulbs for heat treatment. Powder samples for Debye-Sherrer anal- ysis were prepared by filing and screening through 325 mesh. The —325 mesh fraction was taken in an effort to concentrate a microstructural constituent assumed to be less ductile than the a phase. The powder samples were stress relieve2 by annealing

Jan 1, 1958

By S. Jana, C. M. Wayman

The reverse transformation of bcc martensite to the fcc phase was studied in an Fe-33.95 wl pct Ni alloy by nzeans oj dilatometry, melallography, and electron microscopy. Upon "slozc" heating (-1°C per min) length cJmnge us temperature plots showed u gradual contracLion over the temperature range 200" to 280"C ,followed by a more abrupt contraction beginning a1 -280°C. Howet,ev, zchen the heating rate was increased -4°C per tnin, no gradual contraction was observed and only the abrupt contraction starting at -2BO"C was found. Thus on slower heating- the AS "temperature" for the subject alloy, unlike the MS temperature, is better defined as a range of temperatures. Both optical and transmissiorl electron microscope observations showed that some of the martensite plates exizibited a partial loss of transformation twins during reversal. The midvib region of the martensite plates disappeaved relatively early duirng the reversal. Metallographic observations slowed that the earliest detectable stage of the rezlerse tvansforrvration begins (axd Moues inulardly) at The Martensens i te - parent interface. At higher temperatirres, the. formation of martensitically reversed jcc plates within the bcc martensite plales was observed. It is concluded that the reverse transformation consists of a diffusion less process (martensitic); but this is ps-obably aided by a prior or simultaneous dijjusiorz-comltvolled process, at leasl in the case of slower heat-ing' experiments. ALTHOUGH numerous investigations have dealt with the parent-to-martensite ("forward") transformation (fcc — bcc) in Fe-Ni alloys, comparatively little is reported on the ("reverse7') martensite-to-parent transformation.'-4 Even though such reverse transformations have been studied in detail in some nonferrous systems, one of the difficulties of studying the reverse transformation in most ferrous mar-tensites is that the martensite decomposes by tempering during heating. However, carbonless Fe-Ni alloys do not exhibit this difficulty since the transformation in these alloys is completely reversible. The present investigation represents an attempt to shed more light on the nature and mechanism of the martensite-to-parent transformation. 1) EXPERIMENTAL PROCEDURE 1.1) Alloy Prepatation. Fe-Ni alloys of compositions near 34 wt pct Ni were prepared from zone-refined iron (99.994 wt pct Fe) and high-purity nickel (99.999 wt pct Ni) by induction melting in recrystallized alumina crucibles in an argon atmosphere, with prior vacuum evacuation to 10"3 mm Hg. The alloys were homogenized by induction stirring in the molten state for 5 min. After solidification, the alloys were further homogenized in evacuated quartz capsules for 96 hr at 1230°C. 1.2) Dilatometry. Slices of the ingot were hot-forged (750°C in air) into approximate rod form and these specimens were then hot-swaged (750°C in air) into long cylindrical rods 0.55 mm diam. From the rods, specimens about 1 in. long were cut. These were then vacuum-annealed for 24 hr at 1200°C, cooled to room temperature, and subsequently transformed to martensite in liquid nitrogen (whereby about 40 pct transformation was obtained). Dilatation measurements were made by observing length changes in a vacuum dilatometer with an externally mounted LVDT sensing element. 1. 3) Preparation of Electron Microscope Specimens. Slices of the ingots were cold-rolled (with intermediate vacuum anneals) to -0.020 in. Out of these rolled sheets, specimens (about 1 by 1 in.) were cut. These were then vacuum-annealed, transformed to martensite by cooling in liquid nitrogen, and subsequently heated from room temperature to various temperatures to effect either partial or complete reverse transformation. These specimens were then chemically polished to 0.002 in. in l:l HsOz (30 pct) and &PO4 (85 pct) solution, and thinned to electron transparency in an electrolyte consisting of 150 g CraOs, 750 ml glacial acetic acid, and 30 ml ~~0.~ Observations were made with a 100-kv Hitachi HU-11 electron microscope equipped with an HK-2A tilting device. 1.4) Optical Microscopy. Metallographic observations were made with a Leitz MM5 metallograph on the same 0.020-in. sheet specimens as were used for electron microscopy and on bulk specimens which were 0.2 in. or more on a side. The chemical thinning solution when cooled below 20°C also served as an etchant for this alloy. Observations of surface relief were made with a Zeiss interference microscope employing a Thallium light source of wavelength 0.54 p. Specimens for interference studies were prepared by two-stage polishing on Buehler vibromet polishers using 0.3 and 0.05 p alumina abrasives. 2) EXPERIMENTAL RESULTS 2.1) Comparison of the MS,AS, and Af Tempera-tures wTth Previous Re sults. The AS aLd Af tempera -tures of several Fe-Ni alloys were determined dila-tometrically. The MS temperatures of the same alloys were determined by continuously lowering the temperature using a mixture of isopentane and liquid nitrogen and observing the highest temperature at which a prepolished specimen showed surface upheavals. For the present the As temperature is defined as the temperature at which an abrupt decrease in length occurs in the dilatation plot. The Ms,As7 and A determinations in the present investigation and those of Kaufman

