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Characterization and fracture behavior of bismuth–tin thermalfuse alloy wires produced by the Ohno continuouscasting process
Divya Bhardwaj, Hiroshi Soda⁎, Alexander McLeanDepartment of Materials Science and Engineering, University of Toronto, 184 College St., Toronto, Ont. Canada M5S3E4
Article history:Received 11 January 2010Received in revised form 26May 2010Accepted 28 May 2010
Bismuth–tin binary alloys containing high bismuth concentrations of 40 to 77% werecontinuously cast into wires of approximately 2 mm in diameter with casting speedsbetween 15 and 150 mm min−1 using the Ohno Continuous Casting (OCC) process. Themicrostructure was examined and tensile tests were performed for wires cast at variousspeeds. It was found that for slowly cast wires containing large primary bismuth dendrites,bismuth fracture occurring along the (111) plane exerted a key role in wire fracture, whilemicrostructures with refined bismuth dendrites exhibited a mixture of bismuth cracks andinter-phase decohesion, allowing the accommodation of larger strain before wire fracture.For wires with microstructures containing primary tin dendrites, inter-phase decohesionplayed a key role in wire fracture.
Thermal fuses are essential devices in various types ofelectronic equipment for the prevention of fire hazard andcomponent failure when electrical over-loading occurs. Aselectronic devices become more compact in size, there will bean increasing demand for smaller-diameter, thermal fusewires approximately 0.4 mm in diameter with wires of muchsmaller diameter possibly required in the future. Non-leadedlow-melting thermal fuse alloys containing high amounts ofbismuth are prone to segregation when alloys are solidifiedslowly [1–4] and are also strain rate sensitive [5], which willcause brittle fracture when they are strained under a high rate.For these reasons, it is difficult to produce non-leaded micro-thermal fuse wires containing bismuth by the conventionalcasting and extrusion route. The degree of strain ratesensitivity of cast material is microstructure dependent andcast structureswith finer grain size are known to reduce strainrate sensitivity [5] and increase uniformity in microstructure
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thus minimizing segregation. The microstructure modifica-tion can be achieved by the addition of alloying elements andrapid cooling. However, the addition of grain refiners, if itadversely influences the melting behavior, will not be suitablefor thermal fuse alloys. Recent studies [1–4] have focused onmodification of cast structures of low-melting temperaturealloys with rapid cooling using a heated mold continuouscasting process, known as the OCC process. In this process, themold is heated above the solidification temperature of thealloy being cast and cooling occurs outside the mold.Solidification thus takes place at the mould exit, significantlyreducing or eliminating friction between the cast product andmold wall [6–8]. This casting configuration permits thegeneration of small-diameter cast products with a high-quality surface and elongated crystals growing along thecasting direction. These types of products cannot be obtainedthrough traditional processes.
The objective of the present study is to examine the effectof the microstructure modification of non-leaded Bi–Sn alloys
Fig. 1 – Schematic diagram of the melting and castingarrangement [9].
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with respect to structure–ductility relationships of the castwire products produced by the OCC process and to helpdevelop a new process route relevent to the manufacture ofmicro-thermal fuse products.
2. Experimental Aspects
Fig. 1 shows a schematic diagram of the melting and castingarrangement [9]. The equipment consists of a melting apparatuswith a rectangular open cavity constructed out of a 150-mm longand 37-mm diameter graphite rod, a cooling device, a glasssupport tubeplacedonaplatformforminimizing themechanicalinstability, and pinch rolls for withdrawal of cast wires. Therectangular cavity (20-mmW×38-mm L×20-mm H) serves as amelting and holding vessel for the alloy. At the bottomend of thevessel, there is a 16-mm long, 2-mm diameter channel throughwhich the liquid exits into the cooling zone to produce wires.About 25 g of Bi–Sn alloy melt, prepared from 99.99% Sn and99.99%Bi (all inmass%),washeld in thecavity, the temperatureofwhich was controlled and maintained by the thermocouplelocatednear themoldchannelat 207, 179, 194, and237 °Cfor alloycompositions 40, 57, 62, and 77% Bi, respectively. Thesetemperatures were approximately 40 °C above the liquidustemperature of the respective alloy compositions. The meltcomposition was chosen to produce three different cast struc-
Fig. 2 – Continuously cast 57% Bi–Sn wires approximately2 mm in diameter.
tures, containing (a) a primary tin dendrite and eutectic structure(40% and 57% Bi), (b) a eutectic structure only (62% Bi), and (c) aprimary bismuth dendrite and eutectic structure (77% Bi).
