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Institute of Metals Division Iron and Steel Division Extractive Metallurgy Division J C Kinnear, Jr, President A W Thornton, Past President W R Hibbard, Jr, Vice-President R W Shearman, Secretary-Treasurer AIME, 29 West 39th Street New York 18, N Y Board of Directors J B Austin, D L Clark, H B Emerick, T D Jones, H H Kellogg, C T Marshall, E E Schumacher, C S Smith, J D Sullivan

Jan 1, 1957

E. O?C. Acker Jinzoc Adachi Truman H. Aldrich Walter H. Aldridge Anson W. Allen John H. Allen William R. Appleby W. Lawrence Austin W. S. Ayres David Baker John Henry Banks Henry C. Banon Alonzo F. Bardwell Walter A. Barrows H.C. Bellinger Wm de L. Benedict Robert Bentley John F. Berry Andrew A Blair George D. Blood Charles A. Bohn Walter M. Brodie Elwin F. Brown Fredrick Burbidge H.H. Burden Charles Butters

Jan 1, 1944

E. O?C. Acker Jinzoc Adachi Truman H. Aldrich Walter H. Aldridge Anson W. Allen John H. Allen William R. Appleby W. Lawrence Austin W. S. Ayres David Baker John Henry Banks Henry C. Banon Alonzo F. Bardwell Walter A. Barrows H.C. Bellinger Wm de L. Benedict Robert Bentley John F. Berry Andrew A Blair George D. Blood Charles A. Bohn Walter M. Brodie Elwin F. Brown Fredrick Burbidge H.H. Burden Charles Butters

Jan 1, 1943

By Cyril Wells, Robert F. Mehl

The purpose of this investigation was to prepare high-purity iron-carbon alloys, to determine as precisely as possible the A3(GOS), the Acm(SE), and the A1(PSK) transformation temperatures in the meta-stable iron-iron carbide system (up to 1.4 per cent C), and to study the influence of carefully controlled and constant rates of heating and cooling (1/8 to 2o C. per minute) upon the transformation temperature and upon the resulting microstructure. Apart from the general metallurgical importance attending the completion of such a task, it will be noted that highly precise transformation temperatures are of interest and importance in the thermodynamic study of this system1. Several determinations of the iron-iron carbide diagram have been made2. These suffer generally from the low purity of the alloys used, and from the application of methods of investigation of low sensitivity. Attempts to compare and weight the many data available in an effort to construct as accurate a diagram as possible have, in the main, been unsuccessful. In general, the wide differences in the results of different workers, the lack of complete analytical data, the omission of important details relating to the heat-treatment of the alloys, and the meager description of the methods of investigation applied, render such a process of averaging nearly meaningless. It may be seen from the summaries by Epstein2 and Körber and Oelsenl that variations as great as 30' C. exist in the reported data on the A3 temperatures in iron-carbon alloys. Fortunately, the Aa temperature in pure iron is now known to be 910" C. ± 1° C.3 The best recent work on iron-carbon alloys is that of Esser4 on hypoeutectoid steels and that of Sato5 on both hypoeutectoid and hyper-eutectoid steels. Esser made a dilatometric and magnetic study of iron and iron-carbon alloys containing 0.06 to 0.86 per cent C, for which a high purity is claimed. Presumably the electrolytic iron was reasonably pure, but the alloys contained considerable impurities, on the average totaling 0.35 per

Jan 1, 1937

By Edward Keller

ALTHOUGH the subject matter contained in this paper is presented under a new title, its writing was induced by the work of S. Skowronski and K. W. McComas.1 A discussion of the latter is here combined with the presentation of new and independent data. The physical constants quoted by the authors of that paper, such as melting and boiling points, as also the chemical reactions of the substances involved, are so well established that the numerous confirmatory experiments and their application to metallurgical conditions leave little or no room for criticisms. On the contrary, they will greatly aid to dispel any lingering skepticism as to possible considerable losses of copper accruing when the metal is melted in the presence of its own and, probably, other chlorides and when no provision is made for the recovery of volatile products. Exception may be taken, to some degree, to that part of the work under the caption Melting Drillings of Copper Mixed With Sodium Chloride. These experiments should have been extended in order to make the implied reactions clear and the results convincing. The latter are stated to be, "that when sodium chloride is present above certain minimum quantity relative to copper some of it is found to volatilize in the melt while another part reacts with the metallic copper to produce the volatile cuprous chloride, and that When smaller quantities of sodium chloride are present the reaction is complete and all chlorine is found to have volatilized as cuprous chloride." Metallic copper and sodium chloride are mentioned as the sole reagents in this chemical process, acting, according to the authors, in a neutral atmosphere of carbon dioxide. The chemist, quite naturally, will inquire into the formula according to which that reaction takes place and find that it would present itself as follows: 2Cu + 2NaC1 = 2Na + Cu2Cl2 It would appear that the truth of this formula will be challenged and a more rational explanation of the reaction sought. Apparently, carbon dioxide is neutral relative to metallic copper in its molten state; whether it remains neutral in the presence of sodium chloride is quite another question. Does it not at that high temperature become an oxidizer and do not the conditions of the reaction change to those expressed in the following formula? 2Cu + 2NaCl + 2002 = Na2CO3 + CO + Cu2Cl2

