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By Orville R. Lyons, A. C. Richardson

During the past few years there has been a growing demand by coal operators for detailed information about the performance characteristics of the various dryers now being manufactured, preferably in such a form that the information could be used to determine the type of dryer best suited to a particular drying problem. A few operators have been desirous of going even further, because they are of the opinion that present-day dryers are not properly designed and that fundamental studies of dryer design should be undertaken. This growing demand for a method of dryer evaluation and the indicated need for fundamental research resulted in the sponsorship by Bituminous Coal Research, Inc. of a project at Battelle Memorial Institute to investigate the thermal drying of coal. In investigating this problem, the most feasible method of attack appeared to be a survey of active operations to determine the present status of the thermal drying of coal. Data were obtained, usually by actual plant tests with the cooperation of the various preparation engineers and their staffs, for 9 screen-type dryers, 8 rotary-type dryers, 3 continuous-pallet-type dryers, 2 flash-type dryers, I tray-type dryer, and I cascade-type dryer. The testing program provided: (I) data for 24 coal-drying operations, (2) information about the inherent characteristics of the coals dried, (3) information about the previous treatment that the coal had received, and (4) information about operating costs. The data obtained through the dryer survey and relationships determined from the data make it evident that all of the types of dryers included in the survey can be operated with the same degree of effectiveness if each type of dryer is used to dry the sizes of coal for which it is best suited. It should be noted that it is usually undesirable and often impossible to dry different sizes of coal in one type of dryer. By use of the data in the report, any prospective dryer user, having determined the surface moisture percentage and screen-size analysis of the coal to be dried, the tons of water to be removed per hour, and the final surface moisture percentage desired, should be able to determine with a fair degree of accuracy: (I) the difficulty of the drying problem, (2) the approximate cost of drying, (3) the required dryer inlet gas temperature and volume and exhaust gas temperature and volume, (4) the type of dryer to use, and (5) the size of dryer required. The method of testing the dryers and determining the difficulty of drying can be used by operators already having dryers to check the performance of their particular dryer and determine whether they are operating effectively. Considerably more information was obtained from the survey of thermal drying than was anticipated when the survey was first considered. However, a great deal of work still remains to be done in the field

Jan 1, 1949

By L. S. Castleman, H. A. Froot

A study was made of the growth kinetics of B-brass layers in a-brass us y-brass diffusion couples in which both phases were saturated in the 6 phase. It was found that useful estimates can be made of the difhsion coefficients at the ß-phase interfaces from interface displacement data, which are relatively easy to measure experimentally; also, a useful estimate can be made of the mean heat of activation for diffusion from layer growth data. The effects of pressure on layer growth kinetics were explored and found to be consistent with a vacancy model for diffusion in 6 brass in both the disordered and ordered temperature range. 1 HE study of multiphase diffusion receives relatively little attention in comparison with that devoted to diffusion in homogeneous materials. Of the hundred or so papers on diffusion written each year,' only a few are devoted to the subject, and it is the rare review article2,3 that devotes any space to multiphase diffusional phenomena. This is unfortunate, since the metallurgist frequently encounters practical situations involving cementation processes such as cladding, protective coating of metals, and so forth in which he needs guidance of a fundamental nature. The investigation of multiphase diffusion phenomena is complicated by the fact that one must consider reactions at phase interfaces as well as the migration of atoms within the individual phases; these reactions cause the interfaces to be displaced within the diffusion zone. The relative displacement of two adjacent interfaces corresponds to the growth of an intermetallic layer, and it is not difficult to determine experimentally the kinetics of layer growth. In the past few years it has been pointed out4,5 that valuable fundamental information may be gained about diffusion coefficients and the concentrations at interfaces through a study of their motion. Accordingly, it is the purpose of this paper to report the results of an investigation of interface motion in a simple multiphase binary system. In the work reported herein, attention was focussed on the growth of a 6 -brass layer in a diffusion couple prepared from a- and y-brass compositions saturated with respect to the 6 phase. This particular system was chosen for a number of reasons. First, the bcc ß-brass phase is disordered for all compositions at temperatures above 468°C and becomes increasingly long-range ordered for all compositions at temperatures below 454° C;6,7 thus, the effects of a variation in order on interface displacement could be determined. Second, diffusion coefficients have already been obtained as a function of temperature in the ß phase both in the ordered and disordered Re- and a theoretical picture of the mechanism of diffusion in this and other ordered solutions is emerging.11'12 Third, some thermodynamic data are available for the a-, ß-, and ?-brass phases,13,4 and these data could be used to gain insight into questions of phase equilibria and concentrations at phase interfaces.

Jan 1, 1963

By V. H. Gottschalk

SEVERAL years ago a group in the metallurgical division of the U. S. Bureau of Mines began a study of the application of new developments in physics to metallurgical problems'. Among these developments were the extensions of what may be termed by analogy the "sound spectrum." Electromagnetic waves, from X-rays to radio waves, have been quite thoroughly explored for metallurgical use, but the sound spectrum, extending from low-frequency vibrations involved in certain types of metal failure by fatigue through the range of audible sound to inaudible vibrations having frequencies of many kilocycles, has had only meager consideration from metallurgists. Because of the general interest in the process recently demonstrated by St. Clair3 for flocculating smoke by means of sound waves, this article is written to point out possible applications of sound and supersonic (ultrasonic) waves in the metallurgical industry.

