Search Documents

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 H. Margolin, A. Coucoulas

Fused alumina or silica crucibles were used as the containing vessel. X-ray powder photographs were taken with CrKcll radiation and the use of Straumanis type Norelco cameras of 114.6 mm diam. Crystallographic data for these compounds are shown in Table I. These results are in agreement with others of this series reported previousiy by Wernick and Geller.l The abnormally large atomic volume of Eu and Yb metal3 is manifested in EuAl2, and YbAl2, when one compares them with the compounds reported in Ref. 1. However, the behavior of Yb in the Ni compound is the same as that observed in Ni5Yb having the Cu5Ca structure.4 Thus, it appears that Yb has a valence closer to 3 only in the Ni compounds; i.e.. the 2s and Id electrons are lost to the conduction band.1,4 Europium in EuAl2 appears to have a valence closer to 2. In previous work.' X-ray patterns of ingots of NdCo2 indicated a multiphase material with faint indication of the presence of a phase having the MgCu2 structure. It was thought that NdCo2 formed peri-tectically from the melt, resulting in a multiphase system when cooled under nonequilibrium conditions. An X-ray pattern of a sample from a melt cooled more slowly than in past work showed that the major phase was NdCo2. The lattice constant, Table 1, is in excellent agreement with the value one can predict from the data of Ref. 1. One of the other phases was identified as Co5Nd. The author wishes to thank J. H. Wernick for his guidance in this work, and D. Dorsi for the preparation of these compounds. 'J. H. Wernick and S. Geller: AIME Trans., in press. 'J. H. Wernick and S. E. Haszko: to be published. 'F. H. Spedding and A. II. Daane: Progress in Nuclear Energy, Series 5, Eds.: H. M. Finniston and J. P. Howe, McGraw-Hill Book Co., 1956. 'S. E. Haszko: to be published. 'J. H. Wernick and S. Geller: unpublished work. 6J. 11. Wernick and S. Geller: Acia Cryst., 1959, vol. 12, p. '162. A reinvestigation of the lattice parameters of Ti-Sn alloys by the present authors has revealed consistently hiher c and a values than those obtained by Worrier. The difference is most pronounced in the c-parameters which are, therefore, used as a basis of comparison. This discrepancy was not due to air contamination, since precautions were taken to exclude this possibility. Specimens for Debye-Scherrer photograms were prepared by arc-melting 15-g buttons of Bureau of Mines titanium and 99.99 pct Sn, with compositions of 6, 8, 10, and 12 at. pct Sn. These buttons were forged under an oxygen-acetylene torch to 1/8-in. sq rods to produce relatively fine grained material. Scale and some surface were ground off, and the rods, after wrapping in molybdenum and titanium sheet, were vacuum annealed 15 min at 800 °C to remove hydrogen. Heat treatment of one of the samples at 875°C to produce ß and reveal a at the surface showed no surface contamination. Following the vacuum treatment, the specimens were wrapped in molybdenum and annealed for 1 week at 800°C under two atmospheres of argon. After water-quenching, the rods were ground so that the thickness tapered to 0.1 mm at the tips. The ground specimens were then recapsuled as previously and reannealed 1 hr at 800°C to remove cold work. The

Jan 1, 1961

By M. E. Glicksman, G. S. Ansel

It has been shown by Walton and Chalmers,&apos; that the mechanism of the development of solidification textures in castings involves the preferential growth of dendrites along certain crystallographic directions as determined by the influence of prevailing heat-transfer conditions. Specifically, this growth is parallel to the direction of heat flow. The purpose of the present investigation was to determine the applicability of this mechanism to: 1) commercially pure aluminum (99.75 pct) solidified to coarse columnar crystals, 2) an aluminum-silicon casting alloy (5.25 Si, 0.38 Fe, 0.02 Cu, 0.10 Ti, 0.002 S) cast to a fine-grained structure under complex and reproducible heat transfer conditions. PROCEDURE AND RESULTS 1) Specimen Preparation—Using a technique similar to that employed by Northcott and Thomas,2 samples were melted in the directional solidification mold, Fig. 1, by heating the entire assembly to 670° C. The mold was then end cooled using a water jet. From considerations of harmonic analysis, the likely distribution of isotherm normals encountered by the growing crystals is shown in Fig. 1. 2) Orientation Analysis—The pure aluminum sample froze to large columnar crystals, nucleated close to the metal diaphragm, with the growth direction parallel to the mold axis. Using the back-reflection Laue technique, the growth direction of these crystals was determined to be a <100>, which coincided within 53 deg to the mold axis. Metallographic examination of the relatively fine-grained aluminum-silicon sample revealed a dendritic mode of growth having its primary dendrite arms along isotherm normals. Back-reflection Debye-Scherrer patterns, taken with the specimen rotating about its cylinder axis, yielded rings displaying symmetrical arcing, thus proving that the sample did, indeed, possess the solidification texture inferred from the preliminary metallographic examination. To determine the extent and orientation of the texture in the aluminum-silicon alloy, a density distribution of the (111) poles of the sample was plotted using the transmission Debge-Scherrer

Jan 1, 1960

By T. H. Courtney, G. W. Pearsall, J. Wulff

The mechanism, of the decrease in transition temperature of Nb3Sn was studied. Presintered Nb3Sn powder compacts were heated in a dynamic vacuum at elevated temperatures. Loss of tin resulted from these heat treatments and transition temperature as a function of tin loss was discontinuous at a critical value of tin loss. In general, decreases in transition temperature were correlated with deviations from AS a class, the A15 structure-type (also known as Cr3O or ß W) compounds have the highest superconducting transition temperatures.&apos; In this class are stoichiometry and presence of vacancies in niobium sites. The effect of vacancy distribution on the physical properties of Nh3Sn is considered. It is concluded that the close-packed atomic A chain in the A3B A15 structure must be preserved both in coutinuity and direction if good superconducting properties are to he obtained. tempting to associate these high transition temperatures with the most unique crystallographic feature of the A3B structure: the three close-packed, orthogonal chains of A atoms.&apos;0"&apos; There are at least three principal mechanisms by which the perfection of these "A-atom chains" could be interrupted: changing stoichiometry, inducing disorder, or introducing vacancies. Departures from stoichiometry in either direction in V3Ga lower the transition temperature.7&apos;8 In niobium-

