Materials and Design 30 (2009) 2099–2107

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Materials and Design

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Microstructural characteristics and formation mechanism of direct laser-sinteredCu-based alloys reinforced with Ni particles

Dongdong Gu *, Yifu Shen, Zhijian LuCollege of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, 29 Yudao Street, Nanjing 210016, PR China

a r t i c l e i n f o a b s t r a c t

Article history:Received 15 July 2008Accepted 22 August 2008Available online 2 September 2008

Keywords:Non-ferrous metals and alloys (A)Sintering (C)Microstructure (F)

0261-3069/$ - see front matter � 2008 Elsevier Ltd. Adoi:10.1016/j.matdes.2008.08.036

* Corresponding author. Tel.: +86 25 52112904x80E-mail addresses: dongdonggu@nuaa.edu.cn, dongd

Direct metal laser sintering (DMLS) was used to consolidate Cu-based alloy powder (Cu–10Sn and Cu–8.4P)reinforced with Ni particles. Phases, microstructures, compositions, and mechanical properties of laser-sintered part were studied. It showed that particle bonding was through a liquid phase sintering mecha-nism involving the complete melting of matrix alloy powder and the non-melting of cores of Ni reinforcingparticles. A significant smoothening of Ni particles occurred in the liquid and the dissolved Ni elementalloyed with Cu element to form CuNi solid solution, leading to a coherent particle/matrix interface aftersolidification. The dendrites of matrix alloys developed directionally and the primary dendritic spacing washighly refined to �1.5 lm, due to laser-induced super high temperature gradient and solidification rate.The additive P element acted as a localized deoxidizer to prevent the sintering system from oxidation byformation of CuPO3, thereby enhancing liquid–solid wettability and resultant sintering activity. A highdensification level of 95.2% theoretical density was obtained after sintering. The fracture surface oflaser-sintered part was mainly featured by a strong ductile type of fracture. The dynamic nanohardnessof Ni reinforcing phase and Cu-based matrix alloys reached 1.82 GPa and 0.99–1.35 GPa, respectively.

� 2008 Elsevier Ltd. All rights reserved.

1. Introduction

Cu-based alloys are widely used materials due to their excellentelectrical and thermal conductivity combined with outstandingresistance to corrosion and fatigue [1–3]. However, the relativelylow inherent strength limits their practical use. Particle reinforcedmetal matrix composites (MMCs) possess a favorable combinationof metallic matrix and stiffer and stronger reinforcements. Thepreparation of Cu-based alloys reinforced with second-phase parti-cles is expected to improve the overall performance. Typically,ceramic phases in terms of carbide (e.g., SiC [4], TiC [5] and WC[6]), nitride (e.g., Si3N4 [7] and TiN [8]), boride (e.g., TiB2 [9] andB4C [10]), and oxide (e.g., Al2O3 [11] and ZrO2 [12]) are used asthe reinforcement. However, the ceramic reinforcing particles gen-erally have poor wettability with the matrix metal. The deleteriousinterfacial microcracks, consequently, are much easy to form be-tween the reinforcement and the matrix, thus degrading the den-sification level and resultant mechanical properties of finialproducts [13]. In this viewpoint, the selection of strong metals asthe reinforcement is regarded reasonable, due to the favorablewetting characteristics of metal/metal interfaces. Ni, as a typicalnonferrous strong metal, has significantly elevated properties rela-tive to Cu and its alloys (HV 200–250 of Ni vs. HV 85 of Cu [14],tensile strength rb 392 MPa of Ni vs. 196 MPa of Cu [15]). More

ll rights reserved.

517; fax: +86 25 52112626.onggu@hotmail.com (D. Gu).

important, according to Ni–Cu binary phase diagram [16], Ni andCu system exhibits unlimited intersolubility, which favors the for-mation of compatible particle/matrix interfaces when Ni particlesare used to strengthen Cu-alloy matrix.

