Microstructure and deformation behaviour of ultra-fine-grained ODS copper prepared by mechanical alloying

Transcription

1 Microstructure and deformation behaviour of ultra-fine-grained ODS copper prepared by mechanical alloying U. Martin 1, D. V. Kudashov 1, M. Heilmaier 2 & H. Oettel 1 1 Technische Universität Bergakademie Freiberg, Germany 2 Otto-von-Guericke-Universität Magdeburg, Germany Abstract Novel oxide dispersion strengthened (ODS) copper alloys were synthesised by mechanical alloying (MA) at low temperature (cryomilling) and subsequent consolidation via hot pressing in a protective argon atmosphere. In the present study the deformation behaviour of ODS copper with ultra-fine-grained (UFG) structure prepared by MA of pure copper powder and 3 vol.% of incoherent oxide particles, Yttria and Ceria, was examined in a wide temperature range at different deformation rates. The microstructure of UFG ODS copper materials with high yield strength at room temperature was examined by means of highresolution scanning electron microscopy (HRSEM). Compressive deformation experiments at room temperature yield a substantial increase of the 0.2% yield strength of the UFG ODS copper, when compared with pure copper. The overall yield strength could best be modelled by a linear superposition of the Orowan stress for particle hardening and by the contribution from fine grain strengthening after Hall-Petch. The creep behaviour of UFG ODS copper was examined in the temperature interval of 500 to 700 C at different deformation rates. As a result, the oxide particles situated at the grain boundaries inhibit the classical Nabarro-Herring or Coble diffusional creep. With two modifications the creep behaviour can be described within the frame of the Rösler-Arzt diffusional creep model of thermally activated detachment of the grain boundary dislocations from the departure side of particles lying in or adjacent to grain boundaries. Keywords: mechanical alloying, cryo-ball milling, ODS-copper, dispersion strengthening, grain size strengthening, creep behaviour, diffusional creep.

2 342 High Performance Structures and Materials II 1 Introduction Advanced materials for conducting applications require an optimal combination of high mechanical strength, creep and fatigue properties together with superior electrical conductivity. One possibility to approach this target is to combine grain size strengthening of a (pure) copper matrix with dispersion strengthening via a few volume percent of nano-sized, incoherent and insoluble oxide particles (ODS). Therefore, electrical and thermal conductivity of the matrix should not be strongly affected. The present study shows our efforts in applying the technique of MA to synthesize novel oxide dispersion-strengthened copper alloys as potential candidates for functional applications with the main emphasis placed on an optimised combination of mechanical and electrical properties at room and at ambient temperature. Following the principles of particle selection by Groza and Gibeling [1] the optimum choice of oxide type depends on several intrinsic properties of the oxide dispersoids; in particular their thermodynamic stability and a low diffusivity and solubility of their cationic component within the metallic matrix are of major importance. Applying these criteria, a series of oxides were selected to systematically study the validity of the above requirements for ODS metals. 2 Experimental details High purity powders of copper and the two oxides (Y 2 O 3 and CaO 2 ) with a powder particle size smaller than 35 µm and 1µm, respectively, were used as starting materials. A mixture of copper with 3 vol. % of oxide was mechanically alloyed in a high energy planetary ball mill PM4000 (Retsch, Germany) at a rotation velocity of 200 rpm with a powder-to-ball weight ratio of 1:14. Milling and powder handling had to be carried out under protective atmosphere (argon). A fine and homogeneous distribution of the oxides in the copper matrix could be obtained after 20 h only under repeated intensive cooling of the milling vials in liquid nitrogen, see [2,3] for details. Rods with 98% of the theoretical density were produced by uniaxial hot pressing with 650 MPa at 600 C for 20 minutes. The deformation behaviour of the ODS copper alloys was characterised utilising compression tests with cylindrical samples of 4 mm diameter and 6 mm height at a constant deformation rate in a servo-hydraulic, digitally controlled MTS test equipment. Microstructural investigations were carried out in a high-resolution scanning electron microscope LEO 1530 (FEG LEO, Germany). 3 Results and discussion 3.1 Copper grain sizes and dispersoid diameters of Cu-3%Y 2 O 3 and Cu-3%CaO A compilation of particle and grain size parameters of the different ODS copper composites is summarized in Table1. Exemplarily, the microstructure of Cu- 3%CaO after manufacturing and consolidation is displayed in Fig. 1. The oxide

