INFLUENCE OF PRECIPITATION STRENGHTENING PROCESS ON

ADVANCES IN MANUFACTURING SCIENCE AND TECHNOLOGY Vol. 32, No. 4, 2008

INFLUENCE OF PRECIPITATION STRENGHTENING PROCESS ON TENSILE AND FRACTURE BEHAVIOUR

OF THE 6005 AND 6082 ALLOYS

Grażyna Mrówka-Nowotnik

S u m m a r y

The main task of this work was to study the effect of the precipitation hardening on the mechanical properties and fracture toughness of 6005 and 6082 aluminium alloys. The mechanical (Rm and Rp0.2) and plastic (A) properties of the examined alloys were evaluated by uniaxial tensile test at room temperature. Additionally the artificially aged alloys have been tested in tension in order to determine its fracture toughness. Thus, standard ASTM tests were performed on fatigue precracked compact tension (KIc) and sharp-notched specimens ( m ) in both the longitudinal and transverse orientation with respect to the rolling direction. The results show that the microstructure, mechanical properties and fracture toughness changes during artificial aging due to the precipitation strengthening process.

kR

Keywords: 6xxx aluminium alloys, fracture mechanics, mechanical properties

Wpływ procesu umacniania wydzieleniowego na właściwości mechaniczne i odporność na pękanie stopów aluminium 6005 I 6082

S t r e s z c z e n i e

W pracy podjęto badania celem określenia oddziaływania warunków prowadzenia procesu umacniania wydzieleniowego na właściwości mechaniczne oraz odporność na pękanie stopów aluminium 6005 i 6082. Właściwości mechaniczne (Rm i Rp0.2) i plastyczne (A) wyznaczono w próbie statycznej rozciągania w temperaturze pokojowej. W celu określenia odporności na pękanie stopów umocnionych wydzieleniowo przeprowadzono próby rozciągania próbek kompaktowych (KIc) oraz z ostrym karbem ( m ) wyciętych zgodnie z kierunkiem największego odkształcenia plastycznego w kierunku do niego prostopadłym. Na podstawie otrzymanych wyników stwierdzono, że decydujący wpływ na zmianę mikrostruktury, właściwości mechanicznych oraz odporność na pękanie mają warunki prowadzonego procesu umacniania wydzieleniowego.

kR

Słowa kluczowe: stopy aluminium grupy 6xxx, właściwości mechaniczne, mechanika pękania

Introduction

The 6xxx series aluminium alloys are commonly used in extruded form. The main alloying elements in 6xxx series are Si and Mg. These alloying elements are partly dissolved in the primary α-Al matrix, and partly present in

Address: Grażyna MRÓWKA-NOWOTNIK, Ph.D. Eng., Rzeszow University of Technology, Department of Materials Science, ul. W. Pola 2, 35-959 Rzeszów, phone (+48, 0-17) 865 11 24, Fax (+48, 0-17) 854 48 32, e-mail: [email protected]

32 G. Mrówka-Nowotnik

the form of intermetallic phases. A range of different intermetallic phases may form during solidification, depending on alloy composition and solidification condition. Relative volume fraction, chemical composition and morphology of structural constituents exert significant influence on their useful properties [1-5]. Fe is present as an impurity in all commercial alloys [6]. During casting of 6xxx aluminium alloys a wide variety of Fe-containing intermetallics such as Al-Fe, Al-Fe-Si and Al-Fe-Mn-Si phases are formed between the aluminium dendrites [7, 8]. Type of these phases depends mainly on the cooling rate and the Fe to Si ratio in the alloy [9, 10].

The 6xxx aluminium alloys are mostly used in extruded Al products, as well as for construction and automotive purposes. The ease with which these alloys can be shaped, their low density, their very good corrosion and surface properties and good weldability are factors that together with a low price make them commercially very attractive. The aluminium alloys of 6xxx group have been studied extensively because of their technological importance and their exceptional increase in strength obtained by precipitation hardening.

The precipitation of the metastable precursors of the equilibrium β(Mg2Si) phase occurs in one or more sequences which are quite complex. The precipitation sequence for 6XXX alloys, which is generally accepted in the literature [11-14], is: SSSS →atomic clusters →GP zones →β''→ β'→ β (stable). The most effective hardening phase for this types of materials is β''.

