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By J. E. Hilliard, H. E. Cook

It is shown that for systems having a miscibility gap the spinodal composition (c,) in the vicinity of the critical temperature (Tc) is related to the equilibrium composition (c,) by where cc is the critical composition. This expression is exart for the limit T — Tc and is a good approximation down to (T/Tc) - 0.9. If the miscibility gap meets the composition of the pure components THE recent theoretical treatment by cahn1,2 on spinodal decomposition of solid solutions has revived interest in methods of estimating the chemical spinodal, * i.e., the locus on a binary phase By use of the appropriate expression, the spinodal can be drawn on the Phase diagram almost by inspection with an accuracy comparable to that of lengthy calculations in which a free-energy function is fitted to the phase diagram. Satisfactory agveemenf is also obtained with spinodals comPuted from experimental free-energy curves. where G is the Gibbs free energy of the homogeneous solution and c the atomic fraction of component 2, the derivative being evaluated at constant temperature (T) and pressure (P). (These subscripts will hereafter be omitted from the partial derivatives.) If experimental measurements of G as a function of c and T are available in a single-phase region above the spinodal, then these can be extrapolated to obtain a series of metastable free-energy isotherms in the two-phase region, and from these isotherms a series of spinodal compositions can be calculated. Because the value of is very sensitive to experimental errors, an elaborate analytic procedure is required for

Jan 1, 1965

By Sidney Powers

INTRODUCTION THIS paper aims to present the results of an investigation concerning gravitative differentiation in lava flows, based on a quantitative microscopic and chemical study of a Triassic basalt from Nova Scotia, and confirmed by data secured from a study of certain Keweenawan ophites. At Cape d'Or, Nova Scotia, near the head of the Bay of Fundy, marine erosion has exposed a series of Triassic basalt flows in which copper has been sought for a number of years. Studies of the thickest flow of the series with regard to volumetric composition, specific gravity, and grain have been made by the junior author, and measurements by students at Tufts College, together with confirmatory observations of the same nature on this and other flows by the senior writer, and chemical analyses made by the Division of Mines, Mines Branch, Ottawa,1 all show differentiation in lava flows comparable to that in intrusive bodies. The work was made possible by the availability of the drill-cores, the property of Tufts College, and emphasizes the desirability of the study of drill-cores from other localities. Measurements of the specific gravities of certain Keweenawan lava flows of -Michigan by the senior author, from material collected in 1915, will be given as a comparison to those of the Cape d'Or flow. It has been found that the specific gravity of drill-cores, or volume per cubic foot, can be readily obtained by measuring their dimensions. This can be done to an accuracy of within 1 per cent. in pieces of core over 100 min. in length by using a micrometer gage for the diameter, and taking the average of six measurements of the diameter; and the average of four measurements of the length made with a finely graduated ruler. Then, after weighing them, the weight in grains per cubic centimeter, or what is practically the same thing, ounces per cubic foot, is obtained. System-

Jan 2, 1916

By H. O. Hofman

PAGE Part I. Decomposition OF Zinc Sulphate by Heating in Air, . 811 I. Introduction............ 811 II. Zinc Sulphate Used,..814 III. Fumes and Pyrometers,........ 816 IV. Heating Zinc Sulphate—Dehydration,...... 817 V. Heating Zine Sulphate—Decomposition...818 (a.) Beginning of Decomposition..818 (b.) Rapidity of Decomposition,..819 (c.) End of Decomposition..823 (d.) Retardations in Heating.-Curve,..825 (e.) Formation of Basic Salts,..826 Part II. Decomposition OF Zinc Sulphate by Heating with Carbon (Reducing Roast), .830 I. Introduction,........... 830 II. Experiments and Results ,.........832 III. Summay,........... 834 Part 111. Formation of Zinc Sulphate by Roasting,... 834 I. Introduction,........... 834 IT. Description of Blende,......... 836 III. Outline of Operations,.........838 IV. Details of Experiments.......... 839 1. Decrepitation,.......... 839 2. IgnitionPoint,....839 3. Sulphatizing-Roast of Joplin Blende..840 4. Sdphatizing-Roast of Warren Blemde ,..843 5. Sulphatizing-Roast with Addition of Pyrite...848 6. Summary of Sulphutizng-Roasts,....... 854 Part IV. Appendix : Formation of Zinc Ferrate...856 Part I. Decomposition of Zinc Sulphate by Heating in Air. I. Introduction. With the exception of lead sulphate. all common metallic