3 High Performance Structures and Materials II 343 particles are mostly spherical in shape and predominantly located at the grain boundaries, although the finest particles are present in the grain interior. In Fig. 2 the particle size distribution obtained from a series of pictures such as Fig. 1 is plotted for Cu-3%CaO as relative frequency vs. logarithmic classification. To obtain statistically secured particle parameters a minimum number of 500 dispersoids was analysed for each material composition: the mean radii (of circular section in the glide plane) r s of Y 2 O 3 and of CaO were evaluated as 20 nm and 17 nm, respectively. In addition to that, the areal fraction of each class of particles is plotted as dark grey bars in Fig. 2. These distributions clearly show that the areal fraction of dispersoids coarser than 54 nm predominates the one stemming from smaller dispersoids. Figure 1: SEM micrograph of grain structure and dispersoids (dark particles) in consolidated Cu-3% CaO. The copper matrix grain size distribution was evaluated in a similar manner applying the line intersection method: surprisingly, the mean grain size is 150 nm in both alloys. Utilizing the classical Zener-theory for grain growth restricted by the presence of pinning dispersoids (d crit = 2d ox / 3f cited in [4]) grossly overestimates the experimentally determined values, cf. Table 1. Comparison of the particle size distribution with the copper grain size distribution (the latter not shown here) yields that a significant number of oxide particles exhibit a size comparable with the matrix grain size. These particles were disregarded for the following analysis of hardening mechanisms. As a consequence, the amount of dispersoids, which may interact effectively with the dislocations through the Orowan mechanism, is smaller than the entire volume fraction of particles. The areal fraction of oxide dispersoids f A exhibiting diameter values smaller than

4 344 High Performance Structures and Materials II 83 nm and, thus, being attractive from a strengthening viewpoint has been determined to about 70 % for Y 2 O 3 and 50 % for CaO, respectively. Assuming the dispersoids to be randomly distributed we used the relation f A = f D for the following analysis with f D being the revised volume fraction and r D the revised radius of oxides available for dispersion strengthening as tabulated in Table 1. Then, the value of matrix grain size computed through the Gladman equation [5] d crit π r D 3 2 (1) = 3 f 2 d D d 0 correlates well with the effective particle volume fraction and, accordingly, with the reduced particle size, cf. Table 1. The analysis yields d/d 0 = 1.55 concurrently for both ODS alloys, a quantity which agrees well with the theoretical limit of 2 < d/d 0 < 2 [5]. Table 1: Dispersoid parameters and grain sizes of the ODS copper alloys investigated and calculated after Zener- und Gladman equation. Material Cu grain size estimated by HRSEM [nm] Oxide radius r D (HRSEM) [nm] Volume fraction f D of oxide particles Grain size [nm] after Zener Grain size [nm] and d/d 0 value after Gladman, equ. 1 Cu-3% Y 2 O 3 Cu- 3%CaO Room temperature deformation behaviour In Fig. 3 the compressive deformation experiments at room temperature of the UFG ODS materials are compared with commercially available GlidCop Al-15 (Al 2 O 3 dispersion strengthened copper) [13] exhibiting similar overall stress levels and stress strain behaviour. However, the stress-strain curves of the present materials reveal two peculiarities: first, a very high stress level is already observed from the onset of plastic deformation, i.e. the ratio σ 0.2 /σ max is close to unity, which can be rationalized by the fact that dislocation movement in ODS materials is hampered by the higher density of strong obstacles. Second, a constant flow stress over a wide deformation interval is noted without any visible crack formation on the surface of the specimens demonstrating that the addition of a few vol.% of nanoscale oxide dispersoids does not deteriorate the intrinsically high potential of Cu for plastic deformability. As will be treated quantitatively in the following, the very high ambient temperature yield strength (σ MPa) is the result of a simultaneous occurrence of grain boundary strengthening according to Hall [6] and Petch [7] and of particle hardening due to the Orowan mechanism which is

5 High Performance Structures and Materials II 345 the sole overcoming mechanism of incoherent dispersoid particles at room temperature. Figure 2: Oxide particle distribution (light grey bars) and areal fraction of each particle size class (dark grey bars) of Cu-3% CaO, logarithmic classification. Figure 3: True stress σ vers. true plastic strain ε pl for different nanocrystalline ODS copper alloys.

