Materials Science and Engineering A 535 (2012) 22–31

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering A
journal homepage: www.elsevier.com/locate/msea

The effects of Ti and Sn alloying elements on precipitation strengthened Cu40Zn
brass using powder metallurgy and hot extrusion
Shufeng Li a,∗ , Hisashi Imai a , Haruhiko Atsumi a , Katsuyoshi Kondoh a , Akimichi Kojima b ,Yoshiharu Kosaka b , Koji Yamamoto c , Motoi Takahashi cabcJoining and Welding Research Institute, Osaka University, Japan
San-Etsu metals Co. Ltd., 1892, OHTA, Tonami, Toyama, Japan
Nippon Atomized Metal Powders Corporation, 87-16, Nishi-Sangao, Noda, Chiba, Japana r t i c l e

i n f oArticle history:
Received 20 September 2011
Received in revised form
12 December 2011
Accepted 13 December 2011
Available online 21 December 2011
Keywords:
Brass
Spark plasma sintering
Powder metallurgy
Extrusion
Precipitation strengthening
Solid solution
Mechanical property
Phase transformationa b s t r a c t
The effects of Ti and Sn alloying elements on the microstructural and mechanical properties of 60/40 brass
were studied by powder metallurgy processing. The super-saturated solid solution of Ti creates a highprecipitation reaction chemical potential in water-atomized BS40-1.0Ti brass powder. Consequently,
BS40–1.0Ti brass was remarkably strengthened by the addition of Ti. However, Ti readily segregated in
the primary particle boundaries at elevated temperatures, which detrimentally affected the mechanicalproperties of BS40–1.0Ti brass. Accordingly, Sn was proposed as an additive to BS40–0.6Sn1.0Ti to inhibit
the segregation of Ti. Consequently, the Ti precipitate was retained in the form of CuSn3 Ti5 in the interiorof grains and grain boundaries rather than in the primary particle boundaries. This result demonstrates
that the addition of Sn can effectively hinder Ti segregation in the primary particle boundaries. Sn addition
produced signiﬁcant grain reﬁnement and mechanical strengthening effects in BS40–0.6Sn1.0Ti brass. As
a result, outstanding strengthening effects were observed for BS40–0.6Sn1.0Ti sintered at 600 ◦ C, which
exhibited a yield strength of 315 MPa, an ultimate tensile strength of 598 MPa, and a Vickers microhardness of 216 Hv. These values represent increases of 27.5%, 20.1% and 45.6%, over those of extruded
BS40–1.0Ti brass.
© 2011 Elsevier B.V. All rights reserved.

1. Introduction
Brasses are by far the most commonly cast copper alloys.
Because of their excellent castability, relatively low cost, and favorable combination of strength and corrosion resistance, cast brasses
are extensively used in many applications [1–3]. The widely used
␣ + ␤ brasses exhibit the best combination of high technological andservice properties. Solid-solution strengthening and precipitation
hardening are the primary methods used for the development ofhigh-strength brasses. These effects are achieved by the introduction of alloying elements and impurities to improve the mechanicalproperties of ␣ + ␤ brasses [4–6].
The binary phase diagrams of copper alloys [7] show that Cr,
Fe, Ti, Zr, Mg, Sn, etc. can serve as candidate alloying elementsfor precipitation strengthening to develop high-strength copper
alloys because the solid solubility of these alloying elements in copper decreases sharply with decreasing temperature. Titanium is∗ Corresponding author at: 11-1, Mihogaoka, Ibaraki, Osaka 567-0047, Japan. Tel.:+81 6 6879 8669; fax: +81 6 6879 8669.
E-mail address: shufenglimail@gmail.com (S. Li).
0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved.
doi:10.1016/j.msea.2011.12.036

virtually insoluble in copper, and concentrations in excess of
0.2 wt% of Ti can induce precipitation. The phase diagram shows ﬁveintermediate phases in the Cu–Ti system [8]: Cu4 Ti, Cu3 Ti2 , Cu4 Ti3 ,CuTi and CuTi2 . The existence of such a large number of phases suggests Ti as an attractive candidate element that may be expected todevelop high-strength brasses by precipitation hardening.
Powder metallurgy (P/M) techniques are more ﬂexible thancasting and forging techniques, allowing the fabrication of components that would otherwise decompose or disintegrate in the
direct alloying of fused materials [9]. Rapid solidiﬁcation (i.e., ahigh cooling rate) can lead to a super-saturated Ti solid solution
in a brass matrix. Upon heat treatment, very ﬁne dispersoids of thesecond phase are precipitated in the matrix; as a result, the precipitation hardening effect is greatly intensiﬁed. Our previous study[10,11] involved the addition of Ti to 60/40 brass, which favoredgrain reﬁnement by increasing the formation rate of nucleation
and recrystallization centers and by retarding the subsequent grain
growth by ﬁne precipitates. When dissolved in the super-saturated
␣ solid solution and the ␤ phase, Ti increases the ultimate tensile strength, the yield strength and the hardness; it also affects
the phase-transformation temperature. Unfortunately, precipitation strengthening effects in the alloys produced by the powder

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

23

Fig. 1. X-ray diffractograms of the BS40–1.0Ti and BS40–0.6Sn1.0Ti raw powder and SPS compacts that were prepared at different temperatures in the range of 400–600 ◦ C.

