Strength and Toughness of Materials || Titanium Alloys

9

Titanium Alloys

In Japan, because there is no munitions industry or aircraft industry owing to the social conditions of the country, it is imperative to find applications for titanium alloys in the other fields. For that reason, there has been a fluctuation in the research and development of titanium alloys. However, attention has been paid to the application of titanium alloys as structural materials for transport machines especially on the automobile industry, as a biomaterial for dentistry, and as a material for instruments used at cryogenic temperatures. Increasing efforts are advancing toward the study of titanium alloys.

Therefore, high performance of titanium alloy is expected as a lightweight material with a high strength-weight ratio. To understand this, strength and toughness of titanium alloys are important as the fundamentals for these materials.

This chapter presents discussion of the relationship between microstructure and strength or toughness of titanium alloys; strengthening and toughening by microstructure control of titanium alloys will be mentioned as well.

9.1 Mechanical Characteristics of Titanium Alloys

9.1.1 Mechanical Properties of Titanium Alloys

To compare strength and toughness of titanium alloys with other high-strength materials, correlations among yield strength (O'y) and fracture toughness (KId of titanium alloys, high strength aluminum alloys, high strength steel, and maraging steel are summarized in Fig. 9.1. The yield strength and toughness band of titanium alloys is situated between that of high strength aluminum alloy and that of high strength steel. Recently, from the viewpoint of global environmental protection in addition to energy efficiency, weight saving of transport machines is earnestly desired; a material with excellent mechanical properties per unit density is required. Thus, yield strength and fracture toughness shown in Fig. 9.1 are divided by the density of each material; the correlations can then be expressed with specific strength (strength-density ratio) and specific toughness (toughness-weight ratio). Figure 9.2

T. Kobayashi, Strength and Toughness of Materials© Springer Japan 2004

190 9 Titanium Alloys

........ ... E --CD

::E 40 --'" E

cu 30 0..

::E '-'

v> (I) 0)

c ..c CD ~

~ u .. ~

u cu

c!: 0)

t;:

u u Co

(/)

~:- 200 c ...... :tl

::E ...... .. 0..

::E

.- 100 E

... '" CQ

.§ <> I: 0

<> 50 u c.

en

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~ SO~ U •

J: II?" "«.,,. lal5<t)

o SOO 1000 I SOO 2000 250 Yield Ire .0, (MPa) )

200 300 J

pecific yield Ire . (MPa/ Mg/m )

100 150 200 250

pccific ten ile trcngth , (MPa/Mg m1

Fig. 9.1. Comparison of fracture toughness and yield stress among Ti alloy, high strength Al alloy, high strength steel, and maraging steel.

Fig. 9.2. Comparison of specific fracture toughness and specific yield stress among Ti alloy, high strength Al alloy, high strength steel, and maraging steel.

Fig. 9.3. Comparison of specific tensile strength and specific fatigue limit among Ti alloy, high strength Al alloy, and quenched and tempered low alloy steel.

clearly shows that titanium alloy is superior to other high strength materials (although validity of such a comparison is controversial).

Figure 9.3 shows fatigue strength (fatigue limit) at 107 cycles and tensile strength of titanium alloy, high strength aluminum alloy, and low alloying steels in a

9.1 Mechanical Characteristics of Titanium Alloys 191

Table 9.1. Classification of commercial titanium alloys based on microstructure2) •

a. stabilizing element (AI, 0)

Neutralizing element (Sn, Zr)

~-stabilizing

a. type (RCP)

near a. type

a.+~ type

near ~ type

element ~ type (V, Mo, Fe, Cr, Nb) (BCC)

Ti-O (CP-Ti) Ti-SAI-2.SSn

Ti-6Al-SZr-O.SMo-O.2Si (IMI-68S) Ti-S.SAI-3 .SSn-3Zr-O.3Mo-l Nb-O.3Si (IMI-829) Ti-8Al-IMo-1 V Ti-6AI-2Sn-4Zr-2Mo

Ti-6AI-4V Ti-6A1-6V-2Sn Ti-6AI-2Sn-4Zr-6Mo

Ti-SAI-2Sn-2Zr-4Mo-4Cr (Ti-17) Ti-IOV-2Fe-3Al

Ti-I1.5Mo-6Zr-4.SSn (B-m) Ti-ISV-3Cr-3AI-3Sn Ti-ISMo-SZr-3Al Ti-ISMo-SZr Ti-13V-llCr-3Al

quenching and tempering condition. This figure illustrates how the superiority of titanium alloy can be understood from its fatigue aspect [1].