Jan 1, 1965

By J. O. Brittain, E. P. Lautenschlager, D. A. Kiewit

Compression tests were run in the temperature range of 700° to 900°C ox 0&apos; phase NiAl intermetal-lic alloys of several grain sizes. At these temperatures the minimum strengths were observed at the stoichiometric composition. While significant increases in strength occurved in both the low-nickel (vacancy) and high-nickel (substitutional) regions, the highest strengths were found in the high-nickel region. During deformation serrated flow was sometimes observed in the low-nickel alloys. After deformation transgranular cvacking and deformation striations were observed in all compositions tested. AS part of a general investigation of the properties of NiAl inter metallic compounds, a preliminary study of the role of point defects upon plasticity was made by high-temperature compression tests on ß&apos; NiAl specimens of several grain sizes and compositions. ß&apos; NiAl is an intermetallic compound having a CsCl structure and a rather wide range of composition from A1-45 at. pct to 60 at. pct Ni.1 According to Bradley and Taylor2 and to cooper,&apos; it possesses a defect lattice in which departures from stoichiometry in the direction of decreased nickel content lead to the presence of vacant nickel sites (although Cooper&apos;s work indicates that a small amount of substitution also occurs) whereas departures on the high-nickel side lead to substitution of nickel on aluminum sites. NiAl forms congru-ently from the melt at approximately 1650°C,1 and thus has a higher melting point than either of its component elements. Up to this time, although this and other high-melting intermetallic compounds have been suggested for elevated-temperature usage,4 only the hardness4 and a few tensile-strength measurements5 have been reported for NiAl at high temperatures. In the present investigation the effects of composition upon the compressive-strength properties in a range of 700° to 900°C have been measured for NiAl of several grain sizes. EXPERIMENTAL PROCEDURES The alloys were made as described elsewhere6 from an A1-46.8 at. pct Ni master alloy furnished by the International Nickel Co. with additions of high-purity nickel and aluminum. The charges were vacuum-induction-melted in A12O3 crucibles with small amounts of helium added to the atmosphere to suppress vaporization. They were cooled slowly from the melting temperature to achieve uniform grain size. In order to refine the as-grown grain size a special rolling technique was developed. Alloys were packed into 0.10-in. wall-type 302 stainless-steel tubes which were partially filled with magnesium oxide to prevent bonding between the alloy and the steel jacket. The ends of the tubes were closed by hot forging, and the packets were then hot-rolled. The alloys with greater than 50 at. pct Ni were rolled at 1100°C, but it was found necessary to increase the temperature to 1350° C before alloys with less than 50 at. pct Ni would roll without cracking. With these temperatures, reductions as high as 48 pct were achieved in a single pass. The rolled alloys will hereafter be referred to as "fine grained" whereas the as-grown material will be designated "coarse-grained&apos;&apos;. The compression specimens were made by cutting square cross-sectional pieces, approximately 3/16 by 3/16 by 1/2 in., with a water-cooled diamond cut-off wheel from the as-grown or the rolled alloys. Specimens were ground to their final dimensions by polishing through 3/0 grit silicon carbide papers. The final shape was a rectangular parallelepiped of square cross section having a height-to-width ratio of 3:1. Compression testing was carried out in a compression rig of our own design mounted on an In-stron Floor Model. The specimen chamber could be heated to 1000°C and was controlled within ±2°C. The compression rig was enclosed within a bell jar and was maintained at a 50 µ of mercury vacuum throughout the duration of the test. The test cham -ber was heated from room to test temperature within 15 min. Specimens were then held at the test temperature 30 min prior to testing. Previous experiments indicated that no grain growth would occur within this time. An Instron Variable Crosshead speed unit was used to adjust for small variations in specimen lengths in order to have a constant initial strain rate, €, for all specimens of a group. For the fine-grained specimens the strain rate was changed rapidly at constant temperature by a factor of 10 with the speed lever on the Instron. For a given € the compression data was analyzed in terms of true plastic strain (E) and true compressive stress (0).

