Enhancement of strength and ductility of Cu–Sn–Zn alloyby iron addition

Youxiong Ye a, Xuyue Yang a,b,n, Chenze Liu c, Yangzhi Shen a, Xiangkai Zhang a, Taku Sakai d

a Educational Key Laboratory of Nonferrous Metal Materials Science and Engineering, School of Materials Science and Engineering, Central South University,Changsha 410083, Chinab Institute for Materials Microstructure, Central South University, Changsha 410083, Chinac State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, Chinad UEC Tokyo (The University of Electro-Communications), Chofu, Tokyo 182-8585, Japan

a r t i c l e i n f o

Article history:Received 26 December 2013Received in revised form3 June 2014Accepted 14 June 2014Available online 21 June 2014

Keywords:Copper alloysAgingIron precipitatesMicrostructureMechanical properties

a b s t r a c t

Effects of iron addition to Cu–Sn–Zn alloy on the microstructures and mechanical properties wereinvestigated. The Cu–Sn–Zn–Fe alloy, aged at 500 1C for 4 h, showed a peak hardness and an excellentcombination of strength and ductility; e.g. a yield strength of 264 MPa, an ultimate tensile strength of514 MPa and a fracture elongation of around 50%. The precipitates responsible for the strengthening areidentified as bcc Fe with diameters ranging from 20 to 60 nm, and the interfacial relationship betweenthe precipitate and copper matrix is ð002ÞM==ð110ÞP . It is discussed that Orowan dislocation bypassingcan be a dominant strengthening mechanism; quantitative calculations for strengthening due toprecipitation and grain refinement are roughly similar to the experimental results of the present Cualloy. Additionally, the high work-hardening rate in a peak aged condition can be responsible for hightensile ductility of the alloy.

& 2014 Elsevier B.V. All rights reserved.

1. Introduction

Copper and copper-based alloys are one of the most commer-cially important metallic materials because of their excellentphysical and mechanical properties. They are easily fabricatedand thus have been used in numerous applications [1–4]. From anengineering point of view, improvement of the strength andductility for structural materials has been a long-standing missionfor material scientists. Strength and ductility, however, are oftenparadoxical mechanical properties of structural materials includ-ing copper and copper alloys [5–10]. For example, nanostructuredcopper prepared by dynamic plastic deformation [10] exhibitsseveral times higher strength than that of coarse-grained copper,while it shows a very limited tensile ductility.

Recent studies [11,12] show that nano-sized second-phase particlesintroduced in metallic materials can offer a possibility for substantialstrengthening and moderate ductility. Takata et al. [11] fabricated anultrafine-grained Cu–Cr–Zr alloy, including Zr-containing nano-sizedprecipitates, by using an accumulative roll-bonding process andsubsequent aging treatment. The Zr-containing nanoprecipitates pro-moted an improvement of both tensile strength and ductility, while its

total elongation is relatively as low as 14.3% and limits its practicalapplications. It is also reported in [12] that introduction of ironnanoparticles into a Cu–Sn–Zn–Pb alloy by centrifugal casting sig-nificantly improves the tensile strength to reach 420MPa and allowsthe fracture elongation to reach 25.4%. The strengthening mechanismand the reason for the increased ductility, however, were notdiscussed and therefore still remain unclear.

In this paper, we demonstrated a strategy to achieve an excellentcombination of strength and ductility in a copper alloy. A clearimprovement strength without compromising ductility was achievedby introducing iron nanoprecipitates into a Cu–Sn–Zn alloy. It isdiscussed that the enhancement of strength can result from pre-cipitation strengthening and also grain refinement, and that moder-ate ductility can be attributed to high work-hardening rate Θresulting from the fine precipitates. This strategy can be applicableto many other alloys, and paves the way for their large-scaleindustrial applications.

2. Experimental procedures

The equilibrium phase diagram of Cu–10Sn–2Zn–1.5Fe (wt%) isshown in Fig. 1, which is derived from the thermodynamic calcula-tion software Thermo-Calc. The possible existent phases are fcc Cu(α) solid solution, fcc δ phase and bcc Fe one. The δ phase drops to

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Materials Science & Engineering A

http://dx.doi.org/10.1016/j.msea.2014.06.0470921-5093/& 2014 Elsevier B.V. All rights reserved.

n Corresponding author. Tel.: þ86 731 88876470; fax: þ86 731 88830136.E-mail address: yangxuyue@mail.csu.edu.cn (X. Yang).

