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Microstructural characterization of dispersion-strengthened Cu–Ti–Al alloys obtained by reaction milling
Materials Science and Engineering A 454–455 (2007) 183–193
Microstructural characterization of dispersion-strengthened Cu–Ti–Al alloys obtained by reaction milling Rodrigo A. Espinoza a , Rodrigo H. Palma a,∗ , Aquiles O. Sep´ulveda a , V´ıctor Fuenzalida b , Guillermo Sol´orzano c , Aldo Craievich d , David J. Smith e , Takeshi Fujita e , Marta L´opez f Departamento de Ingenier´ıa Mec´anica, FCFM, Universidad de Chile, Beauchef 850, 4◦ Piso, Santiago 8370448, Chile b Departamento de F´ısica, FCFM, Universidad de Chile, Av. Blanco Encalada 2008, Santiago 8370415, Chile c DCMM, PUC - Rio, R´ ua Marques de S. Vicente 225, G´avea, Rio de Janeiro, Brazil d Instituto de F´ısica, USP, Travessa R da Rua do Mat˜ ao, no. 187, Cidade Universitaria, 05508-900, S˜ao Paulo, SP, Brazil e Center for Solid State Science, Arizona State University, Tempe, AZ 85287-1704, USA f Departamento de Ingenier´ıa Metal´ urgica, Universidad de Concepci´on, Concepci´on, Chile a
Received 7 June 2006; received in revised form 2 November 2006; accepted 3 November 2006
Abstract The microstructure, electrical conductivity and hot softening resistance of two alloys (G-10 and H-20), projected to attain Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 nominal composition, and prepared by reaction milling and hot extrusion, were studied. The alloys were characterized by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS) and several chemical analysis techniques. The first alloy, G-10, showed the formation of Al2 O3 nanodispersoids and the presence of particles from non-reacted raw materials (graphite, Ti and Al). A second alloy, H-20, was prepared employing different fabrication conditions. This alloy exhibited a homogeneous distribution of Al2 O3 and Ti–Al–Fe nanoparticles, with the microstructure being stable after annealing and hot compression tests. These nanoparticles acted as effective pinning sites for dislocation slip and grain growth. The room-temperature hardness of the H-20 consolidated material (330 HV) was approximately maintained after annealing for 1 h at 1173 K; the electrical conductivity was 60% IACS (International Annealing Copper Standard). © 2006 Elsevier B.V. All rights reserved. Keywords: Copper alloys; Nanoparticles; Creep; Dispersion strengthening; Reaction milling; Mechanical alloying
1. Introduction Alloys exhibiting high mechanical strength together with high electrical and thermal conductivity at elevated temperatures are in increasing demand. Applications for these materials include high-performance switches, welding electrodes, electromotors, heat exchangers, rotating-source neutron targets and rocket nozzles [1]. Due to its high electrical/thermal conductivity, copper is a most promising metal for all of these applications. Moreover, copper has the advantage of a low elastic modulus, which minimizes thermal stresses in actively cooled structures [2]. However, its strength must be increased in order to meet design requirements for high-temperature applications.
∗
Corresponding author. Fax: +56 2 6896057. E-mail address: [email protected] (R.H. Palma).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.11.042
Conventional cold-worked or age-hardened alloys do not properly satisfy high-temperature strength demands, because of the effects of recrystallization, and of precipitate coarsening and dissolution, respectively. The high-temperature resistance of metallic alloys can be increased by adding a small fraction (e.g., between 2 and 5 vol.%) of ceramic dispersoids. In contrast to solid-solution strengthening, the addition of elements to form insoluble particles has little effect on the electrical conductivity. To be most effective, these dispersoids must be thermodynamically stable, homogeneously distributed in the metal matrix and of nanometer size [1]. On the other hand, upon adding ceramic dispersoids, the high-temperature strength of the material is mainly controlled by two mechanisms: dislocation–particle interaction, and grain boundary–particle interaction. Reaction milling is a modern manufacturing process that uses mechanical alloying for the in situ development of nanometer-sized dispersoids in a metal
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matrix [3]. Thus, dispersion-strengthened copper is obtained with improved mechanical behaviour at high temperature [4,5]. In reaction milling, elemental powders are milled under a certain atmosphere and milling media so that one of the metals reacts with C, N, or O in order to form carbides, nitrides or oxides, respectively. During this process, the elemental powders are subjected to high-energy collisions from balls. This process continuously mixes, cold-welds and fractures the powder mixture. The final product after milling generally consists of agglomerated powders that can be either solid solution or a mixture of phases of micron or nanometer sizes, of crystalline or amorphous nature. Because of attrition, such microcrystalline grains also develop high dislocation density. This process also requires that the powders be consolidated after milling by thermomechanical approaches, such as sintering, hot isostatic pressing or hot extrusion [6]. The mechanical behaviour of copper alloys at high temperature has been studied in the past by some of the present authors [6,7]; reaction milling was employed as the manufacturing technique for developing a copper matrix reinforced by ceramic nanodispersoids. The hypothesis behind these studies was that the combination of two nanometric dispersoids (e.g., TiC and Al2 O3 ), which were resistant to different creep mechanisms, would provide greater hot strengthening, in comparison with the case when the same volume fraction of only one type of dispersoid was considered. The aim of the present work is the study of the microstructure, electrical conductivity and hot-softening resistance of two alloys, here after called G-10 and H-20, both with a nominal composition Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 , produced by reaction milling. The mechanical behaviour of G-10 alloy, and a partial microstructure characterization was previously reported [7]. In the present work, a more complete microstructural study of G-10 alloy is described; these results led us to fabricate the second alloy, H-20, using a modified milling procedure. The microstructural characterization of H-20 alloy, and a comparative study of the electrical and mechanical properties of these two materials, are presented. 2. Experimental procedure Two alloys, here called G-10 and H-20, with nominal composition Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 , were produced by reaction milling. The alloys were prepared starting from elemental powders of the pertinent metallic elements (Cu, Ti and Al), with different methods of adding C, which is the element necessary for TiC formation. Powders of dendritic electrolytic Cu (90 wt.% under 40 m) and of graphite (powders of 6 m medium size) were employed in G-10 alloy. On the other hand, spheroidal powders of atomized Cu (90 wt.% under 45 m) were used for H-20 alloy, while the necessary C was mainly provided by the liquid milling media (hexane). The same types of powders of Ti (<45 m) and Al (80 wt.% under 45 m) were employed for both alloys. The added amount in weight of each powder was that calculated to form the alloy nominal composition already stated in volume percent.
The milling process was carried out in a Szegvari-type attritor of stainless steel (100 mm diameter and 1500 ml volume) containing carbon steel balls of 4.8 mm diameter. The milling experiments were performed with a ball-to-powder weight ratio of 10:1 and a rotational speed of 500 rpm, according to the following two-step procedure. The first step consisted of Cu and Al milling for 1 h under air atmosphere using methanol (CH3 OH) as the milling medium (alloy G-10), or under nitrogen atmosphere using hexane (C6 H14 ) as the milling media (alloy H-20), followed by powder drying, and addition of Ti and C elemental powders. In the second step, this mixture was milled for 10 h (alloy G-10) or 20 h (alloy H-20) under a nitrogen atmosphere using hexane as the milling medium. The aggregates obtained were encapsulated at low vacuum and then consolidated by extrusion at 1023 K, using an extrusion ratio of 10:1. Infrared spectroscopy was employed for the determination of C (LECO TCH-600), and also for the O and N content (LECO CS-444LS). Heavier components (Cu, Ti and Fe) were quantified by X-ray fluorescence (XRF) in a Kristalloflex 710H system, equipped with a Mo tube, and a SL 30165CANBERRA detector. The aluminum content was measured using an OBLF QS 750 optical emission spectrometer. X-ray photoelectron spectroscopy (XPS), using a Physical Electronics 1257 system, was employed to characterize both alloys, after a surface cleaning procedure consisting of 20 min of Argon ion erosion at 4 keV. Using multiplex analysis, some of the elements and compounds present in the alloys were identified. Alloy G-10 was also analyzed by X-ray diffraction (XRD) using the powder XRD beam line at the National Synchrotron Light Laboratory (LNLS), Campinas, Brazil. For this purpose, a powder diffractometer and a monochromatic beam with a wavelength of 0.1608 nm were employed. The synchrotron X-ray source was used to detect very weak Bragg diffraction peaks that could not be observed in diffraction patterns recorded by classical X-ray diffractometers. The alloy microstructure was analysed with: (a) a scanning electron microscope (SEM), LEO 1420 VP, using a tungsten filament operated between 20 and 30 kV, and equipped with an EDS (energy dispersive spectroscopy) system, Oxford 7424 INCA; (b) a focused ion beam (FIB) microscope, Strata DB 235M, at FEI company facilities, Eindhoven, Holland. Transmission electron microscopy (TEM) observations were carried out under different conditions in the following facilities: (a) JEOL 2010 (PUC - Rio de Janeiro, Brazil), operated at 200 kV; (b) JEOL 2010F (ASU, Tempe, USA), operated at 200 kV, and equipped with EDS and STEM (scanning transmission electron microscope), and (c) JEOL 2100 F (JEOL company facilities, Tokyo, Japan), operated at 200 kV with EDS analyzer. The TEM samples were prepared in the normal way to obtain 3-mm diameter disks, while the final thinning was made as follows: (a) G-10 alloy, ion milling using Ar gas; (b) H-20 alloy, electropolishing by twin-jet in a TENUPOL device, using a HNO3 :CH3 OH = 30:70 vol.% electrolyte at 243 K, and a voltage between 5 and 10 V. Grain and precipitate sizes were characterized through the corresponding mean equivalent diameters, as determined from TEM images; more than 25 grains were measured for each material.
