materials Article
Improving the Mechanical Properties of Cu-15Ni-8Sn Alloys by Addition of Titanium Chao Zhao 1 , Weiwen Zhang 1,2, *, Zhi Wang 1,2 and Datong Zhang 1,2 1
2
*
ID
, Daoxi Li 1 , Zongqiang Luo 1,2 , Chao Yang 1,2
Guangdong Key Laboratory for Processing and Forming of Advanced Metallic Materials, South China University of Technology, Guangzhou 510640, China; [email protected] (C.Z.); [email protected] (Z.W.); [email protected] (D.L.); [email protected] (Z.L.); [email protected] (C.Y.); [email protected] (D.Z.) School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou 510640, China Correspondence: [email protected]; Tel.: +86-20-8711-2022
Received: 23 July 2017; Accepted: 21 August 2017; Published: 6 September 2017
Abstract: The effect of Ti addition on the microstructure and mechanical properties of Cu-15Ni-8Sn alloys was investigated. Optical microscopy (OM), scanning electronic microscopy (SEM), and transmission electron microscopy (TEM) were used to determine grain size and distribution of the second phases in the alloys. The results indicate that the tensile properties of Cu-15Ni-8Sn alloys are improved significantly with Ti addition. Tensile elongation increased from 2.7% for the alloy without Ti to 17.9% for the alloy with 0.3% Ti, while tensile strength was maintained and even increased from 935 MPa to 1024 MPa. The improvement of the mechanical properties of Cu-15Ni-8Sn alloys by the addition of Ti is attributed to the grain refinement and suppression of discontinuous precipitation during heat treatment. Keywords: Cu-15Ni-8Sn alloy; titanium; mechanical properties; grain refinement; discontinuous precipitation
1. Introduction Cu-Ni-Sn alloys are considered as an attractive substitute for Cu-Be alloys, and have been widely applied to the manufacture of switches, connectors, and spring components in the electronic industries [1–3]. Among Cu-Ni-Sn systems, Cu-15Ni-8Sn (known as C72900) has the highest strength. Due to its excellent strength and high wear resistance, Cu-15Ni-8Sn alloy is also an ideal candidate for high-performance bearings and highly wear-resistant components for aerospace and mechanical systems [4,5]. It has been found that Cu-15Ni-8Sn alloy can reach a tensile strength in excess of 1000 MPa by suitable thermomechanical treatment [6]. However, such thermomechanical treatment generally results in disappointing low ductility [7–9]. Cribb reported that the tensile strength of the Cu-15Ni-8Sn alloy reached 1140 MPa, but the elongation was only 3% [10]. Discontinuous precipitation has been considered to be one of the main reasons for the loss of ductility, which appeared at grain boundaries in the aging process [11,12]. Many works have attempted to suppress discontinuous precipitation by optimizing thermomechanical treatment. The initial work on thermomechanical processing of Cu-Ni-Sn alloy was done by J.T Plewes [3], showing that ductility response is a consequence of competitive balance between spinodal and discontinuous transformation. Lefevre [13,14] pointed out that spinodal hardening in Cu-15Ni-8Sn alloy was limited by the discontinuous precipitation. Despite a drop of the yield strength, the ductility still remained low in the overaging process. On the other hand, it has been confirmed that the addition of trace elements can effectively suppress the discontinuous precipitation reaction. Insoluble fine particles were observed in the matrix and the Materials 2017, 10, 1038; doi:10.3390/ma10091038
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grain boundaries of the Cu-10Ni-8Sn alloys with different addition of Si, Al, or Cr. These particles occupied the nucleation sites to suppress the nucleation of the discontinuous precipitation and retarded the migration of the advancing boundaries of the cell and matrix by Zener0 s pining effect to inhibit the growth of the cell and refine the grains; meanwhile, the additions barely diminished the maximum hardness of the alloys [15–17]. In the Cu-15Ni-8Sn alloy with the addition of silicon, the nucleation and growth of discontinuous precipitation are suppressed, caused by the formation of Ni2 Si and Ni3 Si particles [18]. The addition of Nb, Ta, V, or Fe also remarkably increased ductility of the Cu-15Ni-8Sn alloy with suitable cold work and aging treatment [8,9]. It has been found that the addition of Ti also has the effect of suppressing discontinuous precipitation and grain refinement in the Cu-10Ni-8Sn alloy, owing to the formation of insoluble particles [17]. As discussed above, Cu-15Ni-8Sn alloys with high strength have an important technical application; however, their low ductility limits their wider application. It is expected that the mechanical properties of the Cu-15Ni-8Sn alloy could be improved by the suppression of discontinuous precipitation and grain refinement by the addition of Ti. Therefore, in the present work, the effect of Ti addition on the high strength Cu-15Ni-8Sn alloy was studied; the ductility of the alloys in the present work was effectively improved, maintaining an ultra-high level of strength. 2. Experimental Procedure 2.1. Materials and Processing Cu-15Ni-8Sn alloy ingots with different Ti contents were prepared by an intermediate frequency induction furnace. The chemical compositions of these alloys are listed in Table 1. The cast ingots were subsequently homogenization treated at 840 ◦ C for 8 h. The hot extrusion process was utilized to fabricate a rod with diameter of 12 mm at an extrusion ratio of 17:1. After solution treatment at 820 ◦ C for 1 h, the extruded rods were quenched into water and then aged isothermally at 400 ◦ C for 4 h. Table 1. Chemical composition of the Cu-15Ni-8Sn alloys with different Ti contents.
Alloy Designation Cu-15Ni-8Sn Cu-15Ni-8Sn-0.02Ti Cu-15Ni-8Sn-0.3Ti Cu-15Ni-8Sn-0.5Ti
wt./% Ni
Sn
Ti
Cu
15.03 15.03 15.23 14.96
8.12 8.14 8.36 8.02
0.02 0.32 0.53
balance balance balance balance
2.2. Microstructure Evaluation The microstructure of each specimen was characterized by optical microscopy (OM, Leika, Microsystems GmbH, Wetzlar, Germany), scanning electronic microscopy (SEM, FEI, Hillsboro, OR, USA), and transmission electron microscopy (TEM, FEI, Hillsboro, OR, USA). Grain structure and fracture surfaces were examined with a LEICA/DMI 5000M optical microscope (Leika, Microsystems GmbH, Wetzlar, Germany) and a FEI NONA430 scanning electron microscope. The mean grain size was obtained from 20 random areas of each specimen by the linear intercept method based on the observation of OM. The representative OM image of determination of grain size by the linear intercept method is shown in Figure 1. The specimens for microstructure observation were prepared by polishing and then etching in a solution of 5 g FeCl3 + 10 mL HCl + 100 mL H2 O. Transmission electron microscope observation was performed on a FEI TECNAI G2 S-TWIN F20 (FEI, Hillsboro, OR, USA). TEM samples were prepared by twin-jet electro-polishing method in 95% alcohol and 5% perchloric acid at −25 ◦ C.
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Figure Figure 1. 1. The The representative representative optical optical microscopy microscopy (OM) (OM) image image of of grain grain size size determination determination by by the the linear linear intercept method. intercept method.
2.3. Resistivity Measurement 2.3. Resistivity Measurement Distribution of atoms was estimated by monitoring the absolute resistance change in different Distribution of atoms was estimated by monitoring the absolute resistance change in different treatment conditions. The electrical resistivity was measured at room temperature by a PPMS-9 treatment conditions. The electrical resistivity was measured at room temperature by a PPMS-9 resistivity analyzer (Qyantum Design Inc., Design Inc., San Diego, CA, USA). The average resistivity resistivity analyzer (Qyantum Design Inc., Design Inc., San Diego, CA, USA). The average resistivity was calculated from at least five independent values for each specimen. was calculated from at least five independent values for each specimen. 2.4. 2.4. Mechanical Mechanical Properties Properties Testing Testing The with aa gauge The aged aged rods rods were were cut cut into into cylindrical cylindrical tensile tensile samples samples with gauge section section of of 55 mm mm in in diameter diameter and 25 mm in length according to Chinese GB/T 228-2002. Tensile tests were performed SANS and 25 mm in length according to Chinese GB/T 228-2002. Tensile tests were performed on on aa SANS CMT5105 material test testmachine machine(SANS, (SANS,Shenzhen, Shenzhen,Guangdong, Guangdong,China). China).Three Three replicates were used CMT5105 material replicates were used to to establish the mechanical property data. establish the mechanical property data. 3. Results and Discussion 3.1. Microstructures Figure 22indicates thethe microstructures of the with different Ti contents differentinconditions: indicates microstructures ofalloys the alloys with different Ti in contents different ◦ ◦ hot-extrusion, solid solution treatment 820 C for 1 h, °C and 400 C for 4 h.°CFor conditions: hot-extrusion, solid solutionat treatment at 820 foraged 1 h, at and aged at 400 forthe 4 h.aged For alloyaged without addition Ti, coarse andgrains discontinuous precipitation in the grainin boundaries the alloythe without theof addition of grains Ti, coarse and discontinuous precipitation the grain (arrow A in(arrow FigureA2i)inwere clearly observed. the aged with the addition 0.02% Ti, boundaries Figure 2i) were clearly For observed. Foralloys the aged alloys with the of addition of discontinuous precipitation was suppressed (Figure 2j) (Figure and no Ti-rich observed in 0.02% Ti, discontinuous precipitation was suppressed 2j) andprecipitate no Ti-richwas precipitate was solutionized alloy (Figurealloy 2f). (Figure When the content equal to or equal more than observed in solutionized 2f).TiWhen thewas Ti content was to or 0.3%, more needle-like than 0.3%, precipitates were observed (arrow C in Figure the discontinuous colony was not needle-like precipitates were observed (arrow2k) C and in Figure 2k) and the precipitation discontinuous precipitation found inwas thenot aged alloys. was found that It the needle-like precipitates were formed in the were solidification colony found in Itthe aged alloys. was found that the needle-like precipitates formed ◦ C for process. These precipitates could be precipitates retained in specimens after solution treatmentafter at 820solution in the solidification process. These could be retained in specimens 2 h (Figureat2g,h). With addition of 0.5% Tithe to the alloy, of the0.5% amount the precipitates increased treatment 820 °C for 2the h (Figure 2g,h). With addition Ti toof the alloy, the amount of the significantly,increased and a small quantity of and precipitates after aging treatment (arrow D precipitates significantly, a small showed quantitycoarsening of precipitates showed coarsening after in Figure 2l). aging treatment (arrow D in Figure 2l).
