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By E. J. Ripling, M. W. Toaz

Tensile tests were conducted on pure magnesium and on four commercial alloys over a variety of temperatures and strain rates. The high positive slope of the ductility vs testing temperature curves that is found over a short range of testing temperatures for these materaturerials was shown to result from the fact that the microstructure was not stable at these testing temperatures. Pure magnesium did not display a transition temperature, although its sub-room temperature ductility was low. This poor ductility resulted from a grain boundary weakness. The aluminum-bearing commercial magnesium alloys had a higher ductility than the pure magnesium. ness.Thealuminum-bearingcommercialThese alloys did exhibit a transition temperature and grain boundaries that were stronger than the grain material. Both pure magnesium and the commercial alloys began initiating cracks at strains equal to one half of their fracture strain. MAGNESIUM base alloys' enormous potential as structural materials lies in their high strength to weight ratios. At present, the high strength magnesium base alloys such as AZ-80, which will be discussed later, have about as high a tensile strength to specific gravity ratio as the highest strength wrought aluminum alloy, 75s-T6, and, of course, a much higher strength to weight ratio than structural steel. Further, since the magnesium base alloys develop their high strength to weight ratios with moderate strengths and extremely light weights, the ratio of load carrying capacity to the weight of many structures is relatively higher for the magnesium base alloys than their strength to weight ratios would indicate. This results because judicious increases in section size, such as increasing the depth of bending members or adding reinforcing ribs, can frequently increase the weight of a structure linearly while the strength and stiffness increase as the second or third power of the added weight. Strength and rigidity must be designed into lightweight structures by those techniques which establish multiaxial stress states. Consequently, strong, lightweight structures frequently also possess a greater propensity toward brittle failures. In order to insure against these catastrophic failures, an extra premium must be placed on ductility and toughness in light alloys. Current knowledge of the ductility and toughness behaviors—notched behaviors—of magnesium base alloys is limited.* Most of the magnesium base alloys * The state of present knowledge regarding the effect of testing temperature on mechanical properties has recently been summar-ized by Pritchard.' crystallize in the hexagonal system, so that they are expected to exhibit a tough and ductile to brittle transition temperature. Actually, a transition temperature behavior has been reported for commercially pure magnesium in the literature, but the characteristics of this transition temperatureh uggest that it may not be a transition in the usual meaning, but simply the onset of some thermal softening, a performance much like that which face-centered-cubic lead might exhibit slightly below room temperature. If transition temperatures do exist, such important questions as the manner in which they are affected by internal variables such as chemical composition and grain size, or the effect of external variables such as rate of straining or notch severity have never been investigated. The investigation reported upon herein was undertaken in order to define some of the mechanical behaviors of the magnesium base alloys better, to gain a better understanding of the influence of some of the above listed variables, and particularly to understand their effect on ductility and toughness better. Material and Procedure Material-—Pure magnesium and four commercial alloys were used in this investigation. The alloys were selected so that chemically they made up a series with approximately the same minor alloying

Jan 1, 1957

By A. J. Griest, P. D. Frost, J. R. Doig

During the past 5 years, research directed toward the development of titanium alloys having improved strength-ductility relationships and heat treatability has been carried out at Battelle for the U. S. Air Force.1"3 In the course of this work, a method of heat treating ß-stabilized titanium alloys was developed which produced a wide range of useful properties.4 The heat treatment consisted of two steps: a) solution treating at temperatures in the a-ß phase region, and b) overaging at a lower temperature in the range 800°to 1100°F. Corollary to the alloy development work, research was conducted on the mechanism of the heat-treatment reaction. The investigations of Frost and co-workers on isothermal transformation in Ti-Cr5 and Ti-Mn8 alloys disclosed the existence of the transition phase, w, during the course of the decomposition of metastable ß phase. Subsequent work at Battelle7,8 'and at Case Institute by Brotzen and co-workerss further clarified the nature of the decomposition process. Beta decomposition was found to proceed by the following series of reactions: 1 2 3 4 a + Ti, My (in eutectoid systems). ßo, ßr, and ßU refer to ß phase of original, enriched, and ultimate alloy contents, respectively. Omega is a transition phase having a complex structure. 10-13 Step 4, of course, occurs only in ß-eutectoid systems. This paper presents data correlating transformation behavior with mechanical properties for the following alloys: Ti-8Mn, Ti-8Mn-2A1, Ti-4Mn-4A1, Ti-6Al-4V, and Ti-4Fe. The isothermal decomposition of ß in titanium-manganese alloys of similar composition has been investigated by Frost and coworkers8 and the effect of various annealing treatments on alloys in this system has been studied by Holden, Ogden, and Jaffee.14 The compositions Ti-8Mn-2Al, Ti-4Mn-4A1, and Ti-6A1-4V, were studied in the present work in an effort to define more clearly the effect of aluminum on the response to solution and aging treatments. Ogden, Holden, and Jaffee15 have invesitgated the effect of aluminum on the response of a similar alloy (Ti-8.8Mn-2.46A1) to annealing treatments, and Luini and Lee16 have made an exploratory investigation of the heat-treatability of the Ti-4Mn-4Al alloy. The Ti-4Fe alloy was studied as an example of a moderate1y active eutectoid system. Holden, Ogden, and Jaffee 17 have studied Ti-Fe alloys and have presented some hardness data for aging at 400°C for alloys of Ti-2.56, 5.01, and 7.35 iron after quenching from the a-ß field (750°C). EXPERIMENTAL PROCEDURES The experimental procedures generally used in the precipitation-hardening study are outlined in the following sections. Departures from these techniques are noted in the discussions of the individual alloys. Table I presents the nominal and actual compositions of the alloys studied, together with the swaging and rolling temperatures. In all cases, these temperatures were in the a-ß field and the microstructures consisted of an equiaxed a-ß mixture. Melting and Fabrication—Except for the commer-

Jan 1, 1960

By L. Burris, G. A. Bennett, P. A. Nelson

Staznless steels ave rrzuch less seuel-ely attacketl by zinc lapov than by molten zinc systems. To determine the applicubility of stainless steels fov equipment items which would be exposed only to zinc vapor, the corvosion resistance of several stainless steels to zinc vapor was investigated in 500-hr- exposuves at 900° C. Two austenilic stainless steels, Types 304 and 3-17, suffered severe intergranular attack. Loss of nickel from these steels and micro-probe crnalyses indicated the probable formation of low-melting zinc-rich phases. Good vesistance to attack by zinc vapor, as determined by metallo-9-rpaphic excr)rtination and physical testing, Loas exhibited by Type 405 stainless steel, a jevvitic steel. Type 440C stainless steel, a tnartensitic steel, although suffering considevable loss in weight, was not structurally weakened. At Argonne National Laboratory liquid zinc and zinc-rich alloys are being employed as solvent media in pyrometallurgical processes under development for recovering and purifying discharged reactor-fuel materials.' Because of the high solubility of most metals in liquid zinc,' there are few metals which can serve as container materials for liquid zinc at temperatures much above the melting point. The poor corrosion resistance of the common metals and alloys to molten zinc has been discussed by Eldred.3 DeKany, Lavendel, and Burris4 tested some of the less-common alloys of the nonrefractory metals and found that they also have poor resistance to attack by liquid zinc and Zn-Mg solutions. Of the metallic materials, only some of the refractory metals such as tantalum, tungsten, and molybdenum and their alloys have shown resistance to attack by liquid zinc.4-7 In the process-development work at Argonne National Laboratory, tantalum and tungsten vessels have been used to hold liquid-zinc alloys. These vessels are housed in stainless-steel equipment, principally to permit maintenance of an inert atmosphere and operation at various pressures. In this experimental work, it has often been observed that Type 304 stainless steel suffers little attack by zinc vapor at temperatures up to about 800 °C. In contrast, severe attack occurs in areas where zinc condenses as a liquid. The extent of attack by zinc vapor on several stainless steels was investigated because there are possible useful applications of stainless steel for equipment items which would be exposed only to attack by zinc vapor. These applications include, for example, housings around processing equipment as indicated above and vapor lines leading from zinc stillpots or retorts to condensers. EXPERIMENTAL EQUIPMENT AND PROCEDURE The experimental procedure consisted in exposing stainless-steel test coupons to static zinc vapor for 500 hr at a temperature of about 900°C. The apparatus, shown in Fig. 1, consisted of a 12-in.-long closed tantalum tube jacketed in Type 304 stainless steel to prevent oxidation of the tantalum. A Zn-10 wtpct Mg* alloy charge of 293 g in the bottom of the