Materials Science & Engineering A 612 (2014) 246–252

zero at around 200 1C, where only α-Cu solid solution and bcc Fephase exist in this system. It is seen in Fig. 1, therefore, that agingtreatment above 200 1C after solution treatment can not onlyprecipitate the iron second phase, but also eliminate the brittle δphase completely. In the present study, two kinds of Cu alloys of thenominal compositions of Cu–10 wt% Sn–2 wt% Zn and 0 or 1.5 wt%Fe were prepared from high purity copper (499.99%), tin(499.99%), zinc (499.9%) and iron (499.99%) by melting in amedium frequency electrical furnace at 1200–1300 1C. The alloyswere solution treated at 750 1C for 1 h in a vacuum chamber, water-quenched and finally aged at various temperatures ranging from 300to 550 1C for 4 h. The chemical compositions of the two alloys wereanalyzed by inductively coupled plasma atomic emission spectro-scopy (ICP-AES) and the results were Cu–10.0 wt% Sn–1.76 wt% Zn(Alloy A) and Cu–10.3 wt% Sn–1.81 wt% Zn–1.26 wt% Fe (Alloy B).

The Vickers hardness was measured on each specimen at roomtemperature under a load of 1 kg with a duration of 30 s. Tensilespecimens with a dog-bone shape (gauge size: 6 mm�1.5 mm�1mm) were machined by an electro-discharge machine.Tensile tests were performed at a strain rate of 3�10�3 s�1 at roomtemperature. Microstructure examination was performed by usingoptical microscope (OM), scanning electron microscopy (SEM, FEIQuantaTM 650 FEG) with energy dispersive spectroscope (EDS) andtransmission electron microscopy (TEM, FEI Tecnai G2 20). Theconstituent phases were identified by X-ray diffraction with Cu Kαradiation (XRD, Rigaku D/max 2550 X-ray diffractometer).

3. Results

3.1. Microstructures of as-cast samples

Fig. 2 shows the back scattered electron images (BSE-SEM) andthe corresponding EDS analysis results for the as-cast samples ofAlloy A and Alloy B. The regions with light contrast in Fig. 2(a) and(b) are the ones between dendrites and at grain boundaries of thedark matrix. The representative dendritic casting structure of AlloyA can be seen as darker regions in gray matrix in Fig. 2(a). Thesedendrite arms show two orders of magnitude larger and areaccompanied by arms of the second dendrite. By contrast, the BSEimage of Alloy B in Fig. 2(b) shows an almost equiaxed dendritic andfiner grain structure with diameters from 10 to 35 μm. This suggeststhat an addition of iron may increase the nucleation sites of thecopper matrix during the solidification [12,13].

On the other hand, atomic number contrasts afforded by BSEimages show that elements of higher atomic numbers than that of

Cu are enriched at grain boundaries [14], as indicated by brightcontrast region. The corresponding EDS analysis results are pre-sented in Fig. 2(c) and (d). The data points in Cu content o84% arefrom light contrast regions between dendrites or at grain bound-aries, and the ones in Cu content 490% are from dark contrastregions within the matrix grains. It is seen that Sn content at neargrain boundaries is around 15–21 at%, much higher than that indark matrix of around 1.5–7 at%. In contrast, the differences of Znand Fe content in the two regions are negligibly small. The averagecompositions of Cu in the Sn rich phase region are around 79–84 at% for both the alloys. According to the Cu–Sn–Zn ternaryphase diagram and Refs. [15,16], the stoichiometry of Sn rich phaseshould be near Cu41Sn11 (δ phase). Fig. 3 shows X-ray diffraction(XRD) patterns of the as-cast specimens. It is observed that Alloy Acontains two different phases, i.e. α-Cu solid solution and δ-Cu41Sn11 phase, and that Alloy B consists of α-Cu solid solution,δ-Cu41Sn11 phase and α-Fe phase. These results correspond well tothe ones of Thermo-Calc software and EDS (Figs. 1 and 2).

3.2. Aging behaviors and peak-aging microstructures

Fig. 4 shows the isochronal aging curves for Alloy A and Alloy B.Solution treated samples were aged at temperatures ranging from 300to 550 1C for 4 h. It is seen in Fig. 4 that the hardness of Alloy B isoverall much higher than that of Alloy A and its peak hardness is123 Hv at 500 1C, above which over-aging occurs rapidly. So thetemperature at a peak age hardening is defined as an optimum agingone. In contrast, Alloy A does not show any appreciable age hardeningresponse. It is concluded that precipitation of secondary phase relatedto Fe addition could bring about such a typical aging behavior of AlloyB. Hereafter, aging 4 h at 500 1C is taken as a peak-aged condition forthe two alloys, though Alloy A has no age hardening response.

