Steel Metal

Commercially pure titanium and most of titanium alloys can be welded by procedures and equipment used in welding austenitic stainless steel and aluminum. Because of the high reactivity of titanium and titanium alloys at temperatures above 550°C, additional precautions must be applied to shield the weldment from contact with air. Also, titanium base metal and filler metal must be clean to avoid contamination during welding.

Unalloyed titanium and all alpha titanium alloys are weldable. Although the alpha-beta alloy Ti-6Al-4V and other weakly beta-stabilised alloys are also weldable, strongly beta-stabilised alpha-beta alloys are embrittled by welding. Most beta alloys can be welded, but because aged welds in beta alloys can be quite brittle, heat treatment to strengthen the weld by age hardening should be used with caution.

Unalloyed titanium is generally available in several grades, ranging in purity from 98.5 to 99.5% Ti. These grades are strengthened by variations in oxygen, nitrogen, carbon, and iron. Strengthening by cold working is possible but is seldom used. All grades are usually welded in the annealed condition. Welding of cold-worked alloys anneals the heat-affected zone (HAZ) and eliminates the strength produced by cold working.

Alpha alloys Ti-5Al-2.5Sn, Ti-6Al-2Sn-4Zr-2Mo, Ti-5Al-5Sn-2Zr-2Mo, Ti-6Al-2Nb-1Ta-1Mo, and Ti-8Al-1Mo-1V are always welded in the annealed condition.

A1pha-beta alloys of Ti-6A1-4V can be welded in the annealed condition or in the solution-treated and partially aged condition, with ageing completed during postweld stress relieving. In contrast to unalloyed titanium and the alpha alloys, which can be strengthened only by cold work, the alpha-beta and beta alloys can be strengthened by heat treatment.

The low weld ductility of most alpha-beta alloys is caused by phase transformation in the weld zone or in the HAZ. Alpha-beta alloys can be welded autogeneously or with various filler metals. It is common to weld some of the lower alloyed materials with matching filler metals. Filler metal of an equivalent grade or one grade lower is used to ensure good weld strength and ductility. Filler metal of matching composition is used to weld the Ti-6Al-4V alloy. This extra low-interstitial (ELI) grade improves ductility and toughness.

The use of filler metals that improve ductility may not prevent embrittlement of the HAZ in susceptible alloys. In addition, low-alloy welds can be embrittled by hydride precipitation. However, with proper joint preparation, filler-metal storage, and shielding, hydride precipitation can be avoided.

Metastable beta alloys Ti-3Al-13V-11Cr, Ti-11.5Mo-6Zr-4.5Sn, Ti-8Mo-8V-2Fe-3Al, Ti-15V-3Cr-3Al-3Sn, and Ti-3Al-8V-6Cr-4Zr-4Mo are weldable in the annealed or solution heat treated condition. In the as-welded condition, welds are 1ow in strength but ductile. Beta alloy weldments are sometimes used in the as-welded condition. Welds in the Ti-3Al-13V-11Cr alloy embrittle more severely when age hardened. To obtain full strength, the metastable beta alloys are welded in the annealed condition; the weld is cold worked by planishing, and the weldment is then solution treated and aged. This procedure also obtains adequate ductility in the weld.

Welding processes

The following fusion-welding processes are used for joining titanium and titanium alloys:

• Gas-tungsten arc welding (GTAW)
• Gas-metal arc welding (GMAW)
• Plasma arc welding (PAW)
• Electron-beam welding (EBW)
• Laser-beam welding (LBW)
• Friction welding (FRW)
• Resistance welding (RW)

Fluxes cannot be used with these processes because they combine with titanium to cause brittleness and may reduce corrosion resistance. The welding processes that use fluxes are electroslag welding, submerged arc welding, and flux-cored arc welding. These processes have been used on a limited basis. However, they are not considered to be economical because they require high-cost, fluoride-base fluxes.

Gas-tungsten arc welding is the most widely used process for joining titanium and titanium alloys except for parts with thick sections. Square-groove butt joints can be welded without filler metal in base metals up to 2.5 mm thick. For thicker base metals, the joint should be grooved, and filler metal is required. The heated weld metal in the weld zone must be shielded from the atmosphere to prevent contamination with oxygen, nitrogen, and carbon, which will degrade the weldment ductility.

Gas-metal arc welding is used to join titanium and titanium alloys more than 3 mm thick. It is applied using pulsed current or the spray mode and is less costly than GTAW, especially when the base metal thickness is greater than 13 mm.

Plasma arc welding is also applicable to joining titanium and titanium alloys. It is faster than GTAW and can be used on thicker sections, such as one-pass welding of plate up to 13 mm thick, using keyhole techniques.

Electron-beam welding is used in the aircraft and aerospace industries for producing high-quality welds in titanium and titanium alloy plates ranging from 6 mm to more than 76 mm thick. Because the welding is performed in a high-vacuum atmosphere, low contamination of the weldment is achieved.

Laser-beam welding is increasingly being used to join titanium and titanium alloys. Square-butt weld joint configurations can be used, and the welding process does not require the use of vacuum chambers; gas shielding is still required. This process is more limited than electron-beam welding regarding base metal thickness, which cannot usually exceed 13 mm.

Friction welding is useful in joining tube, pipe, or rods, where joint cleanliness can be achieved without shielding.

Resistance welding is used to join titanium and titanium alloy sheet by either spot welds or continuous seam welds. The process is also used for welding titanium sheet to dissimilar metals, that is, cladding titanium to carbon or stainless steel plate.

Filler material and electrodes

Filler-metal composition is usually matched to the grade of titanium being welded. For improved joint ductility in welding the higher strength grades of unalloyed titanium, filler metal of yield strength lower than that of the base metal is occasionally used. Because of the dilution that occurs during welding, the weld deposit acquires the required strength. Unalloyed filler metal is sometimes used to weld Ti-5A1-2.5Sn and Ti-6A1-4V for improved joint ductility.

The use of unalloyed filler metals 1owers the beta content of the weldment, thereby reducing the extent of the transformation that occurs and improving ductility. Engineering approval, however, is recommended when using pure filler metal to ensure that the weld meets strength requirements.

Another option is filler metal containing lower interstitial content (oxygen, hydrogen, nitrogen, and carbon) or alloying contents that are lower than the base metal being used. The use of filler metals that improve ductility does not preclude embrittlement of the HAZ in susceptible alloys. In addition, low-alloy welds may enhance the possibility of hydrogen embrittlement.

Shielding gases in welding titanium and titanium alloys are only argon and helium, and occasionally ? mixture of these two gases is used for shielding. Because it is more readily available and less costly, argon is more widely used.

Electrodes. The conventional thoriated tungsten types of electrodes (EWTh-1 or EWTh-2) are used for GTAW of titanium. Electrode size is governed by the smallest diameter able to carry the welding current. To improve arc initiation and control the spread of the arc, the electrode should be ground to a point. The electrode may extend one and a half times the size of the diameter beyond the end of the nozzle.