Jan 1, 1968

By James J. Bean

Although determining and controlling specific gravity of operating medium in a heavy-media plant manually presents no problem, there are advantages to automatic recording and control. The two installations discussed employ the usual method of measurement. Pneumatic control with band control and reset features is used to regulate dilution water and hence specific gravity. Difficulties have been mechanical in nature and associated with obtaining a proper I sample and maintaining tight air connections in bubble-tube piping:

Jan 1, 1950

By C. W. Hill, G. P. Luckey

In a previous article1 preliminary experiments were described, indicating the possibilities of a quantitative spectroscopic method for the determination of small amounts of lead in copper, which would be accurate and rapid, and could be carried out in the refining plant by one not skilled in chemical analysis. The present paper deals with the development of the method in the factory, giving the details of apparatus and its standardization, and presenting a comparison of the accuracy of the method with that of the standard electrolytic determination. (For variations and other applications of the method, the reader is referred to the first article.) A carefully weighed sample (0.4 gm.) of the copper to be tested is placed in a cavity in a graphite positive electrode. An arc capable of

Jan 1, 1919

By J. A. Retty

IN the Labrador iron fields two concessions, totaling nearly 24,000 square miles, have been staked out and commercial-grade deposits delineated. The Newfoundland-Labrador concession, owned by the Labrador Mining and Exploration Co. Ltd., is a 20,000-square-mile tract in the upper basins of the Hamilton and N askaupi Rivers which drain into Lake Melville. The New Quebec licensed area is 3900 square miles, owned by the Hollinger North Shore Exploration Co. Ltd., lying to the northwest of the northwestern part of the Labrador concession. It includes the upper basin of the Swampy Bay River and part of the upper basin of the Whale River, both of which drain into Ungava Bay. Both of the above-named companies are controlled by Hollinger Consolidated Gold Mines Ltd. with The M. A. Hanna Co., of Cleveland, holding a minority interest. When exploration was first undertaken in 1936, the only published map of the area was that prepared by A. P. Low in 1895. It showed only the main watercourses and gave only a rough, general idea of the geology. The numerous lakes of this vast area had to be first sketched and then mapped. Since then, extensive geological and topographical surveys have been conducted by the Labrador company, work which is normally done by Government branches. The object of the work was to examine the area and gradually eliminate unfavorable or less favorable prospecting ground. With this object in view, geological