The microstructures of Alloys A and B in a peak-aged state arerepresented in Fig. 5. It is remarkable to see that the grain size ofAlloy B is much finer than that of Alloy A, i.e. 32.0 mm against209.6 mm of Alloy A. After solution treatment at 750 1C for 1 h andaging at 500 1C for 4 h, the grain size of Alloy A increasedconspicuously, while that of Alloy B scarcely changed comparedwith that for the as-cast one. It is also seen in Fig. 5 that there are aconsiderable amount of precipitates evolved in the matrix of AlloyB. These second-phase particles may act as obstacles to the motionof grain boundaries. Note here that the dendrite structures in theas-cast alloys (Fig. 2) completely disappeared and the δ phaseredissolved into the copper matrix during heat treatment, thusleading to these homogeneous microstructures.

Fig. 6 shows the XRD patterns of the two alloys in the peak-aging condition. It is seen that Alloy A demonstrates only α-Cusolid solution diffraction peaks and Alloy B shows a small butclearly visible diffraction peak (marked by the symbol ▼) for bcc α-Fe phase in addition to α-Cu peaks. The second-phase particlesobserved in Fig. 5(b), therefore, can be identified as bcc α-Fe phase,which is in accordance with the result of the equilibrium phasediagram by Thermo-Calc (Fig. 1).

To investigate the characteristic of the second-phase particles, wehave further examined them by TEM observation. Fig. 7 is (a) atypical TEM micrograph and (b) the interfaces between the Feprecipitate and Cu matrix in high resolution transmission electronmicroscopy (HRTEM) for peak aged Alloy B. It is seen that fineprecipitates with diameters ranging from 20 to 60 nm show lobecontrast and are relatively uniformly distributed in the Cu matrix.The interplanar distance in the precipitate is 0.200 nm, which is closeto bcc Fe plane (110) with a standard d-spacing of 0.202 nm [17]. Theinterplanar distance for Cu matrix of 0.184 nm corresponds to the dvalue of 0.181 nm for (002) plane for fcc copper [18]. Furthermore, asshown in Fig. 7(b), the interfacial relationship between Fe coherentnanoprecipitate and copper matrix is determined as ð002ÞM==ð110ÞP .

0 100 200 300 400 500 6000

20

40

60

80

100

FCC-Cu solid solution FCC-δ BCC-Fe

Mas

s fra

ctio

n (%

)

Temperature (°C)

Fig. 1. The equilibrium phase diagram of Cu–10 wt% Sn–2 wt% Zn–1.5 wt% Fe alloycalculated by Thermo-Calc.

Y. Ye et al. / Materials Science & Engineering A 612 (2014) 246–252 247

There is, however, some disagreement concerning the crystal struc-ture of Fe precipitates evolved in Cu matrix; e.g. Fe precipitates showthe same fcc lattice structure as that of Cu matrix [19,20] or the bcclattice structure of pure Fe [12,21–23]. The latter result is consistentwith the present investigation.

3.3. Tensile strength and ductility

Fig. 8 shows engineering stress–strain curves of the two Cualloys in the as-cast and peak-aged states. It is seen that the

strength of Alloy B, e.g. yield strength (YS) and ultimate tensilestrength (UTS), is much higher than that of Alloy A. The YS and UTSare as high as 264 MPa and 514 MPa for Alloy B and so theincrements of 122 MPa and 159 MPa compared to those of AlloyA, respectively. On the other hand, heat treatment leads to sig-nificant increase of fracture elongations (FE) for both the alloys. TheFEs reach as high as around 50%, which is almost similar to that ofconventional pure copper. This can be attributed to the develop-ment of homogenized microstructures and the disappearance ofbrittle δ phase, resulting in such a remarkable enhancement of

78 80 82 84 86 88 90 92 94 96 98

0

3

6

9

12

15

18

21 SnZn

Sn,

Zn

Con

tent

(at.

%)

Cu Content (at. %)78 80 82 84 86 88 90 92 94

0

3

6

9

12

15

18

21 SnZnFe

Sn,

Zn,

Fe

Con

tent

(at.

%)

Cu Content (at. %)

Fig. 2. SEM-BSE images and EDS results: (a) as-cast Alloy A; (b) as-cast Alloy B; (c) contents of Sn and Zn versus Cu in Alloy A; and (d) contents of Sn, Zn and Fe versus Cu inAlloy B. Light contrast regions are located between dendrites and at grain boundaries, and dark ones within the matrix grains. The data points on the left sides (Cu contentless than 84%) are from light contrast regions, whereas the ones on the right sides (Cu content more than 90%) are from dark contrast regions.