Aluminum and its alloys can be joined by more methods than any other metal, but aluminum has several chemical and physical properties that need to be understood when using the various joining processes. The specific properties that affect welding are its oxide characteristics, its thermal, electrical, and nonmagnetic characteristics, lack of color change when heated, and wide range of mechanical properties and melting temperatures that result from alloying with other metals. Oxide. Aluminum oxide melts at about 2050 oC which is much higher than the melting point of the base alloy. If the oxide is not removed or displaced, the result is incomplete fusion. In some joining processes, chlorides and fluorides are used in order to remove the oxide contain. Chlorides and fluorides must be removed after the joining operation to avoid a possible corrosion problem in service. Hydrogen Solubility. Hydrogen dissolves very rapidly in molten aluminum. However, hydrogen has almost no solubility in solid aluminum and it has been determined to be the primary cause of porosity in aluminum welds. High temperatures of the weld pool allow a large amount of hydrogen to be absorbed, and as the pool solidifies, the solubility of hydrogen is greatly reduced. Hydrogen that exceeds the effective solubility limit forms gas porosity, if it does not escape from the solidifying weld. Electrical Conductivity. For arc welding, it is important that aluminum alloys possess high electrical conductivity — pure aluminum has 62% that of pure copper. High electrical conductivity permits the use of long contact tubes guns, because resistance heating of the electrode does not occur, as is experienced with ferrous electrodes. Thermal Characteristics. The thermal conductivity of aluminum is about 6 times that of steel. Although the melting temperature of aluminum alloys is substantially bellow that of ferrous alloys, higher heat inputs are required to weld aluminum because of its high specific heat. High thermal conductivity makes aluminum very sensitive to fluctuations in heat input by the welding process. Forms of Aluminum. Most forms of aluminum can be welded. All the wrought forms (sheet, plate, extrusions, forgings, rod, bar and impact extrusions), as well as sand and permanent mold castings, can be welded. Welding on conventional die-castings produces excessive porosity in both the weld and the base metal adjacent to the weld because of internal gas. Vacuum die-castings, however, have been welded with excellent results. Powder metallurgy (P/M) parts also may suffer from porosity during welding because of internal gas. The alloy composition is a much more significant factor than the form in determining the weldability of an aluminum alloy. Filler Alloy Selection Criteria When choosing the optimum filler alloy, the application (end use) of the welded part and its desired performance must be prime considerations. Many alloys and alloy combinations can be joined using any one of several filler alloys, but only one filler may be optimal for a specific application. The primary factors commonly considered when selecting a welding filler alloy are: * Ease of welding * Tensile or shear strength of the weld * Weld ductility * Service temperature * Corrosion resistance * Color match between the weld and the base alloy after anodizing * Sensitivity to Weld Cracking. Ease of welding is the first consideration for most welding applications. In general, the non-heat-treatable aluminum alloys can be welded with a filler alloy of the same basic composition as the base alloy. The heat-treatable aluminum alloys are somewhat more metallurgically complex and more sensitive to "hot short" cracking, which results from heat - affected zone (HAZ) liquidation during the welding operation. Generally, a dissimilar alloy filler having higher levels of solute (for example, copper or silicon) is used in this case. * The high-purity 1xxx series alloys and 3003 are easy to weld with a base alloy filler, 1100 alloy, or an aluminum - silicon alloy filler, such as 4043. * Alloy 2219 exhibits the best weldability of the 2xxx series base alloys and is easily welded with 2319, 4043 and 4145 fillers. * Aluminum-silicon-copper filler alloy 4145 provides the least susceptibility to weld cracking with 2xxx series wrought copper bearing alloys, as well as aluminum-copper and aluminum-silicon-copper aluminum alloy castings * The cracking of aluminum-magnesium alloy welds decreases as the magnesium content of the weld increases above 2%. * The 6xxx series base alloys are most easily welded with the aluminum-silicon type filler alloys, such as 4043 and 4047. However, the aluminum-magnesium type filler alloys can also be employed satisfactorily with the low-copper bearing 6xxx alloys when higher shear strength and weld metal ductility are required. * The 7xxx series (aluminum-zinc-magnesium) alloys exhibit a wide range of crack sensitivity during the welding. Alloys 7005 and 7039, with a low copper content (<0.1%), have a narrow melting range and can be readily joined with the high magnesium filler alloys 5356, 5183 and 5556. The 7xxx series alloys that possess a substantial amount of copper, such as 7975 and 7178, have a very wide melting range with a low solidus temperature and are extremely sensitive to weld cracking when are welded. Welding Processes The GTAW (gas-metal arc welding) process has been used to weld thicknesses from 0,25 to 150 mm and can be used in all welding positions. Because it is relatively slow, it is highly maneuverable for welding tubing, piping and variable shapes. It permits excellent penetration control and can produce welds of excellent soundness. Weld termination craters can be filled easily as the current is tapered down by a foot pedal or electronic control. The ac - GTAW process provides an arc cleaning action to remove the surface oxide during the positive electrode half of the cycle and a penetrating arc when the electrode is operated at negative polarity. The dc - GTAW Process. Negative electrode polarity direct current can be used to weld aluminum by manual and mechanized means. Other arc welding processes include shielded metal arc welding (SMAW), as well as electroslag and electrogas welding (ESW, EGW). SMAW with flux-coated rods has been replaced to a very substantial degree by the GMAW process. The oxyfuel gas welding (OFW) process uses a flux and either an oxyacetylene or oxyhydrogen gas flame. When the oxyacetylene flame is used, a slightly reduced flame is required, which causes a carbonaceous deposit that obscures the weld and slows the travel speed. Electron - beam welding (EBW) in a vacuum chamber produces a very deep, narrow penetration at high welding speeds. The low overall heat input produces the highest as-welded strengths in the heat treatable alloys. The high thermal gradient from the weld into the base metal creates very limited metallurgical modifications and is least likely to cause intergranular cracking in butt joints when no filler is added. Laser-beam welding (LBW) is now considered to be a viable fusion joining process for aluminum with the advent of commercially available, stable, high-power laser systems. Because of aluminum`s high reflectivity, effective coupling of the laser beam and aluminum requires a relatively high power density.

Non-heat-treatable aluminum alloys constitute a group of alloys that rely solely upon cold work and solid solution strengthening for their strength properties. They differ from heat-treatable alloys in that they are incapable of forming second-phase precipitates for improved strength. Therefore, non-heat-treatable alloys cannot achieve the high strengths characteristics of precipitation-hardened alloys.

The absence of precipitate-forming elements in these low- to moderate-strength non-heat-treatable alloys becomes a positive attribute when considering weldability, because many of the alloy additions needed for precipitation hardening (for example, copper plus magnesium, or magnesium plus silicon) can lead to liquation or hot cracking during welding. In addition, joint efficiencies are higher in not-heat treatable alloys because the heat-affected zone (HAZ) is not compromised by the coarsening or dissolution of precipitates. This obviates the need for thick joint lands or postweld heat treatment and favors the use of welded structures in the as-welded condition.

Alloy Classification and Typical Applications

Non-heat-treatable wrought aluminum alloys can be placed into one of four groups using standard Aluminum Association designations:

Alloy number Alloy addition
1xxx Al (99% minimum purity)
3xxx Al - Mn
4xxx Al - Si
5xxx Al - Mg

The 1xxx-series alloys are of commercial purity (>99% Al) and are used where thermal/electrical conduction or corrosion resistance becomes paramount over strength in design considerations (for example, alloy 1100 is used for sheet metal work, fine stock and chemical equipment). Alloys with purity levels greater than 99,5% are used for electrical conductors (for example alloy 1350).

The 3xxx-series alloys are used in applications where added additional strength and formability are needed, in addition to maintaining excellent corrosion resistance (for example alloy 3004 is used for sheet metal work, storage tanks, and beverage containers). Typical applications include cooking utensils, pressure vessels, and building products (siding, gutters and so on). These alloys get their strength from cold work and fine (Mn, Fe)Al6 dispersoids that pin grain and subgrain boundaries. There is also a small degree of solid solution strengthening from both manganese and magnesium.

4xxx Alloys. Apart from their use as welding filler material, the 4xxx-series alloys have limited industrial application in wrought form.

The 5xxx-series alloys are used in cases where still higher strengths are required. This strength is achieved from large quantities of magnesium in solid solution. More importantly, magnesium promotes work hardening by lowering the stacking fault energy, thus reducing the tendency for dynamic recovery. Applications for 5xxx-series alloys include automobile and appliance trim, pressure vessels, armor plate, and components for marine and cryogenic service.

While these alloys normally exhibit good corrosion resistance, care must be taken during processing to avoid formation of continuos b-Mg3Al2 precipitates at grain boundaries, which can lead to intergranular corrosion. This can occur in heavily cold-worked, high-magnesium alloys exposed to temperatures from 120 to 200oC. Alloy 5454 possesses the highest magnesium content suitable for sustained elevated temperatures and has become the standard alloy used for truck bodies for hot oil or asphalt applications, and storage tanks for heated products.
Filler Alloy Selection

Filler alloys used to join non-heat-treatable alloys can be selected from one of three alloy groups:

Alloy number Alloy addition
1xxx Al (99% minimum purity)
4xxx Al - Si
5xxx Al - Mg

Commonly used filler alloys include 1100, 1188, 4043, 4047, 5554, 5654, 5183, 5356 and 5556. Selecting the best filler alloy for a given application depends on the desired performance related to weldability, strength, ductility and corrosion resistance.

In general, the filler alloy selected should be similar in composition to the base metal alloy. Thus, a 1xxx filler alloy is recommended for joining 1xxx- or 3xxx-series base metal alloys. An exception to this rule is encountered when weldability becomes an issue. Weldability of non-heat-treatable aluminum alloys can be measured in terms of resistance to hot cracking and porosity.

Hot cracking. Problems with hot cracking are encountered when welding under highly constrained conditions or when welding certain alloys that are highly susceptible to cracking. Similar problems may be encountered when 1xxx fillers are used to join 5xxx alloys (or vice versa) or when welding dissimilar metal alloys such as alloys 1100 and 5083, where mutual dilution may result in low magnesium levels. Electron-beam welding or laser-beam welding can also result in cracking when magnesium, a high-vapor-pressure alloying element, is boiled off. The problem is aggravated when welding in a vacuum environment.

