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Microstructural evolution and thermal stability of precipitation-strengthened Cu8Cr4Nb alloy
Materials Science and Engineering, A169 (1993) 167-175
167
Microstructural evolution and thermal stability of precipitation-strengthened Cu-8Cr-4Nb alloy K e n R. A n d e r s o n a n d J o a n n a R. G r o z a
Mechanical, Aeronautical and Materials Engineering Department, University of California at Davis, Davis, CA 95616-5274 (USA) R o b e r t L. D r e s h f i e l d a n d D a v i d Ellis
NASA Lewis Research Center, Cleveland, OH 44135 (USA) CReceived February 11, 1993)
Abstract The microstructural changes that occur during age hardening of a precipitation-strengthened copper alloy with 8 at.% Cr and 4 at.% Nb (Cu-8Cr-4Nb) alloy were studied by TEM technique and correlated to the corresponding mechanical properties at room and elevated temperatures. The thermal stability of Cu-8Cr-4Nb alloy was attributed to the fine distribution of secondary precipitates and their slow coarsening rate upon aging. These secondary precipitates provide strengthening by the Orowan mechanism and restrict grain growth upon aging. The experimental values of mechanical properties are compared to the results anticipated from Orowan and Hall-Petch strengthening models to account for the main strengthening mechanisms of the present alloy.
1. Introduction
-3 _ ~ 8~'DCo~ r -r~, t
9RT
The need for materials that combine high mechanical strength and high thermal conductivity at elevated temperatures has revived interest in precipitationhardened copper-based alloys. Classical precipitationstrengthened copper alloys, such as Cu-Cr, Cu-Zr, or Cu-Cr-Zr, have good elevated temperature properties up to 673-773 K. Above these temperatures, mechanical strength and thermal conductivity are both compromised by dissolution of precipitates into the matrix and excessive coarsening. Therefore, for temperatures higher than 773 K, new and more stable precipitationhardened copper alloys are desired. C u - C r - N b is a new highly engineered precipitationstrengthened high-temperature alloy system that was developed by NASA [1, 2]. The premises to develop this alloy are primarily based on maximizing its thermal stability while maintaining the excellent thermal conductivity of the copper matrix. According to a recent study of particle-stability criteria in heat-resistant copper alloys [3], thermal stability is mainly related to the retardation of the particle-coarsening process. The coarsening kinetics are described by the LifshitzSlyozov-Wagner formula: 0921-5093/93/$6.00
(1)
where i is the average particle radius at time t, i 0 is the average particle radius at the onset of coarsening, 7 is the particle-matrix interface energy, D is the solute diffusivity in the matrix, C0 is the solubility of the particle, and f~ is the atomic volume of the precipitate phase. It is clear from this formula that the alloy is thermally stable (or the coarsening rate is slow) when D, 7 and C0 values are minimized. In the C u - C r - N b alloy case, both Cr and Nb have negligible solubilities in solid copper up to 1100 K [4]. Nb also has a low diffusivity value in solid copper [5]. At the same time, the complete solubility of Cr and Nb in liquid copper enables the use of conventional melting techniques for alloy preparation. More importantly, Cr and Nb form an intermetallic compound, Cr2Nb, that melts congruently in liquid copper at 2006 K and is not soluble in the solid copper. As a high temperature intermetallic, Cr2Nb is characterized by high hardness and good thermal stability. To take advantage of these properties, the alloy composition was selected such that only the stable Cr2Nb is formed with no excess Cr or Nb. This composition also allows retention of the © 1993 - Elsevier Sequoia. All rights reserved
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high thermal conductivity of the copper matrix since virtually no alloying elements will remain in solid solution. Furthermore, all Nb is tied up to the Cr2Nb compound such that no terminal Nb solid solution is present to adversely affect the hydrogen resistance [6]. Since CriNb is highly soluble in liquid copper above 2000 K, the rapid solidification (RS) technique is a suitable processing route for alloy preparation [2]. RS increases the solubility limits of Cr and Nb in pure copper to produce controlled precipitation upon further aging heat treatment. Rapidly quenched materials also exhibit highly homogeneous and refined structures with minimal segregation, which is critical for the subsequent aging process. The present alloy contains 8 at.% Cr and 4 at.% Nb (Cu-8Cr-4Nb) so as to give a theoretical volume fraction of 13.56% Cr2Nb as the strengthening phase for the copper matrix. This alloy was prepared by commercial gas atomization, followed by hot extrusion and aging, and had excellent mechanical strength up to 1025 K, and conductivities between 72% and 82% of that of pure copper over the temperature range of interest. Preliminary results also indicated a good hydrogen resistance for this alloy [7]. The objectives of the present investigation are to characterize the microstructural evolution of the Cu-8Cr-4Nb alloys upon aging at temperatures up to 973 K, and to correlate these microstructures with mechanical properties.
2. Experimental details Argon-atomized Cu-8Cr-4Nb powders were supplied by Special Metals, Inc. to NASA Lewis Research Center. Melting procedures were based on the methods developed at NASA Lewis for similar C u - C r , N b alloys [8]. The chemical composition of the powder was 6.0 wt.% (7.36 at.%) Cr and 5.8 wt.% (3.98 wt.%) Nb with an oxygen level of 640 p.p.m. Although the chromium content is slightly off stoichiometric, the Cr:Nb ratio (65:35) is still adequate to form only the Cr2Nb (fl) phase, according to the Cr-Nb phase diagram. The atomized powders (less than 106 ktm) were consolidated by extrusion (16:1 reduction in area) in mild steel cans at 1133 K. To minimize the coarsening of the Cr2Nb precipitates, the soak time for the cans was limited to 1 h. The extruded specimens were aged for 1, 5, 10, 50 and 100 h at 773 and 973 K. For the purpose of comparison, one specimen was aged 100 h at 1073 K. The extruded bars and selected aged specimens were cut and machined into tensile samples for testing at room and elevated temperatures. Tensile testing was conducted under vacuum in an Instron frame with a strain
rate of 6.7 x 10-3 min- 1. The pressure in the chamber for the elevated temperature tensile tests was less than 1.33 x 1 0 - 4 MPa. The maximum time excursion for heating specimens was less than 1 h. A minimum of two tests was conducted at each temperature. In addition, Vickers microhardness (HV) with 100 g load and Rockwell B hardness measurements were taken. The penetration depth at this load is 7.1 pro, significantly larger than the Bilby layer, such that the surfacehardening effects resulting from prior mechanical polishing are negligible. The 100 g microhardness indentations produce diagonals of approximately 35 pm. The maximum grain size as well as particle size are in the order of 1 pm, such that the microhardness values average out the mechanical strength of the alloy. Usually ten measurements were taken and analyzed using standard statistical procedures. The as-extruded, and extruded and aged samples for electron microscope examination (TEM) were cut perpendicular to the extrusion direction. They were mechanically thinned to about 100 p m followed by electrolytic thinning in a methanol bath containing 20% nitric acid at 14 V and a temperature below 220 K. Some samples were ion beam cleaned after electropolishing in a cold-state Gatan ion mill using 5 keV argon ions (current 0.5 mA). The specimens were investigated in a Philips 400 T microscope operated at 120 kV. Precipitate particles were counted using a Quantimet 900 unit. Counts were made in at least 30 different regions or a total of at least 280 particles per specimen. To minimize the error in counting particles in the bright field, dislocation contrast was kept low, and various diffraction conditions in the same area were used to ensure the clear distinction of particle boundaries.