30 40 50 60 70 80

α-Cu δ-Cu41Sn11

α-Fe

Alloy B

Inte

nsity

(a. u

.)

2θ (degree)

Alloy A

Fig. 3. X-ray diffraction patterns of as-cast specimens: (a) Alloy A and (b) Alloy B.

80

85

115

120

125

Mic

roha

rdne

ss (H

v)

Temperature (°C)

Alloy A Alloy B

600400 500

Solution treated

300

Fig. 4. Four hours isochronal aging curves for the solution treated alloys.

Y. Ye et al. / Materials Science & Engineering A 612 (2014) 246–252248

tensile ductility. It is surprising to note that the strength of Alloy Bin the peak aged condition is not only improved prominently, butalso the FE maintains a similar high level compared to that of AlloyA. This discovery has a definite advantage for practical applications.

Additionally, a SEM micrograph of the fracture surfaces of thepeak aged Alloy B is shown in Fig. 9. It is clearly seen that typicalductile fracture occurs in the higher strength Alloy B, as char-acterized by ductile dimple features. The reason for the main-tenance of high FE for the peak-aged Alloy B will be discussed inthe next section.

4. Discussion

4.1. Strengthening mechanisms in the Cu–Sn–Zn–Fe alloy

Improved strength of the peak-aged Cu–Sn–Zn–Fe alloy is nowanalyzed and the mechanisms contributing to it discussed here.Isochronal aging curves in Fig. 4 show that only the Cu–Sn–Zn–Fealloy has a significant age hardening response and the Cu–Sn–Znalloy shows no appreciable age hardening. It is found in the peak-aged Cu–Sn–Zn–Fe alloy compared to Cu–Sn–Zn alloy that highdensity of Fe nanoprecipitates is developed by XRD and TEMinvestigations and a much finer grain structure is evolved by OMimages. Therefore, the following two typical strengtheningmechanisms are discussed to operate in the peak-aged Cu–Sn–Zn–Fe alloy; i.e. precipitate strengthening and fine-grained

strengthening (Hall–Petch effect). Quantitative discussion will becarried out below by using some respective strengthening models.

4.1.1. Precipitate strengtheningA typical strengthening mechanism proposed for alloys con-

taining precipitates is usually Orowan dislocation bypassing ordislocation shearing mechanisms. Shearing mechanism can oper-ate actively for coherent precipitates with small radii, and Orowanbypass mechanism occurs when the radius of coherent precipi-tates exceeds a critical value or the precipitates become incoherent[24–26]. During operation of Orowan mechanism, dislocationsbypass precipitates, followed by leaving dislocation loops aroundthe precipitates. The yield strength increment ΔσOrowan is calcu-lated by the following equation [24–28]:

ΔσOrowan ¼M0:4Gbπ

ffiffiffiffiffiffiffiffiffiffi1�v

p lnð2r=bÞλ

ð1Þ

where M¼ 3:06 is the mean orientation factor for fcc polycrystal-line matrix [26,29]; G the shear modulus of 42.1 GPa [30], b thematrix Burgers vector of 0.25 nm [30] and Poisson's ratiov¼ 0:303[30]. The mean radius r of a circular cross-section in a randomplane for a spherical precipitate is described as [28]

r¼ffiffiffiffiffiffiffiffi2=3

pr ð2Þ

where r is the mean precipitate radius of 20 nm obtained frompresent TEM images. The edge-to-edge inter-precipitate spacing λis expressed by the formula [31]:

λ¼ 2rðffiffiffiffiffiffiffiffiffiffiπ=4f

q�1Þ ð3Þ

where f is the volume fraction of Fe precipitates for the peak agedCu–Sn–Zn–Fe alloy. From the analyzed composition and theequilibrium phase diagram, the solubility of Fe in Cu matrix at500 1C is so minimal that it can be taken as zero. Then, the fcalculated is 1.51%.