Another approach to be taken when hot cracking persists is to use 4xxx fillers. These aluminum-silicon alloys have exceptional resistance to cracking, due in part to their abundance of liquid eutectic available for back-filling.

Porosity. Non-heat treatable aluminum alloys are susceptible to hydrogen-induced weld metal porosity, as are all aluminum alloys in general. This porosity forms during solidification due to the abrupt drop in hydrogen solubility when going from liquid to solid. Porosity can best be avoided by minimizing hydrogen pickup during welding.
Weld Properties

When non-heat-treatable alloys are welded, microstructural damage is incurred in the HAZ. Unlike the case of heat-treatable alloys, whose strengthening precipitates may dissolve or coarsen, the HAZ damage in non-heat-treatable alloys is limited to recovery, recrystallization and grain growth. Thus, loss in strength in the HAZ is not nearly as severe as that experienced in heat-treatable alloys. For this reason, 5xxx-series alloys are popular for use in welded pressure vessels where reasonable joint strengths can be obtained in the as-welded condition without the need for post-weld heat treatment.

The weld metal of non-heat-treatable aluminum alloys is typically the weakest part of the joint and is the location of failure when the joint is loaded. This is in contrast to most heat-treatable aluminum alloys, where the heat-affected zone often is the weakest link. The weld metal microstructure of the non-heat-treatable alloys consists of columnar-dendritic substructure that has interdendritic eutectic constituents - primarily (Fe, Mn)Al6, for 1xxx and 3xxx alloys; and Mg3Al2 for 5xxx alloys.

An important application for alloy 5083 is the construction of tactical military vehicles. The hulls and turrets of vehicles such as the M113 armored personnel carrier, the M2/M3 infantry and cavalry fighting vehicles, the M109 self-propelled howitzer, and AAV7A amphibians all consists of welded 5083 aluminum structures. There are also a myriad of brackets, clips and so on, welded to the hulls and turrets, although not normally fabricated to ballistic requirements.

The fast fracture of high-strength aluminum alloys appears to be closely involved with the fracture or cavitations of dispersoid particles, in the size range 0.1-0.5 µm. which are present in high volume fraction and which contribute to grain refinement. Two high-strength Al-Zn-Mg-Cu (7000 series) alloys have been examined: one treated with Zr (7010: Zn=6.2%, Mg=2.3%, Cu=1.9%, Zr=0.11%) and one treated with Cr (7475: Zn=5.6%, Mg=2.2%, Cu=1.6%, Cr=0.12%).

It was found that the crack paths in these alloys were of a "zig-zag" nature due to a "fast shear" fracture process, in which the critical step appears to be the decohesion of dispersoids within narrow shear bands. A model has been used to relate the toughness of the alloys to the basic tensile properties of yield stress and strain-hardening exponent, and to dispersoid parameters such as size, shape, distribution and dispersoid matrix interfacial strength. The model has also been employed to try to explain the difference between the effects of disipersoids on toughness in the Zr-treated alloys and the Cr-treated alloys.

High-strength alloys generally possess poor fracture toughness, which in many cases, appears to be a consequence of a decrease in strain-hardening capacity as the yield strength is raised. Associated with the decrease in strain-hardening capacity, it is observed that plastic flow localizes in narrow bands, and it appears that the critical amount of local deformation needed for fracture processes to operate is more readily attained than when the flow is distributed homogeneously.

The microstructures of commercial high-strength aluminum alloys are rather complex. There are three size ranges of particles:

large, iron- and silicon-rich inclusions;

intermediate chromium-, manganese-, or zirconium-rich particles, termed "dispersoids";

fine-scale aging precipitates, which harden the matrix.

While Fe and Si are present as impurities, elements such as Cr or Zr are added deliberately to control recrystallization and grain growth. The matrix strength in the peak-aged condition is primarily a function of the volume fraction of the aging precipitate, but the toughness. In general, appears to be related to the amount, type, and morphology of all three kinds of particles.

To produce high volume fractions of aging precipitates, the Al-Zn-Mg (7000) series of alloys has been developed. These contain approximately 6Zn-2,2Mg, and additions of approximately 1.7% Cu are also made to give improved resistance to stress corrosion.

The heat treatments of commercial high-strength aluminum alloys are often complicated. Two- or even three-stage aging is frequently employed, and some plastic deformation may additionally be applied before or between aging treatments. Each individual alloy is usually supplied to an optimum "temper".

In this research, simple heat treatments have been used to obtain different micro-structures and properties. Obtained "peak-aged" alloys are not therefore in the same condition as commercially "peak-aged" alloys, although the differences are probably not too large. It is, however, important to contrast the strengths and toughness which was observed with those of the alloys in commercially supplied conditions.

The crack-opening-displacement (COD) method was employed to characterize the fracture toughness, using three-point single-edge-notch (SEN) bend specimens. Testing was according to the British Standard BS 5762. Clip-gauge displacements were converted to crack-tip displacements using the formula

where
K = the stress intensity factor
Vp = plastic component of clip gauge opening displacement
E = Young’s modulus (70GN m^-2)
ν = Poisson’s ratio
σ_γ = yield stress (0,2% proof stress)
W = specimen width
a = crack length
z = knife-edge height.

The optical microstructures of the two alloys are shown in Fig.1 which clearly indicate the grain structure and inclusion distribution in three mutually orthogonal planes, oriented as shown with respect to the rolling direction: in the micrographs, RD = rolling direction, L = longitudinal direction. T = transverse direction, S = short transverse direction.

The grains and inclusions are very elongated parallel to the rolling direction but are also quite elongated in the transverse direction, suggesting that the plates had been cross-rolled.

Fig.1 Optical microstructure of 7010 alloy.

The crack paths and fracture surfaces indicate that shear fracture plays a dominant role in the fracture processes in these alloys, and the following sequence of events is suggested. After a certain amount of plastic deformation has occurred ahead of the crack tip, flow begins to localize within narrow bands.

When a critical strain is achieved in such bands, the interface between the dispersoids and the matrix decoheres. This causes the resistance to flow to drop steeply and leads to a catastrophic "unzipping" of all the dispersoids in the shear band giving total fracture throughout the band.

The critical step is supposed to be the decohesion of dispersoids in the shear bands. Some evidence of this is perhaps given by the shapes of fine dimples formed around dispersoids after large amounts of plastic deformation.

The total amount of plastic deformation is very large. Even so, the fine dimples are still equiaxed. This implies that the foundation of the fine dimples takes place at the end of the whole fracture process at high strain, suggesting that the dispersoid, matrix interface strength is large. It also means that coalescence must have taken place immediately after the voids formed by decohesion, because insufficient deformation to distort the shape of the fine dimples has taken place.

The aging precipitates are very much smaller than the dispersoids and it is supposed that they are not directly involved in the decohesion process. There is, therefore, evidence that the interface strength, size, shape, and distribution of significant factors, affecting toughness in these alloys.

Relationships between δ_i and n² and between estimated K_i and n√a. The relationships are quite linear as observed in a previous study, which collected results from 2014, 2024, 7075, and 2219 commercial alloys.

A simple model has been proposed by Hahn and Rosenfield to relate the toughness of alloys to their basic flow properties (yield stress, strain-hardening exponent).

Conclusions

Shear fracture plays a dominant role in crack extension processes in 7010 (Zr-treated) and 7475 (Cr. treated) alloys. The critical step in the fracture process appears to be decohesion of dispersoids within the shear bands.

The Hahn and Rosenfield model has been applied to relate toughness to flow properties and dispersoid parameters, and this model helps to explain the effects or dispersoids on the toughness of Zr-treated alloys and Cr. treated alloys.

The fracture characteristics of welds of various aluminum alloys were evaluated by means of tear and notch-tensile tests. The tear resistance and notch toughness of welds are generally greater than those of cold-worked or heat-treated base metal, and approach those of annealed base metal.

Subsequent thermal treatment of heat treatable filler metals may appreciably change the fracture characteristics. Except for the 7000 series of alloys, there is no significant decrease in the toughness of the welds between room temperature and -320 or -423°F (-160 or -217°C). Except for 4043 with 6061, the toughness of the weld is equal to or greater than that of the base metal.

Welding is one of the most practical methods of joining aluminum alloys, particularly for applications where components must be leak tight. The welding of almost all commercial aluminum alloys is feasible with one or more of the processes developed since 1950, the two most popular being the gas metal-arc and gas tungsten-arc.

