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By J. W. Faust, H. F. John

Generalized conditions for rod or ribbon growth by the twin plane re-entrant edge mechanism are given. It was shown that this mechanism can result in growth of twinned platelets from dilute metal solutions. The platelrts 01 the fee metals are faceted by (111) planes and contain two or more twin planes. Microsegregation traces were revealed in silver grown from a thallium solution. VOGEL has already set down conditions for dendritic growth;1 however, this was the normal single-crystal dendritic growth experienced during solidification. In this work it is desired to set down the crystallographic conditions that will lead to rod or ribbon growth with smooth facets. These are as follows. a) The growth must be crystallographically limited (i.e., faceted growth). b) In order to have directed growth, some mechanism must allow easier nucleation than two-dimensional nucleation. That mechanism is afforded by the re-entrant edges at twin planes.213 Thus there must be twin planes present in the rods or ribbons. c) The growth direction must lie in the twin plane. If not, the twin plane will grow to the face or side of the ribbon and the site of easy nucleation will be lost. d) The growth direction must be parallel to at least one plane of the faceting planes, and no set of the faceting planes can be perpendicular to the growth direction. A corollary to these conditions involves the number of twin planes necessary to maintain the indestructible re-entrant edge. The number depends on the crystal structure of the material. An analysis of the data available shows that a material whose faceting planes form an open figure requires at least one twin plane while materials whose faceting planes form a closed figure require at least two twin planes in order to form the indestructible re-entrant edges for extended growth. The conditions and an assessment of the possibility of rod or ribbon growth are given in Table I for several crystal systems. As can be seen the fee, diamond and zinc blende, or-thorhombic, rhombohedral, and one form of the hexagonal lattice systems can yield this type of growth, but not the bcc metals nor the hexagonal

Jan 1, 1965

By N. E. Rogen, N. J. Grant

WHILE direct measurements of the growth of precipitate particles during aging is of fundamental value, few such measurements have been made. Direct measurements by means of optical or electron-microscopy and X-ray techniques are preferred, however, many of the aging studies have been based on indirect methods such as electrical resistivity and hardness. Because of the minute size of most precipitate particles (less than 1000A) the light microscope is of limited use while electron microscopy is subject to slowness, etching problems and consequent resolution problems. It is fortunate that alloys of the Nimonic type, based on Ni-Cr-Ti-Al, lend themselves to accurate and relatively easy particle size measurements. With proper balance of titanium and aluminum, the only aging precipitate is the gamma prime (r') Ni3,(Al, Ti)2 phase and this phase is the basis of the high-temperature strength of these allos.3-6 The Ni3(Al, Ti) phase is a face-centered-cubic, ordered structure with nickel atoms occupying the cube faces. Titanium atoms can replace three out of every five aluminum atoms and when substituting do so preferentially.2 Electron microscopy has shown that during aging of most of the commercial alloys of this type, precipitate particles are unusually uniform in size and distribution, and that the volume of (r') phase is large (20 pct).4,7 The aim of this investigation was to measure the growth of (y') phase during aging. In so doing, it was hoped that a clearer understanding of the aging process would be attained. X-ray line broadening was used as the measure of precipitate particle size. This technique was selected because it can measure and average the diameter of thousands of particles in one operation. The variations in hardness and matrix lattice parameter with aging were also studied and related to the particle size. EXPERIMENTAL PROCEDURE An alloy of the following composition was studied: Cr Ti A1 C Ni Wt pct 20.71 2.98 1.43 0.04 bal. At. pct 22.06 3.47 2.95 0.18 bal. The alloy, a 15- lb vacuum melt, was forged to 1/2-in. diam at a temperature of 1900" to 2200°F; was solution heat treated for 4 hr at 2000°F and water quenched; and was then aged in various ways. The ASTM grain size was 4 to 5. No (7') phase could be detected by X-ray analysis after solution treatment and electrolytic extraction. Aging was performed on 2-in. long specimens at 1400°, 1500°, and 1600°F for increasing times. One-half inch sections were removed after each time increment and were then electropolished in a solution of H3P04 saturated with chromium trioxide at a potential of 30 v and rectified with a tantalum electrode. Rockwell hardness measurements were taken using one side of the polished specimen. Matrix lattice parameter determinations were made on the other side using filtered chromium radiation references with the position of the (220) line of a silicon standard. To enable the precipitate particles of the (y') phase to be measured using X-ray line broadening techniques, it is first necessary to separate and concentrate this phase by electrolytic extraction. Extraction is necessary because the difference between matrix and precipitate lattice parameters is only approximately 0.5 pct and intensities of.the (y') phase lines are very weak relative to those of the matrix. The extraction techniques employed have been described before and found to produce a reasonably quantitative separation.398 A water solution of 15 pct H3Po4 is used at a current density of 0.5 amp per sq in. at 6 v. The precipitate particles are separated from the liquor by centrifuging and decanting and are then washed, dried and repowdered. A plot of the X-ray intensity vs diffraction angle for the (111) line of the precipitate particles was used as a measure of particle size. The intensity curve approximated the shape of Gaussian distribution function and its breadth was measured at the half height. The (220) line of a silicon standard was used to correct the measured line breadth for instrumental and other broadening errors. The Scherrer formula was used to calculate particle

