Search Documents

By A. W. Gauger

For perhaps one hundred years scientists have been engaged in attempts at devising a satisfactory method of classification of coal. During this time many charts, tables and graphs have been proposed, some of which have had considerable merit, but no one system has been found to be applicable to all ranks of coal, nor has any one system received general acceptance. In many of the proposals, water eliminated between room temperature and 105" C. has been regarded as extraneous matter and calculations have been made to the air-dry or water-free basis. Such methods of calculation do not lend themselves to satisfactory grouping of the high-moisture coals.1, 2 A consideration not only of the properties of the high-moisture coals and their behavior on drying but of the generally accepted ideas of the coalification process and the role of water in this process indicates that water is an essential constituent of all coals and therefore should not be neglected. For example, in such low-rank coals as the lignites of the Northern Great Plains' deposits the inherent moisture of the lignite as mined plays a very important part in the physical properties thereof. Freshly mined lignite contains upwards of 30 per cent water and is extremely tough and resistant to shock. Upon exposure to the atmosphere, the lump loses part of its moisture and checks and splits, leaving a friable and broken up residue. So far as we have been able to determine, water is the only thing lost in this process and we are forced to the conclusion that moisture plays an extremely important part in the structural properties of the original lignite. With reference to the coalification process, Mack and Hulett3 pointed out in 1917 the important part that water played therein. They observed that water cannot be removed from peat or coal by pressure alone, but only as the result of chemical reaction, and suggested that peat in its early stages may be largely a hydrosol. Slow chemical changes occur over a long period of time until finally the mass becomes more largely a hydrogel. These views are substantiated to some extent by the experi-

Jan 1, 1932

By Oliver C. Ralston, James E. Norman

Separation of minerals by froth flotation is rightly called an art. It can truthfully be said that no two ores separate in the same way. The difference in results obtained when natural and synthetic mixtures of minerals are separated is so great that such tests frequently have been worthless in attempting to select conditions in which a natural mixture will separate. With a few rare exceptions, the natural mixture is much more difficult to treat than a synthetic mixture, and a much poorer separation is obtained. Perhaps several experimenters have wondered why 10 or 20 per cent of a mineral being floated from an ore does not float and is different from the 90 or 80 per cent that is recovered. In experimental tests on pure sphalerite, observers' have noted that varying amounts of sphalerite may be froth-floated with pine oil alone, the amount recovered in the froth depending apparently upon the source of the sulphide. The only explanation of results such as these, barring insufficiency of frother or time of flotation, is that the surface of the mineral floated is different from that of the mineral depressed. This difference often has been found to be caused by natural contamination of mineral surfaces, apparently brought about by weathering and exposure to natural solutions. There is virtually no rock that is absolutely unweathered. One of the writers had an exceedingly instructive experience in 1931 while working with United Verde copper-zinc iron sulphide ore in Arizona. Nonselective flotation of chalcopyrite and sphalerite away from some of the pyrite was all that could be accomplished. All known depressants produced only slight effect. Sodium cyanide showed some effect, but so much had to be used that only a small amount of chalcopyrite could be

Jan 1, 1939

By Oliver C. Ralston, James E. Norman

Separation of minerals by froth flotation is rightly called an art. It can truthfully be said that no two ores separate in the same way. The difference in results obtained when natural and synthetic mixtures of minerals are separated is so great that such tests frequently have been worthless in attempting to select conditions in which a natural mixture will separate. With a few rare exceptions, the natural mixture is much more difficult to treat than a synthetic mixture, and a much poorer separation is obtained. Perhaps several experimenters have wondered why 10 or 20 per cent of a mineral being floated from an ore does not float and is different from the 90 or 80 per cent that is recovered. In experimental tests on pure sphalerite, observers' have noted that varying amounts of sphalerite may be froth-floated with pine oil alone, the amount recovered in the froth depending apparently upon the source of the sulphide. The only explanation of results such as these, barring insufficiency of frother or time of flotation, is that the surface of the mineral floated is different from that of the mineral depressed. This difference often has been found to be caused by natural contamination of mineral surfaces, apparently brought about by weathering and exposure to natural solutions. There is virtually no rock that is absolutely unweathered. One of the writers had an exceedingly instructive experience in 1931 while working with United Verde copper-zinc iron sulphide ore in Arizona. Nonselective flotation of chalcopyrite and sphalerite away from some of the pyrite was all that could be accomplished. All known depressants produced only slight effect. Sodium cyanide showed some effect, but so much had to be used that only a small amount of chalcopyrite could be