Some alloys are, of course, more difficult to weld than others, in that more cracking occurs. Those generally considered "aircraft" alloys, principally the high-strength 2000 and 7000 series, are weldable with special techniques whereas 2219 and the "pressure-vessel" alloys, the 1000, 3000, 5000 and 6000 series, are readily welded. The choice of filler alloy is also an important factor in determining the relative ease of welding.

Welds of almost all of the aluminum alloys are so tough that accurately describing their fracture characteristics is often a problem, particularly in the case of the pressure-vessel alloys. Although data developed in accordance with fracture-mechanics concepts would be of most value to designers, it is practically impossible to develop unstable crack growth under elastic stress conditions in laboratory tests in sound aluminum-alloy welds.

A wide variety of aluminum alloy base sheet and plate and filler alloys are represented by the test data. The sheet and plate range from 1100-H112 (commercially pure aluminum in the as-rolled condition) with an ultimate tensile strength of 15,000 psi (≈110 MPa) to 7178-T6 (Al-Zn-Mg alloy, solution heat treated and artificially aged) with an ultimate tensile strength of 90,000 psi (≈630 MPa).

There are no consistent differences among the unit propagation energies for the individual alloys dependent upon specimen orientation. In the few instances where there appears to be a significant difference, the specimens in which the crack extended along the center of the weld gave the lowest values, so the use of this orientation as the standard is reasonable.

Welds of the high strength 5000 series alloys have the highest tear resistance. There is a trend toward decreasing tear resistant with increasing magnesium content. Filler alloys 2319 and 4043 have less tear resistance than the Al-Mg or, Al-Mn alloys. Aging or heat-treating and aging increases the tear resistance of 2319 welds in 2219, but heat-treating 4043 welds in 6061 markedly reduces the tear resistance.

The tear resistances of welds of the nonheat-treatable alloys are relatively high, in fact in some cases almost as high as those of the base alloys in the annealed condition (0 temper). This is to be expected since welding partially anneals the weld zone (heat-affected zone). In all cases, the unit propagation energies are about equal to or greater than those of the unwelded base metal in the cold-worked tempers.

In the as-welded or aged condition, 2319 has practically the same tear resistance as 2219-T81, T851 sheet and plate. When heat treated and aged after welding, the tear resistance of the weld metal is greater than that of 2219 T62 sheet and plate.

Welds of 4043, without subsequent treatment, have about the same tear resistance as those of 2319, and about 2/3 that of 6061-T6 plate, with which 4043 is commonly used. Heat treatment and aging after welding reduces the tear resistance of 4043 to a much lower level, to only about 15 % of that of 6061-T6 plate.

These variations in tear resistance, indicated by unit propagation energy, show that tear resistance is no simple function of either strength or ductility (elongation), but rather reflects the particular degree to which and ductility combine the material to resist crack propagation. Although it is true that general trend is toward increasing tear resistance with increasing elongation and decreasing strength, the great ductility of 1100 is more than offset by its lower strength; hence, the toughness is at a level lower than that of 5052 or 5154. These latter alloys have both high strength and great ductility.

The practical difficulty of determining empirically the plane-strain fracture toughness (K_1c) of aluminum alloy welds may be seen by spotting the unit propagation energies. Unit propagation energies for all but 4043 and 2319 are above 700in.-lb/in.2. This fact, of course, is a strong indication that problems in unstable crack propagation at elastic stresses in sound welds of these "pressure-vessel" alloys would be quite rare.

Even 2319, in various conditions, and 4043, as-welded, have considerable toughness. Some attempts to measure K_1c for many aluminum alloys with indicated toughness in this range, such as 2219-T851 and 2219-T87, have been fruitless. Values in the range of 30,000 psi√in. might be anticipated. With heat treatment and aging after welding, the toughness of 4043 is low, with a value of K_1c of about 20,000 psi√in indicated.

At -320°F (-160°C), the tear resistances of welds of most aluminum alloys are about as high as or higher than at room temperature. In some cases, notably 2319, as welded and with subsequent aging, the unit propagation energies were appreciably higher at -320°F, while for most of the alloys tested they were in the same general range — all exceptionally high. There are no suggestions of any sudden transitions in fracture behavior.

Tensile tests of severely notched specimens provide information on the ability of materials to deform plastically in the presence of stress raisers, and thereby resist crack initiation. The very sharp notch was selected because it approximates the most severe stress concentration, a crack, in the structure; and studies of various designs of notched specimens showed that the greater range of data from tests of the very sharply notched specimens provided greater discrimination between materials.

Notch-tensile strengths are calculated by dividing the fracture loads by the net area at the root of the notch. These strengths, in themselves, are not very useful as indicators of notch toughness, but their relation to the tensile yield strength provides an indication of whether or not the fracture took place with appreciable yielding.

The greater the ratio of the notch strength to the yield strength (termed the notch-yield ratio), the greater is the ability of the material to deform (yield) in the presence of the stress concentration. The ratio of the notch-tensile strength to the tensile strength (often called the notch-strength ratio or notch-tensile ratio) is not as reliable for evaluating the relative notch-toughness of materials; although, it may be useful as an indication of tensile efficiency for certain specific structural components.

By way of caution, the common practice of setting some arbitrary value of notch-yield ratio or notch-strength ratio as a limiting value may be misleading, since the value of notch-tensile strength and hence, of the ratios, are dependent upon specimen geometry. No specific value of a ratio can have the same significance with different notch designs. The data are useful primarily for rating the alloys.

Tensile tests of notched specimens have been used to evaluate the notch toughness of welds at subzero temperatures as well as at room temperature. Tensile tests of notched specimens were made at -320 and -423°F (-160 and -217°C) in the same manner as at room temperature, except that the specimens and grips were immersed in liquid nitrogen and liquid hydrogen, respectively.

Welds of 5556 filler alloy in 5456 base alloy are relatively tough, but welds of 5556 in high strength 7000-series alloys 7079-T6 and 7178-T6 are less tough. Alloy 4043 shows a similar variability; welds in alloy 6061 are tough in comparison with those in alloys 2014-T6 and 7075-T6.

For both sets of data however, it is clear that the notch toughness of welds in aluminum alloys, as measured by notch-yield ratio, is generally greater than that of cold-worked or heat-treated base metal. Since welding partially anneals the weld zone, the toughness of the joint approaches that of annealed plate. Aging after welding or heal treating and aging after welding can appreciably reduce the toughness of some combinations of alloys.

Data from tests at -320 and -423°F show that the notch-tensile strengths of the welds, as well as the tensile and joint yield strengths, are higher than at room temperature. Notch-yield ratios are generally about the same at -320°F as at room temperature; some show a decrease between -320 and -423°F, but in these cases, the notch-strength ratio remains the same at the two temperatures.

These variations of the notch-yield and notch-tensile ratios suggest that the ultimate strength of the weld influences the weld notch-strength to a greater degree than the yield strength. In all cases, the notch toughnesses of the weld at -320 and -423°F (-160 and -217°C) are superior to that of the base metal.

Summary

The fracture characteristics of aluminum alloy welds have been evaluated from the results of tear and notch-tensile tests. Welds generally have greater tear resistance and notch toughness than potential base alloys in heat-treated or cold worked tempers; they are almost as tough as annealed base alloy plate. This should not be surprising in view of the fact that welding partially anneals the weld (heat-affected) zone.

Welds in the non heat-treatable alloys (1000, 3000 and 5000 series) are exceptionally tough. Comparison of their tear resistances and notch toughnesses with correlations established between these properties and plane-strain stress-intensity factor, K_1c, on the basis of tests of very high strength aluminum alloys, indicates that in many cases K_1c, is greater than 40,000 psi√in., so that unstable crack growth is not likely to ever be a problem in welded structures of these alloys.

The influence of subzero temperatures on the fracture characteristics has been evaluated and, except for welds in the 7000 series of alloys, there is no significant decrease in the toughness at -320 or -423°F (-160 or -217°C). Except 4043 with 6061, the toughness of the weld is greater than that of the comparable base metal.

The high temperature and/or high strength-to-weight requirements of aerospace structures, advanced propulsion systems, high-speed aircraft and deep submergence vessels have stimulated the development of certain nonferrous metals that can be used in these applications.

In addition to the failures that arise because of extreme environmental conditions, e.g., high operating temperatures and corrosive environments, or fatigue, there are also brittle failures that occur in some of these metals because of their inherent low ductility and toughness. These failures can cause problems in fabrication and in certain service applications e.g., at the low temperatures that exist in outer space.

From the point of view of catastrophic fracture, the most important and interesting of the nonferrous metals are the BCC refractories, high strength aluminum alloys and the HCP metals, magnesium, beryllium, and titanium. Although low strength aluminum, copper- and nickel-base alloys are used extensively as structural materials, these FCC metals are quite ductile and tough in the absence of extreme environments and need not to be discussed here.

