7.3 Stainless Steel Weldments

Stainless steels are widely used in various industries where corrosion is of particular concern.These materials can be classified into four main categories, based on their microstructure:

(i) Austenitic stainless steels.(ii) Austenitic-ferritic (duplex) stainless steels.(iii) Ferritic stainless steels.(iv) Martensitic stainless steels.

As shown in Table 7.3, the welding of stainless steels is encumberred by a number of differ-ent metallurgical problems, including solidification cracking, hydrogen cracking, precipita-tion reactions, and grain growth. Some of these problems will be discussed below in the lightof information available in recent literature.75"78

7.3.1 HAZ corrosion resistance

Austenitic stainless steels containing about 0.07 wt% carbon or more are often susceptible tointergranular corrosion in the weld HAZ due to chromium carbide precipitation (see Section6.4 in Chapter 6). This phenomenon is frequently referred to as the 'weld decay' in the scien-tific literature.76

Table 7.3 Characteristics of stainless steel weldments.

Material

Austeniticstainless steels1"7'

Duplex stainlesssteels1"'*''

Ferritic stainlesssteels1"*

Martensiticstainless steels1"'*'5

Chemical Composition

Major Elements

Cr: 17-25 wt%Ni: 8-20 wt%Mn: ~2 wt%C: 0.03-0.15 wt%

Cr: 18-27 wt%Ni: 7-8 wt%Mo: \-A wt%C:

Tem

pera

ture

Tem

pera

ture

Like weld decay, knife-line attack is also related to intergranular C^Q-precipitation, but itdiffers from the former type in that the corrosion attack occurs in a narrow region adjacent tothe fusion boundary following multipass welding of Ti/Nb stabilised stainless steels. As shownin Fig. 7.46, the origin of knife-line attack can be attributed to dissolution of TiC or NbCduring the initial weld thermal cycle in regions of the HAZ where the peak temperature hasbeen above say 12500C. Since the cooling rate within the high temperature regime is veryhigh, reprecipitation of TiC or NbC does not take place on cooling. This leaves carbon free toreact with chromium during subsequent reheating to peak temperatures of about 700-8000Cfollowing deposition of the second layer (filler pass).

Knife-line corrosion attack is usually observed in weldments where the second pass (top)

(a)Dissolution ofTiC or NbC

C-curve for precipitationof TiC or NbC

1st thermalcycle

Time

(b)

C-curve for precipitationof TiC or NbC

C-curve for precipitationOfCr23C6

2nd thermalcycle

TimeFig. 7.46. Mechanisms of knife-line corrosion attack in austenitic stainless steel weldments (schematic);(a) Dissolution of TiC or NbC during the initial weld thermal cycle, (b) Precipitation of Cr23C6 in thelow peak temperature region of the weld HAZ following deposition of the second layer. (Thecorresponding C-curves for precipitation of TiC/NbC and Cr23C6 in the grain growth zone are displacedfar to the right in the diagram.)

overlaps the lower part of the fusion boundary of the first (root) pass, as indicated in Fig.7.47(a). In practice, the problem can be eliminated by simply reversing the welding sequenceor by changing the welding procedure so that the second pass overlaps the middle rather thanthe lower part of the root pass. The latter point is illustrated in Fig. 7.47(b).

732 HAZ strength levelIn single phase materials such as austenitic stainless steels, the primary grain size will make adirect contribution to strength. In general, the grain size dependence of the yield stress (

It follows from the graphical representation of the Hall-Petch relation in Fig. 7.48 that thevariation in the yield strength across the HAZ is rather small under the prevailing circum-stances. In fact, the maximum HAZ strength reduction which may occur in such materialsbecause of grain growth is about 18%, corresponding to high heat input welding conditions.This implies that welding does not impose severe restrictions on the design stress as long as thesteel is used in the fully annealed condition.

7.3.3 HAZ toughness

All body-centred cubic metals (including ferritic and martensitic stainless steels) show a markedtemperature dependence of the fracture toughness, as indicated in Fig. 7.49. At high tempera-tures fracture occurs normally by ductile rupture, whereas cleavage is the dominating fracturemode at low temperatures. This type of behaviour makes both ferritic and martensitic stainless

from which

Readings from the nomograms in Fig. 5.30(b) give the HAZ grain size profile shown in Fig.7.48. In order to obtain the resulting HAZ strength distribution, it is necessary to fix the valueof the friction stress, Oi, in the Hall-Petch relation. In the present example, we have:

Provided that the conditions for one dimensional heat flow are met, estimate on the basis ofthe nomograms in Fig. 5.30(b) (Chapter 5) the variation in the austenite grain size across theHAZ after welding. Calculate then via the Hall-Petch relation (equation (7-21)) the expectedreduction in the HAZ strength level due to this change in the microstructure.

Input data:Base metal yield strength: 300 MPaLocking parameter in Hall-Petch relation: k = 227 MPa (im1/2

SolutionFirst we need to calculate the net heat input per mm2 of the weld:

applied heat input. This means that a permanent soft zone will form adjacent to the fusionboundary after welding, which may reduce the overall load-bearing capacity of the joint.

Example (7.5)Consider plasma arc butt welding of a 5mm thick plate of type 316 austenitic stainless steelunder the following conditions:

Abso

rbed

energ

y

Auste

nite

grain

siz

e, \im

Yield

st

reng

th, M

Pa

Peak temperature, 0C

Fig. 7,48. Computed HAZ grain size and yield strength profiles in a 5mm thick butt weld of type 316austenitic stainless steel (Example (7.5)).

Test temperature

Fig. 7.49. Effect of temperature on notch toughness (schematic).

steels unsuitable for many structural applications where the HAZ toughness is of particularconcern.

In contrast, medium- and high-strength fee metals have usually such high toughness thatbrittle fracture is not a problem at low temperatures, unless there is some special reactivechemical environment. Austenitic stainless steels and aluminium alloys fall within this cat-egory. When it comes to duplex stainless steel weldments, the situation is more complex.Here the HAZ toughness is determined by the austenite/ferrite balance in the weld, which, in

Peak

tem

perat

ure

9II

turn, depends on the steel chemical composition and the operational conditions applied.78"80As shown in Fig. 7.50, complete ferritisation is normally achieved in regions close to thefusion boundary during the initial heating leg of the thermal cycle. Provided that the coolingrate through the critical transformation temperature range is kept reasonably low, a significantproportion of the ferrite may retransform back to austenite on cooling (see Fig. 7.51). This willcontribute to a high HAZ toughness, even at subzero temperatures, as indicated by the CVNtransition curves in Fig. 7.52. In general, an austenite content of about 30 vol% is sufficient toavoid problems with the HAZ toughness in duplex stainless steel welds.

7.3.4 Solidification crackingWhen a high alloy steel solidifies, a variety of microstructures can developed, depending onthe steel chemical composition and the cooling conditions applied. Referring to the Fe-Ni-Crphase diagram in Fig. 7.50, compositions on the Ni-rich side of the peritectic/eutectic liquidussolidify as primary austenite, while those on the Cr-rich side solidify as primary delta ferrite.The former condition is known to promote solidification cracking in welds due to partitioningof low-melting-point segregates of sulphur and/or phosphorus to the ^Fe-IFe grain bounda-ries.76"78

Since the solubility of sulphur and phosphorus is higher in delta ferrite than in austenite(0.18 and 2.8 wt%, respectively vs 0.05 and 0.25 wt% in austenite), the cracking susceptibilitycan normally be reduced if delta ferrite is present in sufficient amounts.76 At the same time,the interfacial Vetting' conditions are significantly improved owing to the formation of a low-

Weld metalPartly melted zone

Delta ferrite grain growth zone

Partly transformed zone

Unaffected base metal

Liquid (L)

Chromium content, wt%

Fig. 7.50. Schematic diagram defining different thermal regions within the HAZ of a single pass duplexstainless steel weld.