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Figure2. SEM 2. SEM micrographs of with alloys with contents differentof Ti contents of conditions: Ti in different conditions: Figure micrographs of alloys different in different (a–d) Hot-extruded ◦ ◦ (a–d) Hot-extruded state; (e–h) Solid solution treatment at 820 °C for 1 h; (i–l) Aging treatment at state; (e–h) Solid solution treatment at 820 C for 1 h; (i–l) Aging treatment at 400 C for 4 h. 400 °C for 4 h.
Energy-dispersive X-ray (EDX) results of the chemical composition of second phases in the alloys Energy-dispersive X-ray (EDX) results of the chemical composition of second phases in the are listed in Table 2. The EDX results show that the needle-like precipitates are Ni3 Ti phase (arrow C alloys are listed in Table 2. The EDX results show that the needle-like precipitates are Ni3Ti phase in Figure 2k and arrow D in Figure 2l), where the ratio of Ni and Ti is approximately 3:1, which is (arrow C in Figure 2k and arrow D in Figure 2l), where the ratio of Ni and Ti is approximately 3:1, further confirmed by the SADP. Figure 3 shows the results of TEM observation of the alloy with the which is further confirmed by the SADP. Figure 3 shows the results of TEM observation of the alloy addition of addition 0.5% Ti. Based onTi. theBased EDX data, analysis of SADP reveals thereveals Ti-rich precipitates were with the of 0.5% on the EDX data, analysis of that SADP that the Ti-rich Ti phase. determined to be Ni 3 precipitates were determined to be Ni3Ti phase. Table by the the arrow arrow A, A, B, B, C, C,and andDDin inFigure Figure2.2. Table2.2.EDX EDXanalysis analysisresults resultsfor forthe the position position indicated indicated by
at./% at./% AA BB C C D D
Cu Cu
Ni Ni
Sn Sn
Ti Ti
47.81 47.81 80.35 80.35 40.23 40.23 7.75 7.75
40.29 40.29 15.56 15.56 45.95 45.95 68.40 68.40
11.90 11.90 4.09 4.09 4.02 4.02 2.98 2.98
--9.80 9.80 20.87
20.87
Figure in different different treatment treatmentstates. states.ItItcan canbebeseen seen Figure4 4shows showsthe thechange changeof ofgrain grain size size of of the the alloys alloys in that with increasing increasingTi Ticontent, content,from from95.9 95.9µm µmtoto5.2 5.2µm. µm. thatthe thegrain grainsize sizeof ofthe theaged aged alloys alloys decreased decreased with The grain refinement effect induced by Ti addition is attributed to two factors. Firstly, the addition The grain refinement effect induced by Ti addition is attributed to two factors. Firstly, the addition ofofTiTistimulates process.In Inthe theprocess processofofhot-extrusion, hot-extrusion, stimulatesrecrystallization recrystallizationnucleation nucleation in in hot-extrusion hot-extrusion process. NiNi precipitates enhance thethe dislocation density, which promotes recrystallization nucleation [19,20] 3 Ti 3Ti precipitates enhance dislocation density, which promotes recrystallization nucleation (Figure 3a). Secondly, the growth of grains during heat treatment is inhibited by Ti addition. For the [19,20] (Figure 3a). Secondly, the growth of grains during heat treatment is inhibited by Ti addition. alloys with addition of 0.3%of Ti0.3% and Ti 0.5% Ni3Ti, Ti Ni precipitates distributed in grain boundaries can For the alloys with addition andTi, 0.5% 3Ti precipitates distributed in grain boundaries retard the grain boundary migration in the process ofofsolid (Figure3b), 3b),which which can retard the grain boundary migration in the process solidsolution solution treatment treatment (Figure decreases decreasesthe thegrain graingrowth growthrate rateremarkably. remarkably.
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Figure 3. TEM images of Ni3Ti precipitates distributed in the grain boundary in 0.5% Ti alloys: Figure 3. TEM images images of of Ni Ni33Ti Ti precipitates precipitates distributed in in the grain grain boundary in in 0.5% 0.5% Ti Ti alloys: alloys: Figure 3. TEM in 0.5% Ti (a) hot-extruded alloy; (b) solid solution alloy;distributed (c) aged alloy;the (d) SADP boundary of the Ni3Ti phase (a) hot-extruded alloy; (b) solid solution alloy; (c) aged alloy; (d) SADP of the Ni3 Ti phase in 0.5% Ti (a) hot-extruded alloy; (b)of solid aged alloy; (e) Schematic the solution SADP in alloy; (d). (c) aged alloy; (d) SADP of the Ni3Ti phase in 0.5% Ti aged alloy; (e) Schematic of the SADP in (d). aged alloy; (e) Schematic of the SADP in (d).
Figure 4. Grain size of the as-extruded and aged alloys. Figure 4. Grain size of the as-extruded and aged alloys. Figure 4. Grain size of the as-extruded and aged alloys.
The discontinuous precipitations are clearly observed at grain boundaries in the alloys without Ti and 0.02% Ti (Figures 2i and 5), which areobserved constituted by γ lamellae (DOin 3 (Cu ,Ni1-x)3Sn) and Thewith discontinuous precipitations are clearly at grain boundaries thexalloys without The discontinuous precipitations arewhich clearly observed at colony grain in the alloys withoutin Ti α lamellae The 2i EDX of Cu and Sn in the of discontinuous precipitation Tidepleted and with 0.02% Ti [21]. (Figures andmap 5), are constituted by γboundaries lamellae (DO 3 (Cu x,Ni1-x)3Sn) and and with Ti presented (Figures 2iin and 5), which constituted by lamellae (DO3 (Cu ,Ni 1-x )3 Sn) and 0.02% Tiα0.02% alloy is Figure 6. 5aSndepicts theγinitiation process ofxprecipitation discontinuous depleted lamellae [21]. The EDX map of Figure Cuare and in the colony of discontinuous in depleted α lamellae [21]. The EDX map of Cu and Sn in the colony of discontinuous precipitation precipitation grain boundary. structure of discontinuous is 0.02% Ti alloyatisthe presented in FigureThe 6. typical Figure lamellar 5a depicts the initiation process ofprecipitation discontinuous in 0.02%inTiFigure alloy is grain presented in Figure Figurelamellar 5a depicts the initiation process ofprecipitation discontinuous shown 5b. precipitation at the boundary. The6.typical structure of discontinuous is precipitation at the grain boundary. The typical lamellar structure of discontinuous precipitation is shown in Figure 5b. shown in Figure 5b.
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Figure 5. TEM images of the morphology of discontinuous precipitation in 0.02% Ti alloy aged at Figure 5. TEM TEM images images of the morphology morphology of discontinuous discontinuous precipitation precipitation in in 0.02% 0.02% Ti alloy aged at 400 ◦°C for 4 h: (a) Initiation of discontinuous precipitation in grain boundaries; (b) Typical colony of 400 °C C for 4 h: (a) Initiation of discontinuous discontinuous precipitation in grain boundaries; boundaries; (b) Typical colony of discontinuous precipitation. discontinuous precipitation.
Figure 6. Energy-dispersive X-ray (EDX) map of Cu and Sn taken from the second phase in 0.02% Ti Figure 6. Energy-dispersive X-ray (EDX) map of Cu and Sn taken from the second phase in 0.02% Ti Figure 6. Energy-dispersive alloy aged at 400 °C for 4 h. X-ray (EDX) map of Cu and Sn taken from the second phase in 0.02% Ti alloy aged at 400 °C for 4 h. alloy aged at 400 ◦ C for 4 h.