Cleavage is the only mode of unstable fracture in the BCC refractory and low strength HCP metals, so that the effect of composition and microstructure on toughness and ductility can be described by variations in impact and tensile-transition temperatures; low-energy tear is the primary mode of unstable fracture in the high strength alloys so that their variations in toughness appear as variations in G1c, DWTT energy, Cv(max), and tensile ductility.

As in the case of steels, alloy content and processing conditions affect the toughness of nonferrous metals by affecting their microstructures, which, in turn, determine their toughness.

The Fracture of Magnesium
The effect of temperature and grain size on the fracture strengths of high-purity magnesium and of a Mg-2wt % Al alloy was examined. These data reveal that each material exhibits two characteristically different types of fracture behavior. Over the higher range of temperatures the fracture stress decreases in a manner that parallels the flow-stress temperature relationship, and the fracture is of a ductile type. Over the lower range of temperatures the true fracture stress is independent of temperature but highly dependent upon grain size. As the grain size is decreased the fracture stress increases, and the temperature of transition to the low temperature fracture behavior range is reduced. The addition of aluminum to magnesium has the important effect of increasing the fracture stress at a given grain size as well as reducing the transition temperature.

In polycrystalline pure magnesium and in commercial magnesium alloys a significant number of microcracks are initiated at a strain which is about equal to one-half of the total strain to fracture. A significant amount of the total plastic deformation therefore occurs in the microcracks grow and link up in the final stages of fracture. The ductility of these alloys increases abruptly with increasing testing temperature at temperatures slightly above ambient. This ductility rise results from a microstructural instability, that is, recovery and recrystallization during straining, of the metals at the testing temperature.

At room temperature, failure of polycrystalline high-purity magnesium occurs on various crystallographic planes of high index as well as by an intergranular mechanism. Fracture at low temperatures is also intergranular along with complex transcrystalline modes.

The addition of lithium to magnesium alloy has an interesting effect on low-temperature fracture behavior. The stress for prismatic slip is reduced relative to that for basal slip, and thereby the low-temperature ductility increased. The greater facility for slip is also reflected in a decrease in strain-hardening behavior as barriers to plastic deformation can be more easily circumvented. In magnesium alloys of high lithium content (14.5%), fracture even at 78°K occurs by necking, and there is no leveling of the fracture stress as in the case of pure magnesium.

The Fracture of Titanium
Titanium alloys can be grouped into three categories according to the predominant phase or phases in their microstructure. In pure titanium the α-phase, an HCP structure, is stable to 1625°F (885°C), and the BCC β-phase is stable from 1625°F to the melting point of 3130°F (1740°C).

Certain alloying additions stabilize the α-phase. Among these are aluminum and the interstitial elements carbon, hydrogen, nitrogen, and oxygen. Most alloying elements, such as chromium, columbium, copper, iron, manganese, molybdenum, tantalum, and vanadium stabilize the β-phase to the extent that a mixed α-β-phase or an entirely β-phase alloy can persist down to room temperature.

Some elements, notably tin and zirconium, behave as neutral solutes in titanium and have little effect on the transformation temperature, acting as strengtheners of the α-phase. The single-phase α-alloys are not heat-treatable, but both the α-β-alloys and the β-alloys can be strengthened by heat treatment. The yield strengths of these alloys range from 30,000 psi for commercially pure titanium up to 250,000 psi for certain α-β-alloys.

One group of elements that affects the low-temperature fracture behavior, especially of notched specimens, are the interstitials. Because of the deleterious effect of interstitials, extra-low interstitials grades (ELI) are available for use where structural reliability at low temperatures is of concern.

The Fracture of Beryllium
Beryllium can be melted and cast but, as the castings are brittle and difficult to machine, practically all the beryllium used in space, nuclear, and other applications is made from beryllium powders and undergoes powder metallurgy processing. The powder is usually fabricated by the mechanical commutation of vacuum-cast ingots of magnesium-reduced beryllium pebbles or electrolytically produced beryllium flakes. Each of the powder particles is enveloped by BeO, which tends to hinder grain growth during processing.

Hydrogen, which can embrittle titanium, is soluble to a negligible extent in beryllium and has no effect on the transition from ductile to brittle behavior. Oxygen is able to embrittle beryllium, but there is no quantitative measure of the effect.

Summary

1. Commercially available group Va refractory metals (V, Nb, Ta) are considerably more ductile and have considerably lower transition temperatures than commercially available group VIa refractory metals (Cr, Mo, W). This results from the fact that the solubility of impurities in the Va metals is much higher than in the VIa metals. Consequently, at the impurity contents that normally exist in commercial materials, impurity atoms and/or second phase particles are segregated at grain boundaries in VIa metals. These lower the plastic work expended in microcrack formation γm, and hence lower the toughness.
2. Oxygen is a particularly strong embrittling agent for molybdenum and tungsten; nitrogen and sulfur are particularly harmful in wrought chromium. Hydrogen is the most effective embrittling agent for Va metals.
3. After impurity content, grain size is the most important variable affecting the ductility and toughness of Va and VIa refractory metals. In general, ductility and toughness are increased by grain refinement.
4. Cold or warm working and alloying provide the most efficient means of increasing the ductility and toughness of refractory metals. The beneficial effects of working are thought to result from the production of a fibrous microstructure, which permits splitting and crack tip blunting when toughness is evaluated in the working direction. Alloying can improve ductility and toughness in VIa metals by lowering the tendency for impurity segregation at grain boundaries or by getting out the impurities.
5. The fracture behavior of the HCP metals beryllium, magnesium, and titanium, together with that of some of their alloys, has been reviewed. Brittle fracture at low temperatures is common to these HCP metals and, in both titanium and beryllium, interstitial elements such as oxygen influence this behavior.
6. In each of these metals the decrease in the number of modes with decreasing temperature is a common feature involved in low ductility fractures.

Aluminum alloys are progressively used in the automobile industry due to several advantages such as low specific weight, good formability, good corrosion resistance and a nice surface appearance. The standard production forming processes such as extrusion and forging, can give rise to large variations in the tensile, fatigue and fracture properties. In AlMgSi alloys (6061, 6062, 6060 and 6082), yield stress have been shown to have only a weak dependence on grain size. However, a large part of the variations in other properties can be traced back to differences in grain size.

Production of Universal Joints. Cold forged products are strong, tolerances are tight and strength/ductility can be optimized by tempering Joints having different grain morphologies are made from two conditions of alloy AA 6082: Alloy A which is rich in Mg and the low Mg containing alloy B also specified with Cr. Three different forging and heat treatments are employed to obtain different grain structures in the alloy A. The alloy B series is similarly produced as Series 1 of the alloy A. All series are artificially aged to the T6 condition. After forging the joints are machined to the shown geometry. A parallel set of joints is machined with a serrated attachment hole.

Characterization of Joints. Metallographic characterization is carried out in scanning (SEM) and transmission electron microscopes (TEM). Tensile specimens are taken from the joint arm, two specimens from each joint. Fatigue life tests are conducted at a fully reversed torsion moment of ±70Nm at 10Hz in moist air environment.

Characterization of laboratory material. Additional laboratory investigations involve five different grain morphologies in the alloy A (T6 temper). Fatigue life is determined from 3-5 electropolished specimens of each condition at R = -1 and 20 Hz. Further, fracture toughness in terms of K, is characterized by loading 9.5 mm diameter short rod specimens (W=14.5 mm, t = 0.2 mm) in a servohydraulic MTS machine. The specimen geometry and test performance is in accordance with the proposed ASTM standard.

Fatigue and mechanical properties of universal Joints. The alloy B series and Series 1 of the alloy A have a larger elongation at fracture (ef), than the two other series (Table 1). The average fatigue life tends to be longer in Series 2 and 3 than in the other series. Joints having serrated attachment holes have the fatigue life lowered by a factor ranging between 2 and 14 when compared to the smooth hole geometry, and all fatigue cracks are initiated in serrations. An important observation is that the microstructures of Series 2 and 3 and the alloy B, seem to be more notch sensitive with respect to fatigue than the Series 1. Furthermore, enhanced fatigue life of the relatively coarse grained joints with serrated holes (alloy B and Series 1), may be due to roughness induced crack closure reducing the crack driving force.

TABLE 1 - Mechanical Properties and Fatigue Life at 70Nm. R = -1L
Aluminum Universal Joints. E-modulus 73 GPa.

