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By W. L. Phillips

Nickel alloys containing approximately 0.5, 2.0, and 6.0 at. pct of Os, Pd, Ru, and Rh were Prepared by vacuum melting. Tension tests were carried out at 25°, 500°, 800°, and 1000°C; stress-rupture tests were carried out at 650° and 800°C. The room- and elevated-temperature tensile strength, as well as the stress-rupture life, increased for the same atomic percent in the order palladium, rhodium, ruthenium, osmium. The ductility decreased with increasing concentration or temperature for a given concentration. The results are discussed in terms of solid-solution theory. BOTH solid-solution strengthening and weakening have been observed in alloys. The simple rules which have been suggested to explain alloy strengthening are subject to several exceptions and deviations based on atom size and valency effects. The purpose of the present investigation was to extend the knowledge of the effect of solid-solution alloying on the mechanical properties of several Ni-Pt group metal alloys at room and elevated temperatures. The platinum group metals, osmium, ruthenium, rhodium, and palladium were chosen as the solute elements because of their different valencies and atom sizes. Nickel was chosen as the base material because of the large amount of previous literature available on nickel alloys. EXPERIMENTAL PROCEDURE Grade A Carbonyl nickel powder having the following composition: C, 0.07 pct; Fe, 0.013 pct; Mn, 0.0005 pct; Si, 0.016 pct; Mg, 0.0003 pct; S, 0.003 pct, was purchased from the International Nickel Co. The platinum-group metal powders were purchased from A. D. Mackay Co. The powders were weighed, cone-blended for 4 hr, and vacuum-melted into 1-in.-diam billets, approximately 2 in. long. The ingots were homogenized for 5 hr at 1800°F before being extruded into 1/4-in. rods at 2100°F. The alloy compositions are shown in Table I. Typical impurity values were not significantly different than that of the as-received nickel. All alloys were single phase. The lattice parameter of each alloy after the heat treatment described below is also included in Table I. Each alloy was annealed at 2200°F for 1 hr and air-quenched to comparable grain sizes (ASTM G.S. No. 4-5). The rods were then machined into test samples 0.160-in. diam and 0.750-in. gage length. Tension testing was carried out in an In-stron tension-testing machine at a strain rate of 0.02 in. per in. per min. An extensometer was used to determine the room-temperature yield stress. Stress-rupture testing was carried out in an Arc Weld stress-rupture unit. All testing was performed in air. EXPERIMENTAL RESULTS A) Room-Temperature Mechanical Properties. The 0.2 pct yield stress, ultimate tensile strength, and elongation for all alloys tested at room temperature are plotted as a function of atomic percent solute element in Fig. 1. The ultimate tensile strength for the same atomic percent increases in the order palladium, rhodium, ruthenium, osmium. The yield strength at low concentrations, -0.5 at. pct, increases in the order osmium, palladium, ruthenium, rhodium. At higher concentrations, 3.0 and 10.0 pct, the yield strength increases in the same order as the tensile strength. Both yield strength and tensile strength increase rapidly ini-

Jan 1, 1964

By George A. Roberts, Arthur H. Grobe

Tensile, hardness and density properties are presented for a new 18-8 stainless steel powder for the —50, —100, and —140 mesh cuts and also for a prepared blend containing 62 pct —325 mesh powder. The data were obtained for sintering temperatures of 2100° to 2350°F and for compacting pressures of 30 to 50 tons per sq in. A NEED has existed for a stainless steel powder of uniform composition from particle to particle, which could be readily molded and sintered, and which would yield moderate to high strength and ductility when processed into parts by standard powder metallurgy techniques. The literature reveals that previously available powders have been lacking in one or several of these desirable characteristics. Originally attempts to produce stainless steel powder by blending and diffusing elemental metal powders were unsuccessfu1.l A truly alloyed stainless steel powder reported by Dale in 1947 yielded reasonable mechanical properties but required an exceptionally high sintering temperature of 2400°F in pure dry hydrogen.' A later development reported in 1950 by Stern provided a stainless type powder of excellent molding properties' but this powder likewise required sintering at temperatures over 2350 °F and did not yield properties as high as those previously reported. The powder was not truly alloyed prior to sintering. A prealloyed stainless steel powder has been developed that is readily moldable at 30 to 50 tons per sq in., can be sintered at low temperatures (2100" to 2300°F), and which yields compacts with both high strength and ductility. The powder is prepared by the liquid disintegration method," direct from molten metal. Early attempts to make 18 pct Cr, 8 pct Ni stainless steel powder by this method did not produce a usable powder for pressing and sintering because of a preponderance of spherical particles with relatively high, surface oxide content. It was learned that silicon additions to iron alloys, powdered by this process, gave bright powder of low-oxide content and such additions were made to 18-8 stainless with similar results. An attendant change in particle shape occurred reducing the number of spheroids materially. The optimum composition which has been studied and is reported here contains approximately 2.5 pct Si. Silicon additions to stainless steel of the 18-8 type improve the resistance to high-temperature oxidation and promote the formation of ferrite. Corrosion resistance is, however, controlled by chromium content and even though ferrite may exist the corrosion resistance is not generally impaired.4 Experimental Procedure A detailed study of the 18-8, 2.5 pct Si powder has been made in accordance with the following outline: 1. Four blends a. — 50 mesh b. —100 mesh