Jan 1, 1937

By Frank R. Corwin, Selden S. Rodgers

PAGE I. Introduction,........... 383 11. Brief History of Present MacDoUgall Plant,.... 384 III. First Experimental Work on Increasing Capacity of Furnaces, 385 IV. Effect On MacDouRall Plant of IncreaBed Draft from New Smelter Stack,,.......... V. Increasing the Capacity of the MacDouGall FurnaceS,.. 391 (A). Test of Arthur Crowfoot on Increasing Capacity,.... (a). Description of Test,........ (6). Discussion of Results Obtained from Test,... 393 (B). Investigation Carried Out by John A. Church, Jr.,... 395 (a). Advantages of High Speed,......398 (C). Effect on Regular Operation of Department by Applying Results Obtained from Experimental Work ..... 399 (D). Fundamental Principles Involved in Increasing the Tonnage, . 400 VI. Treatment of an INcreased Amount of Ore Screenings,.. 401 VII. Maximum Percentage oF Ore Screenings in a MacDoUGall Feed, 403 VIII. Tests to Further Increase Capacity of Furnaces,... 406 IX. Tests to Decrease PercentaGe of Flue Dust made by MacDouGall Furnaces,............. (A). Repath-Marcy Furnace......... 408 (B). Modified Hepath-Marcy Furnace,....... (C). Crouse Furnace Designed for Dust Prevention,.... 413 (a). Tests on Crouse Furnace,....... X. Concrete Hearths,.......... 419 XI. Conclusion,............ I. Introduction. Since the first installation of MacDougall roasters at the Great Falls smelter of the Anaconda Copper Mining Co., the capacity of the furnaces.has been more than doubled. During the first nine or ten years of the operation of the MacDougall department the tonnage treated by the furnaces remained practically unchanged. Experimental work carried out in 1906, however, showed that increasing the draft of the furnaces would bring about an increase in capacity. In 1909 the MacDougall department was connected to the new smelter stack and

Jan 1, 1914

By Francis L. Sperry

Up to within a few years, the consumption of nickel has been more directly dependent upon the available supply than that of any of the other useful metals. The Gap mine, in Lancaster county, Pennsylvania,* has been, for the last quarter of a century, the only property in this country furnishing nickel in paying quantities. Its yearly output was about 300,000 pounds of metallic nickel, or nearly half the amount used annually in the United States. Foreign nickel from mines on the New Caledonia islands, in the South Pacific, found entrance into our markets as the production of the Gap mine fell off. The price of nickel was constantly maintained, and no special effort was made to extend its use. Over-production was cautiously guarded against, and all surplus metal was held by the banking-houses of the Rothschilds, who assumed the bonded indebtedness of the Sociètè le Nickel. The opening of the Ontario nickel-mines has, however, brought about a radical change; and nickel from the Sudbury district can be delivered in New York within four days, and in European markets within two weeks, as against two months consumed in transporting South Pacific ores. Former prices have been irretrievably smashed, and European trade-journals comment favorably on the influence which Canadian nickel has had in making lower prices, and breaking the backbone of the " nickel trust." Production and Costs. The quantity of nickel produced and the prices which it commanded may be briefly summarized as follows: The total production of the world from 1840 to 1860 was about 100 to 250 tons yearly of metallic nickel; from 1860 to

Jan 1, 1896

[PRESIDENT AND DIRECTOR H. G. MOULTON3 NEW YORK, N. Y. PAST PRESIDENTS AND DIRECTORS - D. C. JACKLING1 SAN FRANCISCO, CALIF. DONALD B. GILLIES2 CLEVELAND, OHIO TREASURER AND DIRECTOR KARL EILERS3 SEA. CLIFF, N. Y.]

Jan 1, 1940

By Irving Cadoff, Richard Nolder

An analysis oj a series of samples of single -crystal silicon grown on sapphire shows that ,four distinct orientation relations exist. There are at least thirteen crystallographic planes which serve as substrates ,for epitaxial growth The model which assumes that silicon substitutes for aluminium and bonds with the oxygen atoms of the sapphire to form the first later for subsequent growth can account for the epitaxial geometry in all four cases. Multiplici-ties 01 specific orientations are found which are explained by the symmetry of both the sapphire and silicon crystal structures. It was reported in the previous paper' that single-crystal silicon films can be deposited on single-crystal a A1,03 (sapphire) with a wide variety of orientations. The silicon films, however, were found to deposit in only four different orientation relationships with respect to the sapphire substrate. The X-ray data have been analyzed to determine the plane and direction of atomic fit and an atomic model to account for the epitaxial relationships at the silicon-sapphire interface has been developed. These results are described in the present paper. EXPERIMENTAL The Laue X-ray data described in the previous paper were reduced to stereographic projections in a standard manner. These data were supplemented with X-ray reflection data obtained with -& E and A Full-Circle Goniometer (Electronics and Alloys, Inc.) mounted on a Phillips diffractometer unit. The silicon films were thin enough to result in characteristic reflections from the substrate as well as the silicon. The superimposed stereographic projections were analyzed to determine the orientation relationships. Ambiguities were resolved by referring to constructed models of various lattice planes for both silicon and sapphire. These lattice models were also used to develop the model for atom fit at the interface. DATA AND ANALYSIS 1) Crystal Structures. The crystallography of silico-n, which is of the diamond cubic class with a. = 5.4~, is straightforward. The sapphire structure is more complex and generally difficult to represent in pictorial form. It is of the rhombohedra1 class but usually referred to hexagonal axes. The transformation of indices between these two systems can lead to ambiguities which will have bearing on the analysis presented herein. Therefore, the basis used in this paper is described in the appendix. Referred to the hexagonal system, the parameters of sapphire are co = 12.95.&, a, = 4.75~, c/a = 2.73. All parameters are for room temperature. Figs. 1 and 2 illustrate schematically the hexagonal unit cell model of sapphire used for this analysis. The oxygen atoms are assumed to occupy positions in an idealized hexagonal array on basal planes. Their actual positions are in a slightly distorted hexagonal array. The aluminum-atom positions are scaled to their actual positions in the lattice with respect to the basal layers of oxygen atoms. It was found that an idealized model where the aluminum atoms are assumed to lie exactly at the octahedral interstices of the oxygen lattice, similar to that used by ~ronberg,' was not adequate for this study.