Jan 1, 1965

By F. V. Lenel, J. A. Dromsky

The influence of thermal and mechanical treatments upon the elevated-temperature stability, the microstructure, and the mechanical properties of aluminumaluminum oxide alloys is investigated. Compacts from atomized and flake aluminum powders were sintered at temperatures just below the melting point of aluminum in a vacuum of 10-4 Torr for several hours and were then hot-extruded. During this treatment the hydrated oxide present on the surface of the powder particles decomposes, the water vapor produced reacts with the aluminum forming additional oxide, and the hydrogen developed is removed by diffusion to the surface. Extrusions of compacts so treated are stable at temperatures up to the melting point of aluminum and do not form blisters and internal defects typical of extrusions from untreated powder. Electron micrographs of extrusions from unsintered compacts show long parallel oriented platelets typical of the IN previous investigations1,2 the properties and structure of aluminum alloys were described which had been produced by cold compacting aluminum powder and then hot pressing and extruding the compacts. Both the structure and the properties of these alloys differ significantly from those of commercially produced aluminum-aluminum oxide alloys. When commercial alloys are produced, they are given a sintering treatment in air at about 600°C between cold compacting and hot pressing.3 During this treatment the hydrated oxide, which forms the coating on the aluminum-powder particles from which the alloys are prepared, is decomposed.4 The chemically bound water which is released reacts with the aluminum and forms aluminum oxide and hydrogen. Most of the hydrogen is eliminated by hydrated oxide skins on the original powder particles, while the oxide particles in extrusions from sintered compacts are broken up. The strength of extrusions from treated powder is somewhat lower, but their ductility is considerably improved over that of extrusions from untreated powder. The slip-trace patterns on electron microgvaphs of extrusions which had been strained in tension before replication indicated that the grain size of the extruded alloys was approximately the same as that of the original powder particles. Coarse-grained alloys were produced by severe cold work and annealing of the extrusions. They have significantly lower room-temperature strength, but about the same ductility as the vacuum-sintered fine-grained alloys. The size and shape of oxides in the coarsegrained alloys is similar to those of the fine-grained alloys, but the particles are somewhat more uniformly distributed. diffusing to the interior or exterior surfaces of the compact, but it is difficult to eliminate all of it because the solubility of hydrogen in solid aluminum is very low even near its melting point. Extruded alloys which are not sintered before hot pressing and extrusion appear quite sound upon metallo-graphic examination, but develop severe surface blistering and internal fissuring when they are heated to elevated temperatures after extrusion. Fissures and blisters have been observed even in commercial alloys which had been sintered in air, although they are much less severe.&apos; Because of the importance of this sintering treatment it was studied in this investigation in a modified form, using vacuum sintering rather than sintering in air, to insure more complete elimination of occluded hydrogen. The technique is described and the effect of the treatment upon the structure of the alloys is documented. It has been known for some time that extruded aluminum-aluminum oxide alloys are very finegrained, but an exact delineation of grains in electron micrographs has not been possible. In the conventional metallographic-etching techniques etchants attack the interface between the second-phase particles and the matrix in the alloys much

Jan 1, 1964

By J. S. Kirkaldy, R. J. Brigham

AS part of a program to investigate the diffusive properties of dilute ternary austenites we have made a determination of the effect of chromium on the activity of carbon in ? iron. Samples of pure iron and Fe-Cr alloys were simultaneously equilibrated at 1000°C with atmospheres of controlled CO/CO2 ratio following the procedures of Darken and Gurry 1 and smith.&apos; The 0.020 in. thick samples were analysed for carbon by microcombustion analyses3 to yield a set of isoacti-vity curves. The latter were converted to activity coefficient (standard state at infinite dilution) plots by combination with Smith&apos;s data2 for the activity of carbon in pure iron. Table I shows the compositions of the base alloys used while Fig. 1 shows the experimental results superimposed on a set of curves generated by the relation In?C=6.9NC- 10.7 Ne In the notation of wagner4 the interaction parameter is the second coefficient on the right, Sgr = - 10.7 This can be compared with a rather inaccurate value deduced from the transient equilibrium diffusion data of Kirkaldy et al.&apos; as The authors are indebted to the American Iron and Steel Institute for financial support. This material is condensed from a B.Sc. Thesis submitted to Mc-Master University by R.J.B.

Jan 1, 1963

By L. C. Brown, J. S. Kirkaldy

AS part of a program to investigate the diffusive properties of dilute ternary austenites, we have made a determination of the effects of cobalt on the activity of carbon in y iron. Samples of pure iron and Fe-Co alloys were equilibrated at 1000°C in controlled CO-C02 atmospheres, following the procedure of R. P. smith.&apos; The 0.010-in.-thick samples were analyzed for carbon by microcombustion analysis2 to give a set of isoactivity curves. The latter were converted to activity-coefficient (standard state at infinite dilution) plots by combination with Smith&apos;s data&apos; for the activity of carbon in pure iron measured in a CO-C02 mixture. Table I shows the composition of the base alloys used, while Fig. 1 shows the experimental results superimposed on a set of curves generated by the relation In the notation of Wagner,&apos; the interaction parameter is the second coefficient on the right, Eq. [I] adequately correlates the experimental points for those alloys which were prepared from electrolytic iron by arc-melting (2,4, and 7 pct Co). The results for the 10 pct Co alloy (supplied by Crucible Steel Co.) do not correlate with the others. In fact, the experimental points for this alloy correspond to an interaction parameter of It is possible that the interaction parameter is a function of concentration for large cobalt additions, but we are inclined to attribute the discrepancy to the presence of a significant amount of aluminum added to this alloy for deoxidation purposes (Table I). Preliminary measurements for the system Fe-C-A1 have indicated an interaction parameter of Which implies that the effect of aluminum in solution on the carbon activity is sixty times greater than the effect of an equal weight percentage of cobalt. These observations cast some doubt on the precision of the results of Chandhok, Hirth, and Dulis who used this same 10 pct alloy for their experiments, obtaining an interaction parameter of E = 2.2. There is considerable doubt as to the activity of a plane Fe-C system. Using a CH4-Ha mixture, Smith obtains cg = 6.1 for the effect of carbon on the activity of carbon in austenite, rather than the value of 8.9 given in Eq. [I]. The use of CHa-Ha mixtures has recently been questioned by Kaufman, Radcliff, and cohens who state that there is some possibility that other hydrocarbons are formed. Fortunately, the determination of the interaction parameter is not sensitive to the coefficient of Nc. For example, using the CHI-H2 results we obtain a value which is almost the same as that obtained using the CO-CO, results. The authors wish to thank Mr. E. J. Dulis of the Crucible Steel Co. of America for supplying us with samples of the 10 pct Co alloy used in their original work. Financial assistance from the American Iron and Steel Institute is gratefully acknowledged.

