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M A T E R I A L S C H A R A C T E R I Z A T I O N 9 0 ( 2 0 1 4 ) 6 2 – 7 0
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Strengthening of Cu–Ni–Si alloy using high-pressure
torsion and agingSeungwon Leea,b,⁎, Hirotaka Matsunagaa, Xavier Sauvagec, Zenji Horitaa,b
aDepartment of Materials Science and Engineering, Faculty of Engineering, Kyushu University, Fukuoka 819-0395, JapanbWPI, International Institute for Carbon-Neutral Energy Research (I2CNER), Kyushu University, Fukuoka 819-0395, JapancUniversity of Rouen, CNRS UMR 6634, Groupe de Physique des Matériaux, Faculté des Sciences, BP 12, 76801 Saint-Etienne du Rouvray, France
A R T I C L E D A T A
⁎ Corresponding author at: Department of MateTel./fax: +81 92 802 2992.
E-mail addresses: [email protected]
1044-5803/$ – see front matter © 2014 Publishhttp://dx.doi.org/10.1016/j.matchar.2014.01.00
A B S T R A C T
Article history:Received 8 October 2013Received in revised form8 January 2014Accepted 8 January 2014
An age-hardenable Cu–2.9%Ni–0.6%Si alloy was subjected to high-pressure torsion. Agingbehavior was investigated in terms of hardness, electrical conductivity and microstructuralfeatures. Transmission electron microscopy showed that the grain size is refined to ~150 nmand the Vickersmicrohardnesswas significantly increased through theHPT processing. Agingtreatment of the HPT-processed alloy led to a further increase in the hardness. Electricalconductivity is also improved with the aging treatment. It was confirmed that thesimultaneous strengthening by grain refinement and fine precipitation is achieved whilemaintaining high electrical conductivity. Three dimensional atom probe analysis includinghigh-resolution transmission electron microscopy revealed that nanosized precipitateshaving compositions of a metastable Cu3Ni5Si2 phase and a stable NiSi phase were formedin the Cumatrix by aging of the HPT-processed samples and these particles are responsible forthe additional increase in strength after the HPT processing.
© 2014 Published by Elsevier Inc.
Keywords:High-pressure torsionUltrafine grainCu–Ni–SiElectrical conductivityAPT
1. Introduction
Cu-based alloys containing Ni and Si are well known and calledCorson alloys [1,2]. They do exhibit strong age hardenability[1–8] due to precipitate hardening leading to strength levels upto 800 MPa together with a high electrical conductivity typicallyof about 50% IACS (International Annealed Copper Standard).Up to date applications in the electronic industry, with smallerdevices in size, require further strengthening with reasonableductility while maintaining high conductivity.
Metallic alloys do offer several strengthening possibilities,namely: strain hardening, solid solution, fine precipitation andgrain refinement. Among these possibilities, the precipitation isthe best way to improve the electrical conductivity because
rials Science and Engineerin
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others are accompanied by generation of lattice strain to reducethe electrical conductivity. Recent studies reported that fineprecipitation in ultrafine-grained structures gives rise to highstrength with enhanced ductility in some age-hardenable Alalloys [9–17]. The grain refinement to the submicrometer rangewas achieved by using theprocess of severe plastic deformation(SPD) [18,19]. However, electrical conductivity was decreasedafter processing by any SPD, irrespective of whether the SPD isadopted fromaccumulative roll-bonding (ARB) [20–22], equal-channel angular pressing (ECAP) [23–25] or high-pressuretorsion (HPT) [25], because of enhanced strain due to latticedefects as summarized by Vorobieva et al. [26]. Thus, thecombination with aging should be important at least toreduce the strain due to dislocations. It was shown that
g, Faculty of Engineering, Kyushu University, Fukuoka 819-0395, Japan.
shu-u.ac.jp (S. Lee).
Fig. 1 – Schematic illustration of HPT facility.
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precipitation was feasible within the ultrafine grains withoutsignificant grain coarsening when aging was conducted atlower temperatures [11,17]. Although it was reported [7,8]that the deformation before aging may be effective toenhance the strength of the Corson alloy, there is no attemptso far to examine the combined effect of grain refinementand subsequent aging on the strengthening of the Corsontype Cu–Ni–Si alloy. This study was thus initiated to develophigh strength and high electrical conductivity with reasonableductility by attaining simultaneous strengthening due to grainrefinement and fine precipitation in a Cu–Ni–Si alloy.