Jan 1, 1965

The 114th (New York) Meeting of the Institute will be held in the third week of February, 1917. The Committee on Papers and Publica- tions has set Dec. 1,1916, as the closing date for the receipt of manuscripts intended for this meeting. Manuscripts must be received by the Secretary on or before this date. If they are sent through members of the technical committees they must be sent to those committees enough in advance of the time to be forwarded and in the hands of the Secretary by the time specified. While Dec. 1, 1916, is the final closing date, it is a great advantage to get papers in as much before that time as possible as it enables them to be submitted in proof form to their authors. Papers received on or before

Jan 10, 1916

First Session BRASS Wednesday, A.M., Nov. 29 Co-Chairmen: CARTER S. COLE, Staff Engineer, A.S.T.M. E. A. ANDERSON, Chief of Metals Section, Research Div., New Jersey Zinc Co. (of Pa.) Second Session BRASS Wednesday, P.M., Nov. 29 Co-Chairmen: D. K. CRAMPTON, Director of Research, Chase Brass and Copper Co., Inc. A. W. TRACY, Metallurgist, The American Brass Co. Dinner and Technical Session Wednesday Evening, Nov. 29 Toastmaster: N. L. MOCHEL, Manager Metallurgical Engineering, Westinghouse Electric Corp. Talk on "How Do Things Look Today and For the Future" RALPH KELLY, President, Baldwin Locomotive Works Technical Paper on Stress-Corrosion Cracking E. H. Dix, JR., Assistant Director of Research and Chief Metallurgist, Aluminum Company of America Third Session BRASS Thursday, A.M., Nov. 30 Co-Chairmen: W. B. PRICE, Director of Research, Scovill Manufacturing Co. A. L. JAMIESON, Associate Metallurgist, Frankford Arsenal First Session LIGHT METALS Thursday, P.M., Nov. 30 Co-Chairmen: L. L. WYMAN, Metallurgist, General Electric Co. M. A. HUNTER, Professor of Metallurgical Engineering and Head, Dept. of Metallurgical Engineering, Rensselaer Polytechnic Institute Second Session LIGHT METALS Thursday Evening, Nov. 30 Co-Chairmen: R. T. WOOD, Chief Metallurgist, American Magnesium Corp. J. D. HANAWALT, Director, Metallurgical Dept., Dow Chemical Co. First Session STEEL Friday A.M., Dec. I Co-Chairmen: H. S. RAWDON, Chief, Division of Metallurgy, National Bureau of Standards A. L. FEILD, Director of Research, Rustless Iron and Steel Corp. First Session MISCELLANEOUS METALS Friday P.M., Dec. I Co-Chairmen: E. A. ANDERSON, Chief of Metals Section, Research Div., New Jersey Zinc Co. (of Pa.) CARTER S. COLE, Staff Engineer, A.S.T.M.

Jan 1, 1945

By Mario L. Villa

During 1938, oil activities in the Argentine Republic showed a favorable balance due to the increase in crude-oil production, the incorpora- tion of new producing areas and the discovery of deeper productive horizons. Drilling was as follows:

Jan 1, 1939

By Mario L. Villa

During 1938, oil activities in the Argentine Republic showed a favorable balance due to the increase in crude-oil production, the incorpora- tion of new producing areas and the discovery of deeper productive horizons. Drilling was as follows:

