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By W. A. Boyer

MINE hoisting ropes can be loaded to capacity only when the strength of each component is exactly known. Characteristic curves provide this information. When load and rate of acceleration are specified for an individual rope, the characteristic curve assumes a predetermined shape. The factor of safety, is the value drived by dividing the breaking strength of the rope by the total stress on the rope where it goes over the sheave wheel. Total stress is made up of the weight of skip and cage, weight of ore, weight of rope between sheave and rope clevise or attachment, friction load, and acceleration force. Bending stresses are neglected in this discussion, as sheaves and drum di- ameters are assumed to be large enough to reduce these stresses to a negligible value. Fig. 1 gives a family of characteristic curves for 13/8-in. improved plow steel hoisting ropes of 6 x 19 or 6x21 construction with hemp core. The curves will hold for all rope sizes of this specified quality when they are subjected to an acceleration rate of 1.608 ft per sec2. As acceleration rates of most deep shaft hoists are close to this value, this family of curves is presented as a standard. Deviation is slight for other rope sizes. Note that for 0, which is the most important, the deviation is within the width of the line on the graph. In a comparison of values for ropes of 5/8-in. and 21/4-in. diam the deviation ranges from a maximum 1.2 pct at 0 ft to 0.6 of 1 pct at 5000 ft. The value of 5.55 is the same for all ropes at 2500 ft. The tabulated columns to the left of the curves are values of total connected load for each size of rope from 5/8-in. to 21/4-in. diam inclusive. With each of these connected loads the safety factor characteristic curve for the rope will be as indicated.

Jan 1, 1957

By J. A. Cordier

USINOR has conducted four blast furnace tests with formcoke made from marginal coals in the pilot plant of HBNPC ("H. B. N." process) in four different blast furnaces of various sizes and with medium and high grade burdens. The tests in which up to 30% of the blast furnace fuel burden was replaced by HBN formcoke yielded good results on furnaces ranging in hearth diameters from 5.6 to 9.5 m. The tests demonstrated that: - the slight increase of pressure drop in the stack did not prevent smooth operation at record production rates with the same blowers - the total coke rate was not significantly affected, but there was a slight increase of the calorific value of the top gas - the hot properties of the formcoke are satisfactory as seen from the samples withdrawn at the tuyeres and tapholes. At the end of the two months test on BE nº 2 of Dunkerque, the hearth was as permeable and clean as with good standard coke. Our conclusion is that HBNPC formcoke made from lean coals is an adequate substitute for conventional coke in the range of 20 to 40% replacement, which is of interest to USINOR.

Jan 1, 1977

New Mexico School of Mmes, and State Bureau of Mmes and Mineral Resources, Socorro, N. M E. H Wells, President and Director Available publications of the State Bureau of Mines and Mineral Resources of the New Mexico School of Mines are: Bulletin 4, Fluorspar in New Mexico (1928), 60 cents; 5, Geologic

Jan 1, 1933

By U. C. Tainton

THE world-wide continuation of low prices for zinc in 1934 has militated against any striking changes in the position of the metal. The price of zinc in London at the end of the year, about £11 5/8 per long ton, or 2.5c per pound at the existing rate of exchange, is more than 25 per cent below the average price for last year. In these circumstances probably the most significant event of the year was the putting into operation of the new electrolytic plant of the Giesches company at Magdeburg, Germany. This plant is designed to have an initial capacity of 20,000 long tons per year, and it is planned to increase this to 40,000 tons annually. This development is significant not only as indicating a continued swing of zinc metallurgy towards electrolytic methods, but also because it will make Germany virtually independent of foreign sources of zinc. This will mean that the zinc tonnage normally exported to Germany from Belgium and Poland will. become available for other markets, and it is feared may lead to lower prices unless the production of Bel-

Jan 1, 1935

By F. C. Farnham

In this paper it is shown that the vertical component of the earth's magnetic field for the area of the United States can be very closely approximated by assuming it to be the sum of the vertical components of two magnetic fields. The first of these is the field of the centered dipole described by J. Bartels,l the axis of this dipole intersecting the earth's surface at two antipodal points, the one in the Northern Hemisphere being at latitude 78.5°N. and longitude 69°W., and the dipole having such strength that the vertical component of its field at its axis pole is 0.63 oersted. The second field is that due to a second dipole situated halfway between the center of the earth and a point on the earth's surface at latitude 38.5°N., longitude 91.5°W., the axis of this dipole being along this radius, the south pole being toward the surface, and the dipole being of such strength that its field at the point on the surface mentioned above is 0.069 oersted. Using the field computed on this basis, tables of latitude and longitude correction factors were computed and are included in the paper. It is also shown that the secular variation in the earth's magnetic field may be roughly accounted for by assuming it to be due to motion of the second dipole described above. Although there is an error in the equation used for computing the strength of the field due to the second dipole, the distribution of the field computed is very nearly that of the assumed arrangement of dipoles. Subsequent work by the author reveals that minor changes can be made in the location of the second dipole, which will make the computed field fit the observed field even more closely. It has also been found2 that the vertical component of the field may be represented equally well by a purely empiric relationship.

Jan 1, 1940

By F. C. Farnham

In this paper it is shown that the vertical component of the earth's magnetic field for the area of the United States can be very closely approximated by assuming it to be the sum of the vertical components of two magnetic fields. The first of these is the field of the centered dipole described by J. Bartels,l the axis of this dipole intersecting the earth's surface at two antipodal points, the one in the Northern Hemisphere being at latitude 78.5°N. and longitude 69°W., and the dipole having such strength that the vertical component of its field at its axis pole is 0.63 oersted. The second field is that due to a second dipole situated halfway between the center of the earth and a point on the earth's surface at latitude 38.5°N., longitude 91.5°W., the axis of this dipole being along this radius, the south pole being toward the surface, and the dipole being of such strength that its field at the point on the surface mentioned above is 0.069 oersted. Using the field computed on this basis, tables of latitude and longitude correction factors were computed and are included in the paper. It is also shown that the secular variation in the earth's magnetic field may be roughly accounted for by assuming it to be due to motion of the second dipole described above. Although there is an error in the equation used for computing the strength of the field due to the second dipole, the distribution of the field computed is very nearly that of the assumed arrangement of dipoles. Subsequent work by the author reveals that minor changes can be made in the location of the second dipole, which will make the computed field fit the observed field even more closely. It has also been found2 that the vertical component of the field may be represented equally well by a purely empiric relationship.