Jan 1, 1965

By R. C. Wilcox, R. V. Lawrence, P. A. Calhoun

The V-Rh-Si, Nb-Rh-Si, and Cr-Rh-Si systems were investigated to evaluate the extent of solid solubility between binary Cr3O structures. All Cr3O components lie within the favorable atomic size range1 for the formation of the Cr3O structure. Previous studies indicated complete solid solubility between V3Si and V3Sn2 and no solubility between Sn and cr3si3 or Si and Nb3Rh.4 The Cr3O structure is not found in the Cr-Sn and Nb-Si binary systems. Powder compacts were are-melted under argon, homogenized in a vacuum at 1000°C, and water-quenched. Alloys containing greater than 15 at. pet Si were annealed 100 hr. X-ray diffractometer methods were used to study all systems utilizing copper radiation for the Cr-Rh-Si system and chromium radiation for the other two systems. V-Rh-Si System. The alloys prepared in the V-Rh-Si system and the measured lattice parameters for the Cr3O structure found in each alloy are listed in Table I. The characteristic Cr3O diffraction pattern was found in all binary alloys. The measured lattice parameter of V3Rh (VRS-1) is in agreement with the reported value5 while the value for V3Si (VRS-5) is only in fair agreement.' In

Jan 1, 1965

By A. J. McEvily, R. C. Boettner, C. Laird

The processes leading to fatigue failure in the low-cycle range were studied to obtain an understanding of the basis of Coffin's law. Particular attention was paid to the manner of mack nucleation and to the determination of the portion of the total lifetime spent in crack propagation. It was observed that cracks are nucleated at surface notches meated during reversed plastic deformation, Principally at grain boundaries, and that crack growth occupies greater than 75 pct of the fatigue life. When inclusions are present, they are often the primary sites of crack initiation. From consideration of the incremental distances advanced by a crack in each cycle and of the lifetimes of different materials, it appears that the universality of Coffin's law arises from the similarity of crack-tip strain for a given applied strain amplitude, irrespecti7:e of material. STUDIES of low-cycle fatigue have been particularly rewarding in that a simple correlation between fatigue lifetime and plastic-strain range has been shown to exist. This empirical relationship takes the forml,2 Nner=C [1] where N is the number of cycles to failure, Er is the plastic-strain range defined as the total plastic strain per cycle (the sum of the tensile plastic strain plus the compressive plastic strain), n, the exponent, is a constant independent of the material to a first-order approximation and often has a value of about 0.5, and C is a constant which for a number of ductile and metallurgically stable materials is approximately 0.65 for fully reversed strain.3 This relationship has been shown to be valid for a wide variety of metals and alloys and is often referred to as Coffin's law. However, it is yet not clear why so many different metals and alloys exhibit the same over-all behavior in the low-cycle range. Less is known about the mechanisms of fatigue failure in the low-cycle range than in the high-cycle range, but it is expected that the processes would have much in common. For example, an important characteristic common to both high-and low-cycle ranges is that fatigue crack propagation can occupy a considerable portion of the total lifetime. This has been shown by Laird and smith4 by a count of the number of ripples present on the fracture surface of copper specimens, each ripple corresponding to one cycle of load. On the other hand, some marked differences have been noted. For example, in copper,5 nickel, and aluminum4 there is a greater tendency for grain boundary cracking to occur as the strain amplitude is increased. The purpose of the present work was to study further the processes leading to fatigue failure in the low-cycle range and to obtain an understanding of the basis for the validity of Coffin's law. Particular attention has been paid to the manner of crack initiation and to the determination of the portion of total lifetime spent in crack propagation. MATERIALS A number of metals and alloys were tested, although more attention was focused on a particular few, once general trends had become apparent. Alloys were selected so as to investigate the effects of crystal structure, number of phases, stacking-fault energy, and impurity content. The materials in this grouping were: high-purity copper OFHC copper 1100 aluminum (commercially pure) 70-30 brass (Cu-30 wt pct Zn)

Jan 1, 1965

By R. J. Sherman, M. R. Achter

Creep rates have been measured for nickel at 1200°F in air and in vacuum, and related to the depth of surface cracks in the specimens tested in the two environments. The surface cracks were observed to grow faster in air than in vacuum in both short-time and long-time tests. Crack depths, however, could be correlated with creep rates only in the short-time tests. Although no cracks were observed in a relatively brittle Ni-Cr-A1 alloy creep tested at 1300°F, the appearance of alloy specimens fractured after interruption in creep indicated that there was a decrease of intergranular cohesion resulting from the diffusion of atmospheric gases into grain boundaries. PREVIOUS investigations1-3 have demonstrated that the effect of atmosphere on the high-temperature strength of nickel and nickel-base alloys may be reversed by changes in the stress or temperature of the test. At high stress and low temperature the creep-rupture life is longer in vacuum than in air, while at high temperature and low stress the relative strengths are reversed. A mechanism to explain this behavior was proposed based on two competing processes. Oxidation hardens and strengthens a metal, but adsorption of gas, by lowering surface energy, reduces the work required to propagate an intergranular crack. The controlling process is determined by temperature and stress. Basic to this mechanism is the concept that environment determines the ease of propagation of intergranular cracks and thus affects creep rate and rupture life. To explore the mechanism in greater detail, this investigation was undertaken to determine the effect of atmosphere on crack initiation and propagation rates in nickel1 and a relatively brittle Ni-Cr-A1 alloy.3 EXPERIMENTAL PROCEDURE The materials were the same as used previously. Nickel of 99.8 pct purity had the following impurity content: 0.02 pct Co, 0.03 pct Fe, 0.02 pct Mg, 0.03 pct Si, and 0.008 pct C. Rods were cold swaged from 1/2 in. to 3/8 in. in diam and machined into speci- mens with gage sections 1/4 in. in diam by 1-1/4 in. long. They were annealed in vacuum at 1600" F for 1 hr to achieve a grain size of 0.07 mm. The alloy was composed of 19.1 pct Cr, 3.6 pct Al, 1.7 pct Si, 0.57 pct Fe, 0.10 pct Mg, 0.027 pct C, 0.004 pct S, and the balance Ni. Sheet specimens of 0.500-in. width and 1.25-in. gage length were machined from 0.033-in.-thick strip and annealed in vacuum at 2000° F for 2 hr, resulting in a grain size of 0.036 mm. Large amounts of an intergranular phase were introduced as a result of the anneal. Creep tests on both materials, using equipment previously described by Shahinian; were conducted in air and in a vacuum of 1 x 10-5 mm Hg. They were continued to rupture and interrupted at various stages of the test. Although extensions were measured from dial gage readings of movement of the lever arm, it is considered that valid comparisons of creep rate may be made in this manner.l Two parameters were measured metallograph-ically* on longitudinal sections of nickel specimens interrupted during creep—the mean depth of surface cracks 2, and the density n as the number of cracks per unit length of surface. They were combined into one parameter by the relation p = n-d [1] As used here, p is equal to the total of crack depths per unit length of surface. RESULTS Nickel. It is well known that the incidence of intergranular cracking is increased with increasing rupture life. Indeed, it is evident in Tables I and I1