Temp

eratur

e, 0 C

30-45sat1300C Fe-26.0 wt% Cr - 6.9 wt% Ni

A4- temperature: 1260 0C

% 5 - ferrite at RT

Time, sFig. 7.51. CCT-diagram for a duplex stainless steel. Data from Mundt and Hoffmeister.79

energy interface between yFe and o>e, which prevents spreading of the liquid along the grainboundaries.82 For these reasons, a minimum weld metal delta ferrite content of about 5 to 10vol% is usually specified for austenitic stainless steels.

The quantitative relationship between the delta ferrite content and the weld metal chemicalcomposition in austenitic stainless steels has been determined first by Schaeffler83 and later byDelong et a/.84'85 The constitution diagram of Delong is shown in Fig. 7.53. Here the alloyingelements are grouped into ferrite formers (i.e. Cr, Mo, Si, and Nb) and austenite formers (i.e.Ni, C, N, and Mn) to determine the corresponding chromium and nickel equivalents for agiven alloy. The Delong diagram differs from the Schaeffler diagram in that the importantnitrogen contribution also is included in the former, thus allowing a more accurate predictionof the weld metal delta ferrite content.

Example (7.6)Consider plasma arc butt welding of a 5mm thick plate of austenitic stainless steel under con-ditions similar to those employed in Example (7.5). Data for the base metal (BM) and thefiller wire (FW) chemical compositions are given in Table 7.4. Use this information togetherwith the Delong diagram in Fig. 7.53 to determine which of the two filler wires (I or II) thatprovides the highest resistance against weld metal solidification cracking under the prevailingcircumstances. In the present example we shall assume that the mixing ratio B/(B+D) is 0.57(the mixing ratio is defined in Section 1.10.8, Chapter 1).

SolutionIn the absence of oxidation losses, the weld metal chemical composition is given by the 'ruleof mixtures':

Abso

rbed

en

ergy

, J

3-D heat flowconditions

Peak temperature: 13000C

Test temperature, 0CFig. 7.52. CVN transition curves for a duplex stainless steel after weld thermal simulation. Data fromVidem.81

Table 7.4 Chemical composition of base plate and filler wires used in Example (7.6).

Material1 wt% C wt% Cr wt% Ni wt% Mo wt% Mn wt% Si wt% N

Baseplate 0.03 17.0 12.0 2.0 2.0 0.9

Filler wire I 0.03 18.0 11.5 2.7 0.8 0.9 0.07

Filler wire II 0.03 18.0 15.5 - 0.8 0.9 0.07

Max. 0.015 wt% S and 0.03 wt% P

This leads to the weld metal chemical compositions shown in Table 7.5. We can nowcalculate the Cr- and Ni-equivalents for both welds:

Nick

el eq

uivale

nt (%

Ni+30

x%C+

30x%

N+0.5

x%Mn

)

Austenite

Austenite+ferriteAustenite+martensite

Schaeffler A+M line

Chromium equivalent (%Cr+%Mo+1.5x%Si+0.5x%Nb)Fig. 7.53. The Delong diagram85 showing the relationship between delta ferrite content and weld metalchemical composition for stainless steels.

Table 7.5 Computed weld metal chemical compositions (Example (7.6)).

Weld No. wt% C wt% Cr wt% Ni wt% Mo wt% Mn wt% Si wt% N

I 0.03 17.4 11.8 2.3 1.5 0.9 0.03

II 0.03 17.4 13.5 1.1 1.5 0.9 0.03

Readings from Fig. 7.53 then give:

Weld I : About 93 vol% austenite and 7 vol% delta ferriteWeld IP. 100 % austenite

From this we see that wire I provides the highest resistance against weld metal solidifica-tion cracking for the combination of steel and operational conditions considered above.

It should be emphasised that the Delong diagram gives no information about the real solidi-fication microstructure, since it is based on measurements of retained delta ferrite at roomtemperature. Also, the important effect of cooling rate on the weld metal transformation be-haviour is neglected in the present analysis. Consequently, the use of such empirical diagramsfor selection of steel and welding consumables is a keenly debated question in the scientificliterature.

7.4 Aluminium Weldments

Aluminium alloys are to an increasing extent used as structural components in welded assem-blies because of their high strength, low density, and good resistance against general corro-sion. In certain cases the application of aluminium is restricted by a low HAZ strength leveldue to softening reactions occurring during welding. In other cases the cracking resistance orthe fatigue strength becomes the limiting factor, depending on the design criterion. Table 7.6summarises typical problems associated with welding of aluminium and its generic alloys. Inthe following, we shall focus on the structural and mechanical response of age-hardenablealuminium alloys to the heat of welding processes, with particular emphasis on Al-Mg-Sialloys and Al-SiC metal matrix composites.

7.4.1 Solidification crackingConventional filler metals for aluminium welding are based on the binary Al-Si or Al-Mgsystems. It follows from Fig. 7.54 that maximum cracking sensitivity occurs somewhat be-tween pure Al and 2 to 4 wt% Si and Mg in aluminium, respectively. This behaviour arisesfrom the competitive influence of two different processes.76 In pure aluminium, on the onehand, the solidification cracking problem does not exists, since there is no low-melting-pointeutectic present at the grain boundaries. In solute-rich aluminium alloys, on the other hand, theeutectic liquid is so abundant that it backfills and 'heals' the incipient cracks, which results ina low cracking sensitivity. Somewhere in between the two composition limits, the amount ofeutectic liquid may be just large enough to form continuous films at the grain boundaries. This,in combination with high shrinkage or thermal contraction stresses, may lead to solidificationcracking, as indicated in Fig. 7.55.

Normally, aluminium-based filler metals contain between 4 to 5 wt% Si or Mg. This solutecontent is usually high enough to prevent solidification cracking during welding. Problems

Table 7.6 Characteristics of aluminium weldments.

Material

Al-Mg alloys(5 XXX-series)

Al-Mg-Si alloys(6xxx-series)

Al-Zn-Mg alloys(7xxx-series)

Al-SiC metalmatrix composites(10-15% vol% SiC)

Chemical Composition

Major Elements

Mg: 1-5 wt%

Mg: 0.5-1.3 wt%Si: 0.4-1.4 wt%

Mg: 0.5^4 wt%Zn: 2-8 wt%

Si: -7.5 wt%Mg: -0.5 wt%

Minor Elements

Si, Fe, Mn

Mn, Cr, Fe

Cu, Cr, Mn, Fe

Fe, Ti

Welding Problems

Solidification cracking Reduced HAZ strength level

(work-hardened materials) Solidification cracking Hot cracking Reduced HAZ strength level

Solidification cracking Hot cracking Low HAZ stress corrosion cracking

resistance

Conventional fusion weldingis not recommended

High-quality joints can beproduced by means of frictionwelding or diffusion bonding

Relat

ive cr

ack

sens

itivity

AI-Si system

Al-Mg system

Composition of weld, wt%Fig. 7.54. Solidification cracking sensitivity of binary Al-Si and Al-Mg alloys. Data from Dudas andCollins.86

may arise, however, if the weld metal becomes heavily depleted with respect to Si or Mg dueto dilution with the parent metal. Welds of a high B/(B+D) ratio fall within this category. Insuch cases it may be necessary to use over-alloyed filler wires to obtain crack-free welds.