Jan 1, 1939

A Page Adam. H. F. and Ruggles! G. H., Results with Xanthate at Inspiration............. 337 Adnic proertiea of....................... 474 Africa; Gold Coast, manganese...........abs.474 191 mining districts......................... 254 Aid, H., biography........................ 154 Air compressors. In flotation................ 45 Air-gas lift: gae-oil ratios relation of......abs. 475 Gulf Coast. mechanical installations... .abs. 475 Seminole field........................abs. 476 Air-sand process. cleaning coal............abs. 89 Alabema, Birmingham district. drainage. . abs. 86 " Alclad." aluminum alloy................. 395 Alexander. W. Homer. biomaphy .......... 37 Allen. Andrews. Preparing Coal after Mechanical Loaders......:....T................. 181 Allison. F. H., Comparison of the Effect of Nickel and Cobalt in Steel.............aha. 151 Alloys: See Metals named. Aluminum: heat treatment...............abs. 444 machining...........................abs. 443 Norway............................... *465 Aluminum alloys: " Alclad "................ 395 application of........................abs. 443 deformation..........................abs. 150 Aluminum Bronze Manufacturers Institute. organized.............................. 188 "Alumnum. The Metal and Its Alloys." B.R. 42 Aluminum-iron-silicon alloys, equilibrium relations.............................aba. 444 Aluminum-manganese alloys, equilibrium relations in..........................aba. 87 Aluminum-silicon alloys. heat treatment. .abs. 398 American Association of Petroleum Geologists. annual meeting................... 81 American engineers. war memorial.......... 499 American Institute of Mining and Metallurgical Engineers: Committees: Douglas. James, medal...... 324 Geophysical Prospecting............. E. 204 Mining Medal........................ 324 Nominating: Election................. 235 report............................. 479 Constitution. votes for.................E. 494 Constitution and By-laws. draft of....... 374 Directors: Minutes of meeting.....110. 235. 323 publications.......................... 523 Directors adopt badge................... 324 Dues. increasing........................ 111 Dues are Due.......................... 236 Institute of Metals Division: annual lecture............................ 187 data sheets.......................... 187 Detroit meeting, (September). . .307. 353, 440 lecturer C. H. Desch................. 15 Local Section Committee. members..... 187 New York meeting (February)......... 115 Research Committee.................. 187 research problems..................... 6 Local Sections: Arizona............................. 365 Boston, meetings................40, 197. 486 Page American Institute of Mining and Metallurgical Engineers—Continued Chicago, meetings...40, 94, 156, 196,240,417 Chicago and St. Louis, joint meeting, Danville....................... 486, 527 Colorado, annual meeting.............. 285 Columbia, meetings 39, 93, 156, 196, 238, 286, 327, 365, 486, 527 El Paso, organized.................... 238 Fort Worth, organization meeting...... 485 Joplin-Miami Zinc, meeting.......... 286 Lehigh Valley: annual meeting......... 197 meetings...................... 239, 527 Los Angeles, meeting.................. 285 Mid-Continent, meeting............... 156 Minnesota, annual meeting............ 326 Montana, meeting.................... 40 New York, meetings................240, 284 North Pacific: meetings.... 39, 157, 197, 527 Fieldner, A. C, speaker............. 486 resolution.......................... 439 Oregon, officers....................... 196 Pittsburgh: meetings........... 156, 327, 526 mine transportation conference....... 485 St. Louis and Chicago, joint meeting, Danville........................... 486 San Francisco, meetings............... 92, 197, 239, 285, 439, 486, 527 Southern California, meetings 40, 157,197,283 Upper Peninsula, meeting.............. 486 Utah, meetings..........40, 93, 156, 239, 365 Washington, D. C, meetings 40, 93, 365, 528 Medals: Douglas, James................34, *84 Fritz, John, award, 1928.............. 480 Johnson, J. E., awarded T. L. Joseph... 83 Saunders, William Lawrence........... 34 meetings: Cleveland, April, 1927.......166, 207 Detroit, Sept., 1927............307, 353, 440 Fort Worth, Oct., 1927.........307, 353, 471 • future............................. E. 102 Knoxville, Nov., 1927.............. 480, 500 New York, Feb., 1927, announcements, programs, etc.................33, 50, *106 Salt Lake City, Aug...............306, *422 Members and jobs...................... 477 Membership, benefits.................8, E. 2 Officers and directors.................... *122 Open-hearth Conference................. 282 Petroleum Division: activities............ 186 annual meeting...................... . Fort Worth meeting............307. 353 471 luncheon............................ 91 New York Petroleum Club, officers..... 110 nominations........................23. 519 officers.............................. 117 President's trip. May.................281. 398 Publications: Directors discuss............ 523 Institute of Metals Division Volume... . 178 Petroleum Volume.................... 192 Technical Publications...........E. 294 323 Transactions. Vol. LXXIV. LXXV....00. 524 Review of the year.........:........... ...E. 104 Stockholders. report to..............E. 49

Jan 1, 1927

By G. K. Burgess

§ 22. THe results published in Professor Howe's paper10 of our determinations on the Ac3 and Ar3 points for a series of commercial carbon steels " containing manganese in varying proportion, represent the average temperatures found by taking several (4 or 5) heating and cooling curves. The accompanying curves, Figs. 6 and 7, are, for each steel, the first heating and cooling curves as obtained by the inverse-rate method of Osmond, using a cylindrical chronograph sensitive to 0.1 sec. for recording time. Simultaneous observations were also taken by the differential method with a platinum neutral and plotted by Rosenhain's derived differential method and gave substantially the same results, the curves having sensibly the same contours in both cases.12 The observations were taken in vacuo on cylinders 5 by 1 in. in length bored to center with a 1/8-in. hole to receive porcelain insulated thermo-couple, and inclosed in an out- and in-glazed Berlin porcelain tube heated electrically by platinum foil wound on the tube. The vacuum was usually better than 0.1 mm. The heating and cooling of the furnace was automatically controlled by a salt-water rheostat with shaped electrodes, giving a uniform rate of heating and cooling.

