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By B. Post, F. W. Glaser

Three borides of zirconium have been reported: ZrB,l ZrB2,2,3 and ZrB12, 4 The phase relationships, ranges of stability, and some physical properties of these compounds are described. THE zirconium borides differ considerably from one another, particularly with respect to their relative stabilities at high temperatures in the presence of carbon.4,5 ZrB2 is extremelv stable at all temperatures (to the melting point) in contact with carbon. Both ZrB and ZrB12 tend to decompose at elevated temperatures when carbon is present in excess of about 0.5 pct. ZrB forms ZrB2 and ZrC, and ZrB12 decomposes to form ZrB2 and a boron-carbide phase. In the ZrB12 preparation, boron present in excess of that needed to form ZrB2 reacts with carbon, if present, to form carbides, in preference to forming the higher boride. Similar instability of borides, other than diborides of the transition metals of the 4th and 5th groups of the periodic table, has been reported.5,6 special precautions had, therefore, to be taken to minimize carbon contamination in the preparation of ZrB and ZrB12. Reactions below 1400°C were carried out by heating powdered mixtures of zirconium hydride and boron in quartz tubes in a hydrogen atmosphere. For temperatures above 1400°C, reactions were carried out in graphite molds. Carbon contamination from the molds was minimized by employing extremely rapid heating rates to reduce carbon diffusion into the body of the specimen, and by subsequent "shaving" of sample surfaces that had been in contact with the graphite dies. A heating rate of about 2000°C per min was obtained by heating the dies (by direct conduction), using a current of 40,000 amp. Although this type of heating cycle apparently allowed little time for "homogenization," close approximations to equilibrium conditions were usually obtained as a result of the very strongly exothermic reactions between zirconium hydride and boron at elevated temperatures. Reproducibility of results in repeated runs under varying conditions was taken as a reasonable indication that equilibrium conditions had been effectively achieved. Quenching was accomplished by rapidly immersing the entire die, containing the reaction products, in water. Specimens could also be cooled at controlled rates by adjustment of the rate of flow of water through the water-cooled copper electrodes in contact with the graphite dies.

Jan 1, 1954

By R. F. Domagala, M. Hansen, D. J. McPherson

On the basis of metallographic analysis, incipient melting data, thermal analysis work, and X-ray diffraction, phase relationships in the 0 to 50 atomic pct Cr region were carefully resolved. Phase relationships in the 50 to 100 atomic pct region were outlined. A single intermediate phase, ZrCr,, was established in the system. A eutectic and a eutectoid were found in the zirconium-rich region of the system, while a eutectic was located in the chromium-rich region. Limited solubility of chromium in both a and a zirconium was observed. THIS is another in a series of papers on zirconium-base binary phase diagrams based on a program sponsored by the Atomic Energy Commission. The Zr-Sn,' Zr-Si,' and Zr-Mo and Zr-W" systems have already been published. Previous work on the Zr-Cr system includes a partial diagram by Hayes, Roberson, and Davies' and a complete diagram by McQuillan." The phase relationships reported by these authors were determined with magnesium-reduced zirconium, whereas those reported herein are based on iodide metal. Some X-ray work on the intermediate phase in this system was also reported by Hayes et al., as well as by Wallbaum.' Comparisons with these data are made in the appropriate sections. Thermal analysis, metallographic and incipient melting techniques were employed to accurately resolve phase relationships in the 0 to 50 atomic pct Cr region, while the remainder of the diagram is outlined on a basis of cast structures and thermal analysis. A description and analyses of the metals used in the preparation of alloys for this study are included in Table I. A "low-hafnium" zirconium crystal bar (Westinghouse "Grade 3") was used. This material was produced by the decomposition of a volatile iodide onto a hot zirconium filament. The as-received crystal bar was coated with corrosion product from a standard autoclave test by which its grade designation is determined. A sand-blasting and subsequent HF-HNO, pickling treatment was employed to remove this contamination. The bars were cold-rolled to approximately 1/32-in. sheet, pickled again, sheared to 1/4-in. squares, washed with acetone, and stored for use. Ingot chromium metal from National Research Corp. and electrolytic chromium flakes from Johnson, Matthey and Co., Ltd., were available as alloy additions. The ingot chromium was received in granular form. The large granules were sized by crushing, treated for iron removal, rinsed with acetone, and stored. The flakes were broken to approximately 1/8 in. with mortar and pestle, cleaned with acetone, and stored for charging. Experimental Procedure A nonconsumable electrode arc furnace identical to that described by Hansen, Kessler, and McPher-son,' was used in this work. A 400 amp dc welding generator was the source of power. To insure homogeneity, ingots were melted in a water-cooled spun copper crucible, a total of four times without opening the furnace. No difficulties were encountered in the preparation of these binary alloys. Analyzed chromium contents were found to agree very well with the intended content. Melts of pure zirconium were interspersed in the course of preparing the alloys and their hardness was measured to keep a constant check on the melting technique. Vycor and quartz bulbs were used to contain the samples for isothermal annealing treatments. Vycor bulbs were employed for temperatures up to 110O°C, while quartz was used above this temperature. For treatments up to 950°C the bulbs were evacuated before sealing, but for higher temperatures a partial pressure of argon was admitted to the