Alloy	Series	Serration	σ0.2 (MPa)	σm (MPa)	ef %	Nf (106 cycles)
A	1	No Yes	332	352	11	1.0 +/- 0.4 0.4 +/- 0.3
A	2	No Yes	325	330	6	1.5 +/- 0.6 0.1 +/- 0.02
A	3	No Yes	328	334	8	1.6 +/- 0.7 0.2 +/- 0.09
B	-	No Yes	321	352	12	1.4 +/- 0.2 0.2 +/- 0.08

Alloys with precipitate free zones (PFZ), as with the alloy B and Series 1, have been shown to generate higher closure levels than alloys without PFZ.

Microstructure of universal Joints. The most significant variation in the microstructure due to changes in the heat treatment/deformation procedure is in the grain structure morphology. Series 1 of the alloy A has a partly elongated recrystallized structure, i.e. 40 μm high vs. 60 μm long. Series 2 and 3 are both unrecrystallized having a short and a long fiber-shaped structure respectively. The alloy B series has a recrystallized almost equiaxed grain structure, -140 μm in diameter.

SEM particle analyses show that the alloy B series has less of the small (-1 μm) and significantly more of the coarse (>5 μm) particles than Series 1. This can contribute to the formation of a coarse grain structure in the alloy B.

TEM studies show that Series 2 and 3 have both a higher dislocation density and a coarser precipitate structure than Series 1, i.e. slightly overaged. These two facts may govern the lower the tensile ductility and the increased fatigue notch sensibility in the Series 2 and 3.

Correlation to laboratory material data. In general, the laboratory materials show better tensile properties than the joints. Taking into account the observed coarse precipitate structures in the joints, i.e. Series 2 and 3, the lower strength and ductility of these joints are probably due to overaging effects.

The non-recrystallized fiber structure has the highest fracture toughness. Further, the recrystallized (50 μm) and the deformed fiber structure show a significantly lower K, than the undeformed fiber structure.

From the smooth specimen S-N curves of the five laboratory conditions, it is easily seen that the non-recrystallized fiber structure has better fatigue resistance than the recrystallized structure. This is in agreement with the fatigue life data of the joints.

Conclusions

• The non-recrystallized fiber structure has better fatigue resistance and fracture toughness than recrystallized coarse-grained microstructures.
• Processing cold forged/extruded aluminum automobile components with a property optimized microstructure demands extensive knowledge of the interdependence between alloy composition, thermal treatments, particle and precipitate structures and the deformation procedure.

The engineering tension test is widely used to provide basic design information on images/the strength of materials and as an acceptance test for the specification of materials. In the tension test a specimen is subjected to a continually increasing uniaxial tensile force while simultaneous observations are made of the elongation of the specimen. An engineering stress-strain curve is constructed from the load elongation measurements (Fig. 1).

It is obtained by dividing the load by the original area of the cross section of the specimen.

(1)

The strain used for the engineering stress-strain curve is the average linear strain, which is obtained by dividing the elongation of the gage length of the specimen, d, by its original length.

(2)

Since both the stress and the strain are obtained by dividing the load and elongation by constant factors, the load-elongation curve will have the same shape as the engineering stress-strain curve. The two curves are frequently used interchangeably.

The shape and magnitude of the stress-strain curve of a metal will depend on its composition, heat treatment, prior history of plastic deformation, and the strain rate, temperature, and state of stress imposed during the testing. The parameters, which are used to describe the stress-strain curve of a metal, are the tensile strength, yield strength or yield point, percent elongation, and reduction of area. The first two are strength parameters; the last two indicate ductility.

The general shape of the engineering stress-strain curve (Fig. 1) requires further explanation. In the elastic region stress is linearly proportional to strain. When the load exceeds a value corresponding to the yield strength, the specimen undergoes gross plastic deformation. It is permanently deformed if the load is released to zero. The stress to produce continued plastic deformation increases with increasing plastic strain, i.e., the metal strain-hardens. The volume of the specimen remains constant during plastic deformation, A·L = A₀·L₀ and as the specimen elongates, it decreases uniformly along the gage length in cross-sectional area.

Initially the strain hardening more than compensates for this decrease in area and the engineering stress (proportional to load P) continues to rise with increasing strain. Eventually a point is reached where the decrease in specimen cross-sectional area is greater than the increase in deformation load arising from strain hardening. This condition will be reached first at some point in the specimen that is slightly weaker than the rest. All further plastic deformation is concentrated in this region, and the specimen begins to neck or thin down locally. Because the cross-sectional area now is decreasing far more rapidly than strain hardening increases the deformation load, the actual load required to deform the specimen falls off and the engineering stress likewise continues to decrease until fracture occurs.

Tensile Strength

The tensile strength, or ultimate tensile strength (UTS), is the maximum load divided by the original cross-sectional area of the specimen.

(3)

The tensile strength is the value most often quoted from the results of a tension test; yet in reality it is a value of little fundamental significance with regard to the strength of a metal. For ductile metals the tensile strength should be regarded as a measure of the maximum load, which a metal can withstand under the very restrictive conditions of uniaxial loading. It will be shown that this value bears little relation to the useful strength of the metal under the more complex conditions of stress, which are usually encountered.

For many years it was customary to base the strength of members on the tensile strength, suitably reduced by a factor of safety. The current trend is to the more rational approach of basing the static design of ductile metals on the yield strength.

However, because of the long practice of using the tensile strength to determine the strength of materials, it has become a very familiar property, and as such it is a very useful identification of a material in the same sense that the chemical composition serves to identify a metal or alloy.

Further, because the tensile strength is easy to determine and is a quite reproducible property, it is useful for the purposes of specifications and for quality control of a product. Extensive empirical correlations between tensile strength and properties such as hardness and fatigue strength are often quite useful. For brittle materials, the tensile strength is a valid criterion for design.

Measures of Yielding

The stress at which plastic deformation or yielding is observed to begin depends on the sensitivity of the strain measurements. With most materials there is a gradual transition from elastic to plastic behavior, and the point at which plastic deformation begins is hard to define with precision. Various criteria for the initiation of yielding are used depending on the sensitivity of the strain measurements and the intended use of the data.

1. True elastic limit based on micro strain measurements at strains on order of 2 x 10^-6 in | in. This elastic limit is a very low value and is related to the motion of a few hundred dislocations.
2. Proportional limit is the highest stress at which stress is directly proportional to strain. It is obtained by observing the deviation from the straight-line portion of the stress-strain curve.
3. Elastic limit is the greatest stress the material can withstand without any measurable permanent strain remaining on the complete release of load. With increasing sensitivity of strain measurement, the value of the elastic limit is decreased until at the limit it equals the true elastic limit determined from micro strain measurements. With the sensitivity of strain usually employed in engineering studies (10^-4in | in), the elastic limit is greater than the proportional limit. Determination of the elastic limit requires a tedious incremental loading-unloading test procedure.
4. The yield strength is the stress required to produce a small-specified amount of plastic deformation. The usual definition of this property is the offset yield strength determined by the stress corresponding to the intersection of the stress-strain curve and a line parallel to the elastic part of the curve offset by a specified strain (Fig. 1). In the United States the offset is usually specified as a strain of 0.2 or 0.1 percent (e = 0.002 or 0.001).
   

   (4)

A good way of looking at offset yield strength is that after a specimen has been loaded to its 0.2 percent offset yield strength and then unloaded it will be 0.2 percent longer than before the test. The offset yield strength is often referred to in Great Britain as the proof stress, where offset values are either 0.1 or 0.5 percent. The yield strength obtained by an offset method is commonly used for design and specification purposes because it avoids the practical difficulties of measuring the elastic limit or proportional limit.

Some materials have essentially no linear portion to their stress-strain curve, for example, soft copper or gray cast iron. For these materials the offset method cannot be used and the usual practice is to define the yield strength as the stress to produce some total strain, for example, e = 0.005.

Measures of Ductility

At our present degree of understanding, ductility is a qualitative, subjective property of a material. In general, measurements of ductility are of interest in three ways:

1. To indicate the extent to which a metal can be deformed without fracture in metalworking operations such as rolling and extrusion.
2. To indicate to the designer, in a general way, the ability of the metal to flow plastically before fracture. A high ductility indicates that the material is "forgiving" and likely to deform locally without fracture should the designer err in the stress calculation or the prediction of severe loads.
3. To serve as an indicator of changes in impurity level or processing conditions. Ductility measurements may be specified to assess material quality even though no direct relationship exists between the ductility measurement and performance in service.

The conventional measures of ductility that are obtained from the tension test are the engineering strain at fracture ef (usually called the elongation) and the reduction of area at fracture q. Both of these properties are obtained after fracture by putting the specimen back together and taking measurements of L_f and A_f .