Jan 1, 1962

By A. R. Troiano, W. J. Barnett

IN recent years the demands of space limitations and increased loads, particularly in the aircraft industry, have accelerated the trend toward utilization of ultra-high strength steels. The increased application of such steels has focused attention on the phenomenon of delayed failure or static fatigue. Parts subjected to a static load, after successfully sustaining the load for an extended period of time, may fail in a brittle manner. Subsequent examination of such failed parts by conventional laboratory tests in many cases reveals no apparent signs of embrittle-ment. That is, the material exhibits good ductility and impact resistance. Furthermore, consideration of the service conditions (environment and temperature) under which this brittle failure occurred eliminates its classification as a stress-corrosion or creep type of failure. Recently this unique time dependent fracture process has been the subject of numerous investigations. The most significant result of these studies was the recognition of hydrogen as the source of embrittlement. While the effect of numerous variables on the occurrence of static fatigue failure has been defined, an adequate explanation of these behaviors is lacking. The occurrence of this phenomenon in steels generally has been limited to strength levels in excess of 180,000 lb per sq in. Although steels at high hardness levels have exhibited delayed failure in the absence of hydrogen introduction after heat treat-ment,'.' this behavior was most prevalent when the material in processing had been subjected to an environment conducive to the absorption of hydrogen. Delayed failures have been produced in high strength steels by the introduction of hydrogen

Jan 1, 1958

By R. J. Sherman, M. R. Achter

IT has often been reported that coatings may strengthen single crystals and polycrystalline specimens of coarse grain size. This note reports the effect of a surface layer of fine grains on the creep-rupture properties of coarse grained nickel 99.8 pct pure. Sheet specimens were machined from cold-rolled material, finished by surface grinding and annealed at 1230°C in vacuum for 4 hr. Half of them were re-ground on faces and sides and annealed at 925"C for 1 1/4 hr. The double grind-anneal treatment produced a structure consisting of a surface layer of fine grains of 0.02 mm diam on a coarse-grained substrate of grains of 0.40 mm diam. Comparison with specimens which received only the single grind showed that the second grind-anneal did not alter the structure of the substrate. Several modifications were made in specimen design to eliminate erratic rupture behavior outside of the gage length. In the final design the gage section, oriented parallel to the rolling direction, was 1.00 in. long, 0.4 in. wide, and 0.0625 in. thick. End tabs were pinned in place. Although other configurations (welded end tabs, transverse orientation, and wider gage section) showed similar effects of a layer of fine grains, because of the erratic rupture behavior, only the data from the final design are reported here. Creep tests were performed at 650°C in air and in a vacuum of 10~5 mm Hg. Creep rates were measured from the movements of the lever arm. Since the broken sheet specimens could not be fitted together, elongations at fracture measured by means of gage marks, do not include the necked portions adjacent to the fracture. Results of the creep tests with and without the fine grained surface are given in Table I. An explanation for the lower strength in air than in vacuum has been proposed previously in terms of the reduction of crack propagation energy by gas adsorption.l In both environments the fine-grained layer produced a small decrease in minimum creep rate and, except for the shortest test in air, a substantial increase in rupture life. The elongation values at rupture, as they are measured here, do not demonstrate an appreciable effect of the grain structure but do show much smaller ductilities in air than in vacuum. Greater elongations maybe associated with the finegrained layer, however, especially in vacuum, in creep curves such as those in Fig. 1. Duplicate tests were not run to determine the re-producibility of the data. It is felt, however, that, in view of the consistent behavior, the prolongation of rupture life may be attributed to the effect of the fine-grained surface. In the data not reported be-

Jan 1, 1962

By S. Wiederhorn

The energy and force of interaction between a crack and a slip band have been calculated. When the distance between the crack and the slip band is greater than the dislocation spacing of- the slip band, the force on the cgrack is perpendicular to the slip band. The energy needed to propagate a crack against this force can be lawe enough to pevmit crack propagation parallel to noncleavage planes. SEVERAL recent experiments have demonstrated the occurrence of crack-slip band interaction in ionic crystals having a NaC1-type structure. Stokes, Johnston, and Lil have studied the nucleation and growth of cracks in MgO. From their observations, they conclude that slip bands have the ability to halt and stabilize cleavage crack nuclei. Gilman2 has shown that a crack, propagating along a (100) cleavage plane in a previously deformed LiF crystal, tends, at certain points within the crystal, to switch to (110) planes. He reasons that plastic deformation reduces the cohesive energy of the (110) planes below that of the (1130) planes, making possible (110) crack propagation. Keh3 has demonstrated that when indentations are made in MgO, crkcks form along (110) planes at the indentations. Gilman2 attributes the formation of these cracks to plastic cieformation. There is some doubt as to the formation of cracks on (110) planes. Kuznetsov4 in his discussion of crack propagation in rock salt, shows that splitting along (110) planes requires 1.76 times as much energy as splitting along a stepped (100) surface. Therefore, the observed (110) cracking in LiF and