Jan 1, 1940

By J. Marin, H. A. B. Wiseman

This paper presents results of an experimental study dealing with the plastic stress-strain relations of aluminum alloy 14s-T6 subjected to combined biaxial tension and compression stresses. Plastic stress-strain relations for both constant and various types of variable stress ratios were determined and a comparison made with the simple flow theory of plasticity. THIS investigation was undertaken for two main purposes: 1—to obtain plastic stress-strain relations for Alcoa 14s-T6 when subjected to various combinations of biaxial stresses, and 2—to determine the validity of the flow theory in predicting plastic stress-strain relations for combined stresses. The biaxial stresses used were tension combined with compression. These stresses were produced by subjecting a thin-walled tube to axial tension and torsion. Various types of combined stress conditions were investigated. To provide control test data and information on the influence of biaxial stresses on the strength, the usual constant stress ratio type tests were made. These tests showed that the biaxial yield strength agrees approximately with the distortion energy theory. However, the Prager semi-empirical theory agrees best with the test results. For these constant stress ratio tests the plastic stress-strain relations were found to agree approximately with the flow theory. (For constant stress ratio the flow and deformation type theories are identical.) In view of the fact that the constant stress ratio tests cannot distinguish between the flow and deformation theories, a number of variable and special biaxial stress tests were made. These tests showed that the flow theory is inadequate, since large differences exist between the experimental and theoretical results. Special tests were also conducted to check the validity of the isotropic yielding assumption. This assumption is made in the linear-type flow theory. In these tests specimens were loaded in tension to predetermined values beyond the proportional limit stress. The tension load was removed from the specimen and the specimen was subsequently loaded in torsion to fracture. Other tests were applied with the order of torque and tension loading reversed. The nominal stress-strain curves for these two tests approximately coincide indicating that the assumption of isotropic yielding is valid. In addition, these tests verified the requirement by the slip theory of plasticity that prestraining below 140 pct of the proportioned limit stress in either torsion or tension did not influence the subsequent plastic stress-strain relations in tension or torsion. Another type of test was conducted to check the assumption that the axes of principal stress and strain coincide during plastic flow and variable stress conditions. These tests were conducted by applying combinations of axial tension, internal pressure, and torsional loads such that the axes of maximum principal stress could be rotated through 90" during the loading path. The results of those tests showed that the angle between the axes of principal stress and strain varied greatly and that the assumption of coincidence between axes of principal stress and strain is not valid. Introduction In various metal processes including forming of sheets, rolling of bars, and extrusion of rods, metals are subjected to stresses beyond the yield strength of the material. Often these stresses are not simple stresses acting in one direction, but are combined stresses acting in more than one direction. Structural and machine members are often subjected to combined stresses. To adequately determine the factors of safety for these members, it is necessary to know the plastic stress-strain relations of the material used for various combined stress conditions. Various theories have been proposed for predicting the plastic stress-strain relations for combined stress in terms of the simple tension plastic stress-strain relation for the material. To determine which theory might be adequate, test results must be obtained for various combined stress conditions, in order to compare the actual with the assumed theoretical behavior. Although there is considerable test data available for combined stress conditions in which the principal stress ratios are constant, relatively little test information is available for vari-

Jan 1, 1954

By John G. Roscoe

Continental Materials Corp. operates two under- ground uranium mines in the Crooks Gap district of central Wyoming. Both mines are developed as slusher mining operations through vertical shafts with single level haulage to slusher pocket loading stations. In general, operations at both mines are similar except for depth and some minor details. The shafts are identical in cross-section, each with a 4 ½ x 5-ft hoisting compartment and a 4 ½ x 5-ft manway. Kimberley-type skips are loaded from measuring pockets through air-operated guillotine gates. The Seismic shaft is 245 ft deep and the Reserve shaft depth is 585 ft. The Seismic shaft was sunk in 1957 and the mine has been producing 200 tpd since 1958. The Reserve shaft was sunk in 1960 and the mine is still in the development stage. The hoists of both mines are equipped for semi-automatic operation; that is, they are mechanically controlled to perform a certain cycle of duty and then stop until another cycle is manually initiated. The Seismic hoist, however, was originally a conventional, manually controlled unit; the reserve hoist was installed as an automated system.

Jan 7, 1962

APRIL 10-Officially, spring comes to the Great Lakes on March 21 as it does elsewhere in the country but in the Superior district continued snow and freezing until late in March have caused citizens in Duluth to place their bets on April 15 for the first ore vessels if the ice breakup is left to Mother Nature. North and west of Duluth and near the eastern end of the Mesabilies the little town of Aurora, population 1100. As the little v-nosed snowplow, pride of the community, shuttled back and forth clearing Aurora's few sidewalks, the crowd in Bradich's Cafe were more interested in the impending high school state championship basketball tournament in Minneapolis than in the outcome of Erie Mining Company's petition for water rights for their proposed 10 million ton taconite concentrating plant. Such a plant would employ 4500 workmen, and at the prevailing ratio of 10 people for every mine worker in Minnesota, this would mean a 45,000 increase in population for Aurora and the vicinity in the next few years. Downstate fishermen are alarmed at the flurry of activity by the mining companies in pushing plans for large scale recovery of iron from taconite. Inconvenience to these gentlemen will be inconsequential when compared with the vast increase in payrolls for range communities. Oliver is pushing the construction of a fine ore agglomeration plant at Virginia which will give them valuable experience for projected taconite processing. Oglebay, Norton & Co. is ready to go at Beaver Bay, north of Duluth, as the right of way for a railroad from Babitt to Beaver Bay is secured as well as permission to dispose of tailing into Lake Superior at the site of the proposed concentrating plant. This company is also doing research in agglomeration at Ashland, Ky., in cooperation with other interests. The financing is yet to be worked out. Of the four possibilities for agglomerating taconite concentrate-sintering, pellitizing, nodulizing and briquetting-the prospects for pelletizing appear the most favorable., It seems likely that 60 cents per ton can be saved by this process over sintering because only 1 to 2 pct of coal is required as compared with 7 pct for sintering. One reason for the lower coal consumption is that the chemical transition of magnetite to hematite which transpires in the pelletizing process is exothermic.