Jan 9, 1919

By J. B. Darby

In an investigation of various equiatomic alloys of transition elements, Dwightl did not find evidence for the CsC1-type ordered structure in VMn. More recently waterstrat2 has reported the occurrence of this structure in the alloy his sample had been homogenized at 1150°C and subsequently reannealed at 725°C after crushing to powder. The purpose of this note is to confirm the existence of the CsC1-type ordered structure in VMn and to report a lattice parameter. An equiatomic alloy was prepared by arc-melting under an argon atmosphere vanadium and electrolytic manganese having a purity of at least 99.9 pet. The composition was obtained by charging an excess of manganese and monitoring the weight loss until the stoichiometric composition was achieved; the assumption that the loss during melting was due to the vaporization of manganese appears to be a valid one.' The calculated composition of the sample prepared was 49.85 at. pct V and 50.15 at. pct Mn. Specimens taken from the fractured as-cast buttons were wrapped in molybdenum foil and encapsulated in quartz capsules under vacuum; there was no evidence of a reaction between the alloy and foil. The X-ray patterns were obtained using filtered copper or chromium radiation with a Straumanis-type Debye-Scherrer camera having a diameter of 114.6 mm. The CrKa radiation is best suited for detecting the superlattice lines since it overcomes the slight difference in scattering powers of the components by making use of the anomalies in atomic scattering factor which occur when the frequency of the radiation is close to the characteristic absorption frequency of one of the elements.394 A specimen in the as-cast condition, as well as one which had been annealed for 1 week at 900°C and water-quenched, gave a diffraction pattern for the disordered bcc solid solution. However, when samples were homogenized at 900°C for 3 weeks followed by a subsequent anneal at 700°C for 6 to 7 weeks and then water-quenched or furnace-cooled, the (loo), (1111, and (210) Ka superlattice lines were observed in films taken with chromium radiation. An ao = 2.904 * 0.001 was determined by a Nelson-Riley extrapolation from a film taken with CuKa radiation. The author wishes to thank Mr. J. W. Downey and Mr. L. J. Norton for carrying out the experimental portion of this investigation. This work was performed at the Argonne National Laboratory under the auspices of the U.S. Atomic Energy Commission.

Jan 1, 1963

By M. H. Mueller, P. A. Beck, W. P. Chernock

HARRIS introduced the term inverse pole figure for the diagrams previously used by Barrett' to represent the volume fraction p(a,ß) of material in various orientations with respect to the fiber axis in a specimen with fiber texture. The direction of the fiber axis with respect to the crystal lattice is specified by the two independent angle variables a and ß. In such a diagram p gives the fiber axis density distribution over the various crystallographic directions. The p(a,ß) may be given in a two-dimensional diagram by means of contour lines (p=con-stant) in a stereographic projection. The inverse pole figure incorporates in a single contour line diagram all the information given conventionally in pole charts for various (hkl) reflections and, therefore, it represents a considerable degree of condensation. Also, it gives this information directly in terms of the volume fraction of the specimen in various orientations, and it obviates the necessity of referring back and forth between the various pole charts in an attempt to interpret them in terms of orientations. Inverse pole figures undoubtedly represent the most adequate method for describing a texture in a quantitative manner and, at least in the case of fiber textures, this method is certainly a practicable one. Harris described1 a simple procedure for determining at least the rough outlines of a quantitative inverse pole figure. This method depends on the measurement of the intensity of X-ray reflections from various (hkl) planes lying parallel to the same specimen surface. In the case of a rod, reflections from a plane specimen surface perpendicular to the rod axis (fiber texture axis) are to be used. Harris has shown that the measured intensities I,,, reflected by crystal planes (hkl), aligned perpendicularly to the fiber axis of a specimen with unknown fiber texture, and the corresponding intensity Iu nkl from a random specimen are related to the p(a,ß) value for the unknown texture by the following equation IAkl/I°Aklw=c/c°.p(a, ß) [1] where C and C° re parameters depending on the crystal imperfections (extinction), surface preparation (e.g., thickness of oxide layer), the absorption coefficient in the specimen, effective line breadth, and instrumental factors. The values of the variables a and ß1 defining the crystal orientation, depend only on the indices hkl of the reflecting plane perpendicular to the fiber axis. Since p, as described by Harris1, is normalized in such a way that upon integration over the total orientation range its mean value is 1, the parameter C/C may be calculated approximately by the summation of Eq. l over all available X-ray reflections hkl, corresponding to a set of discrete orientations (rather than by integration over the complete range of orientations represented by the continuous variables a and ß). From the summation of Eq. 1 1/n SI AKL/I° Akl = C/C. SP(a,ß)/n is obtained where, as stated above 2p(a,ß) =1. n Therefore