3. Results 3.1. Precipitate structure
A detailed description of the precipitate structure of extruded and aged Cu-8Cr-4Nb alloy is given elsewhere [9]. Therefore, only a brief outline of the important microstructural features that are relevant to understanding the microstructural changes upon aging will be described here. Typical TEM microstructures of extruded and aged Cu-8Cr-4Nb alloy are shown in Fig. 1. Both large particles (Fig. l(a)), up to about 1 /tm, and small particles in the order of tens to hundreds of nanometers (most particles in Fig. l(b)) are common features observed in extruded and aged specimens. Usually, the large particles are found at grain boundaries (Fig. l(a)) whereas the small precipitates are
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Fig. 1. Typical TEM microstructures of atomized, extruded and aged Cu-SCr-4Nb alloys: (a) aged 10 h at 773 K; (b) aged 1 h at 973 K.
observed both in the interior of the grain (Fig. l(b)) and at grain boundaries (arrows in Fig. l(b)). The volume fraction of small and large precipitates are 2.17 _+0.3% and 10.86 _+ 1.9% respectively. Within the experimental error, the total volume fraction is close to the 12.74% calculated for the alloy chemistry.
3.2. Matrix microstructure Heat treatment after extrusion results in a recrystallized matrix structure with grains containing only few dislocations, and straight boundaries often at equilibrium 120 ° angles (Figs. 2(a) and 2(b)). The grain size ranges from 0.2 to 2.6 /~m, independent of specimen condition, with an average grain size of about 0.9 ktm. No noticeable grain growth has been observed with aging time. Even after aging for 100 h at 1073 K, the grain size remains in the same micrometer range (Fig.
2(c)).
Fig. 2. Matrix structure in Cu-8Cr-4Nb alloys aged: (a) 5 h at 973 K; (b) 50 h at 973 K; (c) 100 h at 1073 K.
3.3. Microstructural changes upon aging A number of changes occur during the aging process in atomized and extruded C u - 8 C r - 4 N b alloy. Few fine particles are found in the as-extruded condition (Fig. 3(a)). The extrusion was performed at sufficiently high temperature (1133 K) to allow the small precipitates to dissolve back into copper solid solution. According to
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Fig. 3. Precipitation sequence of fine Cr2Nb compound in Cu-8Cr-4Nb alloys: (a) as-extruded;(b) 1 h at 973 K; (c) 50 h at 973 K.
Cu-Cr and Cu-Nb phase diagrams, substantial solubility of these alloying elements is seen at temperatures above 1100 K. The cooling rate after extrusion was also sufficiently fast to maintain a supersaturated copper solid solution. After 1 h aging at 973 K, very fine secondary precipitates, tens of nanometers in diameter precipitated out of the supersaturated solid solution (Fig. 3(b)). These fine particles coalesce and form larger particles at longer aging times. The microstructures do not change markedly by aging at the two temperatures used in this research. A typical microstructure of precipitated particles which indicates a bimodal particle distribution is seen in Fig. 3(c). The precipiates are usually of a spherical shape and distributed both inside grains and at grain boundaries (see also Fig. l(b)). As expected, the only change with increased annealing time is in precipitate size. Better evidence for secondary particle precipitation and coarsening is provided by the study of the smallparticle-size distribution as a function of aging time at the two annealing temperatures studied. This size distribution of secondary particles as a function of aging time at 773 and 973 K is presented in Fig. 4. The precipitation of new, fine secondary particles after aging is observed by comparing particle-size distribution in as-extruded and 1 h aged materials. For both aging temperatures (773 and 973K), new, fine particles about 30 nm in size appear after 1 h aging of the as-extruded material. As will be shown later, the nucleation and growth processes for Cr2Nb particles seem to be completed after 1 h aging at 773 K so that only particle coarsening will occur at longer aging times or higher temperatures. The particle-size distribution curves shift towards larger particle sizes with increasing aging time. For instance, all the finest Cr2Nb particles (about 30 nm) that precipitated out after 1 h aging, completely disappear after 100 h aging, thus clearly indicating their growth process. However, even after 100 h aging at 973 K, numerous particles of the order of tens of nanometers still remain, showing the slow coarsening process of Cr2Nb precipitates. Based on the above particle-size distributions, average secondary-particle size was calculated. It is difficult to make a clear distinction between coarse secondary particles and fine primary particles. Therefore, an arbitrary-cut-off value for this differentiation was set at 300 nm. This value will be treated with caution since an overlap of large and fine particles occurs at the right side of the distribution curve. However, since large particles will coarsen at a slower rate than small particles, the error caused by the large particle overlap is about the same for all aging conditions. Similar timedependent particle-size characteristics were obtained regardless of this threshold position within _+ 100 nm. Calculations of the secondary-particle size in the atom-
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Precipitation-strengthened Cu-8Cr-4Nb alloy
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ized powders indicated values between 24 and 76 nm [9]. Based on these values, and if it is assumed that coarsening occurs upon aging, the cut-off value for secondary-particle size of 300 + 100 nm is reasonable. A quantitative variation of average secondary-particle size with aging time at 973 K in selected specimens is given in Table 1. For comparison the average secondary-particle size after 100 h aging at 1073 K is 169+ 17 nm. As expected, this value compares well with the as-extruded particle size since the total thermal excursion of the two specimens may be considered equivalent. The secondary-particle size after 100 h at 1073 K remains in the same order of magnitude as at 973 K, thus emphasizing again the sluggish coarsening process in this temperature range. Calculations of the volume fractions of small particles in aged specimens exhibit a fairly constant value at various aging times thus indicating again that
TABLE 1. Secondary-particle size, spacing and Orowanstrengthening contribution Specimen
Secondary particle size, d (nm)
Particle spacing L* (nm)
A o (MPa)
As-extruded Aged 10 h at 973 K Aged 1 0 0 h a t 973 K
163 + 7 -130 148+7
800 638 727
44.9 54.1 48.6
*L = d(n/6~) 1/~ where ~ is the volume fraction of small panicles.