The initiation process of discontinuous precipitation of the Cu-15Ni-8Sn alloys can be illustrated The initiation process of discontinuous precipitation of the Cu-15Ni-8Sn alloys can be illustrated by the pucker mechanism In the initiation process ofof discontinuous precipitation, is necessary The initiation process [22]. of discontinuous precipitation the Cu-15Ni-8Sn alloys canit by the pucker mechanism [22]. In the initiation process of discontinuous precipitation, itbeisillustrated necessary thatthe some positions of the grain boundary are deformed shapes withprecipitation, special small angle to form by pucker mechanism [22]. In the initiation process ofinto discontinuous is necessary that some positions of the grain boundary are deformed into shapes with special small itangle to form the nucleation position of grain discontinuous precipitation (Figure 5a). In thespecial previous work [15–17], it that some positions of the boundary are deformed into shapes with small angle to form the nucleation position of discontinuous precipitation (Figure 5a). In the previous work [15–17], it wasnucleation reported position that insoluble particles occupy the sites of 5a). discontinuous precipitation, which is the of discontinuous precipitation (Figure In the previous work [15–17], it was was reported that insoluble particles occupy the sites of discontinuous precipitation, which is contributed suppression of the formation of of discontinuous precipitation. Figure 3c shows, no reported thatto particles occupy the sites discontinuous precipitation,As which is contributed to contributed toinsoluble suppression of the formation of discontinuous precipitation. As Figure 3c shows, no facet was found atformation the grain of boundary near Ni 3Ti phase inAs theFigure alloys3cwith addition of was 0.3%found Ti and suppression of the discontinuous precipitation. shows, no facet at facet was found at the grain boundary near Ni3Ti phase in the alloys with addition of 0.3% Ti and 0.5%grain Ti. Therefore, is demonstrated effect ofaddition Ni3Ti phase on grain is not onlyit to the boundary it near Ni3 Ti phase inthat the the alloys with of 0.3% Ti andboundary 0.5% Ti. Therefore, is 0.5% Ti. Therefore, it is demonstrated that the effect of Ni3Ti phase on grain boundary is not only to hinder the grain boundary migration, but also to suppress the puckering of grain boundary. The demonstrated that the effect of Ni Ti phase on grain boundary is not only to hinder the grain boundary 3 hinder the grain boundary migration, but also to suppress the puckering of grain boundary. The number of but nucleation sites of discontinuous precipitation obtainedThe by puckering grain boundary migration, also to suppress the puckering of grain boundary. of of nucleation sites of number of nucleation sites of discontinuous precipitation obtained by number puckering of grain boundary was reduced. As a result, theobtained additionby of puckering Ti had an effect on the suppression of the initiation process discontinuous precipitation of grain boundary was reduced. As a result, the was reduced. As a result, the addition of Ti had an effect on the suppression of the initiation process of discontinuous precipitation. addition of Ti had an effect on the suppression of the initiation process of discontinuous precipitation. of discontinuous precipitation. The effect of Ti addition addition on the growth growth of discontinuous discontinuous precipitation can be discussed by the The effect of Ti addition on the the growth of of discontinuous precipitation can be discussed by the rate of discontinuous discontinuous precipitation [23]. [23]. following following formula formula of of the the growth growth rate rate of discontinuous precipitation precipitation [23]. (1) G M P GG =MM×PP (1) where G corresponds to the growth rate of the cell, M is the mobility of frontier interface of the where G corresponds to the growth rate of the cell, M is the mobility of frontier interface of the where corresponds the of growth rate of the cell, M is of frontier interface the cell, the drivingto force the cell growth. In case ofthe themobility alloys with Ti addition, Ni3Tiofphase in cell, PG is P is the driving force of the cell growth. In case of the alloys with Ti addition, Ni 3Ti phase in cell, P is the driving force of the cell growth. In case of the alloys with Ti addition, Ni Ti phase in the grain 3 the grain boundary produces a back-driving force ( PP ) and decreases M due to the suppression of the grain boundary a back-driving due to the suppression of M suppression PP ) and decreases boundary producesproduces a back-driving force (PP )force and (decreases M due to the of migration migration of the frontier interface. Consequently, the nucleation and growth of discontinuous migration of the frontierConsequently, interface. Consequently, theand nucleation and growth of discontinuous of the frontier interface. the nucleation growth of discontinuous precipitation precipitation are both suppressed by the addition of Ti. In recent years, it has become widely accepted precipitation are both suppressed by the addition of Ti. In recent years, it has become widely accepted that the grain boundary “complexion” transitions could also lead to discontinuous changes in that the grain boundary “complexion” transitions could also lead to discontinuous changes in mobility and diffusivity of grain boundary [24,25]. It is reasonable to speculate that the addition of Ti mobility and diffusivity of grain boundary [24,25]. It is reasonable to speculate that the addition of Ti
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are both suppressed by the addition of Ti. In recent years, it has become widely accepted that the grain boundary “complexion” transitions could also lead to discontinuous changes in mobility and diffusivity of grain boundary [24,25]. It is reasonable to speculate that the addition of Ti elements changes grain boundary complexion, and thus intergranular discontinuous precipitation might be suppressed [26,27]. In order to clarify the solid solubility of Ti in the Cu-15Ni-8Sn alloy and the effect of Ti addition on γ phase, the distributions of Ni, Sn, and Ti atoms were determined. Table 3 lists the values of electrical conductivity used to calculate the distributions of those solute atoms in the base and 0.3% Ti alloys. E0 and E1 are the relative electrical conductivity of the alloys after and before aging at 400 ◦ C for 4 h, respectively. E2 is the relative electrical conductivity of the alloys after solution treatment at 960 ◦ C for 3 h. On the basis of the microstructure observations, γ phase was dissolved in matrix but Ni3 Ti phase still existed in 0.3% Ti alloy after solution treatment at 820 ◦ C for 1 h (Figure 2g). After solution treatment at 960 ◦ C for 3 h, Ni3 Ti phase was entirely dissolved in 0.3% Ti alloy. The change in the electrical resistivity ∆ρ of the alloys on different treatment state can be described as [28]: ∆ρ = ∆ρG.B. + ∆ρ DIS + ∆ρSOL
(2)
∆ρG.B. , ∆ρ DIS , and ∆ρSOL are the resistivity changes contributed by the changes in grain boundaries, dislocation density, and solute atom concentration in the alloys, respectively. For Cu-15Ni-8Sn alloy with addition of 0.3% Ti, the effect of the changes in grain boundary density is considered to be negligibly small, owing to little change in the grain size of the 0.3% Ti alloy. According to a previous study [29], ∆ρ DIS /∆N of copper is about 2 × 10−16 nΩm3 , which is a significantly smaller contribution to resistivity than the resistivity contribution of solute atoms of Ni, Sn, and Ti. Thus, the changes in resistivity primarily contributed to the changes in solute atom concentration. The values of relative electrical conductivity are calculated from the formula: E = ρ0 × 100/ρ
(3)
Here, ρ0 = 17.24 nΩm. In case of 0.3% Ti alloy, the difference between E2 and E1 was mainly caused by precipitation of Ni3 Ti phases, and precipitation of γ phases led to the difference between E1 and E0. The data used to calculate the atomic distributions are reported in Komatsu’s [30] study: ∆ρ Ni = 12.2 nΩm/at%, ∆ρSn = 28.8 nΩm/at%, ∆ρ Ti = 113.3 nΩm/at%. Table 3. Values of the electrical conductivity of the base and 0.3% Ti alloys. Electrical Conductivity (pct IACS 1 )
Alloy Base 0.3Ti 1
E0
E1
E2
41.20 49.72
27.29 27.19
18.57
Relative Conductivity according to International Annealed Copper Standard.
Figure 7 indicates the distributions of Ni, Sn, and Ti atoms in matrix, γ phase, and Ni3 Ti precipitation. In case of 0.3% Ti alloy, approximately 0.177% Ti was dissolved in the α (Cu) matrix, and the remaining Ti formed the Ni3 Ti phase. The solution of Ti in the matrix not only played a role in solid-solution strengthening, but also decreased solid solubility of Tin in the matrix, resulting in promotion of the precipitation of γ phase. The calculation result corresponds well with the result that no Ni3 Ti precipitates are observed by TEM observation of the aged 0.02% Ti alloy.
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Figure distributionof ofNi, Ni, Sn, Sn, and and Ti Ti atoms atoms among 3Ti Ti precipitates Figure 7. 7.The The distribution among the thematrix, matrix,γγphases, phases,and andNiNi 3 precipitates ◦ Cfor base and 0.3%Ti alloysaged agedat at400 400°C in in base and 0.3% Ti alloys for4h. 4 h.
3.2.3.2. Mechanical Properties Mechanical Properties Figure 8a shows the nominal stress–strain 0% TiTi and 0.5% TiTi alloys. It It is is demonstrated that Figure8a shows the nominal stress–straincurves curvesofof 0% and 0.5% alloys. demonstrated thethat addition of Ti had thehad effect improving the plasticity and strength of theof alloys. FigureFigure 8b depicts the addition of Ti theof effect of improving the plasticity and strength the alloys. 8b the mechanical ofalloys the aged withcontents differentofcontents of Ti, indicating that thedepicts mechanical propertiesproperties of the aged withalloys different Ti, indicating that both tensile both tensile strength and elongation were improved byof the Ti. When Ti content increased strength and elongation were improved by the addition Ti.addition When Tiofcontent increased from 0 to 0.3%, 0 to 0.3%, the alloys exhibited significantly increased from tensile from 2.7% tothe 17.9%; thefrom alloys exhibited significantly increased tensile elongation 2.7elongation to 17.9%; meanwhile, tensile meanwhile, the tensile strength increased from 935 MPa to 1024 MPa. Furthermore, the tensile strength increased from 935 MPa to 1024 MPa. Furthermore, the tensile strength was increased to strength was 1080 MPa for while the alloy 0.5elongation Ti addition while the tensile elongation 1080 MPa for theincreased alloy withto0.5 Ti addition the with tensile was slightly decreased to 17.5%. was slightly decreased to 17.5%. The increased strength is largely attributed to grain boundary strengthening and solid-solution The increased strength is largely attributed boundary strengthening and solid-solution strengthening. As shown above, the grain size of to thegrain alloys was significantly reduced by the addition strengthening. As shown above, the grain size of the alloys was significantly reduced by the addition of Ti—from 95.9 µm for the alloy without Ti to 5.2 µm for the alloy with 0.5% Ti. Therefore, the grain of Ti—from 95.9 µm for the alloy without Ti to 5.2 µm for the alloy with 0.5% Ti. Therefore, the grain boundary strengthening may play an important role on strengthening according to the Hall–Petch boundary strengthening may play an important role on strengthening according to the Hall–Petch relationship [31]. The solid-solution strengthening also helped to increase the strength. It is shown relationship [31]. The solid-solution strengthening also helped to increase the strength. It is shown that with 0.3% Ti addition, 0.177% Ti was dissolved in the matrix. that with 0.3% Ti addition, 0.177% Ti was dissolved in the matrix. The improved addition of ofTi Timainly mainlycontributed contributed grain The improvedductility ductilityof ofthe the alloys alloys with with the the addition toto thethe grain refinement and the suppression of discontinuous precipitation. As shown in Figure 9a, the morphology refinement and the suppression of discontinuous precipitation. As shown in Figure 9a, the of morphology discontinuous precipitation at the grain boundaries be highly to becometothe of discontinuous precipitation at the grain would boundaries wouldconducive be highly conducive initiation site of intergranular cracks which reduce the stability of the grain boundary [22]. Therefore, become the initiation site of intergranular cracks which reduce the stability of the grain boundary Materials 2017, 10, effect x FOR PEER REVIEW 9 of 12 the[22]. suppression of the addition of Ti on the discontinuous precipitation at the grain boundaries Therefore, the suppression effect of the addition of Ti on the discontinuous precipitation at the plays a significant in athe improvement ductility of of thethe alloys. grain boundariesrole plays significant role in of thethe improvement ductility of the alloys.