Jan 5, 1950

By D. W. Loucks

There is plenty of evidence to indi-cate that at least one of man's chief interests in life is to make himself as comfortable as possible. If you doubt this, just watch the fellow next to you for the next half hour trying to find the most comfortable position that a hard chair has to offer. Comfort, however, does not always mean an easy chair. To some, it may mean a wealth of money; to another, freedom from worry. But to most of us, it means first of all a comfortable atmosphere in which to live, and to a great many of us it probably also means freedom from that annoying task of firing the furnace. Today more than ever before. automatic heat is one improvement that is placed high on everyone's list. Perhaps this is because automatic heating is becoming relatively cheaper. Perhaps it is because of a good publicity campaign on the part of the oil and gas men or maybe it is just that we are getting lazier day by day. At any rate, almost every issue of Better Homes and Gardens, House Beautiful, or your other favorite home magazine carries an article extolling the virtues of this or that automatic heating system. If I were to ask you to name the first thing that came to your mind when I said automatic heat, you would prob-ably say either gas furnace or oil burner. Or if you had just been studying heating systems, you might possibly say heat pump. But chances are you would not mention anything about coal, and yet coal is the most common source of the greatest automatic heat of them all. I say this because coal is the fuel used almost universally by the district heating industry in producing and delivering to certain heavily populated areas heat ready to use at the touch of a valve or the click of a thermostat. Although the industry is over a half century old, it has not experienced the widespread development of other utility industries because of certain limitations which I believe you will realize from the next few minutes discussion. District Heating Operations We may define district heating as any operation where two or more buildings are heated from a central heating plant. The method of heat transfer may be hot water or in some cases warm air, but generally the medium of heat transfer is steam. So universally is steam used that the industry is frequently referred to as the district steam industry. The Allegheny County Steam Heating Co. which operates the district heating system in downtown Pittsburgh is a subsidiary of the Du-quesne Light Co. Although organized in 1912 primarily as a means of securing the electric load of downtown buildings, the service has now become so valuable and so popular that it is no longer considered a necessary adjunct to the electric business but rather a separate business standing on its own feet. Fig 1 shows the layout of the plants and distribution system of downtown Pittsburgh. Two generating plants, one known as the Stanwix and the other as Twelfth Street, supply the area. Each has two boilers with capacity totaling 1,350,000 lb per hour. The Stanwix Plant is supplied coal by truck. The coal is pulverized at the plant and burned as powdered fuel. Coal is supplied to the Twelfth Street Plant also by truck but the boilers arc stoker fired. Over 1 1/2 miles of tunnel house a portion of our main lines, but it requires over twelve miles of pipeline, ranging in size from 32 down to 1 in. in diameter, to supply all our customers. The distribution system consists of two systems in a sense, one high and one low pressure with certain interconnections between the two. Our high pressure system supplies steam up to 125 Ib to some but not all customers, while the low pressure system operates in the range of 10 to 20 psi. Note that the two plants are tied together through large steam mains and that the system to some extent is a loop system, making it possible to have a portion of the line shut, down without interrupting service to any customer. Fig 2 conveys a picture of the extent to which steam service is used in the downtown triangle. The black area indicates the buildings which now use district steam. The dotted area indi-

Jan 1, 1950

By G. L. F. Powell

g samples of Cu-20 wt pct Ag alloy have been mdercooled to a maximum of 197°C by melting under a slag of commercial soda-lime glass in a vitreous silica crucible. No grain refinement of the primary copper was observed in samples undercooled to the maximum of 197°C. When the samples contained a small amount of oxygen, the copper dendrites were partially recrystallized at undercoolings greater than 97°C. In previous papers'-3 reporting the grain structure of undercooled silver and copper, it was observed that grain refinement was dependent on both undercooling and oxygen content. Grain refinement occurred in undercooled silver when the degree of undercooling exceeded the range 153" to 175"C, while in Ag-0 alloys (0.12 wt pct) fine equiaxed grains were exhibited when undercooling was greater than 50°C. Similarly, copper samples undercooled as much as 208°C displayed fan-shaped growth from a single nucleation site, while the grain structure of Cu-O alloys (0.08 wt pct) was fine and equiaxed at undercoolings larger than 150°C. Thus the presence of oxygen greatly reduced the undercooling at which grain refinement occurred. It was also observed that the change in grain size resulted from recrystallization and was not due to an enhanced nucleation rate in the liquid-solid transformation. It is possible that the influence of oxygen on recrystallization is due primarily to its presence as a solute element. walker4,' reported that, although a grain size change did not occur in pure nickel until the undercooling exceeded 150°C, small grains were observed in samples of Ni-Cu alloy solidified at small and large degrees of undercooling. Jackson et al.6 suggested that the fine grained structure of the Ni-Cu alloy resulted from the melting off of dendrite arms during recalescence. This remelting process may occur in alloys as a result of segregation during freezing which causes a variation in liquidus temperature from point to point within a dendrite. It was therefore decided to undercool copper with a metallic alloying element to ascertain whether the presence of a metallic solute would have a similar effect to oxygen in inducing grain refinement. A Cu-Ag alloy was chosen, since both metals had been shown to behave similarly on undercooling. The alloy Cu-20 wt pct Ag was selected since the eutectic constituent outlines the initial growth form of the primary copper, so that the as-frozen grain structure is not obscured if subsequent recrystallization occurs. This paper describes the results of undercooling experiments carried out with Cu-20 pct Ag samples undercooled to a maximum of 197°C and the effect of oxygen content on the grain structure of the undercooled samples. EXPERIMENTAL Melting was carried out in a small cylindrical resistance furnace using "fine" silver granulate and oxygen-free high conductivity copper. The procedure adopted was to melt the required quantity of silver in air in a clean vitreous silica crucible for approximately 15 min, freeze, and add granulated commercial soda-lime glass to form a complete surface slag cover, after which the sample was melted and frozen several times to reduce the oxygen content. The glass slag cover was approximately 3 in. thick. Pieces of copper (=50 g) were added to the crucible until the required quantity to make 350-g samples of alloy had been charged. Each piece was added quickly to the crucible which was held at a temperature slightly above the melting point of silver. The piece was quickly pushed beneath the glass to minimize oxidation and any oxide coating usually decomposed before the piece had settled down into the silver. After the full quantity of copper had been added, the melt was stirred with a silica rod to hasten homogenization and a Pt/Pt 13 pct Rh thermocouple enclosed in a vitreous silica sheath inserted for temperature measurement. Heating and cooling curves were recorded on a potentiometric chart recorder fitted with a zero suppression unit. The milli-voltage range of the recorder was adjusted so that temperatures could be read to 1°C. Heating and cooling curves were taken every hour until three consecutive readings gave the same solidus-liquidus range, consistent with the solidus-liquidus range for this alloy composition by reference to Hansen and Anderko.7 Metallographic examination of samples frozen at this stage, failed to show any variation in composition from bottom to top of the ingot. Consequently, it was considered that the melt was homogeneous at this stage and undercooling experiments were then car-