High-pressure torsion (HPT) was used as an SPD process forthe production of ultrafine-grained structures [19].
The aging behavior after the HPT process was examinedby hardness measurements, tensile tests and bending tests.Beside, microstructures were characterized using conventionaltransmission electron microscopy (TEM), high-resolution TEM(HRTEM) and Atom Probe Tomography (APT).
2. Experimental Procedures
This study used a commercially available Cu7025 alloy con-taining major alloying elements of 2.91 wt.%Ni (3.11 at.%Ni)and 0.63 wt.%Si (1.41 at.%Si) with additional impurities of0.14 wt.%Mg (0.36 at.%Mg) and 0.03 wt.%Fe (0.03 at.%Fe)). Thealloy was supplied in the form of sheets with 0.85 mmthicknesses. Disks with 10 mm diameters were then cut fromthe sheet by a wire cutting electrical discharge machine andwere subjected to solid-solution treatment (SST) at 850 °C for0.5 h in an argon atmosphere. The microstructure after SST ishomogeneous with an average grain size of ~100 μm.
The solution treated samples were subjected to HPT atroom temperature under a pressure of 6 GPa for 5 revolutionswith a rotation speed of 1 rpm using a facility as schemati-cally illustrated in Fig. 1. The facility consists of upper andlower anvils. Each anvil has a shallow hole with dimensions of10 mmdiameter and 0.25 mmdepth at the center. Each samplewas placed on the hole and the upper and lower anvils wererotated with respect to each other. Equivalent strain created inthese disks were in the range of 1 to 110 according to theequation as
ε ¼ 2πrNffiffiffi
3p
tð1Þ
where r is the distance from the disk center, N is the number ofrevolution and t is the thickness of disk.
The HPT-processed disks were aged at a temperature inthe range of 200 to 450 °C for 2 h. Aging at a temperature of300 °Cwas carried out for different periods of time up to 100 h.All aging treatments were made in an argon atmosphere andwere terminated by quenching into iced water.
Vickers microhardness was measured on the disks usingan Akashi MVK-E3 by applying a load of 200 g for 15 s. Thedisks were polished with abrasive papers and cloths contain-ing alumina powders. Themeasurements weremade in eightradial directions with an interval of 0.5 mm up to 4.5 mmfrom the disk center. The averages were obtained from eightmeasurements at equal distance from the disk center.
Tensile specimens were machined from the disks withboth gauge lengths and gauge widths of 1 mm by wirecutting electrical discharge machine. They were extracted at2.5 mm off-axis position of the disks and were carefullythinned down to the thickness of 0.6 mm. Tensile tests wereperformed with an initial strain rate of 3.3 × 10−3 s−1 at roomtemperature.
In this study, bending tests were also conducted becausethe ductility can be evaluated with a combination of bothtension and compression modes. Three-point bending testswere carried out using rods with cross sectional dimensions of0.5 × 0.5 mm2 and the lengths of 9 mm as illustrated in Fig. 2.They were cut out at 2.5 mm from the center of HPT disks andwere deformed within 8 mm spans at room temperature witha constant displacement speed of 0.5 mm/min. The loadingduring bending tests was made parallel to the pressingdirection in HPT. The Euler–Bernoulli beam theory was usedto estimate the bending stress as [17]
σ ¼ 3Fl2wh2
where F is the bending load, l is the supporting span (8 mm),wis the bending specimen width (0.5 mm), and h is the bendingspecimen height (0.5 mm).
The electrical resistivity/conductivity was measured usingthe miniature rods prepared for the bending tests as illustrat-ed in Fig. 2. Two probe pins 4 mm apart from each other wereplaced on the rods using spring coils and the voltages wererecorded after applying several different currents. Linearregression analysis was adopted for the voltage and currentto determine the electrical resistivity.
TEM samples were prepared as follows: disks of 3 mm indiameters were punched out from the HPT-processed disks at2.5 mmoff-axis positions. The 3 mmdisks were then polishedmechanically to a thickness of 0.1 mm and further thinned forelectron transparency using a twin-jet electropolisher with asolution of 20%HNO3–80%CH3OH at −30 °C. To clean surfaces,samples were subjected to ionmilling at 4.5 kV for 30 min at aglancing angle of 5°. The microstructures were observed witha Hitachi H-8100 at an accelerating voltage of 200 kV. High
Fig. 2 – Dimension of tensile, bending and electricalconductivity specimen.
