Experiment A: Solidification and Casting - Course...
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A1 Figure 1. Grain Structure of a solid casting. Experiment A: Solidification and Casting Introduction: The purpose of this experiment is to introduce students to the concepts of solidification and to study the development of solidification microstructures. The lab is divided into three parts: Part 1: Solidification of a pure element Observation of the dendritic growth in pure lead. Examination of a cross-section of a lead casting. Part 2: Solidification of the ammonium chloride/water system Using this system as a transparent model for metal alloy casting. A saturated solution of ammonium chloride will be used to simulate a superheated liquid metal undergoing different supercoolings. The solidification behaviour will then be observed. Part 3: Solidification microstructures The microstructures of cast alloy systems will be viewed under the light microscope and related to the appropriate phase diagram. Background: Casting The fabrication of most metallic and many nonmetallic materials involves melting the raw materials and pouring the resulting liquid into a mould which produces a solid of manageable size and shape. Solidification usually proceeds inward from the mold wall, as heat is extracted out through the wall. As a result, the grains that form are often columnar or long, narrow and run perpendicular to the mold wall. The grains usually do not grow homogeneously and instantaneously. Each grain forms a skeletal structure of planes first, the remaining liquid between the planes solidifying later. The skeletal framework of a grain is called a dendrite and is similar to the snowflake structure found in nature. A typical casting shows three distinct zones (Figure 1), a thin chill-cast zone adjacent to the mold wall formed by heterogeneous solid nucleation at the mold wall-liquid interface, a columnar zone formed by preferential growth of dendrites and a central equiaxed zone.
Transcript of Experiment A: Solidification and Casting - Course...
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Figure 1. Grain Structure of a solid casting.
Experiment A: Solidification and Casting
Introduction:
The purpose of this experiment is to introduce students to the concepts of solidification and tostudy the development of solidification microstructures. The lab is divided into three parts:
Part 1: Solidification of a pure elementObservation of the dendritic growth in pure lead. Examination of a cross-section of a leadcasting.
Part 2: Solidification of the ammonium chloride/water systemUsing this system as a transparent model for metal alloy casting. A saturated solution ofammonium chloride will be used to simulate a superheated liquid metal undergoing differentsupercoolings. The solidification behaviour will then be observed.
Part 3: Solidification microstructuresThe microstructures of cast alloy systems will be viewed under the light microscope and relatedto the appropriate phase diagram.
Background:
Casting
The fabrication of most metallic and many nonmetallic materials involves melting the rawmaterials and pouring the resulting liquid into a mould which produces a solid of manageablesize and shape. Solidification usually proceeds inward from the mold wall, as heat is extractedout through the wall. As a result, the grains that form are often columnar or long, narrow andrun perpendicular to the mold wall. The grains usually do not grow homogeneously andinstantaneously. Each grain forms a skeletal structure of planes first, the remaining liquid
between the planes solidifying later. Theskeletal framework of a grain is called adendrite and is similar to the snowflakestructure found in nature.
A typical casting shows three distinct zones(Figure 1), a thin chill-cast zone adjacent tothe mold wall formed by heterogeneoussolid nucleation at the mold wall-liquidinterface, a columnar zone formed bypreferential growth of dendrites and acentral equiaxed zone.
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Figure 2. Crystals oriented like (a) will grow furtherinto the liquid in a given time than crystals oriented like(b): (b)-type crystals will get wedged out and (a) -typecrystals will dominate eventually becoming columnargrains.
Figure 3. Dendritic growth of metallic crystals from a liquid state. A through C the dendritesnucleate and grow. The grains of a solid pure metal are depicted in D. Dendritic growth is notevident since all the atoms are identical. E shows an impure metal where the impurities havebeen carried to the regions between the dendrite arms, thus indicating the initial skeleton of themetal structure.
The progressive development of the dendritic structure is illustrated in Figure 2 below.
The cast structure is far from ideal. The firstproblem is one of segregation, as long columnargrains grow they push impurities ahead of them. If, as is usually the case, the alloy is being cast,this segregation can result in big compositionaldifferences and therefore differences inproperties between the outside and the inside ofthe casting. The second problem is one of grainsize. Fine-grained materials are harder thancoarse-grained ones. Indeed, the strength ofsteel can be doubled by a ten-times decrease ingrain-size. Obviously, the big columnar grainsin a typical casting are a source of weakness. But how do we get rid of them?
One cure is to cast at the equilibrium temperature. If, instead of using an undersaturatedsolution, we pour a saturated solution into the mold, we get what is called “big-bang” nucleation. As the freshly poured solution swirls past the mold walls, heterogeneous nuclei form in large
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Figure 4. The microstructure of a cored,cast bronze or copper-tin alloy. 12X.
numbers. These nuclei are then swept back into the bulk of the solution where they act asgrowth centres for equiaxed grains. The final structure is then almost entirely equiaxed, withonly a small columnar region. For some alloys, this technique (or a modification of it called“rheocasting”) works well.
The more traditional cure is to use inoculants. Small catalystparticles are added to the melt just before pouring (or evenpoured into the mold with the melt) in order to nucleate asmany crystals as possible. This gets rid of the columnarregion altogether and produces a fine-grained equiaxedstructure throughout the casting. This important applicationof heterogeneous nucleation sounds straightforward, but agreat deal of trial and error is needed to find effectivecatalysts.
Coring during solidification
If a molten binary alloy solidifies through a liquid plus a solidregion under equilibrium conditions, the compositions of theliquid and solid phases must readjust continuously as thetemperature is lowered. Such readjustments are affected by
the diffusion of both atomic species in both phases. But since the diffusion rate in the solid statetends to be slow, an extremely long time may be required to even out the composition gradients. In practice, cooling rates are almost always so rapid that the composition gradients remain, sucha microstructure is said to be cored because the first region solidify (the “cores”) havecompositions different from those of the last material to solidify. Since a chemical etch oftenattacks regions of different compositions at different rates, cored regions can be delineated in amicrostructure.
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Figure 5 – One way of considering the development of a cored structure. The alloy is not considered to becompletely solid until its composition line crosses the “nonequilibrium solidus” at T5.
Figure 5 shows the process by which a cored structure forms. Consider a molten alloy of over-allcomposition Co at temperature To; as it is cooled, the first solid to form has composition αl. Weassume that the solid forming at the solid-liquid interface at temperatures T2, T3, and T4 hascompositions α2, α3, and α4, that is, that its composition is given by the equilibrium solidus. Ifthe cooling rate is so rapid that each increment of solid formed maintains its initial compositions,we may picture the average composition of all solid formed proceeding along a “nonequilibrium
solidus” from α1’ to α2’ to α3’ and so on. The last liquid disappears only when the averagecomposition of the alloy, that is, when the nonequilibrium solidus crosses the vertical line at Co.
Eutectics
An eutectic reaction represents an easy way by which two (or more) constituents fit togetherduring solidification. During the reaction in a binary alloy, two types of crystal phases intergrowat a constant temperature to give a variety of characteristic patterns. Alloys to left and right of
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Figure 6 - Hypothetical binary equilibrium diagram for elements A and B whichare completely soluble in each other in all proportions in the liquid state but only to a limited extent in the solid state. TA and TB are the melting points of pure A andpure B; Te is the eutectic temperature.
the actual eutectic composition develop primary crystals of the excess phase before the eutecticreaction sets in. One of the types of equilibrium diagrams which may result when there is onlylimited solubility in the solid state is a binary eutectic diagram, illustrated schematically in Figure 6. Consider alloy CO, which exits as a single-phase liquid at point a: when it is cooled topoint b, the composition of the first solid to form is given by the other boundary of the two-phaseregion, Cα1. On further cooling to point c, a solid phase of composition Cα and a liquid ofcomposition C1 are at equilibrium. If we ignore non-equilibrium effects (such as coring) therelative amounts of the two phases in equilibrium may be calculated by the lever rule. At point c,the fraction which is α phase is (Cl - CO)/( Cl - Cα), and the fraction which is liquid phase is (CO -Cα)( Cl - Cα).
If the material is cooled still further below point c, more solid forms, and the composition of theliquid follows the liquidus down to the point e, which is called the eutectic point. With furtherextraction of heat, the eutectic liquid of composition Ce solidifies isothermally at the eutectictemperature Te. This is an invariant of the system; since the three phases are in equilibrium
during solidification of the eutectic liquid, there are no degrees of freedom. The temperature, thecomposition of the liquid phase, and the compositions of both solid phases are fixed.
The solid state microstructure having composition Ce in Figure 6 will be an intimate mixture oftwo phases. The α and β phases in such a eutectic material may be in the form of thin (of the
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Figure 7. Etched cross-section of a lead ingot
order of a micron) plates and rods or tiny particles. A material with composition between Cαeand Ce is called hypoeutectic and, in general, will have a microstructure containing primary α ina matrix of eutectic.
Safety
It is the responsibility of each TA and each student to be aware of the many hazards in thislaboratory and make use the appropriate safety equipment when performing this lab. The mainpotential hazards in this experiment are heat, cryogenic materials and hazardous chemicals. Thefollowing MSDS are available: Lead, Ammonium Chloride and Methanol.
Liquid nitrogen and dry ice expand rapidly at room temperature taking up large volumes of air. Under no circumstances, place solutions containing liquid nitrogen or dry ice in sealed containersor an explosion may result.
An important note about lead: Lead is a designated substance under the Ontario Health andSafety Act, under no circumstances should the lead metal be touched or removed from the
crucibles. Look but don’t touch!
Part 1: Solidification of a pureelement
The samples of lead have alreadybeen prepared in fireclay crucibles. Lead shot was heated to 420oC in thefireclay crucibles until melted usinga muffle furnace. The crucibles werethen air cooled and the oxideremoved from the molten metalsurface. Once a thick skin formedon the molten metal surface, theremaining molten lead was pouredout of the crucible. The molten leadwas allowed to cool to a point wheredendrites have started to form on thecrucible walls. The retaineddendrites on the crucible walls are
clearly visible in the crucibles.
Part 1: Lab Report
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Figure 9. Setup for Part 2 of the experiment.
Figure 8. Early stages of growth of an idealized metallicdendrite.
Observe the lead structures visible in the fireclay pots, sketch of a few of the dendrites as seenunder the stereo microscope. Include a description of the size and direction of growth of thedendrites with your sketches. How does the rate of heat removal from a casting affect the sizeand direction of growth?
Part 1: Lab Report (cont.)
Figure 7 shows the cross-sectionof a cast lead ingot that has beenpolished and etched with anammonium molybdate solution. Sketch the microstructurelabelling the different zones (i.e.,chill zone, columnar zone, andequiaxed zone). If the lead werecast into a chilled mold, howwould the size of the dendrites beaffected? What would be theeffect on the relative sizes of thedifferent zones of the casting?
Part 2: Solidification of theammonium chloride/watersystem
In this part, a transparent analogue
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of metal alloy solidification is used to illustrate the solidification process that occurs when metalsare cast. A slightly saturated solution of ammonium chloride will be made at 50°C. The solutiontemperature is then raised to 75°C (to stimulate superheat) and poured into a mold chilled withliquid nitrogen, minus 196°C (to simulate chill casting). The procedure is then repeated with a mold cooled to minus 50°C (to simulate sand casting).
Part 2: Procedure
Heat 50 ml of water to 50°C (ie. Hot Plate set on low with gentle stirring). Maintaining thetemperature at 50°C, slowly add enough ammonium chloride to make a slightly saturatedsolution (i.e., a few ammonium chloride crystals should remain undissolved). This isapproximately 30g.
Heat the solution to 75°C (25°C of superheat). Meanwhile, cool the solidification cell by pouringthe liquid nitrogen (to a level even with the top of the cell holder). Do not immerse any of theportion of the plastic windows in the liquid nitrogen.
Pour the solution at 75°C into the funnel positioned above the cell until the cell is just filled asdepicted in Figure 9.
Observe the solidification process. Using a magnifying glass and propping a black card behindthe cell will make the process easier to see. Initially and every few minutes squirt a littlemethanol on the windows to keep them frost-free. If the windows frost up, squirt a small amountof methanol on the windows. Once complete, clean the glassware and cell under running water,and dry.
DO NOT PUT THE COLD CELL UNDER HOT WATER – THIS MAY CRACK THEWINDOW!
Repeat the procedure but instead of liquid nitrogen use propylene glycol/dry ice mixture., andadd dry ice to the propylene glycol until the temperature is approximately minus 50°C.