Jan 1, 1914

By F. H. Vitovec, D. W. Hoeppner

Pusk-pull fatigue tests were conducted on copper tricrystals of 99.988 pct purity to ascertain the role of grain boundaries in the initiation and propagation of fatigue cracks. Significant differences in behavior were found for specimens which possessed different transverse-boundary misorientation. In speciwens with low boundary angles cracks initiated within the transverse boundary, while higher angles led to transcrystalline fatigue failure. It is suggested that at low angular misorientation moving dislocations may interact with dislocations of the boundary or dislocations present in adjacent pains on favorably oriented glide planes, thus initiating a fatigue crack. MANY fatigue studies have been concerned with fatigue-crack initiation within grains and the mechanism causing initiation and propagation1-3 Although the initiation of fatigue cracks in or near grain boundaries of pure metals has been observed and reported in the literature, the mechanism of this phenomenon has received little attention.4"11 EXPERIMENTAL PROGRAM Testing Procedure. To investigate the role of grain boundaries regarding initiation and propagation of fatigue cracks, copper tricrystals were tested in push-pull. Axial-stress tests were used to avoid the stress gradients introduced by stressing of some other nature. Copper was selected as the most suitable test material since extensive work has been done on the formation of fatigue-induced slip in copper. Tricrystal specimens were used to provide a grain boundary geometry with one boundary transverse to the principal stress axis and one or more boundaries nearly coincident with the direction of maximum resolved shear stress. The boundary energy and the slip characteristics in the vicinity of the transverse boundary depend on the relative orientations of grains across the boundary. Testing was done at room temperature, at a frequency of 700 cpm, in an atmosphere of either air or argon. It is known that the incidence of grain boundary cracking increases with increasing temperature and decreasing frequency.4 The fatigue machine applied a uniaxial tension-compression load to the specimen by a flat spring which was actuated by an adjustable eccentric. Use was made of an adjustable head and a load cell which had strain gages mounted at 120-deg intervals around its periphery, thus providing a means for eliminating any detectable superimposed bending moment. A clip gage was mounted between the gripping heads to record the cyclic strain amplitude applied to the specimen. Each specimen's hysteresis characteristics were recorded by supplying the load and strain signals to an oscilloscope. Microstructural changes were observed and recorded with an optical microscope which was mounted on the fatigue machine. Immediately prior to insertion in the machine, each specimen was chemically polished. Extreme care was exercised while inserting the specimen in the test machine to avoid either bending the specimen or introducing a mean load. Each specimen was stressed in the tensile direction first and subsequently the load was reversed. Specimen Preparation. Tricrystal fatigue specimens of 99.988 pct Cu were grown from the melt using a modified Bridgman technique. Ingots about 1 in. high were grown to the configuration shown in Fig. 1. Upon sectioning the resultant ingot, several specimens were provided with an identical shape and grain boundary orientation. Except for plane sectioning and final polishing, this method eliminated machining the specimens. The apparatus used for growing the tricrystals is shown in Fig. 2. A spectrographically pure graphite mold, Fig. 3, was inserted in a vycor tube which was mounted on a vertical zone-refining table. Prior to insertion in the tube the mold was assembled as follows. The insert and graphite rods were fixed in position in the mold. The copper was then placed in the mold and the entire mold assembly was positioned in the vycor tube, Fig. 2. At this point the tube was sealed, evacuated, purged with helium gas, and finally placed under a slightly positive helium pressure. The heating coil was then adjusted to melt the copper charge, after which it was raised at a speed of 4 in. per hr. It is also possible to use

Jan 1, 1964

By J. S. Pascover, J. K. Abraham

The influence of the early stages of precipitation on the kinetics and structure of martensite formation in Fe-27Ni and Fe-29.5Ni alloys containing from 0 to 10 pct Ti was examined with X-ray and electron microscopy techniques. The formation of a coherent, ordered preprecipitate had a profound stabilizing effect on the austenite. The Ms was decreased by increased titanium content and aging time up to a critical time. When the critical aging time was exceeded, the Ms was observed to increase markedly. The formation of the clusters was insuppressible and the volume fraction of clusters formed during the quench was a function of the titanium content. Martensite resulting from transformation of the clustered austenite is tetragonal with the c/a ratio increasing with titanium content. A model for the tetragonality is suggested. The morphology and substructure of the m artensite is inter-preted in terms of the above information and the cur-rent models of twinned martensite. ThE ramifications of precipitation in austenite to the properties of austenite have been the subject of numerous investigations. The current research is concerned with the influence of precipitation in austenite on the kinetics and structure of subsequent marten-site formation. In a previous investigation, Abraham et al.1 followed the aging reaction in an Fe-29.5Ni-4.2Ti* (at. pct) alloy using an X-ray diffraction technique. This technique, employing a Guinier camera, provided kinetic measurements through observation of the side band position as a function of aging time. The salient results of this work were: 1) The initiation of precipitation was not suppressed by quenching, i.e., there was a finite cluster zone size at zero aging time; and 2) The hardness of the aged austenite correlated extremely well with the zone size. During the previous work it was noted that the mar-tensite formed after aging was tetragonal, substanti-ating an earlier observation.2 Systematic investiga-tion revealed that the martensite was tetragonal in both the solution-treated then quenched, and the solu-tion-treated, aged, and quenched condition, and, furthermore, that a marked stabilization of the austenite occurred as a function of aging time. The present work is concerned with documenting the tetragonality and the stabilization phenomena as well as the ob- served microstruction with a suggested rationale for the behavior noted. EXPERIMENTAL PROCEDURE The compositions of the alloys are listed in Table I. The analyses were performed after the solution treatment of the strip material. Nickel was determined using the standard dimethy1-glyoxime procedure whereas titanium was determined colorimetrically with hydrogen peroxide and volumetrically by titrating with ferric iron. The materials were melted in a 5-lb vacuum induction furnace, cast into 2-in.-diam ingots, and forged in a temperature range of 950. to 1200°C to 1/2-in. slabs. The three higher titanium containing materials cracked during forging; therefore, to get the alloys into strip form, slices 1/8 in. thick were cut from the slab, homogenized 4 hr at 1150°C, then cold-rolled to a 0.04-in. thickness. The remaining slabs were hot-rolled, homogenized 4 hr at 1065"C, then cold-rolled to a final thickness of 0.03 in. All of the heat treatments were performed under a protective atmosphere of argon. The Ms temperature for most of the alloys is below room temperature; therefore, it was possible to solution treat, quench to room temperature, polish, and then observe the transformation optically on a cold-stage microscope. To determine the effect of aus-tenitizing temperature on Ms, eight of the alloys were treated at two temperatures, 1025" and 1120°C. No measurable variations in Ms were noticed. The remaining alloys were treated at 1025°C. The specimens, : by 5/16 by 0.03 in., were austeni-tized in a vertical tube furnace under a dynamic argon atmosphere. The bottom of the tube was submerged in water for quenching purposes. The question of stabilization that may be operating at room temperature was investigated and found to be negligible. Many of the specimens were held at room