Jan 1, 1905

By Eric Jette

THE alloys of gold and iron were investigated in 1907 by Isaac and Tammann,1 who determined the thermal diagram for the entire system by thermal analysis and microscopic examination. They also reviewed the previous literature. Since that time the present authors have been able to find only two additional references. Guertler and Schulze2 examined the electrical resistance of iron alloys containing up to 40 per cent by weight of gold between 0° and 200° C.; the alloys were in the form of wires. The main result of this work was to confirm the solid solubility limit found by Isaac and Tammann; namely, about 20 per cent Au. More recently, Shih3 has examined the magnetic properties of gold alloys containing up to 10 weight per cent iron and particularly the range below 5 per cent. The alloys in the lower range were found to be paramagnetic; the 10 per cent alloy was ferromagnetic. The present investigation was undertaken to determine the solid solubility limits at various temperatures, particularly below the a-y transformation point of iron, since these solubilities are of some impor-tance in connection with the working of the alloys in the manufacture of certain "grey" and "blue" golds.

Jan 1, 1934

By Elwood Haynes

As in organic nature certain animal and vegetable forms have undergone modifications, and thus, as it were, fitted themselves to live in a new environment, so it has been found possible in certain instances to form new metallic combinations which are practically immune to the natural conditions that exist on the earth at the present time. A few of the metals, such as gold, silver, and the metals of the platinum group, are found native, but the commoner heavy metals are nearly always found in the form of oxides or sulphides. For a number of years I have made experiments with a view to finding certain metallic combinations or alloys which would not only resist oxidation and other harmful influences, but would also possess valuable physical properties, which would render them fit for special service. The first decisive step made in this direction was the discovery of an alloy of nickel and chromium in 1898. The properties of this alloy have already been described in a previous paper.' Immediately following the discovery of the nickel-chromium alloy, I produced an alloy of cobalt and chromium, which has likewise been described in the reference given, but the range of proportion between the cobalt and chromium is so great, and the combination is modified to such a marked degree by the introduction of other substances, that I have felt justified in reading another paper on this subject, describing more fully my researches in this particular direction. As early as 1907 and 1908,I made alloys or combinations of cobalt, chromium, and tungsten; cobalt, chromium, and molyb-

Jan 1, 1913

By E. P. Abrahamson II

The effect of transition element binary solid-solution additions upon the recrystallization temperature of vanadium has been investigated. Cr, Fe, Co, Ni, Mo, Ru, Rh, Os, and IY lower the recrystallization temperature, while Ti, Mn, Zr, Cb, Pd, Hf, Ta, W, Re, and Pt raise it. The majority of elements have their greatest effect in less than 0.15 at. pet. Both the rate of change of recrystallization temperature with atomic percent solute and the limit of initial linearity are found to Correlate with the free atom grozdnd state outer electron configuration of the sollte element. THE work of Abrahamson and Grant1 and Abraham-son and Blakeney2 has shown that the solute's elec- tron configuration appears to be a prime determinant of both brittle-to-ductile transition and recrystallization temperature in dilute binary alloys. It has also been shown by Abrahamson3 that the limit of linearity, i.e., the limit of initial linear change of recrystallization temperature with atomic percent solute, is also a function of the electron configuration of the solute in iron base alloys. Because of the systematic nature of the solute electron configuration correlation, it is to be expected that some systematic variation might be observed upon changing the solvent. Thus, this study of recrystallization in vanadium base alloys is the first of a series aimed at defining the role played by the electron configuration of solute in determining the recrystallization characteristics of dilute transitional alloys. PROCEDURE All alloys were made using 99.8 pct V containing 0.011 C, 0.024 N, 0.08 0, 0.006 H, 0.035 Ta, and