In this paper differential scanning and transmission electron microscopy (SEM, TEM) have been utilized to study the effect of the precipitation hardening on the microstructure of aluminium alloy 6082. The mechanical (Rm and Rp0.2) and plastic (A,Z) properties of the examined alloy were evaluated by uniaxial tensile test at room temperature.

In spite of the improvements which come from advances in the processing of aluminum alloys their high strength is firmly entrenched as the material of construction up to service temperatures as high as 150°C. Further increases in reliability and efficiency of aluminum alloys require increases in strength and toughness. Hence, fracture toughness is a key property for a number of Al alloys utilized in aerospace and process industries. The most frequently used test methods for plane-strain fracture toughness of metallic materials are tests described in ASTM-standards – E399 and E602-78T [15, 16]. Tensile test in the presence of sharp notch as well as the currently most reliable Kk

mR IC measurement techniques where used to determine fracture toughness of heat treated 6082 alloy, on the specimens with various orientations to the rolling direction.

The results show that the microstructure, mechanical properties and fracture toughness are strongly affected by artificial aging due to the precipitation strengthening process. Therefore, the parameters (time and aging temperature) of

Influence of precipitation strenghtening ... 33

precipitation strengthening process that may lead to the most favorable mechanical properties of 6082 alloys were determined. Material and experimental

The investigation has been carried out on the commercial aluminum alloys – appointed in accordance with the standard PN - EN 573-3; 6005 and 6082. The chemical composition of the alloys is indicated in Table 1.

Table 1. Chemical composition of the investigated alloys, %wt.

Alloy Si Fe Cu Mn Mg Cr Zn Others Al

6005 0.60 0.21 0.12 0.15 0.54 0.028 0.01 0.15 bal.

6082 1.2 0.33 0.08 0.50 0.78 0.14 0.05 0.15 bal.

The thermal processing of 6005 and 6082 alloy was started by a heat

treatment at the temperature of uniform α solid solution. Thus the all specimens were heated in a resistance furnace for 4 hours at 575°C and quenched at room temperature. Subsequently the specimens were subjected to artificial aging at four different temperatures: 130oC for 72 h, 160oC for 50 h, 190oC for 42 h and at 220oC for 48 h.

After artificial aging, a set of specimens were prepared for tensile testing to study the effect of T6 heat treatments on mechanical properties of the examined alloys. The specimens were strained by tensile deformation on an Instron TTF-1115 servohydraulic universal tester at a constant rate, in according to standard PN-EN 10002-1:2004 [15] at room temperature. Tensile properties (tensile and yield strength; elongation) were evaluated using round test specimens of 8 mm diameter and 65 mm gauge length (according to ASTM E602-78T [16] standard – see Fig. 1a).The hardness was measured with Brinell tester under 49.03 N load for 10 sec.

In addition, sharp notched tensile and fatigue precracked compact tension specimens C(T) were machined for tension and fracture tests.

To evaluate tensile strength in the presence of a sharp notch standard-sized specimens (12 mm diameter x 55 mm long – see Fig. 1b) containing a sharp notch (1.75 mm deep with a 8.5 mm radius) to localize the stress were used. The tensile specimens were machined such that the gauge lengths were in the transverse and longitudinal direction with respect to the rolling direction. These round tensile specimens were machined to dimensions and tolerances required by ASTM E602-78T [16].

kmR


a)

b)

φ8

65

95

φ1

0

φ1

6

9 ,5 9,5

φ1

2

φ8

,5

φ1

2

13 13

552 2

M1

2 x

1,7

5

M1

2 x

1,7

5

Fig. 1. Dimensions of: a) plain tensile specimens and b) sharp notched tensile specimens

Fatigue precracked compact tension specimens were cut from the 6005 and 6082 alloys plates in longitudinal transverse L-T and transverse longitudinal T-L orientation with respect to the rolling direction. The specimen locations are illustrated in Fig. 2. Following the standard [17], the nomenclature defines applied loading axis by the first letter (L-longitudinal, T-transverse) and the direction of crack advance by the second letter.

Length

Longitudin

al

Rolli

ng dire

ction

TW

idth

Long Transverse

S

Thic

kness

Short T

ransv

erse

L

Fig. 2. Location of the compact tension specimens


The specimen dimensions and testing procedure to determine the fracture toughness KIc (critical stress intensity factor) are in accordance with the following standards: ASTM E399-85 [17], ANSI/ASTM B645-78 [18] and PN-87/H-04335 [19] (Fig. 3).

