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By William Hume-Rothery

I need not say how much I appreciate the honor of being asked to lecture to you, and how much I would thank you for your kind invitation. It is encouraging to feel that the abnormal restrictions of the war have been relaxed sufficiently to allow a free exchange of scientists between our two countries, and I hope that some of you will be persuaded to go to England, and to visit us at Oxford, so that I may repay a little of the kindness I have received from you. As this is one of the first lectures since the end of the war, I have thought it well to describe some of the advances in the theory of alloys which were made in the years between the two World Wars. From the point of view of theoretical science, these years have been of outstanding interest in the history of metallurgy. They have been years in which the development of new and more accurate experimental methods has led to the revelation of many new facts. At the same time, the theoretical developments have enabled some of these facts to be viewed as parts of a general scheme, and we may say that some of the foundations of metallurgical science have been laid. This work was naturally stopped by the Second World War, and it can therefore be reviewed as one definite stage in the development of our Science. To illustrate the new advances, I shall deal mainly with the alloys of copper and silver, since these have revealed many of the general principles. taken, and copper had been found to crystallize in the face-centered cubic structure (Fig I). The determination of this structure naturally was of immense importance in the understanding of our subject, because it meant that we were at last in a position to appreciate the underlying structure out of which the crystal of copper was built. The earlier science of metallography had shown that in many alloys of copper the addition of a second metal in not too large proportions resulted in the formation of a homogeneous alloy which was regarded as a solid solution of the one metal in copper. Fig 2, for example, shows the so-called equilibrium diagram of the copper-zinc alloys, and here you will see that there is a solid solution of zinc in copper extending to something of the order 35 to 40 at. pct zinc.

Jan 1, 1947

By William Hume-Rothery

I need not say how much I appreciate the honor of being asked to lecture to you, and how much I would thank you for your kind invitation. It is encouraging to feel that the abnormal restrictions of the war have been relaxed sufficiently to allow a free exchange of scientists between our two countries, and I hope that some of you will be persuaded to go to England, and to visit us at Oxford, so that I may repay a little of the kindness I have received from you. As this is one of the first lectures since the end of the war, I have thought it well to describe some of the advances in the theory of alloys which were made in the years between the two World Wars. From the point of view of theoretical science, these years have been of outstanding interest in the history of metallurgy. They have been years in which the development of new and more accurate experimental methods has led to the revelation of many new facts. At the same time, the theoretical developments have enabled some of these facts to be viewed as parts of a general scheme, and we may say that some of the foundations of metallurgical science have been laid. This work was naturally stopped by the Second World War, and it can therefore be reviewed as one definite stage in the development of our Science. To illustrate the new advances, I shall deal mainly with the alloys of copper and silver, since these have revealed many of the general principles. taken, and copper had been found to crystallize in the face-centered cubic structure (Fig I). The determination of this structure naturally was of immense importance in the understanding of our subject, because it meant that we were at last in a position to appreciate the underlying structure out of which the crystal of copper was built. The earlier science of metallography had shown that in many alloys of copper the addition of a second metal in not too large proportions resulted in the formation of a homogeneous alloy which was regarded as a solid solution of the one metal in copper. Fig 2, for example, shows the so-called equilibrium diagram of the copper-zinc alloys, and here you will see that there is a solid solution of zinc in copper extending to something of the order 35 to 40 at. pct zinc.

Jan 1, 1947

By R. A. Stauffer

The chief ore of lithium is spodumene, a lithium-aluminum silicate containing up to 3 pct lithium. The preparation of lithium salts from spodumene is costly because of the low concentration of the metal and the difficulty of separating it from the other alkali metals present. Dr. Kroll's and Dr. Schlecten's recent paper1 has reviewed various methods proposed for preparation of lithium. It will be seen from this paper that the methods proposed to date have required one or another of the lithium salts as a raw material. The high cost of these salts has been the chief item of expense in the lithium metal produced. This situation has made it interesting to investigate the possibility of producing the metal directly from spodumene. When spodumene is calcined at 1000ºC in the presence of lime, lithium oxide is replaced in the ore as is evidenced by the fact that it may be leached from the reaction mixture with water.2 Free energy calculations suggested that lithium oxide might be reduced with aluminum or silicon and the lithium produced distilled from the reaction mixture by heating-to about 1000ºC in a vacuum. These facts taken together suggested that it might be possible to prepare lithium metal directly from spodumene by heating a mixture of the ore, lime, and a suitable reducing agent in vacuum and distilling off the lithium metal as formed. The Replacement of Li2O in Spodumene A series of experiments were undertaken to determine the ratio of lime to spodumene necessary to free substantially all the lithium oxide Present in the Ore. The spodumene used in these experiments was supplied by Foote Mineral Co. in the form of powder of about 200 mesh' Analysis of the spodumene was as follows: PCT SiO2......................... 62.7 Al2O3........................ 29.9 Li2o......................... 4.9 Na2O........................ 0.89 CaO......................... 0.75 MgO........................ 0.49 K2O......................... 0.44 Lime (cp —200 mesh) was added in various ratios to spodumene and thoroughly mixed in a ball mill. Two hundred gram samples of spodumene-lime mixtures were charged loosely in fireclay crucibles and held at 1000ºC in a gas-fired furnace for 5 hrs. The product in each case was ground to pass 140 mesh and then placed in a thimble in a soxhlet extractor. Extraction with water was carried on for 6 hrs. After this period the extract was reduced in volume by evaporation in order to precipitate dissolved lime and was then filtered and titrated with 1/4 normal hydrochloric acid. After correcting for residual calcium oxide in solution, the alkali present was calculated as lithium oxide. The results of these experiments are shown in Table I and Fig I. At lime-spodumene ratios less than 1.5, the products were invariably fused. At the ratio of 1.5 and above, the product was loosely sintered with little evidence of fusion. When the ratio of lime to spodumene was three, replacement of alkali oxides in the