Jan 1, 1952

By George A. Roberts, Arthur H. Grobe

H. H. Hausner (Sylvania Electric Products Inc., Bayside, N. Y.)—I tested the 18-8 stainless steel powder described by Grobe and Roberts and the results were excellent. The powder was compacted and sintered in purified hydrogen to a very low density (approximately 20 pct and more porosity). Where the other powders show very little ductility when sintered to such a low density, the described stainless steel was surprisingly high in ductility. I would appreciate it if the authors could give us some data on a similar powder but without silicon and also data on 430 stainless steel powder. H. S. Kalish (Sylvania Electric Products lnc., Bay-side, N. Y.)—I should like to verify the excellent results obtained by Grobe and Roberts with their pre-alloyed stainless steel powder. It seems to me, however, that their data are conservative and that even better densities and higher tensile strength can be obtained with these powders by the selection of proper particle size fractions. The accompanying data were obtained by sintering compacts pressed from the powder made by Vanadium-Alloys Steel Co. to 0.8 in. in diam about % in. high in a cylindrical double action die without the use of lubricant. The compacts were sintered for 1 hr at 2370°F in hydrogen purified by passing it through a Deoxo unit (palladium catalyst) and a Lectrodryer. Even the density obtained using the as-received (—100 mesh) compacted at 40 tsi was quite high as was the hardness. This indicates that by increasing the sintering temperature slightly above 2350°F the coining and subsequent secondary sintering operation can be avoided. With the — 325 mesh fraction, however, very much higher density can be obtained and a high hardness. I have no tensile data to present, but I feel that it is not too optimistic to expect. that unusually high tensile strengths and good ductility can be obtained in this sintered material on the basis of the density and hardness data. The sintered stainless steel made by the above-described methods came out of the furnace very bright and clean. The purified hydrogen atmosphere is important in sintering stainless steel to reduce the oxide present in the powder. The result of such a sintering is good densification and good ductility of the sintered product. G. A. Roberts (authors' reply)—We are greatly indebted to Messrs. Hausner and Kalish for their discussions of our paper. Their encouraging remarks are greatly appreciated and their confirmatory data will do much to add to the value of the paper. At the present time it is impossible to provide complete data on a similar stainless steel powder without silicon, but we do know that, in general, the properties are slightly inferior. Data on such a powder and on other stainless steel types are to be published in the near future.

Jan 1, 1952

By M. Schussler, J. S. Brunhouse

Arc-melted and electron-beam melted tantalum in the cold-worked and the recrystallized conditions showed high strength, good tensile ductility, and excellent notch toughness down to 321°F. Arc-melted material, containing moderate amounts of interstitials, possessed better short-time strength than higher purity electron-beam melted material up to 2000°F. Cold-working substantially strengthened arc-melted tantalum, whereas electron-beam melted material exhibited a low work-hardening rate and possessed better formability characteristics. With identical compositions and microstructures, the mechanical properties of melted tantalum and uacuum-sintered tantalum should be similar. TANTALUM metal usage is increasing rapidly in a number of applications. The outstanding chemical corrosion resistance of tantalum accounts for its applications in chemical processing equipment for severe service conditions and for surgical uses in the medical field. The largest present requirements for tantalum are in the electronics industry, where its special properties are useful in capacitors. The high melting point (5425°F) and inherent ductility of tantalum also suggest that tantalum and tantalum-base alloys might offer an excellent combination of high strength at very high temperatures and good fabri-cability. Pure tantalum has been used in significant amounts in parts requiring high-temperature strength, such as heating elements, reflectors, and heat shields in high-vacuum high-temperature furnaces and in electronic tubes. A few very promising tantalum-base alloys have been developed and are being investigated extensively. The possibilities for tantalum as a high-temperature structural material, either unalloyed or alloyed, will become better defined when more extensive mechanical property data are known. pughl measured the tensile properties of vacuum-sintered tantalum in sheet form for the temperature range of 78" to 1500°K (-321' to 2240°F), and Bechtold reported data obtained at low temperatures.' Drennen3 determined creep and stress-rupture properties of both vacuum-sintered and arc-melted tantalum sheet at 1200 °F. Some additional data on the mechanical properties of tantalum at temperatures up to 5000 °F have been summarized elsewhere.4'5 It is probable that large structural shapes of tantalum, either unalloyed or alloyed, can most readily be produced from ingots prepared by melting. Consequently, the present investigation was undertaken to determine more extensive short-time mechanical properties of tantalum metal consolidated by melting. EXPERIMENTAL PROCEDURE Melting and Fabrication—Two tantalum ingots were utilized in this study to obtain information on the effect of interstitial elements on mechanical properties. Both ingots were melted from Union Carbide Metals Co. tantalum dendrites.

Jan 1, 1961

By R. H. Richman, II Conard G. P.

AS part of a study of deformation twinning in bcc crystals, solid solutions of beryllium in iron have been found to twin profusely when strained slowly at room temperature. This note reports some crystal-lographic aspects of the operative mechanical twin modes in a quenched Fe-5 wt pct Be alloy. After melting twice under a partial pressure of helium, the alloy was homogenized, cut into pieces 7 to 10 mm on an edge, solution treated at 1130°C,1 and quenched. Debye-Scherrer patterns confirmed

Jan 1, 1963

By T. A. Read, H. K. Birnbaum

STRESS-induced twin boundary motion in the AuCd ß'phase (52.5 at. pct Au 47.5 at. pct Cd having an orthorhombic structure (space group D h)' was discussed for the case of transformation twins.' Mechanical twinning, i.e., introduction of new twin orientations by the applied stress was sometimes observed in the ß' phase which contained two transformation twin orientations. (See Ref. 2 for experimental details.) The transformation twins terminated at the mechanical twin boundary; the two initial orientations having been twinned to the new orientation as shown in Fig. l. In a "hard" tensile instrument, formation of each mechanical twin was accompanied by a sharp drop in load (and an audible click) resulting in a typical serrated load-extension curve. The mechanical twins increased in width and additional twins were nucleated as the load increased until the entire specimen was occupied by the mechanical twin orientation. A restoring force2 caused the mechanical twin and transformation twin boundaries to return towards their initial configuration on releasing the load. In a stabilized specimen* the total strain introduced by mechanical twinning was recovered on releasing the applied stress. Initial and final microstructures were identical. After decreasing gradually to a critical width, mechanical twins suddenly disappeared accompanied by audible clicks and sharp increases in load (measured at a constant rate of crosshead motion in a "hard" machine). Mechanical twins could be stabilized in any position by maintaining the applied load for a sufficient time, indicating a decrease in restoring force with time similar to that observed for transformation twins.' Orientation relationships for mechanical and transformation twins were determined and are shown in Fig. 2. The mechanical twin, T3, was twin related to transformation twin, T,, by reflection in (1ll), and twin related to transformation twin, T2, by reflection in (111)T, . T, was related to T, by reflection in(ll1)T1 . The "average" mechanical twin composition plane lay several degrees from (1ll), . In every specimen which mechanically twinned, orientation T3 was most highly favored by the applied stress, i.e., it was the twin orientation which provided the greatest change in length compatible with the direction of applied stress