Example (7.7)Consider plasma arc butt welding of a 10mm thick aluminium plate of type AA 6082-T6(containing 0.7 wt% Mg and 0.9 wt% Si) under the following conditions:

qo = 2OkW, v = 10mm s"1, T0 = 200C

Experience shows that the bead reinforcement amounts to 10% of the groove cross section(details of the groove geometry are given in Fig. 7.56). Two different filler wires are available,wire I with 5 wt% Si and wire II with 5 wt% Mg. Use this information along with the diagramsin Fig. 7.54 to determine which of the two filler wires (I or II) that provides the highest resist-ance against weld metal solidification cracking under the prevailing circumstances. In thesecalculations we shall assume that the temperature field around the heat source is given by thesimplified Rykalin thin plate solution (equation (1-100) in Chapter 1). Thermal data for Al-Mg-Si alloys are contained in Table 1.1.

SolutionFirst we need to estimate the mixing ratio BI(B+D). The total width of the fusion zone (2ym)can be obtained from equation (1-100):

(a)

(b)

Fig. 7.55. Examples of solidification cracking in aluminium weldments (Varestraint test coupons);(a) Al-I wt% Si, (b) Al-I wt% Mg. After Cross.87

Fig. 7.56. Groove geometry for a single pass Al-Mg-Si butt weld (Example (7.7)).

This gives:

and

from which

In the absence of oxidation losses, the weld metal chemical composition is given by the'rule of mixtures':

If wire I is used, the weld metal Mg and Si concentrations become:

Similarly, in the case of wire II, we get:

A comparison with Fig. 7.54 shows that wire I provides the highest safety against weldmetal solidification cracking under the prevailing circumstances, while wire II is unaccept-able.

The lower cracking resistance of the Al-Mg wire compared with the Al-Si wire at highdilution ratios arises from the pertinent difference in the fraction of eutectic liquid which formsduring weld metal solidification. For pure binary alloys the eutectic fraction feuL is given byequation (3-46) in Chapter 3:

Temp

eratur

e

where CeuL is the eutectic concentration, and ko is the equilibrium partitioning coefficient (givenby the binary Al-Si and Al-Mg phase diagrams).

If the contribution from the accompanying alloying element is neglected, the values of feuLbecome:

Weld I (1.80 wt% Si):

Weld II (1.65 wt% Mg):

From this we see that the fraction of eutectic liquid in weld I is so abundant that it backfillsand 'heals' all incipient cracks, while feut in weld II is just large enough to form continuousfilms at the columnar grain boundaries which, in turn, promotes solidification cracking.

7.4.2 Hot cracking

Hot cracking is a phenomenon occurring within the high peak temperature regions of the HAZduring welding of aluminium alloys. In spite of the difference in location, hot cracking likesolidification cracking is intergranular and arises from the combined action of grain boundaryliquation and stresses induced by solidification shrinkage and thermal contraction. The forma-tion of liquid phases within a weld HAZ can readily be explained with the help of a simplebinary phase diagram of the type shown in Fig. 7.57. If the composition of the binary alloy is

Liquid (L)

Composition

Fig. 7.57. Schematic binary phase diagram defining the equilibrium conditions for partial melting duringreheating.

Tem

pera

ture,

0 C

higher than Cmax^ and the alloy is heated to a temperature above TeuU partial melting willoccur. Since the eutectic phase is usually located at the grain boundaries, these sites becomeimmediately covered with liquid films if the wetting conditions are favourable.

7.4.2.1 Constitutional liquation in binary Al-Si alloysExperience shows that even binary alloys with a nominal composition lower than Ceut canundergo incipient melting at the eutectic temperature. This will obviously be the case if thealloy contains segregations so that the concentration locally exceeds the critical composition.Another possibility is melting due to 'constitutional liquation', a mechanism originally pro-posed by Pepe and Savage8889 for the formation of hot cracking in 18-Ni maraging steelweldments at temperatures well below the bulk solidus of the alloy. The same theory has laterbeen applied by Reiso et al.90~93 and Lohne and Ryum94 to explain incipient melting duringhomogenisation heat treatment of aluminium alloys. In the following, the treatment of Lohneand Ryum94 is adopted to illustrate the principles of constitutional liquation in binary Al-Sialloys under conditions applicable to welding. The theory is later extended to ternary Al-Mg-Si alloys.

Figure 7.58 shows a section of the Al-rich corner of the binary Al-Si phase diagram. Con-sider next a pure Al-1.2 wt% Si alloy that first is brought to equilibrium at a lower temperatureto form 10|um large Si particles (Fig. 7.59(a)) and then is rapidly heated to a higher tempera-ture slightly above Teut. If local equilibrium is maintained at the particle/matrix interfaceduring heating, the composition is given by the Al-Si solvus boundary in Fig. 7.58. When Ceutis reached incipient melting will take place by formation of eutectic liquid (according to thereaction Al + Si liquid), as shown in Fig. 7.59(b). The process will continue until the wholeSi particle has dissolved (Fig. 7.59(c)). Depending on the density/location of the Si particlesand the interfacial wetting conditions, grain boundary liquation may eventually occur, as indi-cated in Fig. 7.59(d).

Liquid

Liquid+solid

ss

Al+Si

Silicon content, wt%

Fig. 7.58. Section of the Al-rich comer of the binary Al-Si phase diagram. Data from Ref. 95.

(a) (b)

(C) (cO

Fig. 7.59. Optical micrographs showing the microstructural evolution during homogenisation heat treat-ment of binary Al-Si alloys; (a) Isolated Si particle embedded in a matrix of Al, (b) Partially melted Siparticle surrounded by Al-Si eutectic (25s at 582C), (c) Globular Al-Si eutectic structure formed aftercomplete dissolution of the Si particle (60s at 582C), (d) Spreading of eutectic liquid along a grainboundary (25s at 582C). Courtesy of O. Lohne, Sintef - Division of Metallurgy, 7034 Trondheim,Norway.

7.4.2.2 Constitutional liquation in ternary Al-Mg-Si alloysConstitutional liquation in ternary Al-Mg-Si alloys has been investigated by Reiso et al.96Figure 7.60 shows a section of the Al-rich corner of the quasi-binary Al-Mg2Si phase diagram.When an alloy of composition C < Ceut is heated up rapidly, as in the case of welding, theprimary Mg2Si particles do not have enough time to dissolve. When the quasi-binary eutectictemperature is reached at about 593C, incipient melting will occur, according to the reac-tion:96

liquid (7-22)

Further heating to a higher temperature (T > 593C) provides additional time for dissolu-tion of Mg2Si and formation of more liquid of variable composition. In practice, reaction (7-

Temp

eratur

e, 0C

Mg2Si concentration, wt%

Fig. 7.60. Quasi-binary section of the Al-rich corner of the ternary Al-Mg-Si phase diagram. Data fromReiso et al.96 and Phillips.97

22) may also proceed below the quasi-binary eutectic temperature Teut, since the diffusion rateof Mg is higher than that of Si.96 By considering the kinetics it can be shown that the interfaceconcentration will gradually move towards the silicon side of the quasi-binary line in the phasediagram as the Mg2Si particles dissolve, thereby reducing the temperature at which liquationoccurs (from 593 down to about 5800C). Moreover, in alloys containing excess amounts ofsilicon, two other side reactions can take place:96

(7-23)

(7-24)and

liquid

liquid

The former is analogous to the melting reaction in binary Al-Si alloys (Fig. 7.58), while thelatter corresponds to the eutectic reaction in the ternary Al-Mg-Si system (Teut varies from555 to 559C, depending on the source). Since this eutectic temperature represents the lowesttemperature at which a melt can exist within the system, it means that local melting of secondphase particles cannot take place below, say, 555 to 559C during welding of Al-Mg-Si al-loys.