Jan 1, 1962

By Howard I. Smith

THE mineral trona was discovered on Government land in 1938, about 18 miles west of the town of Green River, Wyo., in the core of the John Hay, Jr., well, a test well drilled for oil by the Mountain Fuel Supply Co. in sec. 2, T. 18 N., R. 110 W., sixth principal meridian. Although no oil or gas was reported, the drilling led to a discovery of great potential economic importance and exceptional scientific interest. Trona is a relatively rare mineral found at various places in playa deposits but not previously found at any great depth. It is composed of hydrous sodium carbonate and sodium bicarbonate (Na2CO3.NaHCO3.2H20). Pure trona contains the equivalent of 70.35 per cent of sodium carbonate. Total sodium computed to carbonate times 1.42 equals the percentage of trona in the absence of other sodium carbonate minerals in the sample. ANALYSES Two short pieces of core sent to the author a few days after the completion of the Hay well were analyzed by R. C. Wells[t] and found to be exceptionally pure trona. A brief article on this discovery has already been published.1 Arrangements were then made through W. T. Nightengale[$] for the author to inspect and sample the core. The company log of the well included notations of sodium salt crystals from a depth of 1190 to 1820 ft. Not being able to identify all the minerals present, the author selected several specimens of the core for petrographic examinations and sampled the sections of trona for analysis. J. J. Fahey, of the Geological Survey, examined the specimens and reported most of the crystals to be a new anhydrous sodium carbonate and calcium carbonate mineral (Na2CO3.2CaCO3), which he named shortite,2 after Maxwell N. Short.* He analyzed the samples of trona, with the results shown in Table I. [ ] In addition to the 10 ft. I in. of trona in the first section, it is reasonable to assume that the 9 ft. 8 in. for which no core was available for sampling is either mostly trona or some other soluble salt, but it should not be included in calculating reserves. In the part of the core between 1590 ft. 3 in. and I600 ft. 4 in. the material insoluble in cold water, amounting to about 5 per cent, was estimated petrographically to contain 70 per cent clay, 25 per cent shortite, and 5 per cent pyrite. ACCOMPANYING MINERALS Fahey later secured about 600 ft. of the core for a detailed examination and study

Jan 1, 1942

By Henry M. Howe

The natural effect of large ingot-size should be to increase segregation. I have previously pointed1 to the excessive segregation in many large ingots as tending to confirm this, but I have shown that in case of ingots less than 16 in. square this expected effect of ingot-size is liable to masked by that of Note.—The "maximum enrichment" of each ingot, i.e., the excess of the richest spot over the average of the whole ingot, is first calculated in percentage of that average. The average of the maximum enrichment of the several ingots of a given group is the abscissa in Fig. 1. The number beside each spot tells the number of cases which that spot represents. FIG. 1.—Influence of Ingot-Size on Maximum EnrichNent in Steel Ingots. other variables. Under these conditions we should expect that, if large ingot-size really does tend to increase segregation, this effect would be shown by taking the average of large numbers of cases, so that the effects of these other variables might off-set and cancel each other.

Jan 1, 1910

By Eugene N. Cameron

Pegmatites occur in abundance in certain areas of South West Africa and in adjacent parts of northern Cape Province in the Union of South Africa. Some of the pegmatite deposits were prospected before World War I, and intermittent mining was done during the period between the two World Wars. The deposits of the Jooste lithium mines became an important source of lepidolite, and beryl, columbite-tantalite, wolframite, and other minerals were produced. Since World War I1 prospecting and mining have been carried on vigorously, with the result that South West Africa produced nearly 3000 tons of beryl during 1949 through 1953, together with more than 40,000 tons of lithium minerals, chiefly lepidolite and petalite. Small amounts of tantalite, columbite, and other minerals have also been produced during the postwar period. Most of this mineral production has come from two pegmatite areas roughly outlined in Fig. 1. One is the Karibib-Omaruru area of north central South West Africa. The second is the northeastern part of the Orange River pegmatite area, which extends into adjacent parts of Namaqualand, in northern Cape Province of the Union of South Africa

Sep 1, 1955

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 W. A. Backofen, R. L. Fleischer