Jan 1, 1970

By J. H. Palm

Experiments have shown12 that the formula -S = St - (St - Se')e ne [1] expresses very well the relation between the true stress S and the true form monotonic deformation of plastic metals in single tension and compression (Fig 1). St (asymptotic or final stress), Se' (threshold stress) and ne (characteristic or specific strain) are constants, the significance of which immediately follows from Fig 1. Only at the initial stages of deformation the experimental curve runs below the theoretical curve. This may be due, at least partly, to the uneven stress distribution caused by the anisotropy of the crystals. It is now obvious that a similar formula H = Ht- (Ht- He)e ne' [2] might be valid for the relation between the hardness and the strain. The hardness can be only a function of the strain and must therefore be independent on the hydrostatic tension present in the neck during straining. Hence, if Eq 2 is correct it will also represent the relation between hardness and strain in the middle section of the neck, the strain of which is generally assumed to be almost or fully homogeneous. For aluminum, copper, and several copper alloys it is indeed observed that Eq 2 agrees very well with the experimental Vickers hardness-strain relation from zero strain on up to the strain at fracture. Moreover nc' is, for several metals fairly equal to nc- Hence H = C1 + C2S [3] C1 and C2 are constants during uniform uniaxial tension. Eq 3 was also obtained for copper by Voce.3 According to Hencky's theory4 the ratio of the Brinell or Vickers hardness and the yieldstress (expressed in kg per mm2) of a non-strainhardening metal is approximately equal to 2.8. If Eq 1, 2 and 3 are correct the ratio Ht/S, must accordingly be the same for all plastic metals and be equal to 2.8. It is now experimentally established that this ratio varies between 2.7 and 3.1 with a mean value of 2.9. Taking into account that Ht and especially St can be obtained only by extrapolation from the range of rather low strains, the agreement with the theory is very satisfactory. References 1. J. H. Palm: Appl. Sci. Res. Al (1948) 198. 2. E. Voce: Jnl. Inst. of Met. (1948) 74, II, 537. 3. E. Voce: Metal Treatment (1948) 15, 53. 4. H. Hencky: Zlsch. f. angew. Math. and Mech. (1923) 3, 241.

Jan 1, 1950

By T. F. Bower, M. C. Flemings

Results are reported of a study of surface dendrilic structure of an Al- Cu alloy solidified against a chill wall. Most primary and secondary "arms " in the surface dendritic structure are arranged orthogonally, giving the impression of strong preferred orientalion on the surface. However, no such preferred orientation exists and it is therefore evident the arms do not represent (100) directions. The primary arms are shown to be interseclions of a (100) plane wilh the chill plane, or, equally often. the projeclion of a (100) direction on the chill plane. Secondary dendrite arms are usually within a few degrees of 90 deg to the primary arm, independent of grain orientalion. Prirary, secondary, and higher-order surface dendrite arms almost always represenl intersections of (100) platzes with the chill surace, or pvojections of (100) direclions. Growlh of secondary arms is favored on the side of the primary arm where a (100) direclion points toward the chill surfAce a1 a Lou, angle. Surface dendrile arms are often observed to be bent. In these cases, the crystal lallice changes orientation; bending is concave to the chill surface. In a previous paper,' a technique was discussed whereby large grains can be obtained at a chill surface. The technique used involves quickly drawing superheated liquid A1-4.5 pct Cu alloy into a thin copper mold, so that the mold is full well before solidification begins. The chill surfaces employed are polished copper blocks coated with amorphous carbon. Shrinkage during solidification between dendrite arms and grains delineates both, without the need for polishing or etching of the cast surface. The grain structure of the chill surface was discussed in a previous paper;' in this paper, the dendrite arms within each grain are examined. Previous work on surface dendrites includes that of Edmunds, who studied the development of preferred orientation in zinc, cadmium, and magnesium.' In zinc and cadmium, he found that the surface region has a (0001) texture (parallel to the chill surface). Walton and Chalmers reasoned that, since the fast growth (1010) directions are in the basal plane, nuclei which have this plane parallel to the mold wall would produce larger grains than nuclei with other orientations. Hence, the texture observed is as expected.3 The same authors, in measurements on aluminum ingots, found no preferred orientation at the mold wall. However, the X-ray technique they used measured the preferred orientation in terms of grain numbers, not grain areas; larger grains were weighted equally with small ones. No preferred orientation is expected on this basis at the chill surface. In a later paper,' Edmunds stated that experiments show a random grain orientation at the surface in die cast aluminum; his technique, also used in his earlier paper, takes account of grain area. Little work has been published on the dendritic structure of metal chill grains. Recent work of Biloni and Chalmers on "predendritic growth" shows the change in morphology from spherical to dendritic during the initial stages of freezing, 5 but this work did not include detailed examination of the fully developed dendrites. Other pertinent work includes that of Lin-denmeyer, who investigated the growth of ice dendrites. 6 When growth was on a substrate, the dendrite axes were bent. The bend corresponded to a change in orientation of the crystal lattice and occurred in such a way as to align the basal plane to the substrate. DENDRITE STRUCTURE Fig. 1 shows the chill surface of a typical casting poured above the critical temperature necessary to produce coarse grains. A cursory examination of these grains shows that the surface dendrite arms within most of the grains are oriented roughly perpendicular to each other. One is tempted to assume that these are (100) directions and that, therefore, marked preferred orientation exists at the chill face. This, however, is not the case. Each of the grains in the casting of Fig. 1 was separately identified, Fig. 2, and its orientation determined by the Laue back-reflect ion method. Results are given in Fig. 3 and it is seen there that no preferred orientation exists. Even when grain area is accounted for, there is no significant preferred orientation. The relationship between surface grain structure and crystal orientation was then obtained by assigning X and Y axes to the casting surface, Fig. 1, and assigning the same axes to the stereographic projections of each grain. Thus, the visible surface structure could be compared readily with grain orientation. This was done for fifty-five of the grains of Fig. 1. Results of this study on three typical grains are described below, and some general observations given subsequently. Fig. 4 shows the structure and stereographic projection of a grain which lies near the (100) zone (with respect to the casting surface). The X and Y directions are marked on the projection, and the photomicrograph mounted with the same orientation. Poles of the stereographic projection represent crys-tallographic directions in the grain which point out of the casting, toward the chill. Two (100) directions are shown in Fig. 4. A line joining the center of the projection and a pole represents the projection of the pole onto the X-Y plane (chill surface). Two such lines are shown in Fig. 4 (solid lines). A line joining the intersection of a great circle with the circumference of the projection gives the trace of a crystallo-graphic plane in the chill surface; two such traces are shown (dashed lines).