450
the precipitation process was completed after 1 h aging at both 773 and 973 K. This conclusion is in agreement with previous resistivity measurements by Ellis [2] that revealed that the copper matrix resistivity does not change after 1 h aging at 773 K, or that the precipitation process is virtually completed. The time for the completion of the precipitation process at 973 K is shorter than 1 h. By considering the observed stability of the volume fraction of the dispersoids in C u - 8 C r - 4 N b alloy, the particle-spacing values were estimated for 10 and 100 h specimens aged at 973 K, as shown in Table 1.
3.4. Mechanical properties variation upon aging of Cu-8Cr-4Nb alloys The tensile strength of as-extruded Cu-8Cr-4Nb alloy vs. temperature is shown in Fig. 5. For comparison, the equivalent tensile properties of NARloy-Z are also included in Fig. 5. The Cu-8Cr-4Nb samples have approximately twice the mechanical strength of NARIoy-Z. Above 773 K, the NARIoy-Z softens dramatically, while Cu-8Cr-4Nb retains good strength up to 973 K. The effects of aging on room temperature properties of Cu-8Cr-4Nb alloy are shown in Fig. 6. No substantial decrease in the tensile strength or ductility is observed for samples aged up to 100 h at 973 K. The variation of hardness values of the C u - 8 C r 4Nb alloy with aging time at two aging temperatures is shown in Fig. 7. Similar to tensile testing, the hardness measurements indicate that there is an increase of
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300
400
500
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700
Temperature
(K)
800
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Fig. 5. Ultimate tensile strength and yield strength as a function of temperature for C u - 8 C r - 4 N b and Narloy-Z alloys.
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Precipitation-strengthened Cu-SCr-4Nb alloy
25 24 23 J
22 f
21
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f
J
J
J
J
i J
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Aging Time
(hrs)
Fig. 6. Effect of high temperature exposure at 973 K on room temperature strength and ductilityof Cu-8Cr-4Nb alloy.
180
68
170
160
66
150 /
65
• 64
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130 0 As-Extruded
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Aging Time (hrs)
Fig. 7. Hardness variation with agingtime at 773 and 973 K.
hardness values upon aging but no significant decrease in mechanical strength of Cu-8Cr-4Nb alloy for aging up to 100 h at both 773 and 973 K. At both aging temperatures, the hardness increased within 1 h of aging. This result is in accordance with the above microstrucrural results and previously mentioned resistivity measurements, and again leads to the conclusion that complete precipitation occurred in the first hour of aging. 4. Discussion
4.1. Microstructural evolution of C u - 8 C r - 4 N b alloy upon aging In the rapidly solidified powders, the primary Cr2Nb particles formed by solidification in the liquid grow very fast and in any direction before the matrix structure is solidified. The resultant microstructure is large primary particles, usually at grain boundaries that may
173
have a very irregular morphology. The solubility range in the solid copper matrix is expected to be higher than the equilibrium value because of rapid cooling rates. Therefore, precipitation of secondary CreNb particles occurs during subsequent cooling of the atomized powder. Calculations of the secondary-particle sizes obtained during atomization with cooling rates of 102-103 K s ~have shown they are in the 24 to 76 nm range [9]. While the primary particles may grow very fast from the liquid solution, all secondary particles are very small because of a substantial diffusion restriction in the solid state. Therefore, the rapidly solidified powders have a bimodal Cr2Nb particle distribution with large primary particles and small secondary ones, as a result of solidification and solid-state precipitation processes respectively. The solubilities of Cr and Nb in copper, exhibit a dramatic increase at temperatures above l l00K. Therefore, the extrusion step at 1 133 K leads some of the fine particles in the initial rapidly solidified material to dissolve back into the copper matrix. Similar to the atomized material, secondary Cr2Nb particles precipitate out of this supersaturated solid solution upon subsequent aging process. This precipitation process is shown by comparing the particle-size distribution diagrams for the as-extruded and 1 h aged alloys at both aging temperatures. This clearly indicates the formation of new small secondary precipitates (Fig. 4). It is assumed that precipitation is complete within 1 h aging. This assumption is consistent with previous measurements done by Ellis, which indicated that the resistivity changes related to the precipitation process in the solid solution in Cu-8Cr-4Nb alloy were completed after 1 h aging at 773 K [2]. 4.2. Precipitate coarsening and thermal stability A detailed study of particle-coarsening kinetics in a Cu-Cr-Nb alloy is in progress, and will be reported later. For the present investigation, attention is focused on the small-particle-coarsening process because this will have the major effect on the mechanical behavior, especially after long-term exposure at high temperatures. Additionally, since the coarsening process is driven by the reduction in total surface area, small secondary particles will coarsen faster than the large primary particles. This coarsening process occurs during aging applied after the extrusion step. The presented microstructural investigations and mechanical tests indicate the high degree of thermal stability of the dispersoids in C u - C r - N b alloys up to 973 K. This temperature is an arbitrary value, but it is considered the upper bound of service temperatures for many new applications of high-strength, highconductivity copper alloys. The above results of secondary-particle-coarsening analysis indicate that
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Precipitation-strengthened Cu-8Cr-4Nb alloy