Jan 1, 1949

By C. S. Barrett, O. R. Trautz

Previous investigations have shown that lithium is body-centered cubic from near its melting point to the temperature of liquid air.1,2,3 Nevertheless there was an incentive to search again for a transformation at low temperatures: C. Zener4.5 has remarked that the low value of the elastic modulus (C11 — Cl2)/z of body-centered cubic metals and alloys implies a rapid increase in free energy with decreasing temperature, which would favor the occurrence of a transformation; also the shear movement to which this modulus applies is one that can transform a body-centered cubic structure into a slightly distorted face-centered cubic structure. It was confirmed that lithium does not transform spontaneously at liquid nitrogen temperature, but it was found that a transformation can be induced at this temperature by cold working the metal,6 and at slightly lower temperatures a different transformation occurs spontaneously. Transformations were also found in solid solutions of magnesium in lithium. The Nature of the Transformations Before presenting a detailed treatment of the technique used and the specific data obtained, it would be well to give a brief survey of the general characteristics of the transformations. A new crystal structure begins to form spontaneously in a specimen when it is cooled below a critical temperature. In lithium the structure appears to be close-packed hexagonal; in theLi-Mg alloys most of the diffraction lines (but not all) are near those for hexagonal lithium. The transformation proceeds as the temperature is lowered and soon stops if the temperature is held constant or raised. The course of the transformation on cooling is indicated by the curve at the lower left of Fig -I. This curve is drawn in stepped fashion because the transformation goes by discrete steps with audible clicks. Because of the similarity to the decomposition of austenite into martensite in steels, it seems appropriate to designate the beginning of the spontaneous transformation on cooling as the Ms point. On heating to room temperature after the low temperature treatment only the usual BCC structure is found. On heating, the low temperature modification begins to disappear at a temperature which will be designated as Ac. The transformation on heating is halted when the rise in temperature is halted. The temperature must be raised 40 to 50°C above Ac before the reversion is complete—an interval which is sensitive to the rate of heating and to the time of holding at temperatures above Ac. This behavior on heating thus has similarities also to the formation of martensite in the cooling of steels.

Jan 1, 1949

By C. S. Barrett, C. T. Haller

Twinning in magnesium is known to occur profusely under certain conditions, and when it occurs in polycrystalline materials it brings about a partial or even a complete change in the preferred orientation. The amount of research on this subject, however, has not been in proportion to its importance as a mechanism of deformation and reorientation. Much of the research has been limited to twinning in single crystals. The present study1 concerns polycrystalline sheet, and is an investigation of the twinning produced by plastic deformation as a function of the following variables: 1. Direction and type of deformation. 2. Amount of deformation. 3. Speed of deformation. 4. Temperature of deformation 5. Composition and temper of the sheet. 6. Twins previously produced. 7. Twins previously produced, then annealed. By proper control of these variables it is shown that, on a laboratory scale, samples can be produced that are free from preferred orientations. The materials used were samples of commercial magnesium-alloy sheets with nominal compositions given in Table I. X-ray diffraction was used in most of this work to detect the twins and to estimate the amount of twinning. Since twinning on (102) planes causes a re-orientation of the crystal lattice by 86", it is very easy to study with X-ray diffraction patterns, particularly if one notices the diffraction ring from the basal plane— the 002 reflection. By properly directing the X-ray beam, it is possible to produce 002 reflections from twinned material at, say, the 6 and 12 o'clock positions on the 002 ring, with the untwinned material reflecting to 3 and 9 o'clock positions; and by judging the relative intensity of the arcs on the diffraction ring at these positions, it is possible to judge the amount of twinned material relative to the untwinned. This X-ray method of estimating the amount of twinning is generally superior to metallo-graphic observation, for it was found

Jan 1, 1947

By C. S. Barrett, C. T. Haller

Twinning in magnesium is known to occur profusely under certain conditions, and when it occurs in polycrystalline materials it brings about a partial or even a complete change in the preferred orientation. The amount of research on this subject, however, has not been in proportion to its importance as a mechanism of deformation and reorientation. Much of the research has been limited to twinning in single crystals. The present study1 concerns polycrystalline sheet, and is an investigation of the twinning produced by plastic deformation as a function of the following variables: 1. Direction and type of deformation. 2. Amount of deformation. 3. Speed of deformation. 4. Temperature of deformation 5. Composition and temper of the sheet. 6. Twins previously produced. 7. Twins previously produced, then annealed. By proper control of these variables it is shown that, on a laboratory scale, samples can be produced that are free from preferred orientations. The materials used were samples of commercial magnesium-alloy sheets with nominal compositions given in Table I. X-ray diffraction was used in most of this work to detect the twins and to estimate the amount of twinning. Since twinning on (102) planes causes a re-orientation of the crystal lattice by 86", it is very easy to study with X-ray diffraction patterns, particularly if one notices the diffraction ring from the basal plane— the 002 reflection. By properly directing the X-ray beam, it is possible to produce 002 reflections from twinned material at, say, the 6 and 12 o'clock positions on the 002 ring, with the untwinned material reflecting to 3 and 9 o'clock positions; and by judging the relative intensity of the arcs on the diffraction ring at these positions, it is possible to judge the amount of twinned material relative to the untwinned. This X-ray method of estimating the amount of twinning is generally superior to metallo-graphic observation, for it was found