Jan 1, 1961

By J. P. Nielsen

When two grains in a polycrystalline specimen meet at a point in the course of grain-boundary movements, and the new boundary created at the point is one of relatively low specific free energy, a nonequilibrium boundary condition occurs. The nonequilibrium is enhanced if the other grain boundaries involved at the point of meeting are relatively high in specific free energy. The nonequilibrium results in a correction by growth of the complex grain (two subgrains) to a large size, sufficient in many cases for it to become a self-propagating unit, i.e., a recrystallization grain. This mechanism, herein called "geometrical coalescence," is proposed as a logical origin for recrystallization nuclei, particularly for secondary recrystallization. RECENT publications1-' indicate that considerable progress has been made toward a complete understanding of the three microstructural processes in metals: recovery, recrystallization, and grain growth. However, a major problem persists, namely, the origin of recrystallization nuclei. It is generally accepted that once a recrystallization nucleus, a secondary recrystallization nucleus particularly, has reached a certain size, it will continue to grow at the expense of its neighbors, simply as a consequence of grain-boundary free-energy considerations. What has remained obscure, however, is the mechanism whereby such a nucleus comes into being and grows to the self-propagating size. Recrystallization Nucleation in a Two-Dimensional Grain-Boundary System In the left portion of Fig. 1 is shown schematically the meeting of grains 1 and 4 on disappearance of boundary 2-3, i.e., the boundary that existed between grains 2 and 3, as might happen in the course of normal grain-boundary migration. If, by chance —and the probability is not necessarily small— grains 1 and 4 are close together in their mutual orientations, or in some other special way mutually oriented so that the new common boundary has relatively a very low specific free energy (hereinafter referred to by s), then the grain boundary arrangement as depicted in the figure is highly unstable. This is particularly so if boundaries 1-3, 3-4, 4-2, and 2-1 are high in their s values. Assuming, for the sake of argument, that boundary 1-4 has a s equal to zero, and the s's of all the other boundaries present are equal to each other and constant, then the boundary arrangement on the right in Fig. 1 will be

Jan 1, 1955

By J. D. Mote, A. Rosen, J. E. Dorn

The effect of strain rate and temperature on the critical resolved shear stress for (1100) [1120] prismatic slip was determined for the intermediate hexagonal phase containing about 67 at. pct Ag and 33 at. pct Al. Whereas the flow stress increased only slightly as the test temperature was first decreased below room temperature, a rapid increase inflow stress was obtained with yet greater decreases in temperature from about 240" to 4°K. The effect of both temperature and strain rate on the flow stress over the low-temperature range could be completely rationalized in terms of the Peierls mechanism when the deformation is controlled by the rate of nucleation of pairs of kinks. A few years ago Mote, Tanaka, and Dorn1 observed that the critical resolved shear stress for prismatic slip of a 67 at. pct Ag plus 33 at. pct A1 hexagonal intermediate phase by the (1100)[1120] mode decreased precipitously as the test temperature was increased from 4" to about 160°K. Their results for both basal and prismatic slip are given in Fig. 1. It is obvious that only the flow stress for prismatic slip is strongly temperature-dependent, whereas the flow stress for basal slip is independent of temperature. If the temperature dependence of the flow stress were primarily due to interstitial impurities, it might be expected that they would influence the flow stress for basal slip as well. Since the activation volume at 115°K for prismatic slip was determined to be about 15b3, where b is the Burgers' vector, this thermally activated deformation process was tentatively ascribed as being due to the Peierls mechanism. But the various analyses of the Peierls mechanism then available as formulated by Seeger2 and also by Lothe and Hirth3 did not agree well with the experimental data. Since that time a more realistic theory for the operation of the Peierls mechanism under conditions where the strain rate is determined by the rate of nucleation of pairs of kinks has been presented by Dorn and Rajnak.4 The present investigation was, therefore, undertaken in order to provide the essential data for a careful check on the agreement between this theory and the low-temperature deformation by prismatid slip of the above-mentioned Ag2A1 intermediate phase. EXPERIMENTAL TECHNIQUE Single-crystal specimens so oriented that the [1120] direction and the normal to the (1100) plane were at 45 * 2 deg to the tensile axis were grown from the congruent melting composition of about Ag2Al for the intermediate hexagonal phase by the