Jan 1, 1961

By J. S. Kirkaldy, D. H. Weichert, G. R. Purdy

Two-phase diffusion couples have been used to simulate the growth of proeutectoid ferrite in ternary Fe-C-Mn austenites. It has been shown, theoretically and expermentally, that the results fall into two classes: a high-super saturation class in which the major role of manganese additions is to influence the boundary conditions for carbon diffusion; and a low-supersaturation class in which the partition of manganese is required. In the latter case, a drastic inhibition of the reaction rate is to be expected. As an aid to the kinetic analysis, a portion of the iron-rich comer of the Fe-C-Mn constitution diagvam has been determined. The conclusions drawn from this work have a direct bearing on the kinetics of the grain boundary ferrite precipitation reaction and thus yield insight into the manner in which alloying elements can influence the rate of austenite-decomposition reactions. THE role of alloying elements in the austenite-decomposition reactions has provoked considerable interest, since it is apparent that an understanding of this phenomenon will yield insight into the harden-ability problem. In the present investigation, attention is focused upon the proeutectoid-ferrite transformation in the system Fe-C-Mn, and upon the manner in which manganese additions affect the growth rate of ferrite. When hypoeutectoid austenite is supersaturated, the ferrite first formed usually exhibits two distinct types of morphology. At low super saturations, grain boundary layers grow into the austenite grains with an approximately planar interface, and usually possess at least one incoherent phase boundary. At high supersaturations, Widmanstiitten plates appear, which bear an orientation relationship' to the parent austenite. In this case, there is a chance of partial lattice matching along the sides of the plates and an attendant semicoherent interface structure. Purdy and Kirkaldy2 have used a two-phase binary Fe-C diffusion couple to demonstrate that carbon volume diffusion controls the incoherent reaction, while Rouze and rube,' using thermionic-emission microscopy, have shown that the thickening of Widmanstatten ferrite plates is appreciably slower than expected for carbon-diffusion control. This suggests that the plates are bounded by partially coherent, low-mobility interfaces. The redistribution of alloying elements during transformation has been the object of studies by Aaronson4 and owmman,5 who have shown that no alloy partition occurs during ferrite precipitation in the systems Fe-C-Cr and Fe-C-Mo, respectively. More recently, Aaronson et al.6 have shown that some manganese partition occurs at low super-saturations in the system Fe-C-Mn, but that no partition occurs at higher super saturations. Pickle-simer et al.7 observed that no manganese partition occurred during the early stages of the pearlite transformation in a manganese steel, and concluded that volume diffusion of the alloying element is not the rate-deter mining factor. The diffusion coefficients pertinent to this investigation have been measured by Wells, Batz, and Mehl8 (carbon in austenite), R. P. smith9 (carbon in ferrite), Wells and Mehl 10 (manganese in austenite), and Kirkaldy and Purdy11 (diffusion of carbon on a manganese gradient in austenite). Kurdjumov 12 has shown that the diffusivity of substitutional elements in ternary austenites up-quenched from martensite may be greatly enhanced, due, apparently, to the presence of a persistent defect structure. The recent work of Krauss 19 showing high densities of tangled dislocations in reversed Fe-Ni martensite is in agreement with this conception. THEORETICAL Diffusion in ternary or higher-order systems may be described with the aid of Onsager's extenbion of Fick's first law,14 in which the flux of each of the n components is assumed to be a linear function of all concentration gradients:

Jan 1, 1964

By H. R. Peiffer, R. Geckle

The silver phase of silver and finely divided alumina composites is shown to grow as single crystabs upon solidification from the melt. These crystals grow without the aid of an externally applied thermal gradient or a mold. The single-crystal growth is more readily disrupted as the alumina content varies from the optimum concentration of 20 wt pct. Pretreatment such as variation of fabrication pressure, and presintering of the specimens appear to play a lesser role in this phenomenon. In actual practice the growth of single crystals of metals from the melt is a tedious, painstaking process.' Care must be taken to insure the proper thermal gradients and rate of growth in a vibration-free environment. Excellent single crystals can be grown when these variables are controlled. When a second insoluble and finely divided phase is present in the melt the growth of single crystals is difficult, if not impossible, even under the most favorable conditions outlined above. During the past few years, a composite of silver and alumina has demonstrated an ability to resist flow even though the major phase, silver, is mo1ten. 2 A more surprising feature of this material is the ease with which single crystals of silver grow upon cooling the composite from temperatures above the melting point of silver to temperatures near room temperature. Observations have been made that demonstrate that single crystals of silver will grow even when conditions most unfavorable for growth of pure silver crystals exist. The crystals grown in this manner are imperfect in that a great deal of subgrain structure exists. It is the purpose of this manuscript to qualitatively describe this process and to attempt to isolate the influencing factors of this behavior. MATERIAL PREPARATION The composite material is prepared by powder-metallurgical techniques. Two items of great importance are the introduction of oxygen into the system and the obvious need for a thorough mixing of the metallic and refractory oxide phases. The oxygen is necessary in order that the silver wets the alumina. Without oxygen the phases separate and the material shows no unusual properties. In order to supply the needed oxygen the silver is added in the form of silver oxide which is readily decomposed to oxygen and silver upon heating. The silver oxide is readily obtained in a finely divided form that allows a uniform mixture of the alumina and silver oxide to be prepared. The correct ratio of alumina to silver oxide is mixed in the presence of alcohol in an ordinary food blender. The mixture has the consistency of a thin slurry and after mixing for approximately 20 min is dried and pressed into billets at various pressures. As is shown in Fig. 1, there is no limit to the shape of the specimens obtainable as single crystals. These billets are heated in air in a predeter-

Jan 1, 1964

By G. R. Speich, Morris Cohen

The growth rate of bainite has been determined by hot-stage metallography in five hypereutectoid high-purity iron-carbon, iron-carbon-chromium, and iron-carbon-nickel alloys. The studies have been confined essentially to lower bainite. The growth rate is independent of time at constant temperature, but both the edgewise and sidewise growth rates increase with increasing temperature. Increasing the carbon content, or adding chromium or nickel, decreases the growth rate without greatly affecting the temperature dependence. The results lead to separate diffusion-controlled models for the edgewise and sidewise growth rates. ALTHOUGH bainite has been subject of many detailed investigations,1,2 the growth mechanism of this intermediate transformation product of steel remains unclear. One difficulty has been that practically all the kinetic measurements have been restricted to studies of the overall reaction rate. Only a limited amount of information3, 4-6 is available on the growth process itself. The purpose of the present work was to investigate the effect of time, temperature and composition on the growth rate of bainite using hot-stage metallography. This method makes it possible to observe continuously the isothermal growth of individual bainitic units because of the surface upheavals accompanying their formation.3 This information is needed to test the existing growth models and to obtain a better understanding of the factors controlling the bainitic transformation. EXPERIMENTAL PROCEDURE Alloys—The compositions of the alloys used in the present work are shown in Table I. The five alloys were vacuum-melted and cast, and were prepared from high-purity materials. The three iron-carbon alloys were available from the work of Roberts et al.7 in the form of 3/8-in. diam rods.* To facilitate machining, the rods were ing. The rods were then tempered for 3 hr at 650°C to allow machining. The iron-carbon alloys in Table I were selected for investigating the effect of carbon on the bainitic growth rate. Due to the rapidity of the transformations involved, it was not feasible to measure the growth rates in iron-carbon alloys having carbon contents much less than 0.9 pct C. The iron-carbon-chromium and iron-carbon-nickel alloys were selected to study the influence of chromium and nickel on the growth rate at a carbon level of about 1.1 pct C, since this carbon content was intermediate among the iron-carbon alloys. The concentrations of chromium and nickel were chosen to match isothermal trans formation diagrams in the literature.8'9 Hot-Stage—Since the iron-carbon alloys had such low hardenabilities, it was necessary to quench the hot-stage specimens drastically from the austentiz-ing temperature to the isothermal transformation temperature to prevent the formation of pearlite or proeutectoid constituents. A rapid-quenching metal-lographic hot-stage was built for this purpose. A maximum temperature of about 1100°C and quenching rates up to 500°C per sec could be attained with this equipment. The details of the apparatus are described in a recent publication. After being polished and mounted on the hot-stage, the specimen (0.020 in. thick) was protected from oxidation and decarburization during austenitizing by a vacuum of 10-4 to 10-5 mm. The quench to the reaction temperature was accomplished with a helium blast, and the isothermal run at this temperature was conducted in a helium atmosphere. A fine-wire thermocouple spot-welded to the sample and a high-speed