9.1.2 Classification of Titanium Alloys and their Mechanical Properties

As shown in Table 9.1 [2], titanium alloys can be classified into an a type consisting of only a phase (HCP structure), a f3 type consisting of only f3 phase (BCC structure) and an a + f3 type consisting of both a and f3 phases, according to their microstructures. More specifically, it can be divided into a near a type alloy, which is between a and a + f3 type alloys, and a near f3 type alloy which is between a + f3 and f3 type alloys. The relations of yield stress and fracture toughness shown in Fig. 9.1 show that a type alloy shows low strength but high toughness, f3 type alloy shows high strength and high toughness. More than 50% of Ti alloys in practical use are presumed to be an a + f3 type alloy of Ti -6AI-4V.

Figure 9.4 shows the equilibrium phase diagram sectioned at A16%. For titanium alloys, martensite a' (hexagonal) and retained f3 microstructure can be obtained when an alloy is quenched from the f3 (BCC) range. As shown in this figure, martensite transformation starting temperature (Ms) is very high. However, this martensite in titanium alloy is too soft to serve as a strengthening phase, in contrast to the martensite phase in steels. Therefore, aging treatment is generally carried out to obtain strengthening by precipitation of fine a (HCP) phase in matrix (Solution Treatment and Aging: STA 1).

For a, near a and a + f3 type alloys, when it is finally worked within a + f3 range, the microstructure obtained will be equiaxed a by this a + f3 treatment; when

1 It is said that the strengthening mechanism differs from that of aging of aluminum alloy.


~

p <I) .... a e 800 <I)

I>. e ~ ~

Ti-6%Al/ /~(mass%~_ Ti-6%AI-4%V

Fig. 9.4. Quasi-binary phase diagram at 6% AI cross section of Ti-6% AI-4% V alloy.

it is finally worked within 13 range, the microstructure obtained will be acicular a (Widmanstiitten a) by this 13 treatment. For a + 13 type alloys, microstructure will develop very intricately and mechanical characteristics vary corresponding to microstructural variation because STA treatment is performed. Nevertheless, with the same heat treatment, the equiaxed a microstructure generally shows excellent tensile strength and fatigue strength, whereas the acicular a microstructure shows superiority in creep strength, fracture toughness, and fatigue crack propagation resistance.

9.2 Influence of Microstructure on Fracture Toughness

9.2.1 Equiaxed a Microstructure

Figure 9.5 shows a typical (a + 13) type titanium alloy of Ti-6AI-2Sn-4Zr-6Mo alloy (for titanium alloy, it is usually designated as Ti-6 . 2 . 4 . 6 alloy) with equiaxed a microstructure. Usually, there are primary equiaxed a (black part in the figure), 13 and precipitated finer a around equiaxed a. For alloys with this microstructure, it is believed that the strength is controlled by precipitated finer acicular a, whereas the fracture is controlled by primary a [3]. The influence of a phase on crack initiation and crack propagation when fracture occurs varies with volume fraction, spacing, and shape (size, aspect ratio). Figure 9.6 shows an example of the influence of the content of primary a on KIC of Ti -6AI-4V alloy, where the KIC value increases with decreasing primary a [4]. In this case, it is assumed that the shape of finer precipitated a is same, but it has been pointed out that the shape of finer precipitated a also plays a role [4].

Based on the correlation between the content of primary a and K1c value of Ti-6AI-4V alloy, the correlation between average spacing (Am) of primary a and

9.2 Influence of Microstructure on Fracture Toughness 193

Equiaxcd a 100

0 E

'" 60 Q..

~

40 :..:

iculUQ II 20

I~ 0

Fig. 9.S. Typical SEM micrograph of equiaxed a structure of Ti-6AI-2Sn-4Zr-6Mo alloy(1123 Kx 1 hAC, 863 Kx 6 h AC) (AC: Air cooling).

Fig. 9.6. Effect of primary a percentage (%) on fracture toughness in Ti-6AI-4V4).

E

~ 100 :2

N 90 E ..... -.

-'" 0 '" E

70 ..... -. -'" -:- 60

0 ~ 50 . -. .; 40 -.

u.f 300

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1" ... "K 0 c) • • 11- :B • •• • * : : •• "' ••• Cf> Q ...,.

... /, ... ..... ... /., , ® . 1l' • ....

........ ............ J IC

9-" 8 ............. ", ... "'"

0.5 1.0 1.5 Mean free path of primary a : Am, !lm

6

4

2

0

E f-

Fig. 9.7. Relation among Et, Jd, JIC , Ko, Tm'l> and mean free path of primary a (equiaxed a structure; A: 978Kx2 hAC, B: 1173Kx 1 hAC, 813K x 4 hAC, C: 1228K x 1 hAC, 813 Kx 4 hAC, AC=Air cooling, WQ=Water quenching).

crack propagation initiation toughness can be deduced as shown in Fig. 9.7. The correlation between crack propagation resistance toughness and Am is also shown in this figure, where good consistence between two kinds of toughness and Am is seen. In this case, correlations between Am and critical opening displacement (oc), which indicates crack propagation initiation toughness, can be projected with Oc = A . Am (A: constant). However, this formula can only be used for cases where the content of primary a is higher than 70% (for finer precipitation a + {3, it is about 30%). On the other hand, Hirth et aI. [5] reported that they found a relationship of cry (MPa) = 850.2 + 69.5A-1(cry: yield stress, A: spacing between finer precipitation a (jim» for a series of alloys. This relationship is considered in the K1C calculation because strength and toughness have an inverse relation; its validity has been reported. In


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~ 40 :::I

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• • • • 0 Pnor (l grain 17C • 0 D i3 .0·25 .2!lm

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7.