Jan 1, 1962

By W. F. Domis, F. von Gemmingen, R. W. Whitmore, F. Garofalo

Constant-load creep-rupture tests at 1100°, 1300° and 1500°F were made on a Type-316, 18 Cr-8 Ni-ZMo, austenitic stainless steel to determine the relationship between ruptzire life and other aspects of creep behavior, The test program was designed on a statistical basis to check the variability of various creep properties and to determine emprically the dependence of minimutn creep rate and ruptzire life on initial stress, The dependence of the rupture life, tn, on the initial stress, a,, is found to be related to the stress dependence of the wrinimum creep rate, i,, and secondary creep strain, The instabilities or breaks found in the conventionual log - logt, plots can be traced to either a change in the linear dependence of log i, on log a, or to a change in secondary creep strain. In the alloy tested, rupture is found to be predowmanantly of the inter crystalline type, For this material, the secondary creep strain is -found to depend on the nature of the grain boundary precipitate, which affects grain boundary migration. It has been shown for a variety of metals and al tested at elevated temperatures under creep conditions that the empirical relationship exists between rupture life, tr, and minimum creep rate, i,. This relationship, in which C and a are in some cases independent of stress or temperature, shows in a general way that creep rupture depends on creep behavior prior to tertiary creep stage. To determine more fully the factors which control creep rupture requires, therefore, more knowledge of the relationship between rupture and creep behavior. To obtain such information, constant-load creep-rupture tests have been made on a Type-316, 18 Cr-8 Ni-2 Mo, austenitic stainless steel at 11000, 1300°, and 1500°F. These tests were designed on a statistical basis to determine the variability of the various creep properties measured and to determine empirically the stress dependence of minimum creep rate and rupture life. As a result of this work, various empirical relations have been established defining more closely the factor, C, in the relation between rupture life and minimum creep rate. This factor is found to be proportional to the amount of strain during secondary creep. The variation of secondary creep strain with rupture life and temperature is determined and its effect on rupture behavior discussed. The variability in minimum creep rate and rupture life is also dis- cussed and the dependence of these quantities on initial stress is interpreted. TEST MATERIAL AND TEST PROCEDURES The material tested is an austenitic stainless steel, AISI Type 316, of the following composition: C, 0.07 pct; Mn, 1.94 pct; P, 0.01 pct; S, 0.021 pct; Si, 0.38 pct; Cr, 18 pct; Ni, 11.4 pct; Mo, 2.15 pct; All 0.003 pct and N, 0.043 pct. The material was received as a hot-rolled 0.5-in. diam bar which was sectioned into 3.25 in. long blanks. Each blank was heat treated by holding 1/2 hr at 2000° F followed by water quenching. A 3/8-in. coupon was removed from each heat-treated blank and the microstructure was examined, the grain size determined, and the hardness measured. The microstructure was found to be uniform from blank to blank and the grain size was found to range from 4 to 6 ASTM numbers. Hardness, measured using a 20-kg load, varied between 127 and 148 Vpn. Tensile creep-rupture specimens having a 0.25 in. diam within a 1.5-in. reduced section were machined from the heat-treated blanks. Constant-load tests were made on these specimens at 1100°, 1300°, and 1500 F. Six specimens were tested at each stress level. Three creep-rupture machines were employed for tests at each temperature, two specimens being tested in each machine for each stress level. To minimize any effect due to local inhomogeneities along the length of the as-received bar the specimens were randomized before test. Each specimen was tested to rupture and in most cases an autographic extension-time curve was obtained. For very low stress levels at 1500°F no auto-

Jan 1, 1962

By N. J. Grant, A. W. Mullendore, Y. Ishida

Creep tests were performed at 500°F on polished specimens of A1-3 pct Cu at stresses of 2000, 4000, and 6000 psi. The effects of the size and distnbution of the second phase were studied in connection with the mechanisms of deformation, especially grain boundary sliding. Although the amount of grain boundary sliding varied considerably for the different heat treatments, yestriction of grain boundary sliding is not due primarily to precipitate on the grain boundary. NUMEROUS prior investigations of grain boundary sliding, notably those of McLean, 1 Rachinger, 2 Rhines, Bond, and Kissel, 3 and Brunner and Grant,4 have established the importance of this phenomenon in the deformation of pure aluminum and its solid-solution alloys at elevated temperatures. An extension of this kind of study to the case of a simple two-phase structure is an important step towards understanding the deformation behavior of more complex alloys. A two-phase A1-3 pct Cu alloy was selected for this study. The creep behavior of alloys in this system has been studied by Sully and Hardy,5 by Underwood, Marsh, and Manning,' and by Pelloux, Chaudhuri, and Grant.7 In these studies, the prior heat treatments were such that some precipitation took place during the tests in most cases. Pelloux, Chaudhuri, and Grant7 studied the creep of A1-2 and 3 pct Cu alloys in the overaged condition. Their results indicated that the structure most effective in promoting high-temperature strength consisted of a fine, closely spaced grain precipitate and large, widely spaced particles on the grain boundary. The present work consists of quantitative meas- urements of grain boundary sliding during creep of an A1-3 pct CU alloy with four different precipitate distributions. The aims were to quantitatively define the role of the second phase on grain boundary sliding, and to obtain information regarding the mechanism of grain boundary sliding. MATERIALS AND PROCEDURE A high-purity A1-3 pct Cu alloy, kindly supplied by Alcoa, was used in this investigation. The composition of the alloy in weight percent was as shown in Table I. This composition is such that A12Cu exists as a second phase at temperatures below 850°F. A homogeneous grain size of 0.6 to 0.8 mm was obtained by annealing the machined test bars for 5 min at 1000°F, furnace cooling to 930°F, and homogenizing for 5 hr. Following the recrystalli-zation and homogenizing anneal, four different heat treatments shown in Fig. 1 were utilized. The F7 and F5 treatments utilize a furnace cool to the aging temperature while Q7 and Q5 have a water quench from the homogenization temperature before aging. Figs. 2 to 5, and Table 11, show the size and shape of precipitate dispersions obtained with the above heat treatments. The heat treatments produce variations in the dispersion both in the grains and along the grain boundaries. Precipitate particles along the grain boundaries are rounded and large (about 1.5 p diameter) in the F5 and F7 structures (Figs. 3 and 5). The grain boundary precipitate was somewhat smaller (0.9 p diameter) in the Q7 structure and no grain boundary precipitate was detected in the Q5 structure using light microscopy. The pre-

Jan 1, 1964

By Nicholas J. Grant, Italo S. Servi

Extensive data of minimum creep rates and rupture times for high purity and commercial aluminum confirm the existence of a transition range from the low temperature-type to the high temperature-type behavior. The data are analyzed in line with the suggested theories of deformation of metals. CONSIDERABLY more work has been done on the rupture and creep behavior of commercial alloys than on the less complex, high purity metals. As a result, certain fundamental behaviors are overlooked. One of the lesser known variables affecting the high temperature behavior of metals is the impurity content (or alloy content). Impurities are known to affect strongly the recrystallization temperature, electrical resistivity, and other properties of metals, but the effect on creep and other high temperature mechanical properties has been determined for very few materials. In view of these facts aluminum appears to be a fairly ideal metal, since it is available in a wide range of purity contents from 99.995+ down through the 2s (99.0 + Al) and 3s (1.2 Mn) grades in small steps. It is highly resistant to oxidation and other surface instabilities, leaving recrystallization and grain growth as virtually the only instabilities. Control of grain size is also well standardized. Limited data on the creep of aluminum have been obtained by Dushman et a1.l for high purity aluminum; by Sherby2 on 2s aluminum of variable grain size; and by Dorn and Tietz³ on cold-worked 3s aluminum. Experimental Procedure The apparatus consisted of a constant stress, creep testing unit similar to the type described by Hop-kin.' The specimens were held at temperature for 30 min before loading. The loading time was less than 2 min. Factors affecting the accuracy of the data were: 1-—the temperature gradient and control, ± 3°F; 2-the stress determinations, ± 1.5 pct; and 3—elongation measurements (1 in. gage), ? 0.001 in. Three grades of aluminum were tested: 1— high purity (99.995 pct Al), 2—2s (99.3 pct Al), and 3—3s (98.2 pct Al). The spectrographic analyses of these materials are reported in Table I. The specimens were annealed then electropolished in a perchloric-acetic mixture before testing. Data on heat treatments and grain size obtained appear in Table 11. The entire creep curve was determined for almost all of the tests. A few tests were interrupted after the beginning of the third stage of creep, and in these cases the rupture time was estimated. Experimental Results Constant stress creep-rupture testing gives a much longer, more accurately measurable second stage of creep than constant load testing. This behavior is in agreement with the results obtained by Andrade for lead (see Fig. 1). If the test is run to completion, the third stage of creep appears as soon as stress intensification occurs. In the case of specimens which fail in a ductile manner, the stress intensification begins when the sample starts "necking". This differs from the "high temperature brittle" specimens in which the stress intensification is caused by intercrystalline cracking. These specimens