Jan 1, 1963

By R. H. Edsal, G. Horvay

The transient one-dimensional heat-conduction equation is solved for two semi-infinite media, at different initial temperatures, brought into (poor) thermal contact. It is shown that the two interface temperatures gradually approach the well-known, instantaneously assumed, common interface temperature of the case of zero contact resistance. The present analysis, with slight modifications, may be used in fair approximation also for the class of problems involving a change of phase. One particular application of the boundary value problem solved in this paper concerns the determination of the temerature distribution in metal castings and molds.P This problem has been treated extensively by analytical methods (see e.g., the literature listed in Refs. 1 to 4) by electrical analogs (see e.g., Refs. 11 to 14) and also by experimental methods (see the literature listed in Refs. 4, 5), but prior to the work of Pellini and Ruddle the experimental results were rather unreliable. Immediate motivation for the work described in this paper was to better understand the results of Pellini,&apos; which are typified by the curves of Fig. 1. Pellini cast 20-in. high, 7-in. sq steel ingots into cast-iron molds of various thicknesses. He measured the temperatures at mid-height, at a number of stations in the casting and the mold, and obtained the cooling curves shown in Fig. 1. The most conspicuous thing about Fig. 1 is the large temperature discontinuity at the interface between casting and mold, due presumably to the shrinking away of the casting from the wall and the attendant creation of an airgap, as the casting cools. Such a contact resistance slows down the anticipated freezing rate; an analytical prediction of this is of much interest in foundry practice. The determination of the effect of contact resistance in the presence of a change of phase will be dealt with, by an approximate method, in a separate paper.&apos; In the present paper the much simpler problem is considered in which there is contact resistance but no change of phase. This problem frequently arises in the analysis of aircraft structures.+ It can initial temperature, and may be regarded as semi-infinite. We suppose the thermal properties in the two media to be time, temperature, and space independent; there are no sources or sinks within the media. In Fig. 2(a) we illustrate, schematically, the case of no contact resistance. The semi-infinite media "0" and "2:&apos; initially at temperatures 8 (- m) and 8(+ -), respectively, are brought into contact at time t =0. A common interface temperature 8,i = is instantaneously established, and thereafter remains constant. Temperature curves are drawn for three instants of time. When there is contact resistance between the two media, the temperature equalization occurs as sketched in Fig. 2(b). Fig. 2(c) illustrates a temperature curve, using the dimension-less notation Barber-Weiner-Boley6 have stated the equations [23a, b] given below, which govern the heat transfer in the two media, and solved them for a rather complicated special case. In their example both media "O" and "2" are finite, are insulated at their far ends, and there is steady heat input into medium 2. In Eq. [27] we give the general solution of the most important special case, where the two media are semi-infinite and there are no sources or sinks. In this case the results may be stated in simple closed form. In Fig. 3, to be discussed in the next section, we illustrate the variation of the interface temperature when steel at 2760" is brought into contact with iron at 100". This example, if we disregard the effect of the heat of fusion, is comparable to Pellini&apos;s experimental results, Fig. 1. (Unfortunately, however, experimental determination of the interface temperatures was difficult to carry out, and hence Fig. 1 is not very accurate near the interface.) In Figs. 4 and 5 of the following section, we similarly compare Pellini&apos;s sandcasting tests with analytical results for the case where steel at 2760" is brought into contact with sand at 100". It will be pointed out that, insofar as the interface temperatures are concerned, heat of fusion has a minor effect, and so the present approximation, with small modifications, may be

Jan 1, 1961

By B. Lustman

IT has been known for a number of years that the addition of certain alkaline-earth and rare-earth metals to nickel-chromium base electric resistance alloys causes marked increase in their oxidation resistance as measured in an intermittent oxidation test similar to the ASTM—B 76-39 test. However, the mechanism whereby such additions confer protection is still unknown. On the basis of the experimental observation that the increase in life in intermittent tests of nickel-base alloys is proportional to the atomic radius of the added element, Hessen-bruchl has proposed that the added metal ions occupy holes in the oxide lattice which prevent diffusion of nickel ions in the oxide. Horn2 has postulated that the effect of such additions is to accelerate the rate of diffusion of chromium atoms in the metal and thus to favor the formation of Cr2O3 scales. Hickman and Gulbransen3 concluded from electron diffraction observation that the addition of the elements in question results in the formation of a barrier to the diffusion of nickel ions to the surface. In connection with some experiments conducted to develop a test for the resistance to oxidation of alloys in massive form under conditions of intermittent heating and cooling, some features of the oxidation of nickel-chromium base alloys were revealed which may serve to elucidate the mechanism of protection conferred by rare-earth and alkaline-earth elements. Specimens of cylindrical shape 0.360 in. diam x 0.37 in. long were placed in alumina thimbles and inserted into and withdrawn from a Globar-heated furnace operating at 2150°F at 7.5-min intervals. The weight of oxygen absorbed as well as the weight of oxide spalled were measured at intervals. With the particular experimental conditions used, the specimens in a 7.5-min cycle were at temperature for 2 min which corresponds to the heating time of the ASTM test. The weight changes of similar specimens under conditions of continuous heating in air at 2150°F were also measured. The specimens after oxidation were examined metal- lographically and samples of the spalled oxide collected for X-ray diffraction identification of the phases. The specimens used were from the same lot of material used by Hickman and Gulbransen3 and had the compositions shown in table 1. Included in table I are results of ASTM life tests on the materials in the form of wires. These materials, their composition and the ASTM life test data thereon were all supplied by F. E. Bash of the Driver-Harris Co. It may be noted that alloy 13246 of 80 pct nickel, 20 pct chromium nominal composition showed the shortest life in the ASTM test; alloy 12246 with the addition of small amounts of the elements zirconium, calcium, and aluminum, intermediate life; and alloy 12046 with the further addition of silicon, the longest life. In fig. 1 are shown the results of intermittent oxidation tests made on bulk specimens at 2150°F as well as continuous tests conducted at the same temperature. Duplicate tests were made on each specimen. The ordinate scale, inches of metal oxidized, was derived from the experimentally observed weight gain of oxygen by assuming the scale to be entirely Cr2O3. It may be noted that the metal-loss curves during intermittent oxidation show a constant or continuously increasing rate of weight