76.9 E

65.9 ~ ~

54 .9

43.9

100 200 300 400

Reciprocal 0. particle ize; j. mm

Fig. 9.8. KQ as a function of reciprocal Ea particle traverse diameter, I/d8).

eicular 0. + J}

• Prior J} grain boundary

.. rain boundary 0. ~tJE~~II~ Acicular a.

Fig. 9.9. Typical SEM micrograph of acicular a structure of Ti-6AI-2Sn-4Zr-6Mo alloy (1338 Kxl hAC, 863 Kx6 hAC).

Fig. 9.10. Schematic presentation of submicrostructure observed by optical microscope in prior f3 grain.

this case, finer precipitation a phase is believed to show great influence on crack propagation initiation toughness.

When primary a is in a spherical particle shape, KIC and crack propagation resistance is inferior to the case of coarsened elongated grain [6], but the latter shape can facilitate void formation leading to decreased elongation [7]. However, for the former discussion, the relationship between a reciprocal number of primary a size and the ~ value (as a reference value of KId (as shown in Fig. 9.8) is reported in positive correlativity [8], implying the significant effect of the shape of primary a grains.

9.2.2 Acicular a Microstructure

Figure 9.9 shows typical acicular a microstructure in (a + 13) type titanium alloy, Ti-6AI-2Sn-4Zr-6Mo, where the a phase (grain boundary a) is widely precipitated on the original 13 grain boundary, whereas acicular a is precipitated inside grain, and 13 or finer aging a are precipitated between the acicular a phases. The acicular a inside original 13 grain almost consists of the same direction groups (colony or packet), where small groups with same orientation in a colony are identified. The author has defined this structure as a sub-colony. This does not require consistency with a sub-unit in the colony structure, defined by Kishi et al. as block, formed during


~

S ~t :d ~C ~Q T~at:G oj 160 e - :H 20 ~ 0 <1>& ® tt :I ~ 140 () ~9- .. * :J

N • •• $ * :K

* S 120 .<l Tmay 15 -- Jd

...... ,.;.:; / . , /

N • 100 / It '<;j s Ko ;" ~ .. ~ s -- ... - .. -~ . f-< Fig. 9.11. Relation among ...... _.$-" - ':7 ,.;.:; 80 10 Er. Jd , Jre , KQ, Tma! and ....... / /

/ / prior f3 grain size (acicular CI 60 / ~ ... / a structure; G: 1338 Kxl

/ u

40 / 5 h WQ, 813 Kx4 hAC, H: ....;' / 1338 Kxl h WQ, 978 Kx4 .; / ...... hAC, I: 1423 Kxl h FC, 20 W 813 Kx4 hAC, J: 1423

0 1.20 Kxl hAC, 813 Kx4 h AC

0.9 and K: 1423 Kxl h WQ,

Prior P grain size, mm 813 Kx4 hAC).

martensitic transformation by rapid cooling [9]. When a colony was excessively coarsened, such a sub-colony can be formed through a nucleus inside the colony (as explained later). At present, although the technical term for the sub-microstructure of acicular a has not been well accepted, this chapter uses terms which are described schematically in Fig. 9.10 [10, 11].

Though ductility or elongation of acicular a microstructure is inferior to that of equiaxed a microstructure, crack propagation initiation toughness and fatigue crack propagation resistance are superior. Such properties increase with microstructure coarsening [12, 13]. Figure 9.11 shows the correlation between original f3 grain size and various toughness values of Ti-6AI-4V alloy. Toughness value (Kq, Et ),

which engenders the main component of crack propagation characteristics, and crack propagation resistance toughness (T mat) can be improved by increasing original f3 grain size. Both Jd (dynamic fracture toughness value, which is not satisfied with a valid criterion) and JIe show a similar dependence on the original f3 grain size. Still, it is difficult to find a good correlation because sub-structure of original f3 grain, spacing of acicular a phase and spacing among sub-colonies also play a role in toughness. In fact, because the treatment after solution is different for present materials, it is hard to claim that original f3 grain microstructures of samples are similar; this point remains an open question.