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To study the microstructure after hot tests at 1123 K, cylindrical specimens of 9.8 mm in length and 6.5 mm in diameter were tested using an Instron TT-DM model machine, under two testing conditions: (a) hot compression tests with a nominal cross-head velocity of 5 × 10−3 cm/min, corresponding to an initial strain rate of 8.3 × 10−5 s−1 ; and (b) constant-stress compression creep tests with a load of 30 MPa. For comparative purposes, the electrical resistivity and hot-softening resistance of both alloys were measured. Electrical conductivity, expressed as % IACS (IACS: International Annealed Copper Standard; 100% IACS being equal to 58.0 m/ mm2 ) was determined from measurements carried out using a double Kelvin bridge with a Digital Low Resistance Ohmmeter (DLRO), AVTM24-1J model; bars of 40 mm long and 6.5 mm diameter were used as samples. On the other hand, to evaluate hot-softening resistance, alloy samples were annealed for 1 h at different temperatures (673, 923, 1023 and 1173 K) and Vickers hardness (100 g) was then measured at room temperature.
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of the mill container and milling balls. Such contamination was also observed in previous studies [6]. 3.1.2. XPS The XPS results are summarized in Table 2. The G-10 alloy showed the presence of Cu, Ti, Al, C and O. Some Fe was also observed, as reported by chemical analysis above (see Table 1). In the G-10 alloy, the C 1s photoelectron peak around 285 eV is from graphite, with the carbide characteristic peak at 282 eV being absent. The O 1s photoelectron energy peak corresponds to metallic oxides, which could be of Cu, Ti or Al. Ti was in the form of oxides (TiO2 and an oxide with lower oxidation state, probably TiO). The Al 2p photoelectron peak position (75 eV) was consistent with alumina (Al2 O3 ), but it was difficult to resolve because it superposes with the Cu 3p photoemission from copper (77 eV), already broadened due to the coexistence of Cu and Cu2 O. These XPS observations of the G-10 alloy confirmed the results from the chemical analysis presented in Section 3.1.1. The TiO2 detected by XPS would be explained by the excess of O (see Table 1), which tended to oxidize the added Ti particles hindering the formation of the less stable TiC. The Cu2 O can be ascribed to oxide on the surface of the milled Cu alloy powders before extrusion, as a product of the O excess in this alloy, as shown by the chemical analysis. The results obtained by XPS (summarized in Table 2) allowed us to confirm the formation of alumina (Al2 O3 ), while Ti showed strong oxidation.
3. Microstructural characterization 3.1. G-10 alloy 3.1.1. Chemical characterization Chemical analysis results are listed in Table 1, where composition values are expressed in weight percent, normalized for the different results. The nominal composition of the target alloy is also included. The measured quantities of Al and Ti in the fabricated G-10 alloy, see Table 1, were considered as appropriate to approximately obtain the target amounts of Al2 O3 and TiC dispersoids (2.5 and 2.5 vol.%). The C content of G-10 alloy was 0.55 wt.%, which was greater than the 0.28 wt.% C added as graphite prior milling. The C excess probably came from the hexane present during the second milling step. The 1.25 wt.% O content of the G-10 alloy was higher than the 0.53 wt.% nominal composition. Finally, as shown in Table 1, small amounts of Fe and Cr contamination were observed, which can be ascribed to steel wear
3.1.3. XRD Different G-10 samples, corresponding to the powders obtained just after milling (i), as-extruded (ii), and as-extruded plus 1-h-annealing at 773 K (iii), 973 K (iv) and 1123 K (v), were studied by powder synchrotron XRD. The weak Bragg peaks detected in the XRD patterns – displayed in Fig. 1 – were assigned to different phases, as shown in that figure and in Table 3. The XRD pattern of a powder just after milling exhibited several Bragg peaks corresponding to metallic Ti, while such peaks were not present in the other patterns (Fig. 1). On the
Table 1 Results of chemical analysis of G-10 and H-20 alloys Alloy
Milling conditions
Nominal composition G-10 H-20
Elements (wt.%)
Step 1
Step 2
Cu
Ti
Al
C
O
N
Fe
Cr
– 1 h Meth-air 1 h Hex-N
– 10 h Hex-N 20 h Hex-N
97.47 99.23 97.09
1.13 0.99 1.13
0.59 0.64 0.57
0.28 0.55 0.47
0.53 1.25 0.72
– 0.001 0.011
– 0.032 0.011
– – 0.009
Table 2 Compounds identified by XPS in the alloys under study Alloy
G-10 H-20
Compounds for elements detected by XPS Cu
Ti
Al
C
O
Cu–Cu2 O Cu–CuO
TiO2 TiC–TiO–Ti2 O3
Al2 O3 Al2 O3
Graphite Carbides and C
Metallic oxides Al2 O3
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Fig. 1. Diffractograms of milled powders and extruded material under different conditions, of G-10 alloy.
Fig. 2. Secondary-electron SEM image of milled powders of G-10 alloy, showing a Ti particle into black circle.
other hand, the pattern of the as-extruded material, and that of the as-extruded plus 1 h annealing at 773 K sample, indicated the presence of titanium carbide (TiC); this phase not being apparent for the same sample after annealing at higher temperatures (973 and 1123 K). Other weak peaks detected in the XRD patterns indicated the presence of titanium and copper-aluminum oxides (TiO2 and CuAl2 O4 ) in all of the studied samples. Several conclusions can be drawn from the results of the XRD study. Metallic Ti peaks detected for the as-extruded material correspond to Ti particles observed by SEM (Section 3.1.4) for milled powders and as-extruded materials. After consolidation, Ti peaks disappeared while TiC peaks appeared, thus demonstrating that metallic Ti reacts with available C to form carbides. The same TiC peaks were also observed for samples after annealing at 773 K, whereas these signals disappeared for materials heat-treated at both 973 and 1123 K. Al atoms are apparent only in a copper-aluminum oxide (CuAl2 O4 ), a phase which was also detected in similar alloys using the same XRD technique [8]. None of the XRD patterns displayed in Fig. 1 exhibited diffraction peaks corresponding to alumina (Al2 O3 ), this phase having been previously detected by XPS (Table 2). A probable explanation for the absence of Al2 O3 diffraction peaks is that aluminium oxide particles either have an amorphous structure or, alternatively, they are composed of very small, nanometer-sized crystals. On the other hand, the apparent absence of Bragg peaks of metallic Al in the XRD patterns displayed in Fig. 1 is probably due to the effect of strong overlapping with Ti diffraction peaks.
3.1.4. SEM 3.1.4.1. Powders after milling. The G-10 alloy powders after milling had a predominantly flaky shape [7], indicating that even though the powders were heavily deformed, they did not extensively cold-weld (Fig. 2). Depending on their size, two types of flaky powders were observed, with the smaller ones being present in larger quantities. EDS analysis revealed that the smaller powders in Fig. 2, were Cu alloy particles with a homogeneous distribution of Ti, Al, C and O. On the other hand, the bigger flaky powders, such as the one marked within a circle, were Ti-rich, which must correspond to the raw-material Ti powders, showing incomplete mechanical alloying. 3.1.4.2. Consolidated G-10 alloy. The G-10 alloy powders were extruded at 1023 K for consolidation. This material, as presented below, was studied under two conditions: (a) as-extruded, and (b) after further annealing at 1123 K for 1 h. As visible in Fig. 3, the extruded material showed Cu alloy powders delineated by surface oxides and pores with sizes
Table 3 Compounds identified by XRD in G-10 alloy for different conditions Condition
Observed phases
(i) Powders after milling (ii) As-extruded (iii) As-extruded and 1 h-annealed at 773 K (iv) As-extruded and 1 h-annealed at 973 K (v) As-extruded and 1 h-annealed at 1123 K
Ti/TiO2 /CuAl2 O4 TiC/TiO2 /CuAl2 O4 TiC/TiO2 /CuAl2 O4 TiO2 /CuAl2 O4 TiO2 /CuAl2 O4
Fig. 3. Secondary-electron SEM image of as-extruded G-10 alloy, showing Ti particle inserted in copper alloy.