As a typical rapid prototyping (RP) technique, direct metal la-ser sintering (DMLS) enables the quick production of complex-shaped three-dimensional (3D) components directly from metalpowder [17–22]. Unlike conventional powder metallurgy (PM)methods, DMLS process creates parts in a layer-by-layer fashionby selectively fusing and consolidation of thin layers of loosepowder with a scanning laser beam [23–25]. This technique com-petes effectively with other conventional manufacturing pro-cesses when the part geometry is complex and the productionrun is not large [26]. Recent research efforts in DMLS have dem-onstrated that this technique gives a great promise for net-shapefabrication of functional prototypes and special tooling for injec-tion molding and die casting [27–31]. DMLS process, due to itsflexibility in materials and shapes, also exhibits a great potentialfor creating complex-shaped MMCs components that cannot beeasily developed by other conventional methods. Nevertheless,the fundamentals of this process, e.g., powder properties, processcontrol, and metallurgical mechanisms, have not been well ex-plored. Process defects such as balling effect, curling deformation,and poor densification are generally associated with DMLS-pro-cessed parts, due to the complex nature of this process includingmultiple modes of heat, mass, and momentum transfer and, insome cases, chemical reactions [32–34]. Actually, not much pre-

Fig. 1. SEM images showing characteristic morphologies of starting powder components: (a) Ni powder; (b) Cu–10Sn powder; (c) Cu–8.4P powder.

2100 D. Gu et al. / Materials and Design 30 (2009) 2099–2107

vious work has been focused on the basic principles regardingDMLS of particle reinforced MMCs. Significant research andunderstanding are still required for the fabrication of high perfor-mance MMCs components with controllable microstructures andproperties.

In this work, DMLS process was used to consolidate a compositesystem consisting of Cu-alloy powder, acting as the matrix metal,and Ni powder, as the reinforcement. The phases, compositions,and microstructures of laser-sintered powder were assessed andthe powder bonding mechanisms during DMLS of a composite sys-tem were elucidated.

2. Experimental

2.1. Materials

The composite powder system as investigated consisted ofthree components (supplier: Haining Feida Metallurgy PowderCo., Ltd. China): 99.9% purity Ni powder with an irregular shapeand a mean equivalent spherical diameter of 52 lm (Fig. 1a),water-atomized CuSn (10 wt.% Sn) powder with a near-ellipsoidalshape and an average particle size of 28 lm (Fig. 1b), and gas-atomized CuP (8.4 wt.% P) powder with a spherical morphologyand a mean particle size of 16 lm (Fig. 1c). Phosphorus element,acting as a fluxing agent, was added in the form of pre-alloyedCu–8.4P to improve the wetting properties of the powder system.The three components were uniformly dispersed according to Ni:-CuSn:CuP weight ratio of 60:30:10 in a Fritsch Pulverisette 6 plan-etary mono-mill at a rotation velocity of 250 rpm for 60 min,without the aid of grinding balls.

2.2. Processing

The used DMLS apparatus mainly consisted of a continuouswave Gaussian CO2 (k = 10.6 lm) laser with a maximum outputpower of 2000 W (type: Rofin–Sinar 2000SM, supplier: Rofin–SinarLaser GmbH), an automatic powder delivery system, and a com-puter system for process control (supplier: Beijing Long Yuan Auto-mated Fabrication System).

When a sample was to be prepared, a 45# steel substrate wasprimarily placed on the building platform and leveled. A thin layerof loose powder (0.20 mm in thickness) was then deposited on thesubstrate by the roller. Subsequently, the laser beam scanned thepowder bed surface to form a layerwise profile according to CADdata of the sample, using a simple linear raster scan (Fig. 2a). Thesimilar process was repeated in a layer-by-layer fashion until amulti-layer (30 layers) sample with dimensions of60 mm � 30 mm � 6 mm was finished (Fig. 2b). The laser sinteringprocess was performed in ambient atmosphere at room tempera-ture. Through a series of preliminary laser sintering experiments,the following optimized processing parameters were chosen: spotsize 0.30 mm, laser power 600 W, scan speed 0.04 m/s, and scanline spacing 0.15 mm.

2.3. Characterization

2.3.1. Phase identificationPhases of laser-sintered powder was identified by a Bruker D8

Advance X-ray diffraction (XRD) analyzer with Cu Ka radiation(k = 0.15418 nm) at 40 kV and 40 mA, using a continuous scanmode. A quick scan (4 degrees per minute) was primarily per-formed over a wide range of 20–100 2h degrees. A slower scan rate

Fig. 2. The real-time laser sintering process (a); The finished laser-sintered sampleon powder bed (b).

Fig. 3. Schematic of specimen for tensile test.

20 30 40 50 60 70 80 90 100

1 Ni2 Cu

3Sn

3 CuNi4 CuPO

3

Inte

nsity

(a.

u.)

2θ (deg)

2

4

31

4231

31 31

42 43 44 45 46

2

Inte

nsity

(a.

u.)