Therefore, by controlling soaking time at solution temperature and fixing treatment parameters after solution, original f3 grain size can be coarsened with heat treatment. Figure 9.12 shows an example of the correlation between original f3 grain and the toughness of Ti-6AI-2Sn-4Zr-6Mo alloy. The elongation, tensile strength, 0.2% yield strength and typical equiaxed a microstructure are shown in that figure for comparison. Compared to that of equiaxed a microstructure, strength and elongation


x

". ...... II Typical equiaxed a I 0

~3orl structure specim~y ~ "! / Jd.~ ~ ~ ~ o~-e-~ ~ • -e-~-e-~IOO ~ ~ ~ (1)20 e-e oB _ "0 --e (1)] .--.~ '" B 50::; 0 Ey ~ Fig. 9.12. Effect of prior f3

grain size on mechanical properties of furnace-cooled acicular a structure specimen (Ti-6A1-2Sn-4Zr-6Mo

~.g --~ ~ '-' <;j 10 ~ 0.52 ·s ;§ -e- ~ ~ r' ~ ';f!, ~ 0 O. I "' ... r'-'-....LJ....LJJ.J.I_--'-'--'-.J...LJUJJI 0 N

Ci Prior ~ grain size, mm 0

o

~V>

120 ,.-., ,.-.,

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! J .. K5

~~

20

0 Fine Coarse

Lamellar

,.-., ~V> V '-'

s :::t

0 0 'D

•

alloy) 10).

Fig. 9.13. Stress intensity factors for Ti-6AI-4Y. Lower line: Ko (where the load versus COD curve deviated from the linear elastic line), Second line: Ks (calculated according to the ASTM procedure), Upper line: Kmax (where unstable fracture occurred), Arrows: sectioned specimen and crack extension in specimen centerI6).

of acicular microstructure are inferior, whereas Jd and Et values are remarkably improved. On the other hand, increasing the original f3 grain size demonstrates that strength decreases slightly, whereas elongation and Jd values increase and the Et

value increases remarkably. Compared with the increase in Jd value, the increase in the Et value is more prominent. Therefore, it is readily inferred that crack propagation resistance toughness is remarkably improved. Coarsening of f3 grain also leads to the increase in the acicular lY width. As shown in Fig. 9.13, toughness is improved with increasing acicular lY width [14]. On the other hand, it has been reported that the i11<Q value, the increment of I<Q, is also influenced by thickness of the grain boundary lY.

Figure 9.14 shows that the i11<Q value increases with increasing thickness of the grain boundary lY until a certain critical thickness is reached [8]. Moreover, Yoder et


30

25 O:"-.-.JUo-=-_..".... o 0

33.0

27.5 c E

20 22 .0 co

'" 15 "'" ':.too 10 <J

16.5 0..

~ 11.0 0

':.to 5 5.5

00 1 2 9 10 11 12 0.0 Fig. 9.14. ~~ as a function of GBa thickness; 18).

Grain boundary (l thickne ,I,!Jm

.... u.r Prior >.

J3 grai b-colon CII .... ., c .,

"Q ., 20 -e 0

'" ~ 10 '" .~ '" '0 I- 0.2 0.3

ub-colony pacing, mm

Fig. 9.15. Relation between total absorbed energy (Et) and sub-colony spacing.

~ I I

- 0 olony 0 E

- '. ub-colony E 1.0 eO

• c: E

0.3 (,)

E co Q.

<i '" >. c:

(I) .2 >. 0 c: • (,) 0 0 .2 .D 0 - -. :;,

V'l

0.5

0 . 1 1 10 Prior ~ grain ize,mm

Fig. 9.16. Relation between sub-microstructure (colony, sub-colony spacing) and prior f3 grain size.

Fig. 9.17. Precipitating process of acicular a structure specimens quenched from (a) 1023 K, (b) 1073 K, (c) 1123 K during furnace cooling process from 1338 K.

al. reported [15] that for annealed Ti-6AI-4V alloy from p range, the fatigue crack propagation rate decreases to less than 1/5 when colony dimensions increases.

Improved toughness as a result of microstructure coarsening was well supported by experimental evidence. Notwithstanding, the author [10] has specifically


N 140 -e ...... ..I<i 120

'" ......