Jan 1, 1964

By G. Derge, Ling Yang, G. M. Pound

The specific conductance and its temperature dependence were measured over the entire composition range of the molten Cu2S-CuCI system. At a typical temperature of 1200°C, 10 rnol pet of the ionically conducting CuCl reduced the specific conductance from about 77 ohm-lcm-l for pure Cu2S to about 32 ohm -1cm -1, and 50 mol pet CuCl reduced the conductance to that for pure CuCI—about 5 ohm 1cm1. The nature of electrical conduction in molten Cu2S, FeS, CuCI, and mixtures was studied by measuring the current efficiency of electrolysis at about 1100°C. The Cu2S, FeS, and mattes were found to conduct exclusively by electrons, but addition of 1 5 wt pet CUS to Cu2S produces a small amount of electrolysis. Addition of CuCl to Cu2S suppresses electronic conduction, and ionic conduction reaches almost 100 pet at a CuCl concentration of about 50 mol pet. These facts are interpreted in terms of electron energy level diagrams by analogy to the situation in solids. RESULTS of electrical conductivity studies on molten Cu-FeS mattes as a function of composition and temperature have been reported.' The specific conductances ranged from about 100 ohm-' cm-' for pure Cu2S to 1500 ohm-' cm-1 for pure FeS. This is in sharp contrast with the low specific conductance of molten ionic salts for which the transfer of electricity is by migration of ions in the field. In general, these ionically conducting molten salts, such as NaC1, KC1, CuC1, etc., have a specific conductance of the order of magnitude of 5 ohm-' cm-'. It was concluded on the basis of this evidence that molten FeS and Cu,S exhibit electronic conduction. Pure molten FeS has a small negative temperature coefficient of specific conductance, resembling metallic conduction, while pure molten Cu2S has a small positive temperature coefficient, resembling semi-conduction. The molten Cu2S-FeS mattes follow a roughly additive rule of mixtures, both with respect to specific conductance and temperature coefficient. Savelsberg2 has studied the electrolysis of molten Cu2S and Cu2S + FeS. He concluded that while molten Cu2S is an electronic conductor, there is some ionic conduction in molten Cu2S + FeS3 owing to the formation of the molecular compound 2Cu2S.FeS and its dissociation into Cu1 and FeS2-1 ions. The present work does not verify his results. Chipman, Inouye, and Tomlinson" have studied the specific conductance of molten FeO and report a high specific conductance, about 200 ohm-1 cm-1 of the same order of magnitude as that found for molten mattes, and a positive temperature coefficient. They interpret these results in terms of p-type semiconduction in the ionic liquid by analogy to the situation in solid FeO.1 imnad and Derne' detected appreciable ionization in molten FeO by means of electrolytic cell efficiency measurements. In order to verify the conclusion that electrical conduction in molten Cu2S and mattes is electronic, and to gain further insight into the structure of molten sulfides, the following investigations were carried out in the present work: 1) The specific conductance, s of the molten system Cu2S-CuC1 was measured as a function of temperature over the entire composition range. As discussed later, molten CuCl is an ionic substance. It was thought that if molten Cu2S were simply ionic in nature, addition of small amounts of CuCl might not have a catastrophic effect in lowering the high conductance of the Cu2S. On the other hand, if much electronic conduction occurs, addition of the ionic CuCl should have a large effect in destroying the electronic conduction. 2) The electrolytic cell efficiency of the following molten systems was measured at about 1100°C in specially designed cells: Cu3; Cu2S + FeS, 50:50 by wt; FeS; Cu2S + CuS, 15 wt pet; Cu2S + CuC1, 5.9 to 46.4 mol pet; and CuC1. This gives a direct measure of the fraction of current carried by ions in these melts. Further, the cell efficiency, extrapolated to zero ionic current, is given by cell efficiency = (s leasile + s elexstronic). [1] s lucile for molten CulS would be expected to be no greater than that for molten CuC1, whose s lonle is about 5 ohm-' cm-1, as will be seen in the following. u,.,,.,.,.......for molten Cu,S is of the order of 100 ohm-' cm-'.' Thus, a large increase in cell efficiency from 0 to values of 10 to 100 pet upon addition of CuCl to Cu2S would indicate destruction of the electronic conductance. Conductance Measurements Experimental Procedure—The apparatus and experimental method were the same as those described in detail in connection with the study of electrical conduction in molten Cu,S-FeS mattes.' A four terminal conductivity cell and an ac poten-

Jan 1, 1957

By W. A. Wood, W. H. Reimann, Maria Ronay

The basic mechanism of fatigue is studied in annealed a brass subjectecl to alternating torsion at room temperature, 100°, 200°, 300°, and 400°C, and in air. It is shown that the slip-zone micro-cracking which characterizes fatigue damage produced by small amplitudes at room temperature is progressively replaced by grain boundary cracking at elevated temperatures, replacement being complete at 400°C. Replacement occurs not because slip activity decreases hut because slip movements at elevated temperatures cease to concentrate in narrow zones and instead disperse. Decrease of amplitude at 400°C through permitting longer lives of specimens before fracture actually causes in- MOST studies of fatigue at elevated temperature have been concerned with data for design. Observations on the basic mechanism of fatigue have been largely incidental, though they have led to the interesting inference that this mechanism has features in common with that of creep.' This paper is concerned primarily with the mechanism itself. Advantage has been taken of mater- crease in pain boundary damage, an anomaly attributed to the greater difficulty at elevated tempwature of starting and propagating a crack. Surface corrosion was most pronounced along slip bands but since at elevated temperatures the slip hands showed no cracking the surface corrosion did not appear to influence onset and early spread of fatigue damage. Similarly the type of fatigue damage produced at room temperature by cyclic strain at large amplitudes, characterized by irregular mi-crocracks at zones between disoriented regions in a grain, was also replaced by mainly grain boundary cracking at elevated temperature, though not so completely. ials and procedures known to show readily how fatigue affects the microstructure of a metal at room temperature, so that direct comparison might be made with what the same procedures show when the material is subjected to fatigue deformation at higher temperatures. 1) MATERIAL This was 70/30 a brass, used extensively in previous studies at room temperature.' Specimens were turned from extruded rod to have parallel test portions 1-1/4 in. long and 7/32 in. diameter, annealed to a grain size -1/10 mm, and electro-