For the shearing mechanism, on the other hand, the increase inyield strength ΔσShearing results from the contributions of coherencystrengthening (Δσcs), modulus mismatch strengthening (Δσms) andorder strengthening (Δσos) [24,25,32]. The larger one of (ΔσcsþΔσms)and Δσos is the total strength increment from the dislocationshearing mechanism [24–26,33]. In coherency strengthening, thestrengthening due to misfitting coherent precipitates Δσcs can resultfrom the interaction between the strain fields of precipitates anddislocations, and be described by the following equation [29,32,34]:

Δσcs ¼ αεMðGεÞ3=2 rf0:5Gb

� �1=2

ð4Þ

where αε ¼ 2:6 for fcc metals [32], and ε is the misfit strainparameter. By far the most thoroughly modeled case of misfit strain

Fig. 5. Optical microstructures of peak-aged specimens: (a) Alloy A and (b) Alloy B.

30 40 50 60 70 80

Alloy B

α−Cu α−Fe

Inte

nsity

(a. u

.)

2θ (degree)

Alloy A

Fig. 6. X-ray diffraction patterns of peak-aged specimens: (a) Alloy A and(b) Alloy B.

Y. Ye et al. / Materials Science & Engineering A 612 (2014) 246–252 249

parameter, ε, is that of dislocations interacting elastically with aspherical coherent precipitate, given by [32]

ε¼ δ½1þ2Gð1�2vpÞ=Gpð1þvpÞ� ð5Þwhere δ¼ jap�amj=am is the misfit lattice parameter of matrix andprecipitates [32]; ap is the lattice parameter of the α-Fe precipitate,0.286 nm [35], and am is the one of copper matrix, 0.361 nm [24];Gp ¼ 80:8 GPa and vp ¼ 0:285 are respectively the shear modulusand Poisson's ratio of the α-Fe precipitate [36].

Next, the strength increment due to modulus strengtheningΔσms is caused by the mismatch between shear moduli ofprecipitate and matrix phases and expressed as [24,25,32]

Δσms ¼M0:0055ðΔGÞ3=2 2fG

� �1=2 rb

� �ð3m=2Þ�1ð6Þ

where m is a constant roughly equal to 0.85, and ΔG¼ jGp�Gj isthe modulus mismatch between the precipitates and the matrix[24,25,32].

Third, strengthening by ordered coherent precipitates occurswhen a matrix dislocation shears an ordered precipitate and formsan antiphase boundary (APB) on the slip plane of the precipitatephase. The value of strength increment Δσos is calculated by thefollowing expression [24,25,32]:

Δσos ¼ 0:81Mγapb2b

3πf8

� �1=2

ð7Þ

where γapb is the APB energy of the precipitate phase.

By inserting the relevant parameters mentioned above into Eqs.(1)–(7), the values for ΔσOrowan, Δσcs and Δσms are calculated as118, 77,220 and 362 MPa, respectively. Since the ΔσShearing is thelarger one of (ΔσcsþΔσms) and Δσos, combining the calculatedresult of ΔσOrowanoΔσcsþΔσms, it follows ΔσOrowanoΔσShearing .Here the specific value of Δσos is not calculated because the valueof γapb is not present. It is concluded, therefore, that Orowandislocation bypassing model is an operative precipitate strength-ening mechanism for the present Cu alloy, and the contribution toyield strength increment ΔσOrowan is 118 MPa.

4.1.2. Fine-grained strengthening (Hall–Petch effect)According to the Hall–Petch relationship, the yield strength

increment ΔσHP can be described as a function of grain size via thefollowing formula [24,26]:

ΔσHP ¼ KHPd�1=2 ð8Þ

where KHP is the Hall–Petch constant, and average grain diameterd is 32 μm for the peak-aged Cu–Sn–Zn-Fe alloy. KHP¼0.13 MPa m1/2 for a stir-cast Cu–11 wt% Sn alloy at room tempera-ture [37] may be reasonably used in Eq. (8) for the present alloy

(002)

(110)

0.184 nm

0.200 nm

Matrix

Precipitate

Fig. 7. Alloy B at peak hardness: (a) TEM image showing the Fe nanoprecipitates; (b) HRTEM image displaying the interface between Fe nanoprecipitate and Cu matrix. g:operating reflexion.

0 10 20 30 40 50 600

100

200

300

400

500

Eng

inee

ring

stre

ss (M

Pa)

Engineering strain (%)

Alloy A, as-cast Alloy A, peak-aged Alloy B, as-cast Alloy B, peak-aged

Fig. 8. Engineering stress–strain curves of Cu–Sn–Zn alloys in different conditions.

Fig. 9. SEM image of fracture surface from peak aged Alloy B showing ductile cups.

Y. Ye et al. / Materials Science & Engineering A 612 (2014) 246–252250

system. Thus, the increment of yield strength ΔσHP due to grainrefinement is 23 MPa.