~ 100 CLl = ~ 80 ::s 2 ~ 60 E ~ 40

<I:: u .§ 20

'" = >. 0 ~ ABCIJK

(a)

70

...... 60

~ ;.:, 50 e? CLl

~ 40 "0

] 30 ... o ., ~ 20 «i ~ 10

o ABCIJK

(b)

Fig. 9.18. Comparison between (a) dynamic fracture toughness (Jd) and (b) total absorbed energy (Et) (specimen notations A-K correspond to ones in Figs. 9.7 and 9.11).

addressed effects of sub-microstructures such as sub-colony spacing. Experimental results have demonstrated that the Et value increases with refining sub-colony spacing, as shown in Fig. 9.15. Consequently, the relationship between sub-colony size or colony size and original f3 grain size were examined; that result is shown in Fig. 9.16. It is seen that sub-colony size decreases when original f3 grain size and colony size increases [10, 11]. This is similar to the situation in steel during martensitic transformation from worked austenite, in which the packet dimension increases with increasing working strain, but the block size decreases [16]. However, when formation of such a sub-colony leads to excessive f3 grain coarsening, nuclear formation of new colony occurs frequently inside a colony. It has been considered to be one kind of known resonance nuclear formation 1 (Fig. 9.17). This implies various possibilities for toughening titanium alloys in the future. Moreover, it can also be understood from Figs. 9.11 and 9.12, that crack propagation initiation toughness (Jd,

Jle) and crack propagation resistance toughness (T mat) do not necessarily show good consistence. This suggests that controlling microstructural factors are different for crack propagation initiation toughness and crack propagation resistance toughness.

9.2.3 Microstmctural Units Controlling Crack Propagation Initiation Toughness

As shown in Fig. 9.18, different toughness evaluations are obtained by V notched Charpy testing and dynamic fracture toughness value (Jd) testing. In this figure, the Charpy test result is shown by the Et value: the absorbed total energy including crack propagation energy. The Et was used to characterize the toughness instead of ~ so that the difference in toughness for different samples can be more clearly seen. In this figure, A, B and D are results for the equiaxed a microstructure and I, J and K are results for acicular a microstructure [17]. By Charpy examination,

1 The precipitated grain has different composition with matrix nucleates on other crystal interface of the same phase, which was proposed by Aaronson.

9.3 Micromechanism of Crack Initiation and Crack Propagation 199

til til 0)

.: ..c: bl) ;:l

£ 0) .... B u oj

-0-Equiaxed a -e-Acicular a

t!:: e----:f-...J U I

. § 0-----:,..... -c:{1 ~ Ifo=lo I;

o

I

;;.., Cl Jd CVN

Notch root radius: p, ~m

Fig. 9.19. Schematic representation of the difference of toughness evaluation between Charpy V notch (CVN) test and dynamic fracture toughness test.

the highest toughness value is obtained with equiaxed a microstructure. The result shown in Fig. 9.18 has also been reported for steels and aluminum alloys [18, 19]. This phenomenon is explained with the difference in tip radius of main crack, that is, the size of the microstructural unit (l~) controlling fracture at crack -tip is different between equiaxed and acicular a.

According to Ritchie's strain-controlled fracture model (refer to 3.2), if assuming fracture starts when strain reaches the critical strain (Sf) in entire l~ range, the Jd

value can be expressed by Jd ex: O'yd . Sf . P (O'yd: dynamic yield stress, p: notch tip radius). Therefore, the Jd value remains almost constant with P being less than 1~/2 (at this moment, p = Po = 1~/2), whereas Jd increases in proportion with p if P further increases from Po. This phenomenon was explained in Fig. 3.8 and Fig. 4.24. Figure 9.19 shows a schematic illustration. In other words, when the notch tip radius is smaller than Po, the Jd value of equiaxed a microstructure is less than that of acicular a microstructure. On the other hand, when radius of the notch tip is larger than Po, the Jd value of acicular a microstructure is higher than that of equiaxed a microstructure because l~ of acicular a microstructure is larger than that of equiaxed a microstructure. For p (CVN) in Charpy testing, it was found that the Jd value (It corresponds to evaluation by Charpy absorbed energy) of equiaxed a microstructure exceeds that of acicular a microstructure.

9.3 Micromechanism of Crack Initiation and Crack Propagation

As shown schematically in Fig. 9.20, macro-crack initiation occurs at the crack-tip in plastic zone with high strain under large-scale yielding condition [20]: formation of a void and a microcrack or accumulation of dislocation in a process zone [21]. When strain in this range reaches fracture strain, the void, microcrack, or dislocation grows and combines with the main crack (refer to 3.4). The process zone size is equivalent to the size of the microstructure unit which controls fracture, which in turn corresponds to the spacing between inclusions or carbides or second phase particles in steel [22] and aluminum alloys [23]. Brittle second phase particles are


Elastic-plastic boundary

Intensely deformed zon

Oy: Plastic zone

Or: Process zone

Voids, microcrack

and dislocation (intensely deformed zone)

ry: Plastic zone size r f: Process zone size

E: Equivalent strain ey : Strain in the plastic zone

er: Strain in the process zone

Fig. 9.20. Schematic explanation of the process zone21 ).

(a)

~~ :~. d'!

(b)

Deformation~Wins or ", /A slip bands ~ /'"

~'.