Fig. 3 – Microstructure (transverse section) of alloy wires castat 30 mm min−1 for bismuth concentration: (a) 57%, (b) 62%,and (c) 77%. Dark phase is tin and light phase is bismuth.
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To begin casting, a stainless steel tube 2 mm in diameterwas brought in contact with the channel exit and themelt wassuctioned into the tube cavity using a pipette bulb attached tothe opposite end to capture the initial melt. The cooling water,with a flow rate of approximately 30mL min−1 was turned onand the casting was initiated by moving the pinch rolls. Thecasting speed was increased gradually to the target speed andthe cooling water, initially positioned approximately 2 mmaway from the channel exit, also readjusted slightly accordingto the casting speed. The mold temperature, cooling distance,and casting speed are all interrelated parameters. Studiespertaining to the effects of these variables are described indetail elsewhere [10,11]. A wire approximately 2 mm indiameter and 30 to 40 cm in length was produced at eachcasting with speeds ranging from 15mm min−1 to 150mmmin−1. An example of cast wire product is shown in Fig. 2. Themicrostructure was examined and tensile tests performed forwires cast at various speedswith a gauge lengthof 100 mmandcross-head speed of 10mm min−1. The effects of microstruc-ture on fracture behavior of the cast wires were examined byobserving the fracture surface and cross-section of the fractureend using an electron microscope.
3. Results and Discussion
3.1. Microstructure and Ductility of Cast Wires
Fig. 3 shows themicrographs of cast structures of wires cast at30 mm min−1 with 57%, 62%, and 77% bismuth concentration.The cast structure for the wire of 57% bismuth concentrationconsisted of primary tin dendrites (wide darker region)
Fig. 4 – Changes in stress–strain curves with alloy composition anrate was 1.67×10−3 s−1.
coupled with coarser bismuth phase (white) and Chinesescript-type eutectic cell occupying the majority of the area.The structure for the wire of 62% bismuth concentration onlyconsisted of eutectic cells containing finer lamellar-typestructure within the cell and coarser irregular structurealong the cell boundaries, and for the 77% bismuth wire thestructure was composed of coarse primary bismuth dendritewith a trifoliate arm and irregular eutectic. Bismuth, which is asoft but brittle material, has almost no solubility for tin and,therefore, solidifies or precipitates as pure bismuth. For thisreason, bismuth concentration affects the mechanical prop-erties of cast wires. For example, as shown in Fig. 4, tensileductility decreased due to an increase in bismuth concentra-tion. However, casting conditions, which exert considerableinfluence on cast structure, also alter mechanical propertiessignificantly. Fig. 5 shows the effect of casting speed onmicrostructure for the alloy with 57% Bi. As can be seen in themicrographs, the wide eutectic areas between the tin den-drites in a specimen produced at a casting speed of 30 mmmin−1 almost disappeared at 150 mm min−1 due to theextensive development of secondary dendrite arms into theeutectic areas, causing the eutectic bismuth to be confinedbetween the regularly spaced dendrite arms. This increasesthe uniformity of the microstructure.
3.2. Fracture Characteristics
3.2.1. Structure of Wire Containing Bismuth DendritesAs shown in Fig. 6a, 77% bismuth alloy wires cast at slowerspeeds of 15 and 30 mm min−1 fractured abruptly withoutcausing apparent necking, resulting in low ductility of about 5to 8% elongation represented by engineering stress–strain
d casting speed for continuously cast wires. The initial strain
Fig. 5 – Longitudinal sections of 57% Bi–Sn alloy cast at speedsof (a) 30 and (b) 150 mm min−1, showing the development oftin dendrites with casting speed: dark phase is tin and lightphase bismuth.
Fig. 6 – Fractured region of 77% Bi–Sn wires cast at (a) 30 mmmin−1 and (b) 150 mm min−1. The initial strain rate was1.67×10−3 s−1.