Fig. 3 – Variation of Vickers microhardness with distancefrom disk center after HPT and after HPT plus aging for 2 hat 200–450 °C. Hardness level of solid-solution treatmentis included.
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resolution and analytical TEM were performed in a probecorrected JEOL ARM-200F (accelerating voltage of 200 kV).Energy Dispersive X-ray Spectroscopy (EDS) was performedwith a JEOL JED2300 detector and a probe size of 0.2 nm.
Nano-analyses were further carried out by 3D-APT using aCAMECA Tomographic Atom Probe detection system (TAP) [27].Specimens were prepared using standard electropolishingmethods [28] and field evaporated in UHV conditions withelectric pulses (20% pulse fraction at 2 kHz and specimentemperature of 40 K) or femtosecond laser pulses (wave length515 nm, energy of 5 · 10−7 J at 10 kHzandspecimen temperatureof 20 K) [29].
3. Results
3.1. Hardness and Electrical Conductivity Measurement AfterHPT Processing and Aging
Fig. 3 showsVickersmicrohardness plottedagainst thedistancefrom the disk center for an HPT-processed sample and samplessuccessively aged for 2 h at 200–450 °C. The hardness increasesas thedistance being away from thedisk center and saturates toa constant level (270 Hv) at a position ~1.5 mm from the diskcenter. It should be noted that the hardness saturation appearsas a consequence of the balance between the hardening due todislocation accumulation and grain refinement and the soft-ening due to dislocation annihilation (recovery) and dynamicrecrystallization as described earlier [30,31]. The hardness at thesaturation is further increased when the HPT-processed sam-ples are aged at 350 °C or lower. The hardness is almostinvariant with respect to the distance for the sample aged at400 °C and it is lowered for the sample aged at 450 °C. It isapparent that hardening occurs for the aging at 350 °C or lower
but not at 400 °C or higher except the region near the diskcenter. The temperatures of 400 °C or higher are too high for thehardening to be pronounced because the softening process dueto recovery and recrystallization exceeds the hardening dueto precipitation. The exception at the disk center is due toinsufficient strain imposed by HPT so that the aging proceededat a reduced rate.
The constant hardness level at the saturation is highest forthe sample aged at 300 °C (320 Hv) and the amount of theincrease in hardness is about 60 Hv above the as-deformedstate. When compared with the hardness level (70 Hv) of thesolution-treated state, the hardness level after the aging at300 °C is enhanced by more than four times. It should benoted that aging in the same conditions (2 h at 300 °C) of thesolution treated sample without HPT processing leads to thehardness of about 80 Hv and thus the amount of the hardnessincrease is only 10 Hv. It is now apparent that the SPD-processed UFG structure including a high density of latticedefects such as dislocations and grain boundaries enhancesthe precipitation kinetic effectively.
Fig. 4 summarizes the hardness variation with the agingtemperature including the hardness level of the solution-treated sample. Here, the hardnesses are plotted from thevalues at the saturation in Fig. 3. Fig. 4 also plots the IACS valuesmeasured after the aging treatment. Thus, the highest hard-ness is achieved at the aging temperature of 300 °C while theconductivity increases continuouslywith the aging temperature.
The aging behavior at 300 °C is then plotted in Fig. 5 wherethe hardness is taken at the saturation. The hardness increaseswith aging time and reaches a maximum after 2 h. Fig. 5 alsoshows the evolution of the electrical conductivity as a functionof the aging time. It turns out that the conductivity keepsincreasing with aging time while the hardness remains as highas ~300 Hv.
3.2. Tensile Tests
The results of tensile tests are shown in Fig. 6 where thenominal stress is plotted against the nominal strain for
Fig. 4 – Variation of Vickers microhardness and IACS withaging temperature for HPT-processed samples. Levels ofhardness and IACS after solid-solution treatment arealso included.
Fig. 6 – Nominal stress and nominal strain curves for HPTand HPT plus aging including solid-solution treatment.