Jan 1, 1954

By R. F. Domagala, M. Hansen, D. J. McPherson

Iodide zirconium was combined with calculated amounts of nitrided zirconium sponge and arc melted to prepare alloys in the 0 to 6 wt pct N region. Annealing treatments were carried out at 21 temperature levels. Metallographic examination of the heat-treated specimens permitted construction of the binary phase diagram from 0 to 6 pct N. Features of the diagram include the peritectic formation of both a and ß solid solutions. The maximum solubility of nitrogen is 0.8 pct in 8 zirconium and 4.8 pct in a zirconium. An X-ray study of nitrided materials was made in the range 6 to 13 wt pct N region because serious nitrogen losses were experienced when attempts were made to arc melt these high nitrogen alloys. PHASE relationships in the Zr-N system have been determined. Due to the inability to retain nitrogen in nitrogen-rich alloys subjected to arc melting, definitive work was possible only in the range 0 to 6 wt pct (0 to 30 atomic pct) N. Sufficient complementary X-ray work was done to permit construction of the binary phase diagram up to 13.3 wt pct (50 atomic pct) N (ZrN). Small ingots were prepared by arc melting master alloys with enough zirconium to produce an alloy of the desired composition. Following pretreatment, alloys were annealed at temperature levels between 600" and 2020°C. Determination of the phase boundaries was then accomplished by metallographic evaluation of specimens quenched from the various temperatures. Incipient melting techniques were used to corroborate solidus curves, and X-ray diffraction was employed to study the nitrogen-rich region of the a + ZrN phase field. Materials Westinghouse Grade 1 iodide zirconium crystal bar served as the base material for this investigation. The as-received bars were lightly sandblasted, pickled in a 20 pct HN0-5 pct HF aqueous solution, rinsed in water and acetone, and dried. They were rolled to about 1/32 in. strip and acid-pickled, followed by water and then acetone rinses. The material was sheared to approximately 1/4 in. squares, cleaned with acetone, and stored for use. Pure zirconium nitride is not commercially available, according to correspondence with possible suppliers. A literature survey indicated that nitrogen may be introduced into zirconium metal by passing nitrogen or ammonia gas over zirconium at an elevated temperature. After considerable experimentation, a train was devised whereby high quality nitrogen was passed through a series of bubblers to remove the last traces of oxygen, through an H,SO, bubbler and a cold trap to remove moisture, and finally over zirconium within a resistance furnace. Hand-picked Bureau of Mines magnesium-reduced zirconium sponge proved more amenable to nitriding than the crystal bar. The nitrogen used was extra pure gas purchased from Linde Air Products Co. The first two bubblers through which the gas was passed contained a solution suggested by L. F. Fieser" as being ideal for removing the last traces of oxygen from nitrogen. The solution contains 20 g KOH, 2 g sodium-anthraquinone ,9-sulphonate, and 15 g NaHSO, dissolved in 100 ml water. It is blood-red and turns brown when contaminated with oxygen. A saturated lead acetate solution, next in the series of bubblers, removed any H2S which might have formed in the 0, removal stage. An H,SO, bubbler and cold trap removed moisture before the nitrogen was passed over the zirconium. The zirconium was contained in a stainless steel screen within a Vycor or quartz tube in a resistance furnace. Unreacted nitrogen passed out of the system through a mercury bubbler. A nitriding run was performed in the following manner: Zirconium was placed in the screen within the furnace tube. The unit was assembled, suitably clamped off, and evacuated to remove the air and moisture within the tube and that trapped by the zirconium sponge. The unit was flushed with nitrogen and evacuated twice. After nitrogen was allowed to pass over the zirconium for about 30 min, the furnace was turned on; several hours were required to reach temperature. The reaction was allowed to continue at temperature for about 4 hr and the screen was then withdrawn from the hot zone of the furnace by means of a Nichrome or Kanthal rod. The screen and its contents were allowed to cool at the end of the tube and were then removed. At 800°C no more than about 5 pct N could be introduced into the sponge zirconium. This material was designated M-1 (master alloy No. 1) and reserved for the preparation of alloys in the dilute nitrogen region. A second set of experiments was conducted at 1000°C; no more than 7.2 pct N could be introduced into the zirconium at this temperature. Enough material (M-2) of this nitrogen content was prepared to produce a second set of alloys to complement the first and to provide alloys in the 0 to 6 wt pct N range. It was hoped that a diffusion anneal plus a re-nitriding might yield a material of higher nitrogen content. The material containing about 7 pct N was therefore given a homogenization treatment at 1000°C for 6 hr under an argon atmosphere. It was

Jan 1, 1957

By R. F. Domagala, D. J. McPherson

Iodide zirconium was combined with calculated amounts of ZrO2 or master alloys and arc-melted. Annealing treatments were carried out at 21 temperature levels. Metallographic examination of the heat treated specimens permitted construction of the binary phase diagram from zirconium to ZrO2. Oxygen additions to zirconium raise the transformation temperature as well as the melting point. Features of the diagram include the peritectic formation of beta, the formation of alpha directly from the melt, an intermediate phase ZrOB with a' range of homogeneity, and a eutectic between alpha and Zr02. PHASE relationships in the Zr-0 system were determined in the range 0 to 66. 7 atomic pct 0 (ZrO,). Arc-melted alloys were annealed at temperature levels between 600" and 2000°C. Determination of the phase boundaries was accomplished by metallographic evaluation of specimens quenched from the various temperatures. Incipient melting techniques were used to determine solidus curves, and X-ray diffraction work was employed to study the lattice parameters of the a solid solution and the phase ZrO,. Experimental Procedures Westinghouse "Grade 1" iodide zirconium crystal bar (approximately 99.8 pct pure) was employed for the investigation. This material is substantially free of the impurity inclusions characteristic of most zirconium metal. Previous work has shown it to be the most satisfactory grade for phase diagram studies. The zirconium crystal bar, as-received, was coated with corrosion product from autoclave tests by which its grade designation is determined. The bars were sand blasted lightly, pickled for 1 min in a 20 pct HN0,-5 pct HF aqueous solution, rinsed in water and acetone, and dried. The bars were rolled to about 1/32 in. strip, cut into 10 in. lengths, and pickled again. The material was sheared to approximately Y4 in. squares, cleaned with acetone, and stored for use. Two pounds of specially prepared "chemically pure," and 100 grams of spectrographic grade ZrO, were obtained for the preparation of alloys. The "Specpure" ZrO, was obtained from Johnson, Mat-they Co.. Ltd.. and Titanium Alloy Manufacturing Div. of National Lead Co. supplied the other grade of oxide. The "chemically pure" material (henceforth referred to as C.P. ZrO,) was used for the preparation of all alloys employed in the determination of phase relationships. A limited number of alloys were prepared with the "Specpure" oxide to recheck certain phase boundaries. Analyses of the two grades of oxide are given in Table I. The C.P. ZrOI was received in powder form, not suitable for arc melting. Therefore, this material was compressed to wafers 1 in. in diameter and about Vn in. thick. These were segmented and stored for use. The "Specpure" material, in the form of small granules, was suitable for use in the as-received condition. A nonconsumable electrode arc-melting furnace was used to prepare alloys. A water-cooled copper crucible and tungsten-tipped electrode were employed. The material was placed in the crucible and rapidly melted under a protective atmosphere of helium gas by striking the arc on a tungsten stud and directing it upon the charge material. Details of operation, including drawings of this type furnace, have been published.'