""

~"/A ~/A ...• ~Offse~ A'"

Voids?"'i

(1) Grain boundary or twin boundary

(2) (3)

(c)

(a) Interaction between slip band and grain boundar/7)

37) (b) Interaction between slip bands Fig. 9.21. Schematic (c) Interaction between twin and matrix or slip band

and grain boundar/8)

presentation on the various void formation models.

found in dimples by fracture surface observation. In this fracture mode, a macrocrack initiation is caused by breakage or failure of second phase particles and interface debonding between second phase particles and the matrix. Usually, for fracture of titanium alloy including (a + f3) type at room temperature, it is hard to find second phase particles in dimples, as shown in Fig. 6.12, where second phase particles were not observed. Therefore, void formation is not associated directly with second phase particles, at which the macro-crack forms. This must be explained

9.3 Micromechanism of Crack Initiation and Crack Propagation 201

with other mechanisms. Therefore, for titanium alloys, void formation sites could be attributed to mechanisms such as vacancy clustering [24], interaction between the grain boundary and slip bands (Fig. 9.21(a) [25], Fig. 9.21(c) [26]), interaction between slip bands (Fig. 9.21(b) [25]) or interface of a twin-matrix (Fig. 9.2 (c) [26]).

However, studies of void formation mechanisms and void formation sites mainly investigated by employing the uniaxial tensile testing with smooth (un-notched) specimens. It will be very helpful to our understanding of fracture toughness if the experiments could be performed using tensile type or bending type specimens with fatigue precracks. Greenfield et al. [27] reported that, for tensile fracture of Ti-5.25AI-5.5V-0.9Fe-0.5Cu, voids form mainly at the interface of primary a and the matrix when a specimen is treated with solution-aging into equiaxed a microstructure. In contrast, voids nucleate at the interface between the a grain boundary and the matrix when microstructure is an acicular a microstructure, whereas the interface between acicular a and the matrix can not be a void nucleation site. Moreover, Margolin et al. [28] studied the Ti-6AI-4V alloy and pointed out that voids also nucleate in primary a of equiaxed a microstructure, at the sub-grain boundary of the primary a or the matrix, at the interface between acicular a and the matrix or inside acicular a. For such cases of void formation at the twin-matrix interface or at the a-f3 interface, the mechanism could be that slip occurs in the a phase first, then a void nucleates because of interaction between the slip and the a-f3 interface. Nevertheless, a void generally nucleates at the a-f3 interface [29]. Nuclear formation sites of voids are determined by mechanical properties such as strength difference between a and f3 or martensite phase, and slip plane orientation between two phases. In other words, nuclear formation is greatly dependent on whether there is a Burgers orientation relationship between the two phases.

There are few detailed and systematic reports on the observation of crack initiation and propagation behavior with fatigue precracked or notched samples, which are presumed to be necessary for this evaluation. With equiaxed a microstructure, crack formation usually occurs at the primary a-f3 interface as an initiation site, where void nucleates and grows along the primary a-f3 interface, and subsequently coalesces with the main crack, leading to the propagation of the main crack along the a-f3 interface. However, Gysler et al. [14] reported that secondary cracks also occur inside the primary a in bimodal microstructure of Ti-6Al-4V alloy. Moreover, when the content of the primary a is less than 20% crack can also nucleate inside the matrix (finer acicular a + f3) surrounding the primary a, and the crack initiation inside matrix takes place as frequent as at primary a-f3 boundaries.

For acicular a microstructure, cracks generally occur at the a-f3 boundary too, and mainly propagate inside the grain; they also sometimes propagate along the original f3 grain boundaries. For equiaxed a microstructure, if it possesses the same strength as that of acicular a microstructure, the elongation is superior whereas fracture toughness is inferior to that of acicular a microstructure. This is because crack deflection in acicular a microstructure is more likely to take place than in equiaxed a microstructure. However, though this is effective for explaining the increase in crack propagation resistance, it can not be used directly to explain increasing crack growth initiation toughness. By way of explanation, Kishi et al. [30]


STEP 1

STEP2

Colony Block

~200-300~m decohesion ~2-4 Ilm

~~>~\ coalescence fracture surface

STEP3 ~ ____ Facet due to Fig. 9.22. Schematic crack growth process '-""''''''''......... colony cracking

~STEP~ -":"~~:'-J11lI~ crac

STEPf~ 1 crac_~racture surface during fracture toughness testing in acicular a Ti-6Al-4V alloy30).