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specimens processed by HPT and subsequent aging. Fig. 6also includes the solution-treated specimen for comparison.It should be noted that the strain includes the elastic straindue to low stiffness of the tensilemachine (no extensometerswere used) so that the initial slopes of the stress–straincurves are much lower than the slope expected from theelastic elongation of the specimen only. It is apparent thatthe HPT processing increases the tensile strength close to1 GPa. The tensile strength is further increased over 1 GPawith appreciable plastic strain after yielding when it is agedat the peak-aging condition (2 h at 300 °C). However, thetensile ductility is lost and the tensile strength decreasesafter aging at higher temperature (2 h at 450 °C).
3.3. Bending Tests
The mechanical property was also evaluated by bending testsand the results are delineated in Fig. 7 for the samples given inFig. 6. The bending strength exhibits the same trends as thetensile testing. However, it is found that the bending ductility
Fig. 5 – Variation of Vickers microhardness and IACS withaging time. Levels of hardness and IACS after solid-solutiontreatment are also included.
is well reserved for all specimens and comparable to thesolution-treated specimen. The good ductility is also con-firmed with the appearance of the specimen shapes after thebending tests shown in the inset. These results are differentfrom those of the tensile tests and it is considered that thedifference is attributed to the specimen size for the tensiletesting which is small so that it is more sensitive to themisalignment along the tensile axis. The ductility may beevaluated more properly in the bending test as it includes notonly a tension mode but also a compression mode which isgood for evaluation of the ductility for less ductile materials.
3.4. Microstructure After Solid-Solution Treatment
APT analyses of the solutionized alloy revealed that alloyingelements (Ni, Si, Mg and Fe) are homogeneously distributedafter quenching (Fig. 8) and the following concentrations weremeasured: 3.3 ± 0.2 at.% Ni, 1.3 ± 0.1 at.% Si and 0.1 ± 0.03 at.%Mg. This measurement is well consistent with the nominalcomposition of the alloy (except for the minor element of Mg).It is then confirmed indicating that the solution treatment wassuccessfully achieved.
Fig. 7 – Bending stress and displacement curves for HPTand HPT plus aging including solid-solution treatment.Appearance of specimens after bending tests is included atcorresponding curves.
Fig. 8 – Elemental distribution analyzed by 3D-APT aftersolid-solution treatment. Ni (blue), Si (green) and Mg (pink).Data collected with electric pulses. (For interpretation of thereferences to color in this figure legend, the reader is referredto the web version of this article.)
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3.5. Microstructures After HPT Processing
TEM micrographs with an SAED pattern after HPT processingof the solution-treated specimen are shown in Fig. 9. Thegrain size is well reduced to the order of ~150 nm. The SAEDpattern exhibits rings, indicating that the microstructureconsists of grains with randomly-oriented grains. Some grain
Fig. 9 – TEM micrographs after HPT: (a) bright-field image,(b) SAED pattern and (c) dark-field image taken by arrow inSAED pattern.
boundaries are not smooth but irregular in nature and manystrained contrasts are visible within grains, which are known tobe a typical feature after severe plastic deformation as reportedearlier by high-resolution electron microscopy [32].
3.6. Microstructures After Peak Aging
Fig. 10 shows the microstructure of the sample subjected topeak-aging at 300 °C for 2 h after HPT processing of thesolution-treated alloy. It should be noted that the observationwas made in the region where the hardness saturation isreached. Microstructural features are very similar to theun-aged sample (Fig. 9), i.e. the mean grain size is still about150 nm with high angle grain boundaries. However, somenanoscaled precipitates are clearly visible as indicated byarrows in Fig. 10, which have nucleated during the agingtreatment. It is important to note that the grain size remainssmall after the aging. It is confirmed that the simultaneousstrengthening due to grain refinement and precipitate forma-tion is attained. APT analyses were carried out to characterizethese precipitates of the same sample and the results areshown in Fig. 11 where (a) is a distribution of Cu, Ni and Siatoms, (b) of Ni, Si and Mg, and (c) is a compositional profileacross the particle. The particle contains mostly Ni and Si asexpected and it is interesting to note that Mg atoms arepreferentially segregated along the interface between theprecipitate and the matrix. This feature is quantitativelyexhibited on the composition profile in Fig. 11(d).