metallurgy route weaken continuously with increased processing
temperature because of the segregation and coarsening of the Ti
precipitates in the primary particle boundaries.
The development of systematic knowledge related to thethermo-dynamic precipitation of Ti in brasses is necessary toinvestigate the feasibility of developing new high-strength brasses
with the combination of the properties obtained from theprecipitation–hardening response of Ti to powder metallurgy processing. In this study, Sn was proposed as an additive for the
Cu40Zn–1.0 wt% Ti brass for the alloy designation and processing
method. As a conventional alloying additive for copper alloys, Sn
is completely soluble in copper. It also forms solid solutions withcopper that can signiﬁcantly increase the strength and corrosion
resistance of the resulting alloys. Alternatively, the interactions of
the binary Cu–Ti [12–14], Ti–Sn [15] and Cu–Sn [16] systems aswell those of the ternary Cu–Ti–Sn system [17,18] are an interesting topic with respect to the possibility of joining materials
by the transition liquid-phase method. The introduction of Sn tobrass is expected to impede the segregation of Ti and to stabilize
the mechanical properties of Ti brass during powder metallurgy
processing at high temperatures. The microstructure, phase transformation, precipitation behavior and mechanical properties of the
60/40 brass with Ti and Sn alloying elements as pre-alloyed P/Mmaterials are investigated in detail.

2. Experimental methods
Water-atomized BS40, BS40–1.0Ti and BS40–0.6Sn1.0Ti prealloyed powders (Nippon Atomized Metal Powders, Japan) wereused as the raw materials in the as-received state. The chemical compositions of the raw powders are shown in Table 1. For
sintering, the raw powder was loaded into a cylindrical graphite
die and consolidated by spark plasma sintering (DR.Sinter/SPS1030; Sumitomo Coal Mining, Japan); sintering temperatures from400 ◦ C to 600 ◦ C were used, with an interval of 100 ◦ C. After holding at the desired temperature for 1.8 ks at pressure of 40 MPa, thesample was cooled in the chamber to less than 150 ◦ C. The resulting sintered billets with diameters of 41 mm were preheated to650 ◦ C for 1.8 ks under a nitrogen gas atmosphere. They were thenextruded using a hydraulic press with a load of 2000 kN (SHP-200450; Shibayama Machine, Japan). The ﬁnal diameter after extrusion
was 7 mm. The extruded round bar was machined into test samples
with diameters of 3 mm, in accordance with ICS 59.100.01. Tensile
tests were conducted on a universal testing machine (Autograph
AG-X 50 kN; Shimadzu, Japan) with a strain rate of 5 × 10−4 s−1 .
The strain was recorded using a CCD camera accessorized to themachine. Three samples were prepared under identical conditions
for the tensile strength tests to obtain the average value. The
micro-hardness was measured by a Vickers micro-hardness tester(HMV-2T; Shimadzu, Japan); 30 measurements were conducted toobtain an average value.
The phase compositions of the samples were identiﬁed using Xray diffraction (Labx, XRD-6100; Shimadzu, Japan). The measured
spectra were compared to the standard ICDD PDF cards available
in the system software. The microstructural evaluations of thepark plasma sintering(SPS)-compacts and extruded samples wereconducted using a ﬁeld-emission scanning electron microscope(FE–SEM, JEM-6500F; JEOL, Japan) and a transmission electron
microscope (TEM; JEM-3010, JEOL, Japan). The phases were examined using an energy-dispersive X-ray spectrometer (EDS) attachedto the SEM. The crystal orientation information, which also contained grain size and phase information, was analyzed using theelectron backscatter diffraction (EBSD) technique. EBSD was performed using a TSL (TSL DigiView IV; EDAX, USA) instrument
attached to the FE-SEM, which was operated at 20 kV. The longitudinal cross-sections of the extruded samples were mechanicallyground to a 5 ␮m silicon-carbide paper ﬁnish and vibratory polished for 30 min to remove any remaining surface damage. The
Table 1
Particle size and chemical compositions
BS40–0.6Sn1.0Ti alloy powders.
Powders

BS40
BS40–1.0Ti
BS40–0.6Sn1.0Ti

of

the

BS40,

BS40–1.0Ti

andParticle size (␮m)

wt%Median
size

Mean
size

Zn

Ti

Sn

O

Cu

248
206
256

279
250
285

40.00
41.19
40.09

–
0.99
1.03

–
–
0.65

0.05
0.23
0.19

Bal.
Bal.
Bal.

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

24

electron beam was moved in 0.2 ␮m steps over a 50 ␮m × 50 ␮m
area.
3. Results and discussion
3.1. X-ray diffraction
The X-ray diffraction data for the raw powders and the SPS compacts of BS40–1.0Ti and BS40–0.6Sn1.0Ti are presented in Fig. 1. The
raw powders were produced by the rapid solidiﬁcation method.
When cooled from the ␤ ﬁeld, the powders retained a single ␤phase state. For the BS40–1.0Ti SPS compacts prepared at 400 ◦ C,
␣ peaks appear at 32.324◦ , 49.52◦ , and 72.54◦ , which indicates thatthe ␣ phase precipitates and parts of the ␤’ phase are transformed tothe ␣ phase after exposure to thermal impact at 400 ◦ C. The phasevolume fractions of the ␣ and ␤ phases in the samples preparedat different temperatures are calculated using the XRD diffraction
peak intensities [19], as shown in the following equation:V␣ =I␣(1 1 1) + I␣(2 0 0) + I␣(2 2 0)
I␣(1 1 1) + I␣(2 0 0) + I␣(2 2 0) + I␤(1 1 0) + I␤(2 0 0) + I␤(2 1 1)

