ce

Hog

rlstad

Univ

epted

line 2

of 250300 and 300350 jC, respectively. For aluminium

Surface & Coatings Technology 19injection and completely and rapidly fill an internally cooled

mould, typically within the order of milliseconds [13]. The

high melt velocity during injection and the continuos

internal cooling of the tool during the process allow pro-

duction of thin-walled and complex near net-shaped cast

products at high manufacturing rates, typically of the order

of 100 castings per hour. When the casting has solidified,

whereas those for copper alloys are about 110 m/s and

970 jC. Hot work tool steels, such as AISI H11, H13, H20,H21, or H22, are commonly used as die materials.

Thermal fatigue cracking (often named heat checking) is

one of the most important life-limiting tool failure mecha-

nisms in aluminium and brass die casting [13]. Other

considerable failure modes that limit the life and perfor-application of pressure to flow with high velocity duringmagnesium, and copper based alloys is forced by the alloys, the entrance velocity of the melt is usually 2060 m/

s and the melt temperature is approximately 700 jC,In die casting, molten metal of e.g. aluminium, zinc,resistance against thermal fatigue is not fully explored.

In this work, a selection of hot work tool steel grades was surface modified and experimentally evaluated in a dedicated thermal fatigue

simulation test. The surface modifications included boriding, nitriding, Toyota diffusion (CrC), and physical vapour deposition (PVD) of

coatings (CrC, CrN and TiAlN), both as single-layers and deposited after nitriding (duplex treatment). Untreated specimens of each tool steel

grade were used as references. The test is based on cyclic induction heating and internal cooling of hollow cylindrical test rods. The surface

strain is continuously recorded through a non-contact laser speckle technique.

Generally, all surface treatments decreased the resistance against surface cracking as compared to the reference materials. The reason is

that the engineering processes influence negatively on the mechanical properties of the tool materials. Of the processes evaluated, duplex

treatment was the least destructive. It gave a lower crack density than the reference steel, but the diffusion layer is more susceptible to crack

propagation. In addition, the single-layered CrN coating showed almost comparable thermal fatigue cracking resistance as the reference

material. Finally, the resistance against thermal crack propagation of surface engineered tool steels is primarily determined by the mechanical

properties of the substrate material.

D 2004 Published by Elsevier B.V.

Keywords: Thermal fatigue; Heat checking; Surface strain; Surface engineering; Hot work tool steel; Die casting

1. Introduction spraying. Prior to casting aluminium and copper alloys, the

die is normally preheated to a temperature within the rangecasting. Surface engineered materials are today successfully appAbstract

Thermal fatigue cracking is an important life-limiting failure mechanism in die casting tools. It is observed as a network of fine cracks on

the surfaces exposed to thermal cycling. The crack network degrades the surface quality of the tool and, consequently, the surface of the

lied to improve the erosion and corrosion resistance. However, theirThermal fatigue cracking of surfa

Anders Perssona,*, StureaDepartment of Materials Engineering, Ka

bThe Angstrom Laboratory, Uppsala

Received 22 October 2003; acc

Available onthe die is opened and the casting ejected. Subsequently, the

tool surfaces may be externally cooled and lubricated by

0257-8972/$ - see front matter D 2004 Published by Elsevier B.V.

doi:10.1016/j.surfcoat.2004.04.053

* Corresponding author. Tel.: +46-54-700-1821; fax: +46-54-700-

1449.

E-mail address: Anders.Persson@kau.se (A. Persson).engineered hot work tool steels

markb, Jens Bergstroma

University, SE-651 88 Karlstad, Sweden

ersity, SE-751 21 Uppsala, Sweden

in revised form 8 April 2004

June 2004

www.elsevier.com/locate/surfcoat

1 (2005) 216227mance are for example gross fracture, erosion, corrosion and

local adherence of the casting alloy (soldering). Thermal

fatigue cracking results from the rapid alternations in die

surface temperature, which may induce stresses high enough

to impose an increment of plastic strain in the tool surface

during each casting cycle. Surface cracks develop generally

within a few thousand cycles, or even earlier, and are,

ed materials [1420], and the mechanisms behind thermal

fatigue cracking of surface engineered materials are not yet

beneath the coated surface at which the hardness was 50 HV

25 gf higher than the substrate hardness. For the specimens

treated by boriding, the depth is given excluding the

outmost compound layer. The PVD coating thicknesses

were determined by light optical microscopy (LOM) on

polished cross-sections.