φ16

x34

Fatigue

precracking

24

.518

85

3

17 40

85

φ5

Fig. 3. Dimensions of a compact tension specimen C(T) used for fracture toughness testing – KIc

The fracture toughness was calculated using the following equation:

KQ = ( 1/ 2QP

BW) f ( a

W) (1)

where:

f( a ) = W

( )( )( )

2 3

3/ 2

2 0,886 4,64 13,32 14,72 5,6

1

α α α α

α

+ + − + −

−

4α (2)

where: α = Wa , B – Sample thickness, mm, W – Sample width, mm, a – Crack

length, mm, KQ – Calculated crack toughness value of a material under plane-strain conditions, MPa·m1/2, PQ – Force, N.

Experimental data were used for evaluation of crack toughness in a plane strain state – KQ. KQ is equal to the plane-strain fracture toughness KIc if the following criteria are satisfied:


a ≥ 2,52

Q

p0,2

KR

B ≥ 2,5

2Q

p0,2

KR

(3) or

Microstructure of characteristic states of examined alloys was observed using an optical microscope – Nikon 300 on polished sections etched in Keller solution (0,5% HF in 50 ml H2O). Thin foils for TEM studies were manufactured by cutting 3 mm diameter discs, followed by grinding manually to a thickness of about 0.1 mm. Finally, the disc was thinned electrolytically using a Struers Tenupol jet polishing machine, with a solution (by volume) CH3OH (84 cm3), HClO4 (3.5 cm3) and glycerin (12.5 cm3), operating at -10°C and U = 28V. The thin foils were examined in a Tesla BS-540 and JEOL-JEM 2010 ARP TEM/STEM operated at 200 kV electron microscope. Post-failure observation of the fracture surfaces of the failed C(T) specimens were made in the scanning electron microscope (SEM), operating at 6-10 kV in a conventional back-scattered electron mode.

Results and discussion

The microstructure of the studied alloys in as-cast state is given in Fig. 4. In the interdendritic spaces of α-Al solid solution one can see the precipitates of the intermetallic phases. The revealed particles of the intermetallic phases were formed during casting of the alloys. The typical as-cast structure of examined alloys consisted of a mixture of β-AlFeSi and α-AlFeMnSi intermetallic phases distributed at cell boundaries, connected sometimes with coarse Mg2Si.

The microstructure of the alloys after hot extrusion forging process is given in Fig. 4c, d. During hot working of ingots, particles of intermetallic phases arrange in positions parallel to direction of plastic deformation (along plastic flow direction of processed material) which allows for the formation of the band structure. As a result, the reduction of size of larger particles may takes place.

The strength (Rm and Rp0.2) and plastic (A) properties of 6005 and 6082 alloys after solutionizing and artificial aging at various temperatures were determined from static tensile test. As can be seen in Fig. 5, the yield strength and tensile strength of the alloys aged at 130°C increase with increasing aging time. The yield strength increases continuously with the time, however a significant increase in mechanical properties was achieved during aging for up to 20 h. Further heating causes a steady increase in the yield strength of the materials. The increment of the alloy strength similarly to the observed increment in hardness can be treated as the effects of initial formation of GP zones followed by precipitation of metastable particles of β’’ and β’ phases (Fig. 7a).


a) b)

c) d)

Fig. 4. Microstructure of examined 6005 and 6082 alloys: a, b) as-cast, c, d) after hot extrusion

a) b)

0 10 20 30 40 50 60 70 8080

120

160

200

240

280

20

25

30

35

40

45

50

Yie

ld s

tre

ng

ht

Rp0

,2,

MP

a

Te

ns

ile

str

eng

th R

m,

MP

a

Aging time, h

A

Rp0,2

Elo

ng

atio

n A

, %

0 1 0 2 0 3 0 4 0 5 0 6 0 7 0 8 0

2 00

2 50

3 00

3 50

4 00

4 50

1 0

1 5

2 0

2 5

3 0

Yie

ld s

tre

ng

th R

p0,2,

MP

a

Te

nsile

str

eng

th R

m,

MP

a

Aging time, h

A

Rp0,2

Elo

ng

atio

n A

, %

Rm

Rm

Fig. 5. Effect of aging time on tensile properties of: a) 6005 and b) 6082 aluminium alloys aged at 130°C