	(5)
	(6)

Because an appreciable fraction of the plastic deformation will be concentrated in the necked region of the tension specimen, the value of e_f will depend on the gage length L₀ over which the measurement was taken. The smaller the gage length the greater will be the contribution to the overall elongation from the necked region and the higher will be the value of e_f. Therefore, when reporting values of percentage elongation, the gage length L₀ always should be given.

The reduction of area does not suffer from this difficulty. Reduction of area values can be converted into an equivalent zero-gage-length elongation e₀. From the constancy of volume relationship for plastic deformation A*L = A₀*L₀, we obtain

(7)

This represents the elongation based on a very short gage length near the fracture.

Another way to avoid the complication from necking is to base the percentage elongation on the uniform strain out to the point at which necking begins. The uniform elongation e_u correlates well with stretch-forming operations. Since the engineering stress-strain curve often is quite flat in the vicinity of necking, it may be difficult to establish the strain at maximum load without ambiguity. In this case the method suggested by Nelson and Winlock is useful.

Precipitation or age hardening was discovered by Alfred Wilm in Germany in 1906. He attempted to harden an alloy of essentially aluminum-2 atom percent copper in an analogous way to steels by a quenching treatment. The specimen was initially soft, but the hardness increased with time at room temperature after the quench.

Merica, Waltenberg and Scott first attributed the hardening to a precipitation effect. This was an extremely important paper for it pointed the way to development of a whole host of precipitation-hardened alloys. They correctly pointed out that the solubility of copper in aluminum decreases markedly on cooling and that quenching gives a supersaturated solid solution.

The change in solubility with temperature in the terminal solid solution is typical for age-hardening systems. At room temperature the stable state of an aluminum-2 atom percent copper alloy is an aluminum-rich solid solution (α) and an intermetallic phase with a tetragonal crystal structure having nominal composition CuAl2 (θ). According to the suggestion of these authors, fine particles of θ (and similar precipitates in other systems) form during aging. These were visualized to lie astride and key the slip planes, a proposal put forth by Jeffries and Archer.

However, with the treatment to give maximum hardness, in general precipitate particles are not visible with the highest-resolution light microscope. When the particles are visible the alloy has overaged. Not only is the hardness less in the overaged state but usually the alloy is very much less ductile compared to maximum hardness.

The mystery was solved first in aluminum-copper alloys independently by Guinier and Preston in 1937 by careful X-ray diffraction work. Diffuse scattering occurs outside but associated with the Bragg reflections of the solid solution, and these are due to regions in the matrix solid solution enriched in solute atoms. Small solute-enriched regions in a solid solution where the lattice is identical or somewhat perturbed from that of the solid solution are called Guinier-Preston zones.

Precipitation in Al-Cu Alloys
In aluminum-2 atom percent copper, four different precipitates occur and each may be made to form by controlling the heat treatment. The occurrence of metastable precipitates is quite common in precipitation-hardening alloys. Typical system for describing in details the sequence of events is aluminum-copper although the phenomena mentioned here occur in many other systems.

Guinier-Preston zones of the first kind (GP-I) are plates of copper atoms one or two atoms thick and commonly 25 atoms in diameter oriented parallel to {100} planes in the aluminum-rich matrix. Since the sizes of copper and aluminum atoms differ by about 12 per cent, the lattice is distorted in the regions of the zones. As a matter of fact, the zones here form as thin platelets to minimize strain energy. They form parallel to {100} because the elastic modulus is least in this direction. GP-I forms at room temperature and is the first precipitate to form at 100°C. It is not stable at 210°C; GP-1 formed at a lower temperature rapidly dissolves in about 30 seconded if the metal is heated to 210°C. This is called reversion or retrogression.

Guinier-Preston zones of the second kind (GP-II) are thicker (10 atoms) and of larger diameter (75 atoms) than GP-I, but they are not just big GP-I precipitates. In GP-II an ordering of aluminum and copper atoms occurs to give an average composition of about Cu2Al5. GP-II is the second precipitate to form at, say, 130°C or the first at 210°C. The strongest aluminum-2 atom percent copper at room temperature contains mainly GP-II.

θ’ is a third metastable precipitate. It has the nominal composition CuAl2 and is tetragonal with a lattice distorted from θ so that it may form nearly coherently or epitaxially with the aluminum-rich matrix. It forms in the matrix in a Widmenstatten pattern. The arrangements of atoms in the interface or habit plane are nearly identical in θ’ and the matrix. θ’ begins to form later than GP-II at 130 or 210°C.

Actually all types of precipitates may give hardening but GP zones and ordinary precipitates with some degree of coherency give greater hardening. In some systems, maximum precipitate on hardening occurs with GP zones (aluminum-copper and aluminum-zinc); in some systems maximum hardening occurs with a coherent ordinary precipitate (nickel-titanium, aluminum; nickel, chromium-titanium, aluminum; and aluminum-silver).

Guinier divides GP zones into two classes: ideal and nonideal. The atomic sizes of zinc and silver differ little from that of aluminum. With these, the lattice of the aluminum-rich matrix is not distorted very much in the region of the zones, which are spherical in shape. These are classed as ideal zones. The zones in aluminum-copper are an example of nonideal zones.

An important consideration is why metastable precipitates such as GP zones and θ’ should form at low temperatures, rather than the stable precipitate. The answer lies in the theory of nucleation and diffusion-controlled growth. In these cases a large surface energy exists between the stable phase and the matrix, but the GP zone is a perturbation of the matrix and the surface energy is small, indention is difficult for the former, easy for the latter. Very little diffusion is required to nucleate a GP zone. A large number of small part ides are able to form during the quench from the solution-treating temperature and on the subsequent low-temperature precipitation heat treatment. Quenched-in vacancies play an important role in facilitating diffusion.

If extra vacancies are present — they may be put in by quenching, irradiation or cold work — diffusion processes take place at very low temperatures. Some age-hardening alloys (copper-beryllium) do not develop full hardness unless they are cold-worked before aging. The free energy of the system may, of course, be further lowered if the metastable precipitate is replaced by the stable (or a more stable) precipitate. This occurs if the temperature is raised.

Interaction of Dislocation and Precipitate
Discussion about the various precipitates in aluminum-copper showed that the problem which must be considered to understand precipitation hardening is the interaction of a dislocation with a field of obstacles. The obstacles of primary interest here are precipitates defined broadly to include Guinier-Preston zones and metastable second phases as well as stable phases.

Consider a row of such obstacles and a dislocation moving on a slip plane. For slip to occur, the dislocation must either move around the particles or through the particles. An active dislocation will select from the various paths available to it the path where the least energy is expanded.

The dislocation may avoid the particles or obstacles by leaving the slip plane in the vicinity of each particle, or it may avoid the particles by the Orowan mechanism. In this mechanism, the dislocation bends between the particles leaving a dislocation ring about each particle. In either case, energy must be supplied to increase the total length of dislocation line; the stress required is, neglecting a numerical factor, roughly (Gb)/ L where G is the shear modulus, b is the Burgers vector, and L is the spacing between obstacles.

Cutting of Particles by Dislocations
In all age-hardened or precipitation-hardened alloys, the particles are generally cut during plastic deformation when the metal is aged to maximum hardness; avoiding of the particles corresponds to the overaged state.

For a dislocation to move through a particle, energy must be supplied for three basic processes, which may be involved in the cutting.

* First, if the particle is formed by a solid-state reaction or if there is a difference in coefficient of expansion, there may be elastic misfit stresses between the particle and the matrix since the particle will generally occupy a different volume than the parent phase it replaced.
* Second, the surface area of the particle is increased by cutting it and slipping the two halves.
* Third, the flow stress for moving a dislocation inside the particle may be larger than that in the matrix.

However, we will first briefly mention dislocation pinning effects due to the presence of particles. Dislocations may catalyze precipitation, and then precipitates may form at dislocations. This is thought to be the source of the yield drop commonly observed in iron and mild steel. Yield-drop effects are not commonly observed in precipitation-hardened alloys with GP zones or with a general distribution of precipitates.

Strengthening by Elastic Misfit Stresses
The importance of elastic misfit strains in age hardening has long been recognized going back to Rosenhain. Transition phases such as θ’ in aluminum-copper or γ’ in nickel-base alloys form with close matching across the habit or interface plane. The lattice must be strained to give matching.

Very large equivalent stresses may be computed if the mismatch is assumed to be taken up completely by strain. Recently was shown that the maximum precipitation hardening in nickel-base aluminum-titanium alloys from γ’ occurs at the proportion of aluminum to titanium giving the largest difference in lattice parameter between the precipitate (at large particle size) and matrix.