Jan 1, 1951

By R. P. Smith, L. S. Darken

FOR a high-purity iron-carbon alloy, the phases existing at equilibrium at temperatures slightly above the temperature of the metastable eutectoid are austenite and cementite (hyper-eutectoid) or austenite and ferrite (hypo-eutectoid) and those below the eutectoid temperature (carbdn content > 0.02 pct) are ferrite and cementite. Such an alloy, held in a carefully controlled and measured temperature gradient (about 1 cm per deg) which includes the eutectoid temperature, then quenched and examined metallographicall.y, provides a precise measure of the eutectoid temperature. The samples were hollow cylinders approximately 1.3 cm OD, 0.63 cm ID, and 2.5 cm long, which had been carburized to the desired carbon content at 1000°C with an appropriate hydrogen-methane mixture. In order to cool the samples to room temperature with a minimum change of carbon content and absence of cracks they were quenched in a Cerro Bend bath at 480°C, followed by a water quench to room temperature. Three cylinders were suspended in the furnace1 in a purified helium atmosphere, by a closed end, thin wall, silica tube which also served as a thermocouple well. The top and bottom cylinders served only as a means of making the temperature gradient along the center cylinder more uniform. The temperature was determined with a platinum-platinrhodium thermel calibrated by comparison with a thermel certified (* 0.5OC) by the National Bureau of Standards. After holding in the temperature gradient for 24 to 72 hr the samples were quenched in Cerro Bend at 480°C, followed by

Jan 1, 1960

By C. F. Floe, M. Cohen, M. B. Bever, V. G. Paranjpe

NITROGEN is becoming recognized as one of the important elements in ferrous physical metallurgy. Several investigations indicate that nitrogen plays a significant part in such phenomena as strain aging and temper brittleness. Nitrogen is known to be a powerful austenite stabilizer and may be a useful alloying element in stainless steels. The hard-enability of steels has been found to be increased by nitrogen. The nitriding and carbonitriding pro- cesses are examples of successful industrial application of nitrogen in ferrous materials. For a complete understanding of the various roles played by nitrogen in ferrous metallurgy a knowledge of the phase relations in the binary iron-nitrogen system is essential. The literature contains a variety of iron-nitrogen phase diagrams. Those suggested by Fry, Sawyer,² Epstein and his coworkers,³,4 Lehrer,5 Eisenhut and Kaupp, and Hägg7,8 differ significantly from one another. Moreover, the most recent of these investigations dates back about twenty years. A re-determination of the entire diagram, therefore, appeared desirable. In addition to the published phase diagrams, investigations have been carried out on portions of the iron-nitrogen system. Bramley and Haywood determined the composition and temperature of the iron-iron nitride eutectoid. The maximum solubility of nitrogen in a iron was investigated by Portevin and Seferian10 and by Diikstra. 11 Several investigators have determined the crystal structure of the various iron-nitrogen phases. In fact, Eisenhut and Kaupp6 and Hägg7 based their entire phase diagrams on the results of such studies. Epstein&apos; and Emmett12 reported lattice-parameter measurements on the E (hexagonal close-packed) phase. Osawa and Iwaisumi13 also studied the change in lattice parameters of the a, ? and E phases with nitrogen content. Brill" and Hägg7 showed that the y&apos; phase has an ordered structure, and Hendricks and Kostingsl5 suggested a structural relation between the E and phases. In a recent investigation Jackl6 arefully studied the super-lattice reflections from the y&apos; and phases. Experimental Procedure A large variety of experimental methods has been used by previous investigators in studying the iron-nitrogen system. These methods include thermal analysis, microscopic examination, X ray diffraction, magnetic measurements, dilatometric analysis and anelastic experiments. Thermal analysis is of low accuracy in determining changes in the solid state. This accuracy is further decreased for iron-nitrogen alloys since composition changes will occur during heating or cooling unless extreme care is taken to use a proper atmosphere. Metallographic methods are of small value . because homogeneous iron-nitrogen alloys of sufficient size can be produced only with great difficulty. The dilatometric method is open to the same objections as thermal analysis in that changes in composition may occur during heating and cooling. Lattice parameter measurements by X ray diffraction can be used satisfactorily for determining phase boundaries if cooling is fast enough to prevent structural changes. In view of the limitations of the various other techniques two methods only were adopted in this investigation, (1) X ray diffraction measurements, and (2) a controlled nitrogenizing method. The second method is an innovation for investigating the iron-nitrogen system, and will be described at length. Materials: All iron-nitrogen alloys were prepared from carbonyl iron powder of the analysis shown in table I. The average particle size was about 20

Jan 1, 1951

By A. H. Willis

In this paper the results of an exploratory study conducted by personnel of the Knolls Atomic Power Laboratory on 18.6, 22, and 40 wt pct uranium-zirconium alloy will be presented. LARGE power output and long life or endurance from small, compact nuclear reactors require the use of enriched uranium for a fuel material. However, the ability to remove heat from the element, the tolerable fission product concentration, or some physical limitation of the system can limit the allowable fuel enrichment, or the volume of the fuel. In such an event, the fuel-element designer must find a suitable diluent for the uranium. Zirconium is an attractive material for such a purpose as it has good nuclear properties and a low absorption cross section. Because it has a good corrosion resistance in water, zirconium is a suitable cladding and the uranium-zirconium alloy makes a good core material. The alloy has a high thermal conductivity, a low elastic modulus, and a low thermal-expansion coefficient, all of which tend to minimize the thermal stresses in the core. Furthermore, the unirradiated material is reasonably ductile and has adequate strength. DESCRIPTION OF SAMPLES Two types of cylindrical samples were used in this study 1) short, 1/2-to 1-in.-long samples to test the irradiation stability of the alloy and 2) 2-ft-long samples to investigate the resistance of an element to longitudinal deformation. Composite samples having 0.060-to 0.140-in. core diameters and clad with a 0.010-in. thick zirconium jacket were fabricated by coextrusion and by hydrostatic pressure bonding of the cladding to the preformed core material. Both fabrication processes produced the uniform metallurgical core-to-cladding bond which was felt necessary to minimize the uncertainty of the central core temperature. Both consumable-arc and induction-melting techniques were used to melt the uranium-sponge zirconium core alloy. Since a graphite crucible was used in induction-melting this alloy contained more carbon impurity, which resulted in a finer grain size than the arc-melted alloy. The core to cladding interface was uniform in elements coextruded from the fine-grain, induction-melted core material, while those elements fabricated using arc-melted core material had very irregular interfaces because of the coarser grain size. Both cold-worked and heat-treated fuel materials were irradiation tested. Cold-worked alloys were swaged to about a 5 pct final reduction in area without subsequent annealing. The heat-treated alloys were heated to 800°C, held at temperature for 15 min and furnace-cooled to about room temperature. This