Relatively high values of correlation coefficient (see equation in Fig. 6) provide evidence that mechanical properties strongly depend on aging time at this temperatures.


a) b)

1 0 1 00

1 20

1 60

2 00

2 40

2 80

Rm

r=0,99 2

Rm

=15 9,8 1+52,3 87 t

Yie

ld s

tre

ng

th R

p0,2,

MP

a

Te

nsile

str

en

gth

Rm,

MP

a

Aging time, h 10 100

240

280

320

360

400

440

r=0,909

Rm=360,36+35,709t

Rm

Yie

ld s

tre

ngth

Rp

0,2,

MP

a

Ten

sile

str

en

gth

Rm,

MP

a

Aging time, h

Rp0.2 r=0,99 2

Rp0.2

=65 ,29 5+90,7 54 t

r=0,991

Rp0.2

=198,56+84,233tRp0.2

Fig. 6. Effect of time on tensile and yield strength of: a) 6005 and b) 6082 Al-alloys aged at 130°C

a) b)

020

020

200

220

220

220

200

220

β - Mg Si2

α - Al

zone axis - 001

0,1 mµ

0,1 mµ

Fig. 7. TEM micrograph of: a) 6082 alloy aged at 160°C for 6 h showing uniform dispersion of fine needle shaped of hardening β’’ particles, b) pinning dislocation by large needle-like particles of α(Al-Fe-Mn-Si) phase

The curves of Rm, Rp0.2 and A obtained at aging at higher temperature (190 and 220°C) are shown in Fig. 8. These results confirm those obtained earlier. After heat treatment at 190°C for the first few hours yield strength and tensile strength increases rapidly. The growing trend of strength continues and it reaches Rm = 265 MPa for 6005 alloy and Rm = 440 MPa for 6082 alloy after aging for 6 hours. Further extension of aging time results in insignificant variations of Rm and Rp0.2 values (Fig. 8a, b). The reason for this is the same as in the case of aging at lower temperature mentioned above. However in case of higher temperature of aging – 190°C the maximum of the strength properties is achieved after shorter time of exposing the alloys to an elevated temperature. On the other hand it was found that the aging temperature has a considerable stronger effect on the plastic properties of investigated alloys aged at higher temperature – 190°C.


a) b)

0 5 10 15 20 25

180

200

220

240

260

280

10

20

30

40

50

Yie

ld s

tre

ng

th R

p0

,2,

MP

a

Te

ns

ile

str

en

gth

Rm,

MP

a

Aging time, h

A

Rp0,2

Elo

ng

atio

n A

, %

0 1 0 2 0 3 0

2 00

2 50

3 00

3 50

4 00

4 50

1 0

1 5

2 0

2 5

3 0

Yie

ld s

tre

ng

th R

p0

,2,

MP

a

Te

ns

ile

str

en

gth

Rm,

MP

a

Aging time, h

A

Rp0,2

Elo

ng

atio

n A

, %

c) d)

0 10 20 30180

200

220

240

260

10

12

14

16

18

20

Yie

ld s

tre

ngth

Rp0

,2,

MP

a

Ten

sile

str

en

gth

Rm,

MP

a

Aging time, h

A

Rp0,2

Elo

ng

atio

n A

, %

0 10 20 30

18 0

20 0

22 0

24 0

26 0

10

12

14

16

18

20

Yie

ld s

tre

ng

th R

p0,2,

MP

a

Te

ns

ile

str

en

gth

Rm,

MP

aAging time, h

A

Rp0,2

Elo

ng

atio

n A

, %

Rm R

m

Rm R

m

Fig. 8. Effect of time on tensile and yield strength of: 6005 aged at 190 (a) and 220°C (c) and 6082 alloy aged at 190 (b) and 220°C (d)

As far as the elongation results obtained for the 6082 alloy aged at lower aging temperature – 130°C (Fig. 5b) changes with aging time irrelevantly, it can be noticed that there is a greater reduction of plastic properties with the time when the 6082 alloy was aged at 190°C (Fig. 8b).

Fig. 8 c and d represents the curves obtained in tensile static tests for the alloys artificially aged at the highest temperature of 220°C. There can be noted roughly linear drop of both yield strength and tensile strength with aging time for the examined alloys. For example, the yield strength of 6005 alloy decreases significantly from its maximum 247 MPa to a value of less than 188 MPa after 30 h of aging. The maximum Rp0.2 value was recorded after 1 h of aging. However, the plastic properties of both alloys increases with increasing time. Initially a minimal deterioration of elongation value was observed but at higher aging time up 5 hours the elongation values increase steadily.