Relative hardness of the surface treatments and coatings

was assessed by micro-Vickers indentations on polished

cross-sections and directly on the top of the surface of the

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227 217fully explored.

In this study, various combinations of surface engineered

hot work tool steel systems were evaluated in thermal

fatigue, including diffusive thermochemical surface treat-

ments (boriding, nitriding, and Toyota diffusion (CrC)), and

physically vapour deposited (PVD) coatings (CrC, CrN and

TiAlN), used as single-layers or on top of a nitrided layer

(duplex treatment). Untreated specimens of each tool steel

grade were used as references. The test is based on cyclic

induction heating and internal cooling of hollow cylindrical

test rods [21]. Additionally, the surface strain is continu-

ously recorded during thermal cycling through a non-con-

tact laser speckle technique.

2. Experimental

2.1. Materials and characterisation

Three Uddeholm hot work tool steels, Orvar Supreme

(premium AISI H13 grade), QRO 90 Supreme and Hotvar,

were used as test material. Their nominal chemical compo-

sitions, and hardening (austenitizing, followed by air

quenching) and tempering treatment are given in Table 1.

The austenitizing treatment gives a nominal austenite grain

size of about ASTM 10 for Orvar, and about ASTM 9 for

QRO 90 and Hotvar. All heat treatments resulted in different

microstructures of tempered martensite. The heat treatments

were followed by grinding to a surface roughness (Ra) of

0.38F 0.05 Am, as obtained by optical surface profilometry.The Orvar specimens were further ground and polished with

1-Am diamond paste in a last step to a surface roughness(Ra) of 20F 14 nm. Test specimens of each tool steel gradeof these conditions were used as references, denoted Orvar,consequently, formed in the low-cycle fatigue range

( < 103104 cycles) [4,5]. Oxidation and creep may signif-

icantly contribute to cracking [48]. Thermal fatigue dam-

age is often observed as a network of fine cracks on the tool

surface, and the cracks penetrate usually only a limited

surface layer. The crack network degrades the surface

quality of the tool, and since it is replicated on the surface

of the cast products, thermal fatigue cracking may ulti-

mately cause higher production costs through rejection of

castings.

Surface engineerings of borides, carbides, and nitrides

are ceramic compounds with high hardness, wear resis-

tance, and chemical inertness. Surface engineering has

proven successful to improve the erosion and corrosion

resistance as well as to reduce soldering of tool materials

in die casting applications and laboratory tests [915].

However, it has been shown that surface engineered

materials may show comparable, increased or decreased

resistance against thermal cracking as compared to untreat-QRO 90, and Hotvar, respectively.Prior to surface engineering, the specimens were ground

and polished as the Orvar specimens above (Ra 20F 14 nm).The QRO 90 specimens were surface treated by

boriding (denoted Q-B), Toyota diffusion to generate

CrC (denoted Q-TDP CrC), or plasma nitriding to pro-

duce a diffusion zone without any iron nitride compound

layer. The first letter in the denotations of the surface

engineered materials denotes always the substrate material

and the others the engineering process. The boriding

process was followed by hardening and tempering (at

1030 jC and 2 2 h at 625 jC, respectively), while theTDP treatment was followed by tempering 2 h at 625 jCand 2 h at 600 jC. All plasma nitrided QRO 90 speci-mens were duplex-treated with a PVD CrN coating

(denoted Q-DCrN) on top of the nitrided layer. Single-

layered PVD CrN coatings (denoted Q-CrN) were also

applied on the QRO 90 specimens. The PVD CrN coat-

ings were produced in a multi-arc process.