Jan 1, 1959

By Gerd L. Klinkenborg

Ever see a curvaceous conveyor belt? Reports from Europe indicate appreciable interest in a relatively new system that has revamped some of the concepts of what single-belt conveyors can and cannot do. Invented about ten years ago by two German engineers, the so-called "Serpentix" single-belt conveyor, patented by the firm of Albert Stuebbe, is now operating at more than 100 foreign installations. The system boasts of three unique capabilities-climbing, turning and side discharge of load-which give it, the advantages of a multiple-unit system. The new conveyor needs only one drive station (which may be combined with tensioning), and except for the immediate area of the drive, is capable of conveying loads over a non-linear route. Transfer points are eliminated, and discharge points, which may be located anywhere along the route, may be operated automatically or manually, by direct or remote control.

Jan 1, 1962

By E. Miller, K. Komarek, W. Johnson

The activity of aluminum in solid Cr-A1 alloys has been measured by an isopiestic technique between Cr-A1890' and 1126" and 13 and 80 at. pct Al. The integral free energy of mixing has a minimum value of —5600 cal per g-atom at 59 at. pct Al. The maximum solid solubility of aluminum in chromium was determined to be 43 at. pct Al, and the composition limits of the compounds CrA14, Cr4A19, and Cr5Al, at 1000"~ were found to be 79 to 80, 66 to 70, and 59 to 63 at. pct Al, respectively. The thermodynamic properties of the Cr-A1 system have been investigated as part of a thermodynamic study of aluminum-transition metal systems.172 Little information is available on the equilibrium properties of the Cr-A1 system. The heats of formation of solid Cr-A1 alloys have been determined by Kubaschewski and Haymer at 600" and low-temperature specific heat data have also been obtained.~ More extensive work has been performed on the phase diagram, and a compilation has been provided by Hansen and Anderko,~ their phase diagram at elevated temperatures being essentially based on the work of Bradley and LU.~ The high-temperature portion of the phase diagram shows an intermediate phase CrA14 decomposing peritectically at 1018°C and existing at 82 at. pct A1 at 1000°C. They also identified the compounds with solubility limits of 72 to 75 at. pct A1 at 1000°C, and Cr5A1,, existing at 61 at. pct A1 at 1000°C. The maximum solid solubility of aluminum in chromium at 1000°C was found to be 46 at. pct Al. These elevated-temperature data were obtained by examination of quenched samples and were considered as less precise than the lower-temperature data. Koester, Wachtel, and Grube7 have revised the phase diagram as a result of their magnetic susceptibility and X-ray study. The results of this work differ appreciably from those of Bradley and Lu at temperatures above 800°C. The CrA1, compound is given as existing between 79 and 81 at. pct A1 at 1000°C, and they do not indicate the presence of a CrA13 phase reported by Bradley and Lu. They also report the compound Cr4Alg as having solubility limits of 66 to 70 at. pct A1 at 1000°C, while Bradley and Lu show this compound stable only up to 870°C. Koester et al. state that the high-temperature modification of the compound Cr5A18 is stable down to 1125"C, and not 980°C as stated by Bradley and Lu, and that the low-temperature modification of Cr5Al, has a range of homogeneity of 58 to 63 at. pct A1 at 1000°C. They also report that the maximum solid solubility of aluminum in chromium is 43 at. pct A1 at 1000°C. APPARATUS AND EXPERIMENTAL PROCEDURE An isopiestic method was employed which has been successfully applied to the determination of aluminum activities in solid ~e-All and Ni-Al alloys. Alloy specimens were held at different positions in a temperature gradient and were equilibrated with aluminum vapor from an aluminum reservoir kept at the temperature minimum of an impressed thermal gradient in a closed alumina system. Diffusion of aluminum into the specimens occurred until equilibrium was reached, at which the partial pressure of aluminum in each of the specimens was given by the vapor pressure of the pure aluminum reservoir. The activity of aluminum referred to liquid aluminum as the standard state in a given equilibrated sample at temperature T could therefore be expressed by: vapor pressure of pure aluminum at _ the temperature of the reservoir Vapor pressure of pure liquid aluminum, at specimen temperature T Since both the temperature of the aluminum reservoir and the specimen temperatures were determined experimentally, and the vapor pressure of pure aluminum is known as a function of temperature,' the activity of aluminum in a given aluminum alloy of known composition could be calculated. Initial runs were made with samples consisting of pure chromium chips placed in alumina crucibles. These runs exhibited large inconsistencies, indicating that equilibrium was not attained. High aluminum content Cr-A1 alloy powders were therefore substituted for the pure chromium specimens. The starting composition of the alloys was adjusted through experimentation until the concentration change necessary to attain equilibrium was small. In this manner, consistent results were obtained in reasonable times. SPECIMEN PREPARATION Alloy specimens were prepared from chromium of 99.997 pct minimum metallic purity: with 0.028 to 0.038 pct H, 0.0002 pct N, and 0.27 to 0.46 pct 0 (Aviquipo, Inc.). The aluminum had a purity of 99.99+ pct and the following impurities: 0.003 pct Cu; 0.002 pct S; 0.002 pct Fe; 0.001 pct Pb; 0.001 pct Ga (Aluminum Corp. of America). Alloy powders were prepared from weighed mixtures of chromium and aluminum by double-arc melt-