Part 2: Lab Report
Include the following:Describe the solidification process for both cases (illustrations would be useful).Why are the dendrites smaller for the liquid nitrogen case?
What effect does the degree of supercooling have on cast structures, in terms of the variouszones?
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What are the limitations of this model? Is cooling with liquid nitrogen a good simulation of chillcasting? Is cooling with dry ice in propylene glycol a good simulation of sand casting?
Part 3: Solidification microstructures
Part 3: Procedure
Each specimen has already been mounted, polished and etched and is designated by the numberon the bottom of the mount. If the specimens require repolishing or etching, please contact thetechnical staff. Each specimen should be observed visually and at high and low magnificationswith a bench microscope. Systematically scan the whole section. Select regions that arerepresentative of the majority of the specimen. Sketch the observed microstructures on blankwhite paper, which will be provided. The sketches indicate whether or not a clear understandingof the basic structures observed has been achieved. Each sketch should show the principalcharacteristics of each specimen. The solidification section of the 3T04 atlas should also beexamined as it contains additional images that may be helpful. The phase diagrams for thespecimen alloys are provided at the end of this write-up. Please return the specimens to thedesiccators after observations are completed. The following three specimens are used toillustrate the coring phenomenon:
Specimen D5. 5% Tin Bronze (chill cast)This specimen was made from cathode copper and high-purity tin. The copper was deoxidizedbefore adding the tin with an addition 0.5% zinc. Pictures of the “as polished” angular oxideinclusions are in the 3T04 Atlas. Etching reveals predominately equiaxed grains. Inside thedifferent coloured grains, a dark pattern is apparent surrounding the dendrites. This representscoring and segregation, or an uneven distribution of tin in the copper. Between the dendrites andinterdendritic regions, a blue-grey delta compound (non-equilibrium) can be observed. Picturesin the solidification section of the 3T04 Atlas show the eutectoid patterns in these particles. Some shrinkage voids are apparent.
Specimen X2. 4% Tin Bronze (sand cast)Slightly elongated or columnar grains withvarying degrees of shading can be seen by eye atthe outer edge of the specimen. With themicroscope this difference in shading betweenthe grains can also be observed. The grainboundaries appear as thin black lines which inthis case follow irregular paths. A dark almostskeleton pattern can be observed inside thegrains. This represents coring and unevennessin composition or uneven distribution of tin in
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the copper. A small amount of blue-grey delta compound (non-equilibrium) can bedistinguished at the interdendritic positions. This specimen also contains shrinkage voids. Pictures in the 3T04 Atlas show the eutectoid pattern apparent in some of the particles.
Specimen X3. 4% Tin Bronze (sand cast, annealed 700°C for 2 hrs)As with X2, the grains show varying degrees of shading. However, the thin black grainboundaries are more regular in appearance. No coring or the blue gray delta compound isvisible. The annealing has allowed the grains to become uniform in composition and themicrostructure is now very similar to a pure metal. The grain size is also significantly smallerthan X2. However, porosity is still apparent within the specimen.
D5, X2 and X3 may show strain markings as a result of deformation during preparation.
The following three specimens are from the eutectic in the Cu-Cu3P system.
Specimen X5 Copper/ 8.4% phosphorus, eutectic alloy (sand cast)The surface of the specimen has an iridescent appearance. The columnar structure and severeporosity of the specimen are visible by eye. Microscopic examination reveals that the grains arecomposed of colonies of fine eutectic structure. High magnification will allow most of thestructure to be resolved. Copper-rich crystals (solid solution alpha) and crystals of copperphosphide have intergrown in a lamellar or laminated pattern. Each lamellar colony has grownradially with quite often a coarsening of the structure at the colony boundaries. The copper-richcrystals appear dark or brown as they are attacked preferentially by the etching solution, whereasthe copper phosphide appears white. Occasional, free pieces of copper phosphide may be seen.
Specimen X6 Copper/ 4.5% phosphorus, hypo-eutectic (sand cast)The specimen is dark in colour. No clear grain structure is apparent to the eye. As this alloycontains excess copper with respect to the eutectic composition (i.e., it is of hypo-eutecticcomposition), there are separate copper-rich crystals, alpha phase, together with the eutectic,which is in a distinctly coarser form than that in X5. Further, the eutectic regions have a fringeof copper phosphide. The copper-rich crystals contain a relatively small amount of phosphorusin solid solution. They are cored, and range in shade from dark blue to light brown or orange. These crystals grew first in the melt and they have developed in characteristic dendritic shape. Infact, the dendritic form is not well developed, the crystals are short and rounded. Some of theapparently isolated, round shapes probably represent regions where the cross-section has passedthrough a dendrite arm. There is a small amount of shrinkage porosity.
Specimen X7 Copper/ 10.5% phosphorus, hyper-eutectic alloy (sand cast)In effect, the reverse of X6, in that rounded dendrites of copper phosphide are set in abackground of eutectic. The dendrites (white in appearance) seem to be better developed thanthose in X6. However, the present dendrites are not cored because copper phosphide does notexist in a range of compositions. The degree of fineness of the eutectic is approximately similarto that of X5. Some porosity is also apparent.
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Part 3: Lab Report
Include labelled sketches of the various specimens indicating the different phases, and/orregions. Indicate the magnification used. Relate what is observed in each specimen to theequilibrium phase diagram.
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Copper - Tin Phase Diagramsα - fcc phase with a maximum solubility for tin of 15.8% at 520oC. An equilibrium state occurs slowly allowing100% α alloys with up to 12 percent or more tin to be created.β - bcc phase formed by peritectic reaction between solid α and residual liquid.γ − bcc phase formed by peritectic reaction between solid β and residual liquid. γ changes to an eutectoid mixture α and δ at 520oC. Beta and Gamma phases are not normal found in commercial alloys at ambient temperatures.δ − intermetallic compound with a γ brass-type structureε - is an orthorhombic structure. The eutectoid transformation of δ phase to α + ε occurs very slowly underequilibrium conditions at 350oC. Usually chill cast tin bronzes will be composed of α + δ.
Equilibrium Phase Diagram Industrial Phase Diagram - Non-Equilibrium(Below 520oC very long annealing times) (Normal annealing times)
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Experiment B: Nucleation and Growth of Crystals inAmorphous Materials
Introduction:
The goal of this lab is to introduce the student to the morphology of polymers as well as to studythe kinetics of polymer crystallization.
Background:
Various substances in nature such as cellulose, starches, proteins and enzymes are polymeric. These natural materials may be used directly from nature as is the cellulose in wood or may beprocessed into other forms such as Rayon, which is also made from cellulose. The developmentof synthetic polymers was possible through research into the various processes with which thesmall organic molecules found in fossil fuels could be polymerized to create plastics such asBakelite and Nylon.
A polymer is a long molecule or macromolecule consisting of many small units called monomersjoined end to end in a chain. The repetition of the monomer units can be linear, branched orinterconnected, resulting in a three-dimensional network.
ϕ1 - [CH2-CH2]n - ϕ2 Polyethylene
Polyethylene, the simplest hydrocarbon polymer, will be used in this discussion. The number ofethylene monomers, n, can vary from 103 to 106. However, all the small end groups ϕ1 and ϕ2occur in very small concentration. These end groups usually have little or no effect on themechanical properties and crystallization behaviour of the polymers except that they mayinfluence the chemical stability of the polymer. For some polymers, an unstable end-group whenheated or irradiated with light can initiate the degradation of the molecule.
Polymers consisting of a single repeating monomer are referred to as Homopolymers. Heteropolymers consist of several repeating monomers. Polymers composed of two differentmonomers are referred to as copolymers. The repeating monomer units can be alternatingrepeats of the two monomers, A and B.
-A-B-A-B-A-B-A-B-A-B-
or random repeats of the two monomers.
-A-A-A-B-B-A-B-B-A-B-
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Figure 1. Different types of copolymers.
Block copolymers are composed of long sequences ofone repeating monomer unit and come in many forms.
-A-B-B-B-B-B-B-B-A-A-B-B-B-B-A-
Thermoplastics (ie. Polyethylene), thermosets (ie.Bakelite) and elastomers (ie. Rubber) are the three major categories of polymers. Thermoplastics can befurther divided into amorphous or glassy andcrystallizable. When heated, thermoplastics flow in themanner of a highly viscous liquid, and do so reversibly,time and time again and on subsequently being heatedand cooled. Thermoplastics usually consist of linear orbranched polymer chains than reversibly melt anddissolve in solvents easily. If the polymer chains in athermoplastic are cross-linked then the heat stability isimproved and the flow during melting is limited.
The cross-links are chemical bonds that create bridgesbetween polymer molecules or chains. Thermosets arethose polymers whose precursors are heated to anappropriate temperature for a short time, so that they
will flow as a viscous liquid; a slow, chemical cross-linking reaction then causes the liquid tosolidify to form an infusible mass. In general, the three-dimensional network formed has shortcross-links and exhibits glassy brittle behaviour. Once cross-linking has occurred the naturalshape of the polymer is fixed and reheating will not cause the polymer to melt. These polymers
tend not to dissolve in solvents. The shape canusually only be changed by degrading thenetwork, for example, by burning.
Elastomers have long flexible chains betweencross-links and like thermosets cannot be re-heated or will not readily melt.
Polymers are considered to be either amorphousor crystalline, although they may not becompletely one or the other. Except forperfectly produced materials, most crystallinepolymers are semi-crystalline due to the varyingcontent of amorphous material. The chemicalcomposition and the arrangement of moleculesin the polymer structure are the two factors thatdetermine the crystallinity. The lack ofregularity in composition and structure reduce
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Figure 3. Schematic view of the stages of molecular arrangement inPE, (a) shows stage one, (b) shows stage two.
Figure 4. Diagram depicting the fringed micelle model.
crystallinity. Polymers with bulky side groups or extensive chain branching are also less likelyto be crystalline. The geometric arrangement of the side groups can also influence the tendencyof the polymer to be crystalline. The three variations in stereochemistry of the polymer chains,or spatial arrangement, are isotactic, syndiotactic and atactic. Polymers will not be 100% oneparticular variation but will have different degrees of tacticity. Atactic materials are difficulty tocrystallize.
Amorphous polymers are glass likeor rubbery at room temperature. Many of the amorphous polymersform brittle glasses when cooledfrom a melt such as atacticpolystyrene. The glass transitiontemperature or glass-rubbertransition is the temperature abovewhich the polymer is rubbery orexhibits elongation and bending andbelow which the polymer behaveslike a glass. Crystalline polymersalso exhibit a crystalline meltingtemperature. Note, however, thatpolymers, unlike metals, find wideapplication as completely amorphoussolids. The two amorphous forms ofgreatest interest are the glasses andthe elastomers or rubbers.
Crystallization from the molten state may be thought of loosely as consisting of two stages. Inthe first, the molecule assumes its lowest energy conformation. In the case of polyethylene (PE),the lowest energy structure is a planar zig-zag as shown in Figure 3(a). In the second stage, the
straight molecules pack together likeparallel rods. Figure 3(b) shows theend view of the packed molecularchains.
The earliest and simplest model of acrystalline polymer is the fringed-micelle model (Figure 4). Itencapsulates the very importantstructural fact that although themolecules lie parallel to each otherwithin the crystal, they do so overlengths far shorter than the total lengthof the molecule: the molecule may
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Figure 7. Transmitted polarized lightmicrograph showing the polyhedralspherulites.
Figure 6. TEM image of polyethylenelamella.
Figure 5. Folded-chain model showing a single crystal ofpolyethylene.
contain 10,000 carbon atoms, but the length of the chain within the crystal may be only 100carbon atoms. In the fringed-micelle model, a molecule meanders out of one crystal into theamorphous fraction and then into a neighbouring crystal. In this way, it passes through manycrystals. All the crystals and the adjacent amorphous regions are thus firmly woven together bythe long, thread-like macromolecules.
The random arrangement of adjacent crystalsin the fringed-micelle model does not accordwith evidence obtained from microscopy. Inthe folded-chain model, the polymer chainextends for long distances within a polymercrystal by folding at the top and bottomfaces of the crystal.
The electron microscope shows the crystals inpolyethylene to be very thin twisted lamellae laid oneupon another. The lamellae are too small to be observedwith the light microscope. Light microscopic examinationof thin films using crossed polarizers reveals complexpolyhedral objects known as spherulites (Figure 7) whichare in fact a complex ordered aggregation of the sub-microscopic crystals. The nucleation and growth of thesespherulites are very similar to that of pearlite in steel.