Jan 1, 1961

By R. V. Sara

Determination of the Hf-C phase diagram was conducted primarily by metallographic and X-ray diffraction studies on appropriate alloys. The only intermediate phase observed in this binary system was HfC. This phase was found to be homogeneous between 34.0 and 48.0 at. pct C at 2200°c and between 36.0 and 49.3 at. pct C at 3150°C. The lattice-parameter variation was also determined for HfCI-, compositions prepared at 2200° and 3150°C. The most most refractory composition, with a melting point of 3830°c, was established at 47.5 at. pct C from melting-point data. Solidus temperatures of 2240' and 3150°C occur on the high-kafnium and high-carbon sides of the monocarbide, respectively. The invariant point between HfC and carbon is located at 66.0 at. pct C, whereas the 2240°C solidus corresponds to the peritectic temperature at which hafnium is formed from HfC and hafnium-rick liquid. Hafnium has a melting temperature of 2208°C and is capable of taking carbon into solution to the extent of 10.5 at. pct at this temperature. ALTHOUGH the Hf-C phase diagram has not been previously evaluated experimentally in its entirety, the belief has been that the general configuration would resemble the chemically similar Group rV carbide systems, Ti-C and Zr-c.' These binaries are characterized by a single carbide phase with a simple NaC1-type structure which is maintained over wide compositional ranges. This family of carbides has high thermal stability that increases substantially as the atomic number of the metal component increases. This trend characterizes HfC as one of the highest melting materials. According to Agte and Alterthum,2 the melting point for this monocarbide is 3890°C, a value that has been quoted quite extensively for the past several decades. Recently, in repeating the work of Agte and Alter-thum, Adams and Beall3 determined the melting point of HfC to be 3895°C. A significant departure from the commonly accepted version of the Hf-C system was reported by Avarbe and his coworkers,4 who proposed that there is an extreme stabilization of a hafnium in a narrow field to 2820°C, above which it melts peritectically to form HfC and liquid. Their study was also concerned with the melting temperature of various HfCI-, compositions, but the peak melting point was taken from the work of Agte and Alterthum. Avarbe and his associates were not concerned with the high-carbon region of the system. However, three widely varying temperatures have been reported for the HfC-C solidus by other investigators. Cotter and Kohn5 observed approximately 2800°C; Portnoy et al.' reported 3260°C; and, more recently, Krikorian7 indicated 2915°C. Equally as uncertain is the solidus between hafnium and HfC. As noted above, Avarbe et al. report a peritectic temperature of 2820°C, Krikorian7 measured 2150°C, and Benesovsky and Rudy' estimated 2000°C on their diagram. I) EXPERIMENTAL PROCEDURE The starting materials for this study consisted of reactor-grade hafnium hydride obtained from Fair-mount Chemical Co., Newark, N.J., hafnium carbide supplied by Wah Chang Corp., Albany, Ore., and Union Carbide spectroscopic-grade graphite, SP-1. The graphite analysis indicated impurities at levels of only 0.5 ppm or less. According to the suppliers, the hydride and carbide contained the typical impurities listed in Table I. The samples required for these studies were prepared by dry blending either graphite and hafnium hydride or hafnium hydride and hafnium carbide powders for approximately 5 min in a "Spex Mixer Mill". The latter combination was used only for preparing several samples in the HfC1-, melting-point studies. Small pellets, varying between 3/16

Jan 1, 1965

By J. L. Giove, J. M. Hodge, R. G. Storm

The hardenability effect of molybdenum has been evaluated by a number of investigators, including one of the present authors.1,2,3,4,6 Considerable discreMncy exists, however, among the results of these various investigators, and two of them, Brophy3 and Kramer,4 have indicated that the apparent hardenability effect of molybdenun would vary, depending upon the other alloying elements in the steels being considered. The results of previous, unpublished work conducted at the Duquesne Works Laboratory of the Carnegie-Illinois Steel Corporation also have indicated significant differences in the hardenability effect of molybdenu~n in nickel and in chromium steels. The present investigation was undertaken in an effort to establish the mechanism governing this difference in the hardenability effect of molybdenum. This work involved both end-quench hardenability and isothermal transformation tests on two series of steels of varying molybdenum content. One steel contained 3 pct nickel and the other 1 pct chromium. It was hoped that such studies would shed further light on the mechanism of this behavior. Materials and Experimental Work MATERIALS The chemical compositions of the steels used in this investigation arc shown in Table 1. These steels were furnished by Battelle Memorial Institutc in the form of 100-lb induction furnace ingots. The top half of each ingot was forged to 1 3/8-in. square billets and then rolled to I-in. round bar stock for isothermal transforma- tion studies. The bottom half of each ingot was forged into approximately 1 1/4-in. rounds for end quench tests. All of the steels were normalized from 1650°F and tempered at 1100°F before nlachming to end quench or isothermal samples. END-QUENCH TESTS standard l-in. diam end-quench tests were made on each of the steels. The test bars were austenitized at 1600°F for 30 min,, quenched in a standard Jominy jig, and hardness surveys were made in all tests. The 95 pct martensite point was deterInined from hardness values5 and cheeked by metallographic examination. The jominy distance values for 95 pct martensite were converted to hardellability values in terms of ideal diameter (Dl), using a revised carve for the relationship between Jominy distance and ideal diameter. This curve was based on the cooling time from the A1 to the bainite nose temperatures rather than on the usual criterion of half-temperature time. The cooling rates as determined by Russell and Williamson9 were used as a basis for this relationship. The derivation of this curve is described in detail in the appendix. ISOTHERMAL TRANSFORMATION STUDIES Samples for the isothermal transformation studies were quarter sectors of 1/16.is-in. thick slices from the 1-in. round bars. All samples were austenitized for a total the of 15 min., transferred rapidly to a lead bath for isothermal transformation, and brine quenched from the lead bath. A time of 3 see was allowed for the sample to come to the temperature of the lead bath, and this time was not included in the isothermal transformation times. Each sample was examined metallo-graphically at a magnification of 100 diam. The per cent transformation was estimated by comparison with standard micrographs, and the values were plotted against the on log/log coordinate paper. At least 4 and as Inany as 10. points were plotted for each steel. A line was drawn through the points for each steel, and the times for 5 pct transformation were obtained from the plot for each steel. Isothermal transformation diagrams in terms of 5 pet total transformation, covering the temperature range of 800 to 1200°F, were determined in this manner for an austenitizing temperature of 1600°F. The times for 5 pct transformation at the bainite nose for austenitizing temperatures of 1800 and 2000°F were also determined in the same manner. These times for 5 pct total transfor-