7.4.2.3 Factors affecting the hot cracking susceptibilityIn addition to the thermodynamic and kinetic effects mentioned above, there are several otherfactors, some interrelated, which play an important part in the formation of hot cracks inAl-Mg-Si weldments. These are:76

(i) The number density and size distribution of Mg2Si and Si particles in the basemetal.

(ii) The total grain boundary area per unit volume available for absorption of eutecticliquid (determined by the HAZ grain size).

(iii) The interfacial wetting conditions.

(iv) The local tensile stress level in the partially melted region.

It follows that the hot cracking susceptibility may be significantly altered by a change inone of these parameters, but, in practice, there is very little that can be done to prevent grainboundary liquation if the base material already contains second phase particles. In fact, theonly useful way of eliminating the cracking problem is to reduce the tensile stress in the par-tially melted HAZ through proper selection of welding consumables. This is because thecomposition of the weld metal can be adjusted so that solidification is completed first in thepartially melted region and then in the fusion zone, thus avoiding hot cracking in the former.

As an illustration of principles, the diagrams of Gittos and Scott98 (reproduced in Fig. 7.61)will be considered. These diagrams show the variation of the weld metal solidus temperaturewith base metal dilution for two commercial filler wires (wire I: Al-5 wt% Si, and wire II: Al-5 wt% Mg). It is evident that the risk of hot cracking is highest when wire II is used, particu-larly at high B/(B+D) ratios, since solidification occurs first in the weld metal and then in thepartially melted region. In contrast, wire I provides a good HAZ cracking resistance over thewhole composition range because of the resulting lower solidus temperature. In the latter casethe weld metal solidification and thermal contraction stresses are imposed on the HAZ at astage where liquid no longer exists at the grain boundaries.

Example (7.8)Consider plasma arc butt welding of a 10mm thick aluminium plate of AA6082-T6 (contain-ing 0.7 wt% Mg and 0.9 wt% Si) under conditions similar to those employed in Example (7.7).Use the diagrams in Fig. 7.61 to determine which of the two filler wires (I or II) that providesthe highest resistance against hot cracking in the partially melted region under the prevailingcircumstances.

SolutionIn the previous example the base metal dilution ratio, BI(B+D), was found to be 0.78. Undersuch conditions wire I provides the highest resistance against hot cracking, since solidificationfirst occurs in the partially melted region and then in the weld metal. This explains why Al-Sibased filler wires are usually recommended for single pass butt welding of Al-Mg-Si extru-sions if the strength level is not of particular concern.

Solid

us te

mpe

ratu

re, 0 C

Solid

us te

mpe

ratu

re, 0 C

(a)

liquid

B/(B+D), %(b)

liquid

B/(B+D), %Fig. 7.61. Variation of weld metal solidus temperature with dilution for 6082 aluminium alloys (Al-0.7wt% Mg-0.9 wt% Si); (a) Wire I (Al-5 wt% Si), (b) Wire II (Al-5 wt% Mg). The melting temperaturefor different base metal constituent phases are indicated by the horizontal broken lines in the diagrams.Data from Gittos and Scott98 and Reiso et al.96

Tem

pera

ture

Har

dnes

s

Fully

re

verte

d re

gion

Una

ffect

ed ba

se m

eta

l

(a) Peak temperature

Partly reverted region:

HAZ-

Distance from fusion line

(b)

C-curve forprecipitationof (3'(Mg2 Si)

Weld thermal cycle

Time

Fig. 7.62. Schematic diagrams showing the sequence of reactions occurring in the HAZ of 6082-T6aluminium welds; (a) Hardness distribution following p"(Mg2Si) dissolution; (b) Precipitation of P'(Mg2Si)at dispersoids during the weld cooling cycle.

Har

dnes

s

7A3 HAZ microstructure and strength evolution during fusion weldingThe age-hardenable Al-Mg-Si alloys have been widely studied, to the extent that most of theunderlying physical processes are well established. They offer tensile strength values higherthan 350 MPa in the artificially aged (T6) condition owing to the presence of very fine, needle-shaped 3"(Mg2Si) precipitates along directions in the aluminium matrix." AlthoughAl-Mg-Si alloys are readily weldable, they suffer from severe softening in the heat affectedzone (HAZ) because of reversion (dissolution) of the (3"(Mg2Si) precipitates during the weldthermal cycle.76100"103 This type of mechanical impairment represents a major problem in en-gineering design, since it reduces the load-bearing capacity of the joint.104

7.4.3.1 Effects of reheating on weld propertiesMicrostructural changes in the HAZ of Al-Mg-Si alloys have been examined by several in-vestigators,100"103 and the main results are summarised in Fig. 7.62.

It is evident from Fig. 7.62(a) that reversion of [3"(Mg2Si) precipitates will occur to anincreasing extent in the peak temperature range from 250 to 5000C. This is associated with acontinuous decrease in the HAZ hardness until the dissolution process is completed. Duringcooling of the weld, some solute recombines to form coarse, metastable (3'(Mg2Si) precipitateswhich do not contribute to strengthening (Fig. 7.62(b)). However, close to the fusion bound-ary a large fraction of alloying elements will remain in solid solution at the end of the thermalcycle, thereby giving conditions for extensive age-hardening at room temperature over a pe-riod of 5 to 7 days (Fig. 7.62(c)).

In general, enhanced HAZ strength recovery can be achieved by the use of artificial ageingin the temperature range from 150 to 1800C,76 but this possibility will not be considered here.

(C) Peak temperature

Contribution;from natural:ageing \

Resulting hardnessprofile

Distance from fusion line

Fig. 7.62. Schematic diagrams showing the sequence of reactions occurring in the HAZ of 6082-T6aluminium welds (continued); (c) Hardness distribution after prolonged room temperature ageing.

7.4.3.2 Strengthen ing mechan isms in A l-Mg-Si alloysDue to the lack of experimental evidence of coherency strains around (3"(Mg2Si) precipitatesin Al-Mg-Si alloys," it has been suggested that the increased resistance to dislocation motionaccompanying the presence of these structures arises from the high energy required to breakMg-Si bonds in the particles as dislocations shear through them. Assuming that this strength-ening effect is associated with order hardening, the net precipitation strength increment, Aop,can be calculated from the equation originally derived by Kelly and Nicholson:105

(7-25)

where 7/ is the internal interface (or antiphase) boundary energy, b is the Burgers vector,/isthe particle volume fraction, and c$ is a kinetic constant.

By introducing the relative particle volume fraction, f/fo, we obtain:

(7-26)

where fo is the initial volume fraction of (3"(Mg2Si) precipitates in the alloy, and C4 is a newkinetic constant (equal to c3/o).