6 346 High Performance Structures and Materials II Also, in other nanocrystalline, particle-strengthened materials the Orowan mechanism was proven to operate for grain sizes on the order of 100 nm by transmission electron microscopy [8, 9]. In its simplest form, the superposition of the individual strengthening mechanisms could be carried out linearly σ 0.2 = σ 0 + k HP d -1/2 + σ OR (2) with σ 0 being the Peierls stress (which is, however, negligible in fcc materials), k HP = 4.5 MPa mm 1/2 [10] is the Hall-Petch-constant. For the Orowan by-passing stress we apply the actual formula suggested by Kocks [11]: 3/ 2 edge [ ln( 8 r )] [ ( )] ( ) s / b K 1/ 2 ln L / b b L 2 rs σ = 0.9 M (3) OR with M = Taylor factor, b = nm is the Burgers vector, L = mean planar dispersoid spacing (= (32 / 3 π f) 0.5 r s ) and K edge = Gb 2 / 4π being the prelogarithmic line tension factor of a straight edge dislocation, G = 42 GPa is the shear modulus at room temperature. As mentioned above, for the calculation of the Orowan stress the interaction of the dislocations with the coarse oxide particles was neglected. Therefore, L was calculated with the reduced value for the volume fraction f D as listed in Table 1. The values of the individual hardening contributions calculated according to eqns. (2) and (3) are compared with the experimental σ 0.2 values and this analysis yields that the grain boundaries provide the largest hardening contribution of roughly two thirds, whereas the remainder of only one third of the entire strength is caused directly by the oxide particles due to the Orowan mechanism. Further, it is nicely supported by experimental data on the maximum stress in pure copper with grain sizes of 150 to 200 nm yielding values of σ max = MPa [14, 15]. However, in both ODS alloys the measured yield strength is slightly underestimated by the sum of both calculated strengthening mechanisms. This may be due to neglecting either the Peierls stress or, more likely, a contribution from solid solution and / or precipitation strengthening from oxygen or iron, dissolved in the copper matrix at a level of around 700 (oxygen) and 500 (Fe) wt.ppm after mechanical alloying and consolidation [12]. 3.3 Creep deformation behaviour The creep behaviour of UFG ODS copper, containing 3 vol.% of Yttrium and Calcium oxides, respectively, was examined in the temperature range from 500 C (0.57 T m ) to 700 C (0.71 T m ) at strain rates ranging from 10-6 to 10-2 s -1. In Fig. 4 typical true stress true plastic strain curves of mechanically alloyed UFG ODS copper at constant compressive strain rate are compared indicating the conventional stages of creep I (transient creep, strain hardening) and II (steady state creep, dynamic equilibrium between dislocation multiplication and annihilation). In the latter stage, the flow stresses defined at 5 % plastic strain were taken to plot the stress sensitivity of strain rate at different temperatures in the form of a logε - log σ representation, see Fig. 5. As characteristic for ODS

7 High Performance Structures and Materials II 347 alloys, the values of the slope of the curves, i.e. the stress exponents n 500 C = 20 and n 700 C = 10 are high and, in particular, near those of coarse grained GlidCop, an ODS Cu- Al 2 O 3 material [16]. Therefore, they are also substantially higher than the values of 3 to 5 reported for pure coarse grained copper [17]. Besides, the stress exponents remain relatively constant over the strain rate range investigated indicating a single deformation mechanism to operate. Figure 4: True stress true plastic strain curves of ODS Copper Cu-3vol.% Y 2 O 3, temperatures and compressive strain rates as indicated. Figure 5: Logε - log σ plot of creep data of ODS copper Cu-3vol.% Y 2 O 3, GlidCop Al-15 and pure copper.