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showed larger pores, as the only significant difference compared with those pores corresponding to the as-extruded condition.
Fig. 4. FIB image showing channelling contrast by ion beam of G-10 alloy after a 1 h-annealing at 1123 K. Milled Cu grains and pores (rounded dark regions of about 10 nm diameter) are observed. Arrows indicate twins in Cu grains.
smaller than 1 m. The Ti particles mentioned above in the milled powders (Fig. 2) remained inserted in the copper matrix. One of these Ti particles is shown in Fig. 3, marked within an ellipse; EDS analysis revealed high-Cu content (39 wt.% Ti–61 wt.% Cu), probably due to high interdiffusion during hot extrusion. EDS performed on the Cu matrix showed a homogeneous distribution of Ti and Al (about 2.7 and 1.4 wt.%, respectively). Fig. 4 shows a FIB image obtained by channelling contrast, for the sample extruded and annealed at 1123 K for 1 h. Cu grains with sizes between 100 and 200 nm are observed, which is consistent with a mechanically alloyed material. Twins were also observed, as indicated by the arrows in Fig. 4. SEM images
3.1.5. TEM 3.1.5.1. As-extruded G-10 alloy. TEM observations of the asextruded G-10 alloy, showed a microstructure consisting of copper grains with a mean size of 183 nm, as shown in Fig. 5(a). A homogeneous dispersion of nanometer-sized dispersoids, with sizes between 2 and 20 nm, was also observed, as shown in Fig. 5(b). Fig. 6(a) shows the pinning effect of dispersoids on dislocations and on grain boundary displacements after hot extrusion at 1023 K (arrows). Thus, dispersoids retain dislocation slip during hot plastic deformation, and grain growth at high temperature and during consolidation. Nanotwins with a thickness of around 10 nm were also observed inside the copper grains, see Fig. 6(b) and (c). Commonly, Cu forms annealing twins, in relation to its low stacking-fault energy. Nevertheless, it is worthy to note that deformation twins were widely observed in a recent work [9] on polycrystalline Cu, with mean grain sizes varying from microns to nanometers and submitted to equal-channel angular pressing at room temperature and low strain rates. EDS analysis showed that both graphite and Al-rich particles remained inserted in the matrix, presenting a micron size after the 10 h milling. The Al-rich particles visible in Fig. 7 appeared to be strongly oxidized during the milling process as verified by EDS mapping. The presence of these particles, and the Ti ones observed by SEM, showed that the milling time of 10 h was insufficient to properly reduce the particle size for this alloy. 3.1.5.2. Extruded and annealed for 1 h at 1123 K material. TEM observations of the as-extruded and annealed for 1 h at 1123 K G-10 alloy, showed a mean grain size of 267 nm, which was only some higher than that of the as-extruded material. Moreover, it can be observed in Fig. 8 that dispersoids exhibited no noticeable size increase, and the pinning effect of disper-
Fig. 5. TEM images of as-extruded G-10 alloy: (a) bright field (BF) and (b) dark field (DF).
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Fig. 6. TEM images of as-extruded G-10 alloy: (a) BF showing dispersoid pinning effect; (b) and (c) BF and DF, respectively, showing nanotwins.
soids on dislocations slip. In this condition, Al2 O3 precipitates and Ti–Al particles were detected by EDS.
in grain boundaries and within grains, such as the one shown in Fig. 10(b), which was identified as an Al and O particle, coherent with the copper matrix. Possible GP zones similar to those reported in [7] were also observed (see Fig. 11).
3.1.5.3. Extruded and tested in compression at 1123 K. After hot compression at 1123 K, the material showed a larger amount of dispersoids, as compared to the as-extruded alloy, effectively pinning dislocations and grain boundaries (see Fig. 9(a)). Coherent precipitates were also detected and nanotwins were again observed (Fig. 9(b)). Thus, the deformation process did not significantly affect grain or dispersoid sizes, because of the pinning effect of the dispersoids. For the compressed condition, C and Al2 O3 rounded precipitates with sizes smaller than 10 nm were observed and identified by EDS (see Fig. 10(a)). Dispersoids were seen both
3.1.6. Comments on characterization of the G-10 alloy The presence of Ti flakes in the milled powders, as shown by SEM (Fig. 2), and the strong Ti peak in the spectra, were attributed to insufficient milling time to impose adequate size reduction and mixing of the original added Ti powders. Moreover, the TiO2 presence in the milled powders, as detected by XPS and XRD, could be due to surface oxidation of these Ti flakes during milling. On the other hand, TEM observations (Fig. 7) showed that part of the added graphite and Al pow-
Fig. 7. STEM image of as-extruded G-10 alloy, with Al and graphite particles.
Fig. 8. BF TEM image showing pinning effect over dislocation in 1 h-annealed G-10 alloy.
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Fig. 9. TEM images of G-10 alloy after compression at 1123 K: (a) BF image of grains, showing precipitates retaining grain growth and (b) DF image of twins and nanometric dispersoids.
ders were also insufficiently reduced in size during milling, since these elements remained as micrometric particles in the Cu matrix. Finally, the chemical analysis revealed excess C in the milled agglomerates, relative to the C added as graphite, which suggested that some C was incorporated from the hexane employed as the liquid milling media. With the above in mind, it was decided to prepare a second Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 alloy, employing a modified milling procedure to obtain a proper structure with final composition closer to the nominal value. Thus, the H-20 alloy was prepared with two procedure modifications: (a) graphite powders were not added, as C coming from the hexane would be enough for TiC formation, and (b) the time duration of the second milling step was increased from 10 to 20 h. The 20 h milling time is the same time employed for the preparation of the alloy with just one type of dispersoid reported in Ref. [7].
The XRD characterization of the milled G-10 agglomerates indicated that the expected precipitates (TiC and Al2 O3 ) were not present. Thus, Ti and Al were present as elemental particles and the rest, possibly, mainly in solid solution. Subsequently, after consolidation by extrusion at 1023 K, different compounds were detected by XRD and XPS, showing that precipitation processes had occurred. The O excess detected previously (Table 1) was also avoided by modifying the first step, changing the milling atmosphere for N and using hexane as the milling medium. 3.2. H-20 alloy 3.2.1. Chemical characterization The results of the chemical characterization for the H-20 alloy (presented in Table 1), showed that the C content obtained after
Fig. 10. Images of G-10 alloy after compression at 1123 K: (a) STEM, precipitated particles of alumina and of carbon and (b) HRTEM, image of a precipitate rich in Al and O.
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Fig. 12. Secondary-electron SEM image of H-20 powder alloy.
detected by EDS, a homogeneous distribution of Ti, Al, C and O. Low Fe contamination was also detected. Fig. 11. HRTEM image of as-extruded G-10 alloy and compressed at 1123 K, with possible GP zones.
modification of the preparation procedure was 0.47%, which is sufficient to obtain the expected 2.5 vol.% of TiC dispersoids. An O content of 0.72 wt.% was established, which was closer to the expected value. 3.2.2. XPS The presence of Cu, Ti, Al, C, O and Fe was observed in the H-20 alloy (Table 2). As for the fabrication of this alloy, graphite was not incorporated to the powder mix prior to milling, the C observed by XPS in the H-20 alloy must correspond to C preferentially incorporated from the milling media (hexane) during the 20 h step. In this case, the identification of carbide is unambiguous from the C 1s peak at 282 eV. The compounds identified in the H-20 alloy indicated that C incorporated into the Cu matrix was in the form of C-precipitated particles and of compounds with alloying elements. The Ti 2p photoelectron peak does not correspond to TiO2 , but to oxides with a lower oxidation state, such as TiO or Ti2 O3 , which cannot be assigned with confidence. Ti could form carbides (TiC) and oxides (Ti2 O3 or TiO). However, the low Ti concentration did not permit a conclusive result to be obtained by this technique. TiC precipitates were previously reported by this group in similar alloys [7], and also detected by EDS in TEM observations. The identification of Al2 O3 has the same limitations as indicated above for the G-10 alloy. 3.2.3. SEM 3.2.3.1. Powder after milling. Cu alloy powders with a rounded morphology were obtained in the H-20 alloy after milling, as seen in Fig. 12. Since these powders were also larger than those of the G-10 alloy, and no Ti-rich powders were observed here, it can be inferred that extensive cold welding of powders was obtained in the H-20 alloy. Moreover, these powders showed, as
3.2.3.2. Consolidated H-20 alloy. The as-extruded H-20 alloy showed a uniform microstructure, and a homogenous distribution of the alloying elements (Ti and Al). In this alloy, an effective size reduction of the added alloying powders, including the disappearance of the Ti powders as such, was observed after milling for 20 h. These observations justified the use of a longer milling time for this material. The lower O content in the H-20 alloy (see Table 1), which could be possibly due to a smaller amount of surface oxide of the Cu alloy powders before extrusion, should have allowed better atomic diffusion between powders during the consolidation densification process. Indeed, the absence of oxides marking the interfaces between Cu powders is an indication of the relatively small amount of initial Cu surface oxide. The H-20 alloy, similar to the G-10 sample, exhibited no significant structural differences compared with that of the asextruded condition after a 1 h annealing at 1123 K. 3.2.4. TEM 3.2.4.1. As-extruded H-20 alloy. Fig. 13(a) shows the microstructure of the as-extruded material, which exhibited high-angle grain boundaries such as the one shown in the diffraction pattern insert, obtained from the central region of that figure. Nanodispersoids were observed all around the material, both at grain boundaries as well as inside the grains; they seem to be very effective as pinning sites. Mean grain and dispersoid sizes were 140 and 6.1 nm, respectively. EDS mapping revealed that Ti and Al2 O3 precipitates were dispersed into the matrix, but TiC particles were not observed at all. In this way, precipitates marked as 1, 2 and 3 in Fig. 13(b) corresponded to Al2 O3 and a small content of Ti and Fe. The latter would come from wear of the milling recipient and balls during the process. The point marked as 4 corresponded to the copper matrix, which showed little Al and Fe content, probably as solid solution. Precipitates of Ti, and of Ti with small amounts of Al and Fe, were also identified.