Jan 1, 1939

By Cyril Stanley Smith

Somewhat over ten years ago the author described studies1'2 on the constitution of the copper-silicon system. The copper-rich portion of this diagram is shown in Fig. 1. The experimental points freely dotted about the diagram seemed to justify an otherwise unconvincing diagram and for several years the curious reverse bend in the alpha-phase boundary was accepted as substantially correct, even by later investigators3 of the system who introduced changes in the high-silicon phases. The binary diagram has been used as a base for several ternary systems5-9 without any suggested changes. Voce10 suspected the existence of another phase from his study of the cast alloys in the copper-manganese-silicon system, but evidently hesitated to extrapolate this to the binary diagram, particularly in view of the too ingenious explanation that was advanced1' to account for the pools of constituent clearly visible in castings; namely, that they were secondary alpha of twinned orientation deposited from beta. Von Schwarz,12 in 1932, called attention to the fact that this micro-constituent in copper-silicon alloy castings was anisotropic under polarized light and asserted that a new phase, not cubic, should be placed in the diagram .adjacent to alpha. Sautner,l3 on the basis of investigations with the polarizing microscope and with X-rays, introduced this phase into the diagram at about 6 per cent Si. The relevant part of his diagram is shown in Fig. 2. Sautner called the new phase beta, renaming the old beta phase zeta. To avoid confusion, in redrawing Fig. 2 the phases have been renamed in accordance with the terminology adopted in the present paper, in which the new phase is called kappa and the other phases retain the designation given them by Arrhenius and Westgren.4 The thermodynamic impossibility of Sautner's diagram and the paucity of experimental data left the exact location of the new phase field most uncertain, although his work proved beyond doubt that an additional noncubic phase exists. Because of the industrial importance

Jan 1, 1940

By Cyril Stanley Smith

Somewhat over ten years ago the author described studies1'2 on the constitution of the copper-silicon system. The copper-rich portion of this diagram is shown in Fig. 1. The experimental points freely dotted about the diagram seemed to justify an otherwise unconvincing diagram and for several years the curious reverse bend in the alpha-phase boundary was accepted as substantially correct, even by later investigators3 of the system who introduced changes in the high-silicon phases. The binary diagram has been used as a base for several ternary systems5-9 without any suggested changes. Voce10 suspected the existence of another phase from his study of the cast alloys in the copper-manganese-silicon system, but evidently hesitated to extrapolate this to the binary diagram, particularly in view of the too ingenious explanation that was advanced1' to account for the pools of constituent clearly visible in castings; namely, that they were secondary alpha of twinned orientation deposited from beta. Von Schwarz,12 in 1932, called attention to the fact that this micro-constituent in copper-silicon alloy castings was anisotropic under polarized light and asserted that a new phase, not cubic, should be placed in the diagram .adjacent to alpha. Sautner,l3 on the basis of investigations with the polarizing microscope and with X-rays, introduced this phase into the diagram at about 6 per cent Si. The relevant part of his diagram is shown in Fig. 2. Sautner called the new phase beta, renaming the old beta phase zeta. To avoid confusion, in redrawing Fig. 2 the phases have been renamed in accordance with the terminology adopted in the present paper, in which the new phase is called kappa and the other phases retain the designation given them by Arrhenius and Westgren.4 The thermodynamic impossibility of Sautner's diagram and the paucity of experimental data left the exact location of the new phase field most uncertain, although his work proved beyond doubt that an additional noncubic phase exists. Because of the industrial importance

Jan 1, 1940

By K. O. Stamm, B. A. Rogers

Because it is the basis of the economically important operation of hardening steel by quenching, the A3 transformation in iron has been the subject of numerous investigations. Although the contributions to this subject are interesting, no general review will be given here. Readers who care to follow the matter further are referred to the comprehensive summary by Cleaves and Thompson,l which covers the literature to about 1935. Their references to 53 papers on the subject suggest that further investigation is hardly justifiable. However, considerable interest has devel oped recently concerning the reversibility of the transformation. From thermo-dynamic considerations, the transition ought to occur at the same point on heating or cooling, if the change of temperature is sufficiently slow. Yet all of the earlier investigators found a difference of several degrees. Burgess and Crowe,2 who used thermal methods, obtained average temperatures of 912" for Ac3 and 900°C. for Ar3 on electrolytic iron. Their data indicate a considerable range of temperature over which the change occurs, both on heating and on cooling. Terry3 measured the susceptibility in the critical range and estimated from the steepest parts of his curves that Ac3 came at 9I8° and Ar3 at 903°C. Ishiwara4 was one of the first investigators to report a shorter interval between the temperature of the change in the two directions. He measured the change of susceptibility of kaolin-coated samples of steel, in a hydrogen atmosphere, at rater of heating and cooling so low that about 3 or 4 hr. were required for the transformation to complete itself. His results show some variability between samples, but indicate that the difference between Ac3 and Ar3 is in the vicinity of 6°C. Roberts and Davey5 measured the space lattice in pure iron in the vicinity of 900°C. and concluded that a change occurs reversibly between 907' 3" and 910" 3°C., probably between 907" and 910°C. Wells, Ackley, and Meh16 made dilatometric determinations of the Ac3 and Ar3 points in samples of hydrogen-treated carbonyl iron maintained under a vacuum. They used three rates of heating and cooling: 2°, ½°, and .?°C. per min. and found the spread between the two points decreased with low rates. At the slowest rate, the Ac3 transition occurred at 909° to 909. 5°C. and the Ar3 change at 907°C. in three of the samples believed to contain the lowest percentages of alien elements. They concluded that if the rate of temperature change and the content of alien eiements were low enough, the A3 transformation would be completely reversible at 909.50 + 1°C. The question of reversibility is an interesting one and when some iron with especially small percentages of alien elements became available the opportunity to conduct an investigation was welcomed. This iron was prepared by the Division of