The Orvar and Hotvar specimens were treated by gas

nitriding (denoted O-GN, H-GN) or plasma nitriding. All

plasma nitrided specimens of these two tool steel grades

were duplex-treated with a PVD CrC (denoted O-DCrC, H-

DCrC) or a multi-layered PVD TiAlN coating (denoted O-

DTiAlN, H-DTiAlN) on top of the nitrided layer. The PVD

CrC coatings were produced by an ion plating process, and

the multi-layered PVD TiAlN coatings were produced by an

arc-evaporation process. The 13 treatments resulted in

different surface characteristics, see Table 2, where also

the deposition time and temperature during the boriding,

TDP, and nitriding and/or PVD coating treatments are

included.

Substrate hardness was assessed by Vickers indentations

on polished cross-sections, using a load of 30 kg. For the

specimens treated by boriding, TDP CrC and nitriding, the

depth of the diffusion zone profiles was assessed by micro-

Vickers indentations on polished cross-sections, using a load

of 25 g. The diffusion depth was defined as the depth

Table 1

Nominal chemical compositions in wt.% and austenitizing (A) and

tempering (T) treatments of the tool steels

Steel grade C Si Mn Cr Mo V A

[jC/min]T

[jC/h]

QRO 90 Supreme 0.38 0.30 0.75 2.6 2.25 0.9 1020/30 625/2 2Orvar Supreme 0.39 1.0 0.4 5.2 1.4 0.9 1025/30 600/2 2Hotvar 0.55 1.0 0.75 2.6 2.25 0.85 1050/30 575/2 2coatings, respectively, using a load of 25 g. For the speci-

ure during the boriding, TDP, and nitriding and/or PVD coating treatments

n

m]

Coating

hardness

[HV 25 gf]

Coating

thickness

[Am]

Time and temperature

during surface treatment

[h/jC], [h/jC]a

1740F 100 f 25/f 8501970F 70 6/10302000F 100 6.1F 0.1 /300400

5

0

0

0

0

ble va

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227218mens treated by boriding, the surface hardness was mea-

sured below the compound layer.

2.2. Thermal fatigue testing

The test equipment is based on cyclic induction heat-

ing and internal cooling of hollow cylindrical test rods

with a diameter of 10 mm, a length of 80 mm, and

having a 3-mm axial hole for internal cooling. Surface

strain measurements through a non-contact laser speckle

technique make it possible to calculate the strains induced

in the specimen surface during thermal cycling. An

induction unit (25 kW, 3 MHz) heats the surface layer

of approximately 20 mm of the middle of the test rod.

Continuous cooling is performed by circulating silicon oil

(flow rate c 2.5 l/min) of 60 jC through the specimen,and also externally with either argon (forced convection),

Table 2

Characteristics of the test materials, including deposition time and temperat

Denotation Substrate

hardness

[HV 30 kgf]

Nitriding

hardness

[HV 25 gf]

Diffusio

depth [A

QRO 90 507F 5 Q-B 520F 5 30F 2Q-TDP CrC 522F 5 30F 2Q-CrN 495F 5 Q-DCrN 507F 5 915F 15 160F 3Orvar 486F 5 O-GN 482F 5 1030F 30 250F 1O-DCrC 485F 5 955F 30 120F 1O-DTiAlN 476F 5 975F 30 115F 1Hotvar 640F 10 H-GN 604F 5 955F 20 295F 5H-DCrC 606F 5 845F 25 175F 1H-DTiAlN 596F 5 910F 25 170F 1

a Single time/temperature values refer to the single treatments, and dou

coating).which also decreases oxidation during the thermal cy-

cling, or air (natural convection).

Argon was used as cooling medium because the tools

are exposed to an environment with reduced oxygen

content during actual die casting. The oxygen in the die

cavity is partly consumed through oxidation of tool mate-

rial and casting alloy. For the tests in argon, the specimen

was contained in a glass tube and argon was flowing

between the specimen and the glass tube wall. Unfortu-

nately, this glass tube makes it impossible to obtain surface

strain recordings by the laser speckle technique. However,

the surface strains could be recorded during the first 100

cycles in a separate test series performed in air. More

information about this test method is presented elsewhere

[21].