Jan 1, 1948

By G. N. Goldberg, Robert A. Rapp

The steady-state kinetics and microstructures of simultaneous internal oxidation and external scale formation were investigated for the oxidation of Cb-Zr and Ch-Zr-Re alloys in pure oxygen at 1000°C. For each binary alloy, a constant rate of scale formation was observed at steady state; the internal oxidatton zone reached a steady-state, time-independent thickness. The dependence of the scaling rate on zirconium content exhibited a maximum at a low zirconium colttent; more concentrated alloys oxidized at a slower rate than pure columbium. This composition dependence of the scaling rate may be partially attributed to the effect of the internal oxidation process as an intenla1 sink for oxygen, with the formation of voluminous, and relatively impermeable ZrO2 precipitates. However, the morphologies of the ZrO2 internal oxide precipitates also affected the scaling rate. The internal oxide precipitates were distributed uniformly in the external scale. The steady-state thicknesses of the internal oxidation zones were not in agreement with those predicted theoretically from an idealized simple model. However, in this particular alloy systen1 the ideal model is not satisfied experitmentally. For ternary Cb- Zr - Re alloys with Nr10e = 0.02 or 0.05, the steady-state thicknesses of the internal oxidation zones were less than those for the corresponding binary Cb-Zr alloys. For ternary alloys with NRe(0) - 0.05, a more adherent and much less porous external scale was formed, and a reduction in the kinetics of scale formation was observed. A number of authors'-7 have reported linear oxidation kinetics for the reaction of pure columbium at 1000°C in air or in oxygen of about 1 atm pressure. A thin and tightly adherent oxide layer is found at the metal/oxide interface beneath a porous external layer of scale. Although the suboxides CbO and CbO2 represent thermodynamically stable phases at 100O°C, generally only Cb2O5* is identified as a reaction product in quenched specimens. Studies at 1000°C at lower oxygen pressures8,10 (Po2 - 10-4 Torr) show that both CbO and CbO2 do form and that CbO2 is a protective oxide. The growth of CbO2 results in parabolic kinetics until Cb2O5 is nucleated and grows. Apparently because of the large volume change associated with its formation, Cb2O5 is porous and does not serve as a diffusion barrier; linear kinetics prevail after the surface is covered with Cb2o5. On the basis of these observations, several authors5,9,10 have suggested that the linear oxidation of pure columbium at 1000°C in air or oxygen of 1 atm is limited by a Loriers-type mechanism,11 whereby the total rate or reaction is controlled by ionic diffusion through a thin layer of CbO2 which is maintained at a constant thickness throughout the linear oxidation. However, the concept of a Loriers model for a three-phase CbOlCbO2 Cb2O5 scale, with the only oxygen activity gradient across the NbO2 phase, is not self-consistent. Further, the suboxides CbO and CbO2 were not observed in the high-temperature X-ray diffraction study of columbium oxidation by Goldschmidt.8 From investigations in which the large PO2, dependence of the linear scaling rate was demonstrated, the importance of an oxygen adsorption or dissolution step has been suggested.4,5,7,10,12 Thus the mechanism for the linear oxidation of pure columbium remains quite controversial. Since further insight into the oxidation mechanism of pure columbium is not provided by this investigation, the authors wish to emphasize at the outset that their experimental results are not interpreted in terms of a particular rate controlling step for the oxidation of pure columbium. Previous investigations13-18 of the oxidation of Cb-Zr solid-solution alloys in air or oxygen at 1000°C suggest a remarkable dependence upon zirconium content. For an alloy of bulk zirconium mole fraction, N2r(0) equal to 0.05 or 0.10, the oxygen uptake is reported to be as high as four times greater than that for pure columbium;13,14 for alloys with NZr(0) 0.05 or 0.10 the oxygen uptake is reported to decrease with increasing NZr(0); to a minimum uptake of about one quarter that for pure columbium at NZr(0): = 0.50.13-15Since zirconium forms a more stable oxide (ZrO2) than the lowest columbium oxide (CbO), and since columbium exhibits a high solubility19,20 and diffusivity21,23 for oxygen, the internal oxidation of the zirconium component to ZrO2 is expected at a reaction front ahead of the advancing metal/scale interface. The external scale is then formed at the metal/scale interface by the inward migration of oxygen through the scale, probably through a series combination of molecular and ionic diffusion. Internal oxidation in conjunction with external scale formation has been investigated by Maak24 for the oxidation of dilute Cu-Be alloys in pure oxygen at 850°C. For cu-Be and many other binary alloys,24-27 very small, essentially uniaxial internal oxide particles are formed in the most dilute compositions; for somewhat more concentrated alloys. the internal oxide particles precipitate as platelets or needles. From pertinent solutions to the diffusion equation, Maak28

Jan 1, 1967

By L. D. Jaffe

THE data recently published1-" on heat treatment of titanium alloys have not been readily systematized into an overall picture. When the ß titanium phase transforms isothermally at low temperatures, the hardness follows a typical age-hardening pattern, going through a peak as schematized by the "hardness" vs "time of isothermal transformation" coordinates of Fig. 1. As the temperature is raised, the peak will decrease in height and occur at shorter times; slowing must however take place as the /3 transus is approached. The brittle microconstituent in the region of peak hardness has been termed /3'. Its structure is still undetermined but it is clearly a transition stage from soft /3 to the £ + a that is found at the right of the peak. On continuous cooling, high cooling rates correspond to short times on Fig. 1; abscissa labeling has been added accordingly. Extremely rapid cooling will retain soft /3; very slow cooling will precipitate soft a; intermediate rates give hard 8'. Increasing distance from the water-quenched end of a Jominy hardenability bar corresponds to slower cooling (right on Fig. I), but the range of cooling rates covered by a Jominy bar is only a small portion of that shown in the figure. The p-stabilizing alloying elements retard the transformation process, moving the material to the left on Fig. 1. A lean alloy is harder water-quenched (P') than air-cooled (/3 +a); a rich alloy is harder air-cooled (/3') than water-quenched (/3). Examples of the position of various compositions with respect to the curve, based on limited data in the literature, are given at the top of the figure. Oxygen and carbon accelerate the transformation process, moving the material to the right on Fig. 1. If an alloy quenched quickly enough to retain soft /3 is reheated below the £ transus (tempered), it ages, going through a hardness maximum as time increases (moving right on Fig. 1). Material which after cooling from the /3 range lies on or right of the hardness peak, will move still further right on tempering, and so soften. Since the rate of the tempering (aging) reaction increases with temperature, the figure also may be used to represent the variation in properties with tempering temperature, at short tempering times. The simple picture given here neglects martensite transformation, as well as differences between decomposition products of /3 formed by direct isothermal transformation and those formed by quenching and tempering. It may nevertheless be of use as a guide for practical heat treatment and for further research.