The dislocation theory of the strengthening from misfitting precipitates is due to Mott and Nabarro. Consider first that particles are so closely spaced that the dislocation must move essentially as a rigid line, that is, the maximum bending of the dislocation between particles is negligible.

The minimum radius of curvature to which a dislocation may be bent under the applied stress is larger than the spacing between particles on the slip plane. Depending on the relative orientation of the dislocation and precipitate, the internal stress will aid dislocation motion or hinder dislocation motion. The sum-total effect on a length of dislocation which is long compared to the spacing is zero; no strengthening is predicted to a first approximation.

The situation is analogous to a solid solution. Second, Mott and Nabarro considered the case where the particle spacing is not negligible compared to the minimum radius of curvature of the dislocation, but are of the same order of magnitude. The dislocation will tend to be wavy, assuming a position in the stress fields of the particles, which minimizes the total self-energy of the dislocation. The dislocation does not assume exactly the position corresponding to the minimum self-energy per unit length, but the dislocation is shortened a little by the line tension of the dislocation.

Strengthening from Modulus Change
It was pointed out that variation of the elastic module between the precipitate and matrix may be a source of strengthening. A general theory of stresses about particles must include misfit as well as variation in the elastic module. Further, the self-energy of the dislocation depends on G; therefore, if G of the particle is larger than that of the matrix, an extra stress will be required to force the dislocation through the particle.

The point values of the module will be functions of local composition since the module are related to the second derivatives of the interatomic interaction energies with respect to interatomic distance. Anisotropy in the elastic constants must also be taken into consideration.

A larger number of tests have been used in an effort to measure or predict the formability of sheet materials. Many of these have been criticized because of cost, complexity, difficulty in the analysis of data, lack of correlation between laboratory results and field forming performance, etc. Some of these problems may be overcome when test procedures have been standardized and a better understanding of the mechanics of the tests is achieved.

There will always be a need for formability tests. The effect of composition and processing modifications on formability must be determined during alloy development, preferably without resorting to expensive field forming trials in the initial stages.

Tests are often needed in the analysis of field forming problems requiring the comparison of problem lots to a data base. Tests are also needed for quality assurance, especially since it appears that many sheet users are working toward the use of test results as acceptance criteria.

Part of the lack of correlation between laboratory test results and field forming performance is due to a misuse of the test results. This lack of correlation leads to a lack of faith in the test procedures. If such a correlation is to be expected, the strain state of the test must match the strain state in which the failure occurred in the field. One should not expect the results of a drawability test to correlate with field failures which occurred in plane strain tension. This is because micro structural features respond differently to different states of stress, which has been demonstrated for both precipitation and dispersion strengthened aluminum sheet alloys.

Numerous questions arise when formability tests are considered as quality assurance tools or when correlations between laboratory and field results are sought. Tool geometry, lubrication, sample thickness and test procedure have all been shown to influence test results. All of these factors contribute to the between-lab variability, which has been shown to be large, and make data analysis difficult.

All sheet metal forming operations are combinations of stretching, bending and drawing. The formability tests available under each of these categories, which is a two-way flowchart to aid in the solution of metal forming problems or to aid in the screening of materials where forming is critical. When is used as a screening aid during alloy development, tests characterizing performance in all three modes may be necessary.

When a field forming problem is encountered, grid strain analysis (GSA) and the limiting strain curve are used to determine the strain state at failure and to decide if the problem is material or tooling related.

Stretching can be subdivided into uniaxial, biaxial and plane strain modes. There is some question as to whether a microstructure yielding good formability in one of these stretching categories will give good formability in the other categories.

Uniaxial tension

The uniaxial tension test is the only commonly used test for the uniaxial stretching mode. The most commonly used parameters calculated from tension test results are: the yield strength, the ultimate tensile strength, the percent total elongation in a standard gauge length (usually 50.8 mm), the percent uniform elongation and the percent reduction in area. These are commonly referred to as the mechanical properties.

The elongation values depend upon the gauge length used. A more fundamental, gauge length independent measure of ductility is the reduction in area, % RA, which is given by

%RA=[1-(A_t/A_o)]·100%

A_o, A_t original and final cross-sectional areas, respectively

The measurement of the final cross-sectional area may be difficult, which has led to another approach to obtaining a fundamental ductility parameter. Elongation surveys consist of measuring the elongation over many gauge lengths and extrapolating to zero gauge length, with this extrapolated value being the fundamental value.

Other parameters can also be calculated from uniaxial tension test results. True stress-true strain data can be fit to the Hollomon equation, where true stress, σ, is given in terms of true strain, ε, the strain hardening exponent, n, and the flow strength, K, by

σ=Kεⁿ

For steel, the strain hardening exponent correlates well with stretch formability. However, for aluminum alloys, the strain hardening exponent alone does not adequately predict formability.

The plastic strain ratio (or normal anisotropy value), r, is often calculated from uniaxial tension test results, and is given by

r=ε_w/ε_t

ε_w, ε_t et true width and thickness strains, respectively.

The plastic strain ratio may be calculated at fracture, at a constant strain or plots of r versus longitudinal strain may be made by continuously measuring ε_w and ε_t.

The prediction of "formability" in modes other than uniaxial stretching from parameters calculated from uniaxial tension test results, often referred to as "forming indices", has often been attempted by many persons in the sheet metal forming industry.

Because parameters such as total elongation, uniform elongation, plastic strain ratio, strain hardening exponent, etc., depend upon microstructure, they cannot be varied independently. It is difficult to assign quantitative values to the relative influence each of these will have on formability, although their qualitative effects are easily rationalized. Regression models can be used to predict formability in terms of these parameters, but these models are specific to given forming operations.

In summary, the uniaxial tension test should not be considered a formability test. Although the results will correlate with field failures in the uniaxial stretching mode, few forming operations result in failures which occur in the uniaxial stretching mode. The test should be used to characterize the mechanical properties of the material and to check for proper temper.

Biaxial tension

The hydraulic bulge test yields stress-strain data in the balanced biaxial stretching mode. The advantage of the hydraulic bulge test is that the strain hardening ability of a material at strains approaching those experienced in actual forming operations can be evaluated.

Strain rate sensitivity is a very important aspect of material behavior. When gradients in strain exist, materials which harden with increasing strain rate will distribute deformation more uniformly because additional deformation in areas of high strain rate, such as neck, will require greater stress. Ductility generally increases with increasing strain rate sensitivity and small changes in rate sensitivity may result in significant changes in the distribution of strain.

Strain rate sensitivity has also been measured using the hydraulic bulge test. Two methods have been used to evaluate strain rate sensitivity. The first involves several abrupt changes in strain rate during a single test. The second requires that two tests be performed at different constant strain rates. The results of the two methods have been shown to be comparable.

In summary, the hydraulic bulge test yields information about the strain hardening and strain rate sensitivity characteristics of the material. In addition, the true strain at fracture is measured. This information is not complicated by the effects of friction.

Plane strain tension

Plane strain tension tests, following the work of Wagoner are in the development stage. Currently, the elongation at fracture in the plane strain state can be evaluated for uniaxially loaded samples. This elongation value can be used in conjunction with those from the uniaxial and biaxial tension tests to plot an approximate limiting strain curve.

A method for obtaining a stress-strain relationship in the plane strain state should be sought. This will allow the strain hardening behavior of aluminum alloys to be studied in uniaxial, biaxial and plane strain tension.

It is unclear as to whether the effect of microstructure on strain hardening will be identical in all three stretching modes. It has been shown for some aluminum alloys that factors reducing the strain hardening capacity in the uniaxial and biaxial modes also reduce strain hardening capacity in plane strain. It may be determined that information obtained from the hydraulic bulge test will adequately predict formability in all three stretching modes.

Attempts should be made to correlate plane strain tension test results with simulative formability test results and field forming test results. Until such work is completed and a method for obtaining a plane strain stress-strain curve is developed, the test should not be considered a formability test and should not be widely used.

Limiting dome height

The limiting dome height (LDH) test has been proposed as a laboratory formability test showing good correlation with press formability. Rectangular blanks of various widths are rigidly clamped in the longer direction and stretched by a hemispherical punch. Transverse constraint, varied by blank width and lubrication, controls the amount of material drawing-in. The height of the dome at peak load, which reflects the combined effects of strain hardening characteristics and limiting strain capability of the material, is used as a measure of stretch-formability.

The effect of friction on the dome height must be held constant in order to evaluate the relative stretchability of alloys. In the past, samples have been solvent cleaned and tested dry in an attempt to hold friction constant at a high value. The test should be continued to be performed dry for the plane strain condition, which is of greatest interest, until that method has been developed.