Jan 1, 1960

By M. Cohen, D. J. Blickwede

Quantitative observations concerning the carbide phases in high speed steel are of importance for two general reasons: (1) the carbides, being inevitable constituents of the final structure, exert a direct influence on the properties of the steel; and (2) a substantial proportion of the total alloy content is tied-up in the carbides, and hence the extent of their solution on austenitizing governs the composition of the steel matrix. The latter relationship has a vital bearing on the response of the steel to tempering as well as on its performance in subsequent service. Accordingly, in the course of a long-term study of the behavior of high speed steels, the authors were confronted with the problem of securing quantitative data on the carbide phases. The obvious method for acquiring such information is to isolate the car-bides from the steel and subject them to chemical, X ray diffraction and other measurements. There are well-known extraction techniques which involve the chemical or electrolytic solution of the less noble matrix (ferrite, marten-site or austenite), thus leaving a residue of the carbide phases. However, the results obtained must be scrutinized carefully1,2 since the carbides may be affected by the chemical or electrolytic action. It is the purpose of the present paper to describe the experiments leading to an electrolytic-extraction technique for quantitatively isolating the carbides from both I and hardened high speed steel. Particular attention is paid to the amount, as well as the composition, of the carbides so that the matrix analysis becomes ascertainable by subtraction from the overall steel composition. Illustrative results are given for the M-2 grade of tungsten-molybdenum steel. Review of the Literature The chemical method of dissolving the matrix selectively with respect to the carbides makes use of dilute non- oxidizing reagents such as hydrochloric or sulphuric acid. Although this simple procedure has led to the determination of the cementite composition,3,4 it achieved only limited success because of the interaction between the acid and the carbide residue. Some of the carbides may not only be destroyed in this way, but the hydrogen released is likely to remove part of their carbon as hydrocarbon gases. The electrolytic technique of isolating carbides has the advantage of rapidly dissolving the specimen (anode) in the presence of less reactive solutions than are practicable with the chemical method. This reduces the possibility of chemical attack on the carbides, and furthermore, the hydrogen evolved during the electrolysis is released at the cathode which is not in close proximity to the carbides. The common electrolytes adopted for this purpose are hydrochloric and sulphuric acids.5-l1 Aqueous solutions of ferrous salts have also been used.12,13 A considerable advance in experimental technique was introduced by Treje and Benedicks14 who developed a double-compartment cell for electrolytic extraction, the anode and cathode chambers being separated by a porous diaphragm. A solution of 15 pct sodium citrate, 2 pct potassium bromide and 1 pct potassium iodide was selected for the anolyte, while the catholyte consisted of a 10 pct solution of copper sulphate, with copper serving as the cathode. This type of cell has a number of desirable characteristics: 1. The anolyte has a pH value close to 7, at least at the beginning of the run. 2. The iron that dissolves from the anode-specimen forms a water-soluble complex ion with the citrate, thereby preventing the precipitation of iron hydroxide (which would contaminate the carbide residue) despite the neutrality of the solution. 3. Copper deposition instead of hydrogen evolution occurs at the cathode, and this avoids an increasing concentration of hy-droxyl ions which (in an otherwise neutral solution) might cause the precipitation of insoluble hydroxides. 4. Contamination of the anode chamber by copper sulphate is inhibited by the porous diaphragm. Houdremont and coworkers15 applied the above method (with the further refinement of excluding oxygen during the electrolysis, washing and drying) to the extraction of carbides from a series of plain carbon steels after various heat treatments. They had quantitative success only with specimens in the annealed condition, and concluded that the size and shape of the carbide particles play an important role in the isolation process, with large spheroids exhibiting the least tendency to decompose during the electrolysis. Up to the present time, the citrate double-cell has not been used to any extent for isolating the carbides of high alloy steels, apparently on the grounds that the complex carbides are more resistant than cementite to attack in the simpler acid electrolytes. In particular, Bain and Grossmann7 and Gulyaev10.10a have employed the hydrochloric acid cell for their investigations of the carbides in high speed steel.* It will be demonstrated here that this type of cell is capable of yielding quantitative results in the case of high speed steel, and actually has certain advantages over the more complicated double cell. However, in order to provide a rigorous test of the quanti-tativeness of electrolytic procedures for the problem at hand, both methods were studied in considerable detail.