Jan 1, 1968

By J. D. Livingston, H. E. Cline

The lamellar-dendritic transition in Sn-Pb alloys near the eutectic composition has been studied at high growth rates. Lamellar structures were found over a substantial range of tin-rich compositions, and this range extended to increasingly tin-rich concentrations as growth rate increased. These results are discussed in terms of stability and competitive-growth arguments. Various experimental and structural limitations to the rate of directional solidification are discussed. The rate of heat removal at the heat sink is the major experimental limitation. ReCENT interet1,2 in the use of fine composite structures produced by directional solidification of eutectic alloys makes it important to determine the range of composition and growth conditions that yield such microstructures. Because increasing growth velocities produce increasingly finer microstructures, it is particularly significant to determine the factors limiting the rate of solidification. Mollard and Flemings3 have shown that composite structures, free of primary dendrites, can be obtained in Sn-Pb alloys of off-eutectic composition. The composition range of composite structures was found to increase with increasing values of G/V, where G is the temperature gradient and V is the growth velocity. These results are in good quantitative agreement with an analysis of the stability of a planar eutectic interface.4 This analysis specifically predicts that over a small range of compositions stable lamellar structures will be obtained even for G/V = 0, hence, even at very high growth rates. The lamellar-dendritic transition in Sn-Pb alloys has also been analyzed with a model based on competitive growth between dendrites and the composite structure.576 This treatment, based on earlier work on organic eutetics,7 predicts that the composition range yielding composite structures in Sn-Pb will increase rapidly at high growth rates. An increase in the composition range of composite structures at high growth rates was recently observed in Cu-Pb alloys near the monotectic composition.8 In view of these results, and the predictions of the stability and competitivegrowth analyses, it was decided to study the lamellar-dendritic transition in Sn-Pb alloys at high growth rates. EXPERIMENTAL Using 99.999 pct pure materials, a series of Sn-Pb alloys were prepared containing 16.8 at. pct to 27.6 at. pct lead. (Eutectic composition is 26.1 at. pct Pb.) Ingots were extruded to 0.175 in. rod, and some rod was drawn to 0.070-in. wire. Directional solidification was accomplished in two different ways, Fig. 1. For growth rates up to 2 x 10-1 cm per sec, a 0.175 in. diam sample was placed in a graphite crucible 5 in. long with 0.250 in. OD and 0.035 in. walls. Samples were melted under flowing argon in a vertical, platinum-wound furnace, and solidified by driving the crucible downwards through a \ in. hole in a water-cooled copper toroid, Fig. l(a). An insulated chromel-alumel thermocouple was imbedded in the center of a representative sample, and moved with the sample during solidification. The local temperature is plotted against the distance travelled by the sample in Fig. 2. As the growth rate increased, the solid-liquid interface moved closer to the water-cooled toroid and the temperature gradient increased. At growth rates above 10-1' cm per sec, heat was not removed fast enough and the sample moved into the toroid in the liquid state. The curve for V = 2 x 10-1 cm per sec shows a plateau caused by incomplete removal of latent heat from the interface, a problem which will be discussed later. To improve the heat removal, the toroid was cooled by nitrogen gas precooled in liquid nitrogen. This allowed successful solidification at rates up to 2 x 10-1cm per sec. Higher solidification rates required still more effective heat removal. Samples 0.070 in. in diam were placed in graphite tubes 0.125 in. in diam with 0.020 in. walls. Instead of cooling by sliding contact with a cooled toroid, these thinner samples were sprayed or directly immersed into water, Fig. l(b). After solidification, samples were stored in liquid nitrogen until they could be examined metallographic-ally. The surface was prepared with a diamond-knife microtome, followed by a light etch. The presence or absence of tin dendrites, Fig. 3, or lead dendrites, Fig. 4, was noted by optical microscopy, usually of a transverse section near the center of the sample. Replicas of the surface were prepared and examined in an electron microscope to resolve the fine lamellar structures, Fig. 5. The structures observed at various compositions and growth rates are summarized in Fig. 6. Composite structures were observed at increasingly cin-rich compositions as growth rate increased. This transition from dendritic to composite structure with increasing growth rate was also demonstrated by solidifying half a sample at a slow rate and then suddenly increasing the growth rate by lifting the furnace and quenching the sample with a water spray. A longitudinal section of this sample, Fig. 7, shows that the tin dendrites, which extended ahead of the slow-moving composite interface, were bypassed by the composite when the growth rate was increased. The range of composite structures at high growth rates was limited by the appearance of primary lead dendrites on the tin-rich side of the eutectic composition. Observation of representative longitudinal

Jan 1, 1970

By Craig R. Barrett, Jack L. Lytton, Oleg D. Sherby

The types of substructures developed during high-temperature creep of (110)[001]-oriented polycrys-talline Fe-3.1 pct Si were examined by electroetching of dislocation sites. Edge dislocations were observed to accumulate adjacent to grain boundaries and poly-gonize perpendicular to their glide planes to form a "Pile-up" of nearly parallel tilt boundaries. The dislocation density developed within the grains during steady-state creep was found to increase with increasing stress. The formation of periodically spaced groups of edge dislocations along grain boundaries was observed, and it was proposed that these formed to relive bending stresses at grain boundary shear .faults. Grain boundary serrations developed at these faults, and it was suggested that this was a result of localized grain boundary migration at the polygonized groups. THE authors have conducted a series of high-temperature creep tests on (110)[001]-oriented poly-crystalline Fe-3.1 pct Si sheet. The creep specimens were examined at various stages of creep by means of electrolytic etching of dislocation sites using the chrome-acetic acid solution developed by Morris.' Hibbard and Dunn2 have demonstrated that the etch pits produced with this technique correspond to single dislocations. It is the purpose of this paper to present and discuss the dislocation substructures which were formed during high-temperature creep at constant stress. EXPERIMENTAL PROCEDURES The polycrystalline Fe-3.1 pct Si sheet employed in this study contained a (110)[001] texture. This texture was quite strong, and the elastic properties of the sheet closely approached those of a single crystal of iron.3,4 A typical analysis in weight percent for this material is 3.1 Si, 0.1 Mn, 0.003 C,