Figure 8. Mechanical properties of alloys with different contents of Ti:(a) Nominal stress–strain Figure 8. Mechanical properties of alloys with different contents of Ti: (a) Nominal stress–strain curves curves of 0% Ti and 0.5% Ti alloys; (b) Ultimate tensile strength (UTS) and tensile elongation of four of 0% Ti and 0.5% Ti alloys; (b) Ultimate tensile strength (UTS) and tensile elongation of four alloys alloys after aging at 400°Cfor 4h. after aging at 400 ◦ C for 4 h.
Figure9 shows the fracture surface morphology of the alloys aged at 400°C for 4h. In the case of the alloy without Ti addition (Figure9a), intergranular fracture dominated. The coarse grains and flat grain boundaries make the cracks easy to propagate along grain boundaries. The insert in Figure9a shows the longitudinal section of the tensile-fractured aged alloy without Ti, indicating that the colony of discontinuous precipitation was observed at the edge of the crack. Figure9b–d
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◦ for 4 h. In the case of Figure atat 400 Figure99shows showsthe thefracture fracturesurface surfacemorphology morphologyofofthe thealloys alloysaged aged 400C°C for 4 h. In the case the alloy without Ti addition (Figure 9a),9a), intergranular fracture dominated. TheThe coarse grains andand flat of the alloy without Ti addition (Figure intergranular fracture dominated. coarse grains grain boundaries make thethe cracks easy to to propagate along Figure 9a 9a flat grain boundaries make cracks easy propagate alonggrain grainboundaries. boundaries.The The insert insert in in Figure shows the longitudinal section of the tensile-fractured aged alloy without Ti, indicating that the colony shows the longitudinal section of the tensile-fractured aged alloy without Ti, indicating that the of discontinuous precipitation was observed at the edge ofedge the crack. typical colony of discontinuous precipitation was observed at the of theFigure crack. 9b–d Figuredepict 9b–d that depict that quasi-cleavage fracture with dimples and tear ridges appeared in the alloys with the addition of typical quasi-cleavage fracture with dimples and tear ridges appeared in the alloys with the addition Ti. With increase of Ti addition, the amount of dimples increased and cleavage surface decreased. of Ti. With increase of Ti addition, the amount of dimples increased and cleavage surface decreased. However, Ti precipitates precipitates are are pulled pulled out, out, the thehollows hollowsappear appearininthe the0.5% 0.5%TiTialloy. alloy. However, since since the the coarse coarse Ni Ni33Ti The decrease of elongation of the 0.5% Ti alloy is considered to be contributed by these hollows. The decrease of elongation of the 0.5% Ti alloy is considered to be contributed by these hollows.
Figure 9. SEM morphologies of the fracture surfaces of the alloys with different contents of Ti: Figure 9. SEM morphologies of the fracture surfaces of the alloys with different contents of Ti: (a) 0% Ti, (a) 0% Ti, the insert image shows the tensile fracture surface after polishing; (b) 0.02% Ti; (c) 0.3% Ti; the insert image shows the tensile fracture surface after polishing; (b) 0.02% Ti; (c) 0.3% Ti; (d) 0.5% Ti. (d) 0.5% Ti.
The of Cu-15Ni-8Sn alloys with with Ti Ti addition addition and and other other The tensile tensile strength strength and and elongation elongation values values of Cu-15Ni-8Sn alloys designed Cu-15Ni-8Sn alloys are summarized in Figure 10 [8–10]. Adding the appropriate amount designed Cu-15Ni-8Sn alloys are summarized in Figure 10 [8–10]. Adding the appropriate amount of of trace elements effectiveway waytotoimprove improvemechanical mechanicalproperties propertiesofofthe the Cu-15Ni-8Sn Cu-15Ni-8Snalloys. alloys. trace elements is isananeffective Compared of of Nb,Nb, V, Fe, of the alloy the present waswork improved Comparedwith withadditions additions V, or Fe,Ta, orthe Ta,ductility the ductility of theinalloy in thework present was effectively while maintaining an ultra-high level of strength by the addition of titanium. improved effectively while maintaining an ultra-high level of strength by the addition of titanium.
(d) 0.5%Ti.
The tensile strength and elongation values of Cu-15Ni-8Sn alloys with Ti addition and other designed Cu-15Ni-8Sn alloys are summarized in Figure10 [8–10]. Adding the appropriate amount of trace elements is an effective way to improve mechanical properties of the Cu-15Ni-8Sn alloys. Materials 2017, 10, 1038 10 of 11 Compared with additions of Nb, V, Fe, or Ta, the ductility of the alloy in the present work was improved effectively while maintaining an ultra-high level of strength by the addition of titanium.
Figure of mechanical properties of of Cu-15Ni-8Sn Cu-15Ni-8Sn alloys alloys with with the the addition addition of of Ti Ti and Figure 10. 10. Comparison Comparison of mechanical properties and other elements. other elements.
4. Conclusion Conclusions In the present and mechanical properties of present work, work,the theeffects effectsofofTiTiaddition additionononthe themicrostructure microstructure and mechanical properties Cu-15Ni-8Sn alloys were investigated. TheThe following conclusions cancan be be drawn: of Cu-15Ni-8Sn alloys were investigated. following conclusions drawn:
alloys withthe theaddition additionofofTiTipossess possessan anexcellent excellentcombination combination of of strength strength and and (1) (1)Cu-15Ni-8Sn Cu-15Ni-8Sn alloys with ductility. Compared to thetoalloy without Ti, the strength and the of theofalloy with ductility. Compared the alloy without Ti,tensile the tensile strength andelongation the elongation the alloy 0.3%Ti increased from 935MPa to 1024MPa and from 2.7% to 17.9%, respectively. with 0.3% Ti increased from 935 MPa to 1024 MPa and from 2.7% to 17.9%, respectively. (2) improvement of mechanical properties by the addition of Ti contributed to grain (2) TheThe improvement of mechanical properties by the addition of Ti contributed to grain refinement refinement and the suppression of discontinuous precipitation. Byaddition the addition Ni3Ti phase and the suppression of discontinuous precipitation. By the of Ti,ofNiTi, 3 Ti phase was was formed in the solidification process, which was undissolved in the following solution formed in the solidification process, which was undissolved in the following solution treatment. treatment. Ni3TiNi phase had a pining effect on grain boundary migration in the process of solution treatment, 3 Ti phase had a pining effect on grain boundary migration in the process of solution treatment, and thus resulted in grain refinement. Ti addition also had and thus resulted in grain refinement. Ti addition also hada asuppressive suppressiveeffect effect on on discontinuous discontinuous precipitation, owing to a reduction of nucleation sites of discontinuous precipitation. precipitation, owing to a reduction of nucleation sites of discontinuous precipitation. Acknowledgments: This work is supported by the Guangdong Natural Science Foundation for Research Team (Grant No. 2015A030312003). Acknowledgments: This work is supported by the Guangdong Natural Science Foundation for Research Team (Grant No. 2015A030312003). Author Contributions: C. Zhao, W.W. Zhang and Z.Q. Luo conceived and designed the experiments; C. Zhao Author Contributions: Zhao, W.W. Zhang Z.Q. Luo and designed the C. experiments; Zhao and and D.X. Li performedC. the experiments; C. and Zhao and Z. conceived Wang analyzed the data; Yang and C. D.T. Zhang D.X. Li performed the experiments; C. Zhao and Z. Wang analyzed the data; C. Yang and D.T. Zhang contributed contributed analysis tools; C. the Zhao wrote paper; Wang and W.W. Zhang revised the manuscript. analysis tools; C. Zhao wrote paper; Z.the Wang andZ. W.W. Zhang revised the manuscript. Conflicts of ofInterest: Interest:The Theauthors authors declare conflict of interest. founding sponsors had in the Conflicts declare no no conflict of interest. The The founding sponsors had no rolenoin role the design design of the study; in the collection, analyses, or interpretation of data; the writing ofmanuscript, the manuscript, in of the study; in the collection, analyses, or interpretation of data; in theinwriting of the and and in the decision to publish the results. the decision to publish the results.