× 100%(1)

where V␣ is the ␣-phase volume fraction, I␣ is the intensity ofthe ␣ diffraction peak, and I␤ is the intensity of the ␤ diffraction
peak. The ␣-phase ratio in the sintered BS40–0.6Sn1.0Ti compact
increases from 50.5% to 52.3% when the temperature is increased
from 400 ◦ C to 500 ◦ C, then decreases to 50.5% as the temperature
is increased to 600 ◦ C. For the sintered BS40–1.0Ti compact, the␣-phase ratio shows a trend similar to that of the BS40–1.0Ti SPS
compact. The ␣-phase ratio reaches a maximum value of 55.9% forBS40–1.0Ti and 52.3% for BS40–0.6Sn1.0Ti at 500 ◦ C. The ␣-phaseratio reaches a maximum value at 500 ◦ C with respect to the Cu–Zn
binary phase diagram [7], which is consistent with that of the developed brasses in this study. The volume fraction of the ␤ phaseshows a corresponding balance variation with the volume fractionof the ␣-phase. A comparison of the materials shows that the ␣phase volume fraction in the extruded BS40 is signiﬁcantly higher:
82.6% at 400 ◦ C, 84.9% at 500 ◦ C and 80.8% at 600 ◦ C. The addition ofalloying elements changes the boundaries of the phase ﬁelds in thephase diagram as follows: an increase in the content of any additive
other than Ni shifts the ␤-phase ﬁeld toward low Zn content [20].
The shift of the maximum value of the ␣-phase ratio from 55.9%
in BS40–1.0Ti to 52.3% in BS40–0.6Sn1.0Ti veriﬁes that the addition of 0.6 wt% Sn is favorable to the formation of ␤-phase. Because
the intermetallic ␤-phase exhibits greater hardness and better hot
workability than the ␣-phase, the appropriate increment of the ␤phase volume fraction created by the addition of Ti and Sn alloying
elements is favorable to the mechanical properties of brass.
Minor peaks at 37.3◦ , 38.1◦ and 39.26◦ are also observed in
the XRD patterns of the BS40–0.6Sn1.0Ti SPS compacts sinteredat different temperatures, as shown in Fig. 1. These peaks areindexed by spacings that are characteristic of the ternary compound
CuSn3 Ti5 . Two ternary compounds, CuSn3 Ti5 and CuSnTi, and theircrystal-structure data have been reported in the Cu–Sn–Ti system[14,15,18]. The ternary compound CuSn3 Ti5 is in equilibrium with aliquid at 900 ◦ C and evolves during cooling [14]. The peak intensityof CuSn3 Ti5 increases with increased sintering temperatures, whichindicates that the volume of CuSn3 Ti5 increases with temperaturefrom 400 ◦ C to 600 ◦ C.
3.2. Microstructure
The optical micrographs of the BS40–0.6Sn1.0Ti and BS40–1.0Ti
SPS compacts sintered at different temperatures in the range of
400–600 ◦ C are shown in Fig. 2(a)–(c) and (a )–(c ), respectively. For
comparison, the BS40 compacts prepared under the same conditions are shown in Fig. 2(a )–(c ). The three raw powders preparedby water atomization consist of fundamentally equiaxial grains ofFig. 2. Optical micrographs of BS40–0.6Sn1.0Ti at (a) 400 ◦ C, (b) 500 ◦ C, (c) 600 ◦ C and BS40–1.0Ti at (a ) 400 ◦ C, (b ) 500 ◦ C, (c ) 600 ◦ C; and BS40 (a ) 400 ◦ C, (b ) 500 ◦ C, (c )600 ◦ C compacts prepared by SPS.

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

25

Fig. 3. Transmission electron micrographs of a BS40–0.6Sn1.0Ti SPS compact consolidated at 600 ◦ C: (a) bright-ﬁeld image and (b) dark-ﬁeld image.

the ␤ phase. After the samples were sintered at 400 ◦ C for 1.8 ks,
the ␣ phase was observed to precipitate along the boundaries of
the initial ␤ phase, which is surrounded by the deeper-etching ␤phase. The ␣ phase is visible as a ﬁne laminar structure (a typicalWidmanstätten structure) that emanates from the primary particle boundaries of the BS40–1.0Ti (Fig. 2(a )) and BS40 (Fig. 2(a ))compacts. The BS40–1.0Ti compacts exhibit a much ﬁner morphology than the BS40 compact. Compared to that of BS40–1.0Ti
compacts, BS40–0.6Sn1.0Ti shows notably ﬁner grains than those of
the binary alloy, as depicted in Fig. 2(a). In addition, ﬁne precipitates
are observed in both BS40–1.0Ti and BS-0.6Sn1.0Ti, distributed uniformly in their matrices. Because Sn addition causes no effective

Fig. 4. SEM micrographs of the longitudinal cross-section of the extruded samples BS40–0.6Sn1.0Ti at (a) 400 ◦ C, (b) 500 ◦ C, (c) 600 ◦ C; BS40–1.0Ti at (a ) 400 ◦ C, (b ) 500 ◦ C,(c ) 600 ◦ C; and BS40 at (a ) 400 ◦ C, (b ) 500 ◦ C, (c ) 600 ◦ C prepared by SPS.