Jan 1, 1965

By J. J. Jonas, G. Gagnon

SHG zinc was extruded in the temperature range 110" to 350°C and the strain rate range 0.05 to 5 sec-1 The strain rate/flow stress/temperature results were analyzed using a power sinh stress relationship. When temperature changes due to the heat of deformation and to heat losses are neglected, the exponent of the sinh function of the stress is 5.6, and the apparent activation energy of deformation is 28 kcal per mole. When these changes are taken into account, the exponent decreases to 4.7 and the activation energy to 23 + 2 kcal per mole. The corrected stress exponent and activation energy are in very good agreement with published values obtained from creep experiments, and suggest that the hot extrusion of zinc is controlled by a mechanism involving self-diffusion. When the extrusion and creep data are compared using a Zener-Hollomon parameter and either a sinh or an exponential stress term, an appreciable offset is observed, which may be due to the difference in impurity content. For a given set of extrusion conditions the ram speed, maximum pressure, and initial temperature can also be correlated using a Zener-Hollomon parameter and a sinh pressure term. THE deformation of metals at temperatures over about one-half the absolute melting temperature has been extensively studied at creep strain rates. By contrast, relatively little work has been carried out on the behavior of metals at hot working strain rates. Most of the latter investigations have been performed using simulative methods, such as hot torsion and hot compression, in which the friction conditions and temperature rise during deformation may differ appreciably from those existing under industrial conditions. Recently, however, Wong and Jonas1 used a scaled-down industrial process to determine the stress and temperature dependence of the strain rate during the extrusion of aluminum. In such tests, the effects of friction and adiabatic heating are closer to those produced in industrial operations. Also, with regard to the testing of materials of limited ductility such as zinc, hot compression and hot torsion do not permit the attainment of true strains as large as the deformations usually applied commercially. The present study was undertaken to investigate the extrusion behavior of Special High Grade (SHG) zinc. The detailed objectives were: 1) to establish the stress and temperature dependence of the strain rate with and without a consideration of adiabatic heating, 2) to study the pressure and temperature dependence of the ram speed, and 3) to investigate the microstruc- tural changes occurring during the deformation. The last aspect of the investigation will be covered in a separate paper. The treatment described below differs from that of Wong and Jonas in that the adiabatic temperature rise during deformation is taken into account, and the calculation of the mean equivalent strain rate is based on the redundant as well as on the homogeneous work. EXPERIMENTAL PROCEDURE Rods from two shipments of continuously cast SHG zinc* were used in the investigation. The composition *Supplied by courtesy of Cominco Ltd. range of the impurities present, as given by the supplier, was: Pb: 0.0013 to 0.0023 pct, Fe: 0.001 pct max Cd: 0.0001 to 0.0006 pct, Cu: 0.0002 to 0.0005 pct, Ti: 0.0001 pct max; thus, by balance, zinc valued from 99.9963 to 99.9966 pct. The as-received rods were machined into billets having a nominal diameter of 1.56 in. and a nominal length of 1½ in.; longer billets up to 4 in. in length, were also used to investigate special aspects. The as-machined rods were annealed at 400°C for 24 hr and slowly cooled. This treatment produced a columnar grain structure, with a grain size of about $ by 2½ cm which was appreciably larger than the as-cast one. A 150-ton, direct extrusion, vertical press was used. Ram displacement and force were recorded continuously against time. A constant flow control valve permitted the maintenance of a range of preselected ram speeds up to in. per sec. The selected speed was held constant, irrespective of the required force, as long as the load developed was below the maximum available. Strain gages were used to determine the force; the gages were calibrated before and after each testing period with a 200,000-lb capacity load cell. Further details of the experimental equipment can be found in Ref. 2. The billets were preheated for 90 min in the extrusion container, which was well insulated so as to minimize temperature gradients. This period was sufficient for the billet to reach a uniform temperature at all temperatures between 110" and 350°C. A square-shoulder die having a 0.290-in. diam central hole was used. The extrusion ratio was 30 to 1. This is equivalent to a reduction in area of 96.7 pct, an elongation of 2900 pct, and a true strain of 3.4. The ram speed was varied over two orders of magnitude from 0.0027 to 0.26 in. per sec. The ram was water-cooled during most of the tests, although some experiments were conducted with a preheated, uncooled ram. All extrusions were run without lubricant, which resulted in conditions of sticking friction. EXPERIMENTAL RESULTS Stress Dependence of the Strain Rate Neglecting Adiabatic Heating. The maximum force required to extrude is given in Table 1 for each of the five initial

Jan 1, 1970

By J. H. Swisher

The rate of internal sulfidation of austenitic Fe-Ti alloys in H2S-H2 gas mixtures is controlled primarily by sulfur diffusion, with counterdiffusion of titanium playing a minor role. At temperatures below 1100°C, enhanced diffusion along grain boundaries becomes important. The rate of internal sulfidation at 1300°C is approximately equal to the rate computed from the sulfur diffusion coefficient. The diffusion coefficient of titanium in y iron has been determined from electron microprobe traces in the base alloy near the subscale interface. The solubility of sulfur in Fe-Ti alloys has been measured in the temperature range from 1150° to 1300°C. The equilibrium sulfur content is found to increase with titanium content, due to the large effect of titanium on the activity coefficient of sulfur. The Ti-S interaction becomes stronger as the temperature decreases. TITANIUM as an alloying element in stainless steels is an effective scavenger for interstitial impurities, carbon in particular. Titanium is known to form stable sulfides; however extensive thermodynamic data on the Ti-S system are not available. Schindlerova and Buzek1 have shown that the Ti-S interaction in liquid iron is moderately strong. There have been no previous studies of the Ti-S interaction in solid iron. Internal sulfidation of Fe-Mn alloys was the subject of a recent investigation by Herrnstein.2 He found the rate of internal sulfidation to be an order of magnitude greater than predicted from available solubility and diffusivity data. A satisfactory explanation for the discrepancy could not be given. In the present study, the solubility of sulfur in austenitic Fe-Ti alloys was measured using a standard gas equilibration technique. Fe-Ti alloy specimens were also internally sulfidized. The rate of internal sulfidation was measured as a function of temperature and alloy composition. Supplementary electron micro-probe measurements were made to provide additional information on the nature of the internal sulfidation process. EXPERIMENTAL The starting materials were alloys containing 0.12, 0.24, 0.38, and 0.54 wt pct Ti. The alloys were made in an induction furnace by adding titanium to electrolytic iron that previously had been vacuum-carbon-deoxidized. The major impurity in the alloys as determined by chemical analysis was carbon. The carbon content of the alloys averaged about 100 ppm; metallic impurities were presented in concentrations of 50 ppm or less. Specimens were made in the form of flat plates, 0.03 by 2 by 4 cm for the equilibrium measurements and 0.5 by 1.5 by 3 cm for the rate measurements. The experiments were performed in a vertical resistance furnace wound with molybdenum wire and containing a recrystallized alumina reaction tube. In the gas train, flow rates of the reacting gases were measured using capillary flow meters. The source of H2S was a mixture of approximately 2 pct H2S in H2, which was obtained in cylinders from the Matheson Co. A chemical analysis was provided with each cylinder. The H2-H2S mixture was diluted with additional hydrogen to obtain the desired ratio of H2S to H2, and the resulting mixture was diluted with 30 pct Ar to minimize thermal segregation of H2S in the furnace. Argon was purified by passage over copper chips at 350°C and subsequently over anhydrone. Hydrogen was purified by passage over platinized asbestos at 450°C and then over anhydrone. The H2-H2S mixture was purified by passage over platinized asbestos and then over Pas. The samples used in the solubility measurements were analyzed for sulfur by combustion and iodometric titration. The subscale thickness in the internally sulfidized samples was measured on a polished cross section, using a microscope with a micrometer stage. Electron microprobe traces for titanium in solution were made on several samples that had been internally sulfidized. A Cambridge microanalyzer was used, and the known titanium content at the center of the specimen was used as a calibration standard. The procedure for the microprobe measurements will be described further when the results are presented. RESULTS AND DISCUSSION Equilibrium Data. Fig. 1 shows the sulfur concentration as a function of gas composition for three alloys equilibrated at 1300°C. The dashed line is based on data published by Turkdogan, Ignatowicz, and pearson3 for pure iron. The breaks in the curves are the saturation points for the alloys. The fact that the initial slope decreases with increasing titanium content indicates that titanium interacts strongly with sulfur in solution. To obtain information on the composition of the precipitating sulfide phase, the measurements described in Fig. 1 were extended to higher sulfur partial pressures. These results are shown in Fig. 2. (The initial portions of the curves are reproduced from Fig. 1.) The highest PH2s /pH2 ratio used is believed to be below the ratio required for the formation of a liquid sulfide phase. Time series experiments were used to study the approach to equilibrium in the samples. It was found that equilibrium with the gas phase was reached in less than 4 hr at 1300°C.