2θ (deg)

1 Ni2 Cu

3Sn

3 CuNi4 CuPO

3

13

4

2

49 50 51 52 53

2

1 Ni2 Cu

3Sn

3 CuNi4 CuPO

3

Inte

nsity

(a.

u.)

2θ (deg)

1

4

3

a

b

c

Fig. 4. XRD spectrum of laser-sintered sample (a); Identified peaks in the vicinity of42–46� (b) and 49–53� (c).

D. Gu et al. / Materials and Design 30 (2009) 2099–2107 2101

of 1 degree per minute was further used over 42–46 and 49–53 2hdegrees to determine phases more accurately.

2.3.2. Microstructural characterizationSpecimens for metallographic examinations were cut, ground,

and polished using standard procedures and etched in a solutionconsisting of HNO3 (10 ml) and (CH3CO)2O (10 ml) for 8 s. Micro-structures were characterized using an optical microscope (OM)and a Quanta 200 scanning electron microscope (SEM) at an accel-erating voltage of 20 kV. An EDAX energy dispersive X-ray spectro-scope (EDX) was used to determine chemical compositions.

2.3.3. Properties testingThe densities of laser-sintered samples were calculated based

on the Archimedes principle. Three samples were sintered underthe same processing conditions and the result of density measure-ment was obtained using the mean value.

Three specimens for tensile test, with individual dimensionsshown in Fig. 3, were prepared according to GB/T 228-2002 (equiv-alent to ISO 6892:1998) standards using a spark-erosion wire cut-ting machine. Prior to tensile test, the specimens underwentstress-relief annealing at 400 �C for 3 h. The tensile strength testwas carried out at room temperature with a universal testing ma-chine (type: CMT5105, supplier: Shenzhen SANS Testing MachineCo., Ltd. China) at a loading rate of 1.0 mm/min. The tensile direc-tion was parallel to the sintered layers. Three readings of tensilestrength were collected and an average value was calculated.

Microhardness of different phases in laser-sintered structurewas measured using a Shimadzu DUH-W201S nanoindentationtester at room temperature. A loading–unloading test mode, a test

force of 100 mN, a loading speed of 2.6478 mN/s, and a hold timeof 1 s were used. In the measurements, the load (P) and the inden-tation depth (h) were displayed. The raw data were then used to

2102 D. Gu et al. / Materials and Design 30 (2009) 2099–2107

construct the loading–unloading plot. The hardness is defined asthe ratio of the peak indentation load (Pmax) to the projected areaof the hardness impression (Ac). The dynamic nanohardness (Hd)is thus calculated by [35]:

Fig. 5. Binary phase diagrams of (a) Cu–Sn

Hd ¼Pmax

AcðAc ¼ 26:43h2

c Þ; ð1Þ

where hc is the contact depth under the maximum indentation load.

, (b) Cu–Ni, and (c) Cu–P systems [16].

Fig. 5 (continued)

D. Gu et al. / Materials and Design 30 (2009) 2099–2107 2103

3. Results and discussion

3.1. Phases

Fig. 4a depicts an overview of XRD pattern of laser-sinteredsample. The expanded views of the spectrums in the vicinity of42–46 and 49–53� (Fig. 4b and c) show the strong diffraction peaksof Ni phase. Moreover, a fraction of Ni element was present in theform of CuNi solid solution. Fig. 4b and c also reveales that a newCu-based phase Cu3Sn was formed after sintering. The P elementwas present as a new phase CuPO3, while no copper oxides, e.g.,CuO and Cu2O, were detected.

It is known that the pre-alloyed Cu–10Sn powder melts incon-gruently and possesses a considerably lower solidus temperature(�840 �C) and a liquidus one (�1020 �C) than the melting temper-ature of the Ni powder (�1455 �C) (Fig. 5a and b). The sinteringtemperature during DMLS of the present Ni–CuSn–CuP systemhas been estimated in literature [36] using the finite elementmethod. Due to a Gaussian laser beam used, the working temper-ature increases gradually from the edge towards the center of themolten pool. Under the given laser processing conditions (Section2.2), the estimated working temperature at the edge of the moltenpool is 1085 �C and the maximum temperature, 1461 �C, is ob-tained at the center of the pool. The minimum operating tempera-ture is well above the liquidus of Cu–10Sn, leading to a completemelting of matrix alloys. Meanwhile, the maximum sintering tem-

Table 1The Gibbs free formation energies (Gf) of pure substances at 298.15 K [37]

Substance Gf (kJ/mol)