Jan 1, 1950

By H. J. Beattie, G. C. Gould

J. B. Seabrook' recently published properties of a low-alloy Ni-A1 age-hardening steel known commercially as "Nitralloy-N". He mentioned three possible mechanisms of age hardening, viz. order-disorder in the matrix, order -disorder in a precipitate, and a precipitate that goes in and out of solution reversibly without losing coherency. The following results indicate that the last - mentioned of these mechanisms is the cause of age hardening in Nitralloy-N. The reversibility of the age hardening with a sub-critical solution temperature was demonstrated strikingly by Seabrook, whose results are reproduced in Fig. 1. Aging for 22 hr at 975°F (525°C) can reharden the fully-tempered martensite to its original as-quenched hardness of R, 45. A Nitralloy steel was obtained from Seabrook' with the following analysis:

Jan 1, 1962

By C. Land, R. Hultgren

Heat contents of four Cr-Ni alloys were determined in a diphenyl ethev calorimeter. Aside from ferromagnetic effects, Kopp's law of additivity of heat capacities is approximately fol-lowed, in disagreement with earlier results of A. W. Foster. In 1934 Foster1 reported heat-capacity measurements of dilute solutions of chromium in nickel, primarily as a study of the effect of alloying on the ferromagnetic Curie temperature. He found that, apart from the Curie temperature effect, the general level of heat capacities was reduced 10 pct by the addition of 1 pct Cr; 20 pct by the addition of 2 pct Cr. His measurements of pure nickel essentially agreed with accepted values. These results indicate that small amounts of dissolved chromium cause an increase in the Debye temperature of nickel, presupposing a change of frequency of interatomic bonding vibrations which is much larger than should be expected. If Foster's results were verified, therefore, they would have important implications for bonding theory. The only other heat-capacity measurements in this system seem to be those of Taylor and Hinton,' who studied anomalies which they attributed to super-lattice formation at 25 at. pct Cr. At temperatures above the anomaly, Taylor and Hinton found heat capacities at this composition were quite normal, approximately obeying Kopp's law of additivity. EXPERIMENTAL WORK In order to verify or disprove the remarkable results of Foster, the heat contents of four dilute solid solutions of chromiunn in nickel were measured in a diphenyl ether calorimeter. This instrument, described elsewhere,, operates on the principle of the Bunsen ice calorimeter, but uses diphenyl ether as the working substance. Diphenyl ether melts at the convenient temperature of 26.S°C and supplies a more sensitive measure of heat than ice. The hot sample, in an atmosphere of argon, is dropped from a furnace at a measured temperature into a chamber which is surrounded by a mixture of solid and liquid diphenyl ether. The heat from the hot specimen melts some of the solid without chang- ing the temperature of the calorimeter. The amount of solid melted, and therefore the number of calories released by the specimen, is measured from the increase of volume of the mixture, which pushes mercury down a graduated capillary tube. The apparatus has been standardized against the known heat content of pure platinum. Heat loss during the dropping of the specimen was estimated by comparing the heat content of an empty platinum capsule with a solid platinum ball of the same radius and degree of polish of the surface. Repeated tests demonstrate the precision of measurement to be closer than *0.3 pct. Smoothed experimental results are listed in Table I. The tabulated values for pure nickel are taken from kelley, several measurements were made which agreed closely with this tabulation. For chromium, however, no similar agreement was found. The values tabulated were taken from a thorough study of chromium with the diphenyl ether calorimeter which will be published elsewhere.= Suitable plots of the data showed the expected ferromagnetic anomalies in the low-chromium alloys, while the 11 pct alloy shows no ferromagnetic anomaly, but a different, small anomaly below 900°K. This is perhaps the same anomaly found in CrNi, which was once attributed to superlattice formation,' but is probably due to some other cause, such as antifer-

Jan 1, 1960

By P. M. Robinson, M. B. Beaer

The heat of formation at 0°C of the intermetallic compound AgMg as a function of composition has been determined by tin-solution calorimetry. In this technique, the heat of formation is determined as the difference in the heat effects, at infinite dilution, on dissolution of the compound and a mixture of its components added alternately to liquid tin from a reference temperature, such as 0°C. The experimental technique and method of calculation have been described elsewhere.' The compound AgMg is a 0 electron compound with a CsCl structure.2 In nonstoichiometric compositions of the compound, atoms of the excess component are substituted on the AgMg lattice.3,4 The homogeneity range is approximately 43 to 52 at. pct Mg at room temperature and 38 to 65 at. pct Mg at 500Oc.2 Single-phase alloys outside the homogeneity range at room temperature can be prepared by quenching. The specimens used in this investigation contained from 39.0 to 54.8 at. pct Mg. They were made available, together with analyses, by the Research Laboratory of the General Electric Co. Metallographic examination of the quenched specimens containing 39.0 and 54.8 at. pct Mg gave no indication of the presence of a precipitated phase. The measured heats of formation are listed in Table I. In a calorimetric run, four to six additions each of a specimen and the corresponding mixture of the components were made. One or two calorimetric runs were carried out with each composition depending on the amount of samples available. The limits stated in the table represent the experimental precision and not the absolute accuracy of the measurements. Thermodynamic investigations of the compound AgMg have been made by Trzebiatowski and Terpilowski 5 and Kachi.6,7 The heats of formation at 500°C derived from the combined electromotive-force values of these investigators8 are more exothermic than the heats of formation at 0°C reported here, Fig. 1. The differences cannot be explained by the difference in temperature. This is shown by the published heat capacity of the stoichiometric compound, which obeys the Kopp-Neumann rule from 0" to 500oC.9 The heats of formation derived from equilibrium data are generally considered to be less accurate than those determined by calorimetric methods because small errors in the change of the equilibrium values with temperature have a large effect on the calculated value of the heat of formation. The heats of formation derived from the combined published electromotive-force values reach a maximum at 56.0 at. pct Mg.8 Kachi reports a maximum of -4.7 kcal per g-atom at the stoichiometric composition.7 The present results also indicate a maximum at or near the stoichio metric composition, Fig. 1.