Jan 1, 1947

By J. B. Hess, R. L. Dietrich

In the mechanical twinning of magnesium on the {1012} planes the crystal-lographic deformation is such that, in the direction of the hexagonal axis [0001], twinning is possible only undef tension stress, while in the directions perpendicular to the hexagonal axis, twinning is possible only under applied compression. Thus in extruded magnesium and magnesium alloys where the extruded textures are best described as either [1010] or [1120] fibers,' the basal planes are all nearly parallel to the extrusion axis and such material is susceptible to ( 1or2 ) twinning by compression in the direction of the extrusion axis. In order to explain the results of cyclic stressing in alternate tension and compression upon the tangent elastic modulus in such material, Dorn and Thomsen2 have postulated that, following twinning under applied compression, a retwinning or un-twinning occurs as a result of residual microtensile stresses immediately upon removal of the applied compressive force. Recently, Carapella and Shaw4 have used this supposition of untwinning in magnesium in discussing their results on indentation hardness and on cold drawing of MI sheet. In trying to confirm this hypothesis by suitable direct experiments, the authors have been unsuccessful in detecting any evidence whatsoever that untwinning follows the release of applied compression and are forced to conclude that this hypothesis is improbable or, at least, that the quantity of material involved is small. For the experiments, small cylindrical specimens 0.040-in. diam. X 0.120- to 0.140-in. long were etched from 0.0625-in. extruded rod. The ends of the specimens were accurately faced off parallel in a jeweler's lathe. The compression jig was simply an ordinary set of micrometer calipers with an anvil interposed between the spindle and the specimen. The anvil was free to move axially but restrained from rotating so that the rotation of the spindle would not introduce a twisting stress on the specimen. With this arrangement the compressive strains could be estimated to 0.0001 in. and X ray pinhole photographs could be taken with or without load as required. Fig I is the pattern of an uncompressed specimen of AZ3I alloy, showing the approximate [1010] extrusion fiber. The X ray beam is normal to the extrusion axis; in this position it is particularly sensitive in detecting twinning if one directs his attention to the diffraction ring from the basal plane. Since twinning on (1072 planes produced an 86" reorientation of the crystal lattice, the original basal reflections occur at the 3 and 9 o'clock positions while the (0002) reflections from any twinned material will occur at the 6 and 12 o'clock positions. Fig 2 shows the same specimen compressed 3.0 pct, photographed while still under load. This strain has been just enough to produce a detectible amount of twinning and the sensitivity is roughly indicated by the fact that no alterations were produced in the intensities or maxima of the other diffraction rings. A strain of

Jan 1, 1949

By M. C. Udy, P. C. Rosenthal

The use of minute boron additions to steel has been given considerable attention in recent years. Comparisons made between boron-free and boron-containing heats of otherwise identical analysis have indicated, beyond a doubt, the potent effect of boron for increasing hardenability. The actual hardenability increase has been expressed as a multiplying factor, or in terms of the percentage of other alloying elements boron might replace. The present study presents the results of a systematic survey of the use of boron as an addition to two types of alloyed steels and as a replacement for various percentages of molybdenum in these steels. Both cast and wrought steels of the two basic compositions were prepared. Molybdenum, because of its scarcity at the time the investigation was started, was selected as the element to be replaced by boron. Attention was also given to the effects of boron on other properties of the steels, particularly their notched-bar toughness, to investigate other advantages or limitations of this element as an alloy addition. Experimental Work Steels Investigated The two basic compositions used in this survey of boron as a replacement for molybdenum are given in Table I. The cast steels were higher than the wrought steels in silicon content to improve the casting quality. Table i.—Nominal Composition of the Two Base Alloy Steels composition. Per Cent. steel Mn Si Ni Cr Xi-Cr-Mo Wrought 0.28-0.70-0.15-o.45—o.45-0.32 0.90 0.25 0.55 0.55 Cast 0.28-0.70-0.35-0.45-o.45- 0.32 0.90 0.45 0.55 0.55 hln-&lo Wrought 0.28-I .go-o.0.15-0.32 1.70 0.15 Cast o. 28- I.40- 0.35- 0.31 1.70 0.45 a Sulphur was 0.040 per cent maximum and phosphorus 0.030 per cent maximum. Each of these base steels, both cast and wrought, was made with molybdenum contents of o, 0.10, 0.20, 0.30, or 0.40 per cent, the total comprising 20 steels. In addition, each of these 20 compositions was prepared with boron additions in amounts of o.oo15, 0.003, and 0.006 per cent as ferroboron; and in amounts of 0.002 and 0.005 per cent as two commercial intensifier alloys, designated alloys A and B. The boron-treated steels totaled 100, and the entire series 120. Test Procedure Each steel was melted as an individual heat in a 50-lb., magnesia-lined induction furnace. Ingot iron scrap or mixtures of ingot iron and steel. scrap were used for melting stock. Deoxidation was carried out in the furnace with 0.10 per cent aluminum. The aluminum addition was