Two temperature cycles were used to simulate aluminium

and brass die casting temperature conditions, see Table 3, in

the following, designated according to their maximum

temperatures, 700 and 850 jC, respectively. The lattertemperature corresponds to the maximum tool surfacetemperature during actual brass die casting [22]. The ther-

mal cycles included a steep ramp to the maximum temper-

ature and subsequent cooling to the minimum temperature,

with or without a short holding time ( < 0.1 s) at the

maximum temperature.

Thermal cycling by induction heating above the Curie

transition temperature of Fe (768 jC) does not ideallygenerate the temperature profile representative for die

casting of, e.g. brass [22]. Thus, the 850 jC heat cycleis not as representative for die casting as that to 700 jC,and only the reference specimens of QRO 90 and those

treated by boriding and TDP were tested to 850 jC. In thefollowing, only the surface strains for the 700 jC tests areconsidered.

Prior to testing, the specimens were pre-oxidised in

order to get a thin oxide layer, which facilitates the

pyrometer temperature control during heating. The refer-

2060F 100 4.5F 0.2 15/480, /300400

48/510

1710F 55 5.5F 0.4 /f 480, /4505002285F 170 5.5F 0.7 /f 480, /450500

48/510

1690F 45 3.8F 0.1 /f 480, /4505002275F 200 2.9F 0.1 /f 480, /450500

lues refers to those for the duplex treatments (nitriding followed by PVDence specimens were pre-oxidised by electrochemical

oxidation in a NaOH-solution at 70 jC for about 5min, followed by 1-h heat treatment at 200 jC in air[21]. All surface engineered specimens were pre-oxidised

0.5 h at 600 jC. This temperature, which is comparableto the tool steel tempering temperature, and the relatively

short time were selected not to affect the mechanical

properties of the tool material. A previous study showed

that 20000 thermal cycles to 600 jC did not affect themechanical properties of these tool steels [21]. A K-type

Chromel-Alumel thermocouple with thin wires (/ 0.13

Table 3

Thermal cycles used in the thermal fatigue tests

Maximum

temperature

[jC]

Minimum

temperature

[jC]

Heating

time [s]

Total cycle

time [s]

External

cooling

700 170 0.30.4 14.314.4 Argon or air

850 170 2.22.5 26.226.5 Argon or air

2.3. Evaluation

During the heat cycling, the surface strain is continuously

obtained by the laser speckle technique from the change in the

specimendimensions, and it is represented as surface strain vs.

temperature. Any thermal fatigue damage of the specimen

surface was revealed using scanning electron microscopy

(SEM). It was further characterised with respect to crack

growth (crack length vs. number of cycles) and crack density

(number of cracks per unit of length) by crack lengthmeasure-

ments on polished axial cross-sections by LOM. All evalua-

tions of cracks are based on cracks larger than about 5 Am,detected along two lines, each of 8-mm length.

Hardness vs. depth profiles and relative coating hardness

after exposure to the heat cycling were assessed by micro-

Vickers indentations on polished cross-sections and directly

on the top of the surface of the coatings, respectively, using

a load of 25 g.

3. Results

3.1. Recorded surface strain during thermal cycling

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227 219mm) was spot welded to the specimen to measure the

surface temperature during testing. The thin wires enable

rapid response of any change in temperature. Finally, to

obtain a good speckle pattern for the surface strain

measurements, an area of approximately 10 10 mmlocated in the middle of the test rods was roughened by

a 1000 mesh abrasive paper. However, for the multi-

layered PVD coatings of TiAlN, it was not possible to

obtain a good speckle pattern and, therefore, no surface

strain recordings for this material are presented.

After about four initial thermal cycles, the starting

Fig. 2. Typical surface cracking after 10000 cycles to 700 jC in argon (Q-B, SEM).

Fig. 1. Example of surface strain recordings (tangential strain) during

thermal cycling to 700 jC in air at equilibrium temperature conditions. (a)Reference material (QRO 90). (b) Surface engineered QRO 90 systems.Fig. 3. Polished cross-section revealing typical thermal cracks (Q-B, LOM).