Jan 1, 1969

C.A. MITKE, *Bisbee, Ariz. (written-discussion-? ).--The advantages of the incline-slicing system have' been ably pointed out by Mr. Scott in his paper in the November, 1917, Bulletin, and also by Mr. P. B. Scotland in the Engineering & Mining Journal, Apr. 7, 1917. From the standpoint of efficiency and economy this method is far superior to the top-slice system as generally practised. However, as regards ventilation, and the danger of mine fires, it is just as necessary to use a mechanical ventilating system in mines worked by this method as it is in those worked by the horizontal top-slice system. Mechanical ventilation not only improves working conditions in top-slice stopes but, if properly designed for the particular mine in question, permits an immediate attack to he made in the event of a mine fire with a greater degree of safety to the men. In the course of time, incline top-slicing will probably take the place of the horizontal top-slicing system just as the incline cut-and-fill has superseded the horizontal cut-and-fill stoping method. While incline top-slicing has been tried in several districts on a very limited scale with inconclusive results, the Coronado is the-only large mine at which it has yet been thoroughly tested, and there it proved to be a complete success. As Mr. Scott's figures show, the tonnage has been more than doubled, with a simultaneous reduction in costs. As has been the history with all other innovations in mining, considerable hesitation will probably be shown in introducing this method in other mines. For example, when the incline cut-and-fill system was put forward as a substitute for the horizontal cut-and-fill stoping method many mining men of good repute were doubtful of its success and post-poned its introduction as long as possible, giving as their reason certain geologic conditions which would make the incline cut-and-fill impracticable. At the present time, most of the mining companies in the South-west, who in the past had used the horizontal cut-and-fill system, have already changed over, or are in the process of changing, and have proved that in the majority of cases horizontal cut-and-fill stopes can be successfully operated under the incline cut-and-fill system. Large horizontal cut-and-fill stopes can now be found in only a few exceptional cases, and

Jan 3, 1918

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 proposed' 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 others.2 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 pressure 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 C. A. Siebert, G. Sandoz

Oil quenching of steel, in good commercial practice, is carried out using a great deal of agitation in the quenching bath. Many of the tests reported in the literature represent the results obtained on steels quenched into still oil. Scott1 describes three stages of cooling in liquid quenching baths, namely: "State 'A'— Vapor Blanket Cooling—In this, the first stage, a thin stable vapor film surrounds the hot metal. Cooling is by conduction and radiation through the gaseous film and is therefore relatively slow. "Stage 'B'—Vapor Transport Cooling— Termination of the 'A' stage is marked by wetting of the metal surface. Vapor forms copiously in bubbles and is carried away by gravity and convection currents to condense in cooler surrounding liquid. This is the fastest stage of cooling. "Stage 'C'—Liquid Cooling—The 'B' stage ends as the surface temperature of the metal approaches the boiling point of the quenching liquid. Vapor no longer forms, so cooling is by conduction and convection and the temperature difference is greatly decreased. This mode of cooling is slower than that of the 'B' stage." Mechanical agitation of the quenching medium will accelerate all three stages of cooling. The greatest effect, however, will be noted on stage ' B' because the removal of the bubbles from the surface of the steel will not be entirely dependent upon gravity and natural convection currents. The method employed in this investigation for evaluating the oils was the end quench test. The various oils were substituted for water as the quenching liquid and hardness determinations were made at various distances from the quenched end. Most of the tests were run using a 3/8-in. orifice with a 21/2-in. stream height. However, to check the effect of other velocities two additional conditions were used, namely: (I) A 3/8-in. orifice and I-in. stream height, and (2) a 1/2-in. orifice and a 31/2-in. stream height. In all cases the samples were suspended so that the quenched end was 1/2-in. above the orifice. Some tests were run on an NE 8740, .3,312, and a 4340 steel. The results obtained on these steels confirmed those obtained on the NE 9450 and have therefore been Omitted from this paper. The test bars were machined from 1/8-in. stock to I-in. diam by 4 in. long.