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curves (Fig. 4), while wires cast at 150 mm min−1 exhibitedmuch higher tensile strength and elongation and producednecking prior to fracture (Fig. 6b). The reduction in area wasapproximately 10% for wires cast at 30 mm min−1, while forwires cast at 150 mm min−1 it was as high as 40% suggestingthe difference in fracture behavior.
Boundary sliding and phase sliding occur due to stressconcentration at locations where continuity in structure wasterminated by large phases or grain boundaries [12]. Fig. 7shows the SEM micrograph of a wire surface near the fractureend of a 77% bismuth specimen cast at 30 mmmin−1. There arevisible deformation bands running approximately 45° to thetensile direction on the primary bismuth crystals and smallsurface rumples within the eutectic region caused by thesliding of the tin and bismuth phases, showing some flexibilityin the material. However, as shown in Fig. 8, the fracturesurface of the wire, cast at 30 mm min−1, exhibited thecharacteristics of brittle failure with trefoil bismuth dendritesexposed throughout the fracture surface. Based on the smalldifferences in angle between the clusters of dendrites, thecross-sectional area contains several grains with slightly
different orientation from each other, causing the differentinclinations in the faceted fracture surface. A close-up of thedendrite and eutectic area indicates cleavedbismuthdendritesand eutectic bismuth.
In order to observe the fracture behavior, the longitudinalcross-section of the wire at the fracture end was examined.Fig. 9 (a) to (c) are the micrographs taken across the diameterof the fracture end, showing the characteristics of fracturemodes such as the cleaved primary bismuth, crack formationwithin the bismuth crystal, the crack path extending from theprimary bismuth crystal to the eutectic region passing alongthe eutectic bismuth (Fig. 9a), the cleaved eutectic bismuth(Fig. 9b), and also the eutectic region, exhibiting the evidenceof decohesion at the tin/bismuth interface (Fig. 9c).
Bismuth crystal cleaves along (111) plane [13,14] andproduces triangular etch pits in the (111) plane of bismuth
Fig. 7 – SEM view of the wire surface near the fracture tip of77% Bi–Sn after tensile testing. Tensile direction was verticaland the initial strain rate was 1.67×10−3 s−1. The wire wascast at 30 mm min−1.
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crystal when the exposed (111) surface is etched [14]. In thisstudy, the deep etching of the fracture surface by a nitric acidsolution also produced triangular etch pits in every fracturedsurface of the primary bismuth dendrites (Fig. 10a). Thisindicates that the fractured surfaces of dendrites are all on the(111) plane. Also seen were perpendicular cracks to the (111)plane, indicating that multiple cracks can easily occur inbismuth crystals (Fig. 10b). Also noticed were the facetedfracture surfaces of the eutectic bismuth phase (Fig. 8c),stemming from the primary bismuth crystal, suggesting thatthe crystal orientation of the fracture surface of eutecticbismuth is in the same orientation as the fractured primarybismuth crystal. As shown in Fig. 11, a deep etched eutecticregion, inwhich the interconnectedeutectic bismuthphasewasexposed, also revealed triangular etch pits, indicating that thecleaved surface of the eutectic bismuth in this region is the (111)plane. Although the fracture surface of some other eutecticareas indicated the evidence of decohesion of phases (Fig. 9c),the microstructure observations of fracture confirm thatbismuth cracking along the (111) plane played a key role inabruptwire fracture and lower elongation values for thosewirescast at slower casting speeds. It was noted that bismuthcracking in this alloy occurs from the early stage of the plasticdeformation process [9]. Casting speed exerts significant effectson microstructure. Fig. 12 shows the necked portion of a wirecast at 150 mm min−1. Much finer dendrite structure wasobserved in comparison with that of wires cast at 30mmmin−1. Primary dendrite spacing, measured by the linearintercept method, decreased (as shown in Fig. 13) fromapproximately 120 μm to 20 μm as casting speed increased
Fig. 8 – (a) Fracture end of 77% bismuth alloy wire cast at30 mm min−1, showing faceted fracture surfaceswith primary tri-petal bismuth dendrites and eutectic, (b) aclose-up of cleaved bismuth dendrite, and (c) cleaved eutecticbismuth (dark arrows).
Fig. 9 – Longitudinal cross-section at the fracture end of a 77%bismuthalloywire cast at 30mmmin−1, showing (a) cracks (lightarrows) in the primary bismuth crystal and crack propagationinto the eutectic region, (b) cleaved eutectic bismuth(lightarrows), and (c) evidenceofdecohesion (lightarrows) in theeutectic region. Tensile direction is vertical. Dark phase is tinand light phase is bismuth.