Jan 1, 1948

By M. C. Udy, P. C. Rosenthal

The use of minute boron additions to steel has been given considerable attention in recent years. Comparisons made between boron-free and boron-containing heats of otherwise identical analysis have indicated, beyond a doubt, the potent effect of boron for increasing hardenability. The actual hardenability increase has been expressed as a multiplying factor, or in terms of the percentage of other alloying elements boron might replace. The present study presents the results of a systematic survey of the use of boron as an addition to two types of alloyed steels and as a replacement for various percentages of molybdenum in these steels. Both cast and wrought steels of the two basic compositions were prepared. Molybdenum, because of its scarcity at the time the investigation was started, was selected as the element to be replaced by boron. Attention was also given to the effects of boron on other properties of the steels, particularly their notched-bar toughness, to investigate other advantages or limitations of this element as an alloy addition. Experimental Work Steels Investigated The two basic compositions used in this survey of boron as a replacement for molybdenum are given in Table I. The cast steels were higher than the wrought steels in silicon content to improve the casting quality. Table i.—Nominal Composition of the Two Base Alloy Steels composition. Per Cent. steel Mn Si Ni Cr Xi-Cr-Mo Wrought 0.28-0.70-0.15-o.45—o.45-0.32 0.90 0.25 0.55 0.55 Cast 0.28-0.70-0.35-0.45-o.45- 0.32 0.90 0.45 0.55 0.55 hln-&lo Wrought 0.28-I .go-o.0.15-0.32 1.70 0.15 Cast o. 28- I.40- 0.35- 0.31 1.70 0.45 a Sulphur was 0.040 per cent maximum and phosphorus 0.030 per cent maximum. Each of these base steels, both cast and wrought, was made with molybdenum contents of o, 0.10, 0.20, 0.30, or 0.40 per cent, the total comprising 20 steels. In addition, each of these 20 compositions was prepared with boron additions in amounts of o.oo15, 0.003, and 0.006 per cent as ferroboron; and in amounts of 0.002 and 0.005 per cent as two commercial intensifier alloys, designated alloys A and B. The boron-treated steels totaled 100, and the entire series 120. Test Procedure Each steel was melted as an individual heat in a 50-lb., magnesia-lined induction furnace. Ingot iron scrap or mixtures of ingot iron and steel. scrap were used for melting stock. Deoxidation was carried out in the furnace with 0.10 per cent aluminum. The aluminum addition was

Jan 1, 1948

By A. Hultgren, B. Herrlander

The original aim of the present investigation was to study the mechanism of cracking on hot-deforming red-short steels. During the microscopical examination of hot-deformed soft steels attention was directed to various patterns of so-called veining in the ferrite, as related to variations in the deformation procedure and the heat-treatment of the steels. Later, a hot-short steel of higher carbon content was also studied. Historical Storey1 suggested in 1914 that veining developed during the gamma-alpha transformation from the growing together of several alpha nuclei within the same gamma grain. Rawdon and Berglund2 found that in soft steel, forged somewhat below A3, the ferrite showed profuse veining, while usually there was very little veining after forging at very high temperatures. They emphasized the similarity between the pattern of veining formed by deformation about 600°C and that of slip lines formed by deformation at room temperature. They also suggested a relation between the gamma-alpha transformation and veining. They became convinced from etching and from the slip-line pattern in cold-deformed material, that the orientation within any ferrite grain showing veining was uniform. In recrystallized alpha grains veining was absent. Veining did not appear materially to affect the properties of ferrite. Independent of veining, continuous networks, attributed to preexisting delta and gamma grain boundaries were sometimes found, the former only in cast steel. Those networks were usually connected with small inclusions. Ammermann and Kornfeld3 confirmed that recrystallized ferrite grains showed no veining. In soft steel annealed between A1 and A3 there was veining only in the ferrite grains formed during cooling. Electrolytic iron deformed at 880' was free from veining. Bannister and Jones' considered veining in ferrite to be a manifestation of microscopical and sub-microscopica~ inclusions. NorthCOtt,5,6 in studies on veining in steel and other metals and alloys, attributed veining to oxide inclusions. It could generally be removed by annealing in hydrogen, a suggestion earlier made by Rawdon and Berg1und. TrittOn (discussion in ref. 5) emphasized the importance of perfect polishing for bringing out veining by etching, and thought the veins or subboundaries were produced during the gamma-alpha transformation, when it was rapid, as a result of the volume change and contraction during cooling. Slight distortion would not allow projecting parts of a growing crystal to join up with atomic symmetry. Hanemann and co-workers showed that veining in ferrite consisted of ridges