CuO �128.29P 0P2O5 �1355.68Cu 0

perature is slightly above the melting point of Ni, favoring the sur-face melting of Ni particles. Therefore, laser sintering of the presentpowder system is based on a liquid phase sintering (LPS) mecha-nism involving complete melting of Cu–10Sn matrix alloy andnon-melting of the cores of Ni reinforcing particles. The Cu–10Snpowder melts to form liquid phase and, subsequently, precipitatesand solidifies again in the form of Cu3Sn on cooling (Fig. 4). Withthe sufficient wetting of the liquid phase during sintering, therough edges on the surfaces of Ni particles tend to melt and dis-solve in the matrix liquid. The Cu and Ni elements, because of theirunlimited intersolubility nature (Fig. 5b), are expected to formsolid solution alloy CuNi after solidification (Fig. 4).

Generally, the surfaces of the commercial metal powder consistof many impurities, especially the oxide film on particle surfaces[3]. In this powder system, a small amount of P element is addedin the form of pre-alloyed Cu–8.4P to act as a deoxidization agent.Above the liquidus temperature of Cu–10Sn powder, the Cu–8.4Ppowder melts completely, due to its low eutectic temperature of�714 �C (Fig. 5c). The dissociative P element tends to act as a local-ized deoxidizer to have a preferential reaction with CuO:

5CuOþ 2P ¼ P2O5 þ 5Cu: ð2Þ

According to Table 1, the change in the Gibbs free energies (DG)of reaction (2) is �714.23 kJ/mol. A negative DG, thus, leads to aspontaneous proceeding of the reaction. Afterwards, the reducingproduct (P2O5) has a further reaction with Cu2O, leading to the for-mation of CuPO3 (Fig. 4) through the following reaction:

Cu2Oþ P2O5 ¼ 2CuPO3: ð3Þ

It is known that metal oxides have a higher surface energy thanthat of the corresponding pure metals [3]. Thus, the presence of Cuoxides in the sintering system always results in a poor wettability.The initiation of reactions (2) and (3), fortunately, yields a cleansintering system free of any Cu oxides, thereby enhancing liquid–solid wetting characteristics and attendant sintering activity.

Table 2EDX results of elemental compositions at different positions in laser-sinteredstructure (Fig. 7a)

Zone Elemental composition (wt.%)

Ni Cu Sn P O

A 100.00B 36.20 52.29 7.78 1.57 2.16

Fig. 6. SEM image showing surface morphology of laser-sintered sample (a); OMimage showing microstructure on polished section of laser-sintered sample (b).

Escaped gas Powder bed

Torque Liquid Ni particle cores

Powder bed

Gas phase CuP CuSn Ni

Laser beam

Fig. 8. Schematic of metallurgical mechanisms during direct laser sintering of a Ni–CuSn–CuP composite system.

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3.2. Densification and microstructures

The characteristic surface morphology of laser-sintered samplewas provided in Fig. 6a. It was clear that the sintered surface wassmooth and dense and almost free of any balling phenomenon.The polished section of the finished sample, as shown in Fig. 6b, re-vealed that the interior structure was near-fully dense. Only asmall amount of small-sized pores were occasionally observed(arrowheaded). In this instance, a careful measurement of the sin-tered density revealed that a high value of 95.2% theoretical den-sity was reached. Thus, it was reasonable to conclude that DMLSprocess possesses a great promise to manufacture MMCs compo-nents with densification level equivalent to or superior than con-ventionally PM-processed Cu-based materials [15].

Fig. 7. SEM images showing (a) etched microstructure of particle/matrix composite

Fig. 7a shows the typical microstructure of the etched laser-sin-tered sample. The EDX results, as depicted in Table 2, revealed thatthe dispersed solid reinforcing particles (Zone A) were Ni and,

system and (b) unidirectionally developed dendritic structure of matrix alloys.

D. Gu et al. / Materials and Design 30 (2009) 2099–2107 2105

meanwhile, Ni element was also detected in the matrix (Zone B). Itwas most likely that a fraction of Ni element was dissolved in thematrix and, subsequently, alloyed with Cu to form CuNi solid solu-tion (Fig. 4). Besides Ni element, the matrix was composed of Cu,Sn, P, and O elements. Combined with XRD results (Fig. 4), it wasconfirmed that the matrix was consisted of the following Cu-basedalloys and compounds in terms of CuNi, Cu3Sn, and CuPO3. Themicrostructural characteristics revealed that the starting irregu-larly shaped Ni particles had a smooth and round shape after sin-tering, exhibiting no edges and corners on particle surfaces.Favorably, the interfaces between the reinforcing particles andthe matrix alloys were clear and continuous (Fig. 7a). Metallo-graphic study of the matrix at a higher magnification showed thatthe dendrites grew directionally along a certain preferred orienta-tion. Moreover, a significantly refined structure with the primarydendritic spacing of �1.5 lm was obtainable after solidification(Fig. 7b).