The experimental results related to cracking sensitivity of the examined alloys show that their tensile strength in the presence of sharp notch R depends mostly on the heat treatment conditions/parameters of hardening process and orientation of the cleavage plane to the direction of maximum plastic deformation (rolling direction) (Table 2).

km


Table 2. Tensile strength in the presence of a sharp notch of 6082 and 6005 alloys

kmR

Alloy Aging oC

Aging Cleavage plane

orientation

kmR ,

MPa

130 70 T-L 505 542

190 5 T-L 553 585 6082

220 1 T-L 538 575

130 70 T-L 375 443

190 5 T-L 409 466 6005

220 1 T-L 389 425

temperature, time, h

L-T

L-T

L-T

L-T

L-T

L-T

The results have shown that the highest tensile strength in the presence of sharp notch is achieved for the sample of 6005 and 6082 alloys with longitudinal transverse L-T with respect to the rolling direction subjected to artificial aging at 190

kmR

oC (Table 2 and Fig. 9).

100 150 200 250350

400

450

500

550

600

6082

alloy

Aging time, Co

Te

nsile

str

en

gth

in

th

e p

resen

ce

of

sh

arp

nn

otc

h R

,

MP

amk

Crack plane

orientation: L -T-

T- L-

6005alloy

Fig. 9. The effect of aging time on tensile strength in the presence

of sharp notch of 6005 and 6082 alloys kmR

Critical stress intensity factor KIC was determined using standard tensile test for 30 samples of 6005 and 6082 alloys subjected to the precipitation hardening process. The crack displacement vs. load relationship, is a curve that actually


represents the results of tensile test performed on the fatigue precracked samples (Fig. 10).

0 1 2 3 40

100

200

300

400

500

PQ

Displacement, mm

Lo

ad

P,

kN

Fig. 10. Exemplary crack displacement vs. load curve of the sample 6082 alloy

a) c)

b) d)

Fig. 11. Typical fracture surface of the samples of 6005 (a, c) 6082 (b, d) alloys: a, b) T-L orientation – type A, c, d) L-T orientation – type C


The ASM has standardized 10 types of curves of varying shapes. Adequate type of a curve provide valuable data on crack growth resistance of material under static loading. According to ASM standard the shape of the curve from Fig. 11 corresponds to type I [17, 18]. Depending upon the crack plane orientation to the rolling direction, the fracture surface appears as a plane – A type, as in Fig 11a, b (T-L orientation) or as C type for L-T orientation, as in Fig. 11c, d.

Crack resistance KQ and critical fracture toughness factor KIc of the examined alloys depend on their respond to the precipitation hardening treatment (Table 3).

Table 3. Average values of KIc of examined alloys (KQ = KIc)

Alloy Aging process Fracture plane orientation

KQ, MPa·m1/2

KIc MPa·m1/2

130oC/17 h L-T - 41,0

130oC/70 h T-L - 33,1 38,52

160oC/10 h L-T - 40,5

190oC/4,5 h T-L - 37,0 43,34

190oC/6 h L-T - 34,5

220oC/1 h T-L - 34,0 38,5

220oC/4,5 h L-T - 36,3 220oC/10 h L-T - 38,0

6082

220oC/17 h L-T - 36,0 130oC/27 h T-L 51,0 - 190oC/5 h T-L 54,81 - 6005 220oC/37 h T-L 46,5 -

L-T

L-T

L-T

Based upon the recent research [1-9] it has been concluded that volume fraction of β-Mg2Si phase play crucial role in hardening and final mechanical properties of the examined alloys. 6082 alloy shows significant better mechanical properties – Rp0.2, Rm, , than the 6005 alloy. This is attributed to higher content of β-Mg

kmR

2Si precipitates in the microstructure of heat treated 6082 alloy. TEM observation showed that fine, spherical in shape particles of quaternary α-Al(FeMn)Si have also influenced strength of the studied alloys considerably (Fig. 7b). Examination of microstructure in the sample with the highest mechanical properties subjected to T6 heat treatment revealed these particles impeding the movement of dislocations. It has been reported that density of the strengthening phase precipitates of α-Al(FeMn)Si in 6082 alloy is higher than in 6005 (Fig. 12) alloy what resulted in a substantial increment of


hardening effect. At the same time its plastic properties and crack toughness have been decreasing, compared to 6005 alloy.