Fig. 10 – SEM micrographs showing (a) triangular etch pits(dark arrows) in the fractured surface of a bismuth dendriteand (b) secondary cracks (dark arrow) perpendicular to (111)plane.
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from 15mmmin−1 to 100 mm min−1. In this alloy, although themicrostructure became significantly finer with an increase incasting speed, it retained similar morphology. However, owing
to finer and more uniform microstructure, wires cast at higherspeeds tended to exhibit higher yield strength and elongationwith the formation of localized necking just prior to fracture.Wire surface near the fracture end, shown in Fig. 14, exhibitedmore extensive deformation and phase sliding in comparisonwith that of wire cast at 30 mmmin−1 (Fig. 7).
Observation of the fracture surface indicates that thefracture mode also involves cleavage of primary bismuth(indicated by dark arrows in Fig. 15a) as observed in specimensproduced at lower speeds. However, as shown in the longitu-dinal cross-section of the fracture end (Fig. 15b), the voidformation was more noticeable at tin/bismuth boundaries.Since the plastic deformation of tin and bismuth does notconform to each other [9], if bismuth cracking plays a lessactive role in wire fracture owing to its finer and uniformstructure, there will be more chance of cavity formation at the
Fig. 11 – Cleaved eutectic bismuth, showing triangularetch-pits (indicated by arrows), revealed by deep etchingwitha nitric acid solution.
Fig. 13 – Change in primary dendrite spacing with castingspeed.
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tin/bismuth boundaries during plastic deformation, resultingin more accommodation of strain.
3.2.2. Structure of Wire Containing Tin DendritesA 57% bismuth alloy, despite its eutectic composition,normally produces primary tin dendrites. Owing to ductiletin dendrites in place of brittle bismuth dendrites, wires casteven at slower speeds exhibited higher ductility than 77%bismuth alloywires (Fig. 4). As can be seen in the stress–strain
Fig. 12 – Longitudinal section of a 77% bismuth alloy wireafter a tensile test. The wire was cast at 150 mm min−1.
curves, the profile was similar regardless of different castingspeeds, indicating that all wires fractured in a similarmanner. Fig. 16 shows the fracture surface of a wire cast at30 mm min−1 with 57% bismuth concentration. The nature ofthe fracture surface indicates the ductile fracture modeexhibiting a spongy appearance in the region shown bythe white arrow, indicating the material separation began inthis region before wire fracture. Outside this region, forexample, a micrograph (Fig. 16b) observed at the locationindicated by the black arrow in Fig. 16a also shows small voidsand decohesion (indicated by light arrows) in addition toareas of cleaved bismuth (indicated by dark arrows). Thefracture surface of wires cast at 150 mm min−1 exhibitedmuch wider areas of spongy regions as shown in Fig. 17, in-dicating more extensive phase separation occurred beforefracture. Fig. 18a shows the wire surface at the fracture end
Fig. 14 – SEM view of the wire surface near the fracture tip of77% Bi–Sn after tensile testing. Tensile direction was verticaland the initial strain rate was 1.67×10−3 s−1. The wire wascast at 150 mm min−1.
Fig. 15 – (a) Fracture surface and (b) longitudinal section at thefracture end of 77% Bi–Sn wire cast at 150 mm min−1.
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for wires cast at 150 mm min−1. Wire surfaces exhibitedextensive phase shifting, causing cracks, cavities, and deco-hesion occurring approximately 30° to the tensile direction.The longitudinal cross-section of the fracture end shown inFig. 18b clearly depicts the location of voids and also similardecohesion occurring at 30°. It is evident in the micrographthat voids occurred at the tin/bismuth phase boundaries andthese voids developed into cracks by linking up with each
Fig. 16 – (a) Fracture surface of a 57% Bi–Sn alloy wire cast at30 mm min−1, (b) large magnification of the region indicatedby black arrow in (a), showing the areas of cleaved bismuth(dark arrows) and voids and decohesion (light arrows), and(c) longitudinal section of fracture end, showing the voids anddecohesion (light arrows).
other. Thus, the failure of 57% Bi–Sn alloy wires was largelydue to inter-phase separation between the tin and bismuthphases. From the observations of the longitudinal sections of
Fig. 17 – Fracture surface of a 57% Bi–Sn alloy wire cast at150 mm min−1.