the coarsening rate is very slow at both aging temperatures used in this study. The slight variation of hardness and tensile strength values with time at both aging temperatures confirm the microstructural results of secondary-particle coarsening. As already shown, the slow-coarsening kinetics are consistent with the low diffusivity of Nb, and the low solubilities of Cr and Nb in copper up to 973 K. The observed matrix and precipitate microstructures can be used to explain the good thermal stability of Cu-8Cr-4Nb alloys. Cr2Nb precipitates probably strengthen the alloy by grain-boundary-pinning effects, and by providing efficient Orowan-type dislocation obstacles. Experimental results indicate that grain sizes remain in the micron range for all aging treatments. Even after 100 h at 1073 K, the grain size does not become larger than 2.6/~m. This clearly indicates that the pinning effect resulting from precipitates does not diminish when they grow in the size range encountered in these alloys. Calculations of the maximum grain size according to Hillert-Gladman [10, 11] expression, dcalc= 4 r / 3 f (where r is the particle radius and f is the volume fraction) were made to distinguish further between the pinning effect of small vs. large particles. The calculated grain sizes resulting from small particles (3.5/~m ~
0.84 M G b r 2:r(1 - v)l/2Llnb
where M is the Taylor factor, G is the shear modulus of the copper matrix, b is the matrix Burgers vector, r is the particle radius, v is the Poisson's ration of the matrix. The results of Ao calculations (considering the
Taylor factor equal to three) are given in Table 1 for small particle sizes, after extrusion and aging, for 10 h and 100 h at 973 K. Satisfactory correspondence is found between the calculated (10%) and experimentally observed (about 11%, see Fig. 6) decrease in strength when small particles coarsen from 130 nm at 10 h aging to 148 nm after 100 h aging. As expected, the contribution to strengthening of the large dispersoids (average size 1/~m) in the same specimens is one order of magnitude less than that induced by small particles and, therefore, may be considered as negligible. Given the observed stability of the large dispersoids in Cu-8Cr-4Nb alloy, their minor contribution to strengthening is constant regardless of the aging conditions. This result supports the validity of the initial assumption that large particles will have a negligible contribution to strengthening in atomized C u - C r - N b alloys. A comparison of the calculated strengthening effects indicates that the Orowan strengthening has a smaller contribution than grainboundary strengthening in Cu-8Cr-4Nb alloys at room temperature. The above calculations based on secondary-particle coarsening are in good agreement with mechanical test results for the aging conditions. The yield strength after aging 100 h at 973 K remains at 88.6% of the yield strength value for 10 h at the same temperature, despite microscopically observed coarsening. The particle-strengthening effects, and the larger and constant grain-boundary-strengthening component, support the experimentally observed fact that there is no substantial change in the mechanical strength of Cu-8Cr-4Nb alloy aged at 973 K for 100 h.
5. Conclusions
An investigation was conducted to understand the thermal stability of Cu-8Cr-4Nb alloy upon aging at 773 and 973 K for up to 100 h. Coarsening of the secondary Cr2Nb particles in C u - 8 C r - 4 N b alloys has been observed during aging at 773 and 973 K, but the secondary-particle sizes remain around 150 nm even after 100 h aging. Although secondary-particle coarsening occurs, the particle distribution after longtime annealing is still favorable to maintain the good mechanical properties of this alloy. The main strengthening contributions arise from the Hall-Petch and Orowan mechanisms. No significant strengthening contribution was found for large particles. The small-particle volume fraction of Cu-8Cr-4Nb alloy is sufficient to provide the stabilization of a fine grain size during heating up to 973 K. This way, the grain-boundary-strengthening contribution is maintained after long-time annealing at high
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Precipitation-strengthened Cu-8Cr-4Nb alloy
temperatures, and constitutes the predominant strengthening contribution in C u - 8 C r - 4 N b alloys at room temperature. Both experimental and calculated values of Orowan and grain boundary strengthening caused by small panicles indicate that coarsening of the secondary precipitates results in only a minor decrease in mechanical strength by aging 100 h at 973 K.
Acknowledgments This work was performed under a N A S A Summer Fellowship Program (Materials Division at N A S A Lewis Research Center), which is gratefully acknowledged. T h e valuable discussions and experimental assistance of many N A S A colleagues--Drs. Miner, Dickerson, Farmer and Mackay and D. Hull-are greatly appreciated. T h e authors are grateful to Jeff Gibeling for critical reading of the manuscript.
References 1 D. L. Ellis and R. L. Dreshfield, Preliminary evaluation of a powder metal copper-8Cr-4Nb alloy, Proc. 1992 Conf. on Advanced Earth-to-Orbit Propulsion Technology, May 19-21, 1992, Marshall Space Flight Center, Alabama,
NASA, Washington, DC, 1992.
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2 D. L. Ellis, Precipitation strengthened high strength, high conductivity Cu-Cr-Nb alloys produced by chill block melt spinning, Ph.D. Thesis, Case Western Reserve University, Cleveland, OH, 1989. 3 J.R. Groza and J. C. Gibeling, Principles of particle selection for dispersion strengthened copper, Mater. Sci. Eng., to be published. 4 D. J. Chakrabati and D. E. Laughlin, Bull Alloy Phase Diagrams, 2 (1982) 507; 4 (1984) 59 and 99. 5 D. B. Butrymowicz, J. R. Manning and M. E. Read, Diflusion Rate Data and Mass Transport Phenomena for Copper Systems, Diffusion in Metals Data Center, Metallurgy
Division, Institute for Materials Research, National Bureau of Standards, Washington, DC, 1977. 6 D.G. Ulmer, P. D. Krotz, J. L. Yuen and R. P. Jewett, Cryogenic hydrogen effects in copper matrix microcomposites, Proc. 120th Ann. Meeting of TMS, New Orelans, LA, Februa O' 17-21, 1991.
7 R. L, Dreshfield, NASA Lewis Research Center, personal communication, August 1992. 8 D. L. Ellis, G. M. Michal and N. W. Orth, Scripta Metall., 24 (1990) 885. 9 J. R. Groza, Microstructural features of a new precipitation strengthened Cu-8Cr-4Nb alloy, submitted for publication. 10 M. Hillert, Acta Metall., 13 (1965) 227. 11 T. Gladman, Proc. Roy. Soc. London, A 294 (1966) 298. 12 E. O. Hall, Yield Point Phenomena in Metals and Alloys, Plenum, NY, 1970, p. 38. 13 M. F. Ashby, Oxide dispersion strengthening, AIME Conf. Proc., New York, Met. Soc. AIME, 1966, p. 143.