It is evident from equation (7-26) that Aop = Aap (max) = c4 when/7/o = 1. Hence, thisequation can be rewritten as:

(7-27)

Here amin denotes the intrinsic matrix strength after complete particle dissolution, whiledmax is the original base metal strength in the artificially aged (T6) condition.

Provided that a linear relationship exists between yield strength and hardness, equation (7-27) can be rewritten as:102

(7-28)

Equation (7-28) provides a basis for assessing the reaction kinetics through simple hardnessmeasurements. Typical values for amax, HVmax, omin, and HVmin are given in Table 7.7.

7.4.3.3 Constitutive equationsThe kinetics of (3"(Mg2Si) dissolution during reheating of 6082-T6 aluminium alloys havebeen considered in Section 4.4.2 (Chapter 4).Table 7.7 Properties of some Al-Mg-Si and Al-SiC weldments.

HVmax HVmin o m a x

Reversion modelWhen the 3"(Mg2Si) particles dissolve, the volume fraction falls from its initial value f0, ac-cording to equation (4-36). A combination of equations (4-36) and (7-28) leads to the follow-ing expression for the dimensionless strength parameter within the partly reverted region ofthe HAZ:

(7-29)

where t\ is the maximum hold time required for complete particle dissolution at a given tem-perature (defined by equation (4-31) in Chapter 4), and n is a time exponent (< 0.5). Thevariation of n with/ / /o is shown in Fig. 4.20.

Natural ageing modelIn general, the fraction of hardening (3"(Mg2Si) precipitates which forms during natural ageing(flfo) depends on the amount of remnant solute present in the matrix material after cooling ofthe weld. By considering the kinetics of the C-curve for precipitation of essentially non-hardening 3'(Mg2Si) particles at dispersoids during the weld cooling cycle, Myhr and Grong102arrived at the following relationship for flfo'

(7-30)

(7-31)where

Here O is a material constant (defined in Table 7.7), and t\ is the critical time required toprecipitate a certain fraction of $'(X = Xc) at an arbitrary temperature (T). The variation oft*2 with temperature is given by equation (6-55) in Chapter 6.

The net precipitation increment following natural ageing (a2) can therefore be written as:

(7-32)

Coupling of modelsBased on equations (7-29) and (7-32) it is possible to calculate the HAZ strength distributionafter welding and subsequent natural ageing when the weld thermal programme is known.Figure 7.63 shows a sketch of the superimposed hardness profiles, as evaluated from theseequations. Since the resulting strength level in the partly reverted region depends on the inter-play between two competing processes (i.e. dissolution and reprecipitation), it is convenient todefine the 'boundary' between the two models on the basis of the intersection point in Fig.7.63 where a\ = a2 , i.e.:

(7-33)

(7-34)and

Hardn

ess

Fully

re

verte

dre

gion

Peak temperature

Partly revertedregion

Reversion model

Intersectionpoint

, Natural ageing model

Distance from fusion lineFig. 7.63. Coupling of reversion and natural ageing models.

It follows that this locus also defines the minimum HAZ strength level, which is an impor-tant parameter in engineering design.

7.4.3.4 Predictions of HAZ hardness and strength distributionThe predictions are based on computer programmes which utilise the medium thick plate heatflow solution described in Section 1.10.4 (Chapter 1) and the kinetic models outlined above tocalculate the HAZ hardness distribution for specific welding conditions (details are given inRef. 106).

Accuracy of predictionsExamples of measured and predicted HAZ hardness profiles are shown in Figs. 7.64 and 7.65.

When stringer bead welding is carried out on a plate of medium thickness, the hardnessdistribution in the transverse y direction will vary with distance from the plate surface due to acontinuous change in the heat flow conditions. A comparison between observed and predictedhardness profiles in Fig. 7.64 shows that such effects are readily accounted for in the presentmodel. In contrast, a full penetration butt weld will always reveal a similar HAZ hardnessdistribution in the transverse section of the weld, as shown in Fig. 7.65. This situation arisesfrom the lack of a temperature gradient in the through-thickness z direction of the plate. Moreo-ver, it is evident from Figs. 7.64 and 7.65 that the final dimensions of the HAZ are stronglyinfluenced by variations in welding parameters and operational conditions. Hence, it is diffi-cult to justify the use of a constant safety factor for the width of the HAZ as recommended incurrent design rules for welded Al-Mg-Si alloys.104

Aptness of modelsBased on the kinetic models described in the proceding sections, it is possible to constructtwo-dimensional (2-D) maps which show characteristic hardness and peak temperature con-tours in the HAZ of 6082-T6 aluminium weldments. Examples of such diagrams are given in

Hardn

ess,

VPN

Peak

tem

perat

ure,

0 CPe

ak tem

perat

ure.

0 C

Hardn

ess,

VPN

(a)

(b)

HV (predicted)

Aym, mm

HV (predicted)

ym,mm

Fig. 7.64. Comparison between measured and predicted HAZ hardness profiles in a stringer bead GMAweld; (a) Upper plate surface, (b) Lower plate surface. The peak temperature distribution is indicated bythe broken lines in the graphs. (Operational conditions: qo = 9.1 kW, v = 5.1 mm s"1, d = 15mm). Datafrom Myhr and Grong.102

Fig. 7.66. Included is also a 3-D plot of the HAZ hardness distribution in the transversesection of the weld.

The results in Fig. 7.66 reveal a direct relationship between the HAZ isothermal contourson the one hand and the resulting HAZ hardness/strength distribution on the other. In thisparticular example the soft zone closely follows the contour of the 4000C isotherm. This, inturn, implies that the minimum HAZ strength level is fairly constant and virtually independentof choice of welding parameters (i.e. close to 60 VPN for single pass welds).

Hard

ness

, VP

N

Peak

te

mpe

ratur

e. 0 C

HV (predicted)

A y m , mmFig. 7.65. Comparison between measured and predicted HAZ hardness profiles in a single pass plasmaarc butt weld. The peak temperature distribution is indicated by the broken line in the graph. (Opera-tional conditions: qo = 14.0 kW, v = 5.8 mm s"1, d = 13mm). Data from Myhr and Grong.102

In practice, the HAZ hardness can be converted into an equivalent yield or ultimate tensilestrength through the following regression formulae:102

(7-35)

(7-36)and

From equation (7-35) we see that a minimum HAZ hardness of about 60 VPN correspondsto a strength reduction factor of:

This value is in good agreement with the recommended strength reduction factor of 0.49incorporated in many welding specifications and standards.104'107

Process diagrams for single pass butt weldsBecause of the complex temperature-time pattern in aluminium welding, it is not possible tocondense general information about the HAZ strength distribution into 2-D process diagrams.An exception is single pass butt welding of plates, where the medium thick plate solutiondescribed in Section 1.10.4 (Chapter 1) can largely be simplified if the net arc power is keptsufficiently high compared with the plate thickness (e.g. qold > 0.5 kW mm"1). Under suchconditions the mode of heat flow becomes essentially one-dimensional, and the temperaturedistribution is determined by the ratio qo/vd, kJ mm"2 (see equation (1-100) in Chapter 1).

z, m

m

(a)

y, mm

Hardness code

Fusion zone

Fig. 7.66. Computed HAZ hardness and peak temperature contours in the transverse section of a stringerbead GMA weld; (a) 2-D graphical representation, (b) 3-D graphical representation. (Operational condi-tions as in Fig. 7.64). Data from Myhr and Grong.102

Figure 7.67 shows plots of the HAZ hardness/strength profiles for different values of qo Ivd.It follows that a narrow width of the HAZ requires the use of a low energy input per mm2 of theweld. In practice, this can be achieved by the choice of an efficient welding process (e.g.electron beam or laser welding) which allows deposition of a full penetration butt weld with-out employing a groove preparation (i.e. eliminates the need for filler metals).