4.1.3. Summary of the strengtheningThe overall yield strength for the peak aged Cu–Sn–Zn–Fe alloy

σy;B is estimated on the basis of the calculated results mentionedabove. Assuming that the precipitation and fine-grained strength-ening mechanisms contribute additively to the yield strength [24],the following equation has been proposed:

σy;B ¼ σy;AþΔσOrowanþΔσHP ð9Þ

where σy;A is the yield strength of Alloy A after aging at 500 1C for4 h, 142 MPa. According to Eq. (9), the yield strength for the peakaged Alloy B with a fine-grained structure containing high densityprecipitates is predicted to be 283 MPa. This value is roughlysimilar to the experiment value of 264 MPa. The slight over-estimate of the yield strength may result from an overestimationof precipitation strengthening. This may be because the volumefraction of Fe precipitates f for the peak aged Alloy B is estimatedby assuming that the solubility of Fe in Cu matrix is zero and so f isan overestimate. It is concluded, therefore, that both precipitationand grain refinement strengthening mechanisms can be operativeand reasonably estimate the yield stress increment of Cu–Sn–Zn–Fe alloy. In addition, Orowan dislocation bypassing serves as amain strengthening mechanism operating in the present Cu alloy.

4.2. Influence of work-hardening rate on the ductility

It is well-known that high strength and excellent ductility ofstructured materials are crucial for industry applications, butunfortunately ductility is usually compromised considerably withincreasing strength. For the present Cu–Sn–Zn–Fe alloy, however,the tensile ductility under the peak-aged condition is able to bepreserved at a high level with substantial strengthening in both YSand UTS, as shown in Fig. 8. Such a high level of ductility could beexplained by Considère criterion expressed as Eq. (10) whichdetermines the uniform elongation (elongation before plasticinstability) through governing the onset of localized deformation[6,38]:

∂σ∂ε

� �_ε

¼ σ ð10Þ

where σ is the true stress, ε the true strain and _ε the strain rateapplied. Hence, uniform elongation is determined by work-

hardening rate Θ expressed by Eq. (11),

Θ¼ 1σ

∂σ∂ε

� �_ε

ð11Þ

Fig. 10 shows normalized work hardening rate against truestrain and true stress of Alloys A and B in the peak-aged condition.It is seen that Alloy B offers a higher Θ value at higher stresses, andthe stress range is beyond that of Alloy A. Both the alloys exhibitpositive and similar work-hardening rate at true strains tested. Ahigh value of Θ can play an important role in sustaining uniformelongation because it helps to delay necking or localized deforma-tion under tensile deformation [6,38]. According to the Considèrecriterion, uniform elongation could be as high as 39% and 37%(true strain) for Alloys A and B in the aged condition. Hence, theappearance of high work-hardening rate Θ in the peak-aged Cu–Sn–Zn–Fe alloy is a main reason for high ductility.

High work-hardening rate Θ of the present Cu alloy can resultfrom the development of second-phase particles within matrixgrains [39]. When dislocations intersect or by-pass these second-phase particles, they facilitate dislocation accumulation, resultingin a significant work-hardening rate and consequently maintaininga high tensile ductility. Thus, it is effective to achieve substantialstrengthening while preserving high tensile ductility by introdu-cing iron nanoprecipitates into Cu–Sn–Zn–Fe alloy. Such anapproach should also be applicable to other alloy systems.

5. Conclusions

Effects of Fe addition to Cu–Sn–Zn alloy on the microstructuresand mechanical properties were investigated and the main resultsobtained are summarized as follows:

(1) The Cu–Sn–Zn–Fe alloy, aged at 500 1C for 4 h after solutiontreatment at 750 1C for 1 h, shows a peak hardness and anexcellent combination of strength and ductility; e.g. the YS,UTS and FE reach 264 MPa, 514 MPa and 50.0%, respectively.

(2) The precipitates, having diameters ranging from 20 nm to60 nm, are identified as bcc Fe phase with an interfacialrelationship of ð002ÞM==ð110ÞP .

(3) Two strengthening mechanisms, i.e. high density precipitatesand grain refinement, can operate and contribute the incre-ment of 118 MPa and 23 MPa to yield strength, respectively.

(4) High tensile ductility of the peak-aged Cu–Sn–Zn–Fe alloy canresult from its high work-hardening rate Θ, which is caused byFe nanoprecipitates that promote dislocation accumulation.

Acknowledgments

The authors gratefully acknowledge support from the NationalNatural Science Foundation of China (No. 51174234).

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