Jan 1, 1955

By C. E. Lundin, M. Hansen, D. J. McPherson

PRIOR work on the Zr-Cu phase diagram by Alli-bone and Sykes,' Pogodin, Shumova, and KUGU cheva,' and Raub and EngeL3 as confined largely to copper-rich alloys. The investigations of Raub and Engel were the most recent and seemingly the most complete of these. Alloys from 0 to 68.3 pct Zr were studied principally by thermal analysis and microscopic examination. These authors reported an inter metallic compound ZrCu, (1116°C melting point) and two eutectics, one at 86.3 pct Cu (977°C mp) and the other at 49 pct Cu (877°C mp). The solubility of zirconium in copper was reported to be less than 0.1 pct at 940°C. The zirconium melting stock consisted of Westing-house "Grade 3" iodide crystal bar (nominally 99.8 pct pure). It was treated by sand blasting and pickling (HF-HNO, solution) to remove the surface film of corrosion product, resulting from grade designation tests. The crystal bar was cold rolled to strip, lightly pickled again, and cut into pieces approximately 1/32 in. thick and 1/4 in. square. These were cleaned in acetone, dried, and stored for charging. The high-purity copper (spectrographic grade) was supplied by the American Smelting and Refining Co. with a nominal purity of 99.99 pct. These copper rods were rolled to strip, cut into squares the same size as the zirconium platelets, cleaned in acetone, dried, and stored. Equipment and Procedures The equipment used for melting and annealing the zirconium binary alloys and for the determination of solidus curves has been described in connection with previous work on the Ti-Si system' and in recent papers in this series describing the studies on eight binary zirconium systems.5-' Techniques employed for preparing and processing the alloys were also similar to those used in the above references. Ingots of 20 g were melted under a protective atmosphere of helium on water-cooled copper blocks in a nonconsumable electrode (tungsten) arc furnace. The ingots were homogenized and cold-worked prior to isothermal annealing to aid in the attainment of equilibrium. The specimens were heat-treated in Vycor bulbs sealed in vacuo or under argon, depending on the temperature of the anneal. Quenching was accomplished by breaking the Vycor bulbs under cold water. Temperature control was within ±3OC of reported temperatures. Thermal analysis was primarily relied on to determine eutectic levels, peritectic levels, and compound melting points. The induction furnace incipient melting technique was also used but did not provide the accuracy obtained by thermal analysis in this system, which involves much lower solidus temperatures than the other zirconium systems. A special technique for the determination of characteristic temperatures was employed in the case of several intermediate phases and their eutectics which displayed very small differences in melting temperatures. Specimens were sealed in Vycor bulbs and annealed at a series of very accurately controlled temperatures. Metallographic examination was then employed to reveal incipient melting. Furnaces and techniques in general were described previously.' The echant used was 20 pct HF plus 20 pct HNO3 in glycerine unless otherwise stated. Results and Discussion The chemical analyses of the majority of alloys prepared for the determination of phase relationships in this system are given in Table I and a brief summary of the equilibrium anneals employed is given in Table 11. In a preliminary program, alloys containing 1, 4, and 7 pct Cu were annealed for three different times at each of the temperatures 700°, 800°, and 900°C. No change in the relative amounts of phases present was detected after 350, 150, and 75 hr at the above temperatures, respectively. The times listed in Table II were accordingly chosen as a result of these preliminary tests. Zirconium-rich alloys containing from 0.1 to 10 pct CU were reduced by cold pressing from 58 to 8 pct, depending upon thk alloy content, homogenized for 7 hr at 900°C, and then reduced 80 to 13 pct by cold rolling, again depending upon copper content. Other alloys were studied in the cast, or cast and annealed conditions. The contracted scope of investigation for this system included the range 0 to 50 atomic pct Cu. This approximate region is shown in Fig. 1. Due to evidence of phase relationships departing considerably from those proposed by Raub and Engel" in the 50 to 100 atomic pct range, the investigation was extended to cover this composition area rather thoroughly also. Fig. 2 is a drawing of the entire diagram. The labeling of some phase fields was omitted in Fig. 2 for the sake of clarity. An expanded view of the zirconium-rich region, with the experimental points necessary for its construction, is given in Fig. 3. The generally accepted value of Vogel and Tonn8 or the allotropic transformation a + 862' ±5OC, was employed in the construction of these diagrams. A careful study revealed that the "Grade 3" crystal bar used in this investigation actually transforms over the approximate range 850" to 870°C, due to impurities. It must be expected that this two-phase field in unalloyed zirconium will cause some departures from binary ideality in the very dilute alloys. Zirconium-rich Alloys: The a + ß transformation temperature is decreased from 862" to about 822°C by increasing amounts of copper. Thus, a eutectoid reaction, fi ß a+ Zr,Cu, occurs at a composition of about 1.6 pct Cu. The eutectoid level was determined to lie between the alloy series annealed at 815" and 830°C. The placement of this eutectoid temperature

Jan 1, 1954

By E. L. Kamen, H. D. Kessler, M. Hansen, D. J. McPherson

The highly reactive Ti-Mo and Ti-Cb alloys were prepared and heat treated under protective conditions. Phase diagrams were established based on micrographic and X-ray diffraction analysis and detection of incipient melting. ß titanium forms a continuous series of solid solutions with both molybdenum and columbium, whereas these metals are only slightly soluble in the a titanium modification. PREVIOUS work on the Ti-Mo and Ti-Cb systems was very limited and is discussed in succeeding sections of this paper relating to the respective phase diagrams. The diagrams in the present work were determined principally by micrographic analysis, because other methods such as thermal analysis and dilatometry are not suitable, due to the low diffusion rates of these alloys in the transformation range. X-ray analysis was believed less useful than micrographic analysis because: 1—On quenching certain high titanium alloys from above their transformation range, the body-centered cubic ß phase decomposes into acicular a (hexagonal close-packed) which room temperature X-ray diffraction techniques cannot differentiate from isothermal a. 2— High temperature X-ray techniques might be used, but are difficult to apply and are subject to error because of the high reactivity of titanium with gases such as oxygen and nitrogen. Although metallo-graphic study was the basic method used, X-ray diffraction analysis was applied to a limited extent, to further substantiate the findings. The transformation ranges on the titanium side of the Ti-Cb and Ti-Mo systems were studied both with DuPont Process A (magnesium-reduced) titanium alloys and with iodide titanium alloys. The Process A metal was used chiefly to bracket the transformation range in order to reduce the number of alloys given a particular treatment for the more accurate final study with iodide titanium alloys. Experimental Procedure The analyses of the metals used in the preparation of alloys for this study are presented in Table I. Melting Practice: Early in the work, a consumable electrode, arc melting furnace was used in the preparation of the Process A Ti-base Mo and Cb alloys. This type of furnace has been described by other investigators.', ' One of the major difficulties of the consumable electrode melting process was the prep- aration of homogeneous electrodes (with the mixture of coarse titanium sponge and fine powder alloying addition) which would give homogeneous ingots. Composition gradients encountered may have been due to a tendency for powdered alloying metals to drop away from the electrode in handling. This difficulty in the preparation of homogeneous alloys was overcome by the use of a nonconsumable electrode furnace for later work. This furnace is similar to the one used by Mc-Pherson and Fontana a Ohio State University. A diagram of the nonconsumable electrode arc furnace is given elsewhere.' The source of power was a 400 amp dc welding generator. Electrode tips of tungsten (% in. in diam) were used for all columbium melts and for a limited number of molybdenum melts. Molybdenum tips were used for the majority of molybdenum alloy melts. Columbium tipped electrodes were tried, but disintegrated rapidly at currents above 300 amp. In the Process A titanium melting practice the metals were charged into the water-cooled spun copper crucible in the following forms: l—titanium: sponge (—4 + 16 mesh); 2—molybdenum: sections of 0.003 in. sheet (approximately ½ in. sq); and 3— columbium: powder compacts (? in. lumps broken from large section). After alternately evacuating and purging the furnace several times with argon or helium, the arc was struck against the metal charge,