Some polymers remain completely amorphous even whencooled from the melt extremely slowly. The main effectof cooling the melt is to decrease the violence of thethermal agitation of the molecular segments. In the melt,segments of each molecule change place by thermallyactivated jumps. As cooling is continued, a temperature isreached at which the rate of segmental movement isextremely sluggish, and then, on further cooling, themovement finally stops. Polymers consist of longmolecules tangled in a liquid-like manner, but with thecomplete absence of the rapid molecular motion which istypical of a liquid. This is the glassy state and thetemperature at which the molecular motions stop is calledthe glass transition temperature. When a polymer isheated through its glass transition temperature, there is anincrease in thermal expansion coefficient and a drop in themodulus of the order of 103.
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Safety:
It is the responsibility of each TA and each student to be aware of the many hazards in thislaboratory and make use the appropriate safety equipment when performing this lab. The mainpotential hazards in this experiment are intense light, heat, cryogenic materials and hazardouschemicals. The following MSDS are available: Methanol, 95%Ethanol, Polyethylene glycol.
Procedure:
The crystallization of thin films of polyethylene glycol will be observed with a light microscope . The apparatus consists of a light transmission microscope equipped with hot and cold stages witha digital camera capable of video capture. The two stages have through-thickness holes in thecentre to allow the passage of light and side holes in which a thermocouple can be inserted. Thehot stage is a copper block equipped with a resistance heater; it works simply by turning thepower supply on and off manually to maintain the desired temperature. It requires some gettingused to and should be practised until you are comfortable with the system. The cold stage isanother copper block with channels running through it. Cooling is achieved by immersion of thecopper block in cryogenic liquids either liquid nitrogen or methanol and dry ice.
It is necessary to do a calibration of the microscope system for each magnification that is usedduring the experiment in order to determine the relationship between actual lengths in thecrystallizing polymer and lengths measured on the computer screen. A scale with very finedivisions will be provided, which can be observed in the microscope for this purpose.
Crystallization
The polymer will be melted using the hot plate provided in a water bath in order to avoid burningthe polymer. The water must be kept boiling during the entire experiment to ensure that themolten polymer stays at a constant temperature. The basic procedure for the hot stage is asfollows:
Keep the slide and the cover glass to be used on the stage until both at the desired temperaturethen start recording the video file.
Dip a clean glass rod into the polymer melt then remove it and let a few drops fall from the rodonto the glass slide.
Quickly cover the drops with the cover glass and press down firmly in order to get a thin film;return the slide to the stage as quickly as possible. This will require some practise.
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On the monitor screen, spherulite growth can be observed. Perform the above steps a few timesfor each temperature and try to record the growth of at least a couple spherulites from nucleationto the end.
When using the cold stage, do not cool the glass slide down prior to preparing the thin film ascrystallization occurs very quickly and it will be impossible to transfer the slide to the stagebefore the phase transformation is complete. Prepare the thin films as above on slides at roomtemperature (or slightly warmer) then transfer the slides to the cold stage.
The spherulite growth should be observed from 45°C down to minus 30°C at 5°C intervals. Again, it is important to perform the experiment a few times at each temperature.
Once the experiment is finished, measure the rate of growth of the spherulites by plottingspherulite diameter vs. time. This can be done simply by placing a ruler on the monitor screen andmeasuring the diameter of a spherulite at different times. It may not be possible to do this accuratelyat temperatures where the spherulites nucleate profusely and are thus very small. In these cases, therate of spherulite growth can also be measured by observing how much the boundary of a spherulitemoves in a given time.
Lab Report
The report should contain the following:
A diagram of the crystallization behaviour of the polymer at each temperature. Consider theconcept of nucleation and growth that have been studied in pearlite transformations experiment.
Plots of the spherulite diameter vs. time for a couple temperatures. Are these plots linear?
Plot a spherulite growth rate (μm/sec) vs. temperature. Discuss why the rate of growth varies asit does in terms of the melting point and the glass transition temperature.
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Experiment C: Phase Transformations and Age Hardeningin Non Ferrous Alloys
Introduction:
This lab is divided into two parts:
Part 1: Diffusion Controlled Solid State Phase Transformations and MicrostructuresIn the first part of the experiment, the transformations in the solid state of some typical binaryalloys involving diffusional nucleation and growth will be examined.
Part 2: Strengthening by Precipitation In this part of the experiment, the intention is to demonstrate the metallographic characteristicsof alloys which have been subjected to precipitation heat treatments.
Background:
The majority of phase transformations that occur in the solid state take place by thermallyactivated atomic movements. One of the most fundamental processes that controls the rate atwhich many transformations occur is the diffusion of atoms. While the phase diagramssummarize the thermodynamic stability of a phase or phases coexisting for specific temperatureranges, the kinetics of diffusion controlled nucleation and growth determine the rates of thesephase transformations.
The different types of phase transformations that are possible can be broadly divided into thefollowing groups: (a) precipitation reactions (b) eutectoid transformations (c) ordering reactions (d) massive transformations and (e) polymorphic changes. Both precipitation and eutectoidtransformations involve the formation of new phases with a composition different from that ofthe matrix or parent phase and therefore long range diffusion is required. The remaining reactiontypes involve the formation of a new phase with the same composition as the parent phase, but adifferent crystal structure.
The decomposition of a phase into one or more phases may be divided into three stages:
1) the formation of nuclei of the new phase
2) the growth of these nuclei
3) the coarsening of the precipitate without changes in its volume fraction
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Nv = CD expT
−+⎡
⎣
⎢⎢
⎤
⎦
⎥⎥
A cg egk
3 2σ / ( )Δ
Figure 1. Rate of transformation curve compared with TTT curve.
The rate of homogeneous nucleation of a precipitate can be expressed in a simplified form:
where
Nv is the nucleation rateD is the diffusion coefficientC is a constantA is a geometrical constantσ is the interfacial energyΔgc is the specific chemical free energyge is the elastic strain energy associated with specific volume change
This expression accounts for the observed minima in the incubation time in the Time-Temperature-Transformation (TTT) curve (see Figure 1). Because of the high value of energy ofactivation for diffusion of substitutional solutes, the value of D and thus the rate of nucleationbecomes very low at low temperatures. At temperatures close to equilibrium, Nv becomes lowbecause there is an inadequate contribution from the chemical driving force to overcome thework of nucleation involved in the formation of the surface of a new phase. Nucleation in solidsis invariably heterogeneous. Suitable nucleation sites are non-equilibrium defects such as excessvacancies, dislocations, grain boundaries, stacking faults, twin boundaries and interphaseboundaries.
The term phase signifies crystallites have the same crystal structure and similar chemical
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Figure 2. Some preferred nucleation sites.
composition. Pure metals are single-phase. Alloys can be single-phase but more often than notconsist of more than one phase. The various phases form in an alloy during thermal coolingfrom the molten state or during thermal treatment as a consequence of the temperaturedependence solubilities of component elements. If the metal lattice contains more foreign atomsthat can dissolve at a given temperature the alloy becomes supersaturated resulting in formationof finely dispersed particles of a second phase known as precipitates.
Some preferred heterogeneous nucleation sites are indicated in the schematic shown in Figure 2.
The crystal structure of the matrix and the precipitating phase influence the interfacial energyand three important cases arise:
Coherent nucleation arises if the crystal structures andlattice parameters of both phases are closely similar.Coherent precipitation will occur, e.g., Ni3Al in Ni, if thedispersed particles are in parallel crystallographicorientation with the matrix.
Semi-coherent nucleation arises if for example the α andβ phases are related in such a way that an interface can bebuilt from well-defined line defects such as interfacialdislocations, e.g. θ’ phase in the Al-Cu system
Non-coherent nucleation arises if for example the structures of the α and β phases are sodifferent that the interface has a structure similar to that of a high angle grain boundary e.g., θphase in the Al-Cu system. In this case, a random orientation between the two phases can beexpected.
In the analysis of precipitate growth from supersaturated solid solutions, the possible limitingfactors to consider are the rate at which atoms are brought to or removed from the interface bydiffusion and the rate at which they cross the interface. The interface reaction is likely to be therate limiting step during the early stages of growth since the diffusion distance tends to zero inthis situation. As the particle grows to about a micron in size, lattice diffusion is likely to beslower step as the matrix is depleted of solute.
Widmanstatten Structures and Precipitation
Solid state transformations involve, over a temperature range during cooling, the formation of asecond phase among the crystals of an earlier formed phase. It takes place because of changeswith temperature in the range of existence of solid solution or compound phases. Frequently,this type of transformation gives rise to a Widmanstatten structure. The characteristic feature ofthe true Widmanstatten structure is that needles or thin plates of the second phase precipitateintrude on a system of planes in the matrix crystals, so that random cross-sections reveal ageometric arrangement of variously inclined sets of essentially parallel precipitates sometimes
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taking the form of rosettes or small skeletons, which arrange themselves with some degree ofsymmetry. In addition, precipitation usually occurs at the grain boundaries of the original phase,these regions being the easiest place for precipitation to start. Under some circumstances,precipitation may occur only at grain boundaries. The boundary precipitates may adopt the sameshape as those inside the grain; alternatively, they may occur as partial or complete networks. Ifprecipitation is limited, the boundary precipitates may be merely small lens-shaped particles.
Eutectoid Transformations
In the eutectoid reaction, a solid solution or a compound decomposes in the solid state(theoretically at a constant temperature) to form a mixture of two other phases. The eutectoidstructure often shows a relatively fine, laminated or lamellar aggregate of two constituents, butmore irregular patterns may occur. At the actual eutectoid composition, the alloy is involvedcompletely in the reaction, but with other alloys in the range, the eutectoid reaction is precededby the separation of one of the constituents as relatively massive precipitates.
The formation of eutectoid structures starts at a number of points on the grain boundaries of theoriginal phase and gradually spreads inwards in the shape of nodules or colonies. The eutectoidin steel is the most important example and the following discussion refers to plain carbon steel. At elevated temperatures, the carbon is dissolved in the iron giving the gamma solid solution(fcc, austenite). At lower temperatures, the iron changes its atomic pattern and can hold verylittle carbon in solution. On cooling, the steels change into a mixture of almost pure iron (bcc,alpha phase or ferrite) and iron carbide, Fe3C (known as cementite). The actual eutectoidcomposition occurs at 0.8% carbon and austenite of this composition on cooling changescompletely to the eutectoid mixture of ferrite and cementite. The structure is lamellar and isknown as pearlite.
In steels, with less carbon than 0.8% (hypoeutectoid alloys) ferrite is first precipitated until thematrix reaches a carbon content of 0.8% when it transforms to pearlite. In a correspondingmanner, hyper-eutectoid steels containing more than 0.8% carbon involve the previousprecipitation of iron carbide. The formation of pearlite, as with other eutectoid structures, isaffected by the actual temperature of transformation. Thus, just below the theoreticaltransformation temperature, pearlite forms very slowly and the structure is relatively coarse. Lower temperatures result in more rapid transformation and finer structures.
In continuously cooled specimens, the pearlite spacing is variable as the transformation occursover a range of temperature. The effective range becomes lower as the cooling rate is increased,until eventually the rate becomes too high for pearlite to form at all. It should be appreciatedthan the spacing of lamellae in pearlite always appears variable in a cross-section because of therandom orientation of the various colonies of pearlite.
Part 1: Procedure
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Each specimen has already been mounted, polished and etched and is designated by the numberon the bottom of the mount. If the specimens require repolishing or etching, please contact thetechnical staff. Each specimen should be observed visually and at high and low magnificationswith a bench microscope. Systematically scan the whole section. Select regions that arerepresentative of the majority of the specimen. Sketch the observed microstructures on blankwhite paper, which will be provided. The sketches indicate whether or not a clear understandingof the basic structures observed has been achieved. Each sketch should show the principalcharacteristics of each specimen. Refer to the questions in the Lab Report section as a guide towhat you should be thinking about when observing the various specimens. Phase diagrams forthe binary alloys are provided at the end of this write-up. Please return the specimens to thedesiccators after observations are completed.
Part1: Binary Alloys
Specimen X10, Brass with 58% Copper and 42% Zinc (sand cast)This shows a well defined “Widmanstatten” structure in which crystals of a solid solution haveprecipitated at the boundaries of the original grain of beta compound, and as needles inside thegrains. After etching, the alpha phase usually appears light against the darker beta phase, thelatter showing some contrast from grain to grain. The alpha phase appears slightly pink orcopper coloured, whereas the beta is yellow in appearance. It should be noted how the shape ofthe precipitated alpha varies according to whether the cross-section cuts along the length of theneedles or transversely.