167
Microstructural evolution and thermal stability of precipitation-strengthened Cu-8Cr-4Nb alloy K e n R. A n d e r s o n a n d J o a n n a R. G r o z a
Mechanical, Aeronautical and Materials Engineering Department, University of California at Davis, Davis, CA 95616-5274 (USA) R o b e r t L. D r e s h f i e l d a n d D a v i d Ellis
NASA Lewis Research Center, Cleveland, OH 44135 (USA) CReceived February 11, 1993)
Abstract The microstructural changes that occur during age hardening of a precipitation-strengthened copper alloy with 8 at.% Cr and 4 at.% Nb (Cu-8Cr-4Nb) alloy were studied by TEM technique and correlated to the corresponding mechanical properties at room and elevated temperatures. The thermal stability of Cu-8Cr-4Nb alloy was attributed to the fine distribution of secondary precipitates and their slow coarsening rate upon aging. These secondary precipitates provide strengthening by the Orowan mechanism and restrict grain growth upon aging. The experimental values of mechanical properties are compared to the results anticipated from Orowan and Hall-Petch strengthening models to account for the main strengthening mechanisms of the present alloy.
1. Introduction
-3 _ ~ 8~'DCo~ r -r~, t
9RT
The need for materials that combine high mechanical strength and high thermal conductivity at elevated temperatures has revived interest in precipitationhardened copper-based alloys. Classical precipitationstrengthened copper alloys, such as Cu-Cr, Cu-Zr, or Cu-Cr-Zr, have good elevated temperature properties up to 673-773 K. Above these temperatures, mechanical strength and thermal conductivity are both compromised by dissolution of precipitates into the matrix and excessive coarsening. Therefore, for temperatures higher than 773 K, new and more stable precipitationhardened copper alloys are desired. C u - C r - N b is a new highly engineered precipitationstrengthened high-temperature alloy system that was developed by NASA [1, 2]. The premises to develop this alloy are primarily based on maximizing its thermal stability while maintaining the excellent thermal conductivity of the copper matrix. According to a recent study of particle-stability criteria in heat-resistant copper alloys [3], thermal stability is mainly related to the retardation of the particle-coarsening process. The coarsening kinetics are described by the LifshitzSlyozov-Wagner formula: 0921-5093/93/$6.00
(1)
where i is the average particle radius at time t, i 0 is the average particle radius at the onset of coarsening, 7 is the particle-matrix interface energy, D is the solute diffusivity in the matrix, C0 is the solubility of the particle, and f~ is the atomic volume of the precipitate phase. It is clear from this formula that the alloy is thermally stable (or the coarsening rate is slow) when D, 7 and C0 values are minimized. In the C u - C r - N b alloy case, both Cr and Nb have negligible solubilities in solid copper up to 1100 K [4]. Nb also has a low diffusivity value in solid copper [5]. At the same time, the complete solubility of Cr and Nb in liquid copper enables the use of conventional melting techniques for alloy preparation. More importantly, Cr and Nb form an intermetallic compound, Cr2Nb, that melts congruently in liquid copper at 2006 K and is not soluble in the solid copper. As a high temperature intermetallic, Cr2Nb is characterized by high hardness and good thermal stability. To take advantage of these properties, the alloy composition was selected such that only the stable Cr2Nb is formed with no excess Cr or Nb. This composition also allows retention of the © 1993 - Elsevier Sequoia. All rights reserved
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high thermal conductivity of the copper matrix since virtually no alloying elements will remain in solid solution. Furthermore, all Nb is tied up to the Cr2Nb compound such that no terminal Nb solid solution is present to adversely affect the hydrogen resistance [6]. Since CriNb is highly soluble in liquid copper above 2000 K, the rapid solidification (RS) technique is a suitable processing route for alloy preparation [2]. RS increases the solubility limits of Cr and Nb in pure copper to produce controlled precipitation upon further aging heat treatment. Rapidly quenched materials also exhibit highly homogeneous and refined structures with minimal segregation, which is critical for the subsequent aging process. The present alloy contains 8 at.% Cr and 4 at.% Nb (Cu-8Cr-4Nb) so as to give a theoretical volume fraction of 13.56% Cr2Nb as the strengthening phase for the copper matrix. This alloy was prepared by commercial gas atomization, followed by hot extrusion and aging, and had excellent mechanical strength up to 1025 K, and conductivities between 72% and 82% of that of pure copper over the temperature range of interest. Preliminary results also indicated a good hydrogen resistance for this alloy [7]. The objectives of the present investigation are to characterize the microstructural evolution of the Cu-8Cr-4Nb alloys upon aging at temperatures up to 973 K, and to correlate these microstructures with mechanical properties.