Jan 1, 1949

By G. Sandoz, C. A. Siebert

Oil quenching of steel, in good commercial practice, is carried out using a great deal of agitation in the quenching bath. Many of the tests reported in the literature represent the results obtained on steels quenched into still oil. Scott1 describes three stages of cooling in liquid quenching baths, namely: "State 'A'— Vapor Blanket Cooling—In this, the first stage, a thin stable vapor film surrounds the hot metal. Cooling is by conduction and radiation through the gaseous film and is therefore relatively slow. "Stage 'B'—Vapor Transport Cooling— Termination of the 'A' stage is marked by wetting of the metal surface. Vapor forms copiously in bubbles and is carried away by gravity and convection currents to condense in cooler surrounding liquid. This is the fastest stage of cooling. "Stage 'C'—Liquid Cooling—The 'B' stage ends as the surface temperature of the metal approaches the boiling point of the quenching liquid. Vapor no longer forms, so cooling is by conduction and convection and the temperature difference is greatly decreased. This mode of cooling is slower than that of the 'B' stage." Mechanical agitation of the quenching medium will accelerate all three stages of cooling. The greatest effect, however, will be noted on stage ' B' because the removal of the bubbles from the surface of the steel will not be entirely dependent upon gravity and natural convection currents. The method employed in this investigation for evaluating the oils was the end quench test. The various oils were substituted for water as the quenching liquid and hardness determinations were made at various distances from the quenched end. Most of the tests were run using a 3/8-in. orifice with a 21/2-in. stream height. However, to check the effect of other velocities two additional conditions were used, namely: (I) A 3/8-in. orifice and I-in. stream height, and (2) a 1/2-in. orifice and a 31/2-in. stream height. In all cases the samples were suspended so that the quenched end was 1/2-in. above the orifice. Some tests were run on an NE 8740, .3,312, and a 4340 steel. The results obtained on these steels confirmed those obtained on the NE 9450 and have therefore been Omitted from this paper. The test bars were machined from 1/8-in. stock to I-in. diam by 4 in. long.

Jan 1, 1949

By G. Sandoz, C. A. Siebert

OIL quenching of steel, in good commercial practice, is carried out using a great deal of agitation in the quenching bath. Many of the tests reported in the literature represent the results obtained on steels quenched into still oil. Scott1 describes three stages of cooling in liquid quenching baths, namely: "State `A'-Vapor Blanket Cooling-In this, the first stage, a thin stable vapor film surrounds the hot metal. Cooling is by conduction and radiation through the gaseous film and is therefore relatively slow. "Stage `B'-Vapor Transport Cooling- Termination of the `A' stage is marked by wetting of the metal surface. Vapor forms copiously in bubbles and is carried away by gravity and convection currents to condense in cooler surrounding liquid. This is the fastest stage of cooling. "Stage `C'-Liquid Cooling--The `B' stage ends as the surface temperature of the metal approaches the boiling point of the quenching liquid. Vapor no longer forms, so cooling is by conduction and convection and the temperature difference is greatly decreased. This mode of cooling is slower than that of the `B' stage." Mechanical agitation of the quenching medium will accelerate all three stages of cooling. The greatest effect, however, will be noted on stage `B' because the removal of the bubbles from the surface of the steel will not be entirely dependent upon gravity and natural convection currents. The method employed in this investigation for evaluating the oils was the end quench test. The various oils were substituted for water as the quenching liquid and hardness determinations were made at various distances from the quenched end. Most of the tests were run using a 3/8-in. orifice with a 2 ½ -in. stream height. However, to check the effect of other velocities two additional conditions were used, namely: (I) A 3/8-in. orifice and 1-in. stream height, and (2) a ½ -in. orifice and a 3 ½ -in. stream height. In all cases the samples were suspended so that the quenched end was ½ -in. above the orifice. The steel used was an NE 9450 having the following analysis: [ ] Some tests were run on an NE 8740, 3312, and a 4340 steel. The results obtained on these steels confirmed those obtained on the NE 9450 and have therefore been omitted' from this paper. The test bars were machined from 1 1/8 -in. stock to 1-in. diam by 4 in. long.