26

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

Fig. 5. The image-quality (IQ) map of the longitudinal cross-section of the extruded BS40–0.6Sn1.0Ti sintered at (a) 400 ◦ C, (b) 500 ◦ C, and (c) 600 ◦ C by SPS.

grain reﬁnement of yellow brass [21], Ti is identiﬁed as a beneﬁcial element for grain reﬁnement. The ␣ and ␤ phases graduallygrow coarser with elevated temperature. The increased volume
fraction of the ␣ phase after sintering at 500 ◦ C accompanies thecoarser microstructure. The ﬁne precipitates in both BS40–1.0Ti
and BS40–0.6Sn1.0Ti grow coarser after sintering at 500 ◦ C. The
microstructure of the BS40–1.0Ti sintered at 600 ◦ C (Fig. 2(c ))exhibits a different morphology from that of the samples sintered
at temperatures less than 600 ◦ C. At this temperature, the ␣ phasegrows larger and develops an island shape in the ␤ matrix. The
Ti precipitate is completely concentrated in the primary particle boundaries when the temperature is increased to 600 ◦ C. In
contrast, BS40–0.6Sn1.0Ti exhibits a distinctly different morphology than BS40–1.0Ti: the Ti precipitates grow coarser, and the
majority of the precipitates remain in the interior of the particles
instead of concentrating in the primary particle boundaries after
sintering at 600 ◦ C (Fig. 2(c)). The precipitates are conﬁrmed to be
CuSn3 Ti5 IMC according to their XRD pattern, as shown in Fig. 1.
Fig. 3(a)–(b) shows the TEM micrographs of the BS40–0.6Sn1.0Ti
SPS compact sintered at 400 ◦ C. Ti precipitates with a size of 50 nm
were observed in the interiors of the grains and at the grain boundaries, as shown in the bright ﬁeld (BF) image in Fig. 3(a). These

precipitates were also discernible in the dark-ﬁeld (DF) image in
Fig. 3(b).
Fig. 4 shows micrographs of the longitudinal cross-sections of
the extruded BS40–0.6Sn1.0Ti, BS40–1.0Ti and BS40 consolidated
at different temperatures. After extrusion with a reduction of 37%,
the BS40–1.0Ti microstructure is characterized by a continuous ␣phase elongation along the extrusion direction with corrugatedboundaries and observable effects of plastic strain in the form ofparallel intersection slip bands (Fig. 4(a )–(c )). At the lower sintering temperatures, such as 400 ◦ C, the ␣ phase consists of ﬁne
globules distributed along a different band after extrusion. The Ti
precipitates assume the form of tiny second-phase particles dispersed uniformly throughout the microstructure (Fig. 4(a )). Sucha microstructure is expected to increase the mechanical propertiesof brass alloys. The morphology of the extruded sample sintered
at 600 ◦ C differs signiﬁcantly from that of the samples sintered atlower temperatures (Fig. 4(c )). The ␣ phase becomes progressively
coarse and elongates successively along the extrusion direction.
The Ti precipitates distribute along the plastic ﬂow line between the␣/␤-phase bands. This microstructure is explained by the tendency
of Ti precipitates to segregate at the primary particle boundaries as
the sintering temperature increases, which consequently weakens

Fig. 6. The grain size in the longitudinal direction of the extruded BS40–0.6Sn1.0Ti, BS40–1.0Ti and BS40 as function of sintering temperature in the range of 400–600 ◦ C.

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

27

Table 2Mechanical properties of the extruded BS40, BS40–1.0Ti and BS40–0.6Sn1.0Ti brasses prepared using SPS in the temperature range of 400–600 ◦ C.
Materials

Temp. (◦ C)

YS (MPa)

Deviation

UTS (MPa)

Deviation

Elongation (%)