Jan 1, 1952

By C. S. Roberts

The creep mechanism and kinetics of fine-grained magnesium have been studied over the temperature range 200&apos; to 600°F. As a result of a photographic study of microstructural changes, transient and steady-state creep components have been correlated with slip, subgrain formation, and cyclic deformation at the grain boundaries. THE approach of this research has been the blend of a quantitative study of the creep strain of polycrystalline magnesium as a function of time, stress, and temperature with direct microstructural observations of the operative deformation processes. The validity of the conclusions is dependent on the condition that the microstructural changes seen on the polished surface qualitatively represent those occurring in the bulk of the metal. The work was intended as much to lay a background to a study of highly creep-resistant magnesium alloys as to provide a description of the behavior of the base metal itself. The spectroscopic analysis of the electrolytic magnesium used in this study is as follows: Al, 0.009 pct; Ca, <0.01; Cu, 0.0011; Fe, 0.021; Mn, 0.012; Ni, 0.0004; Pb, 0.0012; Si, <0.001; Sn, <0.001; and Zn, <0.01. The impurity level is approximately that of commercial magnesium alloys. The original ingot was melted under Dow type 310 flux and cast as a 3 in. diam billet. It was extruded into 1 in. flat stock under the conditions: billet preheat 800°F (1 hr), container and die temperature 800°F, speed 3 ft per min, and area reduction ratio 45:1. The extrusion process was chosen in preference to rolling and recrystallization because it allowed easier grain size control from specimen to specimen. The grains of the extruded metal were fairly equi-axial and uniform in the size range of 4 to 6 thousandths of an inch. The preferred orientation of basal planes about the transverse direction was determined by an X-ray diffraction surface reflection method. A beam of filtered copper radiation was directed at an angle of 17" to both the transverse direction and the surface yet perpendicular to the extrusion axis. Analysis of the (002) diffraction arcs in the resulting photographic patterns gave an approximate intensity distribution along the great circle which extends through the center of the basal plane pole figure and to the extrusion axis poles. Successive layers of metal were removed by macro-etching between exposures. The extruded texture is relatively sharp, but the most significant point is the position of the maximum basal plane pole density and its variation with depth below the surface. Fig. 1 shows that this maximum is rotated 15" from the normal at the surface toward the extrusion direction. Such an inclination has been reported for extruded 1 pct Mn and 8 pct A1-0.5 pct Zn alloys.&apos; The inclination decreases until the maximum splits at about 0.025 in. depth into two elements of equal and opposite rotations from the ideal. The double texture persists to as great a depth as was experimentally convenient to examine. It probably continues to the very center of the extrusion. There is no great change in the sharpness of the individual elements of the texture with depth. A plate of metal about 0.015 in. thick at the surface of the extruded stock was produced by etching. A transmission diffraction pattern was made for the purpose of determining any preferred orientation of a direction in the basal planes. Relatively uniform {loo) and {101) rings were produced. There is little tendency for parallelism of a given direction in the plane with the projection of the extrusion axis on it. The creep specimens were machined from 6¼ in. lengths of the extruded stock. Creep was measured on the reduced section, ½x1/8X2¼ in. long. This section was electropolished on one side for the studies of microstructural changes during creep. An orthophosphoric acid-ethyl alcohol electrolyte was used under the conditions recommended by Jacquet.² Hand polishing was used for previous mechanical preparation. Electropolishing was continued until all mechanical twins had been removed. The electro-polished surface was protected from oxidation during creep testing by a thin layer of silicone oil. All micrographs were taken at room temperature on conventional metallographic equipment and after removal of the oil film. The creep tests were performed with machines which have been described in detail by Moore and McDonald." Five testing temperatures, 200°, 300°, 400°, 500°, and 600° ±3°F were used. Difference in temperature between the two ends of the specimen reduced section was 2°F or less. The testing was done at constant load. Strain readings were taken as frequently as necessary to develop usable creep curves. Tensile Creep vs Time, Stress, and Temperature A definition of terms is necessary. Whenever successive sections of a creep strain-time curve show decreasing, constant, and increasing slope with time they will be termed primary, secondary, and tertiary

Jan 1, 1954

By C. S. Roberts

Four binary alloys in this system were creep tested at 300°&apos; to 600°F. A photographic study of microstructural changes showed that the outstanding creep resistance results primarily from a potent precipitation hardening locally at grain boundaries. Twinning was correlated with the primary stage and nonbasal slip with the tertiary stage of creep. THE utility of the rare earth metals as ingredients of magnesium alloys destined for elevated temperature application has been recognized for about 15 years. A review of the early development of these alloys has been given by Leontis. 1,2 It has been only recently that a desire has arisen to understand basically their remarkable creep resistance. The present research had that object. The work of Leontis showed that the differences between the behavior of the individual magnesium-rare earth element binaries was of commercial significance. However, the variations were small compared to the overall difference in creep resistance between the group as a whole and the magnesium alloys of the Al-Zn type. It was decided to conduct this investigation with one binary system which would be representative of the magnesium-rare earth element alloy group. Cerium was chosen as the solute because of its near-middle position in the series and because of its availability in the commercially pure form. A previously published description of the creep behavior of electrolytic magnesium" serves as the background to this research. The experimental approach used in that investigation has been partially repeated here. A series of binary alloys was planned to include the following: 1—A dilute alloy which would contain the cerium almost entirely in solid solution. This would assess the solid solution effect. 2—An alloy which is nearly saturated at the selected solution heat treatment temperature, 1070°F. Here the effect of nearly maximum precip- itation would be obtained without the influence of undissolved Mg-Ce intermetallic. 3—An alloy containing slightly more cerium than is soluble at 1070°F. This would allow the investigation of the effect of a small amount of excess intermetallic. 4—A high cerium alloy in which the influence of a large volume of undissolved intermetallic may be examined. Experimental Details In order to specify compositions of the binary alloy series a knowledge of the solubility of commercially pure cerium in solid electrolytic magnesium and also of the composition of the terminal intermetallic phase was necessary. The latter has been proven to be approximately Mg,Ce by Voge14 and Haughton and Schofield.5 A series of both cast and extruded alloys were solution heat treated 24 hr at 1070°F for a solubility determination. Lineal analyses of the polished and etched microstructures were accomplished with the aid of a Hurlbut