2. Experimental details Argon-atomized Cu-8Cr-4Nb powders were supplied by Special Metals, Inc. to NASA Lewis Research Center. Melting procedures were based on the methods developed at NASA Lewis for similar C u - C r , N b alloys [8]. The chemical composition of the powder was 6.0 wt.% (7.36 at.%) Cr and 5.8 wt.% (3.98 wt.%) Nb with an oxygen level of 640 p.p.m. Although the chromium content is slightly off stoichiometric, the Cr:Nb ratio (65:35) is still adequate to form only the Cr2Nb (fl) phase, according to the Cr-Nb phase diagram. The atomized powders (less than 106 ktm) were consolidated by extrusion (16:1 reduction in area) in mild steel cans at 1133 K. To minimize the coarsening of the Cr2Nb precipitates, the soak time for the cans was limited to 1 h. The extruded specimens were aged for 1, 5, 10, 50 and 100 h at 773 and 973 K. For the purpose of comparison, one specimen was aged 100 h at 1073 K. The extruded bars and selected aged specimens were cut and machined into tensile samples for testing at room and elevated temperatures. Tensile testing was conducted under vacuum in an Instron frame with a strain
rate of 6.7 x 10-3 min- 1. The pressure in the chamber for the elevated temperature tensile tests was less than 1.33 x 1 0 - 4 MPa. The maximum time excursion for heating specimens was less than 1 h. A minimum of two tests was conducted at each temperature. In addition, Vickers microhardness (HV) with 100 g load and Rockwell B hardness measurements were taken. The penetration depth at this load is 7.1 pro, significantly larger than the Bilby layer, such that the surfacehardening effects resulting from prior mechanical polishing are negligible. The 100 g microhardness indentations produce diagonals of approximately 35 pm. The maximum grain size as well as particle size are in the order of 1 pm, such that the microhardness values average out the mechanical strength of the alloy. Usually ten measurements were taken and analyzed using standard statistical procedures. The as-extruded, and extruded and aged samples for electron microscope examination (TEM) were cut perpendicular to the extrusion direction. They were mechanically thinned to about 100 p m followed by electrolytic thinning in a methanol bath containing 20% nitric acid at 14 V and a temperature below 220 K. Some samples were ion beam cleaned after electropolishing in a cold-state Gatan ion mill using 5 keV argon ions (current 0.5 mA). The specimens were investigated in a Philips 400 T microscope operated at 120 kV. Precipitate particles were counted using a Quantimet 900 unit. Counts were made in at least 30 different regions or a total of at least 280 particles per specimen. To minimize the error in counting particles in the bright field, dislocation contrast was kept low, and various diffraction conditions in the same area were used to ensure the clear distinction of particle boundaries.
3. Results 3.1. Precipitate structure
A detailed description of the precipitate structure of extruded and aged Cu-8Cr-4Nb alloy is given elsewhere [9]. Therefore, only a brief outline of the important microstructural features that are relevant to understanding the microstructural changes upon aging will be described here. Typical TEM microstructures of extruded and aged Cu-8Cr-4Nb alloy are shown in Fig. 1. Both large particles (Fig. l(a)), up to about 1 /tm, and small particles in the order of tens to hundreds of nanometers (most particles in Fig. l(b)) are common features observed in extruded and aged specimens. Usually, the large particles are found at grain boundaries (Fig. l(a)) whereas the small precipitates are
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Fig. 1. Typical TEM microstructures of atomized, extruded and aged Cu-SCr-4Nb alloys: (a) aged 10 h at 773 K; (b) aged 1 h at 973 K.
observed both in the interior of the grain (Fig. l(b)) and at grain boundaries (arrows in Fig. l(b)). The volume fraction of small and large precipitates are 2.17 _+0.3% and 10.86 _+ 1.9% respectively. Within the experimental error, the total volume fraction is close to the 12.74% calculated for the alloy chemistry.
3.2. Matrix microstructure Heat treatment after extrusion results in a recrystallized matrix structure with grains containing only few dislocations, and straight boundaries often at equilibrium 120 ° angles (Figs. 2(a) and 2(b)). The grain size ranges from 0.2 to 2.6 /~m, independent of specimen condition, with an average grain size of about 0.9 ktm. No noticeable grain growth has been observed with aging time. Even after aging for 100 h at 1073 K, the grain size remains in the same micrometer range (Fig.
2(c)).
Fig. 2. Matrix structure in Cu-8Cr-4Nb alloys aged: (a) 5 h at 973 K; (b) 50 h at 973 K; (c) 100 h at 1073 K.
3.3. Microstructural changes upon aging A number of changes occur during the aging process in atomized and extruded C u - 8 C r - 4 N b alloy. Few fine particles are found in the as-extruded condition (Fig. 3(a)). The extrusion was performed at sufficiently high temperature (1133 K) to allow the small precipitates to dissolve back into copper solid solution. According to
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Precipitation-strengthened Cu-8Cr-4Nb alloy
Fig. 3. Precipitation sequence of fine Cr2Nb compound in Cu-8Cr-4Nb alloys: (a) as-extruded;(b) 1 h at 973 K; (c) 50 h at 973 K.
Cu-Cr and Cu-Nb phase diagrams, substantial solubility of these alloying elements is seen at temperatures above 1100 K. The cooling rate after extrusion was also sufficiently fast to maintain a supersaturated copper solid solution. After 1 h aging at 973 K, very fine secondary precipitates, tens of nanometers in diameter precipitated out of the supersaturated solid solution (Fig. 3(b)). These fine particles coalesce and form larger particles at longer aging times. The microstructures do not change markedly by aging at the two temperatures used in this research. A typical microstructure of precipitated particles which indicates a bimodal particle distribution is seen in Fig. 3(c). The precipiates are usually of a spherical shape and distributed both inside grains and at grain boundaries (see also Fig. l(b)). As expected, the only change with increased annealing time is in precipitate size. Better evidence for secondary particle precipitation and coarsening is provided by the study of the smallparticle-size distribution as a function of aging time at the two annealing temperatures studied. This size distribution of secondary particles as a function of aging time at 773 and 973 K is presented in Fig. 4. The precipitation of new, fine secondary particles after aging is observed by comparing particle-size distribution in as-extruded and 1 h aged materials. For both aging temperatures (773 and 973K), new, fine particles about 30 nm in size appear after 1 h aging of the as-extruded material. As will be shown later, the nucleation and growth processes for Cr2Nb particles seem to be completed after 1 h aging at 773 K so that only particle coarsening will occur at longer aging times or higher temperatures. The particle-size distribution curves shift towards larger particle sizes with increasing aging time. For instance, all the finest Cr2Nb particles (about 30 nm) that precipitated out after 1 h aging, completely disappear after 100 h aging, thus clearly indicating their growth process. However, even after 100 h aging at 973 K, numerous particles of the order of tens of nanometers still remain, showing the slow coarsening process of Cr2Nb precipitates. Based on the above particle-size distributions, average secondary-particle size was calculated. It is difficult to make a clear distinction between coarse secondary particles and fine primary particles. Therefore, an arbitrary-cut-off value for this differentiation was set at 300 nm. This value will be treated with caution since an overlap of large and fine particles occurs at the right side of the distribution curve. However, since large particles will coarsen at a slower rate than small particles, the error caused by the large particle overlap is about the same for all aging conditions. Similar timedependent particle-size characteristics were obtained regardless of this threshold position within _+ 100 nm. Calculations of the secondary-particle size in the atom-
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Precipitation-strengthened Cu-8Cr-4Nb alloy
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ized powders indicated values between 24 and 76 nm [9]. Based on these values, and if it is assumed that coarsening occurs upon aging, the cut-off value for secondary-particle size of 300 + 100 nm is reasonable. A quantitative variation of average secondary-particle size with aging time at 973 K in selected specimens is given in Table 1. For comparison the average secondary-particle size after 100 h aging at 1073 K is 169+ 17 nm. As expected, this value compares well with the as-extruded particle size since the total thermal excursion of the two specimens may be considered equivalent. The secondary-particle size after 100 h at 1073 K remains in the same order of magnitude as at 973 K, thus emphasizing again the sluggish coarsening process in this temperature range. Calculations of the volume fractions of small particles in aged specimens exhibit a fairly constant value at various aging times thus indicating again that
TABLE 1. Secondary-particle size, spacing and Orowanstrengthening contribution Specimen
Secondary particle size, d (nm)
Particle spacing L* (nm)
A o (MPa)
As-extruded Aged 10 h at 973 K Aged 1 0 0 h a t 973 K
163 + 7 -130 148+7
800 638 727
44.9 54.1 48.6
*L = d(n/6~) 1/~ where ~ is the volume fraction of small panicles.