Fig. 18 – (a) Wire surface and (b) longitudinal cross-section ofthe fracture end of 57% Bi–Sn wire cast at 150 mm min−1.Tensile direction is vertical. Dark phase in (b) is tin and lightphase is bismuth.
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micrographs (Figs. 16c and 18b), the voids form at tin/bismuthboundaries, however the fissures appear to develop morefrom the coarse tin phase boundaries as indicated by thewhite arrows in Fig. 16c. The coarse tin phase is most likelythe primary tin dendrite, although it is difficult to differen-tiate the primary tin phase from eutectic tin phase. In order toconfirm if the fissure forms preferentially, an alloy ofcomposition 40% Bi–Sn containing more primary tin wascast at a slow speed of 30 mm min−1 to produce largedendrites and subjected to tensile testing. Fig. 19 shows thefracture surface and longitudinal cross-section of the fractureend of the wire. Voids again clearly formed both at theprimary tin/precipitated bismuth boundaries and the primarytin/eutectic bismuth. However, fissures tended to developmore from the primary tin/eutectic bismuth boundaries.
3.3. Fracture Behavior
The deformation behavior of cast 62% Bi–Sn wires has beeninvestigated in a previous study [9]. For 62% Bi–Sn wires castat higher speeds, the occurrence of inter-phase cracksbetween the bismuth and tin phases became predominantin the latter part of tensile deformation. For example, itsharply increased after the amount of strain exceededapproximately 2/3 of the uniform strain regime, while anincrease in the number of bismuth cracks remained linear.This sharp increase contributed to the declining of the load-bearing capability within the uniform strain regime of thetensile specimen.
Fig. 20 shows the fracture surface of a 62% Bi–Sn wire,produced at a casting speed of 150 mm min−1. It exhibitedlarge areas of spongy appearance, indicating that materialseparation due to phase decohesion which led to wire fracturebegan in these areas. The associated stress–strain curve forthis wire is shown in Fig. 21. Since the visible neck formation
occurs at the end of tensile deformation, leading to significantdecrease in stress values, the section of curve where the stressvalue decreases linearly, up to about 15% extension in Fig. 21,can be considered in macroscopic terms as a uniform strainregime. In order to observe the change in the load-bearingcapability within the uniform strain regime, the values of truestress (σ) were calculated at any given extension for theuniform strain part of the stress–strain curves by assumingthe constancy of the volume relationship and using therelation σ=S (ε+1) where S and ε are engineering stress andstrain respectively. These were plotted together with thevalues of the corresponding engineering stress in Fig. 21. The
Fig. 19 – (a) Fracture surface of 40% Bi–Sn wires producedat a casting speed of 30 mm min−1 and (b) longitudinalcross-section of the wires at the fracture end. Dark arrowsindicate eutectic bismuth and white arrows indicate primarytin dendrites. Tensile direction is vertical.
Fig. 20 – Fracture surface of the 62% Bi–Sn wire after tensiletesting at the initial strain rate of 1.67×10−3 s−1. The wire wascast at 150 mm min−1.
Fig. 21 – Uniform strain part of engineering stress–straincurve, represented by line A (approximately up to 0.15 strain)and corresponding true stress value for the 62% Bi–Snwire cast at 150 mm min−1, exhibiting slope change atapproximately 0.1 strain.
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true stress–displacement curve for the uniform strain part,approximately up to 15% strain, can be roughly represented bylines (1) and (2) and slope changes occurring at the vicinity of0.1 strain. This transition point is situated at about 70% of theuniform strain part of the curve and was in accord with theresults obtained previously [9], confirming that the accelerat-ed inter-phase decohesion starts to occur at around this pointof decrease in load carrying ability.