Jan 1, 1952

By R. F. Domagala, M. Hansen, D. J. McPherson

On the basis of metallographic analysis, incipient melting data, thermal analysis work, and X-ray diffraction, phase relationships in the 0 to 50 atomic pct alloy content were carefully resolved. Phase relationships in the higher alloy content regions were outlined. A single intermediate phase, peritectically formed, of the form ZrX2 was established in each system. A eutectic and a eutectoid were found in the zirconium-rich region of both systems. Appreciable solubility of molybdenum and wolfram was found in ß zirconium, but only a slight solubility in a zirconium. PREVIOUS work on these systems was limited to X-ray investigations of the intermediate phases1, 2 and preliminary surveys of the zirconium-rich alloys." The diagrams presented were determined principally by metallographic, thermal analysis, and incipient melting techniques. In accordance with the Atomic Energy Commission contract requirements, phase relationships up to 50 atomic pct molybdenum or wolfram were carefully resolved, while only a limited amount of work was done to outline the 50 to 100 atomic pct alloy region. A description and analysis of the metals used in the preparation of alloys for this study are included in Table I. A "low-hafnium" zirconium crystal bar, produced by the decomposition of a volatile iodide onto a hot filament, was employed for these studies. Westinghouse "Grade 3" crystal bar was used for all alloys. The zirconium, as received, was coated with corrosion product from a standard autoclave test by which its grade designation is determined. This was removed by a sand-blasting and HF-HNO, pickling technique developed by the Atomic Energy Commission. The bars were cold rolled to approximately 1/32 in. sheet, pickled for iron removal, sheared to 1/4 in. squares, washed with acetone, and stored for furnace charge. Molybdenum sheet. 0.003 in. thick, 0.001 to 0.005 in. wolfram wire, and M in. wolfram rod crushed to granular form, from Fansteel Metallurgical Corp., were used as alloying stock. Melting Procedure A nonconsumable electrode arc furnace identical to that described in detail by Hansen et al.4 was used. The source of power for melting Zr-Mo alloys was a 400 amp dc welding generator, while a 900 amp dc generator was used for preparing Zr-W alloys. Electrode tips of the same material as the alloy addition were employed in both systems, i.e., molybdenum tips were used for Zr-Mo alloys, and wolfram tips for Zr-W alloys. This precluded any possible electrode contamination during arc melting. Ingots of 20 to 30 g of Zr-Mo alloys were melted under high purity helium in the cavity of a copper block insert in a spun copper crucible. Homogeneity was insured by remelting all ingots four to six times without opening the furnace. Due to the large difference between the melting points of wolfram and zirconium, direct melting of the charged metals as above generally yielded ingots containing unmelted wolfram particles. This was overcome, after many variations in technique, by melting a pure wolfram charge at high power levels (700 to 900 amp) and slowly adding zirconium to form a master alloy. Such master alloys were sectioned and remelted several times, then crushed to serve as melting stock for alloys of lower wolfram content. Master alloys containing 26, 50, 52, 70, and

Jan 1, 1954

By R. W. Kraft

A back- reflection precession-type X-yay camera for determining the crystallographic orientation of the crystallites of both phases in small areas of thick specimens of polyphase alloys is described ad the geometry, advantages, and limitations of the apparatus are discussed. Metallographic obseg-vations of interfacial angles (or habit planes or growth direction are made on the same specimen so that the crystallographic and metallographzc orientation data call be directly correlated. An example of the application of tile technique to a unidi-~~ectionally solidified CuA12-A1 eutectic specimen is presented. The interfacial planes ad growth directions of each phase were established, and the orientation relationship between the phases was observed and found to be approximately intermediate between two previously reported relationships. The crystallographic interfacial relationship between two solid phases is an important parameter in a variety of metallurgical phenomena because of the energy associated with the interface and because this energy is at least partly associated with the way in which the two space lattices are in contact along the interface. In order to describe the interfacial planes in the crystallites of both phases it is necessary to determine the crystallographic orientation of each phase in a specimen relative to each other, and to directly correlate these data with met-allographic measurements from which the interfacial planes can be determined on the same specimen. The back-reflection Laue technique is the easiest way to determine the orientation of thick single crystals or large grains and the method can be combined with metallographic techniques to provide the necessary data, provided the reflections from each phase can be distinguished from one another. However, if the crystallites are small or of varying orientation or if one or both phases has a unit cell with less than the maximum symmetry, many overlapping and complex Laue patterns are recorded simultaneously and it becomes impractical if not impossible to interpret the photographs. All of these complexities were present in a unidirectionally so- lidified A1-CuA12 eutectic alloy1 for which a crystallographic analysis was desired. Consequently the method described here was developed since no known X-ray or electron diffraction technique had all of the following attributes which were required. 1) Method should yield data from which the crys-tallographic orientation of every crystallite or diffracting unit in the irradiated area can be determined. 2) Method should be adaptable to the study of small areas. Depending upon the degree of preferred orientation in the specimens, it should be possible to obtain reliable data in a reasonable length of time on irradiated areas as small as 1.0 or even 0.1 mm in diam. 3) Method should permit direct correlation of crystallographic data with microscopic orientation data pertaining to crystallite axes, habit, morphology, and growth directions as determined by optical microscopy. 4) method should be such that selected areas of large specimens can be easily studied. Trepanning of a small specimen (such as an electron microscope specimen) from larger specimens was undesirable since it would greatly increase the problem of correlating the crystallographic and microscopic directional data. Requirements 3 and 4 dictated that an X-ray back-reflection pinhole technique should be used on large samples, such as metallographic specimens which had previously been examined and analyzed for their microscopic characteristics. Similarly, requirement 2 could be satisfied by choosing a collimator of appropriately small diameter. Exposure times would not be excessive provided the detector were close enough to the specimen. Because of space limitations film was chosen in preference to a Geiger counter. Requirement 1 was fulfilled by using monochromatic radiation and providing enough additional degrees of freedom in specimen rotation to compensate for Bragg law restrictions on diffracted beams. DESCRIPTION OF APPARATUS The apparatus shown in Fig. 1 consists of a pin-hole collimator, (a), projecting through a rotatable circular film in a holder, (b), which records X-ray reflections in the back-reflection region. The specimen, (c), is mounted on the end of a shaft, (d), which is provided with an adjustment, (e), so that the specimen surface can be placed accurately on

Jan 1, 1962

By C. A. Siebert, S. H. Bush

Isothermal temper-embrittlement studies were conducted on a 5140 steel at various temperatures for times as long as 3000 hr. Specimens from the embrittled steel were subjected to impact tests, metallographic examination, microhardness, and lattice-parameter measurements. SEVERAL of the steels in the low alloy group are subject to a pronounced decrease in impact energy if held within, or cooled slowly through, a range of temperatures below the lower critical. This phenomenon, known as temper embrittlement, is quite pronounced in the chromium steels of which AISI 5140 is typical. This study considered the response of one heat of 5140 steel when embrittled isothermally at various temperatures for times as long as 3000 hr. Excellent reviews of the work done in the field of temper embrittlement have been written by Hollomonl and Woodfine.' Both authors list the various theories which have been advanced to explain temper embrittlement. Some of the mechanisms considered were: 1—a transformation below 700°C; 2—an allotropic modification of iron; 3—precipitation from the ferrite of such compounds as carbides, nitrides, phosphides, or chromium oxides; 4—decomposition of retained austenite; 5—modification of special carbides; 6—grain-boundary segregation; and 7—selective distribution of carbides. The present study investigated the changes in hardness, Meyer's number, impact properties, micro-structure, and lattice parameter which occurred on embrittlement. Previous studies considered the influence of microstructure and of prior austenitic grain size3 on the response of 5140 steel to temper embrittlement. Equipment and Procedure All samples were taken from one heat of 5140 steel; the chemical composition of this steel was: 0.43 pct C, 0.86 pct Mn, 0.020 pct P, 0.019 pct S, 0.27 pct Si, 0.84 pct Cr, 0.10 pct Ni, 0.06 pct Mo, 0.10 pct Cu, and the remainder Fe. The material was received as % in. rounds. This stock was cut down to blanks measuring approximately 0.425~0.425~2.196 in. Heat Treatment: The rough blanks were normalized at 1600°F (870°C) for 1 hr, then austenitized at 1600°F (870°C) for 1 1/2 hr, and quenched into an agitated oil bath. The specimens were tempered at 1275°F (690°C) for 5 hr and water quenched. These tempered specimens were designated as the toughened state. Embrittling Treatments: The tempered bars were divided into sets and subjected to the various embrittling treatments listed in Table I. Temperatures from 750" to 1050°F were investigated for times of 1 to 3000 hr. Special tests to determine the effect of heating a short time at a higher temperature after long-time embrittlement at some lower temperature were also conducted. All specimens were water quenched at the end of the embrittling cycle, then ground and notched to obtain standard Charpy V-notch specimens. Impact Tests: The bars were fractured in a 260 ft-lb Charpy impact unit; specimens were fractured at temperatures ranging from —315°F to room temperature. Fixed temperatures used consisted of —315°F with liquid nitrogen, —115°F with dry ice and acetone, and 32°F with ice and water. Intermediate temperatures were obtained by suspending the bars in a jig over the liquid nitrogen, dry ice and acetone, or ice and water.