Jan 1, 1942

By B. A. Rogers, K. O. Stamm

Because it is the basis of the economically important operation of hardening steel by quenching, the A3 transformation in iron has been the subject of numerous investigations. Although the contributions to this subject are interesting, no general review will be given here. Readers who care to follow the matter further are referred to the comprehensive summary by Cleaves and Thompson,l which covers the literature to about 1935. Their references to 53 papers on the subject suggest that further investigation is hardly justifiable. However, considerable interest has devel oped recently concerning the reversibility of the transformation. From thermo-dynamic considerations, the transition ought to occur at the same point on heating or cooling, if the change of temperature is sufficiently slow. Yet all of the earlier investigators found a difference of several degrees. Burgess and Crowe,2 who used thermal methods, obtained average temperatures of 912" for Ac3 and 900°C. for Ar3 on electrolytic iron. Their data indicate a considerable range of temperature over which the change occurs, both on heating and on cooling. Terry3 measured the susceptibility in the critical range and estimated from the steepest parts of his curves that Ac3 came at 9I8° and Ar3 at 903°C. Ishiwara4 was one of the first investigators to report a shorter interval between the temperature of the change in the two directions. He measured the change of susceptibility of kaolin-coated samples of steel, in a hydrogen atmosphere, at rater of heating and cooling so low that about 3 or 4 hr. were required for the transformation to complete itself. His results show some variability between samples, but indicate that the difference between Ac3 and Ar3 is in the vicinity of 6°C. Roberts and Davey5 measured the space lattice in pure iron in the vicinity of 900°C. and concluded that a change occurs reversibly between 907' 3" and 910" 3°C., probably between 907" and 910°C. Wells, Ackley, and Meh16 made dilatometric determinations of the Ac3 and Ar3 points in samples of hydrogen-treated carbonyl iron maintained under a vacuum. They used three rates of heating and cooling: 2°, ½°, and .?°C. per min. and found the spread between the two points decreased with low rates. At the slowest rate, the Ac3 transition occurred at 909° to 909. 5°C. and the Ar3 change at 907°C. in three of the samples believed to contain the lowest percentages of alien elements. They concluded that if the rate of temperature change and the content of alien eiements were low enough, the A3 transformation would be completely reversible at 909.50 + 1°C. The question of reversibility is an interesting one and when some iron with especially small percentages of alien elements became available the opportunity to conduct an investigation was welcomed. This iron was prepared by the Division of

Jan 1, 1942

By Walter Crafts, C. M. Offenhauer

IN the course of study of the heat-treatment of low-alloy steels, the behavior of alloy carbides at subcritical temperatures was found to vary from that indicated by published investigations. In order to establish the ranges of formation of carbides in low-chromium steels, a series of steels containing up to 0.60 per cent carbon and 7.5 per cent chromium has been investigated with respect to the type of carbide formed after quenching and tempering at sub-critical temperatures. The types of carbide formed as products of austenite transformation to pearlitic and pseudomartensitic structures have also been determined. There have been numerous investigations of the carbides in chromium steels and diagrams of the iron-chromium-carbon system.' Westgren, Phragrnén and Ne-gresco2 identified the following carbides in chromium steels: Cementitc............ Fe3C (containing up to about 15 per cent Cr) Trigonal carbide...... Cr7C3 (containing up to about 55 per cent Fe) Cubic carbide........ Cr4C (or Cr23C6) (con- . taining up to about 25 per cent Fe) Orthorhombic carbide , cr3c2 (containing only a few per cent Fe) The most complete information on the iron-chromium-carbon system was supplied by Tofaute, Sponheuer and Bennek3 and by Tofaute, Küttner and Büttinghaus,4 By means of dilatometric and microscopic stzdies, sections of the diagram were established. This diagram showed the presence of the special carbide CI7C3 in steels containing about 2.5 per cent to 12 per cent chromium in low-carbon steels and from about 5 to 20 per cent chromium in I per cent carbon alloys from A1 to room temperature. The present investigation has shown that a carbide inversion from Fe3C to Cr7C3 occurs in chromium steels containing more than about I per cent chromium after hardening and tempering at subcritical temperatures above about 500°C. The type of carbide observed to be more stable in quenched and tempered specimens was also found to be produced by direct austenite transformation in corresponding ranges of temperature and composition. Experimental Procedure The steels used in these experiments were made in magnesia-lined crucibles in high-frequency furnaces and were poured into 2-in. ingots. The ingots were forged into I-in. round bars, and all heat-treatments were conducted on specimens of this section size. Sections of the bar stock 8 in. long were hardened in air, oil or water, depending on the composition, after heating to a temperature sufficiently high to dissolve the carbides. The low-carbon and low-chromium steels were not entirely martensitic after quenching, and contained some ferrite and pseudomartensite. No pearlite was observed in the microstructures after quenching. The quenched bars were cut into