a) b)

0,2 mµ

0,2 mµ

Fig. 12. Microstructure of 6005 (a) and 6082 (b) alloys after precipitation hardening process, showing spherical particles of α(Al-Fe-Mn-Si) phase

The obtained results allow to conclude that the crack resistance of the alloys is strongly influenced by the condition of heat treatment – aging time and temperature. Analysis of the data obtained by static tensile tests revealed that 6005 and 6082 alloys had nearly the same mechanical properties after aging at various temperatures for the same time.

Table 4. The effect of the temperature and aging time on the mechanical properties of 6082 and 6005 alloy

Mechanical properties Alloy Points in the

Fig. 11 Aging

process Rp0.2, MPa KIc MPa·m1/2 1 130 oC/17h

220 oC/17h 298 41,0

36,0 2 160 oC/10h

220 oC/10h 327 40,5

38,0 3 190 oC/4,5h

220 oC/4,5h 357 37,0

36,0

6082

4 220 oC/1h 190 oC/6h

380 34,5 34,0

6005 5 130 oC/27h 220 oC/27h 190 51,0

46,5 After tension of the compact fatigue prenotched samples it was found that

the samples of both alloys exposed to aging process at various temperature, but for the same period of time (e.g. 130oC and 220oC – point 1(17 h) and 5(32 h);


160oC and 220oC – point 2 (10 h), Fig. 13) even if they are characterized by practically the same yield stress value, they showed differences in their crack resistance, what resulted in different values of the critical stress intensity factor KIc (Table 4).

a) b)

1 10 10080

120

160

200

240

280

5

220 °C

130 °C

Yie

ld s

tren

gth

Rp

0,2,

MP

a

Aging time, h 1 10 100

200

250

300

350

400

450

2

1

3

4

160 °C

220 °C

130 °C

Yie

ld s

tren

gth

Rp

0,2,

MP

a

Aging time, h

190 °C

190 °C

Fig. 13. The effect of the temperature and aging time on yield stress of: a) 6005 and b) 6082 alloy

The samples of 6005 and 6082 alloy aged at lower temperature – 130oC and 160oC have much higher crack resistance compared to the samples aged for the same period of time at 220oC (Fig. 14 – point 1 and 2, Table 4). Only in the case of shorter aging time – 1 h at 220oC and 6 h at 190oC similar values of KIc parameters were obtained (Fig. 14 – point 4, Table 4). Prolongation of aging time at the highest temperature of 220oC results in overaging (Fig. 15) indicated by a significant drop in KIc values. 6005 and 6082 alloys are underaged when

6005 alloy

6082 alloy

30

55

35

40

45

50

200 400300 350250

Yield strength R , MPap0.2

Cri

tica

l str

ess

in

ten

sit

y f

acto

r K

, M

Pa

*mIc

1/2

Fig. 14. The influence of yield strength Rp0.2 on the critical stress intensity factor KIc of the examined alloys – 6005 and 6082


undergoing heat treatment at lower temperature (130oC, 160oC) for 10 and 17 hours – what results in the higher value of KIc (Fig. 14 – points 1, 2, 3, 5, Table 4).

0,2 mµ

Fig. 15. Overaging alloy (300oC/20h) – pinning dislocation by large needle-like particles of β and α(Al-Fe-Mn-Si) phase

Crack resistance of 6005 and 6082 alloy – similarly to the tensile strength in the presence of sharp notch , is anisotropic property – depends on fatigue-cracking plane orientation to the rolling direction. The values of K

kmR

Ic parameter are higher for the samples with the crack plane perpendicular to the direction of greatest plastic deformation – L-T.

0,2 mµ

Fig. 16. TEM micrograph of the 6082 alloy subjected to the precipitation hardening process, exhibiting large

iron particle


Metallographic and fractografic observation of the microstructure and the fracture of the aged samples of 6005 and 6082 alloys with the highest tensile stress in the presence of sharp notch (Fig. 15 and 16) confirmed that fracture usually initiates within void clusters as a result of a sequence of void nucleation, void growth, and void coalescence.