Jan 1, 1970

By A. E. Dwight, Paul A. Beck

Known CsCl-type phases in transition metal alloys are shown to be formed preferentially at an electron concentration of approximately 6; a tetragonally distorted version of the CsCl structure is stable at -e/a = 6.5. The relationship of the distortion to electron concentration is verified by experimental data for the quasibinary systems Time ,Os) , Ti(Re,Ir), and Ti(Re,Pt) and the ternary system Ti-Ru-Rh. It was noted some years ago1j2 that the occurrence of CsC1-type ordered structures in equiatomic binary alloys, AB, of transition metals correlates with the location of the component elements in the periodic table. One of the components, A, must be to the left and the other component. B, must be to the right of the chromium group. It was pointed out2 that the CsCl structure does not occur in CrFe. is metastable in VFe,3 and is stable in TiFe up to high temperatures, and that this progressively increasing stability may correspond to an increase in electron transfer between the two components and to a correlated decrease of the magnetic moment associated with the iron atoms. Such a decrease was in fact later demonstrated by Nevitt. 4 The binary CsC1-type phases that are formed by a transition metal with either another transition metal or a lanthanide are listed in Table I. Thirty-three of these phases occur at an average electron concentration (average number of electrons outside of closed shells) of e/a = 6. The five phases that are formed at an electron concentration of e/a = 5.5 appear to be borderline cases. This is suggested for instance by the fact that TiRe is bcc. without any long-range order, while TiTc (at the same electron concentration) has a CsCl ordered structure. Six of the undistorted CsCl phases formed by transition metals with one another have an average electron concentration of e/a = 6. 5. Table I. With the exception of Zr4,Rhs4 and Hf48Rh52. the B component in these phases is a first long-period transition metal, i.e.. iron or cobalt. TiNi. with an electron concentration of e/n = 7, appears to undergo transformation at around room temperature to a more complex ordered struct~re.~ The five other CsCl structures known to occur at e/a = 6.5 have either scandium or a lanthanide metal as the A component; the difference in the elec-tronegativity and in the e/a ratio between the component elements is in these phases very large, A(e/a) - 7, and they may be expected to have a rather strong ionic bond component. Presumably the ionic bond stabilizes the CsCl-type structure at the equiatomic composition and it overrides the electron concentration requirement. When the B element is not from the first long period and the A component is not a rare earth metal. equiatomic (AB) alloys of transition metals with an average electron concentration of e/a = 6.5 in most cases have a crystal structure closely related to the CsC1-type, but slightly distorted tetragonally, Table I. This was first found6 for VRu, NbRu. and TaRu. It was noted6 that the undistorted CsCl structure is also stable in the three systems. but at a lower electron concentration, near e/a = 6. In these systems this requires a deviation from equiatomic stoichiometry. which is otherwise ideally suited for the CsC1-type structure. A similar situation was later found in the Ti-Rh7 and the Ti-Ir systems. 8 The Ti-Rh alloys undergo a second distortion and become orthorhombic when the rhodium content becomes somewhat larger than 50 pct. As mentioned previously. TiRe does not have the CsC1-type ordered structure. In order to investigate how close TiRe is to the borderline of stability of the

Jan 1, 1970

By R. F. Mehl, R. F. Bunshah

A fast amplifier technique has been developed for the measurement of the rate of propagation of martensite in an Fe-29.5 pct Ni alloy. The time of formation of one plate of martensite is 3x10 sec and the rate of propagation is 3300 ft per sec approximately. IT has been known for some time that the plate-like structural unit of martensite forms from austenite with great rapidity. Wiester1 and Hane-mann, Hofmann, and Wiester took motion-pictures of the transformation as it occurs in a 1.65 pct C steel; they demonstrated that a single plate formed fully in the time interval between successive frames, viz., 1/20 sec, thus setting an upper limit. Forster and Scheil,3 using an Fe-Ni alloy with 29 pct Ni, recorded the sonic characteristics of the process electrically, upon an oscillograph, setting the upper time limit at 0.002 sec. Forster and Scheil,~ measuring the change in electrical resistance in the same alloy upon a cardiograph, set a limit of 0.02 sec. Forster and Scheil5 later, employing the same alloy, improved their technique, reporting an upper limit of 7.10 sec. In studying signals of such short duration, it is an important question whether the frequency response of the electrical system used is high enough compared to frequency of the pulse measured, or, put differently, whether the system is able to reproduce without distortion the signal arising, in this case, from the formation of a single martensite plate. Forster and Scheil (referring only to their last paper) obtained signals of a frequency of 30 kilocycles (hereinafter kc); this was about the frequency response of the equipment used; thus, if the signal had a frequency higher than 30 kc, it would still appear as a signal of frequency 30 kc. All of these results thus provided upper limits only. Recent developments in electronics have made available equipment with very high frequency response, very high sweep-speeds, high gain, etc. The electrical characteristics of such equipment, used in the present study, are given in Table I. Such equipment offers obvious attraction in the study of the rate of propagation of a martensite structural unit—and perhaps of other structural alterations proceeding at a very high rate. This paper reports an attempt to develop a technique employing such equipment to measure the time of propagation of a martensite structural unit and the variation of this with temperature, with the mode of formation—athermal and isothermal—in both polycrystalline and single-crystal samples; and from such measurements to obtain the rate of propagation. As will be seen, the results obtained are useful theoretically. Materials All data presented here are for an Fe-Ni alloy of the following analysis: 29.5 pct Ni, 0.027 pct C, 0.135 pct Mn, 0.094 pct Si, balance Fe. There were several reasons for choosing this alloy: 1—it is substantially the one used by previous investigators; 2—it exhibits both the athermala and the isothermal' mode of formation of martensite, both studied in detail by Machlin and Cohen; 3—the subzero temperatures of transformation in this alloy are experimentally very convenient; 4—it exhibits the "burst phenomenon";" 5—the change in electrical resistance upon the formation of martensite, a decrease, is great, approximately 50 pct.' The polycrystalline specimens were in the form of wires of 0.025 in. diameter; the single crystals were 1x1/4x1/4 in. Experimental Methods Electrical Apparatus: Fig. 1 is a schematic drawing of the electrical circuit used. The principle used in these measurements is the same as that used by Forster and Scheil." A small direct current, about 1 to 2 amp, is passed through the sample. When a martensite plate forms, the resistance of the sample changes and a high frequency signal is generated. This signal is picked up by the probes attached to the sample, fed into the bank of amplifiers and thence to the vertical deflection plates of a cathode-ray oscilloscope. The signal itself triggers the oscilloscope trace which flashes across the tube face and is photographed by means of a 35 mm movie camera at the end of a light-tight hood. The camera has no shutter. As soon as the signal flashes