Jan 1, 1942

By Walter Crafts, C. M. Offenhauer

IN the course of study of the heat-treatment of low-alloy steels, the behavior of alloy carbides at subcritical temperatures was found to vary from that indicated by published investigations. In order to establish the ranges of formation of carbides in low-chromium steels, a series of steels containing up to 0.60 per cent carbon and 7.5 per cent chromium has been investigated with respect to the type of carbide formed after quenching and tempering at sub-critical temperatures. The types of carbide formed as products of austenite transformation to pearlitic and pseudomartensitic structures have also been determined. There have been numerous investigations of the carbides in chromium steels and diagrams of the iron-chromium-carbon system.' Westgren, Phragrnén and Ne-gresco2 identified the following carbides in chromium steels: Cementitc............ Fe3C (containing up to about 15 per cent Cr) Trigonal carbide...... Cr7C3 (containing up to about 55 per cent Fe) Cubic carbide........ Cr4C (or Cr23C6) (con- . taining up to about 25 per cent Fe) Orthorhombic carbide , cr3c2 (containing only a few per cent Fe) The most complete information on the iron-chromium-carbon system was supplied by Tofaute, Sponheuer and Bennek3 and by Tofaute, Küttner and Büttinghaus,4 By means of dilatometric and microscopic stzdies, sections of the diagram were established. This diagram showed the presence of the special carbide CI7C3 in steels containing about 2.5 per cent to 12 per cent chromium in low-carbon steels and from about 5 to 20 per cent chromium in I per cent carbon alloys from A1 to room temperature. The present investigation has shown that a carbide inversion from Fe3C to Cr7C3 occurs in chromium steels containing more than about I per cent chromium after hardening and tempering at subcritical temperatures above about 500°C. The type of carbide observed to be more stable in quenched and tempered specimens was also found to be produced by direct austenite transformation in corresponding ranges of temperature and composition. Experimental Procedure The steels used in these experiments were made in magnesia-lined crucibles in high-frequency furnaces and were poured into 2-in. ingots. The ingots were forged into I-in. round bars, and all heat-treatments were conducted on specimens of this section size. Sections of the bar stock 8 in. long were hardened in air, oil or water, depending on the composition, after heating to a temperature sufficiently high to dissolve the carbides. The low-carbon and low-chromium steels were not entirely martensitic after quenching, and contained some ferrite and pseudomartensite. No pearlite was observed in the microstructures after quenching. The quenched bars were cut into

Jan 1, 1942

By L. S. Darken

The manufacture and treatment of metals comprises operations whose effectiveness depends in large measure upon diffusion phenomena. The significance of such phenomena has, for a few simple cases, long been understood, but in more complex cases it has not been adequately recognized. This paper presents a mathematical analysis of some of the cases in which diffusion leads to the formation of a new phase within the metal, and this analysis provides a basis for interpreting quantitatively many apparently diverse experimental observations recorded in the literature, some of which are outlined in the following pages. OBservt1ons Recorded in Literature An outstanding instance is the general type of phenomenon in which diffusion of an element into an alloy is followed by a phase change, which most frequently is a reaction (precipitation) of the diffusing element with one or more components of the alloy to form a compound. Thus in certain alloys heated in contact with oxygen not only the usual oxide scale appears but, in addition, beneath the surface, a well-defined zone in which free oxide particles are dispersed with notabIe uniformity. This zone, commonly referred to as the subscale or zone of internal oxidation, is remarkable in that the visible boundary between it and the unaltered metal is usually sharp. Such a subscale, which appears to have been recognized first in oxidized copper alloys, has been described and studied by C. S. Smith,l by Rhines,2 and by Rhines, Johnson and Anderson;3 but it appears also in other alloys when they are exposed to conditions that bring about the precipitation of some compound, not necessarily an oxide. Another clear case of this same kind is the diffusion of carbon into a steel containing an element such as molybdenum, chromium, vanadium or tungsten, whose carbide is less soluble in iron than iron carbide, the result being that some part of that element is precipitated as a dispersed carbide. A similar phenomenon may occur on the welding of two alloys each of which contains one of a pair of elements that react to form a compound or a new phase of any sort. But in addition to the precipitation of a compound formed from the diffusing element, other effects may result merely from the presence of the diffusing element. For instance, if homogeneous austenite, nearly saturated with respect to carbon, is heated in contact with silicon in the absence of oxygen, one would expect to find, behind the several iron-silicon phases analogous to a scale, a zone of austenite containing dispersed carbide, because of the fact that silicon decreases the solubility of carbon in iron. As still another example, oxidation of an alloy may result in diffusion toward the surface of one or more of the constituents; for instance, during oxidation of a steel containing an element such as "copper, which is less prone to oxidation than iron, there is a notable enrichment of copper just beneath the scale-metal interface. AS an example of a yet more

Jan 1, 1942

By L. S. Darken

The manufacture and treatment of metals comprises operations whose effectiveness depends in large measure upon diffusion phenomena. The significance of such phenomena has, for a few simple cases, long been understood, but in more complex cases it has not been adequately recognized. This paper presents a mathematical analysis of some of the cases in which diffusion leads to the formation of a new phase within the metal, and this analysis provides a basis for interpreting quantitatively many apparently diverse experimental observations recorded in the literature, some of which are outlined in the following pages. OBservt1ons Recorded in Literature An outstanding instance is the general type of phenomenon in which diffusion of an element into an alloy is followed by a phase change, which most frequently is a reaction (precipitation) of the diffusing element with one or more components of the alloy to form a compound. Thus in certain alloys heated in contact with oxygen not only the usual oxide scale appears but, in addition, beneath the surface, a well-defined zone in which free oxide particles are dispersed with notabIe uniformity. This zone, commonly referred to as the subscale or zone of internal oxidation, is remarkable in that the visible boundary between it and the unaltered metal is usually sharp. Such a subscale, which appears to have been recognized first in oxidized copper alloys, has been described and studied by C. S. Smith,l by Rhines,2 and by Rhines, Johnson and Anderson;3 but it appears also in other alloys when they are exposed to conditions that bring about the precipitation of some compound, not necessarily an oxide. Another clear case of this same kind is the diffusion of carbon into a steel containing an element such as molybdenum, chromium, vanadium or tungsten, whose carbide is less soluble in iron than iron carbide, the result being that some part of that element is precipitated as a dispersed carbide. A similar phenomenon may occur on the welding of two alloys each of which contains one of a pair of elements that react to form a compound or a new phase of any sort. But in addition to the precipitation of a compound formed from the diffusing element, other effects may result merely from the presence of the diffusing element. For instance, if homogeneous austenite, nearly saturated with respect to carbon, is heated in contact with silicon in the absence of oxygen, one would expect to find, behind the several iron-silicon phases analogous to a scale, a zone of austenite containing dispersed carbide, because of the fact that silicon decreases the solubility of carbon in iron. As still another example, oxidation of an alloy may result in diffusion toward the surface of one or more of the constituents; for instance, during oxidation of a steel containing an element such as "copper, which is less prone to oxidation than iron, there is a notable enrichment of copper just beneath the scale-metal interface. AS an example of a yet more