Deviation

BS40

400
500
600

303.3
288.0
251.3

0.58
1.66
1.09

500.1
493.7
483.4

0.99
0.27
1.37

41.4
39.7
41.9

1.5
3.76
1.36

BS40–1.0Ti

400
500
600

390.2
292.5
247.8

12.94
12.24
11.27

617.0
530.7
498.6

12.55
32.71
34.42

26.6
26.1
26.5

2.72
4.11
2.62

BS40–0.6Sn1.0Ti

400
500
600

398.4
355.2
315.2

5.43
24.84
15.57

615.1
601.4
597.7

5.15
27.98
10.96

21.5
21.9
25.5

1.87
2.45
0.81

the constraint on the grain growth provided by the precipitates in
the SPS compacts. The Ti precipitates are redistributed with theﬂow of the large plastic deformation of the soft brass matrix by
the high extrusion pressure during the subsequent hot extrusion.
Compared to BS40–1.0Ti, the BS40–0.6Sn1.0Ti exhibits a much ﬁner
morphology (Fig. 4(a)–(c)) in which ultraﬁne Ti precipitates are
distributed uniformly throughout the structures, even at 600 ◦ C.
The ﬁner microstructure of the BS40–0.6Sn1.0Ti is attributed tothe uniform distribution of Ti precipitates in the matrix. After the
addition of Sn, the segregation of Ti is inhibited by the formationof CuSn3 Ti5 IMCs, which are retained in the interior of the particles instead of the particle boundaries at 600 ◦ C. Consequently, thegrain reﬁnement effect is greatly intensiﬁed by the Ti precipitates.
Fig. 5 presents the image quality (IQ) map of the longitudinal cross-section micrographs of the extruded BS40–0.6Sn1.0Ti
as detected by the electron backscatter diffraction (EBSD) technique. The grain boundaries are identiﬁed by lines, and the phaseboundaries are identiﬁed by colors. The light-gray area represents the ␣-phase grains, and the dark-gray area represents the␤-phase grains. After extrusion with a reduction of 37%, theBS40–0.6Sn1.0Ti microstructure is characterized by a discrete ␣phase distributed along the extrusion direction with corrugatedboundaries (Fig. 5(a)). At the lower sintering temperatures, the ␣phase consists of ﬁne globules distributed along a different band
after extrusion. Recrystallization occurs in both the ␣ phase and the␤ phase after hot extrusion, and the ␤ phase exhibits a much ﬁner
grain size than the ␣ phase. Twins can be observed in the ␣ phase,
which are a typical feature of dual brass. Grains in both the ␣ and ␤
phases grow moderately at 500 ◦ C. The Ti precipitates (black dots)
appear as tiny second particles that are not only dispersed along
the plastic ﬂow line between the ␣/␤ phase bands but also in thegrain boundaries (Fig. 5(b)). The morphology of the extrusion sample sintered at 600 ◦ C differs greatly from the morphologies of thesamples sintered at lower temperatures. The ␣ phase becomes progressively coarser and elongates successively along the extrusion
direction (Fig. 5(c)).
The grain size in the longitudinal direction of the extruded samples is plotted in Fig. 6 as a function of the sintering temperature in
the range of 400–600 ◦ C. The extruded BS40–0.6Sn1.0Ti exhibitedmuch smaller grain size than the extruded BS40–1.0Ti and BS40 atdifferent sintering temperatures. From the longitudinal grain sizeof the extruded samples detected by EBSD, the relationship curves
between the YS and the grain size (the so-called Hall–Petch relationship curves) are plotted in Fig. 7. The correlations between the
YS and the grain size for the extruded BS40–1.0Ti and BS40 show
good linear-ﬁt relationships, which indicate ideal grain growthin both BS40–1.0Ti and BS40. However, the R-squared value forBS40–0.6Sn1.0Ti shows a slight deviation. Based on Fig. 5(a) and(b), grain growth is well restrained by the pinning effect of the ﬁne
Ti precipitates in the grain boundaries at 400 ◦ C and 500 ◦ C. Accompanying the reduction of the Ti precipitate population in the grainboundaries with increasing temperature, BS40–0.6Sn1.0Ti exhibitsa quick grain growth at 600 ◦ C, as shown in Fig. 5(c). Non-ideal grain
growth occurs when grain growth is inhibited by the presence ofa second phase that causes the growth to deviate from the ideal
case. In comparison, BS40–1.0Ti exhibits a grain growth behavior
similar to that of BS40 because the super-saturated Ti in BS40–1.0Ti
segregates readily at primary particle boundaries at elevated temperatures. This segregation depresses the grain reﬁnement effect.
Based on these results, the deviation behavior of the grain growth in
BS40–0.6Sn1.0Ti veriﬁes that the addition of elemental Sn inhibits
Ti segregation at high temperatures.

3.3. Mechanical properties
Fig. 8(a) shows the true stress versus true strain curves of the
extruded BS40, BS40–1.0Ti and BS40–0.6Sn1.0Ti brasses, which areconsolidated at different temperatures in the range of 400–600 ◦ C.
A photograph of the tensile sample used in the present study ispresented in Fig. 8(a). All measurements of mechanical properties
represent those of the extruded materials. The average values of
the ultimate tensile strength ( UTS ), the yield strength ( YS ), andthe elongation are summarized in Table 2. For the samples sintered
at 400 ◦ C, the  UTS of the BS40–1.0Ti is 617 MPa, which is 23.4%
higher than that of the BS40 prepared under identical conditions.
When the sintering temperature is increased from 400 ◦ C to 600 ◦ C,
the  UTS is signiﬁcantly reduced, decreasing by 19.2% from 617 MPa
to 498 MPa, as shown in Fig. 8(b); this latter value is approximately
the same as that of BS40 at 600 ◦ C. The  UTS of the BS40–0.6Sn1.0Ti
exhibits a distinctly different tendency than that of the BS40–1.0Ti.
It has nearly the same  UTS value as the BS40–1.0Ti sintered at
400 ◦ C, and it decreases by 2.8%, from 615 MPa to 598 MPa, when
the sintering temperature is increased from 400 ◦ C to 600 ◦ C. These
results indicate that the addition of elemental Sn has a signiﬁcantly
protective effect on the  UTS of BS40–0.6Sn1.0Ti at elevated temperatures.
Fig. 8(c) shows the effect of the sintering temperature on the
0.2% yield strength ( YS ) of the three brasses. The values of  YS for
all brasses decrease gradually when the sintering temperature isincreased. BS40–1.0Ti and BS40–0.6Sn1.0Ti sintered at 400 ◦ C have
 YS values of 390 MPa and 392 MPa, which are 28.7% and 31.3%
higher than the 303 MPa value of BS40, respectively. These results
verify that the addition of Ti improves the yield strength of brass,while the addition of Sn does not. Interestingly, the behavior of
BS40–1.0Ti is different from that of BS40–0.6Sn1.0Ti prepared at
an elevated sintering temperature. The  YS of BS40–1.0Ti drops
rapidly from 390 MPa to 248 MPa, losing 36.4% of its starting
value, when the sintering temperature is increased from 400 ◦ C
to 600 ◦ C. This result indicates that BS40–1.0Ti has approximately
the same  YS value as that of BS40. In contrast, the  YS valueof BS40–0.6Sn1.0Ti exhibits a moderate reduction of 20.8%. The
decreased rate of  YS degradation at elevated temperatures isclearly attributable to the addition of Sn.