Single crystal, bicrystal, and polycrystal tensile tests of aluminum at 4.2°K, 77°K, and 300°K have been used to examine the role of grain boundaries in the deformation process. Results indicate that a grain boundary may affect the extent and slope of easy glide. The stage II hardening rate, on the other hand, is independent of the presence or absence of grain boundaries. This conclusion allows the size of the region of multiple slip caused by an incompatible grain boundary to be determined. For the size of bicyystal sample used in this study, multiple slip occurs in about half of the cross section. PREVIOUS studies of the stress-strain characteristics of bicrystals of face-centered-cubic metals have been limited to aluminuml-5 at room temperature. Recent results, however, indicate that the stress-strain curves of single crystals of such metals may be separated into at least three stages6 in which different deformation processes are occurring7 provided testing is done at sufficiently low temperatures.' Since for aluminum a well-defined stage II develops only below room temperature, previous studies have not been able to relate effects of grain boundaries to all of the three stages of deformation. It is therefore to be expected that low-temperature deformation of aluminum single crystals and bicrystals should clarify the effects of grain boundaries on the different processes of deformation. EXPERIMENTAL PROCEDURE Single crystals and bicrystals were grown from the melt by the standard techniqueg with aluminum reported by Alcoa to be 99.993 pct pure. Ridges in the boat were used to guide the grain boundary during growth, assuring that the boundary would bisect the sample.10 The rate of furnace motion during growth was 1.0 cm per hr. During growth zone purification resulted, as evidenced by the ability of the first material to freeze to recrystallize at room temperature following severe deformation. Samples were approximately 4.4 X 6.6 mm in cross section and 103.5 mm in length between grips. Samples were annealed at 635" i 5°C for 40 hr and furnace cooled over a 7-hr period. They were then electropolished in a solution of 5 parts methanol to 1 part perchloric acid at a current density of 15 amp per sq dm for about 30 min at temperatures below 0°C. Tensile testing was performed at 295" (room temperature), 77" (sample in liquid nitrogen), and 4.2"K (sample in liquid helium) on the hard-type machine indicated schematically in Fig. 1. The machine con- sists basically of a tube surrounding a rod; one end of the sample is attached to each member, and the rod is pulled up the tube to extend the sample. The rod is rigidly mounted and is moved vertically by a system described by asinski." The pulling force is measured continuously by an electrical strain gage load cell, and the relative displacement of the tube and rod is also recorded continuously by a soft cantilever beam with electrical strain gages. Maximum stress and strain sensitivities were ±2g per sq mm and * 3-10-5. In all tests the strain rate was approximately 5.10-5 per sec. The thin wires in the tensile apparatus introduce softness, which may be corrected for, however, by measuring load vs displacement with the sample replaced by an elastic member. For loads greater than 15 kg the spring constant is 1.875.106 g per cm. The flexible wires also served to reduce substantially the large shearing forces which may arise in the case of grips having horizontal rigidity.'' As in any gripping system, however, bending moments will arise in the course of deformation by single slip. Engineering stress, s = (load)/(original cross-sectional area), and strain, E = (increase in length)/ (original length), are used for stress-strain curves unless otherwise indicated. Tables list resolved shear stress, T=mo and shear strain ? = dm, where m is the usual Schmid resolved shear stress factor for the primary slip system at the start of deformation. The first group of samples to be described forms an isoaxial set, all of the crystals making up the single crystals or bicrystals having the same tensile axis, the orientation of which is indicated by the cross in Fig. 2. For this orientation the primary slip plane and slip direction make angles of 45 deg with the tensile axis and the Schmid factor m has its maximum possible value of 0.5. Rotations about the tensile axis are indexed by means of an angle 0 between the small-area surface of the samples and the projection of the primary slip direction onto the cross section, as defined in Fig. 3. In single crystals, values of 0 were 0 and 90 deg, while in bicrystals 0 values were (0 deg, 180 deg), (90 deg, 270 deg), and (0 deg, 90 deg) as indicated in Fig. 4.