Jan 1, 1942

By M. A. Grossman

The harden ability of most steels can be predicted within 10 to I5 per cent provided the complete chemical composition is known, including "incidental" elements; and provided the as-quenched grain size is known; and provided, finally, that the composition and heating temperatures for hardening are such as to result in austenite free from carbide particles. In the method proposed herein, the steel is considered as having a base hardenability due to its carbon content alone (the hardenability of a "pure steel" of the given carbon content, without any other elements), and this base hardenability is multiplied by a multiplying factor for each chemical element present. After multiplying all these together, the final product is the hardenability. Grain size may be taken into account either in the base hardenability or after the multiplication. Hardenability is stated in terms of "ideal critical diameter", namely, the diameter of bar, in inches, that will just harden all the way through (absence of un-hardened core) in an "ideal" quench, and the calculation may also be related to the Jominy test The data bring out certain features rather clearly. For example: i. It is quite useless to attempt to predict hardenability unless all elements, including "incidentals,,) are known; thus an "incidental,, chromium content of 0.20 per cent increases the hardenability by about 50 per cent. 2. The relative effectiveness of different alloys is sometimes unexpected; thus molybdenum when calculated in this way appears to be of the same order of effectiveness as manganese, rather than much more powerful, as would appear to be common experience; observe that an increase from no manaanese to 0.20 per cent Mn provides a multiplying factor of 1.67, and an increase from no molybdenum to 0.20 per cent Mo provides a factor of 1.63, an increase of over 60 per cent in each case. However, if to a steel containing 0.50 per cent Mn there is added "20 points of manganese," the factor is raised from 2.65 (for 0.50 per cent Mn) to 3.35 (for 0.70 per cent Mn), or an increase of only 26 per cent. Thus the fist small addition of an element has a much more powerful percentage effect than an equal further addition when some is already present, and in most cases the effect of molybdenum is considered in relation to a steel in which molybdenum is absent. 3. If two elements are equally effective, a greater hardenability will be obtained by using, for example, 0.5 per cent of each than by using I.o per cent of either of them alone. 4. A knowledge of the as-quenched grain size is essential for precise work, since a difference of only one grain size number (say No. 7 instead of No. 6) makes a difference of almost 10 per cent in hardenability (in these units); this, however, does not apply to certain It should be emphasized that in chromium steels (over 0.30 per cent Cr) and chrome-molybdenum and chrome-vanadium steels, un-dissolved carbides are likely to be present in the steel as quenched, and that in such cases the charts can indicate only a maximum Passible hardenability, whereas the extent of hardening actually obtained may be much less. Thus tests on a number of chrome-molybdenum steels have indicated a degree of hardening amounting to only 50 to 65 per cent of the maximum possible, and in chromium steels from full hardening (100 per cent) down to as low as 70 per cent. On the other hand, when the amount of such elements is small (Cr under 0.30 per cent, Mo up to 0.25 per cent in the