450
the precipitation process was completed after 1 h aging at both 773 and 973 K. This conclusion is in agreement with previous resistivity measurements by Ellis [2] that revealed that the copper matrix resistivity does not change after 1 h aging at 773 K, or that the precipitation process is virtually completed. The time for the completion of the precipitation process at 973 K is shorter than 1 h. By considering the observed stability of the volume fraction of the dispersoids in C u - 8 C r - 4 N b alloy, the particle-spacing values were estimated for 10 and 100 h specimens aged at 973 K, as shown in Table 1.
3.4. Mechanical properties variation upon aging of Cu-8Cr-4Nb alloys The tensile strength of as-extruded Cu-8Cr-4Nb alloy vs. temperature is shown in Fig. 5. For comparison, the equivalent tensile properties of NARloy-Z are also included in Fig. 5. The Cu-8Cr-4Nb samples have approximately twice the mechanical strength of NARIoy-Z. Above 773 K, the NARIoy-Z softens dramatically, while Cu-8Cr-4Nb retains good strength up to 973 K. The effects of aging on room temperature properties of Cu-8Cr-4Nb alloy are shown in Fig. 6. No substantial decrease in the tensile strength or ductility is observed for samples aged up to 100 h at 973 K. The variation of hardness values of the C u - 8 C r 4Nb alloy with aging time at two aging temperatures is shown in Fig. 7. Similar to tensile testing, the hardness measurements indicate that there is an increase of
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Precipitation-strengthened Cu-SCr-4Nb alloy
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hardness values upon aging but no significant decrease in mechanical strength of Cu-8Cr-4Nb alloy for aging up to 100 h at both 773 and 973 K. At both aging temperatures, the hardness increased within 1 h of aging. This result is in accordance with the above microstrucrural results and previously mentioned resistivity measurements, and again leads to the conclusion that complete precipitation occurred in the first hour of aging. 4. Discussion
4.1. Microstructural evolution of C u - 8 C r - 4 N b alloy upon aging In the rapidly solidified powders, the primary Cr2Nb particles formed by solidification in the liquid grow very fast and in any direction before the matrix structure is solidified. The resultant microstructure is large primary particles, usually at grain boundaries that may
173
have a very irregular morphology. The solubility range in the solid copper matrix is expected to be higher than the equilibrium value because of rapid cooling rates. Therefore, precipitation of secondary CreNb particles occurs during subsequent cooling of the atomized powder. Calculations of the secondary-particle sizes obtained during atomization with cooling rates of 102-103 K s ~have shown they are in the 24 to 76 nm range [9]. While the primary particles may grow very fast from the liquid solution, all secondary particles are very small because of a substantial diffusion restriction in the solid state. Therefore, the rapidly solidified powders have a bimodal Cr2Nb particle distribution with large primary particles and small secondary ones, as a result of solidification and solid-state precipitation processes respectively. The solubilities of Cr and Nb in copper, exhibit a dramatic increase at temperatures above l l00K. Therefore, the extrusion step at 1 133 K leads some of the fine particles in the initial rapidly solidified material to dissolve back into the copper matrix. Similar to the atomized material, secondary Cr2Nb particles precipitate out of this supersaturated solid solution upon subsequent aging process. This precipitation process is shown by comparing the particle-size distribution diagrams for the as-extruded and 1 h aged alloys at both aging temperatures. This clearly indicates the formation of new small secondary precipitates (Fig. 4). It is assumed that precipitation is complete within 1 h aging. This assumption is consistent with previous measurements done by Ellis, which indicated that the resistivity changes related to the precipitation process in the solid solution in Cu-8Cr-4Nb alloy were completed after 1 h aging at 773 K [2]. 4.2. Precipitate coarsening and thermal stability A detailed study of particle-coarsening kinetics in a Cu-Cr-Nb alloy is in progress, and will be reported later. For the present investigation, attention is focused on the small-particle-coarsening process because this will have the major effect on the mechanical behavior, especially after long-term exposure at high temperatures. Additionally, since the coarsening process is driven by the reduction in total surface area, small secondary particles will coarsen faster than the large primary particles. This coarsening process occurs during aging applied after the extrusion step. The presented microstructural investigations and mechanical tests indicate the high degree of thermal stability of the dispersoids in C u - C r - N b alloys up to 973 K. This temperature is an arbitrary value, but it is considered the upper bound of service temperatures for many new applications of high-strength, highconductivity copper alloys. The above results of secondary-particle-coarsening analysis indicate that
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Precipitation-strengthened Cu-8Cr-4Nb alloy
the coarsening rate is very slow at both aging temperatures used in this study. The slight variation of hardness and tensile strength values with time at both aging temperatures confirm the microstructural results of secondary-particle coarsening. As already shown, the slow-coarsening kinetics are consistent with the low diffusivity of Nb, and the low solubilities of Cr and Nb in copper up to 973 K. The observed matrix and precipitate microstructures can be used to explain the good thermal stability of Cu-8Cr-4Nb alloys. Cr2Nb precipitates probably strengthen the alloy by grain-boundary-pinning effects, and by providing efficient Orowan-type dislocation obstacles. Experimental results indicate that grain sizes remain in the micron range for all aging treatments. Even after 100 h at 1073 K, the grain size does not become larger than 2.6/~m. This clearly indicates that the pinning effect resulting from precipitates does not diminish when they grow in the size range encountered in these alloys. Calculations of the maximum grain size according to Hillert-Gladman [10, 11] expression, dcalc= 4 r / 3 f (where r is the particle radius and f is the volume fraction) were made to distinguish further between the pinning effect of small vs. large particles. The calculated grain sizes resulting from small particles (3.5/~m ~