The precipitate contains slightly more Ni than Si andthe amount of Cu is very low. The exact composition wasmeasured by filtering the data (Ni threshold 45%, samplingvolume 1.5 × 1.5 × 1.5 nm3), and it is as follows (at.%): 56.3 (±1)Ni, 41.8 (±1) Si, 1.0 (±0.3) Cu and Mg is below the detectionlimit of the instrument. Thus, the Ni/Si ratio in this particle isabout 1.3.
Smaller particles only few nanometers in diameter werealso detected (Fig. 12(a)). As shown on the typical composi-tional profile computed across such a cluster (Fig. 12(b)), theycontain Ni and Si with a Ni/Si ration of ~2.5. Because they alsocontain a significant amount of Cu (~30 at.%), it is consideredthat these particles are some metastable phases.
Fig. 10 – TEM micrograph after HPT plus peak aging at 300 °Cfor 2 h: arrows indicating nanoscaled particles.
Fig. 11 – Results of 3D-APT analyses for particle with size of~15 nm after HPT plus peak aging at 300 °C for 2 h:(a) distribution of Cu, Ni and Si atoms, (b) distribution of Ni,Si and Mg, and (c) compositional profile across particle. Datacollected with laser pulses.
Fig. 12 – Results of 3D-APT analyses for particle with size of~4 nm after HPT plus peak aging at 300 °C for 2 h:(a) distribution of Cu, Ni and Si atoms and (b) compositionalprofile across one particle. Data collected with electric pulses.
Fig. 13 – TEMmicrographs after HPT plus overaging at 450 °Cfor 2 h: (a) low and (b) high magnification views.
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3.7. Microstructures After Overaging
The microstructure after aging at 450 °C for 2 h is shown inFig. 13. Apparently, grain growth took place as in (a) and thegrain size increased up to ~300 nm during the aging withdislocation density fairly reduced within the grains. Closeexamination in the grains reveals that precipitate particlesare present with a size of ~25 nm as magnified in (b). Fineparticles with the similar sizes are also detected by elementalmapping using energy dispersive spectrometry (EDS) as shownin Fig. 14(a). The line profile analysis that was performed acrossthe precipitate (Fig. 14(b)) shows that a significant amount of Cuis detected. However, it is not certain from this analysis if theparticle contains Cu since it may be embedded in the matrix.The yield of the Si Kα being lower than that of Ni Kα, theintensity ratio seen on the line profile analysis does directlyreflect theNi/Si in theparticle. Froma spectrum recordedon the
Fig. 14 – (a) Elemental mapping using STEM-EDS after HPTplus overaging at 450 °C for 2 h. (b) X-ray intensity profileacross one particle as indicated in (a).
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particle and using theoretical k-factors (kSiCu = 0.4946, kNiCu =0.8822), the Ni/Si ratio was estimated close to 2/1. It should benoted that the energyof Si KαX-rays is five times lower than theenergy of Ni K [33] leading to a stronger attenuation byabsorption of the Si signal and thus a possible over-estimationof theNi/Si ratio during suchEDSanalysis. These fineprecipitatesare responsible for high strength even at overaging conditionsas 450 °C.
4. Discussion
This study clearly showed that strengthening occurs duemostly to grain refinement through HPT processing butadditional strengthening by subsequent aging. Peak hard-ness including tensile and bending strengths was attainedwhen aging was undertaken at 300 °C for 2 h. However,this aging temperature is fairly low and the aging time isshortened when compared with the aging conditions report-ed in other studies [1–8]. It is considered that this is due to thepresence of many lattice defects produced by HPT so that theatomic diffusion is enhanced. In fact, this is well supportedby several direct and indirect experiments such as in-situ
operation of HPT using a high-energy synchrotron light [34],time-differential dilatometry of HPT-processed nanocrystallineFe, Ni and Cu [35] and softening of HPT-processed Ag and Auafter exposure to room temperature [36]. Nevertheless, the peakhardness is achieved as high as 320 Hv and the tensile strengthmore than 1 GPa with good bendability and the electricalconductivity of ~30% IACS. Suzuki et al. [7] reported the peakhardness as 330 Hv comparable to the present peak hardnessbut aging and cold rolling were repeated at ~450 °C for ~3 h andthen ~360 °C for ~74 h, suggesting that the deformation beforeaging should be effective to enhance the hardness. The presentstudy has proved that severe plastic deformation throughHPT is well effective with the aging condition at the reducedtemperature as 300 °C and for the accelerated time as 2 h.