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Fig. 13. BF TEM images of as-extruded H-20 alloy: (a) general image showing grains, with diffraction pattern inserted and (b) nanometer sized particles analyzed by EDS.
3.2.4.2. Creep tested material. The TEM observations of the H-20 alloy, in the condition where the as-extruded material was tested in creep at 1123 K with a load of 30 MPa, for about 5 h to get a real deformation of 0.12, showed no important differences of microstructure compared with that of the abovedescribed as-extruded condition. As seen in Fig. 14(a), there was no significant growth of mean grain size (here measured as 156 nm). High-angle grain boundaries were observed, while the dispersoids presented a mean size of 6.8 nm. EDS analysis of dispersoids showed similar composition as the as-extruded condition. Thus, the dispersoids marked in Fig. 14(b) were composed of Ti, Al, Fe and O. An EDS analysis of the copper matrix, far from any dispersoid, showed very little content of alloyed elements.
Comparison of the microstructure of the tested material with that of the as-extruded one, showed that the H-20 alloy was stable even under the longer creep test at 1123 K. The grain growth was avoided by the pinning effect of the dispersoids during the deformation process at high temperature. 4. Electrical conductivity and hot softening resistance of alloys 4.1. Electrical conductivity The %IACS values of G-10 and H-20 alloys under specific conditions are presented in Table 4. It can be observed that, the H-20 alloy for the as-extruded condition, exhibits lower con-
Fig. 14. BF TEM images of H-20 alloy after creep testing at 1123 K: (a) general image showing grains with diffraction pattern inserted and (b) dispersoids analyzed by EDS.
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Table 4 Electrical conductivity values for the alloys under study and for two reference commercial alloys Alloy
%IACS
Cu–2.5 vol.% Al2 O3 –2.5 vol.% TiC G-10 H-20 H-20 (with thermal treatment)
77 59 62
Cu–1.1 vol.% Al2 O3 GlidCop® Al-25
87
Cu–2.6 vol.% Al2 O3 GlidCop® Al-60
78
ductivity than the G-10 material. Note that the H-20 alloy was submitted to a longer milling time, as compared with the G10 sample, which is consistent with a much-developed milled microstructure, as shown in previous sections. Thus, the lower electrical conductivity of the H-20 alloy, should be related to a higher density of crystal defects, such as dislocations and solid-solution atoms, induced by milling. Moreover, the conductivity of H-20 alloy increased when the material was heat-treated (20 min, 773 K). This increase can be mainly ascribed to a reduction in the alloying element content in solid solution due, for example, to secondary precipitation. The %IACS values of two dispersion-strengthened commercial alloys, GlidCop® Al-25 and Al-60 [10], are given in Table 4, as reference; these alloys are fabricated by internal oxidation of atomized powders, which produces a dispersion of nanometric alumina particles in a copper matrix. For the two GlidCop© alloys it is observed that, as expected, when the particle content increases (from 1.1 to 2.6 vol.%) the electrical conductivity decreases (from 87 to 78% IACS). Compared with the behavior of these commercial alloys, see Table 4, the H-20 alloy, which effectively exhibited ceramic nanodispersoids, follows a rough linear tendency: for higher particle content (5 vol.%) a lower conductivity (about 60% IACS) is obtained. 4.2. Hot-softening resistance The microhardness of G-10 and H-20 alloys after 1 hannealing at different temperatures is presented in Fig. 15. For the initial condition (i.e., without annealing), H-20 alloy shows a higher hardness than G-10 alloy, behaviour that could be explained as a consequence of longer milling time applied to the first alloy (20 h against 10 h). This modification to the preparation procedure produced larger deformation in the copper matrix, and thus, higher dislocation density and lower grain size. Also, the raw materials (Ti and Al) suffered larger deformation and size reduction in the H-20 alloy, as shown by SEM and TEM observations, permitting greater precipitation of nanodispersoids. The hardness values of the H-20 alloy are greater than those of the G-10 alloy over all of the annealing temperature range. In relation to the softening resistance behaviour of each alloy, the G-10 alloy presented a lower hardness relative to the ini-
Fig. 15. Room-temperature hardness of G-10 and H-20 alloys, after annealing of 1 h at different temperatures.
tial value after heat treatments at 673 and 923 K; however, the hardness increased after annealing at 1023 and 1173 K, attaining values close to the initial one. On the other side, the H-20 alloy showed a hardness increase after the treatments at 673 and 923 K, which could be explained by a secondary precipitation of new dispersoids. 5. Final comments and conclusions The microstructural characterization of the original G-10 alloy showed a stable microstructure with nanodispersoids, but with the presence of micrometric raw-material particles (Ti, Al and graphite), which did not undergo the necessary size reduction as to be adequately incorporated in the copper powder matrix. This result led us to modify the fabrication process in order to obtain the same nominal composition alloy (Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 ) with improved microstructure and performance. The material thus obtained, H20 alloy, showed stable microstructure with Al2 O3 and Ti–Al–Fe nanodispersoids embedded in the copper matrix. The dispersoids were very efficient in limiting grain growth, as demonstrated by TEM. The employed two-step milling procedure was attempted with the initial aim of consuming the O available in the copper powders and avoiding the oxidation of Ti, while attaining the formation of TiC in the second step. In fact, the presence of high quantity of TiO2 was observed, in a clear impairment of the TiC formation. After 1 h-annealing, the microhardness at room temperature of the H-20 alloy showed higher values than those of the G-10 alloy, over all of the temperature range. This is explained by a higher milling time of the H-20 alloy (20 h against 10 h), which, in turn, produced higher dislocation density and grain size refining. Also, the alloying elements in H-20 alloy were sufficiently incorporated in the Cu matrix as to precipitate as nanoparticles, probably mainly through hot extrusion. Moreover, during heat treatment applied to the consolidated material, a secondary precipitation occurred, as inferred from an increase of electrical conductivity and hardness. It is concluded that the H-20 alloy is a good candidate for engineering applications at high temperature.
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Acknowledgements
References
The authors acknowledge the financial support of project Fondecyt No. 1011024. R. Espinoza acknowledges Conicyt for the scholarship and financial support for his Doctoral Thesis. They thank Prof. A. Z´arate (U. Cat´olica del Norte) and PhD student D. D´ıaz (U. de Chile) for their assistance with XPS analysis. Moreover, they gratefully thank to ECKA Granulate Micromet GmbII for providing part of the Cu powders used in this study, to Laborat´orio Nacional de Luz S´ıncrotron (LNLA, Campinas, Brazil) for its full cooperation in XRD analysis, to JEOL Company Laboratories in Tokyo for TEM support, and to FEI Company in Eindhoven for FIB microscopy support. We acknowledge use of facilities at the John M. Cowley Center for High Resolution Electron Microscopy at Arizona State University, and partial support from NSF Grant DMR030342.