Fractografic examination has been carried out on the fractures of the samples after: static tensile tests, crack resistance tests and tensile test in the presence of sharp notch k

mR ig. 17-19). (F

a) b)

Fig. 17. Fracture surface of the 6082 alloy after static tensile test in the presence of sharp notch a) shear oval dimples formed after coalescence of the linear void sequence, b) large dimples

around hard intermetallic α(Al:k

mR

8Fe2Si) and β(Al5FeSi) precipitates and smaller around dispersive hardening β-Mg2Si and α-Al(FeMn)Si precipitates

a) b)

Fig. 18. Morphology of the fracture surface in the deformed 6082 alloy. Typical traces of fatigue cracking are shown

Analysis of the literature [20-22] and experimental data confirmed that in the materials consisting second-phase particles, voids firstly heterogenically


nucleates at precipitates of the intermetallic phases. Decohesion process takes place first of all around the non-metallic inclusions and second-phase particles, which is around sites of the interface matrix-particle. a) b)

c) d)

Fig. 19. Typical fracture surface of the tensile specimens observed in the examined alloys (a), debonding of the intermetallics precipitates in the matrix (b), cracked α-Al(FeMn)Si particles

(c) and (d)

Better crack resistance of 6005 alloy than 6082 is due to lower density of intermetallic phases. Volume fraction of the intermetallic phases in 6082 alloy amount VV = 5,4%, thus the number of potential sites (interfaces matrix-particle) for void nucleation and their growth is considerable. This leads to decrease of crack resistance of 6082 alloy.

Conclusion

1. The initial increase in the strength properties of the examined aluminium alloys – 6005 and 6082, is due to initial precipitation of GP zones and then formation of very fine needle-shaped particles of metastable phases – β’’ and β’


in under-aged and peak-aged conditions. These precipitates effectively interfere with the motion of dislocations. A decrease in the hardness and mechanical properties of the alloys in the over-aged conditions (increase in aging time and temperature) has occurred because of coalescence of the precipitates into larger particles of metastable β’’ and β’ phases.

2. Both alloys aged at 190°C for 6 hours exhibit the best combination of performance properties including strength (6005 alloy – Rp0.2 = 255 MPa, Rm = 233 MPa and 6082 alloy – Rp0.2 = 390 MPa, Rm = 441 MPa) with good fracture toughness: 6005 – = k

mR 466, KIc = 54.85 MPa·m1/2 and 6082 – 585, K

kmR =

Ic = 43.34 MPa·m1/2. 3. Fracture toughness of the alloys essentially depends on temperature and

aging time as well as orientation of cleavage surface to the rolling direction. The highest value of tensile strength in the presence of sharp notch k

mR 585 MPa (6082 alloy) and 466 MPa (6005 alloy) and critical stress intensity factor KIc = 54.85 (6005 alloy) and 43,34 MPa·m1/2 (6082 alloy) were achieved for the specimens machined in the L-T orientation and aged at 190oC for 6 hours. Similarly heat treated specimens but with T-L orientation showed following values of: k

mR 409 MPa (6005 alloy); kmR 553 MPa (6082 alloy) and KIc =

46 MPa·m1/2 (6005 alloy); KIc = 37 MPa·m1/2 (6082 alloy).

=

= =

4. Observation of microstructure and fracture surface of the failed C(T) specimens (TEM, SEM) showed that cracking of the examined alloys begin by nucleation and growth of voids. The sites of heterogenic nucleation of voids are the precipitates of intermetallic phases. Subsequent decohesion process initially proceeded at the interface between matrix and particle. The analysis of results revealed that 6005 and 6082 alloys treated at lower temperature (130 and 160oC) for 10 i 17 h were underaged what resulted in the higher value of KIc. Prolongation of aging time at higher temperature 220oC cause overaging of the alloys – lower KIc value.

Acknowledgements

This work was carried out with the financial support of the Ministry of Science and Information Society Technologies under grant No. N507 3828 33.

References

[1] S. KARABAY, M. ZEREN, M. YILMAZ: Investigation of extrusion ratio effect on mechanical behaviour of extruded alloy AA-6101 from the billets homogenised-rapid quenched and as-cast conditions. Journal of Materials Processing Technology, 160(2004),138-147.

[2] G. MRÓWKA-NOWOTNIK, J. SIENIAWSKI: Influence of the heat treatment on the microstructure and mechanical properties of 6005 and 6082 aluminium alloys.


Proc. Int. Conference on Achievements in Mechanical & Materials Engineering, Gliwice-Wisła 2005, 447-450.