The metallurgical mechanisms during direct laser sintering ofthe present Ni–CuSn–CuP composite system are illustrated inFig. 8. When the laser beam scans over the powder bed, the energyis absorbed by powder particles through both bulk-coupling andpowder-coupling mechanisms [38]. In a first step, the energy is ab-sorbed in a narrow layer of individual powder particles deter-mined by the bulk properties of the materials, leading to a hightemperature of the surface of particles during the interaction. Inthis situation, the surfaces of both the high melting point Ni parti-cles and the low melting point CuSn and CuP particles tend tomelt. Afterwards, the heat flows mainly towards the center ofthe particles until a local steady state of the temperature, whichis above the liquidus of Cu–10Sn powder, is obtained within the

Fig. 9. SEM characterization of fracture surface of laser-sintered sample: (a) small-sizedin a small fraction of areas; (c) debonding along particle boundaries (occasionally obser

powder bed. With such a suitable operating temperature tailored,the CuSn and CuP powder melts completely, whereas the cores ofNi particles remain in solid. Thus, laser sintering of this compositepowder system produces a molten pool containing both liquid andsolid phases, inducing a significant gradient in chemical composi-tion at solid–liquid interfaces. On the other hand, since the spatialintensity profile of the laser beam follows a Gaussian distribution,a large thermal gradient is developed between the center and edgeof the molten pool. Either chemical concentration or temperaturegradient at solid–liquid interfaces gives rise to surface tension gra-dient and resultant Marangoni flow [13,32–34]. The presentMarangoni convection induces thermo-capillary forces for liquidflow. The liquid, thus, penetrates sufficiently into the voids be-tween Ni solid particles, forcing the existing gas phase to escaperapidly from the molten pool (Fig. 8). With the sufficient wettingof the liquid, the surfaces of the initially irregularly-shaped Ni par-ticles become smoothened. Meanwhile, due to the non-sphericalshape of Ni solids, there exists a torque because of misalignmentof the particle center [39]. The combined effect of the present tor-que and Marangoni flow tends to rotate the Ni particles in the li-quid, leading to an adequate rearrangement of particles (Fig. 8).With the laser beam moving away, the liquid–solid composite sys-tem enters a rapid solidification process to realize a finial densifi-cation. It is known that a large amount of microscopic pits at thesolid surfaces may act as favorable sites for heterogeneous nucle-ation of the liquid. The required activation energy is, thus, consid-erably low, which highly favors the nucleation of the liquid. Apreferred epitaxial solidification of the molten matrix alloys fromthe surfaces of well-arranged Ni particles tends to initiate, yieldingcoherent metallurgical interfaces between the Ni reinforcing parti-

ductile dimples formed in a majority of areas; (b) large-sized brittle dimples formedved).

1 2 3 4 5 6 7 80.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

2.2

Cu-based matrix alloys

Ni reinforcing phase

Dyn

amic

nan

ohar

dnes

s, H

d (G

Pa)

Measurement point, No.

Fig. 10. Dynamic nanohardness obtained in different positions within matrix alloysand Ni phase by nanoindentation test.

2106 D. Gu et al. / Materials and Design 30 (2009) 2099–2107

cles and the matrix alloy (Fig. 7a). Due to the extremely shortduration of the laser beam on any powder particles (typically lessthan 4 ms [34]), the laser-induced temperature gradient and solid-ification rate are as high as 1.0 � 106 K/m and 6.0 � 103 K/min,respectively [40]. Under this condition, the dendrites of matrix al-loys tend to develop unidirectionally along the heat flow directionand, meanwhile, the primary dendritic spacing is significantly re-fined (Fig. 7b). However, it is noted that due to a considerablyshort lifetime of the molten pool, a small fraction of gas phase isunable to escape completely from the pool and thus collapses,but gets trapped inside as bubbles [41], hence producing a smallamount of residual porosity in the rapidly solidified materials(Fig. 6b).