Jan 1, 1964

By H. R. Gardner

The heat treatmmt of plutonium was studied using the Jominy end-quenching technique commonly used for determining the hardenability of steel. Plutonium specimens were end-guenched from temperatures in each of the ß, y, d, d') and E phases. One series of specimens with a low-iron content, 165 pprn Fe, and another gvoup with a high-iron content, 678 pprn Fe, were used in order to study the effect of a Pu-Pu,Fe eutectic network on the hardness and micro-structure. Hardness traverses indicated no significant variations in hardness with either cooling rate or quench temperature. Metallographic studies indicated major effects on microstructure. Grain size was found to vary markedly with quenching temperature and cooling rate. It was determined that the Pu-PueFe eutectic network could be modified extensively by heat treatment, including spheroi -dization in the y phase and 6 phase below 413°C. An unidentified spheroidal inclusion was observed to go into solution in the delta and higher temperature phases. 1 HE element plutonium is unique among metals in that it has six allotropic forms in the solid state.' These have been designated' as the a, ß, y, d, d', and phases. The respective phase transformations occur at approximate temperatures of 122", 210°, 319°, 450°, and 480°C with a melting point at 640°C. With this number of phase transformations it becomes pertinent to consider the effect of heat treatment and cooling rate in the various phases on the microstructure and hardness of plutonium. To determine; the effect of a wide range of cooling rates, the Jominy end-quenching technique was applied to cylindrical plutonium specimens. In addition, since the presence of iron in amounts greater than 500 pprn is common in plutonium and results in a network of the Pu-Pu6Fe eutectic, it was decided to study heat treatment effects on two bomb reduced plutonium buttons with different iron contents. A low iron button containing 165 ppm Fe and a high iron button containing 678 ppm Fe were selected for this comparison. EXPERIMENTAL PROCEDURE Experimental Material. The cylindrical bars for Jominy quenching were cast from button stock in vacuo in MgO coated graphite molds. Metal pouring temperature was approximately 950°C and the molds were preheated to 300°C. The castings were machined to 0.5 in. diam. by 2.5 in. long cylinders. Chemical analysis and density data for the two groups of Jominy specimens containing different iron contents are presented in Table I. Except for iron, heats 19-12-1 and 20-2-1 are comparable within the limits of analytical accuracy. Representative specimens from the low-iron and high-iron bars were taken for metallographic examination. The low iron specimen was found to have extensive microcracking, Fig. 1. In addition, numerous unidentified spheroidal inclusions were present, Fig. 17. The average grain size of the low-iron plutonium is 0.068 mm, Fig. 4. In the high-iron plutonium, a Pu-Pu6Fe eutectic network is prominent, Fig. 7. The average size of the network is 0.100 mm. Unidentified spheroidal inclusions were also common in the high-iron plutonium. The average grain size of the high-iron plutonium is 0.036 mm. Experimental Technique. A Jominy end-quenching fixture was fabricated for glove box use. A 2.5 in. water height was used with an orifice of 0.250 in. ID. The orifice to specimen distance was maintained at 0.5 in. Annealing temperatures of 160°, 265°, 400°, 465°, and 535°c were chosen for the study of the effect of quenching from ß, y, d, d', and e phases on microstructure and hardness. During quenching, cooling curves were obtained from the Jominy specimens at distances from the quenched end of 1/16, 1/8, 3/8, 3/4, 1-1/4, and 2 in. Cooling rate calculations were made from the cooling curves for temperature intervals of 100° to 110° 160" to 170°, 330" to 340°, 420° to 430°, and 490° to 500°c. These temperature intervals were chosen

Jan 1, 1962

By P. M. Robinson, M. B. Bever

THE heats of formation at 273°K of the compounds InBi, In2Bi, and TIBi2 have been determined by metal solution calorimetry with bismuth as solvent. The published information on the thermody-namic properties of these compounds is limited. The structure of the compound InBi (mp, 383°K) is tetragonal of the PbO (B10) type.1 The compounds In2Bi (mp, 362°K) and TIBi2 (mp of highest-melting composition, 486°K) are now believed to have structures of the NizIn (B8) type273 and not the A1B2 (C32) structures reported earlier.4, 5 The compounds In2Bi and TIBi2 have limited homogeneity ranges at room temperature, while the compound InBi has no detectable homogeneity range.' Samples of the compounds were prepared from 99.99 pet Bi (Mallinckrodt Chemical Works), 99.999 pet in (American Smelting and Refining Co.), and 99.99 pet Tl (Amend Drug and Chemical Co.) by melting the components in sealed, evacuated Vycor tubes. The compounds InBi and in2Bi were of the stoichiometric compositions. The compound TIBi2 contained 40 pet T1; this composition was chosen because at room temperature the alloy of stoichiometric composition does not lie within the homogeneity range.' The melts were held approximately 50°C above the liquidus for 8 hr and then quenched into iced brine. The samples were homogenized for 48 hr at temperatures 20°C below the temperature of complete solidification. Metallographic examination of sections taken from the center and the ends of the ingots did not reveal second phases or segregation . The heats of formation were determined as the difference in the heat effects on solution of a compound and of a mechanical mixture of its components added from the reference temperature of 273°K to liquid bismuth at 623°K. The details of the calori-metric procedure and method of calculation have been described elsewhere? The calorimeter was calibrated by adding bismuth at 273°K to the bismuth bath at 623°K. The calculated values of the heats of formation of the compounds, listed in Table I, are based on a value of 4.96 keal per g-atom for the difference between the heat contents of bismuth at 273" and 623oK7 No published values are available for the heats of formation of the compounds InBi and In2Bi. Hultgren et a1.7 have calculated a value of -0.50 i 0.10 keal per g-atom for the heat of formation at 423°K of the compound TIBi2 containing 40 at. pet T1 from heat-content measurements in the range from 398" to 700°K and the heats of formation of the liquid alloy. The difference between this value and the value of -0.66 * 0.01 keal per g-atom at 273°K is too large to be attributed to a difference between the heat capacities of the compound and the components over the temperature range 273o to 423oK7 However, the direct determination of the heat of formation by a calorimetric method should give a more accurate value. The heats of formation at 273°K of the compounds InBi and In2Bi lie on a straight line when plotted as a function of composition, Fig. 1. The compound In2Bi, therefore, appears to be barely stable with respect to its neighboring phases. In order for the free energy vs composition curve to be concave towards the composition axis, the entropy change on formation of In2Bi must be more positive than that on formation of InBi. As the compounds are ordered,2, 3, 8 there is no configurational entropy change on formation. The difference between the entropy changes on formation of the two compounds, therefore, is probably associated with the vibrational entropies. In view of the low heats of formation, the changes in vibrational entropy, due to changes in bond strength on formation of the two compounds, are likely to be small. Owing to the volume contractions on formation of the two compounds,5 the vibrational entropies probably decrease slightly but the decrease of the vibrational entropy of InBi is expected to be larger than that of In2Bi.