Jan 1, 1955

By Milton R. Pickus, Earl R. Parker

THE modern theories of creep¹-4 in general have been based upon the concept of generation and migration of dislocations, with the generation process normally assumed to be rate controlling. The theories are generally deficient in that they fail to take into account many factors that are known to influence creep. The influence of the state of the surface of the test specimen has been almost completely overlooked; yet the present report shows that the nature of the surface may, in certain cases, govern the creep characteristics of a specimen. In the period since Taylor" applied the concept of dislocations to a study of metals, a school of thought has developed that closely relates the plastic deformation of metals to the generation and migration of dislocations through the crystal lattice. It might be expected that the thermal energy required for the generation of a dislocation would be different from that for migration of the dislocation through the lattice. Furthermore, the activation energy for generation would be expected to vary for different parts of the solid metal. It has been predicted that dislocations would be generated most easily at external surfaces, but could also be activated at certain internal surfaces such as grain or phase boundaries. Within the body of the metal a range of values for the activation energy might be expected because of different degrees of disorder at such regions as grain boundaries, impurities, and second-phase particles. The particular value of the activation energy that was rate determining could then depend on the specific conditions of a test. If, for example, the surface atoms were by some means constrained, the generation of dislocations in the body of the metal might become the important factor. On the other hand, other conditions may favor generation at the surface. It is possible then that the creep behavior may not be completely determined by the inherent properties of the metal. Even the environment in which a test is carried out could have a significant effect. In fact it is conceivable that in order to obtain the maximum creep resistance from a given alloy, the surface atoms must be so constrained that the activation energy for generating dislocations on the surface is at least equal to that required for generation in the body of the metal. On the basis of such considerations, and in view of the limited number of publications discussing this subject, it seemed that an investigation of the influence of the state of the surface on creep might yield information of both theoretical and engineering interest. Experiments on single crystals, demonstrating a variation in the mechanical properties due to alterations in the surface layer, have been reported by several investigators.6-13 he results of these experiments have been briefly summarized;14 consequently, the earlier work will not be reviewed here. As an example of these findings the observations of Cottrell and Gibbons may be cited. They reported the critical shear stress of a lightly oxidized cadmium single crystal is greater by a factor of 2½ than a specimen with a clean surface. Materials and Methods Single crystals M in. in diam and 8 in. long were prepared from Horse Head Special zinc, melted under an atmosphere of helium in a large pyrex test tube, and drawn up into a long ½ in. diam pyrex tube by means of a vacuum pump. The cast zinc rods thus produced were cut into convenient lengths and sealed in evacuated pyrex tubes. Single crystals were grown by gradual solidification of the remelted rods. Cleaving the ends of the single crystal specimens chilled by liquid nitrogen proved a simple method for determining orientations from the exposed basal plane from the markings left on the cleaved surface that gave the slip directions with sufficient accuracy for the experimental work. The specimens chosen for the experiments were those having the angle between the basal plane and the specimen axis within the range of 15" to 65". Since zinc single crystals are quite delicate, it was necessary to devise an appropriate method of gripping the specimens in order to suspend them in the furnace and apply the load. Stainless steel collars were prepared having an inside taper, the smaller end of the taper being of such a size that the specimen could just pass through freely. The tapered hole did not extend the full length of the collar; a sufficient thickness of metal remained so that a hook could be attached to provide a means of applying the load and suspending the specimen. One of the collars was slipped over the upper end of a specimen which was supported vertically in a steel jig. The collar was then heated electrically until the end of the crystal melted and filled the collar with molten zinc. At this point the application of heat was discontinued, whereupon the molten zinc quickly solidified, due to the chilling effect of the jig. The specimen was then inverted and the second collar applied in a similar manner. The jig served several purposes: limiting the length of specimen that was melted, providing excellent alignment of the collars with respect to the specimen axis, and protecting the specimen from mechanical damage. Once the specimen was suspended in the furnace and loaded, it was desired to accomplish the surface treatment with a minimum of disturbance of the specimen. Around the specimen was a long pyrex tube, the upper portion of which was approximately 1 in. in diam, and in it was a copper coil of such a diameter to fit snugly against the tube. A specimen, approximately ½ in. in diam and 4 in. long, was suspended by means of a stainless steel rod so that it hung within the copper coil. The lower portion of the glass tube was approximately ¼ in. in diarn, and passing through it was a 5/32 in. diam stainless steel rod which hung from the lower specimen collar. This portion of the glass tube and the stainless steel rod extended through the bottom of the furnace. A T-connector, with suitable packing, was attached to the lower end of the stainless rod to provide a water-

Jan 1, 1952

By ED. E. Furman, R. H. Atkinson

HITHERTO the practical creep testing of precious metals has received little or no attention. The only previous creep tests of precious metals have been made with wires under conditions such as to yield much more rapid rates of creep than in engineering tests.&apos;, &apos; Up to the present time the value of creep bars of adequate size, in the absence of real need for engineering data, has deterred investigators. However, the increasing use of platinum at high temperatures has demonstrated the need for reliable creep data for the guidance of engineers, especially those engaged in designing certain specialized chemical plant equipment. In order to supply this need, creep tests were conducted at 1382°F (750°C) on 0.290 in. diam specimens of platinum, 90 pct Pt, 10 pct Rh and palladium. The platinum was high purity, nominally 99.95 pct Pt. The 90 pct Pt, 10 pct Rh was of the same high quality as is used for making gauzes for the catalytic oxidation of ammonia. The palladium was also of high purity; two batches of palladium bars were tested, one deoxidized with calcium boride and the other with aluminum. Spectrographic examination of the palladium confirmed its good quality; the only significant impurities apart from the residual deoxidizers were traces of silicon and lead. Procedure The creep bars, which were furnished by Baker and Co. to our specification, were 6 ¾ in. in overall length with a 4½ in. (4 in. gage length) reduced section 0.290 in. in diam and had the ends threaded (?-NC16). It may be of interest that the bars were valued at up to $600 each. The specimens were supplied in a 50 pct cold-worked condition to facilitate attachment of the creep extensometer, which was of the push rod type. Because of the softness of the platinum and palladium, the extensometer rings were secured to the test section by means of circular knife edges instead of the usual pointed set screws. The extensometer rods extended through the bottom of the furnace and readings were taken with a 0.0001 in. "Last Word" dial gage fastened to the rods for the duration of the test. The bars were directly loaded by hanging weights from the lower specimen grip. All tests were conducted at 1382°F ± 2°F, and an effort was made to maintain the temperature gradient over the test section within 2°F. The ends of the furnace tube were packed with asbestos wool, which allowed a very slow circulation of air through the tube. Annealing was accomplished in the creep furnace before the load was applied. The platinum and palladium specimens were annealed at the test tem- perature for about 17 and 24 hr respectively; in the case of the rhodioplatinum it was found expedient to anneal for 1 hr at 1922°F (1050°C). Pilot samples cut from the same stock as the bars were used to check annealing procedures. Pertinent measurements of grain size and hardness were recorded. Results and Discussion The creep data obtained are given in Table I and the creep curves are plotted in Figs. 1, 2, and 3. Two platinum specimens, tested under a stress of 250 psi, had almost identical creep rates at 2000 hr, namely 0.000008 and 0.000009 pct per hr. A third platinum specimen, stressed at 400 psi, had a creep rate at 2000 hr of 0.000026 pct per hr; the reason for a rather sharp decrease in creep rate during the period from 1200 to 1600 hr is unknown. As it was thought that 90 pct Pt, 10 pct Rh would have a lower creep rate than platinum, the first sample was tested at 400 psi; however, the creep rate was approximately 50 pct greater. Microex-amination revealed that differences in grain size might be responsible for the unexpected result, as annealing at 1382°F developed an average grain diameter of 0.0021 in. in the rhodioplatinum specimen compared with 0.004 in. in the platinum bar. Annealing the alloy for 1 hr at 1922°F (1050°C) increased the average grain diameter to 0.0032 in. and materially improved the creep resistance, making it much better than platinum. A second specimen annealed at 1922°F (1050°C) and tested under a stress of 550 psi had a creep rate of 0.000022 pct per hr at 2000 hr, which was still substantially lower than that shown by the specimen annealed at 1382°F (750°C) and stressed at only 400 psi. In contrast to the creep behavior of the platinum and rhodioplatinum specimens, the palladium bars, whether deoxidized with calcium boride or aluminum, were characterized by high first stages of creep. However, after about 1200 hr of test, the creep