3.3. Sintered properties

The characteristic microstructures regarding the fracture sur-face of laser-sintered sample are provided in Fig. 9. It was ob-served that in a majority of areas on the fracture surface (over80%), a number of small-sized dimples were generally present,showing a typical ductile type of fracture (Fig. 9a). It is knownthat DMLS is a layer-by-layer additive manufacturing process.Consequently, the oxidation occurred in the previously sinteredlayer is a severe impediment to liquid–solid wetting propertywhen a fresh layer is being processed, resulting in delaminationcaused by poor interlayer bonding. In the present study, theabove-mentioned chemical reactions (2) and (3) between phos-phorus and oxygen allow the formation of a copper phosphate.Such kind of reducing substance is generally lighter than the li-quid metal and, thus, tends to float on the top of the melt underthe action of Marangoni convection within the molten pool [3].After solidification, the phosphate covers the surface of the sin-tered layer, forming a protective film between the laser-processedpowder and the ambient atmosphere. The effective protection ofsuch a film prevents the sintered layer from oxidation and,accordingly, ensures a good wettability and a coherent interlayerbonding during laser sintering of the subsequent layer. Also, asufficient degree of alloying between the two basic elements(i.e., Ni and Cu) in the sintering system favors an improvementin the inherent strength of laser-sintered powder. Under theseconditions, a strong ductile type of fracture is much easy to pres-ent in the finally finished part (Fig. 9a). The ultimate fracturestrength of the sintered part reached 198.2 MPa, showing a signif-icant improvement upon the strength of laser-sintered Cu-basedalloys without reinforcement reported in literature [1,42,43].However, a very small fraction of areas on the fracture surfaceexhibited a debonding along particle boundaries, showing a rela-

tively weak intercrystalline fracture (Fig. 9b). Correspondingly,some large-sized and round-shaped particles were occasionallyfound on the other side of the fracture section (Fig. 9c). It is mostlikely that the rapid and localized nature of DMLS prevents a suf-ficient removal and diffusion of the molten liquid in some laser-irradiated positions, thereby weakening the bonding coherenceof some relatively large Ni reinforcing particles by the solidifiedliquid.

Fig. 10 depicts the profiles of dynamic nanohardness measuredon different phases in laser-sintered structure. According to Eq.(1), the obtained dynamic nanohardness (Hd) in different positionsof Cu-based matrix alloys was generally between 0.99 and1.35 GPa. Thus, it was reasonable to conclude that laser-sinteredCu-based alloys possessed an equivalent hardness property relativeto conventionally processed copper materials [44]. Also, Fig. 10 re-vealed that the nanohardness in various positions of the matrixshowed some fluctuation. This is because the matrix alloys are com-posed of different Cu-based phases (Fig. 4) which possess differentinherent hardness values. The nanohardness distribution of the rein-forcing Ni phase was uniform and the average Hd reached 1.82 GPa,showing an improvement upon the obtained hardness of the matrixalloys.

4. Conclusions

Direct metal laser sintering (DMLS) was used to consolidate Cu-based alloy powder reinforced with Ni particles, and the followingconclusions were draw:

(1) Particle bonding was via a liquid phase sintering with thecomplete melting of matrix alloys (Cu–10Sn and Cu–8.4P)and the non-melting of cores of Ni reinforcing particles.

(2) A significant smoothening of Ni reinforcing particlesoccurred in the wetting liquid and the dissolved Ni elementalloyed with Cu element to form CuNi solid solution, leadingto a coherent particle/matrix interface after solidification.

(3) Laser-processed Cu-based matrix alloys mainly consisted ofCuNi and Cu3Sn. The dendrites of matrix alloys developeddirectionally and the primary dendritic spacing was signifi-cantly refined to �1.5 lm, due to laser-induced super hightemperature gradient and solidification rate.

(4) The additive phosphorus element acted as a localized deox-idizer to prevent the sintering system from oxidation by for-mation of a copper phosphate CuPO3, thereby enhancingliquid–solid wettability and resultant sintering activity.

(5) A high densification response of 95.2% theoretical densitywas obtained after sintering. The fracture surface of laser-sintered part was mainly featured by a strong ductile typeof fracture. The dynamic nanohardness of Ni reinforcingphase reached 1.82 GPa, showing an improvement relativeto that of Cu-based matrix alloys 0.99–1.35 GPa.

Acknowledgements

The present work is financially supported by the National Nat-ural Science Foundation of China (Grant No. 50775113). One ofthe authors (Dongdong Gu) gratefully appreciates the financialsupport from the Scientific Research Foundation for Newly Em-ployed Talents in Nanjing University of Aeronautics andAstronautics.

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