8 348 High Performance Structures and Materials II Figure 6: Comparison between compressive creep data on Cu-3vol.%CaO and calculated creep rate. Parameters: k gb = 0,74; f D = 0,015; r D = 14 nm; α, b gb, D eff, and σ OR values at 500 and 700 C see Fig. 5. In UFG material deformation at elevated temperatures may be controlled by diffusional creep. However, classical creep models for diffusional creep, see e.g. [18], cannot account for the observed creep behaviour in Figs. 4 and 5, since they yield a stress-dependency of strain rate of n = 1. Motivated by experimental observations on creep in fine-grained ODS Al alloys Rösler & Arzt extended their original creep model for thermally activated detachment of dislocations from incoherent particle / matrix interfaces (in essence in single crystals) [19] to incorporate the effect of grain size d g via the interaction between grain boundary dislocations and dispersoids lying in or adjacent to grain boundaries [20,21]. Following the kinetic equation ( ) σ σ 2 ρ, gb G bgb rd 1 k gb 1 2 D L OR k eff D ρ σ (1 gb ε = exp d b k T g gb was modified in a twofold manner (L D and r D are mean interparticle spacing and radius, respectively, k gb is the relaxation factor of the dislocation segment in the grain boundary): 1. instead of the grain boundary diffusion coefficient D gb we prefer to use an effective diffusion coefficient D eff incorporating the simultaneous effects of volume (D v ), grain boundary and pipe diffusion (D c ) on dislocation movement. b 2 ) 3 2 (4)

9 High Performance Structures and Materials II the applied creep stress σ is substituted for by the reduced creep stress σ - σ ρ,gb necessary for the sole particle overcoming process of a grain boundary dislocation [22]. σ ρ,gb can be calculated in analogy to Taylor [23]. The resulting continuous modelled curves in Fig. 6 agree well with the measured creep data (triangles and squares). The relatively low values of k 0.8 for both materials investigated give evidence for a strong interaction of grain boundary dislocations with the oxide particles and coincide nicely with k-values found for Y 2 O 3 [24]. 4 Conclusions Some of the more important conclusions are listed below: 1. The homogeneous distribution of the dispersoids prevents the copper matrix from significant grain growth during exposition to high temperatures (consolidation) yielding a sub-micron grain size of 150 nm. 2. Only one third of the overall 0.2% yield strength of the ODS copper alloys is attributed to direct strengthening of the oxide dispersoids via the Orowan mechanism [11]. The remaining two thirds of the strength stem from fine-grain strengthening according to Hall [6] and Petch [7] caused by the (indirect) effect of the oxide particles on stabilizing a sub-micron grain structure. This result coincides nicely with experimental data for nano-crystalline pure copper of comparable grain size [14, 15]. 3. Creep in UFG copper alloys dispersion strengthened by 3 vol.% of Yttria or Calcia, respectively, can be described in a unified way on the basis of the dislocation detachment creep model by Rösler and Arzt [19] extended for the influence of grain size through the interaction between grain boundary dislocations and dispersoids lying adjacent to grain boundaries. Acknowledgement The present work was funded by the Deutsche Forschungsgemeinschaft within the frame of the special project Graduiertenkolleg Werkstoffphysikalische Modellierung at the TU Bergakademie Freiberg. We are grateful to Prof. Dr. L. Schultz and Prof. Dr. J. Eckert from IFW Dresden for continuous support. References [1] J.R. Groza, J.C. Gibeling, Mater. Sci. Eng., A171, 1993, 115. [2] U. Grundmann, M. Heilmaier, L. Schultz, german patent DE C1. [3] U. Grundmann, M. Heilmaier, U. Martin, H. Oettel, L. Schultz, Z. Metallkd., 94, 2003, 587. [4] C.S. Smith, Trans. AIME 175, 1948, 151. [5] T. Gladman, Proc Royal Society, A294, 1966, 298. [6] E.O. Hall, Proc. Roy. Soc., B 64, 1951, 474.

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