28

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

Fig. 7. The Hall–Petch relationship between the yield strength and the grain size in the longitudinal direction of the extruded BS40–0.6Sn1.0Ti, BS40–1.0Ti and BS40 atdifferent sintering temperatures in the range of 400–600 ◦ C.

Fig. 8(d) presents plots of the elongation versus sintering temperature for the three brasses. The addition of Ti adversely affects
the ductility of BS40. The elongation of BS40–1.0Ti is 35.8% lower
than that of BS40 at 400 ◦ C. An increase in the sintering temperature shows little improvement on the ductility of either BS40 orBS40–1.0Ti. With regard to BS40–0.6Sn1.0Ti, the Ti and Sn additions induce a strong adverse effect on the ductility of BS40, whichshows a 48.2% lower value than that of BS40 at 400 ◦ C. However,
the elongation of BS40–0.6Sn1.0Ti improves with elevated temperature. It exhibits approximately the same value as BS40–1.0Ti

Fig. 8. Mechanical properties of the extruded BS40, BS40–1.0Ti and BS40–0.6Sn1.0Ti brasses; the billets are prepared by SPS at different temperatures in the range of
400–600 ◦ C. (a) The true stress–strain curves of the extruded brasses. The inset shows a photograph of the tensile test sample. (b) Yield strength; (c) ultimate tensile strength;and (d) elongation of the extruded brasses as a function of sintering temperature.

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

29

Fig. 9. Vickers micro-hardness of the extruded BS40, BS40–1.0Ti and BS40–0.6Sn1.0Ti brasses as a function of sintering temperature in the range of 400–600 ◦ C.

at 600 ◦ C, which is attributed to the impedance of Ti segregation atparticle boundaries at high temperatures by added Sn, as shown in
Fig. 8(c).
The decrease in the ductility after the addition of Ti and Sn is
attributed to the segregation of the alloying elements in the grainboundaries; such imperfections always embrittle a material. The
reduction of the ductile ␣-phase volume fraction is another essential factor. The ␣-phase ratio in BS40–0.6Sn1.0Ti is 52.3% at 500 ◦ C,
which is 38.4% lower than the 84.9% ratio in BS40. In addition, theelongation of BS40 decreases slightly at 500 ◦ C instead of increasing
concomitantly as the temperature increases. The ductile ␣-phasevolume of BS40–0.6Sn1.0Ti increases to its maximum of 52.3% at
500 ◦ C, as calculated using Eq. (1). This ␣-phase volume increase ispresumably attributable to the variation of the chemical composition of the ␤ phase, which exists as an intermetallic compound.
The intermetallic ␤ phase usually contains stoichiometric ratiosof 45–49% Zn at 500 ◦ C [2]. The increase of the Zn content in the␤ phase depresses its ductility because the ␣-phase volume ratio
reaches a maximum value at 500 ◦ C.
Fig. 9 shows plots of the Vickers micro-hardness versus the sintering temperature for the samples sintered at temperatures in
the range of 400–600 ◦ C. The hardness value of BS40–0.6Sn1.0Ti
is 47.1% greater than that of BS40 at 400 ◦ C; this result indicatesthe remarkable hardening effects of the addition of Ti and Sn. In
addition, the hardness of both BS40–0.6Sn1.0Ti and BS40 decreaseslightly with increasing sintering temperature. Only 5.3% of thehardness of BS40–0.6Sn1.0Ti is lost when the sintering temperature is increased from 400 ◦ C to 600 ◦ C. In contrast, the hardness of
BS40–1.0Ti decreases sharply as the temperature is increased from400 ◦ C to 500 ◦ C, and it exhibits approximately the same value asthat of BS40 at 500 ◦ C and 600 ◦ C. These results verify that the segregation of Ti precipitates at the primary particle boundaries reduces
the hardening effect imparted by the addition of Ti in BS40–1.0Ti.
The results also verify that Sn addition effectively retains the hardening effect of BS40–1.0Ti at high temperatures. The often-reported
variation in the hardness with different grain sizes is a result of thevariation in the prior dislocation densities of the samples [22]. Variation in the dislocation densities of samples with different grain
sizes may be a result of the method used for sample preparation. In solid mechanics, the material hardness is described as theresistance of a crystalline solid to deformation, which is inherently
size-dependent over a scale that ranges from a fraction of a micron
to approximately one hundred microns, in which smaller sized particles are harder [23]. In addition, the signiﬁcant increase of theharder ␤-phase volume fraction produced by the Ti and Sn alloying
elements is another essential factor that increases the hardness of
BS40–0.6Sn1.0Ti.
3.4. Discussion
Encounters between dislocations and obstacles, such as grainboundaries, are known to impede dislocation motion, which causes
the stress that is required for the deformation process to proceed
increase [24]. Dislocations may be pinned by stress ﬁeld interactions with other dislocation and solute particles or with physicalbarriers from grain boundaries and second precipitates. Rapid cooling rates retain the single ␤ phase, and the super-saturated solid
solution of 0.99 wt% Ti in the BS40–1.0Ti raw powder is formed.
A high cooling rate can lead to the formation of a further reﬁned
microstructure, extend the solid solubility and even produce metastable phases [25,26]. One source of this size dependence is thepresence of internal obstacles that restrict glide, such as incoherent
precipitates, grain boundaries and phase boundaries; the obstacle leads to dislocation pile-ups that further restrict their motion.
As a consequence, the mechanical properties increase as the scaleof the microstructure decreases. The solid solubility of Ti in brassdecreases rapidly with increasing heat-treatment temperature, asreported in our previous study [10]. Titanium addition causes alarge precipitation–hardening response in 60/40 brass because thesuper-saturation of Ti in the 60/40 brass matrix creates a highdegree of thermodynamic meta-stability. This meta-stability provides a high chemical potential for the precipitation reaction of Ti.
The super-saturated solid solution of Ti results in strengthening of
the solid-solution. Furthermore, precipitate strengthening derived
from the ﬁne precipitates causes signiﬁcant grain reﬁnement,
which enhances the mechanical strength. As a result, BS40–1.0Ti
exhibits signiﬁcantly stronger mechanical properties than BS40 at
low temperature. Unfortunately, Ti precipitates coalesce, coarsen,
and segregate at the primary particle boundaries at higher temperatures. However, the possible recovery or dynamic recrystallization