[3] M. WARMUZEK, J. SIENIAWSKI, A. GAZDA, G. MRÓWKA: Analiza procesu powstawania składników fazowych stopu AlFeMnSi w warunakch zmiennej zawartości metali przejściowych Fe i Mn. Inżynieria Materiałowa, 137(2003), 821-824.

[4] G. MRÓWKA-NOWOTNIK, J. SIENIAWSKI: Influence of heat treatment on the micrustructure and mechanical properties of 6005 and 6082 aluminium alloys. Journal of Materials Processing Technology, 162-163(2005), 367-372.

[5] S. ZAJAC, B. BENGTSSON, Ch. JONSSON: Influence of cooling after homogenization and reheating to extrusion on extrudability and final properties of AA 6063 and AA 6082 alloys. Materials Science Forum, 396-402(2002), 675-680.

[6] M. WARMUZEK, G. MRÓWKA-NOWOTNIK, J. SIENIAWSKI: Influence of the heat treatment on the precipitation of the intermetallic phases in commercial AlMn1FeSi alloy. Journal of Materials Processing Technology, 157-158(2004), 624-632.

[7] R.A. SIDDIQUI, H.A. ABDULLAH, K.R. AL-BELUSHI: Influence of aging parameters on the mechanical properties of 6063 aluminium alloy. Journal of Materials Processing Technology, 102(2000), 234-240.

[8] N.C.W. KUIJPERS, W.H. KOOL, P.T.G. KOENIS, K.E. NILSEN, I. TODD, S. VAN DER ZWAAG: Assessment of diffrent techniques for quantification of α-Al(FeMn)Si and β-AlFeSi intermetallics in AA 6xxx alloys. Materials Characterization, 49(2003), 409-420.

[9] G. SHA, K. O’REILLY, B. CANTOR, J. WORTH, R. HAMERTON: Growth related phase selection in a 6xxx series wrought Al alloy. Materials Science and Engineering, A304-306(2001), 612-616.

[10] A.K. GUPTA, D.J. LLOYD, S.A. COURT: Precipitation hardening in Al-Mg-Si alloys with and without excess Si. Materials Science and Engineering, A316(2001), 11-17.

[11] G.A. EDWARDS, K. STILLER, G.L. DUNLOP, M.J. COUPER: The precipitation sequence in Al-Mg-Si alloys. Acta Materialia, 46,11(1998), 3893-3904.

[12] W.F. MIAO, D.E. LAUGHLIN: Precipitation hardening in aluminum alloy 6022. Scripta Materialia, 40,7(1999), 873-878.

[13] G. BIROLI, G. CAGLIOTI, L. MARTINI, G. RIONTINO: Precipitation kinetics of AA4032 and AA6082 a comparison based on DSC and TEM. Scripta Materialia, 39,2(1998), 197-203.

[14] PN-EN 10002-1+AC1: Metals. Standard method of tensile test at room temperature.

[15] ASTM E602-78T: Tentative method for Sharp-notch tension testing with cylindral specimens.

[16] ASTM E399-85: Standard method of test for plane-strain fracture toughness of metallic materials.

[17] ANSI/ASTM B645-78: Standard Practice for plane-strain fracture toughness testing of aluminium alloys.

[18] PN-87/H-04335: Standard method for plane-strain fracture toughness. [19] Z. KĘDZIERSKI: Role of second-phase particles in the fracture of sites with

narrow plastic zone. Metalurgia i odlewnictwo, z. 117, Wyd. AGH, Kraków 1998.


[20] J.W. WYRZYKOWSKI, E. PLESZAKOW, J. SIENIAWSKI: Deformation and Fracture of Metals, WNT, Warszawa 1999.

[21] Z. LI, A.M. SAMUEL, C. RAYINDRAN, S. VALTIERRA, H.W. DOTY: Parameters controlling the performance of AA319-type alloys: Part II. Impact properties and fractography. Materials Science and Engineering, A367(2004), 111-122.

[22] F.J. MACMASTER, K.S. CHAN, S.C. BERGSMA, M.E. KASSNER: Aluminium alloy 6069 part II: fracture toughness of 6061-T6 and 6069-T6. Materials Science and Engineering, A289(2000), 54-59.

Received in December 2008

Documents

INFLUENCE OF PRECIPITATION STRENGHTENING PROCESS ON