Jan 1, 1965

By L. F. Mondolfo, F. A. Crossley

The mechanism of grain refinement by the addition of small amounts of titanium, molybdenum, zirconium, tungsten, and chromium to aluminum was investigated. The results indicate that the grain refinement is caused by the peritectic reaction which, by transforming the intermetallic compound into crystals of aluminum solid solution, seeds the melt with nuclei above the freezing point of aluminum. THE feasibility of reducing the grain size of aluminum alloy castings by the addition of small amounts of titanium, zirconium, molybdenum, tungsten, columbium, boron, and chromium has been known for a long time, and the use of some of these elements as grain refiners is common practice in the aluminum foundry industry. The purpose of the present investigation was to determine the mechanism by which these additions produce the grain refinement. Asato et al., on the basis of extensive work on copper, antimony, and silver alloys, presented a theory of grain refinement based on the occurrence of the peritectic reaction. Briefly, this theory states that during cooling of the melt crystals of the primary phase form, which react peritectically with the liquid upon further cooling. The peritectic reaction transforms at least partially the primary crystals into crystals of the secondary phase, which then act as nuclei for solidification of the remaining melt. As the peritectic reaction takes place the primary crys- tals break in pieces, and each piece acts as a nucleus. Fine dendritic crystals which can easily break are the most effective for grain refinement. Sicha and Boehm," in investigating the refining effect of titanium on an AI-Cu alloy, concluded that TiA1, crystals serve as nuclei to initiate crystallization, but did not offer further information on the mechanism of refinement. Bonsack in the discussion of their paper suggested that refinement was due to clouds of some solid in the melt, probably TiO² which hindered grain growth. Eborall,' after studying the effect of titanium, zirconium, vanadium, tungsten, columbium, chromium, and boron on some alloys, concluded that the peritectic reaction was not an essential feature of grain refinement. Her conclusion was based mainly upon the observation that under her experimental conditions (sand casting) the Al-Cr alloys, which undergo a peritectic reaction, did not show appreciable refinement. Cibula" investigated the mechanism of grain refinement and reported that two types of grain refinement exist: 1—Grain refinement produced by restriction of grain growth by the concentration gradient mechanism proposed by Northcott.' 2— Grain refinement produced by the presence of particles in the melt, upon which the melt crystallizes easily. He further concluded that very marked grain refinement results from a combination of the two mechanisms. Since Tic was found present in aluminum alloys refined by the addition of titanium, it was identified as the nucleating agent. In the case

Jan 1, 1952

By L. F. Mondolfo, F. A. Crossley

SURFACE effects in the brittle fracture of materials such as glass and in the plastic slip of zinc and cadmium crystals are well known.' Recently, another surface effect has been found for zinc monocrystals. It has been observed that certain normally ductile zinc monocrystals become ex- tremely brittle when their surface is coated with a layer of oxide or copper plate.

Jan 1, 1952

By Nicholas J. Grant, H. C. Chang

Microscopic observations during creep tests were made on AI-20 pet Zn, 80 pet Ni-20 pet Cr, and 25 and 3S aluminum specimens. All these materials failed in an inter-crystalline manner under certain stress and temperature conditions. Particular atten-tion was given to the initiation and propagation of intercrystalline cracks in coarsegrained AI-20 pct Zn specimens. The geometry of the grains and grain boundaries in these specimens was known before the tests. On the basis of the observations and a knowledge of the interaction between grains and grain boundaries, a mechanism of intercrystalline fracture and propagation of the fracture has been proposed. METAL, and alloy fracture occurs essentially in two ways: transcrystalline and intemrystalline. In the past most of the work done in this field was centered on the determination of stress criteria for fracture. The specimens used for these studies were tested under varying conditions of stress state, temperature, and strain rate. Unfortunately, very little attention was given to the initiation and propagation of the fracture, whereas the nature of fracture was examined closely only after failure. Several extensive reviews of these studies have been published by Orowan, Smith, Barrett, and Petch.1, 4 It is well known that a material which exhibits a ductile and transcrystalline fracture can be made to exhibit a brittle and intercrystalline fracture by modifying its composition and by heat treatment. This modification may or may not result in microscopically observable precipitation in the grain boundaries." In the creep of any one alloy the transition of transcrystalline to intercrystalline fracture is a function of stress, strain rate, and temperature." The tendency for intercrystalline fracture to occur increases as the strain rate decreases and the testing temperature increases. Extensive work on high purity aluminum (99.995 pet) under creep conditions showed the importance of the interaction between grains and grain boundaries, and the different responses of grain and grain boundaries to the variables: stress, strain rate, and temperature.7, 8 It was also shown that grain boundary deformation necessarily has to be accommodated in the grains by various deformation means as. for example, fold formation. The type of fracture resulting from the creep of high purity aluminum was invariably found to be transcrystalline. On the other hand, 2S and 3s aluminum were, under certain conditions, found to fracture in an intercrystalline manner." One reason given for this was that the grains could accommodate to a lesser extent the deformation of the grain boundaries.19 It was decided to examine this line of reasoning in greater detail and to investigate the particular interactions between grains and grain boundaries that result in the incidence and progression of inter-ct,ystalline fracture. It is also the purpose of this paper to show where and how intercrystalline cracks start and how they propagate to result in fracture. Several alloys—Al-20 pct Zn, 80 pet Ni-20 pet Cr, 2S and 3s aluminum-— were used for this study, but attention was paid mainly to the behavior of an A1-20 pet Zn alloy, since its general creep behavior has been established." From the results of this study a mechanism of intercrystalline fracture is proposed. Experimental Procedure The experimental procedure was similar to that used in previous work and has been detailed elsewhere.' Two parallel flat surfaces were milled from an originally round specimen, giving a gage portion which was 1x0.09X0.17 in. Some round specimens, 1/4 in. round by 1 in., were also tested. The A1-20 pet Zn specimens. which were annealed at 1030°F for 24 hr, showed two to four grains across the width of the test bar, and most of the grains occupied the whole thickness of the specimenu. Before some of the specimens were subjected to creep, the grain arrangement was mapped out in a development drawing, Fig. 1. The initiation and