Jan 1, 1948

By A. E. Powers, J. F. Libsch

THE simultaneous influence of time and temperature upon the tempering process in steel is well known. Hollomon and Jaffe¹ have expressed the effect of time and temperature upon the progress of tempering by a single parameter, T (c + log t); where T is the absolute temperature; t, the tempering time; and c, a constant characteristic of the steel and treatment. It was found that when hardness data for plain carbon and low alloy steels were plotted vs. this parameter, a master tempering curve was obtained which satisfied all the data for a given steel and a given hardening treatment. Roberts, Grobe, and Moersch² successfully applied the tempering parameter to various tool steels, including high speed steel for tempering times ranging from 6 min to 100 hr. The influence of tempering time and temperature upon the physical properties of high speed steel has also been extensively investigated.'" Recently, the use of induction heating for very rapid tempering treatments has received attention. The authors% have shown that tempering treatments of only a few seconds duration may be successfully applied to plain carbon and low alloy steel. However, the influence of such short tempering treatments in high speed steel has not been explored. Since high speed steel is a highly alloyed steel exhibiting physical reactions during tempering that are often relatively slow and large in magnitude, it is an interesting steel for the study of tempering reactions. It was the purpose of the present investigation, therefore, to study the early progress of reactions leading to secondary hardness in high speed steel, type 6-5-4-2, and to determine if induction-tempering treatments of a few seconds duration are able to produce results equal to those of more conventional long-time tempering cycles. Experimental Details The chemical compositions for two different heats of W-Mo high speed steel used in this study are given in Table I. Table II summarizes the heat treatment given various groups of specimens cut 21/2 in. long and centerless ground to ¼ in. diam. Hardening treatments consisted of preheating to 1600°F (870°C) for V2 hr and then oil quenching from 2250°F (1230°C). Heat I was subjected to the high temperature in a salt bath for 3 min, while specimens of heat II were austenitized in a carbon-block muffle furnace for 10 min. Specimens of series E were immediately refrigerated to — 318°F (— 195°C) in liquid nitrogen following the hardening treatment to transform the major portion of the retained aus-tenite. Specimens of series F and G were subzero cooled to — 100°F (-75°C) in a dry ice-alcohol mixture. Induction tempering was performed with a Lepel high frequency converter operating at a frequency of approximately 300,000 cycles per sec. Heating cycles were controlled and recorded for each specimen using a chromel-alumel thermocouple, percussion welded to the surface of the specimen, and a modified high speed instrument with a timer previously used and described.' Temperature oscillation of the specimens, while being held at the tempering temperature, was ± 15°F (8°C). Tempering Treatments: After hardening, specimens were tempered between 900" and 1500°F (480" and 815°C) to study the progress of retained austenite decomposition and secondary hardening. The specimens tempered in salt for 1 hr (series C, G, and I) were considered to represent conventional treatment and are used for comparison with rapid tempering treatments. Several induction-tempering cycles were employed: 1—flash tempering, i.e., induction heating to the tempering temperature and immediately cooling (series A, E, and F); 2—induction heating for 60 sec at temperature (series B); and 3—double-induction tempering by two successive heating and cooling cycles (series D). All specimens tempered by induction were heated at an initial rate of 450°F (250°C) per sec except those of series E, which were heated at an initial rate of 1300 °F