Jan 1, 1950

By J. M. Lommel, B. Chalmers

The transfer of material from the solid to the liquid states can be accomplished in several ways. It occurs by the application of heat in the more familiar metallurgical operations of melting but it can also be accomplished by solution in a liquid, as can occur in systems using liquid metals for heat-transfer media. The kinetics of this reaction were studied to learn more of the separate steps in the process. In the case of a pure material at its melting temperature, Te,A,, thermal energy must be supplied to the solid in an amount equal to the latent heat of fusion in order for the solid-to-liquid transfer to occur. The rate at which heat is supplied, and hence the rate of melting, is a function of the temperature gradients in the system. This phase transfer can also occur in the absence of temperature gradients in the system illustrated in Fig. 1, which is a partial phase diagram between the components A and B. A solid of composition CI will

Jan 1, 1960

By R. H. Fillnow, D. J. Mac

IT has been demonstrated that alloys in systems structurally analogous to steel undergo reactions during heat treatment similar to those of steel, and yet very little work has been done on such systems, particularly from the standpoint of isothermal transformation. The literature shows that most of the isothermal work has been done on copper-alu-minum alloys.&apos;-3 One paper has been reported on a eutectoid copper-silicon alloy,&apos; and one on iron-nitrogen alloys.&apos; There are no published reports of previous investigations on the isothermal transformation of a eutectoid beryllium bronze. At least two papers have been reported on isothermal age-hardening phenomena in beryllium-copper alloys of lower beryllium content.5,6 No information on eutectoid transformations in this alloy is given in Rimbach&apos;s book, "Beryllium." &apos; This paper reports the results of an investigation on the isothermal transformation characteristics of a eutectoid beryllium bronze. Materials Used and Procedure The eutectoid beryllium bronze used in this investigation was centrifugally cast at the Ampco Metal Co., Inc.. Milwaukee. Wis. A master beryl-lium-copier alloy, purchased from the Brush Beryllium Co., was used to make the beryllium addition to the copper. The alloy was melted under a protective carbon cover and Centrifugally cast into a graphite mold preheated to 920°F. Some trouble with segregation of beryllium in the melt was encountered but was eliminated by violent stirring during melting and pouring. Since very thin .specimens were used for the various heat treatments, only one casting of the four prepared was necessary for the complete isothermal study, thus eliminating any discrepancies caused by a change in composition. The analysis of this casting was 94.035 pct copper and 5.965 pct beryllium. A slight trace of iron was found. The castings were sawed into specimens approximately 3/32 in. in thickness. The specimens were first given an "austenitizing" heat treatment at 788°C, well within the single phase beta region, in a vertical tube furnace. The specimens were then quenched individually into an adjacent molten salt bath and held there for a definite time, after which they were quenched into water at room temperature. Temperature was held to ± 2°C in the salt bath. No homogenizing treatment was given the specimens other than the soak at 788°C. Despite the fact that this was a cast alloy, no evidence of segregation or coring was found in the microstructure of any specimen, either as cast or as heat-treated. Upon completion of the heat treatment, the speci-mens were subjected to metallographic examination

Jan 1, 1951

By L. E. Tanner

A study was made of the transformation kinetics of the commercial titanium-base alloys, Ti-2.5Al-16V and Ti-4A1-3Mo-1V, using two different heat treatment cycles: 1) step-quenching to aging temperatures from a ß solution anneal and 2) water-quenching from an a+ß solution anneal followed by reheating to aging temperatures. Metallographic examination and hardness testing provided the major portion of the data, while electrical resistivity, dynamic elastic modulus, and X-ray diffraction techniques were also used to obtain critical or confirnzatory data. TTT diagrams were constructed from these results. The sheet alloys, Ti-2.5Al-16V and Ti-4Al-3Mo-lV, are of current interest to the Department of Defense. As an aid to the planning of their respective heat treatments, a program was carried out to determine the transformation kinetics of these alloys using ß and a + ß reference states. This paper presents the TTT diagrams constructed from data obtained by metallographic examination and hardness testing, as well as by X-ray diffraction, electrical resistivity, and dynamic elastic modulus determinations. EXPERIMENTAL PROCEDURE Materials—Sheets of the alloys were obtained from the producers in the "mill anneal&apos;&apos; condition and were tested prior to shipment to assure their meeting mill specifications. Chemical analyses were also supplied and are presented in Table I. Heat Treatment—The majority of the data presented in this study were obtained from hardness testing and metallographic examination. Sample coupons were cut from the alloy sheets and heat treated on a mass scale. The annealing cycles were operated in the following manner: 1) The ß solution anneals were carried out in a horizontal tube furnace with the bare samples protected by a dynamic helium atmosphere. Following this treatment, the samples were quickly transferred to lead, solder, or oil-bath furnaces for aging. After prescribed times at the various reaction temperatures they were water-quenched. 2) The second cycle simulated commercial practice, which requires the alloys be quenched to room temperature following the a + ß sotution anneal. This anneal was carried out in a muffle furnace with a dynamic helium atmosphere. The specimens were then reheated from room temperature to the reaction temperatures in lead, solder, or oil-bath furnaces, followed by water-quenching after prescribed times at these temperatures. Information concerning the critical temperatures of the alloys was obtained from the literature and solution treatments were based on these values. A summary of the solution annealing treatments is found in Tables II and 111. The aging temperatures for the (a - ß alloys ranged in 50°C intervals from 200°C (392oF) to approximately 50°C below the ß/a + ß transus in the first cycle and approximately 50°C below the a + ß solution temperature in the second. The transformations in these alloys were known to be rather rapid, and thus aging times started at 0.5 min and went to a maximum of 7000 min. Hardness Testing and Metallographic Preparation— Heat-treated samples were sheared and mounted in Bakelite with their cut edges exposed. After a flat surface was ground, the hardness was determined in Vpn on an Armstrong-Vickers machine utilizing a 20-kg load. Three to four impressions were made on each specimen. After this procedure, the hardness impressions were removed and the edge surfaces of the specimens were prepared for metallographic examination in the usual manner. The etchant was 20 pct HF and 20 pct HNO3 in glycerine. Electrical Resistivity—It has been shown that the isothermal reactions in titanium-base alloys can be followed by measuring the changes in e1etrical resistivity attending these transformations. As mentioned earlier, the reaction kinetics of a-ß alloys are rather rapid; thus, it would be expected that resistivity data obtained from these alloys would be more significant if measurements were made continuously at the respective aging temperatures. To accomplish this end a special apparatus, based on the current-potential principle, was assembled. Its design was derived from earlier devices constructed by Colner and zmeska1 4 and Levinson.5 The set-up provided for the heat treatment zof machined specimens (approximately 2 x l0-4in. in cross section by 3 in. in length) and the determination of their relative