Jan 1, 1969

By J. O. Betterton, D. S. Easton

A detailed investigation was made of the phase diagram of silver-zirconium, particularly in the region 0 to 36 at. pct Ag. The system was found to be characterized by two intermediate phases Zr2Ag and ZrAg and a eutectoid reaction in which the -zirconium solid solution decomposes into a-zirconium and Zr2Ag. It was found that impurities in the range 0.05 pct from the iodide-type zirconium were sufficient to introduce deviations from binary behavior, and that with partial removal of these impurities an increase in the a-phase solid solubility limit from 0.1 to 1.1 at. pct Ag was observed. The phase diagram of the silver-zirconium system is of interest as an example of alloying a transition metal from the left side of the Periodic Table with a Group IB element. Silver would normally act as a univalent metal, its filled 4d-shell remaining undisturbed during the alloying. However, there is a possibility that some of the 4d electrons might transfer to the zirconium. An insight into such a question can occasionally be obtained by comparison of phase diagrams. The silver-zirconium system forms part of a more complete review of various solutes in zirconium in which these valency effects were studied.' Earlier work on the silver-zirconium system was done by Raub and Enge1,2 who investigated the silver-rich alloys. After the start of the present experhents, work on this system was reported by Kemper3 and by Karlsson4 which for the most part agrees with the phase diagram presented here. EXPERIMENTAL PROCEDURE The alloys were prepared by arc casting on a water-cooled, copper hearth with a tungsten electrode and in a pure argon atmosphere. Uniform solute composition was attained by multiple melting on alternate sides of the same ingot. Progressive improvements in the vacuum conditions inside the apparatus during the course of the experiments reduced the Vickers hardness increase of the pure zirconium control ingot from 10 to 20 points, observed initially, to negligible amounts at the end of the experiments. Such hardness changes in zirconium are a well known indication of purity. For example, -01 wt pct additions of oxygen, nitrogen, and carbon increase hardness by 6, 10, and 3 VPN respectively. '9' Further verification that the final casting technique did not add a significant quantity of impurities was obtained when pure zirconium was arc cast and then isothermally annealed in the vicinity of the allotropic transition. The transition was always observed to take place over the same temperature range as in the original crystal bar. The alloy ingots were annealed in sealed silica capsules for times and temperatures which varied between 1 day at 1300°C and 60 days at 700°C. The best method found to prevent the reaction of the zirconium with the silica was foil wrapping of molybdenum or tantalum. With this method, samples of pure zirconium were found to be unchanged in hardness after annealing for 3 days at 1200°C. In most of the experiments the protection of these foils was supplemented by an additional layer of zirconium foil inside the molybdenum or tantalum foil. The alloys, foil, and the capsule were outgassed at pressures in the range 10 to l0-7mm Hg in the temperature range 800" to 1100°C before each anneal in order to remove hydrogen and other impurities, and to provide a suitable container for the high purity, inert atmosphere, which is essential in the annealing of zirconium. The temperature measurements were made with Pt/Pt + 10 pct Rh thermocouples calibrated frequently during the experiments against the melting points of zinc, aluminum, silver, gold, and palladium. For the longer anneals the sum of various temperature errors was generally well within ± 2°C. For short-time anneals and during thermal analysis the overall temperature error is considered to be within ± 0.5°C. The compositions of the alloys from the quenching experiments were determined by chemical analysis at Johnson Matthey and Company, Ltd., under the direction of Mr. F. M. Lever. The actual metallo-graphic samples were individually analyzed in every case, and prior to the analyses two or more sides of each specimen were examined to insure that the specimen was not segregated. The sum of the solute and solvent analyses was in each case within the range 99.9 to 100.1 pct. In the course of the experiments, minor impurities in the range 0 to 500 ppm were found to have significant effects on the zirconium-rich portion of the phase diagram. Similar effects had been encountered previously in other zirconium phase-

Jan 1, 1959

By M. J. Richards, A. Szirmae, G. R. Speich

The formation of austenite from ferrite, ferrite plus retastable carbide, spheroidite, and pearlite has been studied in a series of irons, Fe-C alloys, and plain-carbon steels using fast heating techniques. In the absence of carbide, austenite nucleates at ferrite/ferrite grain boundaries; nucleation is followed by the rapid growth characteristic of a massive transfornation. The trarnsformation occurs at 950°C at heating rates of 106º C per sec and cannot be suppressed. Metastable carbide dissolves before austenite forms and does not influence the transformation kinetics. For spheroidite structures, austenite nucleates preferentially at the jinction between carbides and ferrite grain boundaries. Growth from these centers proceeds until the carbide is completely enveloped; subsequent growth occurs by carbon diffusion through the austenite envelope. For pearlite structures, austenite nucleates preferentially at pearlite colony intersections. Carbide la)?zellae dissolve at the advancing austenite interface but complete solution of carbide does not occur; the residtial carbide is eventually dissolvled or spheroid-ized depending on the carbon cuntent. The magnitude and temperature dependence of the austenite growth rate into Fe-C pearlite when incomplete carbide dissolution is assumed are satisfactorily explained by an approximate colume diffusion model. The impurities present in plain-carbon steel reduce the growth rate of austenite in comparison to that jound in an Fe-C alloy. The formation of austenite has been studied in much less detail than the decomposition of austenite. This is primarily a result of the importance of harden-ability in determining the mechanical properties of steel. Recently, more interest in the kinetics of austenite formation has resulted from the discovery by Grange1 that rapid heating techniques strengthen steel by refining the austenite grain size. Although the strengthening effect is not large, it is accompanied by no loss in ductility. In addition, interest continues in rapid heat treatment of low-carbon steel sheet for tin plate applications.2,3 Among the few systematic studies of austenite formation are the early work of Roberts and Mehl4 on formation of austenite from pearlite and recent work of Molinder5 and of Judd and paxton6 on formation of austenite from spheroidite. Also, Boedtker and Duwez7 and Haworth and paar8 have recently studied the formation of austenite from ferrite in relatively pure iron, Kidin et al.9,10 have studied the formation of austenite in 8 pct Cr steels, and Paxton has recently discussed various aspects of austenite formation in steels." The present work was undertaken to determine the kinetics of austenite formation for a variety of starting structures including ferrite, ferrite plus metastable carbide, ferrite plus spheroidal cementite, and ferrite plus pearlitic cementite. Emphasis was placed on determining the active sites for austenite nucleation, determining the temperature and time range of austenite formation, and in the case of pearlite a careful study of the growth rate of austenite was made in the absence and presence of impurities. By using a variety of heating techniques including laser-pulse heating, it has been possible to study austenite formation in an isothermal fashion over a wide range of temperatures. EXPERIMENTAL PROCEDURE The alloys studied in the present work are a zone-refined iron with 4 pprn C, an Fe-C alloy with 130 pprn C, 2 Fe-C alloys with 0.77 and 0.96 wt pct C, and a plain carbon steel with 0.96 wt pct C. The zone-refined iron and Fe-C alloys contained 60 pprn and 200 pprn total substitutional impurities, respectively. The plain carbon steel contained 2400 pprn Si, 2000 pprn Mn, and 900 pprn Cr. Various heat treatments were given to these alloys to produce different starting structures of equiaxed ferrite, ferrite plus metastable carbide, fine pearlite, and spheroidite. These heat treatments are given in Table I. A wide range of heating rates were employed in this work because many of the reactions occur so quickly at temperatures in the austenite range that they are completed during the initial heating cycle unless very fast heating rates are used. Essentially the same heating techniques employed by Speich et a1.12 and Speich and Fisher13 were used in this work. For time intervals of 2 sec to 20 hr, simple hand immersion of 0.010-in. thick specimens in a Pb-Bi bath was employed. These specimens were quenched in a 10 pct NaC1, 2 pct NaOH aqueous bath. For time intervals of 100 m-sec to 2 sec, an automatic dunking and quenching device was employed with 0.002-in. thick specimens. Again, liquid Pb-Bi baths were used for a heating medium but now helium gas quenching was employed. For time intervals of 2 to 100 m-sec a laser heating device was employed with 0.002-in. thick specimens; a helium plus fine water-droplet spray was now used for quenching. Additional information on heating times shorter than 2 m-sec was obtained by study of the zones around the centrally heated laser spot. Here diffusion of heat from the centrally heated zone raises the temperature of the specimen locally to all temperatures between ambient and the peak temperature, but for times of the order of microseconds. All the heat-treated specimens were examined by