Specimen X13, same composition as X10 (sand cast, reheated to 800°C for 1 hour, furnace-cooled to room temperature)Note the change in grain size and the size of the alpha phase. This alloy system permits thestructure and mechanical properties to be varied by heat treatment in the solid state. The usualprocedure is to quench the alloy from the single-phase region and then to release the secondphase as extremely tiny precipitates either by aging at room temperature, or by reheating to amoderate temperature. The temperature required varies with the alloy system. Fine precipitationgives greater hardness and strength than that obtained from coarse precipitates. However, theeffect on properties is not very significant with this brass.
Specimen B17, Brass with 49% Copper and 51% Zinc (chill cast)This is a pure binary alloy made from cathode copper and 99.99% zinc which was cast into 3/4"fingers. A large amount of gamma phase is present in this alloy, especially as a grain boundarynetwork. In addition, clusters of angular inclusions of ZnO may be detected. Etching revealsradical columnar grains. The grains consists of the beta phase which has an almost continuousgrain boundary network of gamma. Numerous star-shaped crystals of the latter are precipitatedinside the grains. The gamma is somewhat darkened by the etching. Gamma precipitation is inhibited in a narrow rim around the exterior surface of the casting.
Part1: Eutectoid Transformations
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Specimen X17, 0.8% carbon steel, eutectoid composition (rolled bar, heated for 1 hour at800°C then furnace cooled) At lower power, small grains of varying shades of grey, blue and brown are seen. High powerexamination resolves all the structure as relatively coarse pearlite. The outer surface of thespecimen has been decarburised, that is, has lost carbon, during the annealing and the extremesurface may be almost pure ferrite.
Specimen X19, 0.35% carbon steel bar ( furnace-cooled from 870°C)X19 illustrates a hypo-eutectoid structure in comparison with a pure eutectoid structure. Areasof coarse pearlite are arranged in a background of ferrite. The grain boundaries in the ferrite arevisible. With less carbon in the steel, the structure would be expected to consist of ferritecontaining small areas of pearlite. With greater amounts of carbon, there would be less ferriteand it would show more definitely as a network around the pearlite. In such steels, the excessphase does not precipitate in a Widmanstatten form, unless the material has a very coarse grainsize, as a result of casting, overheating, or is cooled relatively rapidly through the transformationrange.
Specimen X24 – 0.35% carbon steel (normalized from 870°C)This structure should be compared with that of X19. The most apparent difference is the shapeof the ferrite, which in X24 occurs in block-like and needle formation, and approaches aWidmanstatten pattern. The difference between the normalized and annealed structures of theferrite is rather noticeable in the present steel (possibly because of the relatively hightemperature of treatment). However, it should be appreciated than the structure in hypoeutectoidsteel is governed by the carbon content and the specific rate of cooling. In some cases, thedifferences may be less pronounced. It is to be expected that less ferrite will occur in thenormalized steel because less time for separation is allowed by its more rapid cooling. However,in the present specimens, comparison is difficult because of the different type of formation and adifference in quantity is not readily apparent. Finally, pearlite is usually finer in normalizedsteels, but grain comparison is difficult with the present specimen as the growth pattern has beensomewhat different and especially as fine ferrite needles are present in the pearlite regions ofX24 and give an impression of coarseness. X24 emphasizes primarily the overall pattern of theferrite formation. Alloys undergoing eutectoid reactions also lend themselves to heat treatment. The most important example is the quench-hardening and tempering of steel. In general, aircooling (normalizing) from the austenite region gives finer structure than slow cooling(annealing), and in consequence the hardness and strength are greater in the normalizedcondition.
Another well-known eutectoid occurs in the copper-aluminum system, the eutectoid alloycontains 11.8% aluminum. The reaction involves the breakdown of a beta compound to the
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alpha solid solution of aluminum in copper and another compound, gamma. This alloy alsoundergoes a martensitic transformation, if rapidly cooled from a high temperatures.
Specimen X21, Copper with 11.8% aluminum, eutectoid alloy (chill cast, reheated for 1hour at 900°C and slowly cooled)The surface of this specimen has an iridescent appearance. The material has completely transformed to the eutectoid, which consists mainly of a fine lamellar arrangement of the alphasolid solution (light) and the gamma compound (dark). However, in some regions, the alpha hasa coarser more irregular pattern. Some porosity and cracking is evident.
Specimen X23, Copper with 11.3% aluminium, hypo-eutectoid alloy (sand cast, annealedat 900°C, slowly cooled)Rather jagged needles of alpha are set in a background of eutectoid. The eutectoid is mainlyirregular and very coarse. The gamma particles are large enough to be clearly seen as grey incolour. Sometimes, small brown regions are found in the eutectoid areas, which represent areaswith a very fine lamellar structure.
Part 2: Strengthening by Precipitation
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Part2: Additional Background
Solid state precipitation is the most common strengthening mechanism employed in non-ferrousalloys. One of the classical examples is the aluminum-copper alloys developed for the aircraftindustry. The specimens observed in this experiment have been taken from other alloy systemswhich also have important technological applications. The size of the precipitates formed inalloys hardened by heat treatments that provide optimum properties are generally too small toobserve by light microscopy. Electron microscopy is usually necessary to resolve the fineparticles. However, it is possible to appreciate the existence of unresolved precipitates sincethey invariably alter the etching characteristics of an alloy. Such effects are demonstrated inmany of the specimens to be observed in this experiment. In general, first look for evidence thatprecipitation has occurred and second for where that precipitation has taken place, e.g. at grainboundaries or grain interiors. Since precipitation is carried out primarily for strengthening, it isimportant to relate the precipitation heat treatments to the improved (or degraded) mechanicalproperties achieved. Tables of typical hardness values have been included to assist you in thisevaluation. Before commencing this laboratory study, you should read the appropriate referencematerials listed for this course so that a good appreciation of the phenomenon of precipitationhas been achieved.
Part 2: Procedure
Each specimen has already been mounted, polished and etched and is designated by the numberon the bottom of the mount. If the specimens require repolishing or etching, please contact thetechnical staff. Each specimen should be observed visually and at high and low magnificationswith a bench microscope. Systematically scan the whole section. Select regions that arerepresentative of the majority of the specimen. Sketch the observed microstructures on unruledwhite paper, which will be provided. The sketches indicate whether or not a clear understandingof the basic structures observed have been achieved. Each sketch should show the principalcharacteristics of each specimen. Refer to the questions in the Lab Report section as a guide towhat you should be thinking about when observing the various specimens. Some observationsmay only be possible by high resolution light microscopy, the 3T04 atlas contains additionalimages that may be helpful. Please return the specimens to the desiccators after observations arecompleted.
Description of Specimens:
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Set sets of alloys are available for observation: A cast aluminum alloy with 10% magnesiumA cast copper alloy with 1.8% beryllium and 0.3% cobalt A brass with 20% zinc, 6.5 % nickel and 1.75 % aluminum
Cast Aluminum-Magnesium AlloyAnalysis Mg 10%/ Fe 0.25%/ Si 0.1%/ Mn 0.08% / Cu 0.01% / Ti 0.1%/ Be 0.005%
Specimen F14. (Sand cast, as-cast condition)In the polished condition, numerous needles of FeAl3 as well as a considerable quantity of largerand often angular Mg2Si can be seen. The FeAl3 is purple-grey in colour, and the Mg2Si is eitherblue or rather dark, depending on the time of contact with water during polishing (water tends todarken Mg2Si). Porosity is also apparent. The angular-shaped pieces of a beta Al-Mg phase are pale grey in colour and well outlined. This phase occurs at the grain boundaries of the matrixphase, and at interdendritic positions inside the grains. Etching tends to attack Mg2Si, but has nosignificant effect on the FeAl3. Both these constituents are found to occur mainly at the grainboundary positions. The grain boundaries of the solid solution matrix phase show up as thindark lines and some coring may be evident.
Specimen F15. (Same as F14 then solution treated for 8 hours at 430°C, quenched in hotwater)The purpose of the heat treatment is to homogenize the material, that is to remove the coring, andto dissolve the beta phase which is introduced by reason of coring. The alloy is normally used inthe solution-treated condition, although the material is losing favour because of its susceptibilityto stress-corrosion. The FeAl3 and Mg2Si and matrix grains can be seen as before. In thesolution-treated condition no beta phase should be present, but in practice, a varying proportionmay remain. In the present material, there is residual beta, probably because of the relatively short heating time. However, no coring can be detected even after long etching in the mixedreagent.
Specimen F16. (Same as F15 then aged 4 years at 250°C)The alloy is never used in this condition (when it is relatively brittle) but the treatment has beenapplied to demonstrate the precipitation habits. FeAl3 and Mg2Si are again present as in theother specimens. Etching reveals the cellular-shaped grains marked out by fine precipitation. Inside the cells, fine but noticeable precipitation is also to be found. The extent of the precipitation is somewhat variable, in that some regions are light in appearance and contain littleprecipitation, whereas many other regions are dark in appearance as a result of extensiveprecipitation.
Cast Beryllium BronzeAnalysis Be 1.8% / Co 0.3%
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Specimen S1. (Chill cast)In the polished condition, purple-grey particles of the Co-Be compound may be seen, and in thebackground an interdendritic constituent is faintly visible. Patches of interdendritic porosity arealso detected frequently. On etching, the overall grain pattern may be columnar, alternativelythere may be a preponderance of equiaxed grains, depending on the position of the specimen inthe original bar. The grains consist of cored alpha with an interdendritic constituentcorresponding to the original beta phase. Mostly the latter constituent has a fine granularappearance, but sometimes there is a lamellar pattern, it is not possible to assess the degree oftransformation which it has undergone. In addition, it should be noted that considerable fineprecipitation has occurred in the alpha phase in the form of veining. This precipitation isparticularly marked alongside the regions of “beta” phase. The Co-Be constituent is somewhatlightened in colour by etching, and is found to be mainly in interdendritic positions.
Note about Wrought and Heat Treated Specimens:Fabrication defects in the form of elongated oxide and even voids are present in some of thespecimens. These oxides are associated with regions containing very small grains. This effect isusually attributed to the mechanical inhibition of grain growth by the oxide particles. Thegeneral grain size also tends to be somewhat variable. A general feature of all the wroughtspecimens is that in the polished condition, elongated particles of the Co-Be constituent (purplegrey in colour) may be seen.
Specimen S3. (Hot rolled, Solution treated for 1 hour at 790°C, oil quenched)The etched structure is found to consist of twinned grains of alpha phase, showing somevariation in size across the section. Relatively small particles of the Co-Be phase now becomeapparent, in addition to the larger pieces which can be detected in the polished state.Some grains in the body of the material may show evidence of strain lines, but as this featurediminishes in extent with repeated polishing and etching, it is probably attributable todeformation introduced during sectioning, rather than to that resulting from a straighteningoperation. However, such an operation has probably been responsible for the bent twins andstrain lines which may be detected near the exterior surfaces.
Specimen S4. (Same as S3 then aged at 315°C for 2 hours)The aging treatment applied is that recommended commercially. The precipitation is manifestedin the grains by criss-cross markings of diffuse appearance. The precipitate is probably of anintermediate type and may be designated gamma prime. In addition, a small amount of grainboundary precipitation has occurred. This takes the form of thin networks, of brown colouration,at some grain boundaries. In these preferred regions, the structure is unavoidably overaged. Theprecipitation here is the type described as discontinuous, and the strains set up in earlier stageshave caused local recrystallization with the precipitation forming a type of eutectoid structure. The precipitate in these regions is probably of the equilibrium type, namely gamma.
Specimen S5. (Same as S3 then over-aged at 350°C for 6 hours)This shows a similar type of structure to that of S4 but the markings in the grain are more
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noticeable and there is a greater quantity of discontinuous precipitation at grain boundaries.
Specimen S9. (Same as S3 then over-aged at 350°C for 24 hoursThe structure is essentially similar in form to that of S5.
Variations in the grain size of the original S3 stock may result in grain size differences inindividual specimens of S3-S5 and S9.
Results of Hardness Tests on the Heat Treated Beryllium BronzesNote: Variations in grain size may cause some scatter in the results; moreover, care should betaken to avoid regions containing defects.
The figures in brackets represent values taken from the manufacturer’s literature.