Multipass weldingThe present process model can also be extended to multipass welding if it is assumed thatreversion of indigenous (3'(Mg2Si) precipitates occurs instantaneously on reheating above thephase boundary solvus temperature (here taken equal to 5200C).

(b)

Hardn

ess,

VPN

Scale:10 mm

V 2 [ M P a ] Rm [MPa]

Distance from fusion lineFig. 7.67. Process diagram for single pass 6082-T6 butt welds. Data from Myhr and Grong.102

Figure 7.68 shows 2-D and 3-D plots of computed HAZ hardness and peak temperaturecontours for a simulated two-pass butt weld. This model system consists of two imaginarystringer beads which are placed symmetrically on each side of a 15mm thick plate. Compara-ble hardness data for a single pass weld of same thickness are contained in Fig. 7.66. It isevident from Figs. 7.66 and 7.68 that deposition of a second pass will neither increase thewidth nor reduce the strength of the HAZ to any great extent. The explanation lies in the factthat the temperature field around each heat source tends to overlap with increasing distancefrom the fusion boundary due to symmetry effects (see Fig. 7.68(a)), which prevents excessivestrength loss during reheating of the weld. Similar observations have also been made fromactual testing of multipass Al-Mg-Si weldments.108

Example (7.9)Consider plasma arc butt welding of a 10mm thick aluminium plate of type AA 6082-T6 underconditions similar to those employed in Example (7.7). Estimate on the basis of the processdiagram in Fig. 7.67 and the simplified Rykalin thin plate solution (equation (1-100) in Chap-ter 1) both the minimum HAZ strength level, the total width of the reduced strength zone afterwelding, as well as the lower temperature limit for dissolution of the (3"(Mg2Si) precipitatesduring the weld thermal cycle. Thermal data for Al-Mg-Si alloys are given in Table 1.1(Chapter 1).

SolutionFirst we need to calculate the net heat input per mm2 (q

ol vd). In the present example, wehave:

z, m

m

(a)

Hardness code

Fusion zone

y, mm

(b)Fig. 7.68. Computed HAZ hardness and peak temperature contours for a simulated two-pass butt weld(the second pass is deposited immediately after cooling of the first pass); (a) 2-D graphical representa-tion, (b) 3-D graphical representation. (Operationalconditions: qo = 9A kW, v = 5.1 mms~ l,d- 15mm).Data from Myhr and Grong.102

Readings from Fig. 7.67 then give:

and

Similarly, by considering the extension of the HAZ and the corresponding length of thescale bar in Fig. 7.67, the total width of the reduced strength zone becomes:

The relationship between peak temperature T- Tp and distance y = ym from the heat sourcecan now be obtained by differentiating equation (1-100) with respect to time. After some ma-nipulation, we obtain:

By substituting qo/vd = 200 J ram"2, pc = 0.0027 J mm"3 0C"1, and Tm = 652C into theabove equation, the following temperature for incipient dissolution of the p"(Mg2Si) precipi-tates is obtained:

It is obvious from the above calculations that the degree of HAZ softening occurring duringwelding is substantial under the prevailing circumstances. This explains why, for instance,high heat input deposition is usually not recommended for Al-Mg-Si alloys.

7.4.4 HAZ microstructure and strength evolution during friction weldingFriction welding is a solid state joining process that involves both heating and plastic deforma-tion of the parent material under extreme thermal and strain rate conditions.109"113 During thewelding operation the material is deformed at high temperatures, initially at low strain rates,but due to the axial displacement of the specimen, the material is subsequently brought into thefully plasticised region where the strain rate may exceed 103 s"1 (e.g. see Fig. (6.79) in Chapter6). Although the resulting microstructural changes are similar to those observed in many hotworking processes, it is obvious that transient effects must play a more dominant role in fric-tion welding because of the rapid temperature and strain rate fluctuations. This, in turn, in-creases the complexity of the analysis.

7.4.4.1 Heat generation in friction weldingA major problem in modelling of heat flow phenomena in friction welding is to obtain anaccurate description of the energy input at the faying interface. This is because the frictioncoefficient |LL is changing continuously during the welding cycle from |i > 1 at the dry slidingstart, towards zero when the temperature for asperity melting is reached at the interface.113However, for the ideal case considered in Fig. 7.69, the torque required to rotate two circularshafts relative to one other under the action of an axial load pressure is given by:114115

(7-37)

where M is the interfacial torque, R* is the surface radius, and P(r) is the pressure distributionacross the interface (here assumed constant and equal to P).

Pressure Pressure

Fig. 7.69. Schematic arrangement of friction welding of a solid rod.

If all the shearing work at the interface is assumed to be converted into frictional heat, theaverage heat input per unit area and time becomes:114

where qo is the net power (in W), P is the friction pressure (in N mm"2), A is the cross section(in mm2), and wmax. is the maximum surface velocity at the outer edge (in m s"1).

Equation (7-38) provides a basis for estimating the heat generation at the interface duringcontinuous drive friction welding in the absence of asperity melting.

7.4.4.2 Response of Al-Mg-Si alloys and Al-SiC MMCs to friction weldingIn friction welding of aluminium alloys and Al-SiC metal matrix composites, the deformationis maintained throughout the steady state period.113 The microstructural changes that takeplace at this stage of the process can be classified as dynamic changes, since they occur underthe action of variable plastic straining. Similarly, those taking place after the forge operation(when the weld cools) are referred to as static changes due to the lack of external plasticdeformation. The former includes the generation of recovered subgrains and reversion ofconstituent precipitates in the area of high temperature and high deformation, whereas thelatter involves precipitation reactions, e.g. formation of hardening [3"(Mg2Si) particles in theheat affected zone following natural ageing.109"113

In order to assemble an adequate model for the microstructural evolution during frictionwelding, it is necessary to divide the HAZ into different reaction zones. Referring to Fig. 7.70,the following three main regions are of specific interest in this context:

(i) The fully plasticised region, ZpL, where the material is able to accommodate theplastic strain by dynamic recovery (or recrystallisation) of the microstructure.

(7-38)

Fig. 7.70. Schematic diagram showing the three main reaction zones within a friction welded component(Zpi/. fully plasticised region, Zpd\ partly deformed region, Zud: undeformed region).

(ii) The partly deformed region, Zpd., where the degree of plastic deformation is ac-commodated by an increase in the dislocation density of the matrix grains. In thisregion the temperature is sufficiently high to facilitate dissolution of the base metalhardening precipitates.

(iii) The undeformed region, Z1^., characterised by partial reversion of the base metalprecipitates.

Aspects of HAZ subgrain evolution during continuous drive friction welding have beendescribed in Section 6.5.2 (Chapter 6). In the following, the structural and mechanical re-sponse of T6 heat treated Al-Mg-Si alloys and Al-SiC metal matrix composites to the im-posed heating and plastic deformation will be considered more in detail.