Jan 1, 1955

By Charles L. Guettel

The temperature coefficients of electrical resistance of 31 alloys in the nickel-rich corner of the Ni-Cr-Fe system were determined. The results indicate that a range of binary Ni-Cr alloys has lower temperature coefficients than a specific region of Ni-Cr-Fe alloys which has been reported and quoted by others as providing the lowest coefficients in the ternary system. DESIRABLE properties of resistivity and temperature coefficients of resistance have been the basis for the successful commercial application of various Ni-Cr and Ni-Cr-Fe alloys. In particular, low or negligible temperature coefficients of resistance are necessary for wire used in precision wire wound resistors and potentiometers involved in closely controlled electrical and electronic circuitry. Alloys most successfully applied to this purpose basically have not been ternary Ni-Cr-Fe alloys but rather heat-treated Ni-Cr alloys in the region of approximately 73 to 79 pct Ni and 20 pct Cr with significant additions of other elements such as aluminum, silicon, iron, copper, manganese, and cobalt. While the addition elements play a vital role in the nature and control of the temperature-resistance curve of any one of these alloys, the inherent relatively low temperature coefficient of the binary Ni-Cr base composition itself in the approximate region of 80 pct Ni and 20 pct Cr contributes fundamentally to the eventual negligible temperature coefficient obtained with the applied alloys. W. A. Dean 1 has reported mean temperature coefficients (20' to 100oC) of as-cast samples for the Ni-Cr-Fe system, the nickel corner of which is given in Fig. 1. Dean had to use as-cast samples because alloys from a considerable region of the system were unworkable. Dean refers to the work of Hunter and sebast2 who used annealed wire samples to show that, for chromium contents of 10 to 25 pct, the temperature coefficient increases as the iron content is increased at the expense of nickel. Although the data of Hunter and Sebast were limited in the nickel-rich areas to eight samples of 60 to 91 pct Ni with only four samples containing iron, the possibility that no Ni-Cr-Fe alloy has a lower temperature coefficient than certain binary Ni-Cr alloys should have been given serious consideration. However, Dean's work has been referred to by J. S. Marsh,3 who, quoting Dean, states that for the temperature coefficients "the smallest values-about 0.0001 per oC (100 ppm)—were found for alloys containing about 55 pct Ni and 20 pct Cr" (balance iron). A region such as that reported by Dean would be significant to those trying to develop alloys for the applications described. Therefore, a series of three-pound laboratory open-air induction melts of Ni-Cr and Ni-Cr-Fe alloys were melted, conventionally processed to 0.005 in. diam wire, and continuously strand annealed at 1900° F. The basic Ni-Cr compositions, with balances essentially iron other than usual minor impurities, and the corresponding temperature coefficients of resistance and resistivities are given in Table I. The resistivities are approximate and are not to be considered as a basis for relating them to their position of composition in the ternary system. The wire diameters were determined with a micrometer, and not by weight and density measurements which would be required for accurate resistivity values. However,

Jan 1, 1963

By W. G. Pfann

Under certain conditions, a molten zone can be made to move through a solid by impressing a stationary temperature gradient across the solid. This phenomenon can be utilized in fabricating semiconductive devices, growing single crystals, joining, boring fine holes in solids, measuring diffusivities in liquids, small scale alloying, and purification. Fundamentals and exemplary applications are outlined. ANEW aspect of the travel of a molten zone through a solid is considered in this paper. Whereas, in previously described zone melting techniques,'-' zones were usually caused to move by changing the position or temperature of a heat source, in the present technique a zone is made to move by impressing a temperature gradient across it. For example, a thin layer of molten A1-Si alloy, sandwiched between two silicon slabs having a temperature gradient normal to the layer, will travel through the hotter slab. One feature of temperature gradient zone melting, as this technique will be called here, is that zones of unusually small dimensions can be maintained. This leads to such diverse uses as have been listed in the abstract. In this paper fundamentals of the movement of a molten zone in a temperature gradient, and some practical applications thereof, are outlined. Assumptions The movement of a molten zone through a solid body of solvent substance in which a temperature gradient exists will be discussed. The length of the zone will be its dimension in the direction of motion. The direction of the temperature gradient will be toward the region of higher temperature. To be considered is a binary or higher order solute-solvent system in which the concentration of solute in the molten zone is sufficient to produce a lowering of the freezing temperature of the solvent which is at least of the order of the temperature range impressed on the charge. While we discuss primarily molten zones in a solid matrix, the principles are applicable to solid or vaporous zones in a solid matrix, a requirement being that the diffusivity in the zone be substantially greater than in the matrix. Diffusion of solute into the surrounding solid will be assumed negligible. In general, its presence will not basically alter the conclusions to be drawn. For illustration, a system represented by the partial constitutional diagram of Fig. la will be discussed, in which k, the distribution coefficient defined in the figure, is constant and less than unity. However, the technique is applicable to any solute- solvent system in which one component lowers the melting point of another component. Movement of a Molten Zone in a Temperature Gradient Consider, in the system AB of Fig. 1, where B is designated the solute and A the solvent, the effect of sandwiching a thin layer of solid B between blocks of solid A and placing the whole in a temperature gradient such that the temperature of the layer is above the lowest melting temperature of the system. Surrounded by an excess of A, the layer will dissolve A, becoming molten, and expand in length." As so- lution of A continues, at both interfaces, the mean solute concentration in the zone will move to the right in Fig. la, until at temperature T, the liquid at the cooler interface reaches the liquidus concentration C,. Solution of A thereupon ceases because the liquid T1 is saturated with A. The liquid at the hotter end of the zone, being at temperature T2, is not saturated with A at concentration C1. Hence solution of A continues there and concentration C2 is approached. A concentration gradient therefore is set up in the zone, causing A to diffuse toward the cool end and B toward the hot end. As a result, the liquid at the cool interface becomes supersaturated and a layer of crystalline A containing concentration kc, of B in solid solution freezes. Since a source of A