Specimen # Condition Average Vickers Hardness Number [30 kg load]
S3 Solution treated 92 (100)S4 S3 then aged 315oC, 2 hours 388 (370)S5 S3 then aged 350oC, 6 hours 335 (320)S9 S3 then aged 350oC, 24 hours 330
Precipitation-hardening brass (20% Zn, 6.5% Ni, 1.75% Al)
Note: particles of blue-grey inclusions are also present in this material, being of elongated shapein the wrought specimens. Their appearance is essentially unchanged by etching. It has beensuggested that these represent relatively massive particles of Ni-Al precipitate responsible forthe hardening effects in this alloy, and that their presence in wrought material results from incomplete solution. However, in the cast condition, they do not appear to be part of theessential microstructure. Moreover, in the wrought and heat treated alloys, many of the particlesare angular in shape, whereas it would be expected that they would become rounded if they weresoluble in the matrix. It is, therefore, concluded that the inclusions are of foreign materials. Afrequent feature of the wrought specimens also is presence of central, elongated ‘strings’ of darkinclusions, which probably represent an extrusion defect. Segregation, presumably of theslowly-diffusing nickel, is sometimes indicated by a light- and dark-banding effect in the etchedsections, apparent on the macro scale in V10. The grain size is medium to large and essentiallyuniform.
A final point to note with this material is that, whereas the etched structure of solution-treatedspecimens is relatively bright, the surface of specimens in which precipitation has occurredbecomes dark and often of tarnished appearance on etching.
Specimen V6. Extruded and solution treated for 1 hour at 850°C, water quenched
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The structure shows regular twinned grains of the alpha phase with no evidence of residual coldwork, except possibly for slight curving of the twins and occasional strain lines near the exteriorsurfaces. These features are consistent with the small increase in hardness from the interior tothe outside of the section (see table), probably resulting from slight cold straining duringstraightening operations after solution treatment and quenching.
Specimen V7. As V6 but precipitation treated for 2 hours at 500°CThis represents the optimum precipitation treatment. On etching, the surface of the specimenbecomes dark in appearance. Microscopic examination shows relatively dark grains, but withgrain boundary regions appearing light in contrast. High-power examination (especially with anoil-immersion objective) reveals very fine grain boundary precipitation of the Ni-Al complex. The light regions could represent impoverishment or denudation of the grain boundary zones ofthe solute atoms, as a result of the localized precipitation. However, as the effect is on arelatively gross scale, this seems unlikely, and it may represent regions which have acted ascathodes to the boundary precipitates, during etching.
Apart from the general darkening of the grains, no other evidence of precipitation is to be foundin the grains, except for occasional needle markings which are seen more strongly in certainsubsequent specimens. The overall effect of the light grain boundary regions is best observedafter relatively deep etching. However, the boundary precipitates and occasional needlemarkings in the grains are clearer after shorter etching times, when it will be observed that thereis an inner white zone closely associated with the boundary precipitates. This zone is most likelyis a solute-denuded zone. In addition, it will be noticed that near the original bar surfaces strainlines are apparent. Some degree of precipitation has clearly occurred on active slip planes andcaused them to etch up. This effect is demonstrated more markedly in later specimens whichhave been definitely cold worked before precipitation treatment.
The observation of surface strain lines in the present specimen confirms the conclusion drawn inrespect of V6 that slight cold work has been introduced subsequent to solution treatment,probably during straightening operations. The hardness difference between surface and interioris apparently no longer maintained.
It will be seen that the gross light-boundary zones do not seem to be present near the externalsurface of the specimen, that is in the region of the strain lines. Thus, the slip-plane precipitationapparently affects the etching characteristics and the grain boundary regions. A similar effect isfound throughout the structure of specimen V10 which has been definitely cold worked beforeprecipitation treatment.
The following point is mentioned as a matter of general interest. When material, such as thepresent brass, has been strained after solution treatment, there is a possibility of matrixrecrystallization during precipitation at temperatures of the order of 500°C. Clearly, this willonly occur at certain critical combinations of strain, precipitation rates, and temperature. Thereseems no evidence of this occurring during precipitation treatment of V6 (in which the surface is
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slightly strained) or V9 (which has been definitely cold worked). However, in some materialstrained to a degree of intermediate between V6 and V9, matrix recrystallization was detected(after treatment at 500°C) in some regions, especially at the outer surfaces, but also in theinterior. These recrystallized regions were lighter etching than the remainder, and often retainedthe lines of precipitates that formed in the original matrix, that is, at the positions of originalgrain and twin boundaries.
Specimen V8. As V6, but precipitation treated for 18 hours at 500°CThis treatment corresponds to over-precipitation although the overall effects are not verynoticeable. The specimen becomes dark on etching. Compared with V7, the grain boundaryprecipitates in V8 are possibly larger, although it is difficult to be positive about this. There is awhite zone at grain boundaries but this is more irregular in contours and is not so thick as that inV7. The inner light zone associated with the boundary precipitates is absent in V8. The needleprecipitates in the grains are more numerous than in V7, although they are still not prolific. Occasionally, small needles are clustered together. In respect of the larger needle markings, itshould be noted that sometimes these have lighter surrounds. In addition, it should be observedthat the grains generally have a granular appearance suggesting the occurrence of general over-precipitation. As in V7, strain lines are apparent near the outer surfaces of the bar, and again, thelight boundary zones diminish markedly in these regions.
Specimen V12. As V6, but precipitation treated for 18 hours at 600°C(note break in number sequence)The etched surface is lighter in appearance than that of V8. The structure shows generally darkgrains, with light boundary regions somewhat as in V8, and clear grain boundary precipitates. The earlier mentioned coarse precipitates of needle-like form are present in the grains, and thesemay have light surrounds. Examination a very high magnifications indicates that general precipitation (as detected at an earlier stage in V8) has developed to resolvable sizes. This fineprecipitation is often rather irregular form, but sometimes as fine needles. Some strain lines stillapparent in the outer regions of the section and occasionally twins may be curved, but in thisspecimen, the light zones do extend to the bar surfaces.
Specimen V9. As V6, but cold drawn after solution treatment (to “half-hard” condition)Consideration of the various hardness results suggests a degree of cold working in the surfacelayers corresponding to 10-15% reduction. The structure consists of twinned grains with slightcurving of the twins as a result of the final cold working operation. There are occasional tracesof strain lines at the extreme outer surfaces of the bar. As in V6, there seem to be someprecipitates at grain boundaries.
Specimen V10. As V9, but precipitation-treated for 2 hours at 450°C
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A lower precipitation temperature has been used than for V6, because the present material is coldworked, and the optimum temperature decreases with the degree of previous cold work. Theetched surface has a relatively bright appearance but marked banding is apparent on the macroscale. Extensive strain markings are visible throughout the twinned grain structure. As a resultof precipitation on the slip planes, these regions have become particularly responsive to etching. Examination at high magnification shows the presence of fine grain boundary precipitates, butthe light zones at grain boundaries (seen in V7., V8., and V12.) are absent.
Specimen V11. As V9, but precipitation treated for 18 hours at 450°CThe strain line markings are still present; the grain boundary precipitates are more definite, andoften associated with very thin light zones running along the rain boundaries; these zones are ofthe form typical of denuded zones. The general appearance of the slip traces in the grainsindicates that precipitation is relatively well advanced in these regions, although the precipitateshave not reached a resolvable size. he etched surface is dark in appearance comparable with thatof V8.
Specimen V13. As V9., but precipitation treated for 18 hours at 500°CThe general appearance of the etched surface is comparable with that of V12. At low power, thegrains appear moderately dark, with light grain boundaries. Strain lines as such are absent butthey appear to have imposed some residual pattern on he structure. In addition, twin boundariesmay be curved. High power examination shows resolvable particles of precipitates at grainboundaries. Twin boundaries having associated light zones may contain precipitates; and someapparently well cured twin boundaries appear to contain almost continuous lines of precipitate. Whether the latter lines, in fact, represent twin boundaries is not quite certain, because theysometimes occur as single lines of a form not commonly associated with twins. In addition tothe foregoing effects, general precipitation in fine form is apparent throughout the grains, and thepattern adopted appears to have been affected by the initial strain lines.
Hardness Values of V-series Specimens
Specimen # Condition Average Vickers Hardness Number [30 kg load]
V6 solution treated 70 (centre); 80 (1.6 mm from surface)V7 V6 then precipitation 2 hrs, 500oC 220V8 V6 then precipitation 18 hrs, 500oC 193V12 V6 then precipitation 18 hrs, 600oC 157
V9 V6 then cold drawn 115 (centre); 132 (1.6 mm from surface)V10 V9 then precipitation 2 hrs, 450oC 210(centre); 245 (1.6 mm from surface)V11 V9 then precipitation 18 hrs, 450oC 214 (centre); 225 (1.6 mm from surface)V13 V9 then precipitation 18 hrs, 600oC 160
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Lab Report
Part1:
A) Explain what is meant by the following heat treatments and list them in order of the rate ofcooling achieved:
(i) sand casting(ii) chill casting(iii) normalizing(iv) furnace cooling(v) quenching
B) Include sketches of all ten microstructures. In addition, sketch the relevant part of thecorresponding phase diagram and explain how the observed microstructures developed duringcooling.
C) The lab report should address the following points:
Specimens X10 and X131) Compare and contrast these structures and explain the reasons for the different
morphologies in terms of nucleation and growth concepts.2) The alpha phase is in the form of Widmanstatten precipitates in X10. What is the
history of the term “Widmanstatten”?3) Why do Widmanstatten precipitates show preferred directionality?
Specimen B171) The precipitated phase (what is it?) is present in what two forms?2) Explain what is meant by phrase “precipitation occurs when a solvus line is
crossed”.
Specimens X21 and X171) Describe the morphology of the eutectoid.
Specimens X23, X24 and X19
1) What is the difference between a hypoeutectoid, a hypereutectoid and aneutectoid alloy?
2) What is meant by the term “pro-eutectoid constituent”? What are they in thesethree specimens?
3) What two differences exist between X19 and X24? Explain why they haveoccurred.
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Lab Report (cont):
Part2:
A) Explain what is meant by the following terms:
(i) precipitation(ii) inter- and intra-granular precipitation(iii) again(iv) over-aging(v) twinning(vi) coring(vii) solution treatment(Viii) hot and cold working
B) Examine all of the specimens provided but only include sketches for the following andexplain the differences between specimens in the same series:
All 3 of the “F” series specimens (F14, F15, and F16)S1, S3, and one of (S4, S5, S9)V7 or V8 and (V9, V10, V12)
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EXPERIMENT D: PEARLITE TRANSFORMATIONS NUCLEATION AND GROWTH IN A EUTECTOID STEEL Materials 3TO4 Laboratory, 2000
1. Introduction
Part of the Fe-C phase diagram is shown in figure (1). When the austenite phase is cooled below 727˚C (A1 temperature), it transforms into ferrite and cementite. The reaction may be written as:
( → " + Fe3C
This reaction is an example of a eutectoid transformation. In this lab we are interested in studying the nucleation and growth of the new-phases.
Figure 1
2
2. BACKGROUND 2.1 Pearlite Consider, the specific example of a steel containing 0.78 wt% C. When the temperature drops below 727˚C, the austenite phase is no longer stable with respect to ferrite and cementite. The new phases will nucleate on grain-boundaries. Consider the formation of a ferrite nucleus. The solubility of C in ferrite is less than 0.05 wt%. Consequently, carbon must diffuse out of the areas that form ferrite (i.e. the carbon is rejected from the ferrite). This will increase the carbon content of the surrounding areas leading to the nucleation of cementite. This then sets up the process of cooperative growth where carbon is partitioned away from the growing ferrite and into the growing cementite by diffusion in the austenite ahead of the reaction interphase. The resulting microstructure is shown in figure (2). The alternating layers of ferrite and cementite are collectively referred to as pearlite. The pearlite appears to grow into the adjoining grains as nodules as shown in figure 2b. Each nodule contains one or more colonies: a colony refers to adjacent ferrite and cementite sheets having the same orientation and growing in the same direction. In this lab we will study the kinetics of the ( → pearlite transformation.
Figure 2a
Figure 2b
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2.2 The Avrami (JMAK) Equation Consider a Fe-0.78%C alloy that was cooled rapidly from 900ºC (austenite range) to 700ºC ("+ cementite range). If we measured the pearlite volume fraction as a function of time, a sigmoidal curve is obtained (figure 3). The initial period where no pearlite is formed is referred to as the incubation period. It reflects the fact that diffusion is needed in order to form the nuclei. Following the incubation period, nucleation and growth occur rapidly, leading to an exponential increase in the pearlite fraction with time. Finally, impingement occurs and rate of transformation is reduced.