The APT analysis including the STEM-EDS analysis furthershowed the formation of nanosized-precipitates during agingof the HPT-processed samples. Three types of the precipitatesare detected with the sizes of ~4 nm and ~15 nm after peak-aging and of ~25 nm after over-aging. The smallest precipitatehaving the size of ~4 nm (Fig. 12) contains Cu with a Ni/Siratio of 2.5, suggesting the formation of a metastable phase asCu3Ni5Si2. The precipitate with the size of ~15 nm (Fig. 11)consists of Ni and Si with a Ni/Si ratio of ~1.3, suggestingthe formation of a NiSi or Ni3Si2 phases. The largest one withthe size of ~25 nm (Fig. 14) takes also a Ni/Si ratio of ~2,suggesting the formation of Ni2Si. It should be noted that IACSincreased with aging time because of the precipitation ofsuch particles, which reduces the dissolved alloying elementsout of the matrix and makes electron movement easier inthe sample.
According to an equilibrium phase diagram of the Ni–Sisystem [37], six intermetallic phases are present, such asNi3Si, Ni5Si2, Ni2Si, Ni3Si2, NiSi and NiSi2. Among them, asreviewed by Schlesinger [38], the Ni2Si phase has the lowestformation enthalpy based on experimental measurements.However, the formation enthalpy is the lowest for the Ni5Si2,Ni3Si2, or NiSi phase when the reference is included fromtheoretical modeling. Because the difference between theformation enthalpies for Ni5Si2, Ni2Si, Ni3Si2 and NiSi isminimal, it is reasonable that phases with varying Ni/Si ratiosmay form in the matrix depending on local energy variation,which is more likely to the case in a heavily deformed non-equilibrium state produced by HPT processing.
When compared with the past studies on the Cu–Ni–Sisystem, three types of precipitates were reported, which areNi2Si [2,4–6], N5Si2 [39] or Ni3Si [3,6]. Consistency with thepresent study is thus Ni2Si, but the observation of Cu3Ni5Si2and Ni3Si2 is unique to this study. The former particle con-taining Cu is considered to be a metastable phase formed inan early stage of the precipitation as the size is very smallas ~4 nm,whichmay then be rather considered as clusters. It isexpected that this cluster evolves to Ni3Si2 while Cu is replacedwith Ni and Si and finally into Ni2Si, the most stable phase forthe largest particles. A similar evolution was also reportedrecently in an APT analysis of a Cu–1 mass%Cr–0.12 mass%Zralloy [40]. Aging of the alloy was accompanied by Cr precipita-tion but the precipitate contains a fairly large amount of Cu(~40 at.%) when the size is small at an initial stage of aging.Thus, the Cu content in the Cr precipitate gradually decreasesas the precipitate grows with aging.
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5. Summary and Conclusions
(1) Ultrafine-grained structure with a grain size of ~150 nmwas obtained in a Cu–2.91 wt.%Ni–0.63 wt.%Si alloy(7025) through application of HPT.
(2) Such a fine grain structure was retained during agingat 300 °C for up to 100 h, while age hardening occurs toincrease the hardness well above the hardness levelattained by the HPT processing. A peak hardness wasobtained after aging at 300 °C for 2 h. Thus, simulta-neous hardening due to aging and grain refinementwas achieved in the Cu–Ni–Si alloy.
(3) Electrical conductivity increases continuously with agingand reaches 40% IACS after aging for 100 h.
(4) Three Dimensional AtomProbe Tomography (APT) includ-ing the STEM-EDS analysis revealed the formation ofnanosized-precipitates during aging of the HPT-processedsamples, which have the sizes of ~4 nm and ~15 nm afterpeak-aging and of ~25 nm after over-aging, the formercorresponding to a metastable phase of Cu3Ni5Si2 and thelatter two to a stable phase of Ni2Si.
Acknowledgments
One of the authors (ZH) would like to thank the Universityof Rouen for having an opportunity to participate in theinternational invited program and carry out the APT analyses.This workwas supported in part by a Grant-in-Aid for ScientificResearch from the MEXT, Japan, in Innovative Areas “BulkNanostructured Metals” (22102004) and in part by KyushuUniversity Interdisciplinary Programs in Education and Projectsin Research Development (P&P).
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