Jan 1, 1952

By O. D. Sherby, J. E. Dorn

SEVERAL years ago Zener and Hollomon1 suggested that the flow stress of metals might be related to the temperature and strain rate in accord with the functional equation: s=s(eeh/rt) [1] at the same state. The Zener-Hollomon relation also contains the significant tacit implication that the energy for activation, AH, is substantially independent of the state of the material. The utility of this method for correlating creep data with tensile data was illustrated in a preliminary reportL on the effect of alloying additions on the secondary creep rates of binary a solid solutions in aluminum as shown for Al-Mg alloys in Fig. 1. Not only do the secondary creep and ultimate tensile properties for each alloy correlate well on this plot but in addition the activation energy, AH = 17,900 R cal per mol. is identical for the various compositions. Similar types of curves correlating the secondary creep rates at 422°K (300°F) with ultimate tensile data were also obtained for a series of dilute a solid solutions of copper, germanium, zinc, and silver in aluminum. In all cases the activation energy for creep, AH, was found to be equal to about 17,900 R cal per mol independent of the type of alloying element or its concentration. The coincidence of the activation energies for these alloys is probably due to the fact that the activation energies for creep are rather insensitive to small composition changes. Alloying additions, as shown in Fig. 1, however, increased the stress necessary to obtain equivalent values of ee?H/RT. Thus the stress level in curves of the type represented by Fig. 1 gives the relative creep strengths of the various alloys. Inasmuch as these correlations between creep and tensile data were obtained for creep tests at 422°K, it appeared advisable to ascertain the range of secondary creep rates and temperatures over which correlations between creep and tensile data could be made on the basis of the Zener-Hollomon relationship. If, for example, the activation energy changes with different ranges of creep temperatures, the utility of the proposed analysis would be severely weakened. But if AH is constant not only for all dilute alloys, but also over wide ranges of temperature and secondary creep rates, confidence will be developed not only in the broad utility of the Zener- Table I. Chemical Analysis&apos;:&apos; and Grain Size of Alloys Mean Grain Residual Impurities. Wt Pct Diam- Alloying Atomic-eter, Element Pct Si Fe Cu Mg Mn mm Aluminum 99.987 0.003 0.003 0.006 0.001 0.25 Magnesium 1.617 0.003 0.004 0.006 0.26 Copper 0.101 0.003 0.003 0.0006 0.001 0.29 Zinc 1.616 0.003 0.005 0.007 0.001 &apos;0.26 * The authors express their appreciation to the Aluminum Company of America for the preparation and chemical analyses of these alloys. Hollomon relationship, but also in its theoretical justification. The following investigation was instituted in order to ascertain the range of validity of the Zener-Hollomon relationship when it is applied to creep and tension data. Materials, Equipment, Technique In view of rather extensive correlatable data now available on the plastic properties of a series of binary a solid solutions of magnesium, copper, germanium, zinc, and silver in aluminum, these alloys were again selected for the present investigation. Previous reports have already covered the tensile properties of these alloys at subatmospheric temperatures," and at elevated temperature," as well as their fatigue properties at atmospheric temperatures, nd their creep properties at 422°K.&apos; Resistivity measurementsa as well as precision lattice constant determinations? evealed that these alloys are a solid solutions. After homogenization. 0.100 in. thick sheets were rolled to 0.070 in. in thickness and recrystallized to about the same grain size. Since the details of these treatments were reported previously," they will not be repeated here. In order to appropriately reduce the scope of creep tests in this investigation only the typical alloys shown in Table I were investigated. All creep specimens were machined such that their tensile axes were in the rolling direction of the 0.070 in. sheet stock. The gage section was 2 in. long and 0.500 in. wide with a reduced section of 3.50 in. Creep strains were measured by means of rack and pinion type extensometers with a sensitivity of 0.005 in. per division. Readings were estimated to 1/10 of

Jan 1, 1953

By R. L. Orr, O. D. Sherby, J. E. Dorn

Creep data for pure metals at temperatures above those at which rapid recovery occurs (above about 0.45 the melting temperature) are correlatable by means of the equations and These correlations were applied successfully to data for aluminum, iron, nickel, copper, zinc, platinum, gold, and lead as well as for simple alloys. For a given metal, AH is a constant about equal to the activation energy for self-diffusion. THE early recognition that creep is stimulated by thermal activation prompted numerous investigators&apos;, &apos; to apply the same laws that are valid for the viscous behavior of liquids to analyses of creep data. To a good first approximation these laws are summarized by the equation were i is the creep rate; T, the absolute temperature; R, the gas constant in cal per degree per mol; T, the applied stress; AH, the activation energy in cal per mol; S&apos;, a constant dependent on the entropy of activation, the frequency of activation, and the contribution of each activation to the strain; and B&apos;, a constant dependent upon the size of the flow unit that is activated. Although Eq. 1 describes the flow of viscous fluids very accurately, its application to the creep of solids has been disappointing. Two reasons for the failure of Eq. 1 for creep are now known: First, the decelerating creep rates during primary creep demand that the internal structures of the metals are changing and that these changes might be reflected by changes in one or more of the three creep parameters S&apos;, AH, and B&apos;. Consequently when the conventional methods of evaluating these parameters are employed by comparing the secondary creep rates at a series of temperatures and at a series of stresses, the true effects of temperature and stress are masked because of simultaneous changes in one or more of the creep parameters. Second, as will be illustrated more fully later, three distinctly different types of investigations have shown that the stress term for high temperature creep enters the analysis as and not as t/T. This rather unexpected result points sharply to a significant fundamental difference between the mechanisms for the viscous behavior of fluids and the high temperature creep of metals. More recent extensive investigations",&apos;4 on the creep of high purity aluminum and several of its dilute a solid solutions have shown that where E is the total creep strain; temperature compensated time = for a constant temperature test; t, is the initial stress in a constant load test or the constant true stress in a constant stress test;" and t is the duration of test. • Eq. 2 is valid for either constant stress or constant load tests; but the total creep curve is difierent (i.e. the function, f, is differentl for each test. The fundamental origin of the validity of Eq. 2 is thought to arise from some equivalence of substructures in constant load creep tests at equal values of 0 or E. Both X-ray and metallographic examination revealed that practically identical lattice distortions and subgrain sizes are obtained for the same creep strain under the same load over wide ranges of test temperature.&apos; Furthermore the room temperature tensile properties following equal creep strains at different temperatures under constant loads were also identical, illustrating that identical substructures were indeed obtained.&apos; " A second type of correlation was obtained by differentiating Eq. 2 with respect to time, giving For the secondary creep rate, denoted by i,, But Eq. 2 reveals that 0, is a function of a, alone because the secondary creep rate in a given constant load test is always reached at a fixed value E,". Therefore where is the function that was introduced a number of years ago by Zener and Hollomon." Eq. 4 is not only valid for creep but will permit correlations between constant load creep and tensile data where e, is the rate of tensile straining and (TC is the engineering ultimate tensile strength.