Figure 3
In the late 1930's, Johnson, Mehl, Avrami and Kohnogorov modelled the transformation kinetics. The model is now commonly referred to as the INIAK or Johnson-Mehl model. We will derive the model for the simple case of a constant nucleation rate N (nucleus/sec/m3) and a constant growth rate G (m/sec): The volume of a single new grain as a function of time is: V = 4B/3 (G t)3. The total volume fraction of new phase is obtained by summing up (integrating) the
volumes of all new grains within one m3:
dntG3
4X 33t
0total ∫π
=
where n is the number of grains. We can express n as a function of time; n = N t.
Therefore:
NdttG3
4X 33t
0total ∫π
=
The integration is straightforward because G and N are both assumed to be independent
of time. One therefore obtains,
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NtG3
X 43total
π=
The volume fraction is always between 0 and 1. If you examine the above expression, however, Xtotal → ∞ as t → ∞. To understand the mistake that we have made consider a sample where 75% of the austenite has already transformed to pearlite. We assumed a constant nucleation rate of N nuclei per sec per m3. These nuclei can only form in the untransformed material (i.e. in austenite). If the total volume is 1 m3, the volume of austenite is (100% - 75%) or 25% of 1 m3. More generally, the volume of austenite is (1-X) * 1 m3. Our previous analysis ignored the correction factor of (1-X). To obtain the real volume fraction (X) form the total (extended) volume fraction we simply apply a correction of (1−X):
dX = (I−X) dXtotal ,
Rearranging this expression we get:
totaldXX1
dX=
−
integrate
totalXX1
1ln =⎟⎠⎞
⎜⎝⎛−
exponentiate
)Xexp(X1
1total=
−
rearrange
X = 1−exp(−Xtotal)
substitute Xtotal
⎟⎠⎞
⎜⎝⎛ π
−= 43 NtG3
4exp1X
Which is the JMAK equation. More generally, the equation is written as:
X= 1−exp (−b tm)
where b and m are constants. The exponent, m, is known as the Avrami or the JMAK exponent. Typical values of m are between 1 and 4. Try plotting X with b = 1 and m = 4 to convince yourself that this expression does indeed give a sigmoidal curve.
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2.3 Isothermal Transformation (IT or TTT) Diagram The eutectoid transformation of steel is an extremely important practical reaction used by the steel industry to impart specific properties on steels. As such, it is of great importance to understand and be able to predict the speed with which this reaction will occur (if this reaction will occur at all). For example, if we want to produce a steel with the optimum hardness we would want to convert all of the austenite to martensite. Martensite may be formed by rapidly cooling a sample by quenching such that the above described pearlite reaction does not have time to occur. For us to know whether this is possible we must have a means of knowing the amount of pearlite transformation as a function of steel composition, temperature and time. This information is typically plotted as a time, temperature, transformation (TTT) diagram (figure 4).
Figure 4
The first step in constructing a TTT diagram is plotting the transformed fraction as a function of time for different transformation temperatures. The information from the X-t plots is then condensed into a single TTT diagram as shown in figure 5. Note how each X-t diagram contributes one point to the TTT diagram. Typically, this point corresponds to the time needed to achieve a transformed fraction of 1% (start) or 99% (finish). Knowing the time spent at a given temperature one is able to predict the microstructure of the alloy. The TTT diagram has a "C" shape reflecting the fact that the kinetics are determined by two factors, namely, the driving force and the diffusion kinetics. Just below the A1 temperature, diffusion is fast, but the driving force is small. As a result, the transformation time is very long. At very low temperatures, the driving force is large, but now the diffusion coefficient is very small. Again, the transformation is very slow. There is an intermediate temperature range (nose of the C-curve) where the transformation is very fast because the driving force and the diffusion coefficient are both large.
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Figure 5
2.4 Continuous Cooling (CCT) Diagrams The use of isothermal transformation (TTT) diagrams to predict the kinetics of the eutectoid transformation is relatively simple. In industrial practice, however, most heat treatments are far from being isothermal. Instead, most treatments involve the continuous cooling of a sample. In this case, the TTT diagram is not expected to predict the phase transformations accurately. To deal with this, a lot of time and effort has been put into the development of continuous cooling transformation (CCT) diagrams like the one shown in figure 6. To a first approximation the CCT diagram is just a TTT diagram shifted to lower temperatures and longer times. To a first approximation, this may be attributed to the fact that in a continuously cooled sample, fewer nuclei form initially compared with an isothermally transformed sample since few nuclei can form just below the A1 temperature. Finally, note that whereas the TTT diagram is interpreted by reading from left to right at constant temperature, the CCT diagram is read along the cooling curve going from top left to bottom right. A number of ways have been used to construct the CCT diagram. Experimentally, CCT diagrams may be determined from the Jominy end-quench test (figure7). In this test a bar of steel is austenitized for a period of time before being quickly removed from the furnace and placed in a rig which sprays water on the end of the sample. The sample then experiences different rates of cooling along its length. A CCT diagram may be roughly produced by plotting cooling curves corresponding to different positions along the sample length and observing the microstructure at these positions.
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Figure 6
Figure 7
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The Jominy test has a second, perhaps more common, use in determining the hardenability of a sample. The hardenability may be defined as the ability of a steel to harden by the formation of martensite on quenching. This must not be confused with the hardness of a sample. A sample which has a large hardenability is one which may be cooled relatively slowly and still produce a fully matensitic structure. In other words, such a material would be one that has its CCT start curve (X = 1%) shifted far to the right so that pearlite is not formed if the sample is quenched from above the Al temperature. In a Jominy test the cooling rate decreases along the length of the sample in a way that is well know. Therefore, it may be used to determine the hardenabiltiy of a given steel. This is normally determined by plotting the hardness of the sample as a function of the distance along the its length (top of figure 6). If the hardness of a fully martensitic sample is known, then the critical cooling rate may be determined from knowledge of the distance along the sample at which the hardness drops below this level. Finally, the Jominy test is frequently used by steel producers and heat treaters as a means of predicting the hardness profile across heat treated and quenched rods. Standard chats exist which allow the cooling rate along a Jominy sample to be related to the cooling rates within quenched rods. In this way the hardness data determined from the Jominy test may be used to predict the hardness profile across industrially heat treated bars of various diameters. 3. PROCEDURE This lab is made up of two separate experiments; isothermal transformation and continuous cooling transformation. 3.1 Isothermal Transformation For this portion of the lab we will use a 1070 plain carbon steel with the composition given below.
C Mn S P Cu Si Cr 0.72 0.76 0.02 max 0.02 max 0.01 max 0.24 0.06
The material is in the form of drawn rod approximately 20 mm in diameter with a hole drilled through it so that a wire can be threaded through for handling. The heat treatment will be performed using two salt baths. The one on the left is used to austenititze the sample while the second bath is used for the isothermal annealing treatment. The procedure is as follows:
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Important: Extreme caution should be exercised when working with molten salts. An obvious concern is the high temperature. A second concern is the reactivity of the salt. Contact with the skin will result in severe burns. It is mandatory to wear a full face mask and heat-resistant gloves when working around the salt baths. Keep water, dust and paper way from the bath as these will react with the salt.
1. When you arrive at the lab, the samples will be in the austentitizing bath at 900°C. Each
sample will have a wire attached.
2. When the one hour austenitizing period has passed, remove one sample from the first bath and place it quickly in the second salt bath. The transfer must be done very quickly so as to avoid cooling the sample below 700°C. Also make sure the sample does not touch anything on its way into the salt pot.
3. Leave the sample in the salt bath at the transformation temperature for the required length of time, then quickly remove it and drop into the water bucket.
4. Repeat this procedure for each sample. The transformation times are: 100, 200, 300, 400, 600 seconds.
5. Once all the samples have been heat treated they must be mounted, polished and etched with 2% nital. It is very important that the polishing is done very well to ensure good results from the image analyzer.
6. Use the image analysis system to determine the total area fraction of pearlite as well as the number and size distribution of the pearlite colonies for each specimen.
3.2 Continuous Cooling - The Jominy Test In this test we will use an alloy produced by Atlas Steels. The alloy is designated as SPS, it is essentially a 1040 steel with Cr added to improve the hardenability. The chemistry of the alloy is as follows:
C Mn S P Cu Si Cr 0.72 0.76 0.02 max 0.02 max 0.01 max 0.24 0.06
The procedure is as follows: 1. The Jominy bars will be placed in an Inconel heat treat box filled with pitch coke. The
coke is needed to minimize the decarburization of the steel. The box is then placed in a furnace at 900°C for 1/2 hour.
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2. Quench the sample in the Jominy tank. It should take no more than 5 seconds to transfer
the sample from the furnace to the quench tank and begin quenching. When setting up the Jominy tank make sure that the height of the water jet is 6 cm. The water should be turned on immediately following the placement of the sample in the tank.
3. Leave the sample in the quench tank for at least 30 minutes. 4. Grind two parallel flats 180° apart and 0.015 inch deep. Avoid overheating the sample
during grinding. 5. Hardness testing is done by mounting the sample bar in a special jig which mount on the
elevating screw of the Rockwell harness tester. The jig is calibrated in inches. 1. Measure at 1/16" intervals for the first inch.
2. Measure at 1/8" intervals for the second inch. 3. Measure at 1/4" for the remainder of the bar.
4. LAB REPORT Please follow the following format: 1. One lab is to be handed per person two weeks after the second lab period. 2. The lab write up should include:
a) Purpose: (Short, 3-4 sentences explaining what we are trying to do).
b) Observations: Include: - a plot of the pearlite fraction as a function of time, - briefly described the microstructure observed on the microscope, - plot of the hardness as a function of distance, - include error bars for all plots.
c) Discussion: Answer the questions given below:
d) Sources of Error: (describe any sources of errors that may have influenced your results).
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QUESTIONS FOR DISCUSSION
A. Isothermal Transformation Experiment
Al: From your X-t plot determine the values of B and m in the JMAK model.
A2: If possible try to obtain a rough value of the growth rate (G) of the pearlite colonies.
A3: Why is it necessary to use thin samples for this experiment? Can you see any advantages for using a salt bath rather than a normal furnace?
A4: The addition of solutes (e.g. Mn, Cr, Si.... etc) will affect the shape and position of the C- curve. Describe any two mechanisms by which the solutes are able to bring about such changes.
B. Continuous Cooling Experiment
B 1: From your plot of hardness vs. distance, estimate the minimum cooling rate necessary to obtain a fully martensitic structure. How can you increase the minimum cooling rate a) by changing the composition and b) without changing the composition of the steel.
B2: Sometimes in the Jominy test, one initially observes an increasing hardness as one moves away from the quenched end. Under these circumstances, the maximum hardness does not coincide with the maximum cooling rate. Why?
C. Other Questions
Cl: Explain the differences between the TTT and CCT curves. Suppose you are concerned with the formation of martensite following welding. Which curve will you use to analyse the process?
C2: How do you explain the sigmoidal shape of the X vs. t curve? Do you expect a plot of the recrystallized fraction vs. t to look the same?
C3: Consider the precipitation of the θ-phase in Al. The overall reaction is:
Al (supersaturated) → (Al) + θ.
A plot of the precipitate fraction as a function of time will show sigmoidal kinetics. In this case, however, the maximum precipitate fraction is a few percent (say 3%) and the physical impingement of the precipitates will not take place. How can we explain the levelling off of the curve?
C4: How do you explain the "C" appearance of the TTT-diagram? Give an example of a phase transformation where C-Curve kinetics are not expected. Explain your example.
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C5: The nucleation of pearlite is known to occur on grain-boundaries. What effect would a coarser grain size result in? Do you expect the austenitizing temperature to affect the C- curve kinetics? How?
C6: A heat treating company wishes to predict the hardness profile across a bar of SPS steel. The bar was austenitized for 1 hour at 900°C. It was then quenched into water at a rate of 60 m/min. If the bar is 62.5 mm in diameter, plot the hardness profile using your experimentally determined data along with figure 8 below. How would the profile change if oil (faster cooling rate) was used instead of water?
Figure 8
REFERENCES D.A. Porter and K. E. Easterling, Phase Transformations in Metals and Alloys, Chapman Hall, New York, 1992. (Pearlite, TTT, CCT). A.K. Sinha, Ferrous Physical Metallurgy, Butterworth, Stoneham, MA, 1989. (Pearlite, TTT, CCT). ASM Handbook, volume 9, p. 37 (Quantitative Metallography).