30

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31

induced by hot deformation at higher temperatures might reducethe dislocation density [27,28]. The recrystallization of the ␣ phasethat occurs during hot extrusion decreases the dislocation density, which strongly affects the mechanical behaviors of BS40–1.0Ti.
The recrystallization therefore engenders the segregation of Ti
precipitates at primary particle boundaries, which decreases themechanical properties and hardness of the alloy. Consequently, themechanical properties of BS40–1.0Ti continue to weaken as a resultof the coarsening and clearing away of Ti precipitates in the primaryparticle boundaries at elevated temperatures.
Sun et al. [29] have established a model (SMJ model) for the
nucleation of MnS particles near boundaries and dislocations inwhich all atoms in the boundaries and dislocations are presumed
to be potential nucleation positions. This model can be used tointerpret the precipitation mechanism of Ti in this study. The consumption of Cu and Zn atoms around an atom of Ti in boundaries
or dislocations engenders the formation of nuclei and growth of Ti
precipitates. The high-angle boundaries between the matrix grains
act as rapid-diffusion channels at high temperature. These channels
favor the coarsening of Ti precipitate particles in the boundaries,
which weakens the pinning effects against boundary migration.
According to the theory of precipitation hardening [30], a large
number of nuclei appears in the matrix because of the severe
driving force at the early aging stage, and the size of precipitatedparticles is small. The particles then grow with time, and the spacing between the precipitates becomes large. The precipitation andcoarsening processes of the secondary phase are accompanied by
the long-distance diffusion of the second-phase atoms. The evolution of the microstructures during the heat-treatment processcauses the hardness of the materials to increase, then to decrease.
Therefore, the mechanical properties continue to weaken as a resultof the coarsening and clearing away of Ti precipitate particles in theBS40–1.0Ti alloy.
A super-saturated solid solution of 1.03 wt% Ti in
BS40–0.6Sn1.0Ti raw powder is formed by the water-atomization
method. The difference between the atomic radii of Ti, Sn, Cu,
and Zn provides a strain ﬁeld that favors the strengthening of thesolid solution. Furthermore, rapid cooling from high temperatures
applies stress that causes dislocation motion [31], thereby creatinga high dislocation density and a high precipitation potential in theraw powder. The higher dislocation density of the microstructure
provides more sites for precipitation, which generates a thermokinetic advantage for the precipitation reaction [32], and thesuper-saturated solution of Ti in the brass matrix creates a highchemical potential for the precipitation as well. The dislocation
density increases gradually as the deformation strain increases
during the hot extrusion, which supplies more potential nucleationsites for Ti precipitation. In fact, Ti precipitates in the form of an
ultraﬁne dispersoid of CuSn3 Ti5 with the addition of Sn during sintering. The precipitates distribute in grain boundaries and provide
a pinning effect, which increases the resistance to grain boundary
slip and produces marked work-hardening effects. However, thetiny CuSn3 Ti5 precipitates induce signiﬁcant reﬁnement of the
microstructure, which hinders the movement of dislocations sothat stronger mechanical properties and higher hardness areachieved at lower sintering temperatures. The BS40–0.6Sn1.0Ti
sintered at 600 ◦ C exhibits a yield strength of 315 MPa, a tensile
strength of 598 MPa and a micro-hardness of 216 Hv. These values
are much higher than those of BS40–1.0Ti. These results show thatthe addition of Sn has remarkable positive effects on the alloy. Tin
increases the hard strength by increasing the volume fraction of theharder ␤ phase. Furthermore, it also can sustain the Ti precipitate
strengthening by retaining the Ti precipitates in the interior ofits grains and at grain boundaries in the form of CuSn3 Ti5 IMC
particles, as opposed to allowing them to segregate at the primaryparticle boundaries. For the development of high-strength brasses