Jan 1, 1942

By M. A. Grossman

The harden ability of most steels can be predicted within 10 to I5 per cent provided the complete chemical composition is known, including "incidental" elements; and provided the as-quenched grain size is known; and provided, finally, that the composition and heating temperatures for hardening are such as to result in austenite free from carbide particles. In the method proposed herein, the steel is considered as having a base hardenability due to its carbon content alone (the hardenability of a "pure steel" of the given carbon content, without any other elements), and this base hardenability is multiplied by a multiplying factor for each chemical element present. After multiplying all these together, the final product is the hardenability. Grain size may be taken into account either in the base hardenability or after the multiplication. Hardenability is stated in terms of "ideal critical diameter", namely, the diameter of bar, in inches, that will just harden all the way through (absence of un-hardened core) in an "ideal" quench, and the calculation may also be related to the Jominy test The data bring out certain features rather clearly. For example: i. It is quite useless to attempt to predict hardenability unless all elements, including "incidentals,,) are known; thus an "incidental,, chromium content of 0.20 per cent increases the hardenability by about 50 per cent. 2. The relative effectiveness of different alloys is sometimes unexpected; thus molybdenum when calculated in this way appears to be of the same order of effectiveness as manganese, rather than much more powerful, as would appear to be common experience; observe that an increase from no manaanese to 0.20 per cent Mn provides a multiplying factor of 1.67, and an increase from no molybdenum to 0.20 per cent Mo provides a factor of 1.63, an increase of over 60 per cent in each case. However, if to a steel containing 0.50 per cent Mn there is added "20 points of manganese," the factor is raised from 2.65 (for 0.50 per cent Mn) to 3.35 (for 0.70 per cent Mn), or an increase of only 26 per cent. Thus the fist small addition of an element has a much more powerful percentage effect than an equal further addition when some is already present, and in most cases the effect of molybdenum is considered in relation to a steel in which molybdenum is absent. 3. If two elements are equally effective, a greater hardenability will be obtained by using, for example, 0.5 per cent of each than by using I.o per cent of either of them alone. 4. A knowledge of the as-quenched grain size is essential for precise work, since a difference of only one grain size number (say No. 7 instead of No. 6) makes a difference of almost 10 per cent in hardenability (in these units); this, however, does not apply to certain It should be emphasized that in chromium steels (over 0.30 per cent Cr) and chrome-molybdenum and chrome-vanadium steels, un-dissolved carbides are likely to be present in the steel as quenched, and that in such cases the charts can indicate only a maximum Passible hardenability, whereas the extent of hardening actually obtained may be much less. Thus tests on a number of chrome-molybdenum steels have indicated a degree of hardening amounting to only 50 to 65 per cent of the maximum possible, and in chromium steels from full hardening (100 per cent) down to as low as 70 per cent. On the other hand, when the amount of such elements is small (Cr under 0.30 per cent, Mo up to 0.25 per cent in the

Jan 1, 1942

By G. V. Smith, R. F. Mehl

The decomposition of austenite in steels, because of its immerse practical importance, has been subjected to extensive study in recent years from the point of view of the mechanism of the process.1"3 Although much information has been acquired, a number of questions, principally of theoretical interest, remain. One is that of the atomic-crystallographic mechanism by which each of the various reaction products forms. The isothermal reaction products of the eutectoid decomposition in plain carbon steels (to which this study is restricted) have conveniently been classified in three general groupings—pearlite, bainite, and martensite. Of these several products, pearlite and martensite have been studied extensively, but bainite has received only scant attention. -As a result, the mechanism of the formation of pearlite and that of martensite are now fairly well understood while a great deal of uncertainty exists regarding bainite. Pearlite Pearlite forms from austenite by the precipitation of alternate platelets of cementite and ferrite, which grow into the austenite matrix.2--5 The rate of reaction may be described in terms of a rate of nucleation and a rate of growth.2.3 Mehl and Smith6 have attempted to determine the lattice relationships existing between the ferrite in pearlite and the' parent austenite. Using a pole-figure method of analysis, it was found that a limited number of ferrite orientations resulted from the decomposition of a single grain of austenite; it was not possible, however, to decide between two assumed ideal relationships, both of which appeared to account for the experimentally determined pole figure. It was definitely shown, however, that the relationship obtaining was quite different from that determined for the proeutectoid separation of ferrite.7 No attempt was made to determine the habit plane (matrix plane to which the precipitate lies parallel) of the pearlite lamellae. Belaiew8 earlier had stated that the number of directions assumed by the pearlite lamellae in a single grain of austenite was 12 (no details were given) and concluded wrong!:!, as Mehl and Smith6 have pointed out, that the lamellae delineate the {2I0} planes of ferrite.* Martensite It is known that martensite is a transiti0nal constituent of body-centered tetragonal lattice,' and that it forms suddenly

Jan 1, 1942

By R. F. Mehl, G. V. Smith

The decomposition of austenite in steels, because of its immerse practical importance, has been subjected to extensive study in recent years from the point of view of the mechanism of the process.1"3 Although much information has been acquired, a number of questions, principally of theoretical interest, remain. One is that of the atomic-crystallographic mechanism by which each of the various reaction products forms. The isothermal reaction products of the eutectoid decomposition in plain carbon steels (to which this study is restricted) have conveniently been classified in three general groupings—pearlite, bainite, and martensite. Of these several products, pearlite and martensite have been studied extensively, but bainite has received only scant attention. -As a result, the mechanism of the formation of pearlite and that of martensite are now fairly well understood while a great deal of uncertainty exists regarding bainite. Pearlite Pearlite forms from austenite by the precipitation of alternate platelets of cementite and ferrite, which grow into the austenite matrix.2--5 The rate of reaction may be described in terms of a rate of nucleation and a rate of growth.2.3 Mehl and Smith6 have attempted to determine the lattice relationships existing between the ferrite in pearlite and the' parent austenite. Using a pole-figure method of analysis, it was found that a limited number of ferrite orientations resulted from the decomposition of a single grain of austenite; it was not possible, however, to decide between two assumed ideal relationships, both of which appeared to account for the experimentally determined pole figure. It was definitely shown, however, that the relationship obtaining was quite different from that determined for the proeutectoid separation of ferrite.7 No attempt was made to determine the habit plane (matrix plane to which the precipitate lies parallel) of the pearlite lamellae. Belaiew8 earlier had stated that the number of directions assumed by the pearlite lamellae in a single grain of austenite was 12 (no details were given) and concluded wrong!:!, as Mehl and Smith6 have pointed out, that the lamellae delineate the {2I0} planes of ferrite.* Martensite It is known that martensite is a transiti0nal constituent of body-centered tetragonal lattice,' and that it forms suddenly