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Improving the Mechanical Properties of Cu-15Ni-8Sn Alloys by Addition of Titanium Chao Zhao 1 , Weiwen Zhang 1,2, *, Zhi Wang 1,2 and Datong Zhang 1,2 1
2
*
ID
, Daoxi Li 1 , Zongqiang Luo 1,2 , Chao Yang 1,2
Guangdong Key Laboratory for Processing and Forming of Advanced Metallic Materials, South China University of Technology, Guangzhou 510640, China; [email protected] (C.Z.); [email protected] (Z.W.); [email protected] (D.L.); [email protected] (Z.L.); [email protected] (C.Y.); [email protected] (D.Z.) School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou 510640, China Correspondence: [email protected]; Tel.: +86-20-8711-2022
Received: 23 July 2017; Accepted: 21 August 2017; Published: 6 September 2017
Abstract: The effect of Ti addition on the microstructure and mechanical properties of Cu-15Ni-8Sn alloys was investigated. Optical microscopy (OM), scanning electronic microscopy (SEM), and transmission electron microscopy (TEM) were used to determine grain size and distribution of the second phases in the alloys. The results indicate that the tensile properties of Cu-15Ni-8Sn alloys are improved significantly with Ti addition. Tensile elongation increased from 2.7% for the alloy without Ti to 17.9% for the alloy with 0.3% Ti, while tensile strength was maintained and even increased from 935 MPa to 1024 MPa. The improvement of the mechanical properties of Cu-15Ni-8Sn alloys by the addition of Ti is attributed to the grain refinement and suppression of discontinuous precipitation during heat treatment. Keywords: Cu-15Ni-8Sn alloy; titanium; mechanical properties; grain refinement; discontinuous precipitation
1. Introduction Cu-Ni-Sn alloys are considered as an attractive substitute for Cu-Be alloys, and have been widely applied to the manufacture of switches, connectors, and spring components in the electronic industries [1–3]. Among Cu-Ni-Sn systems, Cu-15Ni-8Sn (known as C72900) has the highest strength. Due to its excellent strength and high wear resistance, Cu-15Ni-8Sn alloy is also an ideal candidate for high-performance bearings and highly wear-resistant components for aerospace and mechanical systems [4,5]. It has been found that Cu-15Ni-8Sn alloy can reach a tensile strength in excess of 1000 MPa by suitable thermomechanical treatment [6]. However, such thermomechanical treatment generally results in disappointing low ductility [7–9]. Cribb reported that the tensile strength of the Cu-15Ni-8Sn alloy reached 1140 MPa, but the elongation was only 3% [10]. Discontinuous precipitation has been considered to be one of the main reasons for the loss of ductility, which appeared at grain boundaries in the aging process [11,12]. Many works have attempted to suppress discontinuous precipitation by optimizing thermomechanical treatment. The initial work on thermomechanical processing of Cu-Ni-Sn alloy was done by J.T Plewes [3], showing that ductility response is a consequence of competitive balance between spinodal and discontinuous transformation. Lefevre [13,14] pointed out that spinodal hardening in Cu-15Ni-8Sn alloy was limited by the discontinuous precipitation. Despite a drop of the yield strength, the ductility still remained low in the overaging process. On the other hand, it has been confirmed that the addition of trace elements can effectively suppress the discontinuous precipitation reaction. Insoluble fine particles were observed in the matrix and the Materials 2017, 10, 1038; doi:10.3390/ma10091038
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grain boundaries of the Cu-10Ni-8Sn alloys with different addition of Si, Al, or Cr. These particles occupied the nucleation sites to suppress the nucleation of the discontinuous precipitation and retarded the migration of the advancing boundaries of the cell and matrix by Zener0 s pining effect to inhibit the growth of the cell and refine the grains; meanwhile, the additions barely diminished the maximum hardness of the alloys [15–17]. In the Cu-15Ni-8Sn alloy with the addition of silicon, the nucleation and growth of discontinuous precipitation are suppressed, caused by the formation of Ni2 Si and Ni3 Si particles [18]. The addition of Nb, Ta, V, or Fe also remarkably increased ductility of the Cu-15Ni-8Sn alloy with suitable cold work and aging treatment [8,9]. It has been found that the addition of Ti also has the effect of suppressing discontinuous precipitation and grain refinement in the Cu-10Ni-8Sn alloy, owing to the formation of insoluble particles [17]. As discussed above, Cu-15Ni-8Sn alloys with high strength have an important technical application; however, their low ductility limits their wider application. It is expected that the mechanical properties of the Cu-15Ni-8Sn alloy could be improved by the suppression of discontinuous precipitation and grain refinement by the addition of Ti. Therefore, in the present work, the effect of Ti addition on the high strength Cu-15Ni-8Sn alloy was studied; the ductility of the alloys in the present work was effectively improved, maintaining an ultra-high level of strength. 2. Experimental Procedure 2.1. Materials and Processing Cu-15Ni-8Sn alloy ingots with different Ti contents were prepared by an intermediate frequency induction furnace. The chemical compositions of these alloys are listed in Table 1. The cast ingots were subsequently homogenization treated at 840 ◦ C for 8 h. The hot extrusion process was utilized to fabricate a rod with diameter of 12 mm at an extrusion ratio of 17:1. After solution treatment at 820 ◦ C for 1 h, the extruded rods were quenched into water and then aged isothermally at 400 ◦ C for 4 h. Table 1. Chemical composition of the Cu-15Ni-8Sn alloys with different Ti contents.
Alloy Designation Cu-15Ni-8Sn Cu-15Ni-8Sn-0.02Ti Cu-15Ni-8Sn-0.3Ti Cu-15Ni-8Sn-0.5Ti
wt./% Ni
Sn
Ti
Cu
15.03 15.03 15.23 14.96
8.12 8.14 8.36 8.02
0.02 0.32 0.53
balance balance balance balance
2.2. Microstructure Evaluation The microstructure of each specimen was characterized by optical microscopy (OM, Leika, Microsystems GmbH, Wetzlar, Germany), scanning electronic microscopy (SEM, FEI, Hillsboro, OR, USA), and transmission electron microscopy (TEM, FEI, Hillsboro, OR, USA). Grain structure and fracture surfaces were examined with a LEICA/DMI 5000M optical microscope (Leika, Microsystems GmbH, Wetzlar, Germany) and a FEI NONA430 scanning electron microscope. The mean grain size was obtained from 20 random areas of each specimen by the linear intercept method based on the observation of OM. The representative OM image of determination of grain size by the linear intercept method is shown in Figure 1. The specimens for microstructure observation were prepared by polishing and then etching in a solution of 5 g FeCl3 + 10 mL HCl + 100 mL H2 O. Transmission electron microscope observation was performed on a FEI TECNAI G2 S-TWIN F20 (FEI, Hillsboro, OR, USA). TEM samples were prepared by twin-jet electro-polishing method in 95% alcohol and 5% perchloric acid at −25 ◦ C.
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Figure Figure 1. 1. The The representative representative optical optical microscopy microscopy (OM) (OM) image image of of grain grain size size determination determination by by the the linear linear intercept method. intercept method.
2.3. Resistivity Measurement 2.3. Resistivity Measurement Distribution of atoms was estimated by monitoring the absolute resistance change in different Distribution of atoms was estimated by monitoring the absolute resistance change in different treatment conditions. The electrical resistivity was measured at room temperature by a PPMS-9 treatment conditions. The electrical resistivity was measured at room temperature by a PPMS-9 resistivity analyzer (Qyantum Design Inc., Design Inc., San Diego, CA, USA). The average resistivity resistivity analyzer (Qyantum Design Inc., Design Inc., San Diego, CA, USA). The average resistivity was calculated from at least five independent values for each specimen. was calculated from at least five independent values for each specimen. 2.4. 2.4. Mechanical Mechanical Properties Properties Testing Testing The with aa gauge The aged aged rods rods were were cut cut into into cylindrical cylindrical tensile tensile samples samples with gauge section section of of 55 mm mm in in diameter diameter and 25 mm in length according to Chinese GB/T 228-2002. Tensile tests were performed SANS and 25 mm in length according to Chinese GB/T 228-2002. Tensile tests were performed on on aa SANS CMT5105 material test testmachine machine(SANS, (SANS,Shenzhen, Shenzhen,Guangdong, Guangdong,China). China).Three Three replicates were used CMT5105 material replicates were used to to establish the mechanical property data. establish the mechanical property data. 3. Results and Discussion 3.1. Microstructures Figure 22indicates thethe microstructures of the with different Ti contents differentinconditions: indicates microstructures ofalloys the alloys with different Ti in contents different ◦ ◦ hot-extrusion, solid solution treatment 820 C for 1 h, °C and 400 C for 4 h.°CFor conditions: hot-extrusion, solid solutionat treatment at 820 foraged 1 h, at and aged at 400 forthe 4 h.aged For alloyaged without addition Ti, coarse andgrains discontinuous precipitation in the grainin boundaries the alloythe without theof addition of grains Ti, coarse and discontinuous precipitation the grain (arrow A in(arrow FigureA2i)inwere clearly observed. the aged with the addition 0.02% Ti, boundaries Figure 2i) were clearly For observed. Foralloys the aged alloys with the of addition of discontinuous precipitation was suppressed (Figure 2j) (Figure and no Ti-rich observed in 0.02% Ti, discontinuous precipitation was suppressed 2j) andprecipitate no Ti-richwas precipitate was solutionized alloy (Figurealloy 2f). (Figure When the content equal to or equal more than observed in solutionized 2f).TiWhen thewas Ti content was to or 0.3%, more needle-like than 0.3%, precipitates were observed (arrow C in Figure the discontinuous colony was not needle-like precipitates were observed (arrow2k) C and in Figure 2k) and the precipitation discontinuous precipitation found inwas thenot aged alloys. was found that It the needle-like precipitates were formed in the were solidification colony found in Itthe aged alloys. was found that the needle-like precipitates formed ◦ C for process. These precipitates could be precipitates retained in specimens after solution treatmentafter at 820solution in the solidification process. These could be retained in specimens 2 h (Figureat2g,h). With addition of 0.5% Tithe to the alloy, of the0.5% amount the precipitates increased treatment 820 °C for 2the h (Figure 2g,h). With addition Ti toof the alloy, the amount of the significantly,increased and a small quantity of and precipitates after aging treatment (arrow D precipitates significantly, a small showed quantitycoarsening of precipitates showed coarsening after in Figure 2l). aging treatment (arrow D in Figure 2l).