Jan 1, 1955

By R. M. N. Pelloux

This study of aluminum-copper alloys had two aims: 1) To determine the effect of the amount and distribution of a second phase, CuAl2, on the creep-rupture strength, ductility, and fracture characteristics of the alloys. The work of Gemmell and grant&apos; on single-phase aluminum-copper alloys provided a basis for the present study. 2) To attempt to provide a relation between the microstructure and strength of two-phase alloys deformed at comparatively high temperatures (greater than 0.45 T,). The creep behavior of the two-phase magnesium alloys Mg;Ce and Mg-A1, has been treated by Roberts. 9 In work reported by Sully and Hardy4 and by Underwood, Marsh, and Manning~ on A1-Cu two-phase alloys, and also in a portion of the work of Gemmell and grant,&apos; aging took place during the creep tests. &apos;In the present work, overaged aluminum-copper alloys were subjected to creep deformation at two temperatures, 500 and 700OF. The overaging treatments to produce the precipitate dispersions were performed prior to the test, at temperatures which were the same as the ultimate respective test temperatures. Hence, the present study is concerned with the creep behavior of stable (in contrast to "underaged") two-phase A1-Cu alloys, i.e., alloys in which no depletion of the solid-solution matrix occurred during: creep. MATERIALS AND PROCEDURE A 2 and a 3 pct Cu alloy, supplied by Alcoa, were studied; Table I shows the compositions. The grain size of the specimens, 0.9 to 1.0 mm, was achieved by annealing the machined test bars for 2 hr at 1000°F, followed by furnace-cooling to 900°F, and homogenizing for 4 hr. Both compositions are in the single-phase condition at 900°F. Two types of overaged dispersions, termed "over-aged I" and "overaged 11" were prepared after the grain size and the homogenization anneals. The overaged I alloys were prepared by quenching to room temperature after homogenization, aging at either 500&apos; or 700°F for 72 hr, and air-cooling to room temperature. The overaged I1 alloys were prepared by furnace-cooling after homogenization, to either 500&apos; or 700°F, and holding at the necessary temperature for 72 hr. The times for furnace-cooling from 900°F to 700° and 500°F were, respectively, about 2 and 3.5 hr. The structures of the A1-2 pct Cu alloys are shown in Fig. 1 (overaged I) and in Fig. 2 (overaged 11). A comparison of Figs. 1 and 2 shows that in the overaged I alloys the precipitate particles along the grain boundaries are much more closely spaced and the grain boundaries are straighter than they are in the overaged I1 alloys. Further, the particle size and spacing in the grains are smaller than in the overaged I1 alloys. Microstructures of the A1-3 pct Cu alloys have not been shown; the structures of these alloys differed from those of the A1-2 pct Cu alloys only in that the distribution density of the particles was greater. The creep specimens had a gage section 1 in. long with a diameter of 0.155 to 0.195 in. Before testing, specimens were electropolished in a solution of 100 cc glacial acetic acid, and 30 cc of 60 pct perchloric acid, at 40&apos; to 50°F, at 24 v. The specimens were kept refrigerated prior to testing to avoid precipitation at room temperature. Constant stress creep tests were conducted; the load was applied 75 min after heating of the specimens to test temperature had begun. EXPERIMENTAL RESULTS Creep-Rupture—Fig. 3 shows the log stress-log rupture life plotfor the overaged alloys studied, and also for the underaged A1-2 pct Cu alloy.&apos; At 700°F, the stress-rupture life relationship is nearly independent of both the amount of precipitate (i.e., composition) and of the type of overaging treatment (i.e., precipitate dispersion). At 500°F it is apparent that the increase in rupture life that was ex-

Jan 1, 1960

By H. C. Chang, N. J. Grant, A. R. Chaudhuri

During the creep of coarse-grained polycrystalline magnesium at elevated temperatures, a nonbasal type of slip was found to play an important role in the deformation processes. The nonbasal slip traces were examined metallographically in eight specimens (13 grains) and the observed glide plane located stereographically for each grain. The tests were run at 500&apos; and 700°F at stresses of 148 to 786 psi. Based on these measurements and theoretical calculations, the crystallographic elements for nonbasal slip were determined. WHEN a metal deforms by slip, the conventional appearance of the slip bands on the surface of a specimen is that of straight lines. Such is the behavior of the face-centered-cubic metals, as for example, aluminum. The particular slip system on which slip occurs in these metals is governed by the criterion of resolved shear stress. Work on the body-centered-cubic metals a-iron,&apos;. &apos; molybdenum, and columbium4 has shown that the process of slip is somewhat more complicated than the simple gliding of one close-packed plane of atoms over another in a close-packed direction and that it results in the formation of wavy and irregular slip traces on the surface. The recent reviews and experiments of Vogel and Brick and Chen and Mad-din3 show that, for the body-centered-cubic metals, the resolved shear stresses along the planes and the degree of close-packing do not necessarily determine the slip system. Theirs and other investigations indicate that a complete understanding of the slip process in the body-centered-cubic metals is not yet possible. However, one behavior displayed by all these metals is that the slip direction is invariably one of the close-packed directions <11l>. The wavy nature of the slip-band traces has been explained on the basis of cooperative slip by planes sharing this same slip direction.1&apos;3l&apos; It is generally acknowledged that, when the hexagonal-close-packed metal magnesium (axial ratio c/a = 1.624) deforms by slip at room temperature, it does so by basal slip in the close-packed direction <1I%>. This type of slip has the conventional appearance of straight lines and is governed by the critical shear-stress law. Schmid and coworkers" showed that basal slip was operative in the temperature range of —185" to 300°C (—300" to 572°F). Work by Servi, Norton, and Grant6 has shown that, during the creep of coarse-grained aluminum at high temperatures, new slip systems come into operation. The existence of a new high temperature slip system at temperatures greater than 225 °C (437°F) for magnesium was suggested also by Schmid. This was later indicated by Bakarian and Mathewson&apos; to be slip along the pyramidal planes (10i1) in the close-packed direction <11%>.* They * Mention will be made in the text of pyramidal planes and prismatic planes. The pyramidal planes referred to are the type I. order 1 of the form {10il); and the prismatic planes of the type 1 of the form (10i0). For the sake of brevity, they are termed in the teat as pyramidal (10il: and prismatic (10i0) planes, respectively. Where reference is made to the pyramidal plane type 1, order 2 of the form {10i2), they are referred to as the pyramidal planes of the type (10i2:. observed that this kind of slip resulted in irregular markings on the specimen surface. According to them, a possible cause for the appearance of the markings was due to "limited accessory slip on the pyramidal {10i2) planes." Burke and Hibbard," on deforming a single crystal of magnesium at room temperature, found evidence of slip on a pyramidal plane {1011). They explained it as being due to the effect of grip restraints. It is pertinent to note at this point that, although basal slip has received a fair amount of attention, only two specimens have been investigated, that of Bakarian and Mathewson and of Burke and Hibbard, to establish the elements of nonbasal slip in magnesium. During the study of the deformation mechanisms

Jan 1, 1956