Jan 1, 1952

By A. J. Lyon

The stability and permanence of any structural material used in aircraft are of paramount importance. The spontaneous hardening, or age-hardening, which takes place in some of the aluminum alloys under ordinary conditions of temperature and atmosphere, is accompanied by an increase in tensile strength and proportional limit and a decrease in elongation or ductility. Within the last few years the introduction of a number of strong and ductile casting alloys led to a study by the Materiel Division, Air Corps, of the permanence of the physical properties in several of these alloys. It has been found that even though some alloys have good properties directly after casting or heat treatment, they are not satisfactory for stressed parts that require a combination of strength and ductility. It has also led to the conclusion that all alloys that can be hardened by artificial aging—heating in a temperature range below 170" C. (340" F.)—will harden spontaneously if allowed to stand for a sufficient length of time. It follows that, in aircraft work where the factors of safety are limited on account of the necessity of keeping the weight of the structures as low as possible consistent with their safe operation, the properties of the alloys in the stable or final condition should be used as a basis for design rather than the properties obtained on the material directly after casting. The work of Wilm in Germany, Hanson and Gayler at the National Physical Laboratory in England, and Merica, Jeffries, and others in the United States, has shown that the spontaneous hardening is characteristic to a greater or less extent in all aluminum-base alloys in which the solubility in the solid state of the alloying constituents decreases with the temperature. It has been shown that, by heating to a temperature as near as possible to the melting point of the material and soaking at that temperature for a sufficient length of time and then quenching or cooling rapidly, a maximum amount of the soluble constituents can be retained in solution and the condition of maximum hardening effect upon aging can be obtained. Also, in ordinary sand and chill castings, without heat treatment, a sufficient amount of the alloying constituents is held in solution to cause important changes in the final physical properties.

Jan 1, 1929

By A. J. Lyon

The stability and permanence of any structural material used in aircraft are of paramount importance. The spontaneous hardening, or age-hardening, which takes place in some of the aluminum alloys under ordinary conditions of temperature and atmosphere, is accompanied by an increase in tensile strength and proportional limit and a decrease in elongation or ductility. Within the last few years the introduction of a number of strong and ductile casting alloys led to a study by the Materiel Division, Air Corps, of the permanence of the physical properties in several of these alloys. It has been found that even though some alloys have good properties directly after casting or heat treatment, they are not satisfactory for stressed parts that require a combination of strength and ductility. It has also led to the conclusion that all alloys that can be hardened by artificial aging—heating in a temperature range below 170" C. (340" F.)—will harden spontaneously if allowed to stand for a sufficient length of time. It follows that, in aircraft work where the factors of safety are limited on account of the necessity of keeping the weight of the structures as low as possible consistent with their safe operation, the properties of the alloys in the stable or final condition should be used as a basis for design rather than the properties obtained on the material directly after casting. The work of Wilm in Germany, Hanson and Gayler at the National Physical Laboratory in England, and Merica, Jeffries, and others in the United States, has shown that the spontaneous hardening is characteristic to a greater or less extent in all aluminum-base alloys in which the solubility in the solid state of the alloying constituents decreases with the temperature. It has been shown that, by heating to a temperature as near as possible to the melting point of the material and soaking at that temperature for a sufficient length of time and then quenching or cooling rapidly, a maximum amount of the soluble constituents can be retained in solution and the condition of maximum hardening effect upon aging can be obtained. Also, in ordinary sand and chill castings, without heat treatment, a sufficient amount of the alloying constituents is held in solution to cause important changes in the final physical properties.

Jan 1, 1929

By A. J. Lyon

THE stability and permanence of any structural material used in aircraft are of paramount importance. The spontaneous hardening, or age-hardening; which takes place in some of the aluminum alloys under ordinary conditions of temperature and atmosphere, is accompanied by an increase in tensile strength and proportional limit and a decrease in elongation or ductility. Within the last few years the introduction of a number of strong and ductile casting alloys led to a study by the Materiel Division, Air Corps, of the permanence of the physical properties in several of these alloys. It has been found that even though some alloys have good properties directly after casting or heat treatment, they are not satisfactory for stressed parts that require a combination of strength and ductility. It has also led to the conclusion that all alloys that can be hardened by artificial aging-heating in a temperature range below 170° C. (340° F.)-will harden spontaneously if allowed to stand for a sufficient length of time. It follows that, in aircraft work where the factors of safety are limited on account of the necessity of keeping the weight of the structures as low as possible consistent with their safe operation, the properties of the alloys in the stable or final condition should be used as a basis for design rather than the properties obtained on the material directly after casting. The work of Wilm in Germany, Hanson and Gayler at the National Physical Laboratory in England, and Merica, Jeffries, and others in the United States, has shown that the spontaneous hardening is characteristic to a greater or less extent in all aluminum-base alloys in which the solubility in the solid state of the alloying constituents decreases with the temperature. It has been shown that, by heating to a temperature as near as possible to the melting point of the material and soaking at that temperature for a sufficient length of time and then quenching or cooling rapidly, a maximum amount of the soluble constituents can be retained in solution and the condition of maximum hardening effect upon aging can be obtained. Also, in ordinary sand and chill castings, without heat treatment, a sufficient amount of the alloying constituents is held in solution to cause important changes in the final physical properties.