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Figure2. SEM 2. SEM micrographs of with alloys with contents differentof Ti contents of conditions: Ti in different conditions: Figure micrographs of alloys different in different (a–d) Hot-extruded ◦ ◦ (a–d) Hot-extruded state; (e–h) Solid solution treatment at 820 °C for 1 h; (i–l) Aging treatment at state; (e–h) Solid solution treatment at 820 C for 1 h; (i–l) Aging treatment at 400 C for 4 h. 400 °C for 4 h.
Energy-dispersive X-ray (EDX) results of the chemical composition of second phases in the alloys Energy-dispersive X-ray (EDX) results of the chemical composition of second phases in the are listed in Table 2. The EDX results show that the needle-like precipitates are Ni3 Ti phase (arrow C alloys are listed in Table 2. The EDX results show that the needle-like precipitates are Ni3Ti phase in Figure 2k and arrow D in Figure 2l), where the ratio of Ni and Ti is approximately 3:1, which is (arrow C in Figure 2k and arrow D in Figure 2l), where the ratio of Ni and Ti is approximately 3:1, further confirmed by the SADP. Figure 3 shows the results of TEM observation of the alloy with the which is further confirmed by the SADP. Figure 3 shows the results of TEM observation of the alloy addition of addition 0.5% Ti. Based onTi. theBased EDX data, analysis of SADP reveals thereveals Ti-rich precipitates were with the of 0.5% on the EDX data, analysis of that SADP that the Ti-rich Ti phase. determined to be Ni 3 precipitates were determined to be Ni3Ti phase. Table by the the arrow arrow A, A, B, B, C, C,and andDDin inFigure Figure2.2. Table2.2.EDX EDXanalysis analysisresults resultsfor forthe the position position indicated indicated by
at./% at./% AA BB C C D D
Cu Cu
Ni Ni
Sn Sn
Ti Ti
47.81 47.81 80.35 80.35 40.23 40.23 7.75 7.75
40.29 40.29 15.56 15.56 45.95 45.95 68.40 68.40
11.90 11.90 4.09 4.09 4.02 4.02 2.98 2.98
--9.80 9.80 20.87
20.87
Figure in different different treatment treatmentstates. states.ItItcan canbebeseen seen Figure4 4shows showsthe thechange changeof ofgrain grain size size of of the the alloys alloys in that with increasing increasingTi Ticontent, content,from from95.9 95.9µm µmtoto5.2 5.2µm. µm. thatthe thegrain grainsize sizeof ofthe theaged aged alloys alloys decreased decreased with The grain refinement effect induced by Ti addition is attributed to two factors. Firstly, the addition The grain refinement effect induced by Ti addition is attributed to two factors. Firstly, the addition ofofTiTistimulates process.In Inthe theprocess processofofhot-extrusion, hot-extrusion, stimulatesrecrystallization recrystallizationnucleation nucleation in in hot-extrusion hot-extrusion process. NiNi precipitates enhance thethe dislocation density, which promotes recrystallization nucleation [19,20] 3 Ti 3Ti precipitates enhance dislocation density, which promotes recrystallization nucleation (Figure 3a). Secondly, the growth of grains during heat treatment is inhibited by Ti addition. For the [19,20] (Figure 3a). Secondly, the growth of grains during heat treatment is inhibited by Ti addition. alloys with addition of 0.3%of Ti0.3% and Ti 0.5% Ni3Ti, Ti Ni precipitates distributed in grain boundaries can For the alloys with addition andTi, 0.5% 3Ti precipitates distributed in grain boundaries retard the grain boundary migration in the process ofofsolid (Figure3b), 3b),which which can retard the grain boundary migration in the process solidsolution solution treatment treatment (Figure decreases decreasesthe thegrain graingrowth growthrate rateremarkably. remarkably.
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Figure 3. TEM images of Ni3Ti precipitates distributed in the grain boundary in 0.5% Ti alloys: Figure 3. TEM images images of of Ni Ni33Ti Ti precipitates precipitates distributed in in the grain grain boundary in in 0.5% 0.5% Ti Ti alloys: alloys: Figure 3. TEM in 0.5% Ti (a) hot-extruded alloy; (b) solid solution alloy;distributed (c) aged alloy;the (d) SADP boundary of the Ni3Ti phase (a) hot-extruded alloy; (b) solid solution alloy; (c) aged alloy; (d) SADP of the Ni3 Ti phase in 0.5% Ti (a) hot-extruded alloy; (b)of solid aged alloy; (e) Schematic the solution SADP in alloy; (d). (c) aged alloy; (d) SADP of the Ni3Ti phase in 0.5% Ti aged alloy; (e) Schematic of the SADP in (d). aged alloy; (e) Schematic of the SADP in (d).
Figure 4. Grain size of the as-extruded and aged alloys. Figure 4. Grain size of the as-extruded and aged alloys. Figure 4. Grain size of the as-extruded and aged alloys.
The discontinuous precipitations are clearly observed at grain boundaries in the alloys without Ti and 0.02% Ti (Figures 2i and 5), which areobserved constituted by γ lamellae (DOin 3 (Cu ,Ni1-x)3Sn) and Thewith discontinuous precipitations are clearly at grain boundaries thexalloys without The discontinuous precipitations arewhich clearly observed at colony grain in the alloys withoutin Ti α lamellae The 2i EDX of Cu and Sn in the of discontinuous precipitation Tidepleted and with 0.02% Ti [21]. (Figures andmap 5), are constituted by γboundaries lamellae (DO 3 (Cu x,Ni1-x)3Sn) and and with Ti presented (Figures 2iin and 5), which constituted by lamellae (DO3 (Cu ,Ni 1-x )3 Sn) and 0.02% Tiα0.02% alloy is Figure 6. 5aSndepicts theγinitiation process ofxprecipitation discontinuous depleted lamellae [21]. The EDX map of Figure Cuare and in the colony of discontinuous in depleted α lamellae [21]. The EDX map of Cu and Sn in the colony of discontinuous precipitation precipitation grain boundary. structure of discontinuous is 0.02% Ti alloyatisthe presented in FigureThe 6. typical Figure lamellar 5a depicts the initiation process ofprecipitation discontinuous in 0.02%inTiFigure alloy is grain presented in Figure Figurelamellar 5a depicts the initiation process ofprecipitation discontinuous shown 5b. precipitation at the boundary. The6.typical structure of discontinuous is precipitation at the grain boundary. The typical lamellar structure of discontinuous precipitation is shown in Figure 5b. shown in Figure 5b.
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Figure 5. TEM images of the morphology of discontinuous precipitation in 0.02% Ti alloy aged at Figure 5. TEM TEM images images of the morphology morphology of discontinuous discontinuous precipitation precipitation in in 0.02% 0.02% Ti alloy aged at 400 ◦°C for 4 h: (a) Initiation of discontinuous precipitation in grain boundaries; (b) Typical colony of 400 °C C for 4 h: (a) Initiation of discontinuous discontinuous precipitation in grain boundaries; boundaries; (b) Typical colony of discontinuous precipitation. discontinuous precipitation.
Figure 6. Energy-dispersive X-ray (EDX) map of Cu and Sn taken from the second phase in 0.02% Ti Figure 6. Energy-dispersive X-ray (EDX) map of Cu and Sn taken from the second phase in 0.02% Ti Figure 6. Energy-dispersive alloy aged at 400 °C for 4 h. X-ray (EDX) map of Cu and Sn taken from the second phase in 0.02% Ti alloy aged at 400 °C for 4 h. alloy aged at 400 ◦ C for 4 h.