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Figure 1. Plate martensite in high carbon steel.
Experiment E: Martensitic Transformations
Introduction:
The purpose of this experiment is to introduce students to a family of phase transformationswhich occur by shear rather than diffusion. In metals, these transformations are referred to asmartensitic transformations. Each student will study the characteristics of the martensitetransformations in an eutectoid steel (the same steel which is used in the nucleation and growthexperiment), aluminum bronze alloys and a series of Fe-Ni alloys of varying compositions.
Background:
Certain metal and ceramic systems experience transformations in which shape changes occurwithout diffusion. These shape changes involve large shear forces and dilatation of the crystal
structure. In metals, these phase transformationsare referred to as martensitic. They areimportant in both ferrous and non-ferrous alloys(e.g. aluminum bronzes and shape memoryalloys) and are utilized in order to improveproperties such as strength and hardness. Inmetals, the magnitude of these changes can belarge enough to allow for observation throughthe light microscope. In certain ceramics (e.g.partially stabilized zirconia, barium titanate),these phase transformations are referred to asdisplacive and are exploited in order to improveproperties such as strength and fracturetoughness. With these types of ceramics,displacive transformations may significantlydecrease the energy available for crack
propagation thus reducing or halting the growth of cracks. The changes that occur within theceramic structure are small in nature and are best examined by various diffraction methods, forexample x-ray diffraction. It is highly recommended that the student review the section of thetextbook on martensite transformations prior to the experiment.
Safety:
It is the responsibility of each TA and each student to be aware of the many hazards in thislaboratory and make use the appropriate safety equipment when performing this experiment. The main potential hazards in this experiment are extreme heat, infrared radiation, high voltage,hot gases, cryogenic materials and hazardous chemicals. The following MSDS are availablePitch Coke (Calcined Petroleum Coke), Barium Chloride (Heat Treating Salts), Sodium Chloride(Heat Treating Salts), Potassium Chloride (Heat Treating Salts), Methanol, 95% Ethanol, Nitric
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Figure 2. TTT diagram with continuous cooling curves.
Acid and Sodium Metabisulphite.
Procedure:
Part One:
In the first part of this experiment,the formation of martensite in aeutectoid steel is investigated. Thesteel must first be austenitized at900°C for 1 hour using either afurnace or a salt bath in JHE 244. The furnace temperature or salttemperature will be preset by theTA. The samples with small holesdrilled through them will beprovided to each student along withwire for holding the specimens. Thread and twist the lengths of wirethrough the holes for easy handling. If a furnace is used for theaustenizing heat treatment, pack the
samples with pitch coke in the Inconel box provided. Place the box in the pre-heated furnace. Inthe salt bath, the samples can be hung by their wires. When the austenizing treatment is done,use tongs to remove the samples from the furnace or salt bath and transfer them as quickly aspossible to a bucket of cold water. Rapid transfer is critical to this procedure. Mount, polish andetch each sample and observe it in the microscope.
Part Two
Five samples consisting of Fe-Ni alloys of varying compositions are provided. The nickelcompositions in weight percent are as follows:
Alloy 1: 25.54 % wt. Ni Alloy 2: 28.30 % wt. Ni Alloy 3: 29.23 % wt. Ni Alloy 4: 31.09 % wt. Ni Alloy 5: 33.88% wt. Ni
In addition, these samples are low in carbon and other residual elements. A table of Ms, and Mffor Fe-Ni alloys is shown on the next page (Trans. AIME, 206 pp.1393). Also included are theAs and the Af temperatures at which austenite starts and finishes forming upon heating themartensitic microstructure, respectively.
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Based on the results provided in this table, the lab group has to determine which samplecorresponds to which alloy and to carefully observe and sketch the type of martensite, whichforms.
The following information should help with the analysis. The initial quenching is done in waterat or slightly below room temperature, a thermometer is available to measure the temperature. The following mediums are available for subsequent quenching: water, methanol, ice, dry iceand liquid nitrogen (do not put thermometers in the liquid nitrogen). Use appropriatethermometers for all solutions (except liquid nitrogen). The freezing point of dry ice is -77°C. Quenching mediums below 0OC can be made by dissolving the dry ice in methanol (use aNalgene beaker - NO GLASS beakers). Liquid nitrogen has a temperature of -196°C (use aNalgene beaker – NO GLASS beakers). Do not put the methanol/dry ice mixture or liquidnitrogen in a sealed container, it will explode.
The martensite can be revealed by etching with 2% or 5% nital after polishing. Aqueous 10%sodium metabisulphite can be also used to darken the martensite, if necessary, but the samplewill have to be lightly repolished before further quenching. It is important to follow good
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metallographic procedures in order to see the martensite clearly.
Transformations of austenite to martensite occur almost instantaneously when the austenite isquenched to a temperature below the Ms temperature. When martensite forms at a lowtemperatures, it is clearly visible without etching because of the considerable surface relief whichit produces.
Before coming to the lab, students should consider how they will go about differentiating thesamples. Before doing anything, a plan should be discussed with the TA Note that the TA canonly offer suggestions. However, the TA will help with observations of the martensite structuresin the light microscope to make sure the salient features of the microstructures are noted for thesamples produced.
Part Three
Three specimens of copper with between 11 and 12 wt.% aluminum labeled X21, X22 and X23and two specimens of a hypoeutectoid plain carbon steel labeled X24 and X25 are available inthe microscope room desiccators for observation. The copper aluminum alloy is commonlyknown as aluminum bronze. Both alloy systems contain an eutectoid and undergo martensitictransformations when rapidly cooled from high temperatures. In the copper - aluminum alloysystem, this eutectoid occurs at 11.8 wt. % Al and is created by the breakdown of a compound,beta (β), to a solid solution of aluminum in copper called alpha (α) and another compound calledgamma (γ).
Each specimen has already been mounted, polished and etched and is designated by the numberon the bottom of the mount. If the specimens require repolishing or etching, please contact thetechnical staff. Each specimen should be observed visually and at high and low magnificationswith a bench microscope. Systematically scan the whole section. Select regions that arerepresentative of the majority of the specimen. Sketch the observed microstructures on unruledwhite paper, which will be provided. The sketches indicate whether or not a clear understandingof the basic structures observed have been achieved. Each sketch should show the principalcharacteristics of each specimen. Refer to the questions in the Lab Report section as a guide towhat you should be thinking about when observing the various specimens. The 3T04 atlascontains additional images that may be helpful. Please return the specimens to the desiccatorsafter observations are completed.
Specimen X21, Copper with 11.8% aluminum, eutectoid alloy (chill cast, reheated for 1hour at 900°C and slowly cooled)The surface of this specimen has an iridescent appearance. The material has completelytransformed to the eutectoid, which consists mainly of a fine lamellar arrangement of the alphasolid solution (light) and the gamma compound (dark). However, in some regions, the alpha hasa coarser more irregular pattern.
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Specimen X23, Copper with 11.3% aluminium, hypo-eutectoid alloy (sand cast, annealedat 900°C, slowly cooled)Rather jagged needles of a are set in a background of eutectoid. The eutectoid is mainlyirregular and relatively coarse in pattern; the gamma particles are large enough to be clearly seenas blue in colour. Sometimes, small brown regions are found in the eutectoid areas; theserepresent regions where there is a fine lamellar structure that is unresolvable.
Specimen X22 (as X21, but water quenched after slowly cooling to 530°C)This specimen has been partially transformed. Nodules of eutectoid are found at the grainboundaries of the original beta phase. The structure of the eutectoid is lamellar except near theoriginal beta boundaries, where it is coarser and more irregular. The remainder of the structureis light with a faint acicular pattern. This is known as beta prime and is a martensitetransformation product of beta. Note that the beta phase is body-centered cubic, while the betaprime and (Cu) solid solution phases are distorted face-centered cubic and face-centered cubic,respectively.
Specimen X24 – 0.35% carbon steel (normalized from 870°C)This structure of X24 consists of fine pearlite and ferrite almost approaching a Widmanstattenpattern. Alloys undergoing eutectoid reactions also lend themselves to heat treatment. The mostimportant example is the quench-hardening and tempering of steel. In general, air cooling(normalizing) from the austenite region gives finer structure than slow cooling (annealing) resulting in higher hardness and strength in the normalized condition. However, considerablymore marked effects are obtained by quenching the steel and then modifying or tempering thestructure to give the final desired result.
The above specimen illustrates some difference in structure as a result of air-cooling instead ofslow cooling. But, more pronounced differences occur when the steel is water quenched fromthe austenite condition. This treatment prevents the formation of ferrite and pearlite. However,on reaching a relatively low temperature, the austenite structure is so unstable that a rapidmartensitic change takes place, involving adjustment of atomic positions. Diffusion does notcontrol this type of solid state reaction. The microstructure then consists of many interlacingmartensitic needles, with a very small amount of residual austenite. There are no sharpboundaries between the needles and the matrix and the structure has a diffuse acicularappearance. The carbon does not escape from solution, the overall atomic pattern is renderedmore complex, it is distorted and in a state of stress. For these reasons, the material becomeshard and brittle (unless the carbon content is low). Reheating, that is tempering, is then used torelease the carbon as fine carbide particles, to a degree determined by the final propertiesrequired.
Specimen X25 – 0.35% carbon steel, water-quenched from 870°C (1 hour treated)A typical martensitic structure. Away from the edge of the specimen, small precipitates of ferritemay be found. These have either an angular shape, or sometimes, they have ‘saw-teeth’
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contours. They have formed where the rate of cooling has not been so high as at the surface.
Lab ReportThe lab report should address the following points:
Why does the steel not transform to pearlite on quenching to room temperature? Refer to theTTT diagram(s) in the response.
What is the difference between lath and plate martensite?
Why are martensite particles plate- or lath-shaped?
Relate the observed microstructures in the X21-X23 series to the Cu-Al phase diagram (nextpage). Describe the transformations which occur on quenching and on slow cooling.
Specimens X22 and X251) What are the metallographic characteristics of martensite?2) In X22, there are nodules of eutectoid. Explain why.3) How could equilibrium structures be obtained in these alloys?
Make sure the sketches clearly show examples of the different accommodation mechanismsobserved.
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Experiment F: Failure Analysis
Introduction:
In this experiment, a component that has failed in service will be analyzed to determine theprobable cause or causes of the failure. Each group will give a short oral presentation on theirfindings.
Procedure
For this laboratory, a broken component is required, any failed component can be used eitherbring something from home or work or a selection of components are available in thedepartment. The component should have failed in service and be in a condition suitable foranalysis. If in doubt about the suitability of the component ask the TA. Remember, don’t fittogether the fracture parts unless necessary, delicate features on the fracture surfaces can bedamaged. Cleaning can remove important evidence and may contaminate or corrode surfaces.
Each group must perform metallurgical analysis, and if needed a chemical analysis and SEM. Based on these results and any observations, each group will present its findings to the class in aBRIEF (5-7 minutes) oral presentation. In the presentation (and the report), the following pointsmust be addressed:
a. Function of the part in serviceb. Results of the chemical analysis (if appropriate)c. Description of the microstructured. Description of the failuree. Why did the part fail (design, material, operating conditions?)f. What can be done to rectify the problem?
An “industrial” presentation should be prepared, i.e., suppose that the finding are beingpresented to corporate engineers who are responsible for the design of the component. The duedate will be discussed in the lab.
Background:
What is a failure? The fracture of a shaft in a machine is an oblivious failure. But is thisexample what the definition for a failure should be based on? What about a plane crash? Duringthe course of a normal day how many situations occur where things or procedures don’t workproperly? A possible definition for a failure might be “ the inability of a component, machine orprocess to function properly”.
In this experiment, the focus is on failures of components in service. Failure analysis will beused to determine the factors responsible for the failure of the component examined. The properapplication of failure analysis techniques can provide valuable feedback to design problems andmaterial limitations. In the real world, there may also be legal reasons for conducting a failureanalysis: if a failure has caused injury or damage to expensive equipment, it will be necessary todetermine where to lay the blame!
Based on the results of a careful failure analysis, it is generally possible to attribute the problemto one of the following:
Poor design (stress raisers, etc…) or improper selection of materialImproper testingService and operation (use, misuse, etc...)Improper fabrication or manufacturing procedures (casting, machining, improper heattreatment etc…)Improper assemblyImproper maintenance
Due to the very wide variety of factors which can contribute to an engineering failure, this typeof analysis requires considerable knowledge in metallurgy, corrosion and machine design. Inaddition, there is no substitute for experience.