[1] J.R. Groza, J.C. Gibeling, Mater. Sci. Eng. A 171 (1993) 115–125. [2] I. Anzel, A.C. Kneissl, L. Kosec, Z. Metalkd. 90 (1999) 621–635. [3] L. L¨u, M.O. Lai, Mechanical Alloying, Kluwer Academics Publishers, Massachusetts, EEUU, 1998. [4] C. Sauer, T. Webg¨arber, G. Dehm, J. Mayer, W. P¨usche, B. Kieback, Z. Metallkd. 89 (1998) 119–125. [5] R. Palma, A. Sep´ulveda, A. Z´un˜ iga, R. P´erez., Molienda reactiva de aleaciones Cu–Ti–C, Anales de las Jornadas SAM-CONAMET-AAS 2001, Posadas, Argentina, Septiembre (2001) 317–324. [6] R.H. Palma, A.O. Sep´ulveda, Mater. Sci. Forum 416–418 (2003) 162–167. [7] R.H. Palma, A.O. Sep´ulveda, R.G. Espinoza, A.P. Z´un˜ iga, M.J. Di´anez, J.M. Criado, M.J. Sayagu´es, Mater. Sci. Eng. A 384 (2004) 262–269. [8] J. Kaneko, M. Sugamata, O. Aoya, PM World Congress – Mechanical Alloying, 1998, pp. 323–328. [9] C.X. Huang, K. Wang, S.D. Wu, Z.F. Zhang, G.Y. Li, S.X. Li, Acta Mater. 54 (2006) 655–665. [10] GlidCop technical data sheet: http://www.spotweldingconsultants.com/ glidcop.htm.
Microstructural characterization of dispersion-strengthened Cu–Ti–Al alloys obtained by reaction milling Rodrigo A. Espinoza a , Rodrigo H. Palma a,∗ , Aquiles O. Sep´ulveda a , V´ıctor Fuenzalida b , Guillermo Sol´orzano c , Aldo Craievich d , David J. Smith e , Takeshi Fujita e , Marta L´opez f Departamento de Ingenier´ıa Mec´anica, FCFM, Universidad de Chile, Beauchef 850, 4◦ Piso, Santiago 8370448, Chile b Departamento de F´ısica, FCFM, Universidad de Chile, Av. Blanco Encalada 2008, Santiago 8370415, Chile c DCMM, PUC - Rio, R´ ua Marques de S. Vicente 225, G´avea, Rio de Janeiro, Brazil d Instituto de F´ısica, USP, Travessa R da Rua do Mat˜ ao, no. 187, Cidade Universitaria, 05508-900, S˜ao Paulo, SP, Brazil e Center for Solid State Science, Arizona State University, Tempe, AZ 85287-1704, USA f Departamento de Ingenier´ıa Metal´ urgica, Universidad de Concepci´on, Concepci´on, Chile a
Received 7 June 2006; received in revised form 2 November 2006; accepted 3 November 2006
Abstract The microstructure, electrical conductivity and hot softening resistance of two alloys (G-10 and H-20), projected to attain Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 nominal composition, and prepared by reaction milling and hot extrusion, were studied. The alloys were characterized by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS) and several chemical analysis techniques. The first alloy, G-10, showed the formation of Al2 O3 nanodispersoids and the presence of particles from non-reacted raw materials (graphite, Ti and Al). A second alloy, H-20, was prepared employing different fabrication conditions. This alloy exhibited a homogeneous distribution of Al2 O3 and Ti–Al–Fe nanoparticles, with the microstructure being stable after annealing and hot compression tests. These nanoparticles acted as effective pinning sites for dislocation slip and grain growth. The room-temperature hardness of the H-20 consolidated material (330 HV) was approximately maintained after annealing for 1 h at 1173 K; the electrical conductivity was 60% IACS (International Annealing Copper Standard). © 2006 Elsevier B.V. All rights reserved. Keywords: Copper alloys; Nanoparticles; Creep; Dispersion strengthening; Reaction milling; Mechanical alloying
1. Introduction Alloys exhibiting high mechanical strength together with high electrical and thermal conductivity at elevated temperatures are in increasing demand. Applications for these materials include high-performance switches, welding electrodes, electromotors, heat exchangers, rotating-source neutron targets and rocket nozzles [1]. Due to its high electrical/thermal conductivity, copper is a most promising metal for all of these applications. Moreover, copper has the advantage of a low elastic modulus, which minimizes thermal stresses in actively cooled structures [2]. However, its strength must be increased in order to meet design requirements for high-temperature applications.
∗
Corresponding author. Fax: +56 2 6896057. E-mail address: [email protected] (R.H. Palma).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.11.042
Conventional cold-worked or age-hardened alloys do not properly satisfy high-temperature strength demands, because of the effects of recrystallization, and of precipitate coarsening and dissolution, respectively. The high-temperature resistance of metallic alloys can be increased by adding a small fraction (e.g., between 2 and 5 vol.%) of ceramic dispersoids. In contrast to solid-solution strengthening, the addition of elements to form insoluble particles has little effect on the electrical conductivity. To be most effective, these dispersoids must be thermodynamically stable, homogeneously distributed in the metal matrix and of nanometer size [1]. On the other hand, upon adding ceramic dispersoids, the high-temperature strength of the material is mainly controlled by two mechanisms: dislocation–particle interaction, and grain boundary–particle interaction. Reaction milling is a modern manufacturing process that uses mechanical alloying for the in situ development of nanometer-sized dispersoids in a metal
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matrix [3]. Thus, dispersion-strengthened copper is obtained with improved mechanical behaviour at high temperature [4,5]. In reaction milling, elemental powders are milled under a certain atmosphere and milling media so that one of the metals reacts with C, N, or O in order to form carbides, nitrides or oxides, respectively. During this process, the elemental powders are subjected to high-energy collisions from balls. This process continuously mixes, cold-welds and fractures the powder mixture. The final product after milling generally consists of agglomerated powders that can be either solid solution or a mixture of phases of micron or nanometer sizes, of crystalline or amorphous nature. Because of attrition, such microcrystalline grains also develop high dislocation density. This process also requires that the powders be consolidated after milling by thermomechanical approaches, such as sintering, hot isostatic pressing or hot extrusion [6]. The mechanical behaviour of copper alloys at high temperature has been studied in the past by some of the present authors [6,7]; reaction milling was employed as the manufacturing technique for developing a copper matrix reinforced by ceramic nanodispersoids. The hypothesis behind these studies was that the combination of two nanometric dispersoids (e.g., TiC and Al2 O3 ), which were resistant to different creep mechanisms, would provide greater hot strengthening, in comparison with the case when the same volume fraction of only one type of dispersoid was considered. The aim of the present work is the study of the microstructure, electrical conductivity and hot-softening resistance of two alloys, here after called G-10 and H-20, both with a nominal composition Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 , produced by reaction milling. The mechanical behaviour of G-10 alloy, and a partial microstructure characterization was previously reported [7]. In the present work, a more complete microstructural study of G-10 alloy is described; these results led us to fabricate the second alloy, H-20, using a modified milling procedure. The microstructural characterization of H-20 alloy, and a comparative study of the electrical and mechanical properties of these two materials, are presented. 2. Experimental procedure Two alloys, here called G-10 and H-20, with nominal composition Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 , were produced by reaction milling. The alloys were prepared starting from elemental powders of the pertinent metallic elements (Cu, Ti and Al), with different methods of adding C, which is the element necessary for TiC formation. Powders of dendritic electrolytic Cu (90 wt.% under 40 m) and of graphite (powders of 6 m medium size) were employed in G-10 alloy. On the other hand, spheroidal powders of atomized Cu (90 wt.% under 45 m) were used for H-20 alloy, while the necessary C was mainly provided by the liquid milling media (hexane). The same types of powders of Ti (<45 m) and Al (80 wt.% under 45 m) were employed for both alloys. The added amount in weight of each powder was that calculated to form the alloy nominal composition already stated in volume percent.
The milling process was carried out in a Szegvari-type attritor of stainless steel (100 mm diameter and 1500 ml volume) containing carbon steel balls of 4.8 mm diameter. The milling experiments were performed with a ball-to-powder weight ratio of 10:1 and a rotational speed of 500 rpm, according to the following two-step procedure. The first step consisted of Cu and Al milling for 1 h under air atmosphere using methanol (CH3 OH) as the milling medium (alloy G-10), or under nitrogen atmosphere using hexane (C6 H14 ) as the milling media (alloy H-20), followed by powder drying, and addition of Ti and C elemental powders. In the second step, this mixture was milled for 10 h (alloy G-10) or 20 h (alloy H-20) under a nitrogen atmosphere using hexane as the milling medium. The aggregates obtained were encapsulated at low vacuum and then consolidated by extrusion at 1023 K, using an extrusion ratio of 10:1. Infrared spectroscopy was employed for the determination of C (LECO TCH-600), and also for the O and N content (LECO CS-444LS). Heavier components (Cu, Ti and Fe) were quantified by X-ray fluorescence (XRF) in a Kristalloflex 710H system, equipped with a Mo tube, and a SL 30165CANBERRA detector. The aluminum content was measured using an OBLF QS 750 optical emission spectrometer. X-ray photoelectron spectroscopy (XPS), using a Physical Electronics 1257 system, was employed to characterize both alloys, after a surface cleaning procedure consisting of 20 min of Argon ion erosion at 4 keV. Using multiplex analysis, some of the elements and compounds present in the alloys were identified. Alloy G-10 was also analyzed by X-ray diffraction (XRD) using the powder XRD beam line at the National Synchrotron Light Laboratory (LNLS), Campinas, Brazil. For this purpose, a powder diffractometer and a monochromatic beam with a wavelength of 0.1608 nm were employed. The synchrotron X-ray source was used to detect very weak Bragg diffraction peaks that could not be observed in diffraction patterns recorded by classical X-ray diffractometers. The alloy microstructure was analysed with: (a) a scanning electron microscope (SEM), LEO 1420 VP, using a tungsten filament operated between 20 and 30 kV, and equipped with an EDS (energy dispersive spectroscopy) system, Oxford 7424 INCA; (b) a focused ion beam (FIB) microscope, Strata DB 235M, at FEI company facilities, Eindhoven, Holland. Transmission electron microscopy (TEM) observations were carried out under different conditions in the following facilities: (a) JEOL 2010 (PUC - Rio de Janeiro, Brazil), operated at 200 kV; (b) JEOL 2010F (ASU, Tempe, USA), operated at 200 kV, and equipped with EDS and STEM (scanning transmission electron microscope), and (c) JEOL 2100 F (JEOL company facilities, Tokyo, Japan), operated at 200 kV with EDS analyzer. The TEM samples were prepared in the normal way to obtain 3-mm diameter disks, while the final thinning was made as follows: (a) G-10 alloy, ion milling using Ar gas; (b) H-20 alloy, electropolishing by twin-jet in a TENUPOL device, using a HNO3 :CH3 OH = 30:70 vol.% electrolyte at 243 K, and a voltage between 5 and 10 V. Grain and precipitate sizes were characterized through the corresponding mean equivalent diameters, as determined from TEM images; more than 25 grains were measured for each material.