Jan 1, 1942

By Robert A. Colten, Frederick C. Hull, Robert F. Mehl

It is known that pearlite forms from austenite by a process of nucleation and growth, and that the rate of formation of pearlite may be described by a rate of nucleation and a rate of growth.l,2 The manner in which these two quantities operate to produce the observed isothermal reaction curve has been analyzed in a paper published in I939.3 Two types of process were described: (I) that in which nucleation occurs at the grain boundaries of the parent phase, with the nodules growing toward the centers of the grains, initially as hemispheres and then after impingement as roughly radial-columnar structures; and (2) that in which nucleation occurs generally, with the nodules growing to full spheres until impingement with other growing nodules restricts growth. Both cases are met in the formation of pearlite from austenite (Fig. I);2,4 the former, occurring near the knee of the S-curve, has been named "grain-boundary transformation," and the latter, occurring at temperatures near Ael, has been named "group-nodule transformation." In developing our knowledge of the factors that determine the rate of formation of pearlite from austenite, it is important1,2 to obtain data on the rate of nucleation and the rate of growth of pearlite. Such data, when obtained near the knee of the S-curve, will furnish information of use in understanding the factors that affect hardenabil-ity; when obtained at higher temperatures, such data will be useful in understanding the behavior of steels that during heat-treatment react to pearlite.2 It is the purpose of this paper to present data on the rate of nucleation and the rate of growth of pearlite in dependence upon composition, grain size, temperature, and the degree of homogeneity of the parent austenite. The studies are restricted chiefly to plain carbon commercial steels." Previous Data on Rate of nucleation and rate of Growth Data on the rate of nucleation and the rate of growth of pearlite have been published by Scheil and Lange-Weise,5 by Mirkin and Blanter,6,7 by Meh1,l and by Dorn, De Garmo and Flanigan.2, It is difficult to interpret the data, for in most cases too little information is given on austenitizing temperatures and times, the state of homogeneity of the austenite (presence or absence of carbide, etc.), austenite grain size, and mode of formation of pearlite at the temperatures given. Except for the approximations of Mehl and the work of Dorn, austenite grain size was disregarded, though it is well known

Jan 1, 1942

By Robert A. Colten, Frederick C. Hull, Robert F. Mehl

It is known that pearlite forms from austenite by a process of nucleation and growth, and that the rate of formation of pearlite may be described by a rate of nucleation and a rate of growth.l,2 The manner in which these two quantities operate to produce the observed isothermal reaction curve has been analyzed in a paper published in I939.3 Two types of process were described: (I) that in which nucleation occurs at the grain boundaries of the parent phase, with the nodules growing toward the centers of the grains, initially as hemispheres and then after impingement as roughly radial-columnar structures; and (2) that in which nucleation occurs generally, with the nodules growing to full spheres until impingement with other growing nodules restricts growth. Both cases are met in the formation of pearlite from austenite (Fig. I);2,4 the former, occurring near the knee of the S-curve, has been named "grain-boundary transformation," and the latter, occurring at temperatures near Ael, has been named "group-nodule transformation." In developing our knowledge of the factors that determine the rate of formation of pearlite from austenite, it is important1,2 to obtain data on the rate of nucleation and the rate of growth of pearlite. Such data, when obtained near the knee of the S-curve, will furnish information of use in understanding the factors that affect hardenabil-ity; when obtained at higher temperatures, such data will be useful in understanding the behavior of steels that during heat-treatment react to pearlite.2 It is the purpose of this paper to present data on the rate of nucleation and the rate of growth of pearlite in dependence upon composition, grain size, temperature, and the degree of homogeneity of the parent austenite. The studies are restricted chiefly to plain carbon commercial steels." Previous Data on Rate of nucleation and rate of Growth Data on the rate of nucleation and the rate of growth of pearlite have been published by Scheil and Lange-Weise,5 by Mirkin and Blanter,6,7 by Meh1,l and by Dorn, De Garmo and Flanigan.2, It is difficult to interpret the data, for in most cases too little information is given on austenitizing temperatures and times, the state of homogeneity of the austenite (presence or absence of carbide, etc.), austenite grain size, and mode of formation of pearlite at the temperatures given. Except for the approximations of Mehl and the work of Dorn, austenite grain size was disregarded, though it is well known

Jan 1, 1942

By J. K. Stanley, R. F. Meel

The recrystallization of cold-worked metals is studied ordinarily by determining the temperatures required for complete recrystallization to occur within a given arbitrary time period, usually within 15 min. to 2 hr. The results obtained are conventionally assembled in a three-dimensional diagram relating this temperature to the amount of cold-work (percentage deformation) and the resultant grain size, though sometimes other modes of representation are emp1oyed.l Such three-dimensional diagrams are not complete, for the number of variables exceeds the number represented. The process of recrystallization is characterized by a time-rate, a fact that is not portrayed upon ordinary recrystallization diagrams. It is known that recrystallization consists in the formation of new nuclei and the growth of these nuclei, a process that continues until the matrix is completely consumed and the original set of grains entirely replaced by the new. This knowledge is old.2 Truly heterogeneous reactions invariably proceed by nucleation and growth; the decomposition of eutectoid solid solutions provides an example about which much is known,3 and freezing, originally studied by Tammann, provides the classic example. Although recrystal- lization is not a heterogeneous reaction within the orthodox meaning of the term, the mechanism of the process is the same. As in eutectoid inversion and freezing, it should be possible to study the rate of isothermal recrystallization as a function of time, and it should also be possible to analyze the isothermal reaction curve into the component rates of nucleation and growth that determine them.* Other methods, not employing the rates of nucleation and growth, have been used to represent the rates of such reactions, but these in general have been artificial and have not led to progress. 4,5 The theory of recrystallization is not in advanced state; if recrystallization data were obtained in terms of the rate of isothermal recrystallization, and in terms of the rate of nucleation and the rate of growth of recrystallized grains, and in particular if the temperature coefficients of these rates were measured accurately, an adequate theory should be more readily developed.27 Such an effort requires proper experimental technique, but especially requires suitable methods of mathematical analysis to convert mcasurements made upon a surface of polish into values for the rate of nucleation and the rate of growth and to combine these into a rate of isothermal recrystallization. These methods are now available for the case of three-dimensional recrystallization (recrystallization of