Jan 1, 1948

By B. Herrlander, A. Hultgren

The original aim of the present investigation was to study the mechanism of cracking on hot-deforming red-short steels. During the microscopical examination of hot-deformed soft steels attention was directed to various patterns of so-called veining in the ferrite, as related to variations in the deformation procedure and the heat-treatment of the steels. Later, a hot-short steel of higher carbon content was also studied. Historical Storey1 suggested in 1914 that veining developed during the gamma-alpha transformation from the growing together of several alpha nuclei within the same gamma grain. Rawdon and Berglund2 found that in soft steel, forged somewhat below A3, the ferrite showed profuse veining, while usually there was very little veining after forging at very high temperatures. They emphasized the similarity between the pattern of veining formed by deformation about 600°C and that of slip lines formed by deformation at room temperature. They also suggested a relation between the gamma-alpha transformation and veining. They became convinced from etching and from the slip-line pattern in cold-deformed material, that the orientation within any ferrite grain showing veining was uniform. In recrystallized alpha grains veining was absent. Veining did not appear materially to affect the properties of ferrite. Independent of veining, continuous networks, attributed to preexisting delta and gamma grain boundaries were sometimes found, the former only in cast steel. Those networks were usually connected with small inclusions. Ammermann and Kornfeld3 confirmed that recrystallized ferrite grains showed no veining. In soft steel annealed between A1 and A3 there was veining only in the ferrite grains formed during cooling. Electrolytic iron deformed at 880' was free from veining. Bannister and Jones' considered veining in ferrite to be a manifestation of microscopical and sub-microscopica~ inclusions. NorthCOtt,5,6 in studies on veining in steel and other metals and alloys, attributed veining to oxide inclusions. It could generally be removed by annealing in hydrogen, a suggestion earlier made by Rawdon and Berg1und. TrittOn (discussion in ref. 5) emphasized the importance of perfect polishing for bringing out veining by etching, and thought the veins or subboundaries were produced during the gamma-alpha transformation, when it was rapid, as a result of the volume change and contraction during cooling. Slight distortion would not allow projecting parts of a growing crystal to join up with atomic symmetry. Hanemann and co-workers showed that veining in ferrite consisted of ridges

Jan 1, 1948

By R. S. Dean

The commercial production of electrolytic manganese on a small scale commenced in 1939. The writer made a short report on the progress of production and utilization in Mining and Metallurgy for January 1941. Progress during the last two years, naturally, has been more rapid. In June 1940, Congressman (now Senator) Scrugham presented a proposal to Congress for an appropriation authorizing erection of a pilot plant for " the production of metallic manganese by electrolytic or other means." As a result of this, a pilot plant, having a capacity of a few hundred pounds per day was built at Boulder City, Nev. Subsequent appropriations permitted expansion of the plant to about a ton per day. Much has been learned in this plant about electrolytic manganese, and the product even in its small way has found war uses of considerable importance. Meanwhile, the Knoxville plant of the Electro Manganese Corporation, the sole commercial producer, was expanded to about 4 tons per day to supply needed material for the war program. Other proposals for commercial plants have not been favorably received by the War Production Board. The Bureau reported on one such proposal by the American Alloys and Chemical Corporation. Practice at BouldeR City The present practice, as carried out at Boulder City, which is being described in more detail in a current paper by the Bureau staff at Boulder City, starts with a manganese dioxide ore containing about 20 per cent manganese. Currently this ore comes from the Three Kids deposit near Las Vegas. The steps for the production of electrolytic manganese, in general, are as follows: 1. The manganese in the ore is reduced to MnO. 2. The MnO is dissolved from the ore in spent electrolyte, which contains about 38 to 47 grams per liter of free sulphuric acid, 135 grams per liter of ammonium sulphate, and 10 to 12 grams per liter of Mn as sulphate. In the leaching step the manganese is built back up to 32 to 36 grams per liter of Mn as sulphate; the pH is adjusted to neutral by means of gaseous ammonia and passed through thickeners. 3. The electrolyte is purified by adding H2S, which precipitates the heavy metals. After filtering, ferrous sulphate is added to the solution and oxidized with air and the solution is filtered and clarified on a precoat filter. 4. The purified catholyte is electrolyzed in a diaphragm cell using stainless-steel cathodes and lead-silver anodes. The current density is about 45 amp. per sq. ft. Current efficiencies of 60 to 65 per cent are regularly obtained.

Jan 1, 1944

By R. S. Dean

The commercial production of electrolytic manganese on a small scale commenced in 1939. The writer made a short report on the progress of production and utilization in Mining and Metallurgy for January 1941. Progress during the last two years, naturally, has been more rapid. In June 1940, Congressman (now Senator) Scrugham presented a proposal to Congress for an appropriation authorizing erection of a pilot plant for " the production of metallic manganese by electrolytic or other means." As a result of this, a pilot plant, having a capacity of a few hundred pounds per day was built at Boulder City, Nev. Subsequent appropriations permitted expansion of the plant to about a ton per day. Much has been learned in this plant about electrolytic manganese, and the product even in its small way has found war uses of considerable importance. Meanwhile, the Knoxville plant of the Electro Manganese Corporation, the sole commercial producer, was expanded to about 4 tons per day to supply needed material for the war program. Other proposals for commercial plants have not been favorably received by the War Production Board. The Bureau reported on one such proposal by the American Alloys and Chemical Corporation. Practice at BouldeR City The present practice, as carried out at Boulder City, which is being described in more detail in a current paper by the Bureau staff at Boulder City, starts with a manganese dioxide ore containing about 20 per cent manganese. Currently this ore comes from the Three Kids deposit near Las Vegas. The steps for the production of electrolytic manganese, in general, are as follows: 1. The manganese in the ore is reduced to MnO. 2. The MnO is dissolved from the ore in spent electrolyte, which contains about 38 to 47 grams per liter of free sulphuric acid, 135 grams per liter of ammonium sulphate, and 10 to 12 grams per liter of Mn as sulphate. In the leaching step the manganese is built back up to 32 to 36 grams per liter of Mn as sulphate; the pH is adjusted to neutral by means of gaseous ammonia and passed through thickeners. 3. The electrolyte is purified by adding H2S, which precipitates the heavy metals. After filtering, ferrous sulphate is added to the solution and oxidized with air and the solution is filtered and clarified on a precoat filter. 4. The purified catholyte is electrolyzed in a diaphragm cell using stainless-steel cathodes and lead-silver anodes. The current density is about 45 amp. per sq. ft. Current efficiencies of 60 to 65 per cent are regularly obtained.