Jan 1, 1956

By R. Maddin, A. Lawley

Single crystal Mo-Re alloys (99.99+ purity), grown by electron bombardment floating zone heating, were deformed in tension at temperatures from -196° to +200°C. At -196oC, Mo-6 pet Re and Mo-20 pct Re alloys deform by slip on systems of the type (101) [1ll]. Above-78oC, the operative slip system is of the type (112) [111], (101) [1ll], or (211) [1ll], depending upon orientation. Mo-39 pct Re alloys deform by twinning and slip at all test temperatures with a (112) [111] type twinning system. Critical resolved shear stress for slip and twinning is calculated as a function of composition and temperature, and the mechanical properties compared with those of less Pure polycrystalline alloys. HEXAGONAL close-packed rhenium, c/a = 1.615, is highly soluble in molybdenum. The bee molybdenum-rich solid solution extends to approximately 42 at. pet Re* at 2500°C,' Fig. 1. Recent work on polycrystalline alloys has shown conclusively that the addition of rhenium to molybdenum improves both the fabricability and the low-temperature ductility.2"5 Similar behavior is observed for W-Re and Cr-Re alloys.5 Initially, Geach and Hughes2 found that the Mo-30 pct Re alloy could be cold rolled to approximately 90 pct reduction before the onset of cracking. The remarkable effect of the rhenium was later confirmed by Jaffee et al. 3-5 with an optimum condition of maximum strength combined with maximum ductility at the Mo-35 pet composition. At higher rhenium contents, the ductility was severely impaired due apparently to the presence of brittle tetragonal s phase, Fig. 1. In the Mo-35 pet Re alloy, at least 110 ppm (parts per million) oxygen could be tolerated before cold workability was reduced. By comparison, the corresponding tolerance figure in unalloyed molybdenum is ~ 10 ppm. The effect of rhenium was considered in terms of: 1) a reduction in oxygen solubility in the lattice; 2) a redistribution of intergranular oxide from grain boundary films to isolated spheres; and 3) the promotion of twinning as a further mode of plastic deformation. Paradoxically, twinning is often associated with brittle fracture in iron and the silicon irons.6'7 Work to date has been restricted to are-melted and powder metallurgy polycrystalline material of ~ 99.9 pet purity. The present investigation was initiated for the purpose of determining the mode and crystallography of plastic deformation of a series of bee Mo-Re alloy single crystals over a wide temperature range from -196o to +200°C. Using high-purity zone-melted material (~ 99.99 + pet purity), the specific role of rhenium in the molybdenum lattice may be understood more fully since effects due to grain boundaries and oxide redistribution will be absent. At the same time, the study provides a comparison of single and polycrystalline deformation behavior. EXPERIMENTAL TECHNIQUES A) Preparation and Purity of Single Crystals. Alloys in the form of either are-cast and swaged or sintered and ground bar, 0.120 in. diam, were obtained from the Chase Brass and Copper Co. Alloy single crystals ~ 8-in. long were grown from the

Jan 1, 1962

By R. P. Carreker

ONE of a series, this report is concerned with the experimental documentation of the deformation behavior of pure metals over a wide range of temperature. Previous reports in the series describe the creep behavior of platinum1 and aluminum2 and the tensile behavior of copper,&apos; silver,& and molybdenum.5 Subsequent reports will describe the creep behavior of copper and silver and the tensile behavior of platinum, thus providing extensive creep and tensile data on four face-centered-cubic metals tested under comparable conditions. Material Two lots of high-purity aluminum were used in this investigation. One lot, designated AC, was obtained direct from the Aluminum Co. of America as high-purity aluminum in notch-bar form. A typical analysis of this type of aluminum is 0.006 pct Si, 0.015 pct Cu, 0.006 pct Fe, and 99.975 pct Al. Lot BS was obtained from J. I. Hoffman of the Bureau of Standards as a portion of the Aluminum A described in his report of the redetermination of the atomic weight of aluminum.6 Lot BS originated with the Aluminum Co. of America also, and was received in the standard notch-bar form. A detailed chemical and spectroscopic analysis of lot BS has been published.6 Principa1 impurities were 0.006 pct Si, 0.003 pct Fe, 0.002 pct Cu, and >99.987 pct Al. Samples were processed identically by room-temperature swaging and drawing to 0.030 in. diam wires from approximately ½ in. diam rods that were machined from the notch-bar pigs. Samples were cut from the cold-drawn wire, placed in a grooved graphite block in groups of thirteen, and annealed in air for 1 hr at each of several temperatures, with the results shown in Table I. The two lots of aluminum responded quite differently to annealing, the behavior of aluminum AC suggesting that its grain growth was controlled by impurities. The recrystallization texture of aluminum of similar purity comparably treated has been reported as predominantly <111> fiber with some <100> fiber present.7 Testing Procedure The testing procedure and method of analyzing the data have been described previously.2,3 Annealed specimens of 5 in. gage length were tested in an Instron tensile testing machine at a constant head motion corresponding to a strain rate of 0.04 min-&apos; at a number of temperatures. Instantaneous tenfold rate changes were employed in some tests to deter-

Jan 1, 1958

By J. W. Spretnak, G. W. Powell, J. H. Bucher

Tlze room-temperature tensile fracture oj smooth, round specitnens of three ultrnhigh- strength steels tempered to a wide range of strength levels was studied by means by light and electron-microscopic examination of the fracture surfaces. The fracture of AISI 4340 and 300 M at all the strength levels studied, and H-11, except after tempering at 1200° and 1300°F, occurs in three stages. The initiation of fracture is internal (except in some lightly tcmpeved specimers in which fracture is initiated at surface flaws), and is nucleated largely by separation at metal-second phase intevjaces. TIze voids grow and, coalesce to form a crack. When the crack has reached a sufficienl size, rapid propngutio~z ensues. Failure in this stage of fracture usually occurs by dimpled rupture of inicroshear stefis. In the case of H-11 tempered in the 1125° to 1300°F range, fracture in the shear steps is predominantly by concentrated deformation without void formation. The termination of fracture is usually occomplished by the formation of a shear lib in which fracture occurs by shear dimpled rupture. In the case of H-11 tempered at 1200° and 1300°F, no shear lip was obserued, and the radial elelments extend to the surface—a true termination slage does not exist. ThE tensile fracture of several metals and alloys has been investigated.2-4 In the case of polycrystal-line materials, cup-cone fracture usually results. The mechanism of cup-cone fracture may be summarized as follows.5 Cavities are formed in the necked region of the specimen. They usually are initiated by inclusions or second-phase particles. The cavities extend outwards by means of internal necking, and a crack lying about perpendicular to the length of the specimen is formed in the necked region. Subsequent crack growth occurs by the spread of bands of concentrated plastic deformation inclined at an angle of 30 to 40 deg to the tensile axis. Cavities are formed in the bands of concentrated deformation. The deformation bands zigzag across the bar with the net result that mac-roscopically the crack extends about perpendicular to the specimen axis. The final separation, or cone formation, appears to occur by continued crack propagation along one of the deformation bands out to the surface of the specimen. The micromechanics of the tensile fracture of ultrahigh-strength steels have not been thoroughly investigated. Larson and carr6,7 studied the tensile-fracture surfaces of AISI 4340 with a low-power microscope and reported that three stages of fracture could be observed in general. A centrally located region characterized by circumferential ridges, an annular region characterized by radial surface striations, and a peripheral shear lip were found. It was first pointed out by 1rwin8 that the central region is very probably one of fracture initiation and slow growth, and that the annular, radially striated region is one of rapid crack growth. Presumably the crack grows slowly, assuming roughly a lenticular shape, until it is large enough for the initiation of rapid propagation. In this investigation, it was attempted to determine the fine-scale aspects of the room-temperature tensile fracture of some ultrahigh-strength steels, and to relate the variation in fracture mode with microstructure. The steels studied were AISI 4340, 300M, and H-11 tempered to a wide range of strength levels. I) EXPERIMENTAL PROCEDURE The compositions of the steels studied are given in Table I. The steel was received in the form of hot-rolled bar stock 5/8 to 1 in. in diameter from which oversized specimens were machined and heat-treated. The heat treatments employed are given in Table 11. Subsequent to heat treatment, the specimens were ground to the final dimensions and stress-relieved by heating for 1 hr at 350°F (with the exception of the as-quenched steel). Standard smooth round specimens of 0.252-in. diameter and 1-in. gage length were tested in a Tinius Olsen Universal Testing Machine using a cross-head speed of 0.025 in. per min. The relatively coarse aspects of the fracture topography were determined by light-microscopic examination of sections through the fracture surface of nickel-plated specimens. A direct carbon-replication technique9 was used in the electron-microscopic study of the fracture surfaces. The replicas were examined in the electron microscope, and stereo pairs of electron micrographs were taken. The stereo pairs were then examined using a Wild ST4 Mirror Stereoscope. Carbide and inclusion particles extracted in the replicas were analyzed by selected-area electron diffraction. II) EXPERIMENTAL RESULTS The mechanical testing data are summarized in Table 111. The values reported are the average of