Jan 1, 1954

By D. W. White

The kinetics of isothermal transformation of ß-to-u uranium have been studied over a broad temperature range in alloys containing from 0.3 to 4.0 atomic pct Cr. Two modes of transformation are indicated by the existence of two C-curves in the TTT diagram. The upper temperature mode is regarded as a nucleation and growth mechanism, whose rate is controlled by diffusion of chromium in the ß phase matrix. The lower temperature mode is martensitic in nature. The M, temperature increases with decreasing chromium content, suggesting that the two transformation processes become synonymous in unalloyed uranium. URANIUM metal undergoes two allotropic transformations in the solid state. The a phase, orthorhombic in crystal structure,' is stable from room temperature up to about 665°C. The ß phase, characterized by a complex tetragonal structure,' prevails from 665" to about 770°C. The y phase is body-centered-cubic3 and is the stable modification from 770°C up to the melting point (about 1130"). In uranium of reasonable purity, neither of the two high temperature phases can be retained by quenching. However, the addition of certain alloying elements to uranium makes it possible to retain either the y-uranium phase or the ß-uranium phase at room temperature. Chromium alloyed in small amounts with uranium will permit retention of the ß-uranium phase in a metastable state at room temperature upon quenching from a ß-phase temperature.' From available information' on the equilibrium phase diagram for the U-Cr alloy system (Fig. I), it is to be expected that, however sluggish in its rate, the ß phase in such alloys should decompose eutectoidally to a phase and elemental chromium. It was the aim of this investigation to measure the rate and study the nature of this decomposition as a function of temperature and of chromium content. The investigation was reported in classified literature about five years ago and has recently been declassified for publication. In the meantime, there have appeared the papers of Holden,5 Mott and Haines,".' and Butcher and Rowe8 ealing with the metallography and the crystallography of the ß-to-a transformation in U-Cr alloys. These investigators have confirmed several of the phenomenological observations that will be described in the present paper and have examined in considerable detail certain aspects of the transformation and its mechanism. Although all of these investigations have concerned themselves experimentally with U-Cr alloys for the most part, an important consequence has been a clearer understanding of the nature of the ß-to-a transformation in uranium metal itself. Experimental Procedure This investigation dealt with a series of uranium alloys varying in chromium content from 0.3 to 4.0 atomic pct (0.066 to 0.90 weight pct). On five of the alloys, rates of isothermal transformation from the ß to the a phase were measured over a wide temperature range, leading to the development of TTT (time-temperature-transformation) diagrams. Transformation rates were measured over certain narrow temperature ranges on additional alloys. The alloys were prepared by vacuum melting and casting, using zircon or magnesia crucibles and graphite molds. Electrolytic chromium was used as the alloying addition, and the uranium was Mallin-ckrodt biscuit metal that had been vacuum remelted and cropped to remove many of the nonmetallic impurities that had floated to the top of the ingot. The ingots, 3/4 or 1 in. diam, were reduced in size by swaging. Alloys containing less than about 2 atomic pct Cr were swaged at 250° to 275°C, with initial and intermediate anneals at 550°C after every 75 pct reduction in area. Alloys with higher amounts of chromium were swaged at 550" to 600°C, although at the smaller sizes some of them were reduced by the procedure used on the more dilute alloys. Before use as test specimens, the swaged rods were annealed at 700" to 720°C for several hours, followed by slow furnace-cooling. The purpose of the anneal was to achieve the maximum amount of solution of the available chromium into the ß phase, as well as to remove extensive preferred orientation. The isothermal transformation rates were measured dilatometrically, using a quenching dilatometer and an experimental technique similar to those employed by Davenport and Bain in their original work on the transformation kinetics of austenite in