to satisfy the requirements of new applications during their service
period, more detailed studies and systematic analyses must be
undertaken to elucidate the interactions between the Ti and Sn
alloying elements with respect to their precipitation behaviors,
dynamic recrystallization, and segregation mechanisms in thebrass matrix.
4. Conclusion
The addition of Ti and Sn alloying elements is found to increasethe ␤ phase volume fraction of 60/40 brass, which is favorable
to its strength and hardness properties. The results of this study
show that the super-saturated solid solution of Ti in the wateratomized 60/40 brass powder creates a high chemical potential forthe precipitation reaction of the Ti. Remarkable grain-reﬁnement
and strengthening effects are achieved by the addition of Ti
to 60/40 brass in the powder metallurgy route. However, Ti isfound to be readily segregated in the primary particle boundaries at elevated temperatures, which signiﬁcantly deteriorates themechanical properties of the BS40–1.0Ti. To inhibit the segregation behavior of the Ti in the BS40–1.0Ti, Sn is proposed as anadditive. The Ti precipitates in the form of CuSn3 Ti5 in grain boundaries, rather than segregating at the primary particle boundaries;
this result leads to signiﬁcant grain reﬁnement and mechanicalstrengthening effects on the BS40–0.6Sn1.0Ti brass. These results
show that the introduction of Sn to BS40–1.0Ti brass can effectively impede the segregation of Ti and stabilize the mechanicalproperties at high processing temperatures via powder metallurgy.
Acknowledgments
This work was supported as a project of the Japan Scienceand Technology Agency (JST). The authors sincerely thank theresearchers of the Joining and Welding Research Institute (JWRI)
at Osaka University as well as Mr. Yoshinori Muraki and Mr. Kyugo
Inui for assistance in carrying out the extrusion experiments andtensile tests.
References[1] J.R. Davis, Alloying, Understanding the Basics, ASM International, Materials
Park, OH 44073-0002, 2001.
[2] J.R. Davis, Copper and copper alloys, in: Handbook Committee, ASM International, Materials Park, OH 44073-0002, 2001.
[3] E.G. West, Copper and its Alloys, Ellis Horwood Ltd., 1982.
[4] S. Li, H. Imai, H. Atsumi, K. Kondoh, J. Alloys Compd. 493 (2010) 128–133.
[5] S.V. Smirnov, N.B. Pugacheva, A.N. Soloshenko, A.V. Tropotov, Phys. Met. Metallogr. 93 (2002) 584–593.
[6] T. Nakagawa, T. Tanaka, T. Amano, J. Mech. Working Technol. 2 (1978) 179–195.
[7] Metals Handbook, Alloy Phase Diagrams, ASM Handbook, vol. 3, 10th ed.,
1990.
[8] K.C.H. Kumar, I. Ansara, P. Wollants, L. Delaey, Z. Metallkunde 87 (1996)
666–672.
[9] S. Li, H. Imai, H. Atsumi, K. Kondoh, J. Mater. Sci. 45 (2010) 5669–5675.
[10] S. Li, H. Imai, H. Atsumi, K. Kondoh, Mater. Des. 32 (2011) 192–197.
[11] S. Li, H. Imai, A. Kojima, Y. Kosaka, K. Yamamoto, M. Takahashi, H. Atsumi, K.
Kondoh, Adv. Mater. Res. 233–235 (2011) 2732–2735.
[12] W.A. Soffa, D.E. Laughlin, Prog. Mater. Sci. 49 (2004) 347–366.
[13] A. Poter, A.W. Thompson, Scripta Metall. 18 (1984) 1185–1188.
[14] S. Nagarjuna, K. Balasubramanian, D.S. Sarma, J. Mater. Sci. 34 (1999)
2929–2942.
[15] C. Kuper, W. Peng, A. Pisch, F. Goesmann, R. Schmidht-Feetzer, Z. Metallkde 89(1998) 855.
[16] N. Saunders, A.P. Miodownik, Bull. Alloy Phase Diagrams 11 (1990) 278–287.
[17] S. Hamar-Thibault, C.H. Allibert, J. Alloys Compd. 317–318 (2001) 363–366.
[18] X. Zhang, Y. Zhan, Q. Guo, G. Zhang, J. Hu, J. Alloys Compd. 480 (2009) 382–385.
[19] G. Chu, Y.F. Cong, H.J. You, ACTA Metall. Sin. (English) 16 (2003) 489–494.
[20] N.B. Pugacheva, A.A. Pankratov, N. Yu Frolova, I.V. Kotlyarov, Russ. Metall. 3(2006) 239–248.
[21] M. Sadayappan, J.P. Thomson, M. Elboujdaini, G. Ping Gu, and M. Sahoo, MTL
Report 6 (2004) (TR-R).
[22] L.Y. Yee, C.M. Munawar, Philos. Mag. A 82 (2002) 2071–2080.
[23] L.Y. Zhang, Y.H. Jiang, Z. Ma, S.F. Shan, Y.Z. Jia, C.Z. Fan, W.K. Wang, J. Mater.
Proc. Technol. 207 (2008) 107–111.

S. Li et al. / Materials Science and Engineering A 535 (2012) 22–31[24] E.O. Hall, Proc. Phys. Soc. B 64 (1951) 747–753.
[25] H. Jones, Mater. Lett. 26 (1996) 133–136.
[26] L. Katgerman, F. Dom, Mater. Sci. Eng. A 375–377 (2004) 1212–1216.
[27] W. Mao, Z. An, Y. Li, Frontiers Mater. Sci. China 2 (2008) 233.
[28] W. Mao, Z. An, W. Guo, P. Yang, Steel Res. Int. 81 (2010) 6.

31[29] W.P. Sun, M. Militzer, J.J. Jonas, Metall. Trans. 23A (1992) 821–830.
[30] D.A. Porter, K.E. Esterling, Phase Transformations in Metals and Alloys, Van
Nostrand Reinhold Company Ltd., 1981.
[31] J.W. Hall, H.F. Rase, I & EC Fundamentals 3 (1964) 158–167.
[32] H. Fernee, J. Nairn, A. Atrens, J. Mater. Sci. 36 (2001) 5497–5510.