Jan 1, 1965

By A. H. Duane, A. E. Miller

The high-temperature allotropy of some heavy rare-earth metals and their alloying behavior with magnesium in the 0 to 50 at. pct Mg region was studied by thermal, microscopic, and X-ray methods. Examination of a hcc phase retained upon quenching alloys of different magnesium content confirmed the existence of a bcc high-temperature allotrope in pure gadolinium, terbium, dysprosium, holmium, erbium, thulium, and lutetium. The lattice constants of the pure metals were determined from a Vegard's Law extrapolation and .found to be 4.05, 4.02, 3.98, 3.96, 3.94, 3.92, and 3.90 + 0.02 A, respectively. In all of the systems studied, the high-temperature bcc phase decomposed eutec-toidally, when slowly cooled, into a hexagonal rare earth-Mg solid solution and a simple-cubic peritec-tic AB-type compound. In 1956, Spedding et al.' observed high-temperature resistivity anomalies in the light rare-earth metals and made several unsuccessful attempts to retain the high-temperature form of these materials at room temperature by quenching the pure metals from temperatures near their melting points. High-temperature X-ray techniques were used by Spedding et a1.' in 1959, to investigate the high-temperature allotropy of the rare-earth metals. This study showed that lanthanum, cerium, praseodymium, neodymium, ytterbium, and possibly gadolinium are bcc at temperatures near their melting points. Discontinuous changes in resistivity at elevated temperatures were also observed for pure gadolinium, terbium, dysprosium, holmium, and lutetium, thus indicating the possible existence of a high-temperature crystalline transformation in these metals. However, due to experimental difficulties encountered in the high-temperature X-ray analysis of these metals, their allotropic form was not identified. Spedding et a~.~ in 1960 observed that yttrium displayed a continuous solid solubility with lanthanum at elevated temperatures, thus showing the high-temperature structure of yttrium to be the same as that of lanthanum, i.e., bcc. Similarly, the high-temperature alloying behavior of gadolinium with yttrium showed gadolinium to be bcc at elevated temperatures. The bcc nature of the high-temperature form of yttrium was again recognized in 1960 by Eash and carlson4 in their study of the alloying behavior of yttrium with thorium. Gibson and Carlson, in a similar study of the alloying behavior of yttrium with magnesium, found that by quenching Y-Mg alloys from within the 0-solid solution region, Fig. l, they were able to retain the allotropic form of yttrium. It was the present authors' intention to use this method to show the existence of a bcc high-temperature form of the heavy rare-earth metals gadolinium, terbium, dysprosium, holmium, erbium, thulium, and lutetium. Simultaneous with the present work, Beaudry and Daane' have shown the bcc nature of the high-temperature form of scandium. EXPERIMENTAL PROCEDURES Materials. The rare-earth metals used in this study were prepared by the calcium reduction of the fluoride and were further purified by vacuum distillation. The magnesium was obtained from the New England Lime Co. and was also purified by distillation prior to its use in the preparation of the rare earth-Mg alloys. The results of chemical and spec-

Jan 1, 1964

By D. A. Kraai, S. Floreen

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

Jan 1, 1964

By A. H. Daane, J. F. Haefling

The limits of miscibility in some of the uranium rare-earth alloy systems have been determined in the temperature range 1000°to 1250°C. The solubilities of lanthanum and cerium in uranium are greater than those of the remaining rare earths by a factor of more than two. The solubility of uranium is greater in cerium, braseodymium, and neodymium than in the other rare-earth metals studied. The values found in this study are in qualitative agreement with those which might be expected if the solubility rules of Hildebrand and Scott are applicable. AS interest in nuclear reactors intensifies, many new types of fuels are being suggested in attempts to improve the economics of some of the proposed reactor schemes. To remove some of the difficulties inherent in the use of solid-fuel elements and their reprocessing, many types of liquid-metal reactors have been suggested. One of the more attractive features of several of these reactor concepts is that they include a continuous or semicontinuous process for the extraction of fission products and "bred" fissionable materials from the fuel, utilizing immiscible metal extractants. This would enable a much higher burn-up of fissionable material to be achieved and would present a very attractive economic picture. Several studies have been reported on equilibrium systems in which there exists a high degree of immiscibility between uranium and another metal that might be used as an extractant in such a processing scheme.' Two of these systems in which a high degree of immiscibility exists are those of uranium with the two rare-earth metals, lanthanum, and cerium. Since the rare earths constitute a significant fraction of the fission products, their removal is of prime importance. It is reasonable to believe that this might be accomplished by equilibrating a rare-earth phase with the contaminated uranium fuel in the liquid state. In order to make a more complete study of those systems which would be of interest either as extractants in a liquid-liquid extraction process, or as fission products formed in the fuel, the alloy systems of uranium with lanthanum, cerium, praseodymium, neodymium, and samarium were studied in some detail in the temperature range 1000" to 1250°C; less detailed studies were made with the other rare earths. In addition to being of value to the reactor program, the data obtained in this study should be of help in making a study of the role played by the electronic structures of metals in determining the nature of metallic solutions. The unique electronic structures of the rare-earth elements make them particularly interesting in this respect. EXPERIMENTAL The usual procedure for a solubility determination was to seal equal volumes of uranium and the particular rare earth in a tantalum crucible under an atmosphere of helium; this crucible was then sealed in a stainless steel jacket in an atmosphere of helium. These samples were equilibrated by repeated inverting of the crucibles in a furnace for 15 min at the desired temperature, left in an upright position for 15 min to permit separation of the two phases, and then quenched under a stream of water. In some runs the temperature of the furnace was held 50' to 100°C above the desired quenching temperature while inverting in order to insure good mixing. However, it was found that above 1200°C the crucibles were subject to failure and for these runs the furnace temperature was not raised above the desired quenching temperature. A small amount of tantalum was dissolved in the uranium and the rare earths in these runs, a maximum of 3 wt pct in the uranium phase at 1250°C and up to 1 wt pct in the rare-earth phase at this temperature. On cooling, the major portion of this tantalum precipitated as primary tantalum crystals. Any residual tantalum would probably have a negligible effect on the mutual solubility of uranium and the rare earths in each other. Samples for analysis were cut from each phase with an abrasive cutting wheel; the region near the interface between the two metals was carefully avoided. In the case of the rare earths with melting points above 1250°C no solubility data were taken on the rare-earth phase since this phase could not have achieved equilibrium in a reasonable length of time. (For the same reason no data were taken on the uranium phase below its melting point of 1132°C.) Equilibrium appeared to have been reached in the uranium phase in these cases although the rare-earth phase had not melted. To verify this, samples were melted together in an arc furnace similar to that described by Kroll.2 These samples were sub-