Jan 1, 1957

By J. J. Gilman

The dependence of ortho kink-band formation on crystal orientation, on temperature, and on the conditions at the ends of a specimen is described. Load-compression curves for crystals that kink are presented. The reversibility of ortho kinking is demonstrated, and motion of kink planes through crystals is shown. Finally, a phenomenological explanation of ortho kink-band formation is given. ORTHO compressive kink bands (Fig. la) were discovered for the case of metals by Orowan,1 and they have been discussed by Jillson2 and in some detail by Hess and Barrett." However, their history in nonmetallic crystals is considerably older. It is not known when they were first observed in mineral crystals, but descriptions of them may be found as early as 1885.* They were sometimes mistaken for twins.6 Mügge seems to have been first to realize their true nature and importance. He reviewed the knowledge of them for a wide variety of crystals: anhydrite, antimonite, kyanite, vivianite, gypsum, mica, graphite, molybdenite, barium bromide, cal-cite, sodium nitrate, and barite. Also, the "twins" observed by Bridgman7 in sapphire crystals resemble kink bands more than twins. The terms ortho and para kink bands have been coined in order to distinguish between two types of kink bands which seem to represent the same geometrical configurations, but which have different origins and sizes. Ortho kink bands are the usual type that form at low loads when zinc crystals are compressed parallel to the basal plane. Para kink bands form at much higher loads in highly deformed crystals. These terms refer to compression phenomena and are not to be confused with analogous tension phenomena such as tensile kink bands and deformation bands. Mügge described the geometry of ortho kink bands quite completely for kyanite crystals, Fig. lc. The geometry observed for cadmium by Orowan1 and for zinc by Hess and Barrett3 is very similar to that described by Mügge. (See Hess and Barrett for a complete discussion.) Mügge associated kinking with bending combined with basal slip. He pointed out that the kink angles

Jan 1, 1955

By F. D. Rosi, F. C. Perkins, L. L. Seigle

An investigation was made of the mechanism of plastic flow in coarse grained specimens of both sponge and iodide titanium at low (-196°C) and high (500° and 800°C) temperatures. Deformation by slip occurs predominantly on a {1011} plane and in a <1120> direction over this entire temperature range, and secondary slip on {101-1} planes becomes more prevalent with overthisentireincreasing temperature.range, The crystallographic habit of the predominant twin planes becomes more prevalent with type is temperature-dependent, and deformation by twinning increases as the temperature of testing is lowered. Kink bands were not observed at -196OC and rarely at the high temperatures. A LTHOUGH the deform a tional mechanisms in . -titanium at atmospheric temperature have been extensively investigated," there is little information regarding the influence of testing temperature on these mechanisms. The only available data come from the recent work by McHargue and Hammond, who extended single crystals of iodide titanium at a temperature of 815°C. No significant changes in the types of slip and twinning planes were observed, but increased activity on the type I, order 1 pyramidal slip planes and on the type 11, order 2 twinning planes was reported. These planes contribute very little to the deformation process at room temperature.&apos;&apos; &apos; Since there is interest in titanium as a material for use at both low and high temperatures, the present investigation was undertaken to examine the slip and twinning behavior in coarse grained specimens of titanium at -—196", 500°, and 800°C in order to determine the general effect of temperature on the mechanism of plastic deformation. Experimental Material Used and Specimen Preparation—Both sponge and iodide titanium were used in the present investigation; chemical analyses appear in Table I. These materials were arc-melted under argon into 25 g slugs, which were then cold-rolled approximately 90 pct to a sheet thickness of 0.05 in. Since the strain-anneal technique was used for obtaining coarse grains, the cold-rolled sheets were made into tensile specimens which were then annealed to produce a relatively coarse uniform grain structure prior to critical straining. The details of the strain-anneal treatment. size of grains, and the method for preparing the specimen surface for optical microscopy have been described in a previous report.&apos; In order to obtain a wider spread in crystal orientation for the present work, it was necessary to vary the rolling schedule to reduce the degree of preferred orientation in the annealed sheets prior to critical straining. Methods for Tensile Testing—The apparatus for experiments at the temperature of liquid nitrogen (—196°C) consisted of two concentric stainless steel

Jan 1, 1957

By F. D. Rosi, C. A. Dube, B. H. Alexander

The slip and twinning planes have been determined in deformed crystalsof titanium by an X-ray method of analysis. The slip planes are of the type {1010} and {1011}, while the twinning planes are of the type {1072}, {1121}, and (1122). In the case of the predominant (1010) slip, a type I digonal axis of indices <1120> was the effective slip direction. Microscopic observations on the appearance of the three twin types disclosed a distinct difference in their shapes, and an attempt was made to correlate this difference with the amount of twinning shear. The slip mechanism was discussed on the basis of the difference in c/a ratios for the common hexagonal close-packed metals. IT is a general rule that slip in metals occurs most easily on the planes of greatest atomic density and of the largest interplanar spacing, and that the direction of slip is a close-packed direction. Thus, of the hexagonal close-packed metals which have been investigated by the German school in the decade 1925 to 1935 (cadmium, zinc, and magnesium), it has been found that slip occurs on the basal plane (0001) and in a <1120> direction.&apos; In the case of magnesium above 225oC, slip occurs on a first-order pyramidal plane (1011) as well as on the basal plane.&apos; The hexagonal metals, moreover, deform in part by twinning, and, again with the exception of magnesium at elevated temperatures, the twinning plane has always been reported to be a second-order pyramidal plane {1012).&apos; In view of this consistency in the plastic behavior of these metals, it has been supposed that the same slip and twinning elements would be operative in the other important hexagonal close-packed metals (titanium, zirconium, and beryllium), although there are some reasons why this may not be so. Referring to the tabulation in Table I of the axial ratios of the important metals in this structural class, it may be seen that they can be categorized into three distinct groups depending on the value of their c/a ratios. In the first group are cadmium and zinc with large positive deviations from the ideal ratio for closest packing. This expansion in the direction of the c-axis would be such as to accentuate the basal plane as the slip plane, because it increases the interplanar spacing. The second group includes cobalt and magnesium, which closely approximate the ideal ratio. In the third group are zirconium, titanium, and beryllium whose axial ratios are decidedly less than the ideal. The lattice of the elements in this group is compressed along the c-axis which, in effect, tends to make the basal plane less favorable for slip inasmuch as it reduces the interplanar spacing and the atomic density of this plane. Thus, for metals in the third group, it may be said that there is a certain indefiniteness about the glide plane, in that planes other than the basal plane may be expected to operate. The fact that magnesium begins to exhibit planes other than the normal slip and twinning planes at elevated temperatures may indicate that this behavior will become more pronounced in metals of lower axial ratios. For this reason it might be expected that the slip and twinning behavior of zirconium, titanium, and beryllium would be different from that found in cadmium and zinc. One indication that there is a difference is reflected in the recent work which has been done on the deformation textures of these metals."&apos; As will be shown later, there is a significant difference between the rolling textures of cadmium, zinc, and magnesium on the one hand, and zirconium, titanium, and beryllium on the other, and this difference suggests that the mechanism of plastic flow is not the same in these two groups of metals. From these considerations, therefore, it may be seen that there are good reasons why it cannot be concluded that titanium will have the same crystal-lographic elements of slip and twinning as those exhibited by zinc and cadmium. For the complete determination of these elements it would be most advantageous to have single crystals. However, much valuable information can be obtained by using coarse grained polycrystalline specimens, provided the grains are sufficiently large to permit X-ray