Jan 1, 1928

By H. C. Chao, Y. E. Smith, L. H. Van Vlack

ThE phase relationships for the MnO-MnS system have been investigated only in the eutectic region. wentrupl reported a eutectic at 1280°C (2345°F) with approximately 50 wt pct of each component, as based on the earlier work of Andrew, Maddocks, and Fowler.2 More recently silverman3 corrected these figures to 2275° ±10°F (1246°C) with 63 wt pct MnS and 37 pct MnO. Neither Wentrup nor Silverman investigated the solid-solution limits in this system. However, Wentrup did suggest with his phase diagram sketch that the maximum solubility of MnS in MnO is 10 wt pct, and the solubility of MnO in MnS is nearer 20 wt pct, Fig. 1. These figures have been utilized in subsequent publications without further verification.4 There had been evidence found in this laboratory that the above solid-solution figures were high. An investigation was therefore conducted using the following procedure. Samples containing MnO and MnS were contained in platinum foil, heated just above the eutectic temperature in a purified nitrogen at- mosphere, and quenched into distilled water. The MnO had been prepared by calcining reagent grade MnCO3 in nitrogen at 1400°C. Only MnO was detected in subsequent X-ray diffraction analyses. The MnS had been prepared by sulfur deoxidation of reagent grade MnSO4 according to the reaction: MnSO4 + S2 ? MnS + 2SO2 The details of the sulfide preparation procedure, which utilized atmospheric sulfur pressures at 950°C, are discussed elsewhere.5 The sulfur content of the sulfide was determined to be 36.8 * 0.2 pct, which matches the stoichiometric value of 36.9 pct. No evidence of either MnO or MnSO4 was detected microscopically or by X-ray diffraction. The single-phase combinations of MnO and MnS at the eutectic temperature indicate the limit of solid solution as shown in Fig. 2(a) and (c), and the revised phase diagram is shown in Fig. 2(d). The melting temperature for MnO is that provided by Glasser.6 The eutectic temperature of 1232° ± 5°C and eutectic composition of 64 wt pct MnS are in close agreement with Silverman's correction of Wentrup's work. The maximum solubility of MnO in MnS was determined to be 1.7 ± 0.2 wt pct. The maximum solubility of MnS in MnO was determined to be 1.8 ± 0.2 wt pct. These figures represent the lower limit for the presence of liquid at 1260°C and are significantly lower than the earlier estimated solid-solution figures. The authors wish to acknowledge the support of U.S. Steel's Edgar C. Bain Laboratory for Fundamental Research. The comments and critical suggestions of Drs. A. S. Keh, R. Rickett, and R. B. Snow are appreciated.

Jan 1, 1963

APPROXIMATELY 3500 people thronged the Hotel Statler from February 18 to 24 for the 173rd general meeting of AIME. It was a technical extravaganza in that 82 technical sessions were held, at which 500 authors presented some 350 papers. Fortunately, Jerry Peirce, outgoing President, didn't have, to preside at all the events although after his year's stint as President, he has more than a nodding acquaintance with some of the latest problems in the mining, metals, and petroleum industries. Although the Welcoming Luncheon on Monday noon was the official opening of the meeting, Local Section Delegates and the Board of Directors convened on Saturday and Sunday, respectively. The large attendance and business like demeanor of these two groups was in accord with the energetic and competent service they have rendered the Institute through the year. To the applause of engineers and their wives who thronged the grand ballroom of the Statler for the Welcoming Luncheon, President Peirce introduced Oliver B. Hopkins, president of the Canadian Institute of Mining and Metallurgy and C. Gerow, secretary of CIM. Dr. Hopkins brought greetings from the CIM and expressed his satisfaction in the splendid spirit of cooperation that exists between the two Institutes. Dr. Hopkins is president, Interprovincial Pipe Line Co., Toronto, and counts as a close personal friend M. L. Haider, who is the 1952 President of AIME and lately of Toronto where he was a director and general manager of Imperial Oil Ltd. Mr. Haider will be known shortly as Mike to thousands of AIME members as he makes his Local Sections presidential tour. There was a pause in luncheon proceedings when Judge Harold R. Medina arrived amid the applause of the audience and somewhat breathless from his trip via motorcade from an antitrust trial against investment bankers at the courthouse downtown. With the Judge at his elbow, President Peirce took the opportunity in his introduction to good humoredly chide the Judge for the vixenish habits of public servants in their treatment of astute businessmen. The successful inventor who patents his inventions is a monopoly; the copper manufacturer who buys his competitors a drink is in collusion; but the business failure need only look to the government for a subsidy, Mr. Peirce pointed out. Judge Medina, quickly captured the admiration of his listeners. He spoke about justice, deploring partisan deals on nominations for judge-

Jan 1, 1952