0.84 M G b r 2:r(1 - v)l/2Llnb
where M is the Taylor factor, G is the shear modulus of the copper matrix, b is the matrix Burgers vector, r is the particle radius, v is the Poisson's ration of the matrix. The results of Ao calculations (considering the
Taylor factor equal to three) are given in Table 1 for small particle sizes, after extrusion and aging, for 10 h and 100 h at 973 K. Satisfactory correspondence is found between the calculated (10%) and experimentally observed (about 11%, see Fig. 6) decrease in strength when small particles coarsen from 130 nm at 10 h aging to 148 nm after 100 h aging. As expected, the contribution to strengthening of the large dispersoids (average size 1/~m) in the same specimens is one order of magnitude less than that induced by small particles and, therefore, may be considered as negligible. Given the observed stability of the large dispersoids in Cu-8Cr-4Nb alloy, their minor contribution to strengthening is constant regardless of the aging conditions. This result supports the validity of the initial assumption that large particles will have a negligible contribution to strengthening in atomized C u - C r - N b alloys. A comparison of the calculated strengthening effects indicates that the Orowan strengthening has a smaller contribution than grainboundary strengthening in Cu-8Cr-4Nb alloys at room temperature. The above calculations based on secondary-particle coarsening are in good agreement with mechanical test results for the aging conditions. The yield strength after aging 100 h at 973 K remains at 88.6% of the yield strength value for 10 h at the same temperature, despite microscopically observed coarsening. The particle-strengthening effects, and the larger and constant grain-boundary-strengthening component, support the experimentally observed fact that there is no substantial change in the mechanical strength of Cu-8Cr-4Nb alloy aged at 973 K for 100 h.
5. Conclusions
An investigation was conducted to understand the thermal stability of Cu-8Cr-4Nb alloy upon aging at 773 and 973 K for up to 100 h. Coarsening of the secondary Cr2Nb particles in C u - 8 C r - 4 N b alloys has been observed during aging at 773 and 973 K, but the secondary-particle sizes remain around 150 nm even after 100 h aging. Although secondary-particle coarsening occurs, the particle distribution after longtime annealing is still favorable to maintain the good mechanical properties of this alloy. The main strengthening contributions arise from the Hall-Petch and Orowan mechanisms. No significant strengthening contribution was found for large particles. The small-particle volume fraction of Cu-8Cr-4Nb alloy is sufficient to provide the stabilization of a fine grain size during heating up to 973 K. This way, the grain-boundary-strengthening contribution is maintained after long-time annealing at high
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Precipitation-strengthened Cu-8Cr-4Nb alloy
temperatures, and constitutes the predominant strengthening contribution in C u - 8 C r - 4 N b alloys at room temperature. Both experimental and calculated values of Orowan and grain boundary strengthening caused by small panicles indicate that coarsening of the secondary precipitates results in only a minor decrease in mechanical strength by aging 100 h at 973 K.
Acknowledgments This work was performed under a N A S A Summer Fellowship Program (Materials Division at N A S A Lewis Research Center), which is gratefully acknowledged. T h e valuable discussions and experimental assistance of many N A S A colleagues--Drs. Miner, Dickerson, Farmer and Mackay and D. Hull-are greatly appreciated. T h e authors are grateful to Jeff Gibeling for critical reading of the manuscript.
References 1 D. L. Ellis and R. L. Dreshfield, Preliminary evaluation of a powder metal copper-8Cr-4Nb alloy, Proc. 1992 Conf. on Advanced Earth-to-Orbit Propulsion Technology, May 19-21, 1992, Marshall Space Flight Center, Alabama,
NASA, Washington, DC, 1992.
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2 D. L. Ellis, Precipitation strengthened high strength, high conductivity Cu-Cr-Nb alloys produced by chill block melt spinning, Ph.D. Thesis, Case Western Reserve University, Cleveland, OH, 1989. 3 J.R. Groza and J. C. Gibeling, Principles of particle selection for dispersion strengthened copper, Mater. Sci. Eng., to be published. 4 D. J. Chakrabati and D. E. Laughlin, Bull Alloy Phase Diagrams, 2 (1982) 507; 4 (1984) 59 and 99. 5 D. B. Butrymowicz, J. R. Manning and M. E. Read, Diflusion Rate Data and Mass Transport Phenomena for Copper Systems, Diffusion in Metals Data Center, Metallurgy
Division, Institute for Materials Research, National Bureau of Standards, Washington, DC, 1977. 6 D.G. Ulmer, P. D. Krotz, J. L. Yuen and R. P. Jewett, Cryogenic hydrogen effects in copper matrix microcomposites, Proc. 120th Ann. Meeting of TMS, New Orelans, LA, Februa O' 17-21, 1991.
7 R. L, Dreshfield, NASA Lewis Research Center, personal communication, August 1992. 8 D. L. Ellis, G. M. Michal and N. W. Orth, Scripta Metall., 24 (1990) 885. 9 J. R. Groza, Microstructural features of a new precipitation strengthened Cu-8Cr-4Nb alloy, submitted for publication. 10 M. Hillert, Acta Metall., 13 (1965) 227. 11 T. Gladman, Proc. Roy. Soc. London, A 294 (1966) 298. 12 E. O. Hall, Yield Point Phenomena in Metals and Alloys, Plenum, NY, 1970, p. 38. 13 M. F. Ashby, Oxide dispersion strengthening, AIME Conf. Proc., New York, Met. Soc. AIME, 1966, p. 143.