Jan 1, 1954

By F. D. Rosi

The slip and twinning behavior in extended titanium crystals were studied in some detail. The formation and appearance of coarse kink bands are discussed. Their crystallographic geometry was determined by X-ray analysis. A phenomenological interpretation of the complexities in kink band development is also presented. The critical resolved shear stress, coefficient of shear hardening, and plane of fracture were determined for several crystals extended at room temperature. THE slip and twinning elements observed in the room-temperature deformation of titanium were enumerated in a previous paper1 in which considerations were advanced regarding the nature and selection of these elements and their effect on the known mechanical properties of this metal. The present study concerns the crystallographic, microscopic, and mechanical aspects of flow in relation to the slip and twinning elements, and includes a prediction of slip systems, nature of slip and twin markings, inhomogeneities of plastic flow, and stress-strain characteristics. Unless otherwise noted, arc-melted titanium sponge (99.77 pct) was used in these experiments. The method of production of crystals, their dimensions, and the surface preparation for micrographic examination have been reported.&apos; For obtaining stress-strain characteristics, only those crystals which traversed the entire width of the specimen and were at least 8 mm in length were used. Tensile deformation of the crystals was performed with conventional grips for sheet specimens and a constant-stress loading beam, designed after the method of Andrade and Chalmers.&apos; The specimens were loaded by allowing sand to flow from a reservoir into a bucket suspended from the longer end of a balanced 6:1 lever arm at a rate controlled to load the specimen approximately 2 kg per min. Strain measurements were made using the Baldwin SR-4, bonded, resistance-wire strain gage, Type A-8, which permitted a reading accuracy of 2 microinches per inch. The formulas used in the evaluation of shear stress and shear strain, as in deriving the coefficient of shear hardening, are given by Schmid and Boasv n terms of the original orientation and change in length of the crystal. For calculating the critical resolved shear stress, the standard equation was used. The crystallographic nature of the unpredictable slip observed in a number of specimens was determined by the single-surface X-ray method of analysis as described in ref. 1. Experimental Results Prediction of Slip System: It was reported&apos; that room temperature slip in titanium takes place predominantly on {10i0} and in a <1120> direction, giving three potential slip systems. Hence, it should be possible to predict the operative primary system in the manner used by Taylor and Elam&apos; for alu- minum (i.e., the one with the greatest component of shear stress in the direction of slip). The stereo-graphic construction in Fig. 1 shows that this is true. In all cases, slip was found to occur initially on the (0110) plane and, in a number of cases, in the [21f0] direction. (The direction of slip was not determined for all orientations.) It follows that duplex slip can be expected when two slip systems are geometrically equally favorable for slip, as demonstrated in Fig. 2. In both crystals slip took place simultaneously on the (1700) and (olio) prismatic planes, as indicated by Fig. 2a and b. In the case of crystal B the movement of the specimen axis with increasing extension was toward the (10i0) direction, which represents the stable end orientation resulting from alternate slip on the [2ii0] and [1120] directions. This is analogous to the duplex slip process in face-centered cubic metals." Unpredictable Slip: In a number of crystals whose orientations were well within the operation of a single slip system, secondary slip occurred on planes not predicted by the criterion of maximum shear stress. Examples are shown in Fig. 3. In addition to

Jan 1, 1955

By J. B. Newkirk

IN 1939 Bitter and Kaufmann1 suggested that iron, precipitating from a copper-rich, Cu-Fe solid solution, appears initially as coherent particles of r-Fe which transform to the body-centered-cubic form of iron when the alloy is cold worked. Since then there has been considerable discussion, supported only by indirect evidence, regarding the validity of the concept. For example, Reekie, Hutchison, and Hether-ington2 interpreted the changes in electrical resistance accompanying the aging of Cu-Fe specimens in terms of a coherent 7-Fe precipitate which transformed to the incoherent a-phase during long aging or cold working. Greninger- ound that Cu-Fe alloys which had been water quenched from 1050" and then held at 650°C for 3 hr changed from weakly to strongly ferromagnetic as the aged wire was subsequently cold worked. He also found diffraction lines corresponding to body-centered-cubic iron in Debye-Scherrer patterns given by aged and worked specimens but lines corresponding to only one face-centered-cubic phase in specimens which had been aged but not worked. Knoppwost,&apos; and later Cech and Turnbull,5 found that aged Cu-Fe specimens showed a slight increase in ferromagnet-ism after they had be? cooled in liquid nitrogen. Very recently Denney6 followed the kinetics of the precipitation reaction as indicated by the saturation magnetization, hardness, and diffuse X-ray effects. The experimental results reported by all of the authors mentioned above are readily explained in terms of the Bitter and Kaufmann concept. However, most of the evidence is indirect when used to define the structural changes involved. It was the aim of the present investigation to determine the mechanism of precipitation of iron from a Cu-Fe solid solution in the light of available published information and with additional new and more direct experimental data. Experimental Method At 1060°C copper will contain up to 3 atomic pct Fe in equilibrium solid solution.7 The solvus slopes rapidly toward lower iron concentration, so that at 600°C only about 0.5 pct Fe remains in solution at equilibrium. Below the eutectoid temperature (850°C), the iron-rich phase in equilibrium with the copper-rich solid solution consists of body-centered-cubic a-Fe. Above 850 °C the equilibrium iron-rich phase is face-centered-cubic r-Fe. The alloy to be studied was made by melting electrolytic iron and OFHC copper by induction under vacuum in an Alundum-lined crucible. The metal was held molten for about 5 min, after which time the power was turned off and the crucible was tipped about 70" to promote directional solidification. The resulting ingot, containing 2.5 wt pct Fe by analysis, was clean, large grained, and free of gross segregation. Slices % in. thick were cut from the ingot, rolled to 1/8 in. in thickness, homogenized in hydrogen at 1060°C for 24 hr and, finally, quenched in cold water. Specimens for all subsequent tests were made from these slices. Squares measuring about 1 cm on a side were cut from the slices to make specimens for studying hardness and microstructure. Wires (0.020 in.) for the Debye-Scherrer survey were drawn from 1/8 in. sq rods which had been cut from the homogenized slices. Large crystals for observing diffuse X-ray diffraction effects were prepared by mechanically