Steps in Failure Analysis
A proper failure analysis involves a systematic study of all the factors which may havecontributed to the failure. The role of an engineer in a failure analysis situation is to answer oneor more of the following questions:
- What is the problem? - What is the main cause of the problem? - What are the potential solutions? - What is the best solution?
It is important to remain objective and have clear information about the failure. In the realworld, a lot of time and effort can be spent on understanding and negotiating the goals,procedures of an analysis with a client (internal or external) and the allowable costs associatedwith the analysis. Establishing clearly defined expectations is very important. Finding out whathappened at the actual location of the failure. Investigate! Visiting the “scene of the crime” canoften be very valuable. Workers on the shop floor may know more about what happened thenthe manager. Note-taking, documentation, drawings, pictures and planning are key parts offailure analysis. Do as much non-destructive analysis as possible before the destructive parts ofthe investigation. During the course of an investigation, it is important to identify all thepossible causes of a failure and eliminate as many as possible either due to testing or the balanceof probabilities. For solutions, look at all possibilities and list criteria for rejecting and acceptingeach solution. Select the best solution and if possible, evaluate the solution.
For the purposes of this lab, the following will service as a guide to the stages, which comprisethe investigation and analysis of a failure:
INITIAL OBSERVATIONSPreliminary visual study of the failed componentLOTS of photographs
COLLECTION OF BACKGROUND DATA (INVESTIGATION/RESEARCH)Component design (specifications, fabrications, repairs, material)How long was the part in service?Nature of stresses at time of failure?Working environment of the part (temperature, medium, corrosion)
MACROSCOPIC EXAMINATION AND ANALYSISDetailed visual study of the failed componentIdentification and interpretation of surface irregularities
SELECTION, IDENTIFICATION, AND CLEANING OF SPECIMENSSelect specimens for failure analysisPrepare metallographic sections
MECHANICAL ANALYSISStress-strain curves, if availableHardness measurements (macro, micro)
CHEMICAL ANALYSISVerify composition of component to see whether it falls within specified limits
MICROSCOPIC EXAMINATION AND ANALYSISMetallographic analysis (correct microstructure)Scanning electron microscopy (wear, fractures, corrosion, microstructures)EDS - Elemental analysis with the SEM (inclusions, corrosion, phases)
TESTING UNDER SIMULATED SERVICE CONDITIONS (if possible)May help to identity cause of failure
SYNTHESIS OF FAILUREAnalysis of all facts and evidenceFormulation of conclusionsWrite report and suggest solution to problemDesign? Normal wear? Material? Etc....
Some Notions about Failures
Figure 1. A - Tensile B - Shear C - Tearing
Figure 2. Ductile “ Cup - Cone” Fracture.
In this section, some notions about failures (and fracture in particular) will be discussed. It is abrief overview and students wishing to look at the subject more in depth are encouraged to dosome research on their own in the library. Most books on introductory metallurgy will have atleast one section on failure analysis. Also, there are quite a few books devoted entirely to thesubject. Figures and micrographs are provided at the end help to illustrate the points raised inthe following discussion. Primarily, the discussion applies to metals, but many non-metallic orcomposite materials will exhibit similar behaviour.
When a part or component has fractured, it is often necessary to determine the nature of thefracture in order to draw any conclusion about the failure.
Was the fracture ductile, brittle, or a combination of the two?Was fatigue, wear, or corrosion involved in the failure?Where did the fracture start at or below the surface?Did the fracture start at one point, or did it originate at several points?What was the loading of the component at the point of fracture? Tensile? Shear? Tearing?
Modes of Fracture - Ductile orBrittle?
The mode of facture of a part may bedifficult to discern if the backgroundinformation of the failure is limited. Mix modes of fracture may occurwithin the same part. Anunderstanding of fracture mechanicsand fractography can be very helpful.
Ductile fractures are high-energy fractures. They are the result of shear forces, which produce
Figure 3. Brittle Fracture
Figure 4. A distinct Chevron pattern is apparent on the right-hand side of the fracture surface.
plastic deformation (slip or twinning). Thus, they occur at stresses above the material’s yieldstrength. They are characterized by stable crack propagation of the load causing cracks topropagate. When the load is removed, the crack stops advancing. Shear fractures caused by asingle load are dull grey and fibrous, with edges that are usually deformed plastically. Examination of the fracture surface under the stereoscope will usually show dimples unless thefracture has been severely smeared at some point during fracture. During ductile fracture, smallcavities or microvoids are initially formed by slip. These microvoids are normally associatedwith matrix/secondary particle (inclusions, precipitates etc) decohesion or cracking. When themicrovoids join together or coalesce, they eventually grow to form a crack under continuedloading. The microvoids appear as dimples after fracture. The crack propagates with the aid ofstress concentration at its tip, generally moving perpendicular to the tensile force and eventuallyforming a ‘shear lip’ at the surface (the plane of the shear lip will be at 45° from the tensile loadaxis). The classic example of a ductile fracture is the cup-and-cone geometry often observedduring a standard tensile test on ductile materials. The reduction in area that occurs prior to thefinal ductile fracture of a component is commonly referred to as necking. If the material is veryductile, the sample may “neck” down to a point prior to final fracture.
Brittle fractures are low-energy fractures. They occur at stresses well below the yieldstrength, are usually associated with flaws, areoften catastrophic and usually occur rapidlywithout warning. They are distinguished fromductile fractures by the absence of gross plasticdeformation and little or no evidence ofreduction in area at the point of fracture. Brittle fractures often appear bright andcrystalline. The two modes of brittle fracture
are transgranular and intergranular. With transgranular fractures, each grain tends to fracture ona single cleavage plane. Thus a brittle fracture surface will often sparkle in the light whenrotated in the hand. Brittle fracture surfaces sometimes have distinctive appearances from theorigin of fracture, a characteristic ‘chevron’ or ‘herringbone’ pattern is formed which points tothe fracture origin. The fracture surfaces are usually smooth compared with ductile failures andthe fractured parts usually fit together well (please resist the temptation to do this unless
absolutely necessary). Brittle failures occur in brittle materials, but can also occur in normallyductile materials under certain loading conditions (e.g. high strain rates, tri-axial stresses, or lowtemperatures).
Figure 6. Quasi cleavage
Figure 8. Locating crack origins.
Figure 6. Quasi-cleavage
Figure 7. Radial and chevrons marks point to theregion from which the crack originated.
In some materials such as certain tempered martensitic steels, the fracture is predominatelybrittle but has elements of ductile fracture. The fracture surface has small facets, dimples andridges which sometimes result in a rosette-like appearance. This type of fracture mode is usuallyreferred to as quasi-cleavage, a mixed mechanism of microvoid coalescence and cleavage.
Point of Origin
Several features exist around fractures or on fractured surfaces, visible to the eye, that indicatethe area where a fracture has initiated or its point(s) of origin. Branching cracks will usuallyoriginate from a single crack or point. A network of cracks will commonly have a larger ordominate crack. On fracture surfaces, chevron marks, radial marks or ridges can point to theregion where the crack initiated. In the case of many fatigue failures, “beach marks” oralternating bands radiate out from the origin like waves on a beach. These marks can be said toidentify the fracture origin as crack length increases, beach marks increase in size and spacing sothat the location of the smallest beach mark indicates the origin of the failure.
Figure 9. Fatigue cracking is comprised of two stages. Figure 10. The first stage of fatigue cracking isreferred to as Stage I.
Fatigue Failure
Fatigue failures are the most common types of fracture in machines. They result from theapplication of a repetitive stress over time, sometimes many years. There are two stages to afatigue fracture. In Stage I, the surface or interface as the result of slip is subjected to intrusionsand extrusions resulting in crack propagation usually at 45o to the tensile stress. In Stage II, thefatigue fracture is progressive, i.e. there is a certain amount of crack growth per load application. The appearance of fatigue failures has often been described as brittle because there is very littlegross plastic deformation associated with them and the fracture surfaces are fairly smooth. However, because fatigue failures are progressive, they usually leave characteristic marks, calledbeach marks as previously described. Note, however, that not all fatigue failures show beachmarks. Therefore, their absence does not necessarily preclude fatigue. At the microscopic level,fatigue striations are created on the fracture surface by each progressive step of the crack and canbe used to estimate the rate of crack growth - low or high cycle fatigue. Although beach marksappear similar regardless of the type of load application, specific features vary, as shown in thetable at the end of the write-up. On fact common to all fatigue failures is that the fractures arecaused by a tensile load generated by tension, rotation, or bending and the fracture path is normalto the tensile load.
Bending Failure
Bending is one of the most common causes of fracture in machine and structural components. Failure can result from a single application of a load greater than the overall strength of a part ofcan be due to a reversing load that results in a bending fatigue fracture.
Torsional Failure
Torsional failure are most common in shafts, including crankshafts, torsion bars and axles. Theappearance of a torsion-fatigue fracture is quite different from that caused by bending. Torsion-fatigue failures occur along the planes of maximum shear or along the plane of maximumtension. Maximum shear stress occurs along the axis of the shaft and at right angles to it, whilethe maximum tensile stress acts at an angle of 45° to the two shear stresses.
In torsion, the maximum shear stress is equal to the maximum tensile stress so that which type offracture occurs will depend on the relative values of the shear strength and the tensile strength ofthe part.
Wear
The quality of most metal products depends on the condition of their surfaces and on surfacedeterioration due to use. Indeed, surface deterioration is often the major factor limiting the lifeand the performance of machine components. Wear may be defined as unintentionaldeterioration resulting from used or environment due to the displacement or detachment ofmetallic particles from a metallic surface.
Several types of wear can be identified:(A) Adhesive or metallic wear (contact with another metal)(B) Abrasion (contact with a metallic or non-metallic abrasive)(C) Erosion (contact with moving liquids or gases)
In adhesive wear (scoring, galling, seizing, scuffing), tiny projections produce friction bymechanical interference, with the relative motion of contacting surfaces increasing resistance tofurther motion. If the driving force is sufficient to maintain movement, the interlocked particlesare deformed. If they are of brittle material, they may be torn off. Thus wear resistance can beimproved by preventing metal to metal contact, increasing the hardness to resist initialindentation, increasing the toughness to resist the tearing out of metallic particles, and increasingthe surface smoothness to eliminate projections. In adhesive wear, the effect of heat produced byfriction between contacting surfaces can reduce war resistance in several ways. It may temperhardened structures, cause phase changes that increase hardness and brittleness, and acceleratecorrosion reactions. If the pressure between two mating parts and/or their temperature issufficiently large, the parts can be welded together they may seize and thus cause completestoppage, or if relative motion is not prevented, pieces of the opposite face may pull out.
Abrasive wear occurs when hard particles slide or roll under pressure across a surface. The hardparticles tend to gouge or scratch the softer material. They may also penetrate the softer material
Figure 11. Corrosion pits on an intergranular fracture surface.
and cause the tearing off of metallic particles.
Corrosion
Corrosion may be defined as the destruction of a material by chemical, electrochemical ormetallurgical interaction between the environment and the material. The basic cause ofcorrosion is the instability of metals in their refined forms. The metals tend to revert to theirnatural states through the process of corrosion.
Corrosion is an electrochemical process resulting in part or all of a metal being transformedfrom the metallic to the ionic state. It requires the flow of electricity between certain areas of ametal surface through an electrolyte, which can be any solution containing ions.
An electrochemical cell is made up of a connected anode and a cathode which are both in contactwith the electrolyte. Material is removed from the anode, the metal which corrodes.
Figure 12 The formation of dimples in a material containing sphericalprecipitates. during the tensile stress and different loading conditions
Figure 13. Ductile Dimples.
Figure 14. Dimples from a less ductile material.
Figure 16. A cleavage facet on the surface of a brittlefracture.
Figure 15. Directional dimples caused by shear forces.
Figure 18. An intergranular fracture.
Figure 16. SEM Micrograph showing cleavage facets on the surface of a brittle fracture. A - a grain, B - another grain, C - a carbide particle and D and E - secondary cracks.
Figure 19. Fatigue fracture showingbeach marks.
Figure 20. Fatigue fracture with internal origin showing beach marks.
Figure 21. Stage I of a fatigue fracture.
Figure 23. SEM Micrograph showing fatigue striations on the surface of a fatiguefracture.
Figure 22. Diagram showing a typical stage II fatigue fracture withcrack paths indicated.