Jan 1, 1944

By E. A. Nesbitt

The important permanent-magnet alloys 15 years ago contained carbon and depended upon it for their permanent-magnet properties. In recent years great advances have been made in a number of new alloys substantially free of carbon. These include Remalloy,' an alloy of iron, cobalt, and molybdenum, the various iron-nickel-aluminum alloys of the Mishima2 type, and the more recent Alnico V3 in which the best properties are developed by heat-treating in a magnetic field. The Mishima and the Alnico types of alloys are very brittle and after casting must be ground to size. This article describes a permanent-magnet material4 that can be machined as cast, and when handled with some care can be rolled to thin sheet and drawn to fine wire. Optimum magnetic properties are developed by a simple heat-treatment, and at this time the ductility of the alloy disappears. If the cast alloy is heat-treated without cold-working the mag netic energy product is about 1.0 X 106. This material is called Vicalloy I (pronounced vicalloy, V for vanadium, I for iron, and C for cobalt). When the alloy is first hot-worked and then considerably reduced in area by cold-swaging, drawing or grooved rolling, it is found that after heat-treatment the magnetic energy product measured in the direction of previous extension can be increased to 2.0 to 3.5 X 106. These higher values have been obtained only after reductions in area of 75 to 95 per cent, and since the material hardens considerably during working, the reduction ordinarily can be accomplished only in fine wires, although larger sizes can be reduced by the use of powerful machinery. When the material has been cold-worked in such a manner as to improve its energy product substantially, it is called Vicalloy 11. Commercial use of Vicalloy has been made in the recording of speech. Tape 0.002 by 0.050 in. has been found superior in quality to all others tried, and is now used in the Western Electric Mirrorphone 5 In a variety of forms Vicalloy I or I1 is useful in many laboratory instruments, and in experiments in which the sizes and shapes required cannot be obtained with the other more brittle permanent-magnet materials. In 1936, during a study of other age-hardening magnetic systems, it occurred to the author that some of the iron-cobalt-vanadium alloys might be suitable for permanent magnets. The proportions of iron, cobalt, and vanadium were varied and a range of alloys of particular interest was found, as shown by the indicated area in Fig. I. The limits of this range of composition are 30 to 52 per cent iron, 36 to 62 per cent cobalt, and 4 to 16 per cent vanadium. The compositions indicated by points outside this area were also tested. Preparation and Properties OF Vicalloy I The alloys were usually cast in graphite molds into bars 3/4 in. in diameter and

Jan 1, 1946

By E. A. Nesbitt

The important permanent-magnet alloys 15 years ago contained carbon and depended upon it for their permanent-magnet properties. In recent years great advances have been made in a number of new alloys substantially free of carbon. These include Remalloy,' an alloy of iron, cobalt, and molybdenum, the various iron-nickel-aluminum alloys of the Mishima2 type, and the more recent Alnico V3 in which the best properties are developed by heat-treating in a magnetic field. The Mishima and the Alnico types of alloys are very brittle and after casting must be ground to size. This article describes a permanent-magnet material4 that can be machined as cast, and when handled with some care can be rolled to thin sheet and drawn to fine wire. Optimum magnetic properties are developed by a simple heat-treatment, and at this time the ductility of the alloy disappears. If the cast alloy is heat-treated without cold-working the mag netic energy product is about 1.0 X 106. This material is called Vicalloy I (pronounced vicalloy, V for vanadium, I for iron, and C for cobalt). When the alloy is first hot-worked and then considerably reduced in area by cold-swaging, drawing or grooved rolling, it is found that after heat-treatment the magnetic energy product measured in the direction of previous extension can be increased to 2.0 to 3.5 X 106. These higher values have been obtained only after reductions in area of 75 to 95 per cent, and since the material hardens considerably during working, the reduction ordinarily can be accomplished only in fine wires, although larger sizes can be reduced by the use of powerful machinery. When the material has been cold-worked in such a manner as to improve its energy product substantially, it is called Vicalloy 11. Commercial use of Vicalloy has been made in the recording of speech. Tape 0.002 by 0.050 in. has been found superior in quality to all others tried, and is now used in the Western Electric Mirrorphone 5 In a variety of forms Vicalloy I or I1 is useful in many laboratory instruments, and in experiments in which the sizes and shapes required cannot be obtained with the other more brittle permanent-magnet materials. In 1936, during a study of other age-hardening magnetic systems, it occurred to the author that some of the iron-cobalt-vanadium alloys might be suitable for permanent magnets. The proportions of iron, cobalt, and vanadium were varied and a range of alloys of particular interest was found, as shown by the indicated area in Fig. I. The limits of this range of composition are 30 to 52 per cent iron, 36 to 62 per cent cobalt, and 4 to 16 per cent vanadium. The compositions indicated by points outside this area were also tested. Preparation and Properties OF Vicalloy I The alloys were usually cast in graphite molds into bars 3/4 in. in diameter and