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To study the microstructure after hot tests at 1123 K, cylindrical specimens of 9.8 mm in length and 6.5 mm in diameter were tested using an Instron TT-DM model machine, under two testing conditions: (a) hot compression tests with a nominal cross-head velocity of 5 × 10−3 cm/min, corresponding to an initial strain rate of 8.3 × 10−5 s−1 ; and (b) constant-stress compression creep tests with a load of 30 MPa. For comparative purposes, the electrical resistivity and hot-softening resistance of both alloys were measured. Electrical conductivity, expressed as % IACS (IACS: International Annealed Copper Standard; 100% IACS being equal to 58.0 m/ mm2 ) was determined from measurements carried out using a double Kelvin bridge with a Digital Low Resistance Ohmmeter (DLRO), AVTM24-1J model; bars of 40 mm long and 6.5 mm diameter were used as samples. On the other hand, to evaluate hot-softening resistance, alloy samples were annealed for 1 h at different temperatures (673, 923, 1023 and 1173 K) and Vickers hardness (100 g) was then measured at room temperature.
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of the mill container and milling balls. Such contamination was also observed in previous studies [6]. 3.1.2. XPS The XPS results are summarized in Table 2. The G-10 alloy showed the presence of Cu, Ti, Al, C and O. Some Fe was also observed, as reported by chemical analysis above (see Table 1). In the G-10 alloy, the C 1s photoelectron peak around 285 eV is from graphite, with the carbide characteristic peak at 282 eV being absent. The O 1s photoelectron energy peak corresponds to metallic oxides, which could be of Cu, Ti or Al. Ti was in the form of oxides (TiO2 and an oxide with lower oxidation state, probably TiO). The Al 2p photoelectron peak position (75 eV) was consistent with alumina (Al2 O3 ), but it was difficult to resolve because it superposes with the Cu 3p photoemission from copper (77 eV), already broadened due to the coexistence of Cu and Cu2 O. These XPS observations of the G-10 alloy confirmed the results from the chemical analysis presented in Section 3.1.1. The TiO2 detected by XPS would be explained by the excess of O (see Table 1), which tended to oxidize the added Ti particles hindering the formation of the less stable TiC. The Cu2 O can be ascribed to oxide on the surface of the milled Cu alloy powders before extrusion, as a product of the O excess in this alloy, as shown by the chemical analysis. The results obtained by XPS (summarized in Table 2) allowed us to confirm the formation of alumina (Al2 O3 ), while Ti showed strong oxidation.
3. Microstructural characterization 3.1. G-10 alloy 3.1.1. Chemical characterization Chemical analysis results are listed in Table 1, where composition values are expressed in weight percent, normalized for the different results. The nominal composition of the target alloy is also included. The measured quantities of Al and Ti in the fabricated G-10 alloy, see Table 1, were considered as appropriate to approximately obtain the target amounts of Al2 O3 and TiC dispersoids (2.5 and 2.5 vol.%). The C content of G-10 alloy was 0.55 wt.%, which was greater than the 0.28 wt.% C added as graphite prior milling. The C excess probably came from the hexane present during the second milling step. The 1.25 wt.% O content of the G-10 alloy was higher than the 0.53 wt.% nominal composition. Finally, as shown in Table 1, small amounts of Fe and Cr contamination were observed, which can be ascribed to steel wear
3.1.3. XRD Different G-10 samples, corresponding to the powders obtained just after milling (i), as-extruded (ii), and as-extruded plus 1-h-annealing at 773 K (iii), 973 K (iv) and 1123 K (v), were studied by powder synchrotron XRD. The weak Bragg peaks detected in the XRD patterns – displayed in Fig. 1 – were assigned to different phases, as shown in that figure and in Table 3. The XRD pattern of a powder just after milling exhibited several Bragg peaks corresponding to metallic Ti, while such peaks were not present in the other patterns (Fig. 1). On the
Table 1 Results of chemical analysis of G-10 and H-20 alloys Alloy
Milling conditions
Nominal composition G-10 H-20
Elements (wt.%)
Step 1
Step 2
Cu
Ti
Al
C
O
N
Fe
Cr
– 1 h Meth-air 1 h Hex-N
– 10 h Hex-N 20 h Hex-N
97.47 99.23 97.09
1.13 0.99 1.13
0.59 0.64 0.57
0.28 0.55 0.47
0.53 1.25 0.72
– 0.001 0.011
– 0.032 0.011
– – 0.009
Table 2 Compounds identified by XPS in the alloys under study Alloy
G-10 H-20
Compounds for elements detected by XPS Cu
Ti
Al
C
O
Cu–Cu2 O Cu–CuO
TiO2 TiC–TiO–Ti2 O3
Al2 O3 Al2 O3
Graphite Carbides and C
Metallic oxides Al2 O3
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Fig. 1. Diffractograms of milled powders and extruded material under different conditions, of G-10 alloy.
Fig. 2. Secondary-electron SEM image of milled powders of G-10 alloy, showing a Ti particle into black circle.
other hand, the pattern of the as-extruded material, and that of the as-extruded plus 1 h annealing at 773 K sample, indicated the presence of titanium carbide (TiC); this phase not being apparent for the same sample after annealing at higher temperatures (973 and 1123 K). Other weak peaks detected in the XRD patterns indicated the presence of titanium and copper-aluminum oxides (TiO2 and CuAl2 O4 ) in all of the studied samples. Several conclusions can be drawn from the results of the XRD study. Metallic Ti peaks detected for the as-extruded material correspond to Ti particles observed by SEM (Section 3.1.4) for milled powders and as-extruded materials. After consolidation, Ti peaks disappeared while TiC peaks appeared, thus demonstrating that metallic Ti reacts with available C to form carbides. The same TiC peaks were also observed for samples after annealing at 773 K, whereas these signals disappeared for materials heat-treated at both 973 and 1123 K. Al atoms are apparent only in a copper-aluminum oxide (CuAl2 O4 ), a phase which was also detected in similar alloys using the same XRD technique [8]. None of the XRD patterns displayed in Fig. 1 exhibited diffraction peaks corresponding to alumina (Al2 O3 ), this phase having been previously detected by XPS (Table 2). A probable explanation for the absence of Al2 O3 diffraction peaks is that aluminium oxide particles either have an amorphous structure or, alternatively, they are composed of very small, nanometer-sized crystals. On the other hand, the apparent absence of Bragg peaks of metallic Al in the XRD patterns displayed in Fig. 1 is probably due to the effect of strong overlapping with Ti diffraction peaks.
3.1.4. SEM 3.1.4.1. Powders after milling. The G-10 alloy powders after milling had a predominantly flaky shape [7], indicating that even though the powders were heavily deformed, they did not extensively cold-weld (Fig. 2). Depending on their size, two types of flaky powders were observed, with the smaller ones being present in larger quantities. EDS analysis revealed that the smaller powders in Fig. 2, were Cu alloy particles with a homogeneous distribution of Ti, Al, C and O. On the other hand, the bigger flaky powders, such as the one marked within a circle, were Ti-rich, which must correspond to the raw-material Ti powders, showing incomplete mechanical alloying. 3.1.4.2. Consolidated G-10 alloy. The G-10 alloy powders were extruded at 1023 K for consolidation. This material, as presented below, was studied under two conditions: (a) as-extruded, and (b) after further annealing at 1123 K for 1 h. As visible in Fig. 3, the extruded material showed Cu alloy powders delineated by surface oxides and pores with sizes
Table 3 Compounds identified by XRD in G-10 alloy for different conditions Condition
Observed phases
(i) Powders after milling (ii) As-extruded (iii) As-extruded and 1 h-annealed at 773 K (iv) As-extruded and 1 h-annealed at 973 K (v) As-extruded and 1 h-annealed at 1123 K
Ti/TiO2 /CuAl2 O4 TiC/TiO2 /CuAl2 O4 TiC/TiO2 /CuAl2 O4 TiO2 /CuAl2 O4 TiO2 /CuAl2 O4
Fig. 3. Secondary-electron SEM image of as-extruded G-10 alloy, showing Ti particle inserted in copper alloy.