Jan 1, 1942

By R. F. Meel, J. K. Stanley

The recrystallization of cold-worked metals is studied ordinarily by determining the temperatures required for complete recrystallization to occur within a given arbitrary time period, usually within 15 min. to 2 hr. The results obtained are conventionally assembled in a three-dimensional diagram relating this temperature to the amount of cold-work (percentage deformation) and the resultant grain size, though sometimes other modes of representation are emp1oyed.l Such three-dimensional diagrams are not complete, for the number of variables exceeds the number represented. The process of recrystallization is characterized by a time-rate, a fact that is not portrayed upon ordinary recrystallization diagrams. It is known that recrystallization consists in the formation of new nuclei and the growth of these nuclei, a process that continues until the matrix is completely consumed and the original set of grains entirely replaced by the new. This knowledge is old.2 Truly heterogeneous reactions invariably proceed by nucleation and growth; the decomposition of eutectoid solid solutions provides an example about which much is known,3 and freezing, originally studied by Tammann, provides the classic example. Although recrystal- lization is not a heterogeneous reaction within the orthodox meaning of the term, the mechanism of the process is the same. As in eutectoid inversion and freezing, it should be possible to study the rate of isothermal recrystallization as a function of time, and it should also be possible to analyze the isothermal reaction curve into the component rates of nucleation and growth that determine them.* Other methods, not employing the rates of nucleation and growth, have been used to represent the rates of such reactions, but these in general have been artificial and have not led to progress. 4,5 The theory of recrystallization is not in advanced state; if recrystallization data were obtained in terms of the rate of isothermal recrystallization, and in terms of the rate of nucleation and the rate of growth of recrystallized grains, and in particular if the temperature coefficients of these rates were measured accurately, an adequate theory should be more readily developed.27 Such an effort requires proper experimental technique, but especially requires suitable methods of mathematical analysis to convert mcasurements made upon a surface of polish into values for the rate of nucleation and the rate of growth and to combine these into a rate of isothermal recrystallization. These methods are now available for the case of three-dimensional recrystallization (recrystallization of

Jan 1, 1942

By J. T. Eash, N. B. Pilling

As a group, the alloys of iron, nickel and carbon are, in application, one of the most versatile of the ferrous alloy family, and while many investigations have been made of their properties and structure only a few attempts have been made to show the microstructural relation over a wide range of composition. It is neccssary at once to be clear about the kind of structural relation with which we are concerned. The equilibrium diagram provides quantitative data regarding structural make-up and its variation with temperature, but equilibrium structures often are far different from the technologically significant ones. This is particularly true of steels and cast irons, and some of the most important cornponents of their structure may have no place in the diagram of equilibrium. The fugitive phases that result from transFormation are dealt with in a highly detailed and descriptive manner in. the S-curves originated by Bain. The present effort lies in a still different field, which is to portray the relation between sweeping changes in composition and in resultant structure, under controlled but nonequilib-rium conditions of thermal treatment. The Guilletl diagram presented in 1903 was the initial attempt at such a correlation and frequent references to this diagram can be found in the literature even though later work has shown it to require considerable modification. Kase2 published in 1925 a diagram that was based upon observations of furnace-cooled melts: At low carbon contents this diagram agreed fairly well with experience in showing pearlite to exist in alloys containing up to 6 per cent nickel, and completely austenitic steels at 24 to 28 per cent nickel. At high carbon contents, however, irons were indicated as being martensitic at around 2 per cent nickel and austenitic at 12 per cent nickel; these interpretations are open to serious question. Based upon experience with steels and cast iron, Wickenden3 suggested a diagram that agreed fairly well with Kase's in the steel range but which presented a wide departure from it in the higher-carbon, high-nickel regions. In its broad picture, the Wickenden diagram is an approach to "practical" structures. Reference may be made at this point to the efforts of Marsh5 to construct an iron-nickel-carbon equilibrium diagram by the coordination of scattered data, which are described at length in the monograph cited. None of these preceding attempts has appeared to meet the need for an exposition of the structural sequence that is encountered in traversing the nickel-steel nickel-cast-iron range under what might be called ordinary departures from equilibrium. ExperimentaT. Conditions Since the intent has been to establish structures developed under arbitrary conditions, experimental details have perhaps more than usual importance. Three of these conditions should particularly be noticed: (I) the material was cast and not subjected

Jan 1, 1942