7.4.4.3 Constitutive equationsSince the HAZ strength level in both types of materials is mainly controlled by dissolution andprecipitation reactions occurring within the aluminium matrix during the weld thermal cy-cle,109"113 we can use equations (7-29) and (7-32) with minor modifications to describe thereversion and the natural ageing characteristics of the weld components (see Table 7.7). Inaddition, it is necessary to allow for the plastic deformation introduced during friction weld-ing, because dislocations will be generated in the matrix material to accommodate the strain.113In general, work hardening of metals and alloys is a very complex problem which has not yetbeen properly solved. However, with the precision aimed at here it is sufficient to assume thatthis strength contribution is given by:

(7-39)

where 1F is a constant which is characteristic of the material under consideration (given inTable 7.7). The form of equation (7-39) indicates that the hardness is essentially constantwithin the plasticised regions of the weld HAZ.

7.4.4.4 Coupling of modelsBased on equations (7-29), (7-32), and (7-39) it is possible to calculate the HAZ strengthdistribution after friction welding and subsequent natural ageing when the weld thermal pro-gramme is known. Figure 7.71 (a) and (b) show schematic representations of the superim-

Contact section

Z

Hardn

ess

Hardn

essposed hardness profiles, as evaluated from these models. The resulting HAZ hardness distri-bution is indicated by the solid curve in the graphs for short and long duration thermal cycles,respectively. Since the present treatment oversimplifies the problem by only considering thestrongest contribution with no interaction, the justification of this assumption relies on a goodcorrelation between theory and experiments.

It follows from Fig. 7.71 that particle dissolution is the major softening mechanism in fric-tion welding of Al-Mg-Si alloys and Al-SiC metal matrix composites. At the same time asubstantial strength recovery will occur as a result of external plastic straining in combinationwith intrinsic precipitation of hardening (3"(Mg2Si) particles following prolonged room tem-perature ageing. Depending on the operational conditions applied, this may give rise to differ-ences in the shape of the HAZ hardness profiles, as shown in Fig. 7.71.

(a)Reversion model

Work hardening model

Natural ageing model

Unaffected base materialAxial distance, Z

(b)Reversion model

Work hardening model

Natural ageing model Unaffectedbase material

Axial distance, Z

Fig. 7.71. Schematic representation of the HAZ hardness distribution after friction welding and subse-quent natural ageing; (a) Short duration thermal cycle, (b) Long duration thermal cycle. The parametersZp/., Zp(L and Zud% are defined in Fig. 7.70.

Hard

ness

, VP

N7.4.4.5 Prediction of the HAZ hardness distributionThe predictions are based on computer programmes which utilise the heat and material flowmodels described in Ref.113 in combination with the constitutive equations given above tocalculate the HAZ hardness distribution for specific welding conditions.

Accuracy of predictionsExamples of measured and predicted hardness profiles are given in Figs. 7.72 and 7.73.

A closer inspection of the graphs reveals a good agreement between theory and experi-ments in all three cases. It is interesting to note that there is no clear distinction in the shape ofthe HAZ hardness profiles between friction welded Al-Mg-Si alloys and Al-SiC metal matrixcomposites when comparison is made on the basis of a similar temperature-time pattern (seeFig. 7.73). However, large diameter weld components will normally reveal a different hard-ness distribution, as shown in Fig. 7.72, because longer welding times will increase the totalheat input. Under such conditions the contribution from the plastic deformation becomesnegligible, which means that the resulting HAZ hardness profile will closely resemble thatobserved during conventional gas metal arc (GMA) and plasma arc welding of Al-Mg-Sialloys (see Figs. 7.64 and 7.65).

Process diagramsBased on the above process model, it is possible to construct a series of diagrams which sum-marise information about the effect of important welding variables in a systematic and illustra-tive manner. Examples of such diagrams for 6082-T6 aluminium alloys and T6 heat treatedAl-SiC metal matrix composites are given in Fig. 7.74(a) and (b), respectively.

PredictedMeasured

Unaffected basematerial

Axial distance, mm

Fig. 7.72.Comparison between measured and predicted HAZ hardness profiles in a

Hard

ness

, VP

NHa

rdne

ss, VP

N(a)

PredictedMeasured

Unaffected base material

Axial distance, mm

(b)

PredictedMeasured

Axial distance, mm

Fig. 7.73.Comparison between measured and predicted HAZ hardness profiles; (a) 4> 16mm Al-Mg-Siweld component. (Operational conditions: qJA - 25W mrrr2 and ts = 0.9s), (b) 16mm friction weldedAl-SiC metal matrix composite. (Operational conditions: qJA = 25W mirr2 and ts= 3.8s). Data fromMidling and Grong.113

Hard

ness

, VP

NHa

rdne

ss, VP

N(a)

Axial distance, mm

(b)

Axial distance, mmFig. 7.74. Process diagrams for friction welding; (a) 6082-T6 aluminium alloys. (Operational condi-tions: umax_ = 2.5m s"1 and [i = 0.5), (b) Al-SiC-T6 metal matrix composites. (Operational conditions:Umax. = 2.5m s"1 and \x = 0.5). Data from Midling and Grong.113

It is evident from these diagrams that the HAZ hardness distribution depends on the totalheat input applied during friction welding. Although the controlling parameters qo IA and ts(welding time), in practice, are kept within relatively narrow limits, it is obvious that a smallwidth of the HAZ requires the use of a high specific power (qol A) in combination with a shortduration heating cycle (ts < 2 s). This is also in agreement with general experience.109"112

Example (7.10)Consider continuous drive friction welding of a T6 heat treated Al-SiC metal matrix compos-ite under the following conditions:

Use the process diagram in Fig. 7.74(b) to estimate the minimum HAZ hardness level aswell as the total width of the strength reduced zone after welding. In this example we shallassume that the friction coefficient Ji is equal to 0.5.

SolutionFirst we need to calculate the frictional heat per unit area of the weld. From equation (7-38),we have:

A comparison with Fig. 7.74(b) shows that a specific power of 25 W mm"2 corresponds toa minimum HAZ hardness of about 90 VPN, i.e. a reduction of 45 VPN compared with thebase material. At the same time the total width of the reduced strength zone is seen to be12 mm.

It should be emphasised that the observed strength loss is not permanent, since the resultingHAZ strength level is mainly controlled by dissolution reactions taking place within the alu-minium matrix during the weld thermal cycle. Consequently, a full HAZ strength recoverycan be achieved by the use of an appropriate post weld heat treatment, as shown by the tensiletest data in Table 7.8.

Table 7.8 Mechanical properties of friction welded Al-SiC metal matrix composites. Data fromMidling and Grong109'113

HV Rp0 2 Rm eBMaterial [VPN] [MPa] [MPa] [%]

Al-SiC-T6 135 315 352 3.6(base material)Al-SiC-T6 90 207 268 3.0(as-welded condition)Al-SiC-T6 135 313 348 3.1(PWHT condition)!

Solution heat treated at 535C for 3 h followed by water-quenching and artifical ageing at 1600C for 10 h.