Jan 1, 1965

By W. V. Kotlensky, H. E. Martens

In an earlier communication, Ref. 1, the authors presented the tensile properties and the changes in crystal structure and microstructure for pyrolytic graphite tested at 2750°C. It was suggested that changes in the structure may have an effect on the tensile properties of this material. Accordingly, tensile tests were run on a series of specimens, from one block of pyrolytic graphite, some of which were hot-worked by being strained at 2750°C, others of which were tested in the as-deposited condition. According to the manufacturer, General Electric Co., the polycrystalline material was deposited at 2100°C on a substrate of synthetic graphite by decomposition of a dilute natural gas mixture. The as-deposited material had a density of approximately 2.19 g per cc, and as shown in Fig. 1 had a structure of medium size growth cones which were regenera-tively nucleated. All specimens had a gage section 0.75 in. long by 0.06 in. wide by 0.1 in. thick with the thickness direction perpendicular to the basal planes. The testing method and equipment used have been previously described, Ref. 2 and 3. For both the hot-working and the testing the load was applied parallel to the basal planes; a cross-head speed of 1.5 x l04 in, per sec was used. All heating was done in an atmosphere of bottled helium.

Jan 1, 1962

By R. L. Smith, J. L. Rutherford

ALTHOUGH considerable effort has been devoted toward the determination of the mechanical properties of pure metals, it is extremely difficult to compare the results of such work. This is because of differences in method and type of test, grain size, heat treatment, and purity of material. There is very little information on the properties of most of the high melting point reactive metals with purities greater than 99.9 pct. Previous evidence concerning the tensile properties of unalloyed metals show that: body-centered-cubic metals are brittle at low temperatures; face-centered-cubic metals remain ductile; some hexagonal metals remain ductile, others become brittle. It may be that traces of impurities decrease the low temperature ductility of the body-centered-cubic lattice. This would suppose that a body-centered-cubic lattice free from impurities would exhibit nearly the same ductility as a face-centered-cubic structure. The present investigation has been an evaluation of the tensile properties of zone refined iron in order to determine whether higher purity material than that previously evaluated behaves differently. Experimental Procedure In this investigation, high purity iron was zone purified using the floating zone method.&apos; The starting stock was obtained from two sources. The first,

Jan 1, 1958

By O. W. Simmons, L. W. Eastwood, C. M. Craighead

The results of a preliminary study of 113 ternary titanium-base alloys are described. The compositions investigated were as follows: 1. Ternary titanium-carbon alloys containing copper, silicon, vanadium, chromium, manganese, iron, or cobalt. 2. Ternary titanium-nitrogen alloys containing chromium. 3. Ternary titanium-chromium alloys containing additions of vanadium, molybdenum, tungsten, cobalt, or nickel. 4. Ternary titanium-manganese alloys containing additions of silicon, chromium, tungsten, or iron. Tensile properties, minimum bend radii, hardnesses, response to heat treatment and aging treatment, and phase relationships for these alloys were determined. THIS is the second in a series of papers describing some of the results of the exploratory work on titanium alloys being conducted at Battelle Memorial Institute for the Wright-Patterson Air Force Base. The work with the binary alloy systems has been described in a previous publication.&apos; This paper describes ternary titanium-base alloys as follows: 1. Ternary titanium-carbon alloys containing copper, silicon, vanadium, chromium, manganese, iron, or cobalt. 2. Ternary titanium-nitrogen alloys containing chromium. 3. Ternary titanium-chromium alloys containing additions of vanadium, molybdenum, tungsten, cobalt, or nickel. 4. Ternary titanium-manganese alloys containing additions of silicon, chromium, tungsten, or iron. As the preparation and fabrication of the 0.5-Ib experimental titanium alloy ingots to 14-ga sheet, as well as the methods of testing used in this work, have been described in the first paper of this series, Titanium Binary Alloys, no further discussion of these techniques will be given. Ternary Alloys Containing Carbon The following alloys of titanium, containing carbon and one other metal, were investigated: 1. 0.25 pct carbon; 1, 2.5, and 5 pct copper. 2. 0.25 pct carbon; 1, 2, and 3 pct silicon. 3. 0.25 pct carbon; 1, 2.5, and 5 pct vanadium. 4. 0.25 pct carbon; 2.5, 3.5, 5, 10, and 15 pct chromium. 5. 0.5 pct carbon; 2.5, 3.5, and 5 pct chromium. 6. 0.25 pct carbon; 1, 1.75, 2.5, 3.5, 5, 7.5, and 10 pct manganese. 7. 0.5 pct carbon; 1.75, 2.5, 3.5, 5, 7.5, and 10 pct manganese. 8. 0.25 pct carbon; 0.25, 0.5, 1, and 2 pct iron. 9. 0.25 pct carbon; 1, 2, and 3 pct cobalt.