Jan 1, 1970

By William M. Boorstein, Robert D. Pehlke

The rates of solution (of hydrogen in liquid pure iron and in several liquid binary iron alloys were meas-ured using a constant volume technique. The rates of absorption and desorption were found to be equal un-der all experimental conditions. increasing concen-trations of S, Si, or Te decrease the rate of hydrogen uptake but additions of Al, B, Cr, Cu, or Ni have no measurable effect up to concentrations normally en-countered in steelmaking practice. No relation ship was found between the effect of an alloying element on the equilibrium solubility of hydrogen in liquid iron and its effect on the solution rate constant. Mathe-rnatical analysis of the data indicates that under the present experimental conditions the rate of reaction of hydrogen with liquid iron is controlled by transport of gas solute atoms in the metal phase. Comparison of the present resuts with data on nitrogen taken un der similar conditions establishes that the hydrody-nurnic conditions which exist near the surface of a metal bath are best approximated mathematically by a surface renewal model for the case of rapid in-ductive stirring and by a boundary layer model for more quiescent melts. HYDROGEN has long been recognized as being a detrimental constituent in steel. If dissolved in the molten metal in excess of its solid solubility, hydro-gen can be evolved during solidification and cause bleeding or porosity in ingots and castings. In the solid metal, lesser amounts play a definite role in causing other defects such as hairline cracks, blisters, and embrittlement. For significant refinements to be made in metallurgical procedures designed to control or eliminate hydrogen from liquid iron or steel dur-ing processing, available equilibrium solubility data must be supplemented with reliable fundamental in-formation pertaining to the kinetic factors involved in the transfer of hydrogen to or from the metal. The scarcity of such information in the literature prompted the present investigation. PREVIOUS RESEARCH Whereas much of the existing data on the solution kinetics of gases such as nitrogen were obtained during the course of thermodynamic investigations, the solu-tion rate of hydrogen has been found too rapid to be accurately determined by conventional solubility meas-urement techniques. Consequently, little work on hy-drogen solution kinetics has been reported in the lit-erature. Carney, Chipman, and crant1 attempted to study the rate of solution and evolution of hydrogen from liquid iron by employing a newly devised sampling method. Although no significant quantitative data could be obtained, it was observed that the rate of solution was approximately equal to the rate of evolution of hy-drogen from the melt. Karnaukov and Morozov2 stud-ied the rate of absorption and Knuppel and Oeters3 the rate of desorption of hydrogen from molten iron by measuring pressure changes with time in a constant volume system. Karnaukov and Morozov determined the hydrogen pressures over their inductively stirred melts with the aid of a McLeod gage and therefore, were forced to work at pressures not in excess of 40 mm of Hg. Their experimental data conformed to a mathematical correlation based on diffusion control: and the rate coefficients calculated on this basis were shown to be independent of the initial absorption pres-sure. These authors reported the solution rate of hy-drogen to be eight-to-ten times higher than they had found for nitrogen in a previous study. They also re-ported that under identical conditions, hydrogen dis-solves somewhat more slowly in iron-columbium alloys than in pure iron. Knuppel and Oeters found that the desorption of hydrogen from pure iron at 1600°C was controlled in all cases investigated by diffusion in the metal bath as long as bubble formation was sup-pressed. This was substantiated by Levin, Kurochkin, and umrikhin4 who studied the kinetics of hydrogen evolution from liquid (technical) iron while applying a vacuum. Salter5 measured the rate of hydrogen ab-sorbed by iron buttons, arc-melted by direct current, as a function of hydrogen partial pressure in a hy-drogen-argon atmosphere. A carrier gas technique was used for analysis of the hydrogen absorbed. The initial rate of absorption was found to increase di-rectly with the square root of the partial pressure of hydrogen. EXPERIMENTAL METHOD Because of the rapid uptake and evolution of hydro-gen by iron-base melts, a constant volume technique was devised in order to obtain meaningful kinetic data over the entire course of the solution process. Apparatus. A schematic view of the experimental apparatus is given in Fig. 1. The hydrogen-liquid iron reaction system consisted of a gas storage bulb con-nected to a meltcontaining reaction chamber through a normally-closed solenoid valve. The gas storage bulb, an inverted 250 ml round-bottomed Pyrex flask was joined to the inlet port of the solenoid valve by a glass-to-metal seal. A more detailed illustration of the reaction chamber is shown in Fig. 2. The design of the Vycor reaction bulb was essentially that de-scribed by Weinstein and Elliott6 with the exception of a shorter, larger diameter gas inlet for this kinetic study. In position, the reaction bulb was closely by an eight-turn coil of water-cooled copper tubing which, when energized by a 400-kc oscillator, provided the inductive heating source. The walls of the bulb were maintained relatively cool by circulating cold water along their outer surface, thus preventing

Jan 1, 1970