Jan 1, 1946

By R. B. Snow

This paper describes a technique of polishing and etching specimens of open-hearth furnace slags or hearth aggregates for identification of the crystalline constituents —lime (CaO), tricalcium silicate (3CaO.SiO2), dicalcium silicate (2CaO.-SiO2), rnonticellite (CaO.MgO.SiO2), or forsterile (2MgO.SiO2), with especial em-phasis on the mineral merwinite (3CaO.-MgO.2SiO2). With proper standardization, this identification does not require the use of the petrographic microscope. The composition of basic open-hearth slags and furnace bottoms falls, almost without exception, within systems containing CaO, MgO, 'IFeO,,, MnO and SiOz, in which the number of basic molecules so greatly exceeds the orthosilicate ratio (two molecules of base to one of silica) that free basic oxides, and combinations between them such as alumi-nates or ferrites, are present in cooled specimens. Orthosilicates of (CaO + NgO) are the most common in such specimens, since in nearly all cases, except premelt slags, the molecular ratio of (CaO + MgO) to SiO, is more than 2 to I. When sufficient lime is available it combines with the silica to form dicalcium silicate (2Ca0.Si02), which contains little, if any, IvfgO, FeO or MnO in solid solution whereas the latter oxides combine to form the oxide solid solution known as periclase. If the lime present is insufficient to form dicalcium silicate (2Ca0.Si02) it combines with Mgo to form either merwillite or moIlticellite (SiOz); these minerals take little if any FeO or MnO into solid solution and the remaining MgO, FeO and hInO combine as periclase. This generalization seems to be valid for basic slags and furnace bottoms, since minerals such as Ca0.MnOSi02 and CaO.FeO.SiOz are found only in slags in which the lime-silica ratio is less than 2 and are not observed in 'pecimens from furnace bottoms. The identification of crystalline constituents in such materials, especially of fine crystals in the groundmass, is difficult under the petrographic microscope. They are often masked by their neighbors because of their small size in relation to the thickness of the thin section and because of the presence of Opaque Or colored constituents. The indices of refraction and the optical sign of the mineral are sometimes difficult to determine because of the small size or because Of twinning or of inclusions within the crystal. Moreover, the positive identification of merwinite (3Ca0.Mg0.2Si02) from its optical properties is usually difficult in the presence of dicalcium silicate (zCaO.SiO2). CaO, MgO, jCa0.Si02 and 2Ca0.SiOz in open-hearth 'lags have been identified for a number of years in the U.S. Steel Corporation Laboratory by the usual

Jan 1, 1948

By R. B. Snow

This paper describes a technique of polishing and etching specimens of open-hearth furnace slags or hearth aggregates for identification of the crystalline constituents —lime (CaO), tricalcium silicate (3CaO.SiO2), dicalcium silicate (2CaO.-SiO2), rnonticellite (CaO.MgO.SiO2), or forsterile (2MgO.SiO2), with especial em-phasis on the mineral merwinite (3CaO.-MgO.2SiO2). With proper standardization, this identification does not require the use of the petrographic microscope. The composition of basic open-hearth slags and furnace bottoms falls, almost without exception, within systems containing CaO, MgO, 'IFeO,,, MnO and SiOz, in which the number of basic molecules so greatly exceeds the orthosilicate ratio (two molecules of base to one of silica) that free basic oxides, and combinations between them such as alumi-nates or ferrites, are present in cooled specimens. Orthosilicates of (CaO + NgO) are the most common in such specimens, since in nearly all cases, except premelt slags, the molecular ratio of (CaO + MgO) to SiO, is more than 2 to I. When sufficient lime is available it combines with the silica to form dicalcium silicate (2Ca0.Si02), which contains little, if any, IvfgO, FeO or MnO in solid solution whereas the latter oxides combine to form the oxide solid solution known as periclase. If the lime present is insufficient to form dicalcium silicate (2Ca0.Si02) it combines with Mgo to form either merwillite or moIlticellite (SiOz); these minerals take little if any FeO or MnO into solid solution and the remaining MgO, FeO and hInO combine as periclase. This generalization seems to be valid for basic slags and furnace bottoms, since minerals such as Ca0.MnOSi02 and CaO.FeO.SiOz are found only in slags in which the lime-silica ratio is less than 2 and are not observed in 'pecimens from furnace bottoms. The identification of crystalline constituents in such materials, especially of fine crystals in the groundmass, is difficult under the petrographic microscope. They are often masked by their neighbors because of their small size in relation to the thickness of the thin section and because of the presence of Opaque Or colored constituents. The indices of refraction and the optical sign of the mineral are sometimes difficult to determine because of the small size or because Of twinning or of inclusions within the crystal. Moreover, the positive identification of merwinite (3Ca0.Mg0.2Si02) from its optical properties is usually difficult in the presence of dicalcium silicate (zCaO.SiO2). CaO, MgO, jCa0.Si02 and 2Ca0.SiOz in open-hearth 'lags have been identified for a number of years in the U.S. Steel Corporation Laboratory by the usual