Jan 1, 1960

By H. R. Gardner

THE wide variation in tension and compression properties' which exist between the a,ß and ? phases of plutonium indicate that a study of the impact properties would give a greater insight into the variables affecting the mechanical properties and formability of plutonium. In tension the a phase is quite brittle with elongations less than 1 pct. The ß phase at normal test speeds is unusually ductile; elongations up to 621 pct have been obtained. However, the ß phase is quite sensitive to strain rate; elongations decrease rapidly as strain rate is increased. The y and 6 phases are less ductile than the ß phase, with the 5 phase the more ductile of the two. Neither are very sensitive to strain rate. Considering the brittle nature of a Pu in tension and compression and the strain rate sensitivity of ß -phase Pu, both unnotched Charpy specimens and standard V-notched Charpy specimens were fabricated for testing. In addition, since iron in amounts greater than 500 ppm forms a Pu-PlgFe eutectic network which inhibits microcracking in as-cast plutonium,' it was decided to prepare impact specimens with low and high iron contents. Chemical analysis and density data for the impact specimens are presented in Table I. A total of 55 impact specimens, 32 unnotched and 23 with the standard Charpy V-notch were tested at temperatures from —45' to 330°C. The test results for both unnotched and notched specimens are illustrated in Fig. 1.

Jan 1, 1962

By G. S. Cole, G. F. Bolling

Recent experitnents on crystal growth have indicated that thermal convection in the liquid influences the conditions of solidification and subsequent crystal characteristics. However, the important columnar-equiaxed transition in castings has not been examined as a function of this free convection. An experiment is presented here to demonstrate a change in the transition as convection is suppressed. Lead- and tin-based alloys were cast in stainless-steel molds either without special procedures or with thin grids placed horizontally across the tnold diameter in portions of the castings. While the grids allowed some diffusion of heat and mass transport to occur during solidification, massive fluid flow was hindered in the vegiotz of the grids. A larger percentage of columnar grains appeared in the grid region than in other parts of the special castings or than in all regions of normal castings. This structural difference can be explained in terms of thermal convection. THERE is little information about the possible effects of natural convection which may occur during the solidification of a casting. Ruddlel-3 has stated that the movement of liquid metal produced by convection might assist in the transfer of heat, but in the absence of confirmatory evidence he suggested that this idea be treated with caution. chalmers4 considers, but does not assess, the importance of fluid motion in casting, while Tiller5 suggests that, in liquid metals which have extremely high thermal diffusivities, the effects of thermal convection may be negligible. In addition, throughout the literature, there are statements indicating the possible presence of thermal convection but without proofs or evaluations of its influence. Natural convection may be divided into two categories: thermal convection, where liquid flow is due to density differences caused by temperature gradients in the liquid (e.g., liquid near a cold wall falling with respect to hotter liquid further away from the wall); and solute convection, where flow is due to density differences as a result of concentration gradients, see Wagner.6 In most instances thermal convection predominates over solute convection. Nevertheless, even thermal convection has not been studied in any detail and is difficult to handle, since it can vary quantitatively from one experiment to another depending upon the exact positions of heat sources and sinks or the attitude of the container with respect to gravity. In spite of these difficulties it is possible to realize that a higher rate of heat transfer must result with convection than in any situation where conduction acts alone. Thus, the degree of increased heat transfer is the important undecided question in considerations of casting. It will be demonstrated here that a reduction in convective flow leads to a suppression of the columnar-equiaxed transition in a casting and that this suppression is related to a decrease in the degree of heat transfer. Fortunately it is possible to do this by comparative rather than quantitative measures. The insertion of horizontal plates or grids into a simple mold does not appreciably influence heat transfer per se, but in reducing the allowable height through which colder liquid may fall freely it reduces natural thermal convection. This leads to a reduced fluid velocity, a different flow pattern, and an altered heat transfer which can be measured. EXPERIMENTS The alloys were prepared from elements of 99.9 pct nominal purity and cast into an uncoated 2-in.-diameter stainless-steel mold insulated at the bottom with a suitable refractory so as to ensure predominantly radial heat flow during freezing. Casting was performed either in empty molds or in molds containing grids or plates in various configurations . The grids were constructed of aluminum screen (1/32 in. thick) which was cut to the desired shape, cleaned in aqua regia, and coated with a mixture of 2 parts aqua dag graphite -1 part Sauereisen cement mixture, thus supplying a hard, protective, and insulating layer approximately 0.006 in. thick. The