Jan 1, 1961

By Harmon D. Nine

DISCONTINUITIES in the strain vs cycles to failure (S-N) fatigue curves have been reported for polycrystalline materials by Porter and Levy for copper,&apos; by Benham and Ford for mild steel,&apos; and by Williams and Stevens for 18/5 stainless steel.3 However, the considerable scatter of the points for polycrystalline materials have made a statistical analysis of the data necessary to confirm the discontinuities. By using oriented single crystals of high-purity material and low grown-in dislocation content, the reproducibility has been improved to a degree which makes statistical analysis unnecessary and clearly shows discontinuities in the S-N curves. A recent publication4 reports an abrupt change of slope in the S-N curve for ( 100) axially oriented copper single crystals fatigued in torsion. The reported discontinuities in the S-N curve, along with a different mode of failure in the two branches of the curve, have been interpreted as being due to two different amplitude-dependent mechanisms. Further work on copper single crystals of (111) axial orientation has now been completed which supports the existence of an amplitude-dependent change in the mechanism of fatigue. In the present study torsional-fatigue experiments were performed at room temperature in air on (111) axially oriented single crystals. The crystals were grown and prepared as reported in Ref. 4. Two different methods of growing the crystals also discussed in Ref. 4 produced crystals of two different perfections: Type I with dislocation density of >l06 per sq cm and Type I1 with dislocation density of <l05 per sq cm. In accordance with the accepted hypothesis that crystal flaws are primarily responsible for the scatter of fatigue lifetimes, particularly at low amplitudes, the values for the three lowest amplitudes for Type I crystal are believed to be uncertain. They are marked in braces in Fig. 1. On the other hand, at amplitudes above a surface shear strain (SSS) of *0.003, the data for both types of crystals are in agreement. Most remarkably the S-N curve for torsional fatigue of (111)-oriented single crystals is composed of three segments separated by rather large discontinuities, Fig. 1. Metallographic and X-ray measurements indicate that the lower branch corresponds to the low-amplitude (F) region of the (100) copper single-crystal results as discussed in Ref. 4, and the two upper branches correspond to the high-amplitude (H) region. The lower discontinuity in the (111) S-N curve and metallographic evidence give further support to the hypothesis that different mechanisms of failure operate for high and low strain amplitudes. Reference to the resolved shear stress diagram, Fig. 2, shows that the highest stressed systems in

Jan 1, 1965

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Table

Jan 1, 1910

Jan 1, 1938

By N. C. Fick, R. L. Rickett

DESPITE the large volume of literature on alloys of iron and silicon,1 there is little published information dealing specifically with the constitution, at various temperatures, of the alloys containing less than 0.10 per cent carbon and from 0.5 to 4.5 per cent silicon that are used extensively in the electrical industry. Because such information was needed by those engaged in the production of these alloys in, sheet and strip form, the constitution was determined of a series of commercial silicon steels containing 0.05 to o.08 per cent carbon and 0.20 to 4.31 per cent silicon, over the temperature range 1300° to 2500°F. These alloys were all from regular production heats and of commercial quality. The results of, the tests are given in Fig. I and for comparison the constitution diagram of iron-silicon alloys with 0.01 to 0.02 per cent carbon, taken from the literature, is also presented. MATERIAL AND PROCEDURE All steels were hot-rolled to a thickness of 0.060 to 0.080 in. and used in this condition. Their chemical compositions, including over-all check analysis, and "center" analysis for carbon and silicon, are given in Table I. The so-called center analysis was determined on samples pickled to remove the outer surface and leave only the center third of the gauge thickness-the portion of the sample on which metallographic ratings were based. In most, but not all, of the alloys, slightly higher carbon was found at the center than when the whole strip was sampled; the two methods of sampling show relatively little difference in silicon. Two of the steels, A and E, contain relatively high phosphorus, while I and K are somewhat higher in manganese than the others. [ ] The procedure consisted essentially of heating specimens (about 1/16 by 3/8 by 3/4 in.), protected from oxidation or change in carbon content, at a temperature in the desired. range and for a selected period, after which they were quenched in iced brine; the degree of transformation that had then occurred was determined by metallographic examination. The soaking periods were selected after preliminary investigation had shown that a con-

Jan 1, 1946

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&apos;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