Jan 1, 1956

By J. W. Spretnak, V. Kandarpa

Stress -induced ordering of carbon atoms is studied in a series of Fe-Mn-C alloys. A prominent peak is found in the vicinity of 280°C at frequencies of the order of 1.0 cps, with an associated activation energy of 37 kcal per mole. The height of the peak is linearly rekzted to the concentration of carbon in solution. The distortion of octahedral holes by manganese atoms appears to be predominant over carbon-carbon pair interactions. RELAXATION by stress-induced ordering of point defects is expected whenever the introduction of these point defects produces distortions which have a lower symmetry than that of the lattice. Under zero stress, the isolated point defects occupy the crystallographic-ally equivalent positions in the lattice, as these represent states of equal energy. However, if the defect sites are asymmetric, application of an uniaxial stress will split the energy states, and a redistribution of the defects among various states will take place. This is the case of the internal friction peak called the Snoek peak,1 resulting from isolated interstitials in bcc metals. The interstitial sites in this case have tetragonal symmetry. In the case of fcc and hcp lattices, such an effect is not expected from isolated point defects because of the symmetrical nature of the interstitial sites. However, internal friction peaks arising from interstitial diffusion have been reported both in hcp2,3 and fcc4-8 lattices. These peaks are often explained on the basis of stress-induced ordering of interstitial solutes, caused by the deviation of interstitial sites from their cubic symmetry through the presence of nearby defects. In the case of fcc lattices, evidence for interactions of both the substitutional-interstitial4,6,13 and interstitial-interstitial types5'798'14 have been presented by various investigators. The purpose of the present investigation was to study the internal friction peak attributed to the diffusion of interstitial carbon atoms in high purity austenitic manganese steels and to account for the peak in the light of the existing models. MATERIALS The Fe-Mn-C alloys used in the present investigation were made in two different ways, designated as Type I and Type 11. Type I alloys were made from high purity Fe-Mn alloys obtained in the form of 0.04- in.-diam wires from Materials Research Corporation, Orangeburg, N.Y. These alloys were carburized to different levels using gas mixtures of H2 and CH4 at 1000°C. Type I1 alloys were made in this laboratory starting with zone refined iron, spectrographically pure manganese, and spectrographically pure carbon. They were melted in an argon arc melting furnace and drawn into 0.04-in.-diam wires. All the wires were annealed at about 900°C for 3 hr prior to the internal friction experiments. After the measurements of internal friction, the phases in the samples were identified by X-ray diffraction and the carbon determined by the combustion method. EXPERIMENTAL PROCEDURE In the present work, a classical Ke-type pendulum was used. The details of the equipment were described previously by D. T. Peters.9 Dry helium at 40 torr was used in all the experiments. The internal friction, measured as the logarithmic decrement of the torsion amplitude of vibration was determined as a function of temperature, from ambient to about 500°C. The background internal friction was assumed to have the form of the exponential of the inverse temperature and was subtracted from the raw data. The height of the peak was measured at the position of the maximum in the plot of the internal friction versus temperature. The activation energies of the peaks were measured by the peak shift method. The internal friction values for an alloy were obtained as a function of temperature at different frequencies of vibration. The position of the peak changes with frequency, the higher the frequency the higher the peak temperature. The activation energy of the process associated with the peak is obtained using the formula

Jan 1, 1970

By H. C. Fiedler

The development of cube-on-edge secondary re crystallization texture in Si-Fe strip depends upon the ability of inclusions, such as manganese sulfide, to restrain nomal grain gvowth. The ability of inclusions to perform this function is shown to depend upon their size and distribution. The desired uniform dispersion of many small inclusions is obtained by controlling the cooling rate from above the solution tempmature of the inclusions or, alternatively, by drastic 'quenching from above the solution temperature followed by heat treating at a lower temperature. The development of the cube-on-edge secondary recrystallization texture in Si-Fe strip depends upon both establishing the proper primary recrystallization texture and having inclusions to restrain normal grain growth. It has been suggested that the first requirement is met with a texture which, though weak, is near (110)[001]. A mechanism for secondary recrystallization originally proposed by ielsen' is consistent with the presence of this primary texture. It is visualized that a secondary grain appears where two adjacent grains happen to be sufficiently similar in orientation to be capable of behaving as a single large grain. This pair of grains will grow at the expense of its smaller neighbors as a consequence of its having lower grain boundary free energy. Other explanations of the origin of the secondary recrystallization texture have been offered,' and it can be concluded that the mechanism is not fully understood. Nevertheless, it is apparent that operationally the necessary primary recrystallization texture is obtained simply by interposing an anneal during the cold reduction to final gage. The role of inclusions in the development of the secondary recrystallization texture has been demonstrated by May and Turnbull4 in a comparison of a high-purity heat and a heat containing manganese sulfide. They showed that the driving force for secondary recrystallization is inversely proportional to the diameter of the matrix grains and that the function of the inclusions is to maintain a small matrix grain size. Other inclusions, such as various sul- fides,')' silicon nitride,8 and vanadium nitride,' titanium carbide,'' and vanadium carbide," have been shown to be effective in promoting secondary recrystallization to the cube-on-edge texture. There is also evidence that silica inclusions in Si-Fe strip made by powder-metallurgy techniques perform this function.' This report is concerned with the second of the two requirements for secondary recrystallization, namely the importance of the metallurgical control of the size and therefore number of inclusions. While the concepts involved are simple, they are fundamental to the making of this material. EXPERIMENTAL PROCEDURE A heat containing 3.27 pct Si, 0.026 pct S, 0.06 pct Mn, and 0.004 pct C was made in an air induction furnace from electrolytic iron, 98 pct ferro-silicon, electrolytic manganese, and iron sulfide. The heat was divided between graphite and baked-sand molds with a cavity 2-5/8 by 5 in. and with walls 2-3/4 in. thick. Slabs, 1 in. thick and 2-5/8 by 5 in., were cut from the ingots and hot-rolled either in the as-cast condition or after one of the heat treatments described later. After pickling, the 80- to 100-mil-thick hot-rolled band was, unless otherwise noted, heat-treated for 5 min at 900°C before being cold-rolled to 25 mils. After a heat treatment of 1 to 3 min at 860°C, the strip was cold-rolled to 12 mils, the final thickness. The percent cube-on-edge texture was determined by dividing the maximum torque as averaged from four samples by the maximum torque of a 3-1/4 pct Si-Fe single crystal having the (110) plane in the plane of the sample. The torque measurements were made in a field of 1000 oe on 1-in.-diam disks. RESULTS AND DISCUSSION The importance of the thermal history of the ingot is evident in Fig. 1 from the appearance of strips heat-treated in a temperature gradient. The sample

Jan 1, 1964