Jan 1, 1962

By G. R. Speich, L. Zwell, H. A. Wriedt

HORNBOGEN1 has recently shown that it is possible to measure the degree of segregation accompanying cellular precipitation in ferritic Fe-Zn alloys by measuring the lattice parameter of the ferrite phase within the cell. In extending this work to study a wider range of composition and to evaluate the free-energy change accompanying the growth of the cells, the lattice parameter and the boundary of the a! phase in Fe-Zn alloys were redetermined to check the earlier work of Schramm.2 This appeared to be necessary since the Fe-Zn diagram has been drastically revised since the work of Schramm. Stadelmaier and Bridgers3 have recently shown that, contrary to prior belief: the y iron field is closed and the a phase boundary extends to an a + T + L peritectic at 782°C where peritectic a contains 42 at, pet Zn. The phase T in equilibrium with the a solid solution has a composition near Fe3Zn10 and a cubic 081 type structure.&apos; The alloys were prepared by two different methods. In the first method carbonyl iron powder (400 mesh, 99.8 pet pure) and high-purity zinc powder (200 mesh, 99.99 pet pure) were mixed in a ball mill and then pressed in a die at 40,000 psi pressure to make compacts with a diameter of 5/8 in. and a thickness of l/4 in. These were sealed in evacuated quartz capsules, gradually heated to 1100°C over a period of 1 week (3 days at 600°C) 3 days at 800°C, 1 day at (1100°C), and then quenched in iced brine. For the alloys containing 23.4 and 30.1 at. pet Zn, the final homogenization anneals were performed at 1000" and 850°C) respectively, to prevent melting. In the second method, alloys were prepared by saturating high-purity iron millings (0.001 in. thick, 99.9 pet pure) in zinc vapor. For each alloy the impregnation was conducted at a fixed vapor pressure of zinc and a fixed temperature between 750&apos; and 800°C. When saturated, the millings were quenched in water. The compositions of the alloys prepared by both methods are shown in Table I. The lattice parameters of the ferritic solid solutions were determined by comparing the 26 values of their 310, peaks (using Co Ka) or of their 211 , peaks (using Cr Ka) with those of corresponding peaks for high-purity a iron. The 20 values were measured at room temperature with an X-ray dif-fractometer (G.E.-XRD5). By using a fiducial value of 2.8664Å for the lattice parameter of pure a iron at 25oC,5 the lattice parameters of the zinc-bearing

Jan 1, 1964

By D. Mizer, J. B. Clark

In a recent investigation of the entire magnesium-yttrium phase diagram, Gibson and Carlsonl report the maximum solid solubility of yttrium in magnesium as 8.0 wt pct Y at 1050°F (565°C). However, independent studies of the precipitation in magnesium-yttrium alloys, 2 conducted prior to the publication of the above phase diagram, indicated a greater solid solubility of yttrium in magnesium. Accordingly, the magnesium-rich region of the diagram was re-investigated to determine more accurately the liqui-dus, solidus, and solvus of the magnesium solid solution in this system. Metallography and thermal analysis were chosen as the best corroboratory methods for accurately determining this region of the diagram. A magnesium-yttrium master alloy was prepared from yttrium of 99.9 pct purity and singly sublimed magnesium of 99.9 pct purity. Subsequently, the twelve magnesium-yttrium alloys shown in Fig. 1 were prepared using standard alloying procedures for magnesium. Metallographic samples of the alloys were heat treated under an SO, atmosphere in an electric furnace capable of control within ±3oF of the desired temperature. Alloy samples for the metallographic determination of the solvus, and for the precipitation studies were solution treated at 1000° F for 16 hr. Subsequent aging times for the solvus determination varied from a minimum of 6 hr for the higher temperatures up to 816 hr for lower temperatures. Periods of 30 min at temperature were used to generate liquation for the metallographic placement of the solidus and eutectic horizontal. The liquidus and the eutectic horizontal were determined by thermal analysis. Simple time-temper- ature heating and cooling curves were obtained from 10-g alloy samples heated and cooled in a well-insulated electric furnace. An Iron-Constantan thermocouple was placed in direct contact with the sample. The thermocouple did not contaminate the alloy sample since metallographic examination showed that the alloy melt did not wet or attack the thermocouple. The heating and cooling rates were approximately 5 to 8OF per min. The standard metallographic procedures for magnesium were used in polishing and etching the magnesium-yttrium alloys studied. Glycol etchant* was etchant was used to reveal precipitate in the solvus determination. The redetermined magnesium-rich region of the magnesium-yttrium diagram is shown in Fig. 1. The

Jan 1, 1962

By K. Winterton

The effect of a near-saturation unidirectional magnetic field was to decrease slightly the endurance life in Armco iron in the annealed condition. This is expected since the energy-absorbing mechanism associated with domain wall movemenfs is inhibited bv the strong field. THE factors determining failure under dynamic loading are not yet clearly recognized. One of the properties of a metal which might be thought to influence fatigue life is its damping capacity. Foppl1 conducted an alternating torsion test at a frequency of 700 cpm continuously for a period of 3 years, without the test piece breaking or showing any sign of deterioration after 1100 Million cycles of stress; the rate of dissipation of energy was sufficient to maintain the test piece at about 100°C above the ambient temperature. cazaud2 concludes from this, and others follow him, that (&apos;the capacity for dissipation in the form of heat by internal damping is unlimited" and that therefore "damping is not directly related with fracture by fatigue." Future work may necessitate qualification of this rather uncompromising view. It is shown later that in all

Jan 1, 1960