Jan 1, 1948

By Shueling Woo, Charles S. Barrett, Robert F. Mehl

When one solid phase is generated from another, fixed and rational orientation relationships are observed to subsist between the parent and the new crystal. The principle has been proposed1 that the relationships obtaining between the same two phases are identical whatever the type of phase change involved, or whichever of the phases may act as parent. Many types of phase changes have been studied: allotropic transformations in pure metals, precipitation from solid solutions, eutectoid decompositions, peritectic reactions, the formation of oxide layers, and others2 Phase layers generated by diffusion of one metal into another constitute an additional type of reaction. This report provides new data on the orientation relationships occurring between phase layers formed by diffusion, in the Cu-Zn system,§ and furnishes further evidence in support of the principle. Materials and Methods The Cu-Zn system is especially suitable for these studies since the high vapor pres- sure of Zn provides an opportunity readily to create one phase poorer in Zn upon a parent phase richer in Zn by removing Zn by vaporization, or to reverse this process by adding Zn from Zn vapor; e.g., the beta phase can be found as a layer upon gamma by dezincing, or gamma can be found on beta by addition of Zn from Zn vapor. Since orientation relationships are to be sought, requiring crystals large enough for orientation determinations, it is convenient to prepare single crystals or large grains of the parent phase. The samples were prepared from oxygen-free high-conductivity copper and Horse Head zinc by melting under a borax flux in a graphite crucible. Large crystals of beta brass were grown by the strain-anneal method, in plates sectioned from the cast ingots, by the following procedure. Rockwell-hardness indents were made at I cm. intervals and with various loads, then the plates were annealed for 4 days at 700°C. in Nichrome boxes packed with beta-brass chips and graphite. Surface layers were removed by grinding and etching to a depth of about 1.0 mm., and individual grains were cut out. Crystals of alpha and gamma brass were prepared by freezing, by lowering a graphite crucible at speeds less than one inch per hour through a furnace held at a constant temperature above the melting point. Single crystals were cut out with a jeweler's saw, ground, polished and etched. Care was taken to remove cold-worked layers from the finished crystals

Jan 1, 1944

By Charles S. Barrett, Robert F. Mehl, Shueling Woo

When one solid phase is generated from another, fixed and rational orientation relationships are observed to subsist between the parent and the new crystal. The principle has been proposed1 that the relationships obtaining between the same two phases are identical whatever the type of phase change involved, or whichever of the phases may act as parent. Many types of phase changes have been studied: allotropic transformations in pure metals, precipitation from solid solutions, eutectoid decompositions, peritectic reactions, the formation of oxide layers, and others2 Phase layers generated by diffusion of one metal into another constitute an additional type of reaction. This report provides new data on the orientation relationships occurring between phase layers formed by diffusion, in the Cu-Zn system,§ and furnishes further evidence in support of the principle. Materials and Methods The Cu-Zn system is especially suitable for these studies since the high vapor pres- sure of Zn provides an opportunity readily to create one phase poorer in Zn upon a parent phase richer in Zn by removing Zn by vaporization, or to reverse this process by adding Zn from Zn vapor; e.g., the beta phase can be found as a layer upon gamma by dezincing, or gamma can be found on beta by addition of Zn from Zn vapor. Since orientation relationships are to be sought, requiring crystals large enough for orientation determinations, it is convenient to prepare single crystals or large grains of the parent phase. The samples were prepared from oxygen-free high-conductivity copper and Horse Head zinc by melting under a borax flux in a graphite crucible. Large crystals of beta brass were grown by the strain-anneal method, in plates sectioned from the cast ingots, by the following procedure. Rockwell-hardness indents were made at I cm. intervals and with various loads, then the plates were annealed for 4 days at 700°C. in Nichrome boxes packed with beta-brass chips and graphite. Surface layers were removed by grinding and etching to a depth of about 1.0 mm., and individual grains were cut out. Crystals of alpha and gamma brass were prepared by freezing, by lowering a graphite crucible at speeds less than one inch per hour through a furnace held at a constant temperature above the melting point. Single crystals were cut out with a jeweler's saw, ground, polished and etched. Care was taken to remove cold-worked layers from the finished crystals

Jan 1, 1944

By David J. Mack

A survey and critical appraisal of published information about Young's modulus was originally made by the writer because of a complete lack of information about this very important quantity in works on. mechanics, physical metallurgy, physics, and other sciences. It was felt that a comprehensive summary of such information about Young's modulus might be of general interest and may suggest new problems or lines of attack on other problems that involve E. Hence this paper has been prepared even though it is only a review, with no new ideas or experimental data. The references are only those used in assembling the paper. No attempt was made to compile a complete bibliography on the subject of Young's modulus. CalculatioX oP YouNg's Modulus from Theoretical or Empirical CoNsidera-tions Many attempts have been made to compute E from theoretical considerations. The earliest calculations were probably those of Tomlenson (1883) and Souther-land (1891), but Fess'enden4',~ made extensive calculations a few years later and arrived at the relationship: £ = 78.TOw(i)* where V is the atomic volume. In 1923, Peczalski, working from the same standpoint but with more information available on atomic structure, arrived at an identical relationship, which he expressed as: where B is a constant of value about 8 . 107, p is density and mis atomic weight; the atomic volume, of course, being ml p. Portevind4 immediately pointcd out that while the equations of Fessenden and Peczalski were true of the common metals, they gave low values for the refractory metals with high moduli (Fig. I). He showed (Fig. 2) that much better concordance was obtained by means of an empirical equation of the form: F- KT" where K is a constant; T the absolute melting point; V, atomic volume, and a and b constants of value approximately I and 2 respectively. This constant K in the Portevin equation is not necessarily the same for all elements, as shown by Thompson. Other equations for E based upon theoretical considerations have been derived by Honda and Yamada49 and Lasareff.50 The most recent calculations of elastic moduli based upon quantum mechanics are probably those of Fuchs He extended the method of Wigner and Seitz for calculating the lattice energy and compressibility of monovalent metals and computed the elastic constants of lithium, 18

Jan 1, 1946