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Appendix 7.1Nomenclature

lattice parameter (nm)

cross section (mm2, m2)

start temperature of ferrite to austenitetransformation (0C, K)

end temperature of ferrite to austenitetransformation (0C, K)

Burgers vector (nm, m)

body-centred cubic

body-centred tetragonal

amount of fused parent metal (mm2)

kinetic constant in Hollomon-Jaffe pa-rameter

base metal

half crack length (jum, m)various kinetic and empirical constants

element concentration (ppm, ml per100 g, |il cm-3)eutectic concentration (wt%)

initial element concentration (ppm,ml per 100 g, |il cnr3)maximum solid solubility (wt%)

base metal or weld metal element con-centration (wt%, ppm, ml per 100 g,jil cm"3)Yurioka carbon equivalent

HW carbon equivalent

crack tip opening displacement

Charpy V-notch

plate thickness (mm)

particle diameter (Jim, m)

inclusion diameter (jJm, m)amount of deposited metal (mm2)

diffusivity (mm2 s"1, m2 s"1)

average grain size ((Xm, m)

lattice diffusion coefficient (mm2 s"1,m V 1 )

elongation (%)

gross heat input (kJ mm"1)

binding energy between hydrogen andtrap site (J mol"1, kJ mol"1)

Young's modulus of matrix (MPa, GNm-2)

Young's modulus of particle (MPa, GNm-2)

Gaussian error function

particle volume fraction

volume fraction of hardening precipi-tates in natural ageing model

initial particle volume fraction

eutectic fraction

flux-cored arc welding

face-centred cubic

filler wire

grain coarsened region

1 gas metal arc welding

grain refined region

gas tungsten arc welding

hydrogen content related to fused metal(ml per 100 g or g per ton)

Rockwell C hardness

Vickers hardness (kg mm"2 or VPN)

Vickers hardness in the artificially agedcondition (kg mm"2 or VPN)

matrix Vickers hardness in the absenceof hardening precipitates (kg mm"2 orVPN)

kinetic strength of thermal cyclewith respect to P' (Mg2Si) formation

intercritical region

locking parameter in Hall-Petch equa-tion (MPa |im1/2)equilibrium partitioning coefficient

density of trap site

half length of fusion zone in hydrogendiffusion model (mm, m)

longitudinal direction (notch perpen-dicular to plate rolling direction)

strain hardening exponent in Ludwikequation

interfacial torque (Nm)

martensite-austenite constituent

metal matrix composite

time exponent

dimensionless operating parameter inheat flow model

friction pressure (N mm"2, MPa)

Hollomon-Jaffe parameter

Dorn parameter (s)

pressure distribution across weld inter-face (MPa)

post weld heat treatment

net power (W)

apparent activation energy for soften-ing reaction (J mol"1, kJ mol"1)

radius vector (mm, m)

isothermal zone width referred to fu-sion boundary (mm)

universal gas constant (8.314 J mol"1K-i)

friction surface radius (mm, m)

implant rupture strength (MPa)

ultimate tensile strength (MPa)

0.2% proof stress (MPa)

scanning electron microscope

shielded metal arc welding

subcritical region

time (s)

maximum hold time required for com-plete particle dissolution at T (s)

critical hold time required to precipi-tate a certain amount of P'(Mg2Si) at T(S)

welding time (s)

cooling time from 600 to 4000C (s)

cooling time from 800 to 5000C (s)

cooling time from 1200 to 8000C (s)

equivalent isothermal hold time at333K (s)

temperature (0C, K)

chosen reference temperature(0C, K)

eutectic temperature (0C, K)

melting point (0C, K)ambient temperature (0C, K)peak temperature (0C, K)

artificially aged condition

transmission electron microscope

transverse direction (notch parallel withplate rolling direction)velocity (m s"1)

surface velocity (ms"1)

area under stress-strain curve (J m~3)

ultimate tensile strength

welding speed (mm s"1)

inclusion volume fraction

Vickers pyramid number

welding direction (mm, m)

transverse direction (mm, m)

y-coordinate at maximum width of iso-therm (mm)

isothermal zone width (mm, m)

through-thickness direction (mm, m)

partly deformed region in friction weld-ing model (mm)

fully plasticised region in friction weld-ing model (mm)

undeformed region in friction weldingmodel (mm)

stress concentration factor

true fracture strain

dimensionless plate thickness

material constants

volume heat capacity (J mm"3 0C"1)

strength (MPa)

yield stress (MPa)

intergranular fracture strength (MPa)

matrix fracture strength (MPa)

friction stress (MPa)

yield strength in artificially aged con-dition (MPa)

matrix yield strength in the absence ofhardening precipitates (MPa)

critical stress for particle cracking(MPa)

radial stress (MPa)

threshold stress for H2S stress corrosioncracking (MPa)

tangential stress (MPa)

effective surface energy (J irr2)

internal interface (or antiphase) bound-ary energy (J m~2)

surface energy of particle (J m~2)

friction coefficient

maximum shear stress (MPa)

dimensionless cooling time from 800to 5000C

dimensionless temperature conformingto 8000C

dimensionless temperature conformingto 5000C

arc efficiency factor

dimensionless supersaturation

) non-hardening precipitates in Al-Mg-Si alloys and Al-SiC MMCs

) hardening precipitates in Al-Mg-Sialloys and Al-SiC MMCs

dimensionless strength parameter

dimensionless strength parameter inreversion model

dimensionless strength parameter innatural ageing model

dimensionless strength parameter inwork hardening model

net precipitation strength increment(MPa)

maximum precipitation strength incre-ment (MPa)

HAZ strength reduction factor

angular velocity (rad/s)

Front MatterTable of Contents7. Properties of Weldments7.1 Introduction7.2 Low-alloy Steel Weldments7.2.1 Weld Metal Mechanical Properties7.2.1.1 Weld Metal Strength Level7.2.1.2 Weld Metal Resistance to Ductile Fracture7.2.1.3 Weld Metal Resistance to Cleavage Fracture7.2.1.4 The Weld Metal Ductile to Brittle Transition7.2.1.5 Effects of Reheating on Weld Metal Toughness

7.2.2 HAZ Mechanical Properties7.2.2.1 HAZ Hardness and Strength Level7.2.2.2 Tempering of the Heat Affected Zone7.2.2.3 HAZ Toughness

7.2.3 Hydrogen Cracking7.2.3.1 Mechanisms of Hydrogen Cracking7.2.3.2 Solubility of Hydrogen in Steel7.2.3.3 Diffusivity of Hydrogen in Steel7.2.3.4 Diffusion of Hydrogen in Welds7.2.3.5 Factors Affecting the HAZ Cracking Resistance

7.2.4 H2S Stress Corrosion Cracking7.2.4.1 Threshold Stress for Cracking7.2.4.2 Prediction of HAZ Cracking Resistance

7.3 Stainless Steel Weldments7.3.1 HAZ Corrosion Resistance7.3.2 HAZ Strength Level7.3.3 HAZ Toughness7.3.4 Solidification Cracking

7.4 Aluminium Weldments7.4.1 Solidification Cracking7.4.2 Hot Cracking7.4.2.1 Constitutional Liquation in Binary Al-Si Alloys7.4.2.2 Constitutional Liquation in Ternary Al-Mg-Si Alloys7.4.2.3 Factors Affecting the Hot Cracking Susceptibility

7.4.3 HAZ Microstructure and Strength Evolution during Fusion Welding7.4.3.1 Effects of Reheating on Weld Properties7.4.3.2 Strengthening Mechanisms in Al-Mg-Si Alloys7.4.3.3 Constitutive Equations7.4.3.4 Predictions of HAZ Hardness and Strength Distribution

7.4.4 HAZ Microstructure and Strength Evolution during Friction Welding7.4.4.1 Heat Generation in Friction Welding7.4.4.2 Response of Al-Mg-Si Alloys and Al-SiC MMCs to Friction Welding7.4.4.3 Constitutive Equations7.4.4.4 Coupling of Models7.4.4.5 Prediction of the HAZ Hardness Distribution

ReferencesAppendix 7.1: Nomenclature

Index