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showed larger pores, as the only significant difference compared with those pores corresponding to the as-extruded condition.
Fig. 4. FIB image showing channelling contrast by ion beam of G-10 alloy after a 1 h-annealing at 1123 K. Milled Cu grains and pores (rounded dark regions of about 10 nm diameter) are observed. Arrows indicate twins in Cu grains.
smaller than 1 m. The Ti particles mentioned above in the milled powders (Fig. 2) remained inserted in the copper matrix. One of these Ti particles is shown in Fig. 3, marked within an ellipse; EDS analysis revealed high-Cu content (39 wt.% Ti–61 wt.% Cu), probably due to high interdiffusion during hot extrusion. EDS performed on the Cu matrix showed a homogeneous distribution of Ti and Al (about 2.7 and 1.4 wt.%, respectively). Fig. 4 shows a FIB image obtained by channelling contrast, for the sample extruded and annealed at 1123 K for 1 h. Cu grains with sizes between 100 and 200 nm are observed, which is consistent with a mechanically alloyed material. Twins were also observed, as indicated by the arrows in Fig. 4. SEM images
3.1.5. TEM 3.1.5.1. As-extruded G-10 alloy. TEM observations of the asextruded G-10 alloy, showed a microstructure consisting of copper grains with a mean size of 183 nm, as shown in Fig. 5(a). A homogeneous dispersion of nanometer-sized dispersoids, with sizes between 2 and 20 nm, was also observed, as shown in Fig. 5(b). Fig. 6(a) shows the pinning effect of dispersoids on dislocations and on grain boundary displacements after hot extrusion at 1023 K (arrows). Thus, dispersoids retain dislocation slip during hot plastic deformation, and grain growth at high temperature and during consolidation. Nanotwins with a thickness of around 10 nm were also observed inside the copper grains, see Fig. 6(b) and (c). Commonly, Cu forms annealing twins, in relation to its low stacking-fault energy. Nevertheless, it is worthy to note that deformation twins were widely observed in a recent work [9] on polycrystalline Cu, with mean grain sizes varying from microns to nanometers and submitted to equal-channel angular pressing at room temperature and low strain rates. EDS analysis showed that both graphite and Al-rich particles remained inserted in the matrix, presenting a micron size after the 10 h milling. The Al-rich particles visible in Fig. 7 appeared to be strongly oxidized during the milling process as verified by EDS mapping. The presence of these particles, and the Ti ones observed by SEM, showed that the milling time of 10 h was insufficient to properly reduce the particle size for this alloy. 3.1.5.2. Extruded and annealed for 1 h at 1123 K material. TEM observations of the as-extruded and annealed for 1 h at 1123 K G-10 alloy, showed a mean grain size of 267 nm, which was only some higher than that of the as-extruded material. Moreover, it can be observed in Fig. 8 that dispersoids exhibited no noticeable size increase, and the pinning effect of disper-
Fig. 5. TEM images of as-extruded G-10 alloy: (a) bright field (BF) and (b) dark field (DF).
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Fig. 6. TEM images of as-extruded G-10 alloy: (a) BF showing dispersoid pinning effect; (b) and (c) BF and DF, respectively, showing nanotwins.
soids on dislocations slip. In this condition, Al2 O3 precipitates and Ti–Al particles were detected by EDS.
in grain boundaries and within grains, such as the one shown in Fig. 10(b), which was identified as an Al and O particle, coherent with the copper matrix. Possible GP zones similar to those reported in [7] were also observed (see Fig. 11).
3.1.5.3. Extruded and tested in compression at 1123 K. After hot compression at 1123 K, the material showed a larger amount of dispersoids, as compared to the as-extruded alloy, effectively pinning dislocations and grain boundaries (see Fig. 9(a)). Coherent precipitates were also detected and nanotwins were again observed (Fig. 9(b)). Thus, the deformation process did not significantly affect grain or dispersoid sizes, because of the pinning effect of the dispersoids. For the compressed condition, C and Al2 O3 rounded precipitates with sizes smaller than 10 nm were observed and identified by EDS (see Fig. 10(a)). Dispersoids were seen both
3.1.6. Comments on characterization of the G-10 alloy The presence of Ti flakes in the milled powders, as shown by SEM (Fig. 2), and the strong Ti peak in the spectra, were attributed to insufficient milling time to impose adequate size reduction and mixing of the original added Ti powders. Moreover, the TiO2 presence in the milled powders, as detected by XPS and XRD, could be due to surface oxidation of these Ti flakes during milling. On the other hand, TEM observations (Fig. 7) showed that part of the added graphite and Al pow-
Fig. 7. STEM image of as-extruded G-10 alloy, with Al and graphite particles.
Fig. 8. BF TEM image showing pinning effect over dislocation in 1 h-annealed G-10 alloy.
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Fig. 9. TEM images of G-10 alloy after compression at 1123 K: (a) BF image of grains, showing precipitates retaining grain growth and (b) DF image of twins and nanometric dispersoids.
ders were also insufficiently reduced in size during milling, since these elements remained as micrometric particles in the Cu matrix. Finally, the chemical analysis revealed excess C in the milled agglomerates, relative to the C added as graphite, which suggested that some C was incorporated from the hexane employed as the liquid milling media. With the above in mind, it was decided to prepare a second Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 alloy, employing a modified milling procedure to obtain a proper structure with final composition closer to the nominal value. Thus, the H-20 alloy was prepared with two procedure modifications: (a) graphite powders were not added, as C coming from the hexane would be enough for TiC formation, and (b) the time duration of the second milling step was increased from 10 to 20 h. The 20 h milling time is the same time employed for the preparation of the alloy with just one type of dispersoid reported in Ref. [7].
The XRD characterization of the milled G-10 agglomerates indicated that the expected precipitates (TiC and Al2 O3 ) were not present. Thus, Ti and Al were present as elemental particles and the rest, possibly, mainly in solid solution. Subsequently, after consolidation by extrusion at 1023 K, different compounds were detected by XRD and XPS, showing that precipitation processes had occurred. The O excess detected previously (Table 1) was also avoided by modifying the first step, changing the milling atmosphere for N and using hexane as the milling medium. 3.2. H-20 alloy 3.2.1. Chemical characterization The results of the chemical characterization for the H-20 alloy (presented in Table 1), showed that the C content obtained after
Fig. 10. Images of G-10 alloy after compression at 1123 K: (a) STEM, precipitated particles of alumina and of carbon and (b) HRTEM, image of a precipitate rich in Al and O.
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Fig. 12. Secondary-electron SEM image of H-20 powder alloy.
detected by EDS, a homogeneous distribution of Ti, Al, C and O. Low Fe contamination was also detected. Fig. 11. HRTEM image of as-extruded G-10 alloy and compressed at 1123 K, with possible GP zones.