Jan 1, 1955

By W. P. Barba

IN THE early summer of 1917, it became evident that the then existing sources of supply of guns and gun forgings were totally inadequate for the enormous and rapidly growing requirements of the Ordnance Department of the Army, so arrangements were made with sixteen makers of other products to take up the manufacture of guns and gun forgings. Nine of these makers were required to produce their own steel from the raw materials and convert it into rough machined gun forgings; the others finished and assembled these forgings into completed guns. In December, 1914, four of these makers of steel and rough gun forgings were instructed to manufacture the same amount of the 6-in. and 9.5-in. guns and howitzers. It, therefore, seemed practical .and profitable to assemble these makers into a small group, together with the officers of the Ordnance Department, and provide standard 'equipment in each of the four works for the production of these two types of guns. The advantages of this cooperation were immediately manifested, and there commenced an interchange of shop practice and a free discussion of facilities and methods that covered not only the equipment but the manufacturing processes.

Jan 1, 1920

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

Alpha solutes increase the strengths of Ti-Mn alloys through solid-solution strengthening. The substitutional a addition, aluminum, decreases, and the interstitial solutes, carbon and nitrogen, increase the rate of nucleation and growth of a from ß. The best combinations of properties of a-ß alloys are obtained when there is a sufficient quantity of a phase in the structure to dissolve the a solutes. OF the many different titanium-base alloy systems, the predominant alloy type is the a-ß alloy. The properties of the a-ß alloys are dependent on solid-solution strengthening and heat-treatment effects involving the a-ß ratio and transformation reactions. Another variable which influences the mechanical properties of a-ß alloys is the a-stabilizer content of the alloy. An a solute may be present as an intentional addition, such as aluminum, or as an impurity element, such as carbon, oxygen, or nitrogen. It is known that these a stabilizers, when added to titanium, form single-phase alloys which are not heat treatable but which obtain their strength from solid-solution strengthening. Thus, it would be expected that a additions to a-ß alloys would increase the strength of the alloys by solid-solution strengthening of the a phase. In addition, they would affect the transformation kinetics of the ß-to-a reactions and other processes based on the instability of the ß phase. The effects of heat treatment and structure on the mechanical properties of Ti-Mn alloys have been shown in a previous paper.6 This system offers a good base to demonstrate the effects of typical a solutes on the properties of a-ß alloys. The three a solutes described in this work are aluminum, representative of a substitutional a solute, nitrogen, representative of an interstitial a solute, and carbon, representative of an interstitial compound-forming element. The effects of heat treatment and microstructure on the properties of a alloys containing these three elements are described in concurrent publications. Some of these data are used for base-line points in several of the curves used for illustration herein. Experimental Procedures Iodide titanium was used as the base for all of the alloys studied in this work. The alloys were prepared as ½ lb ingots by double arc melting in an argon atmosphere. The ingots were forged to ¾ in. rounds, vacuum annealed for 6 hr at 900°C at a pressure of 10 ' to 10-5 mm of Hg to remove hydrogen, and hot swaged to 1/4 in. diam rod. After me- chanical descaling, test specimens were prepared for heat treatment. The alloys used in this study together with the fabrication temperatures are given in Table I. Heat treatments were done in argon. For the most part, the specimens were sealed in Vyeor capsules under a partial pressure of argon. Quenching was accomplished by breaking the capsule under water. Other cooling methods used included oil quench, argon cool (simulated air cool in an argon atmosphere), and furnace cool. The times for the various heat-treating temperatures are given in Table 11. The tests performed on the alloys consisted of tensile tests on ? in. diam specimens, hardness tests, and microimpact tests. Specimen sizes have been adequately described in a previous publication.' The micrographs presented in this paper were taken from specimens cut from the shoulders of broken tensile specimens. Final polishing was done with Linde B on a slow-speed wheel, and the specimens were etched with a 1½ HF — 3½ HNO, solution. Ti-N-Mn Alloys The transformation diagram and microstructures of the Ti-0.1 pct N-Mn alloys used in this investigation are given in Fig. 1. The effect of small nitrogen additions on the binary Ti-Mn diagram is to raise the ß-transus temperature with little effect on the a solubility of manganese. Also, as has been noted previously,' high manganese-content alloys containing nitrogen, when quenched from temperatures high in the ß field, contain a subgrain boundary phase which appears to be nitrogen-rich a. Marten-site is formed when alloys containing less than about

Jan 1, 1956