Lenticular Coarse Alpha

Nucleation-Easy Growth-Difficult

Ductility Toughness

Plastic ~ zone~_ ~ ~ Crack~(J"'::,l:J

a Small Plastic Zone Tortuous Crack Path

Globular Coarse Alpha Nucleation-Difficult

Growth-Easy

Ductility a I

Plasti~ zone a

Large plastic zone

Toughness

_00 Crack 0

Non-Tortuous Crack Path

Fig. 9.23. Schematic representation of effect of coarse alpha shape on ductility and toughness. For a given volume fraction of coarse alpha, the total plastic zone is greater for globular alpha. However, for a propagating crack, the tortuosity factor for lenticular alpha overcomes this effect. Voids are shown as black regions5) .

pointed out that the stress concentration of main crack can be relaxed by microcrack occurrence caused by plate-shaped a-f3 interface debonding at the crack-tip inside the colony (first step); also, large shielding occurs because of microcrack growth (second step) and coalescence to the plate thickness direction (third step), as shown in Fig. 9.22. Figure 9.23 shows a schematic explanation of a key point on crack initiation and crack propagation in a case of equiaxed a or acicular a microstructure, proposed by Hirth et al. [5]. With this microstructure, deflection of the crack growth path occurs considerably with coarse original f3 grain. The decline of fatigue crack propagation rate can be attributed to the same mechanism as previously stated based on the report of Yoder et al. [15].

However, as Fig. 9.16 shows, even if the original f3 grain and colony or packets are coarsened, there are sub-colony formations in many colonies that have a

9.4 Embrittlement and Strengthening of Titanium Alloys by Hydrogen 203

-: 100 >. 00 ........

o ~ o CI)

",:: "'0 os CI) .... .0

§ ~ 0 .~., 5 .... .0 (.) os

rj rprt (j o. AC after soiutionizing " .. I1WQ after soiutionizing

.6-§ ~ ~

020 30 40 50 60 70 80 90 Area fraction of primary a, %

Fig. 9.24. Hydrogen embrittlement ratio vs. area fraction of primary a in Ti-6Al-4V alloy. rj: reduction ratio in crack initiation energy, rp: reduction ratio in crack propagation energy, rt: reduction ratio in total absorbed energy

deflecting effect because of great difference in direction. Therefore, as schematically shown in Fig. 9.23, sub-colony structure does not only introduce deflection to crack propagation path, but also relieves the stress concentration accompanying slip deformation. Moreover, it has also been reported that FCC structure boundary phase exists in a-f3 interface, which can be altered by varying heat-treating conditions [30, 31]. Future investigation should address effects of such microstructural features.

9.4 Embrittlement and Strengthening of Titanium Alloys by Hydrogen

9.4.1 Embrittlement

Titanium alloys readily absorb hydrogen to form FCC structured titaniumcompound, 6(TiH2), when the absorbed hydrogen exceeds the solid solution limit. This hydride is a very brittle phase which therefore may act as a crack initiation site. Even if the hydrogen content is within the solid solution limit, stress at the crack-tip rises and hydrogen concentrates to the crack-tip area by diffusion when a crack exists and stress is applied. Therefore, the hydrogen content at the crack-tip exceeds the solution limit and hydride precipitates, giving rise to possible hydrogen embrittlement [32]. Hydride is more likely to precipitate at the alf3 interface or in the a-phase first for two-phase a + f3 alloys because the hydrogen solution limit in the f3 phase is larger than that in a phase.

By electrolytic charging method, for equiaxed a microstructure Ti-6 . 4 alloy charged with excessive hydrogen over the solution limit, hydrogen embrittlement sensitivity decreases with. increasing primary a phase content, as shown in Fig. 9.24. A decreased hydrogen embrittlement tendency was found in material which was dissolved followed by water-cooling and aging (STA) treatment. This implies that hydrogen diffusion resistance is enhanced in the microstructure with fine and randomly-oriented a precipitation. Compared to equiaxed a microstructure, hydrogenation speed in acicular a microstructure is higher, whereas hydrogen


~>AS received

20 30 50

-6 0: 8.3X 10_7 mls /).:8.3XI0_s m/s 0: 8.3 X 10 mls

70

293 K

90 110 Average hydrogen content, ppm

Fig. 9.25. Relationship between fracture toughness (JIe) and average hydrogen content under various cross-head speed conditions (Ti-6A1-4V alloy).

embrittlement sensitivity of acicular a microstructure is not necessarily higher than that of equiaxed a microstructure [32].

On the other hand, as shown in Fig. 9.25, hydrogen softening occurs when hydrogen is charged within the solution limit (ca. 80 ppm) to Ti-6 . 4 alloy; this improves both toughness and ductility [33].

Fatigue behavior of titanium alloy in a corrosive environment of NaCI water solution was reportedly affected by hydrogen-related cracking [34]; that mechanism remains unclear.