modification of the preparation procedure was 0.47%, which is sufficient to obtain the expected 2.5 vol.% of TiC dispersoids. An O content of 0.72 wt.% was established, which was closer to the expected value. 3.2.2. XPS The presence of Cu, Ti, Al, C, O and Fe was observed in the H-20 alloy (Table 2). As for the fabrication of this alloy, graphite was not incorporated to the powder mix prior to milling, the C observed by XPS in the H-20 alloy must correspond to C preferentially incorporated from the milling media (hexane) during the 20 h step. In this case, the identification of carbide is unambiguous from the C 1s peak at 282 eV. The compounds identified in the H-20 alloy indicated that C incorporated into the Cu matrix was in the form of C-precipitated particles and of compounds with alloying elements. The Ti 2p photoelectron peak does not correspond to TiO2 , but to oxides with a lower oxidation state, such as TiO or Ti2 O3 , which cannot be assigned with confidence. Ti could form carbides (TiC) and oxides (Ti2 O3 or TiO). However, the low Ti concentration did not permit a conclusive result to be obtained by this technique. TiC precipitates were previously reported by this group in similar alloys [7], and also detected by EDS in TEM observations. The identification of Al2 O3 has the same limitations as indicated above for the G-10 alloy. 3.2.3. SEM 3.2.3.1. Powder after milling. Cu alloy powders with a rounded morphology were obtained in the H-20 alloy after milling, as seen in Fig. 12. Since these powders were also larger than those of the G-10 alloy, and no Ti-rich powders were observed here, it can be inferred that extensive cold welding of powders was obtained in the H-20 alloy. Moreover, these powders showed, as
3.2.3.2. Consolidated H-20 alloy. The as-extruded H-20 alloy showed a uniform microstructure, and a homogenous distribution of the alloying elements (Ti and Al). In this alloy, an effective size reduction of the added alloying powders, including the disappearance of the Ti powders as such, was observed after milling for 20 h. These observations justified the use of a longer milling time for this material. The lower O content in the H-20 alloy (see Table 1), which could be possibly due to a smaller amount of surface oxide of the Cu alloy powders before extrusion, should have allowed better atomic diffusion between powders during the consolidation densification process. Indeed, the absence of oxides marking the interfaces between Cu powders is an indication of the relatively small amount of initial Cu surface oxide. The H-20 alloy, similar to the G-10 sample, exhibited no significant structural differences compared with that of the asextruded condition after a 1 h annealing at 1123 K. 3.2.4. TEM 3.2.4.1. As-extruded H-20 alloy. Fig. 13(a) shows the microstructure of the as-extruded material, which exhibited high-angle grain boundaries such as the one shown in the diffraction pattern insert, obtained from the central region of that figure. Nanodispersoids were observed all around the material, both at grain boundaries as well as inside the grains; they seem to be very effective as pinning sites. Mean grain and dispersoid sizes were 140 and 6.1 nm, respectively. EDS mapping revealed that Ti and Al2 O3 precipitates were dispersed into the matrix, but TiC particles were not observed at all. In this way, precipitates marked as 1, 2 and 3 in Fig. 13(b) corresponded to Al2 O3 and a small content of Ti and Fe. The latter would come from wear of the milling recipient and balls during the process. The point marked as 4 corresponded to the copper matrix, which showed little Al and Fe content, probably as solid solution. Precipitates of Ti, and of Ti with small amounts of Al and Fe, were also identified.
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Fig. 13. BF TEM images of as-extruded H-20 alloy: (a) general image showing grains, with diffraction pattern inserted and (b) nanometer sized particles analyzed by EDS.
3.2.4.2. Creep tested material. The TEM observations of the H-20 alloy, in the condition where the as-extruded material was tested in creep at 1123 K with a load of 30 MPa, for about 5 h to get a real deformation of 0.12, showed no important differences of microstructure compared with that of the abovedescribed as-extruded condition. As seen in Fig. 14(a), there was no significant growth of mean grain size (here measured as 156 nm). High-angle grain boundaries were observed, while the dispersoids presented a mean size of 6.8 nm. EDS analysis of dispersoids showed similar composition as the as-extruded condition. Thus, the dispersoids marked in Fig. 14(b) were composed of Ti, Al, Fe and O. An EDS analysis of the copper matrix, far from any dispersoid, showed very little content of alloyed elements.
Comparison of the microstructure of the tested material with that of the as-extruded one, showed that the H-20 alloy was stable even under the longer creep test at 1123 K. The grain growth was avoided by the pinning effect of the dispersoids during the deformation process at high temperature. 4. Electrical conductivity and hot softening resistance of alloys 4.1. Electrical conductivity The %IACS values of G-10 and H-20 alloys under specific conditions are presented in Table 4. It can be observed that, the H-20 alloy for the as-extruded condition, exhibits lower con-
Fig. 14. BF TEM images of H-20 alloy after creep testing at 1123 K: (a) general image showing grains with diffraction pattern inserted and (b) dispersoids analyzed by EDS.
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Table 4 Electrical conductivity values for the alloys under study and for two reference commercial alloys Alloy
%IACS
Cu–2.5 vol.% Al2 O3 –2.5 vol.% TiC G-10 H-20 H-20 (with thermal treatment)
77 59 62
Cu–1.1 vol.% Al2 O3 GlidCop® Al-25
87
Cu–2.6 vol.% Al2 O3 GlidCop® Al-60
78
ductivity than the G-10 material. Note that the H-20 alloy was submitted to a longer milling time, as compared with the G10 sample, which is consistent with a much-developed milled microstructure, as shown in previous sections. Thus, the lower electrical conductivity of the H-20 alloy, should be related to a higher density of crystal defects, such as dislocations and solid-solution atoms, induced by milling. Moreover, the conductivity of H-20 alloy increased when the material was heat-treated (20 min, 773 K). This increase can be mainly ascribed to a reduction in the alloying element content in solid solution due, for example, to secondary precipitation. The %IACS values of two dispersion-strengthened commercial alloys, GlidCop® Al-25 and Al-60 [10], are given in Table 4, as reference; these alloys are fabricated by internal oxidation of atomized powders, which produces a dispersion of nanometric alumina particles in a copper matrix. For the two GlidCop© alloys it is observed that, as expected, when the particle content increases (from 1.1 to 2.6 vol.%) the electrical conductivity decreases (from 87 to 78% IACS). Compared with the behavior of these commercial alloys, see Table 4, the H-20 alloy, which effectively exhibited ceramic nanodispersoids, follows a rough linear tendency: for higher particle content (5 vol.%) a lower conductivity (about 60% IACS) is obtained. 4.2. Hot-softening resistance The microhardness of G-10 and H-20 alloys after 1 hannealing at different temperatures is presented in Fig. 15. For the initial condition (i.e., without annealing), H-20 alloy shows a higher hardness than G-10 alloy, behaviour that could be explained as a consequence of longer milling time applied to the first alloy (20 h against 10 h). This modification to the preparation procedure produced larger deformation in the copper matrix, and thus, higher dislocation density and lower grain size. Also, the raw materials (Ti and Al) suffered larger deformation and size reduction in the H-20 alloy, as shown by SEM and TEM observations, permitting greater precipitation of nanodispersoids. The hardness values of the H-20 alloy are greater than those of the G-10 alloy over all of the annealing temperature range. In relation to the softening resistance behaviour of each alloy, the G-10 alloy presented a lower hardness relative to the ini-
Fig. 15. Room-temperature hardness of G-10 and H-20 alloys, after annealing of 1 h at different temperatures.
tial value after heat treatments at 673 and 923 K; however, the hardness increased after annealing at 1023 and 1173 K, attaining values close to the initial one. On the other side, the H-20 alloy showed a hardness increase after the treatments at 673 and 923 K, which could be explained by a secondary precipitation of new dispersoids. 5. Final comments and conclusions The microstructural characterization of the original G-10 alloy showed a stable microstructure with nanodispersoids, but with the presence of micrometric raw-material particles (Ti, Al and graphite), which did not undergo the necessary size reduction as to be adequately incorporated in the copper powder matrix. This result led us to modify the fabrication process in order to obtain the same nominal composition alloy (Cu–2.5 vol.% TiC–2.5 vol.% Al2 O3 ) with improved microstructure and performance. The material thus obtained, H20 alloy, showed stable microstructure with Al2 O3 and Ti–Al–Fe nanodispersoids embedded in the copper matrix. The dispersoids were very efficient in limiting grain growth, as demonstrated by TEM. The employed two-step milling procedure was attempted with the initial aim of consuming the O available in the copper powders and avoiding the oxidation of Ti, while attaining the formation of TiC in the second step. In fact, the presence of high quantity of TiO2 was observed, in a clear impairment of the TiC formation. After 1 h-annealing, the microhardness at room temperature of the H-20 alloy showed higher values than those of the G-10 alloy, over all of the temperature range. This is explained by a higher milling time of the H-20 alloy (20 h against 10 h), which, in turn, produced higher dislocation density and grain size refining. Also, the alloying elements in H-20 alloy were sufficiently incorporated in the Cu matrix as to precipitate as nanoparticles, probably mainly through hot extrusion. Moreover, during heat treatment applied to the consolidated material, a secondary precipitation occurred, as inferred from an increase of electrical conductivity and hardness. It is concluded that the H-20 alloy is a good candidate for engineering applications at high temperature.
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Acknowledgements
References
The authors acknowledge the financial support of project Fondecyt No. 1011024. R. Espinoza acknowledges Conicyt for the scholarship and financial support for his Doctoral Thesis. They thank Prof. A. Z´arate (U. Cat´olica del Norte) and PhD student D. D´ıaz (U. de Chile) for their assistance with XPS analysis. Moreover, they gratefully thank to ECKA Granulate Micromet GmbII for providing part of the Cu powders used in this study, to Laborat´orio Nacional de Luz S´ıncrotron (LNLA, Campinas, Brazil) for its full cooperation in XRD analysis, to JEOL Company Laboratories in Tokyo for TEM support, and to FEI Company in Eindhoven for FIB microscopy support. We acknowledge use of facilities at the John M. Cowley Center for High Resolution Electron Microscopy at Arizona State University, and partial support from NSF Grant DMR030342.
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