Jan 1, 1951

By J. W. Downey, S. T. Zegler

A study has been made of the occurrence of Phases having the Cr30-type structure in ternary alloys having the general composition where B cept for iron and copper all the B components are known to form binary compounds having the type structure. The occurrence of the struc-ture type in the ternary alloys was investigated using optical and X-ray metallography. The results indicate that iron can replace virtually none of the cobalt in at 900°C, but substitutes for nickel in up to the approximate composition tually insoluble in V3Ni and 2) and exhibit complete solubility. At 100 nickel can be substituted for cobalt in VSCo up to the composition solubility; and 3) in the system , two solid solutions having the type structure occur. At the same temperature, 1000°C, V3Si exhibits limited solubility for V3Co but is completely mis-cible with V3Ga. Lattice parameters have been determined for all the binary and ternary phases. The solid solutions are discussed in terms of the atomic size and electronic factors which are thought to influence the occurrence of Cr30-type phases. The binary compounds V3Co and V3Ni were found to form peritectoidally from o phases and a vanadium, the former compound at a temperature between 1000" and 1050°C, the latter between 800°and 900°C. The substitutions of nickel and rhodium for cobalt in V3Co were determined not to raise the superconducting transition of VSCo to a temperature as high as 2°K. The occurrence and stability of binary intermediate phases having the Cr30- or A15-type structure (space group o;, Pm3n) have been discussed in several recent papers . Vanadium forms the struc -ture at or near V3B stoichiometry with all the mem -bers of the cobalt and nickel groups in the periodic table and also with Au, Ga, Si, Ge, Sn, P, As, and Sb. This paper describes a survey of the occurrence of the CrsO-type structure in ternary alloys having the general composition V3(B, B&apos;) where B and B&apos; are Fe, Co, Ni, Cu, Rh, Ir, Si, and Ga. The purpose of the survey was to clarify the connection between the position of the B partner in the periodic table and its CrsO-type phase-forming tenden~y.~,~ The survey is part of a broader program concerned with the influence of B-partner substitutions on the superconducting transition temperatures of CrsO-type compounds. The transition temperatures, Tc, of some of the ternary Cr,O-type phases found in the survey are also discussed in this paper. EXPERIMENTAL PROCEDURE The alloys were prepared by arc-melting in an A-He atmosphere and were in the form of buttons, weighing approximately 3 g. The metals used in their preparation were received from commercial sources and were in all cases 99.9+ pct pure. The vanadium was spectrochemically pure except for 800 ppm of Si, 200 ppm of Fe, and 200 ppm of Al. For study of each pseudobinary system at least seven alloys were prepared having the general compositions case of alloys containing gallium, melting losses were small, usually less than 0.3 pct of the total weight of the metals charged for melting. Melting losses were somewhat greater for gallium alloys, ranging in general from 0.9 to 5.0 pct. Ternary V-Si-Ga alloys were chemically analyzed and the results are given in Table I. The analyses show most alloys to be within 2 at. pct of the intended composition. Throughout the balance of the paper the intended compositions are used in discussing the alloys. The cast buttons were annealed within evacuated quartz capsules for times ranging from 3 to 6 weeks at temperatures between 800" and 1000°C and water-quenched. The phases present in the annealed and quenched alloys were identified by optical and X-ray metallography. X-ray diffraction photograms were made with powders prepared from the bulk specimens, by the use of a 114.6-mm-diam powder camera with copper irradiation. The lattice parameter of CrsO-type phases and their standard errors

Jan 1, 1963

By E. T. Stephenson

Wert and Marx1 pointed out that a straight-line relationship exists between the activation energy of a relaxation process and the temperature at which the maximum relaxation occurs. The data available at the time indicated that this relationship held for a number of Snoek peaks, several grain boundary peaks, and pair relaxation in a brass. Theoretical justification for the observed relationship was found in the following equation, which Wert and Marx derived from the theory of interstitial atomic diffusion coefficients:&apos; where AH is activation energy, R is the gas constant, T is the internal friction peak temperature, v is the interstitial atomic frequency, f is the frequency of the applied stress, S is entropy of activation. More accurate internal-friction data on interstitial diffusion in bcc metals have become available since Wert and Marx&apos;s paper. Table I gives activation energies calculated by several authors from a combination of elastic aftereffect data and the peak temperatures at several frequencies of measurement. (In the case of Fe-C and Fe-N, the referenced authors5&apos;7&apos;8&apos;10 also used high-temperature diffusion data.) However, diffusion experiments at 150°C and above1&apos;-13—rather than anelastic techniques—were used to determine the activation energy of hydrogen diffusion in Fe-H. In order to plot this recent data as a AH vs T graph we calculated the peak temperatures of the systems listed in Table I (with the exception of Fe-C, Fe-N, and Fe-H) by using published values3&apos;&apos; of AH and Do. These calculated peak temperatures varied by only a few degrees from experimentally determined peak temperatures. The calculated values were felt to represent the better average. A number of investigators have found the peak temperatures in Fe-C and Fe-N to be 38 to 40°C and 20° to 24°C, respectively. The values of 39" and 22°C (312° and 295°K in Table I) were determined by carefully executed experiments.6&apos;8 The hydrogen peak temperature has been measured at 1 cps,12 20 cps,14 and 35,000 cps.15 All hydrogen peak temperatures were adjusted to 1 cps for Table I using AH = 3000 cal per mole. As Fig. 1 shows, the data of Table I fit a straight line at least as well as the earlier data used by Wert and Marx. Linear-regression analysis was applied to the data in Table I. The Fe-H data was omitted because the activation energy has not been experimentally determined in the temperature range of the peak. The analysis yields the equation: with a standard error of 860 cal per mole. The agreement of the data in Table I with the Wert-Marx equation is excellent. Strict application of Eq. [I] forces the linear regression analysis through the origin. In this case, the analysis gives a slope of 62.1 cal per mole per deg, and a standard error of *890 cal per mole. With this value, it follows from internal-friction theory that Hence, Do varies by only 30 pct among these bcc-interstitial systems. Thus, a concept suggested originally by Wert and Marx1 is confirmed by recently published data. The

Jan 1, 1965

By R. P. Smith

The carbon content of a number of alloys ranging (before carburitation) from purified iron to purified cobalt was determined for fixed activities of carbon by equilibration with each of six CO-CO2 atmosphere at 1000°C. At fixed carbon activity, the carbon content decreases continuously with increasing cobalt content. For concentrations of 0 to 50 at. pct Co the results can be expressed by a linear interaction equation of the type In y2 = 7.76n2/(n1+n3)+ 2.30 n3/(n1+n3) where y2 is the activity coefficient of carbon (standard state such that y2— 1 as n2 — 0 in the Fe-C binary) and n1, n2, and n3 are the number of atoms of iron, carbon, and cobalt per unit weight of alloy. The results are also given in terms of the a functions. In the same concentration range, the results are in accord with the regular solution theory of Alcock and Richardson. THE carbon content of Fe-Si, Fe-Mn, and Fe-Ni alloys at fixed activities of carbon has been reported previously.&apos;-, The present work extends R. P. SMITH, Member AIME, is Research Chemist, E. C. Bain Laboratory for Fundamental Research, United States Steel Corp., Research Center, Monroeville, Pa. these measurements to a series of Fe-Co alloys. These alloys, like the Fe-Ni alloys, are fcc ("aus-tenitic") over the entire range of composition at 1000°C. EXPERIMENTAL PROCEDURE The alloys were made from electrolytic cobalt and Plast-Iron in a small vacuum-melting furnace. A small amount of graphite was added to each melt to keep the oxygen content low. The ingots were analyzed for iron and carbon: the amount of cobalt was obtained by difference. Analysis of the top and bottom of the ingots indicated that they were uniform in alloy content. Except for the pure cobalt the ingots were hot-rolled to approximately 0.15 cm thickness, cleaned by sandblasting, then cold-rolled to about 0.045 cm. We were not successful in cold rolling the pure cobalt; therefore it was hot-rolled to the final thickness. The last traces of surface oxide were removed by acid pickel, then the samples were treated in a hydrogen-2 pct water vapor mixture at 1100°C to remove carbon and most of the oxygen. The carburizing procedure was similar to that previously described.1,2,4 The alloys in the form of strips about 1.7 cm wide, 7.5 cm long, and 0.04 cm thick, together with a similar specimen of purified iron, were equilibrated with each of six CO-CO2 mixtures of fixed composition, then quenched in the

Jan 1, 1965