9.4.2 Strengthening

Because hydrogenation and dehydrogenation of titanium alloy is reversible, degassing hydrogen can be easily achieved by heating within high-vacuum. Therefore, mechanical properties of titanium alloy can be recovered to the same level as that before hydrogenation [35]. When hydrogen is absorbed into titanium alloy, f3 transformation temperature (f3 transus) reduces; it gives rise to improvement in hot-working ability because hydrogen is a f3 phase stabilizing element. Moreover, powders can be easily obtained because the material becomes brittle with formation of hydrides. After the dehydrogenation process for a titanium alloy, which absorbs hydrogen and then deforms at high temperature or disintegrates, mechanical properties can be restored to the condition before hydrogenation. Such hydrogenation and dehydrogenation treatment is called the thermo-chemical treatment process (TCP), which can produce fine microstructure of titanium alloy. With this method, a near-net shape process can be applied to provide workability that was not available before. The TCP method has also been applied to refine microstructure of titanium alloy castings which grow easily into coarse acicular a microstructure.

Regarding TCP processes using hydrogen, there are many reports addressing: HVC (Hydrovac), CST (Constitutional Treatment), f3 Q-HDH (Hydrogenation and Dehydrogenation of Beta Quenched Materials), and HTH (High Temperature

"""' o:s 1200

~ 1100 '-'

blOOD r£ rIl 900 e ..... rIl

'0 "ii :;::

800

700

•

9.5 Strain Induced Transformation and Mechanical Properties 205

Ti-6AI-4V

0 • • 0 00 • 0

• 0 0: Equiaxed (l .: Acicular (l c: ~s received .: treated

4 6 8 10 12 14 16 18 Elongation, EI (%)

Fig. 9.26. Relationship between yield stress (Uy) and Elongation (EI) in thermochernically treated Ti-6AI-4 V alloy.

Hydrogenation). Titanium alloy microstructure can be refined by any of the processes; the strength, especially fatigue strength, is improved remarkably. Though it is believed that reduction in elongation is not prominent by TCP processing, the whole process is heated into the f3 range, and f3 grain growth or a grain growth are involved. Accordingly, elongation of the alloy may be decreased whereas the strength is increased. Based on this consideration, the temperature for the whole process is controlled below the f3 transus, as shown in Fig. 9.26 [36]. This enables an improvement in elongation, instead of reduction, whereas the strength is also concurrently improved. Fracture toughness is generally decreased considerably by any TCP.

If a working process is performed with the help of hydrogenation and dehydrogenation treatment and the recrystallization phenomenon is utilized, microstructure can be further refined to a fine equiaxed a microstructure with about 1 f1.m grain size.

Presently, a refining method by TCP with hydrogen is applied mainly for a + f3 type alloy, but this method is also presumed to be effective for f3 type alloys or titanium-based intermetallic compounds [37].

9.S Strain Induced Transformation and Mechanical Properties

For a + f3 and f3 type titanium alloys, when they are rapidly cooled after solution treatment, a large amount of f3 phase will remain in a metastable condition until reaching room temperature. If such a metastable f3 phase is strained, it will transform into martensite, w phase! or twin. This phenomenon is referred to as the so-called strain induced transformation, by which titanium alloys can be strengthened and toughened in the same way as that of steels. For instance, strength or toughness can be improved considerably by the retained f3 phase transformation into martensite in

1 It forms in a martensite-like transformation at quenching, and when retained f3 is heated (low temperature aging) or cold-worked (hexagonal). Generally, heat treatment is carried out to avoid formation of a w phase because it is a brittle phase.


a: STQ1173 R " ... : STA1173 R ~ V:STQI088Ri "

,,"

f./~ ~~

f~ Fig. 9.27. Relationship between fatigue crack growth rate (da/dN) and effective stress intensity factor range (AK.,tr) of solution treated (STQ 1173 R, STQ 1088 R) and solution treated and aged (STA 1173 R) Ti-6A1-4Valloy39l .

Ti-6 . 4 and Ti-6 . 2 . 4 . 6 alloy, or into twins in Ti-15 . 3 alloy [38]. This phenomenon is especially effective at lower temperatures.

Strain induced transformastion not only improves strength and toughness but also improves fatigue characteristics. Fatigue life of Ti-6 . 4 alloy under solution treated condition is longer than that under solution-aging treated condition, which is attributed to strain induced transformation of retained p phase into martensite. Moreover, for Ti-6·4 alloy as shown in Fig. 9.27, fatigue crack propagation resistance under solution treated condition (STQ10088R) is higher than that under a solutionaging treated condition (STAI173R), where strain induced transformation has also been observed [39].

It is well accepted that the martensite formed by strain induced transformation in titanium alloys possesses an orthorhombic (a") crystal structure. However, it has also been reported that the retained p phase may transform into a" + twinned p, face-centered orthorhombic martensite, face-centered cubic martensite (a'" [31]) or hexagonal martensite (a').

References

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2001.

9.5 Strain Induced Transformation and Mechanical Properties 207

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Strength and Toughness of Materials || Titanium Alloys