Jan 1, 1965

By Robert G. Johnson, Richard J. Prosen

THE present work was done in connection with a program for investigating metal-to-ceramic braze techniques using the well-known reactivity of titanium with certain ceramic oxides. The investigation extended over the range of 0 to 30 pct Ti. The indium was obtained from the Indium Corp. of America and was specified to be 99.97 pct pure. The titanium was supplied by Metals Disintegrating Co. and was stated to be 99.5 pct pure. Each 2 to 4 g quantity of the alloy compositions studied was prepared using a mixture of 300 mesh powders pressed into pellet form. The pellets were heated in thin-walled molybdenum crucibles using a resistance-wound furnace in a vacuum less than 10-4 mm of Hg. Little reaction occurred between the molten alloy and the crucibles at the temperatures and short time intervals used. Only slight spectroscopic traces of molybdenum were found in the melt after a typical period of heating. The points shown on the phase diagram of Fig. 1 represent thermal arrest measurements. The phase boundaries shown are consistent with the thermal arrest and X-ray data obtained. The temperatures were determined using a Pt-Pt 10 pct Rh thermocouple with the junction spot-welded to a short molybdenum rod. The rod was almost completely immersed in the liquid metal, so that good thermal contact with the melt was achieved. The small mass of the crucible and its contents made it possible to obtain many heating and cooling curves in a few minutes time, an advantage which helped to minimize the effects of indium evaporation at the higher temperatures. A Brown strip chart recorder was used to display the thermal arrest curves. To check the technique, the melting point of 99.999 pure Al was measured and found to be within 1 of 660.2c, the accepted melting point. The upper limit of temperature with this technique was about 900°C beyond which the evaporation rate of indium became great enough to alter appreciably the alloy composition during the runs. A peritectic transition was found at 796°+ 5°C from 1.5 to 30 wt pct Ti. It was not found at 0.5 or 1.0 pct. A small super-cooling effect was observed. The melting-point transition to the B + L phase was found up to 20 pct Ti but not at 25 pct. Within the 1 deg accuracy of the measurement, the transition temperature was that of pure indium, 156°C. The liquidus line to the left of the 1.5 titanium point was located only approximately. The two crossed circle points at 1.0 and 0.5 pct Ti place a lower limit on the liquidus line and were determined by observing visually the cessation of vibrations in the melt caused by the damping effect as B-phase crystals were formed during cooling. The B-phase boundary was found at 23.8 pct Ti by X-ray analyses of small single crystals of the B phase isolated from 2 pct Ti melts. The compound was found to be Jetragonal with lattice dimensions of c = 3.052A and a = 10.094W. From the unit cell dimensions and the measured density of 6.42 i 0.05 g per cc, the compound was determined to be In4Ti3 which has 23.8 pct Ti. The calculated density of 6.44 g per cc is obtained assuming 2 mol

Jan 1, 1962

By R. H. Frazier, F. W. Boulger

THE performance of welded structures is closely associated with the ductile-to-brittle transition temperature of the steel from which they are made. A low transition temperature is desirable because it indicates that the steel is less likely to fail suddenly at low ambient temperatures. Structures such as bridges, ships, storage tanks, and pipelines are usually made from hot rolled semikilled steel. Changes in rolling practice or chemical composition appear to be the most practical methods for improving the toughness of such materials. This results from the fact that production is likely to be seriously curtailed if improvements were obtained by recourse to heat treatment or complete deoxidation. This paper discusses the effect of variations in carbon and manganese contents on the properties of semikilled steels. The transition temperature, a property to which considerable importance is attached, varies with specimen configuration, testing method, and criterion of performance. Both the Navy tear test' and the keyhole Charpy test were used in the investigation. According to the terms used by Vanderbeck and Gensamer,V he tear test was used to measure a fracture transition and the Charpy test to measure a ductility transition. In either case the specimens absorb considerably less breaking energy in tests below the transition temperature than in tests above the transition temperature. Decreasing the testing temperature of notched-bar specimens seems to be equivalent in its effect to increasing the severity of loading on fabricated structures. Therefore, structures built from steels exhibiting lower transition temperatures in laboratory tests are expected to be less susceptible to sudden brittle fractures in service. Consequently, changes in composition which lowered the transition temperature of the experimental steels were judged desirable. Materials and Methods The steels for this study were made in a laboratory induction furnace and rolled to 3/4 in. plate, using a finishing temperature of 1850°F. Precautions taken to insure reproducible melting and testing practices are discussed in detail elsewhere.' All tests were made on hot rolled steels. The analyses of the experimental steels are given in Table I. The list includes steels with manganese contents ranging from 0.21 to 1.46 pct, at each of five carbon levels. The phosphorus, silicon, sulphur, and nitrogen contents of the steels are reasonably constant. The tensile and notched-bar properties of the experimental steels are presented in Table 11. Both upper and lower yield points are listed, and all values are averages for duplicate specimens. Transition temperatures based on three criteria are given. The keyhole Charpy transition temperatures are the temperatures at which the energy-temperature curves based on averages of four specimens cross the 20 and 12 ft-lb levels. Both definitions of Charpy transition temperatures have been used by other investigators. Charpy tests were made at intervals of 10°F, using a pendulum with an available striking energy of 220 ft-lb and a velocity of 18.1 ft per sec. The specimens were oriented parallel to the rolling direction and notched normal to the surface of the plates. The tear test transition temperatures are based on the usual' criterion of the highest temperature at which at least one of four specimens exhibits brittle behavior. The dimensions of the tear test specimens are shown in Fig. 1. The specimen is loaded eccentrically in tension, with pin and shackle fixtures, through the large holes while submerged in a liquid bath at the proper temperature. A specimen developing a fracture area with less than 50 pct shear or ductile texture is classed as brittle. Tear tests were made at intervals of 10°F.

Jan 1, 1955

By G. R. Kotler, W. A. Tiller, W. C. Johnston

Experiments carried out on 100 g samples of molten tin and Sn-Pb alloys showed that the magnitude of the field strength, H, does not effect the supercooling, aTn, at which the initial nucleus of solid forms. However, the nuwlber of nuclei forming per cc, n, is found to increase greatly with H for fixed ATN and to increase with ATN for fixed H. The value of n increases greatly with the presence of solute and is found to he indefiendent of the cooling rate of the simples. Further, the most effective time for stirring to enhance n is immediately following the initiation of recalescence. The effect of inducing nucleation with the aid of a mechanical impact both in the presence and absence of stirring has also been studied. All proposed mechanisms for the enhancement of n by physical means seem to be inadequate in accounting for these results. IN any real nucleation experiment there are two effects influencing the results that would not be predieted from classical homogeneous nucleation theory. These may be separated into chemical and mechanical effects. The chemical effects are twofold. 1) Nucleation catalysts are present which enhance the nucleation frequency of new grains over that predicted by classical homogeneous nucleation theory.' 2) Solute will restrain the growth of nuclei so that many others may form.2,3 The mechanical effects are also two-fold. 1) Vibration (sonic or ultrasonic) of the freezing liquid enhances the nucleation frequency of new grains over that given by the chemical effects alone.4-6 2) Macroscopic fluid-flow effects produce the same effect as l).'7' It is also possible to influence the nucleation experiment by certain physical effects such as the application of electric and magnetic fields.9 From classical homogeneous nucleation theory one would predict the following relationship between the nucleation rate, N1 (number of nuclei formed per unit volume per unit time), and bath supercooling, AT, for an isothermal experiment.

Jan 1, 1963