The initiation process of discontinuous precipitation of the Cu-15Ni-8Sn alloys can be illustrated The initiation process of discontinuous precipitation of the Cu-15Ni-8Sn alloys can be illustrated by the pucker mechanism In the initiation process ofof discontinuous precipitation, is necessary The initiation process [22]. of discontinuous precipitation the Cu-15Ni-8Sn alloys canit by the pucker mechanism [22]. In the initiation process of discontinuous precipitation, itbeisillustrated necessary thatthe some positions of the grain boundary are deformed shapes withprecipitation, special small angle to form by pucker mechanism [22]. In the initiation process ofinto discontinuous is necessary that some positions of the grain boundary are deformed into shapes with special small itangle to form the nucleation position of grain discontinuous precipitation (Figure 5a). In thespecial previous work [15–17], it that some positions of the boundary are deformed into shapes with small angle to form the nucleation position of discontinuous precipitation (Figure 5a). In the previous work [15–17], it wasnucleation reported position that insoluble particles occupy the sites of 5a). discontinuous precipitation, which is the of discontinuous precipitation (Figure In the previous work [15–17], it was was reported that insoluble particles occupy the sites of discontinuous precipitation, which is contributed suppression of the formation of of discontinuous precipitation. Figure 3c shows, no reported thatto particles occupy the sites discontinuous precipitation,As which is contributed to contributed toinsoluble suppression of the formation of discontinuous precipitation. As Figure 3c shows, no facet was found atformation the grain of boundary near Ni 3Ti phase inAs theFigure alloys3cwith addition of was 0.3%found Ti and suppression of the discontinuous precipitation. shows, no facet at facet was found at the grain boundary near Ni3Ti phase in the alloys with addition of 0.3% Ti and 0.5%grain Ti. Therefore, is demonstrated effect ofaddition Ni3Ti phase on grain is not onlyit to the boundary it near Ni3 Ti phase inthat the the alloys with of 0.3% Ti andboundary 0.5% Ti. Therefore, is 0.5% Ti. Therefore, it is demonstrated that the effect of Ni3Ti phase on grain boundary is not only to hinder the grain boundary migration, but also to suppress the puckering of grain boundary. The demonstrated that the effect of Ni Ti phase on grain boundary is not only to hinder the grain boundary 3 hinder the grain boundary migration, but also to suppress the puckering of grain boundary. The number of but nucleation sites of discontinuous precipitation obtainedThe by puckering grain boundary migration, also to suppress the puckering of grain boundary. of of nucleation sites of number of nucleation sites of discontinuous precipitation obtained by number puckering of grain boundary was reduced. As a result, theobtained additionby of puckering Ti had an effect on the suppression of the initiation process discontinuous precipitation of grain boundary was reduced. As a result, the was reduced. As a result, the addition of Ti had an effect on the suppression of the initiation process of discontinuous precipitation. addition of Ti had an effect on the suppression of the initiation process of discontinuous precipitation. of discontinuous precipitation. The effect of Ti addition addition on the growth growth of discontinuous discontinuous precipitation can be discussed by the The effect of Ti addition on the the growth of of discontinuous precipitation can be discussed by the rate of discontinuous discontinuous precipitation [23]. [23]. following following formula formula of of the the growth growth rate rate of discontinuous precipitation precipitation [23]. (1) G M P GG =MM×PP (1) where G corresponds to the growth rate of the cell, M is the mobility of frontier interface of the where G corresponds to the growth rate of the cell, M is the mobility of frontier interface of the where corresponds the of growth rate of the cell, M is of frontier interface the cell, the drivingto force the cell growth. In case ofthe themobility alloys with Ti addition, Ni3Tiofphase in cell, PG is P is the driving force of the cell growth. In case of the alloys with Ti addition, Ni 3Ti phase in cell, P is the driving force of the cell growth. In case of the alloys with Ti addition, Ni Ti phase in the grain 3 the grain boundary produces a back-driving force ( PP ) and decreases M due to the suppression of the grain boundary a back-driving due to the suppression of M suppression PP ) and decreases boundary producesproduces a back-driving force (PP )force and (decreases M due to the of migration migration of the frontier interface. Consequently, the nucleation and growth of discontinuous migration of the frontierConsequently, interface. Consequently, theand nucleation and growth of discontinuous of the frontier interface. the nucleation growth of discontinuous precipitation precipitation are both suppressed by the addition of Ti. In recent years, it has become widely accepted precipitation are both suppressed by the addition of Ti. In recent years, it has become widely accepted that the grain boundary “complexion” transitions could also lead to discontinuous changes in that the grain boundary “complexion” transitions could also lead to discontinuous changes in mobility and diffusivity of grain boundary [24,25]. It is reasonable to speculate that the addition of Ti mobility and diffusivity of grain boundary [24,25]. It is reasonable to speculate that the addition of Ti
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are both suppressed by the addition of Ti. In recent years, it has become widely accepted that the grain boundary “complexion” transitions could also lead to discontinuous changes in mobility and diffusivity of grain boundary [24,25]. It is reasonable to speculate that the addition of Ti elements changes grain boundary complexion, and thus intergranular discontinuous precipitation might be suppressed [26,27]. In order to clarify the solid solubility of Ti in the Cu-15Ni-8Sn alloy and the effect of Ti addition on γ phase, the distributions of Ni, Sn, and Ti atoms were determined. Table 3 lists the values of electrical conductivity used to calculate the distributions of those solute atoms in the base and 0.3% Ti alloys. E0 and E1 are the relative electrical conductivity of the alloys after and before aging at 400 ◦ C for 4 h, respectively. E2 is the relative electrical conductivity of the alloys after solution treatment at 960 ◦ C for 3 h. On the basis of the microstructure observations, γ phase was dissolved in matrix but Ni3 Ti phase still existed in 0.3% Ti alloy after solution treatment at 820 ◦ C for 1 h (Figure 2g). After solution treatment at 960 ◦ C for 3 h, Ni3 Ti phase was entirely dissolved in 0.3% Ti alloy. The change in the electrical resistivity ∆ρ of the alloys on different treatment state can be described as [28]: ∆ρ = ∆ρG.B. + ∆ρ DIS + ∆ρSOL
(2)
∆ρG.B. , ∆ρ DIS , and ∆ρSOL are the resistivity changes contributed by the changes in grain boundaries, dislocation density, and solute atom concentration in the alloys, respectively. For Cu-15Ni-8Sn alloy with addition of 0.3% Ti, the effect of the changes in grain boundary density is considered to be negligibly small, owing to little change in the grain size of the 0.3% Ti alloy. According to a previous study [29], ∆ρ DIS /∆N of copper is about 2 × 10−16 nΩm3 , which is a significantly smaller contribution to resistivity than the resistivity contribution of solute atoms of Ni, Sn, and Ti. Thus, the changes in resistivity primarily contributed to the changes in solute atom concentration. The values of relative electrical conductivity are calculated from the formula: E = ρ0 × 100/ρ
(3)
Here, ρ0 = 17.24 nΩm. In case of 0.3% Ti alloy, the difference between E2 and E1 was mainly caused by precipitation of Ni3 Ti phases, and precipitation of γ phases led to the difference between E1 and E0. The data used to calculate the atomic distributions are reported in Komatsu’s [30] study: ∆ρ Ni = 12.2 nΩm/at%, ∆ρSn = 28.8 nΩm/at%, ∆ρ Ti = 113.3 nΩm/at%. Table 3. Values of the electrical conductivity of the base and 0.3% Ti alloys. Electrical Conductivity (pct IACS 1 )
Alloy Base 0.3Ti 1
E0
E1
E2
41.20 49.72
27.29 27.19
18.57
Relative Conductivity according to International Annealed Copper Standard.
Figure 7 indicates the distributions of Ni, Sn, and Ti atoms in matrix, γ phase, and Ni3 Ti precipitation. In case of 0.3% Ti alloy, approximately 0.177% Ti was dissolved in the α (Cu) matrix, and the remaining Ti formed the Ni3 Ti phase. The solution of Ti in the matrix not only played a role in solid-solution strengthening, but also decreased solid solubility of Tin in the matrix, resulting in promotion of the precipitation of γ phase. The calculation result corresponds well with the result that no Ni3 Ti precipitates are observed by TEM observation of the aged 0.02% Ti alloy.
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Figure distributionof ofNi, Ni, Sn, Sn, and and Ti Ti atoms atoms among 3Ti Ti precipitates Figure 7. 7.The The distribution among the thematrix, matrix,γγphases, phases,and andNiNi 3 precipitates ◦ Cfor base and 0.3%Ti alloysaged agedat at400 400°C in in base and 0.3% Ti alloys for4h. 4 h.
3.2.3.2. Mechanical Properties Mechanical Properties Figure 8a shows the nominal stress–strain 0% TiTi and 0.5% TiTi alloys. It It is is demonstrated that Figure8a shows the nominal stress–straincurves curvesofof 0% and 0.5% alloys. demonstrated thethat addition of Ti had thehad effect improving the plasticity and strength of theof alloys. FigureFigure 8b depicts the addition of Ti theof effect of improving the plasticity and strength the alloys. 8b the mechanical ofalloys the aged withcontents differentofcontents of Ti, indicating that thedepicts mechanical propertiesproperties of the aged withalloys different Ti, indicating that both tensile both tensile strength and elongation were improved byof the Ti. When Ti content increased strength and elongation were improved by the addition Ti.addition When Tiofcontent increased from 0 to 0.3%, 0 to 0.3%, the alloys exhibited significantly increased from tensile from 2.7% tothe 17.9%; thefrom alloys exhibited significantly increased tensile elongation 2.7elongation to 17.9%; meanwhile, tensile meanwhile, the tensile strength increased from 935 MPa to 1024 MPa. Furthermore, the tensile strength increased from 935 MPa to 1024 MPa. Furthermore, the tensile strength was increased to strength was 1080 MPa for while the alloy 0.5elongation Ti addition while the tensile elongation 1080 MPa for theincreased alloy withto0.5 Ti addition the with tensile was slightly decreased to 17.5%. was slightly decreased to 17.5%. The increased strength is largely attributed to grain boundary strengthening and solid-solution The increased strength is largely attributed boundary strengthening and solid-solution strengthening. As shown above, the grain size of to thegrain alloys was significantly reduced by the addition strengthening. As shown above, the grain size of the alloys was significantly reduced by the addition of Ti—from 95.9 µm for the alloy without Ti to 5.2 µm for the alloy with 0.5% Ti. Therefore, the grain of Ti—from 95.9 µm for the alloy without Ti to 5.2 µm for the alloy with 0.5% Ti. Therefore, the grain boundary strengthening may play an important role on strengthening according to the Hall–Petch boundary strengthening may play an important role on strengthening according to the Hall–Petch relationship [31]. The solid-solution strengthening also helped to increase the strength. It is shown relationship [31]. The solid-solution strengthening also helped to increase the strength. It is shown that with 0.3% Ti addition, 0.177% Ti was dissolved in the matrix. that with 0.3% Ti addition, 0.177% Ti was dissolved in the matrix. The improved addition of ofTi Timainly mainlycontributed contributed grain The improvedductility ductilityof ofthe the alloys alloys with with the the addition toto thethe grain refinement and the suppression of discontinuous precipitation. As shown in Figure 9a, the morphology refinement and the suppression of discontinuous precipitation. As shown in Figure 9a, the of morphology discontinuous precipitation at the grain boundaries be highly to becometothe of discontinuous precipitation at the grain would boundaries wouldconducive be highly conducive initiation site of intergranular cracks which reduce the stability of the grain boundary [22]. Therefore, become the initiation site of intergranular cracks which reduce the stability of the grain boundary Materials 2017, 10, effect x FOR PEER REVIEW 9 of 12 the[22]. suppression of the addition of Ti on the discontinuous precipitation at the grain boundaries Therefore, the suppression effect of the addition of Ti on the discontinuous precipitation at the plays a significant in athe improvement ductility of of thethe alloys. grain boundariesrole plays significant role in of thethe improvement ductility of the alloys.