In order to determine whether the same behavior exist forwires with other compositions, engineering stress–straincurves for 57% Bi–Sn wires cast at 30 and 150 mm min−1
were treated in the same manner and these are shown inFig. 22. The uniform strain part of the engineering stress–strain curves, represented by line A, was up to 0.16 and 0.22
strain approximately for 30 mm min−1 and 150 mm min−1,respectively. True stress–displacement curves for the uniformstrain part were found to be represented in the same way bylines (1) and (2). Transition occurred at about 0.11 strain forwires cast at 30 mm min−1 and at 0.15 strain for wires cast at150 mm min−1. These transition points were approximately at70% of the uniform strain amount for both wires and are alsoin agreementwith that of the 62% Bi–Snwire (Fig. 21). This 70%of the uniform strain amount may be considered as athreshold strain amount beyond which the void formationincreases dramatically.
In order to observe if the same behavior also occurs for77% Bi–Sn wires, the stress–strain curves for wires cast at 30and 150 mm min−1 were treated in the same way in Fig. 23.For a wire cast at 30 mm min−1, only a single straight line
Fig. 22 – Uniform strain part of engineering stress–straincurves, represented by line A (approximately up to 0.15 and0.34 strain for 57% Bi–Sn wires cast at 30 and 150 mm min−1,respectively), and corresponding true stress value.
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was required to represent the curve. This supports thefact that wires cast at slower casting speeds fractured in abrittle manner due to bismuth cracking before significant
Fig. 23 – Uniform strain part of engineering stress–straincurves, represented by line A (approximately up to 0.076 and0.1 strain for 77% Bi–Sn wires cast at 30 and 150 mm min−1,respectively) and corresponding true stress value.
inter-phase decohesion has a chance to take place (Figs. 8and 9). On the other hand, the engineering stress–straincurve for a wire cast at 150 mmmin−1 shows a uniform strainamount of 0.1. Lines 1 and 2, with transition occurringapproximately at 70% of the uniform strain amount, canrepresent the true stress displacement curve. Thus, it canalso be that the accelerated decline in slope, represented byline 2, is due to the development of inter-phase decohesion,which had a significant role in accommodation of strain inthe latter part of the tensile deformation process.
Fracture mode is influenced by strain rate imposed duringtensile testing and also dependent on microstructure [1,5,12]. Itwasnoted [12] that for theSn–37%Pbeutectic alloy cast andagedthat cavities occurred owing to lead-rich particles along thegrain boundaries at the low strain rate of 1.34×10−4 s−1 andcracks progressed inter-granularly linking up the cavities, whileat the high strain rate of 2.78×10−3 s−1, the failuremode becameintra-granular. The bismuth–indium–tin eutectic alloy [1] with acoarse eutectic solidification microstructure exhibited brittleintra-granular fracture when it was tensile tested beyond theinitial strain rate of 3.33×10−3 s−1, while the same alloycomposition, cast using the OCC process that produced a finelymodified microstructure, exhibited a ductile inter-granular(inter-phase) fracture mode even when it was pulled at theinitial strain rate of 6.67×10−3 s−1, indicating that strain sensi-tivitywas greatly reduced forwires cast by theOCCprocess. Thestrain sensitivity alsodecreased for 77%Bi–Snwires producedathigher speed and fracture occurred with both bismuth fractureand cavity formation at tin–bismuth boundaries, leading tointer-phase decohesion.
4. Summary
Bismuth–tin wires with varying bismuth compositions ap-proximately 2 mm in diameter were produced using the OCCprocess at different speeds in order to observe the structure–property relations. The results obtained from tensile tests andmicrostructural observations are summarized as follows:
1. Alloy compositions with varying degrees of liquidus andsolidus ranges can be cast into wires with uniformmicrostructure by the OCC process, which is essential forthermal fuse alloy products.
2. Cast wires containing primary bismuth dendrites, bismuthfracture occurring along the (111) plane exerted a key rolein wire fracture.
3. By modifying microstructure using faster casting speeds,the adverse effect of the bismuth phase was reduced,leading to improved ductility.
4. Wires with microstructures containing tin dendritesexhibited inter-phase voids, leading to decohesion andfracture.
5. Accelerated decrease in load carrying capability due to anincrease in inter-phase decohesion was recognized afterstrain exceeding approximately 70% of the uniform strainregime for wires with microstructures containing tindendrites. This phenomena were also observed for 77%bismuth-containing wires cast at faster speeds.
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Acknowledgements
Appreciation is expressed to the Natural Sciences and Engi-neering Research Council of Canada for financial support of thisproject. The in-kind support provided by Process ResearchOrtech and Electrovaya Inc. is also gratefully acknowledged.
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