Jan 1, 1958

By G. H. Rowe, A. N. Stroh, D. P. Gregory

The magnitude and variation with strain of the parameters activation volume, V*; activation energy, H; and frequency factor, A, in the Arrhenius equation for strain rate are determined for colunlbium single crystals and polycrgstals. Comparison of dislocation structures in deformed columbium crystals as determined by transmission electron microscopy with the rate IF deformation of a metal crystal is a thermally activated process, rate theory can be used as a powerful tool in the study of dislocation mechanisms. seeger1 develops the following expression for plastic deformation governed by a single thermally-activated theory results suggests thai formation of lattice vacancies at jogs in screw dislocations is responsible for thermally activated deformation of poly crystalline columbium. Single crystals appear to deform by a different mechanism. Work hardening in fully recrystallized fine grain columbium is found to result from simultaneous decrease in activation volume and in the number of mobile dislocations with increasing strain. Work hardening in recovered polycrystals is found to result from decrease in activation volume with strain only. dislocation mechanism. His equation for strain rate is where N is the number of mobile dislocations per unit volume held up at energy barriers, v is the frequency with which the dislocations "try" the energy barriers, b is the Burgers vector, a is the area a dislocation loop sweeps out after overcoming a barrier, V* is the activation volume which will be defined more completely later, and t* is the thermal component of stress equal to (Ta-Ti), where Ta is applied shear stress and Ti is internal stress of the dislocation network17&apos; given by

Jan 1, 1963

By B. Chalmers

THE practical importance of the phenomena of melting and freezing must have been recognized for a very long time. The difference between ice and water, for example, has had a profound influence on the history of mankind and the evolution of society. The possibility of melting a metal and allowing it to freeze in a mold of chosen shape has been an essential ingredient in our mastery of the art of shaping metals, and therefore in the evolution of the: machine age in which we find ourselves. The importance of melting and freezing, as applied to metals and alloys, has been so great, in fact, that empirical solutions have been found for the multitude of practical problems that have arisen. This approach has been so successful that relatively little attention has been directed to arriving at an understanding of the fundamentals of the processes. But metallurgy has come to a stage at which we may expect that some, at least, of the more complex problems that have not yet been solved (or perhaps even recognized) may be handled more effectively by scientific study, theoretical understanding, and logical experimentation than by trial and error. In this lecture, therefore, I propose to describe in outline what I think really happens when a metal freezes. In doing so I hope to explain many of the phenomena which have been observed, and in particular to account for the structures that are obtained in actual ingots and castings. The basic problem, to which this lecture represents a tentative partial answer, is this: a mass of metal, containing known proportions of various elements, is melted, heated to a given temperature, and then allowed to freeze under specified conditions. What will be the "structure" of the resulting metal? The term structure includes: 1—crystal size, shape, and orientations, 2—distribution of chemical elements, and 3-—shape, including cracks, cavities, pores, etc. The Solid-Liquid Interface We will first consider what takes place if a single crystal of a metal in the form of a rod is heated, not uniformly, but so that one end is hotter than the other. If this heating process is continued long enough, the hotter end will eventually melt; we will suppose that the rod is in a containing vessel so that the molten metal does not run away, Fig. 1. When some of the metal has melted, we have some solid, some liquid, and an interface or surface of contact between them. If the source of heat is now removed, the interface will move so that some of the liquid freezes, and if the supply of heat is suitably adjusted the interface will remain at rest. This very simple arrangement allows us to study the basic processes of melting and freezing, and if we fully understand this simple case, we may be able to account for what takes place under practical conditions where the heat does not all flow in the same direction, and where the heat flow is determined not by a controllable source of heat but by the heat capacity and temperature of metal and mold, and by the heat loss from the mold surface. The solid-liquid interface is evidently the region of the greatest interest to us; on one side of it there is crystalline solid, and on the other, liquid. In the solid, each atom has a well defined position, around which it vibrates as a result of thermal agitation. It only leaves this position in the relatively rare event of a "diffusion jump." The liquid is much less systematically organized. The atoms are about as far from their neighbors as in the solid, but the arrangement is much less systematic and is continuously changing. The solid and the liquid are represented diagrammatically in Fig. 2. The average energy of the atoms in the liquid is greater than in the solid by an amount that corresponds to the latent heat of fusion, i.e., the amount of heat that has to be supplied to convert unit mass of solid into liquid at the same temperature. The Two Processes As has recently been shown by Jackson and Chalmers,3 many of the features of the processes of freezing and melting can be understood if it is assumed that a continuous and rapid interchange of atoms between solid and liquid always takes place at a solid-liquid interface." It is necessary to con- sider two distinct processes, that of melting, in which atoms leave the surface of the solid and become part of the liquid, and the converse process,

Jan 1, 1955