Figure 8. Mechanical properties of alloys with different contents of Ti:(a) Nominal stress–strain Figure 8. Mechanical properties of alloys with different contents of Ti: (a) Nominal stress–strain curves curves of 0% Ti and 0.5% Ti alloys; (b) Ultimate tensile strength (UTS) and tensile elongation of four of 0% Ti and 0.5% Ti alloys; (b) Ultimate tensile strength (UTS) and tensile elongation of four alloys alloys after aging at 400°Cfor 4h. after aging at 400 ◦ C for 4 h.
Figure9 shows the fracture surface morphology of the alloys aged at 400°C for 4h. In the case of the alloy without Ti addition (Figure9a), intergranular fracture dominated. The coarse grains and flat grain boundaries make the cracks easy to propagate along grain boundaries. The insert in Figure9a shows the longitudinal section of the tensile-fractured aged alloy without Ti, indicating that the colony of discontinuous precipitation was observed at the edge of the crack. Figure9b–d
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◦ for 4 h. In the case of Figure atat 400 Figure99shows showsthe thefracture fracturesurface surfacemorphology morphologyofofthe thealloys alloysaged aged 400C°C for 4 h. In the case the alloy without Ti addition (Figure 9a),9a), intergranular fracture dominated. TheThe coarse grains andand flat of the alloy without Ti addition (Figure intergranular fracture dominated. coarse grains grain boundaries make thethe cracks easy to to propagate along Figure 9a 9a flat grain boundaries make cracks easy propagate alonggrain grainboundaries. boundaries.The The insert insert in in Figure shows the longitudinal section of the tensile-fractured aged alloy without Ti, indicating that the colony shows the longitudinal section of the tensile-fractured aged alloy without Ti, indicating that the of discontinuous precipitation was observed at the edge ofedge the crack. typical colony of discontinuous precipitation was observed at the of theFigure crack. 9b–d Figuredepict 9b–d that depict that quasi-cleavage fracture with dimples and tear ridges appeared in the alloys with the addition of typical quasi-cleavage fracture with dimples and tear ridges appeared in the alloys with the addition Ti. With increase of Ti addition, the amount of dimples increased and cleavage surface decreased. of Ti. With increase of Ti addition, the amount of dimples increased and cleavage surface decreased. However, Ti precipitates precipitates are are pulled pulled out, out, the thehollows hollowsappear appearininthe the0.5% 0.5%TiTialloy. alloy. However, since since the the coarse coarse Ni Ni33Ti The decrease of elongation of the 0.5% Ti alloy is considered to be contributed by these hollows. The decrease of elongation of the 0.5% Ti alloy is considered to be contributed by these hollows.
Figure 9. SEM morphologies of the fracture surfaces of the alloys with different contents of Ti: Figure 9. SEM morphologies of the fracture surfaces of the alloys with different contents of Ti: (a) 0% Ti, (a) 0% Ti, the insert image shows the tensile fracture surface after polishing; (b) 0.02% Ti; (c) 0.3% Ti; the insert image shows the tensile fracture surface after polishing; (b) 0.02% Ti; (c) 0.3% Ti; (d) 0.5% Ti. (d) 0.5% Ti.
The of Cu-15Ni-8Sn alloys with with Ti Ti addition addition and and other other The tensile tensile strength strength and and elongation elongation values values of Cu-15Ni-8Sn alloys designed Cu-15Ni-8Sn alloys are summarized in Figure 10 [8–10]. Adding the appropriate amount designed Cu-15Ni-8Sn alloys are summarized in Figure 10 [8–10]. Adding the appropriate amount of of trace elements effectiveway waytotoimprove improvemechanical mechanicalproperties propertiesofofthe the Cu-15Ni-8Sn Cu-15Ni-8Snalloys. alloys. trace elements is isananeffective Compared of of Nb,Nb, V, Fe, of the alloy the present waswork improved Comparedwith withadditions additions V, or Fe,Ta, orthe Ta,ductility the ductility of theinalloy in thework present was effectively while maintaining an ultra-high level of strength by the addition of titanium. improved effectively while maintaining an ultra-high level of strength by the addition of titanium.
(d) 0.5%Ti.
The tensile strength and elongation values of Cu-15Ni-8Sn alloys with Ti addition and other designed Cu-15Ni-8Sn alloys are summarized in Figure10 [8–10]. Adding the appropriate amount of trace elements is an effective way to improve mechanical properties of the Cu-15Ni-8Sn alloys. Materials 2017, 10, 1038 10 of 11 Compared with additions of Nb, V, Fe, or Ta, the ductility of the alloy in the present work was improved effectively while maintaining an ultra-high level of strength by the addition of titanium.
Figure of mechanical properties of of Cu-15Ni-8Sn Cu-15Ni-8Sn alloys alloys with with the the addition addition of of Ti Ti and Figure 10. 10. Comparison Comparison of mechanical properties and other elements. other elements.
4. Conclusion Conclusions In the present and mechanical properties of present work, work,the theeffects effectsofofTiTiaddition additionononthe themicrostructure microstructure and mechanical properties Cu-15Ni-8Sn alloys were investigated. TheThe following conclusions cancan be be drawn: of Cu-15Ni-8Sn alloys were investigated. following conclusions drawn:
alloys withthe theaddition additionofofTiTipossess possessan anexcellent excellentcombination combination of of strength strength and and (1) (1)Cu-15Ni-8Sn Cu-15Ni-8Sn alloys with ductility. Compared to thetoalloy without Ti, the strength and the of theofalloy with ductility. Compared the alloy without Ti,tensile the tensile strength andelongation the elongation the alloy 0.3%Ti increased from 935MPa to 1024MPa and from 2.7% to 17.9%, respectively. with 0.3% Ti increased from 935 MPa to 1024 MPa and from 2.7% to 17.9%, respectively. (2) improvement of mechanical properties by the addition of Ti contributed to grain (2) TheThe improvement of mechanical properties by the addition of Ti contributed to grain refinement refinement and the suppression of discontinuous precipitation. Byaddition the addition Ni3Ti phase and the suppression of discontinuous precipitation. By the of Ti,ofNiTi, 3 Ti phase was was formed in the solidification process, which was undissolved in the following solution formed in the solidification process, which was undissolved in the following solution treatment. treatment. Ni3TiNi phase had a pining effect on grain boundary migration in the process of solution treatment, 3 Ti phase had a pining effect on grain boundary migration in the process of solution treatment, and thus resulted in grain refinement. Ti addition also had and thus resulted in grain refinement. Ti addition also hada asuppressive suppressiveeffect effect on on discontinuous discontinuous precipitation, owing to a reduction of nucleation sites of discontinuous precipitation. precipitation, owing to a reduction of nucleation sites of discontinuous precipitation. Acknowledgments: This work is supported by the Guangdong Natural Science Foundation for Research Team (Grant No. 2015A030312003). Acknowledgments: This work is supported by the Guangdong Natural Science Foundation for Research Team (Grant No. 2015A030312003). Author Contributions: C. Zhao, W.W. Zhang and Z.Q. Luo conceived and designed the experiments; C. Zhao Author Contributions: Zhao, W.W. Zhang Z.Q. Luo and designed the C. experiments; Zhao and and D.X. Li performedC. the experiments; C. and Zhao and Z. conceived Wang analyzed the data; Yang and C. D.T. Zhang D.X. Li performed the experiments; C. Zhao and Z. Wang analyzed the data; C. Yang and D.T. Zhang contributed contributed analysis tools; C. the Zhao wrote paper; Wang and W.W. Zhang revised the manuscript. analysis tools; C. Zhao wrote paper; Z.the Wang andZ. W.W. Zhang revised the manuscript. Conflicts of ofInterest: Interest:The Theauthors authors declare conflict of interest. founding sponsors had in the Conflicts declare no no conflict of interest. The The founding sponsors had no rolenoin role the design design of the study; in the collection, analyses, or interpretation of data; the writing ofmanuscript, the manuscript, in of the study; in the collection, analyses, or interpretation of data; in theinwriting of the and and in the decision to publish the results. the decision to publish the results.
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