(a) 5000 cycles to 700 jC in argon; (b) 500 cycles to 850 jC in argon.specimen temperature of 60 jC ( = temperature of coolingoil) is raised to 170 jC and cyclic equilibrium tempera-ture conditions are established, cp. Table 3. Thereafter, the

surface strain recordings are obtained during almost

identical thermal cycling. When equilibrium temperature

conditions are established, the surface strain typically

increases with the surface temperature, followed by an

almost constant or slightly increasing strain level during

the first part of the cooling, wherafter it decreases with

temperature, see Fig. 1. After completed cooling, each

surface strain recording forms a closed loop with no

remarkable residual strain at the minimum temperature.

For some materials, the scatter in residual strain at the

minimum temperature was considerable. Note that the

tangential strain eB and axial strain ez is almost identical(see Fig. 1a). Note also that the surface strain recordings

for the surface engineered materials are comparable to

those of the reference material (see Fig. 1b). No remark-

able differences among the tested materials were observed

during the initial 100 heat cycles recorded.

3.2. Surface cracking

Thermal fatigue cracking of a surface engineered tool

steel after thermal cycling to 700 jC is exemplified in Fig.2. Polished cross-sections revealed that the crack path was

strongly dependent on the maximum temperature during

each cycle. Essentially two different types of crack paths

were distinguished. Relatively straight cracks dominated

c) afte

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227220Fig. 4. Maximum crack length (a), mean crack length (b), and crack density (value of two to four specimens of each material, and the error bars indicate the mr cycling to 700 and 850 jC, respectively, in argon. Each pile is the average

aximum and minimum recording.

after heat cycling to 700 jC, cp. Fig. 3a. Both straight andbranched cracks (Fig. 3b) were observed after thermal

cycling to 850 jC.

3.2.1. Crack length and crack density of surface engineered

QRO 90

The crack length was strongly dependent on the number

of cycles and the maximum temperature during each cycle,

see Fig. 4. In addition, the crack density of each material

was almost constant or slightly increasing within the num-

ber of cycles tested, for the two maximum cycle temper-

atures. It is also seen that the crack length and density of

cracks differ significantly among the surface engineered

QRO 90 substrates.

Generally, the boriding and TPD surface treatments show

a tendency to increase the crack length and reduce the

density of cracks as compared to the reference QRO 90

steel, especially after thermal cycling to 850 jC, see Fig. 4.In a relatively early stage of heat cycling to 700 jC, thesingle-layered and the duplex-treated CrN coating shows a

tendency to increase the crack length as compared to the

reference material. Later on, the crack characteristics of the

single-layered CrN coating proved almost comparable to the

reference material, whereas the crack length in the duplex-

treated coating still is longer. In general, it is obvious that

duplex treatment reduces the crack density, as compared to

both the single-layered CrN coating and the reference QRO

90 steel.

(c) aft

C whi

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227 221Fig. 5. Maximum crack length (a), mean crack length (b), and crack density

specimens of each material, except that for O-TiAlN 10000 cycles at 700 j

minimum recording.er cycling to 700 jC in argon. Each pile is the average value of two to fourch is based on one specimen, and the error bars indicate the maximum and

3.2.2. Crack length and crack density of surface engineered

Orvar and Hotvar

The crack length was strongly dependent on the number

of cycles up to 700 jC, and the crack length and density ofcracks varied significantly between the surface engineered

Orvar and Hotvar substrates, see Fig. 5. Again, the crack

density of each material was almost constant or slightly

increasing with the number of cycles.

In general, the gas nitriding treatment increases the

maximum crack length and the crack density as compared

to the reference steels, see Fig. 5. It also decreases the mean

crack length of Orvar and increases that of the Hotvar steel.

In a relatively early stage, all duplex-treated PVD coatings,

except the duplex-treated CrC coating on top of Hotvar,

show comparable crack length or a slight tendency to

increase the crack length as compared to the reference Orvar

and Hotvar steels. Thereafter, there is a tendency that all

duplex treatments increase the crack length as compared to

the reference materials. Generally, all duplex-treated PVD

coatings on top of the Orvar and Hotvar substrates proved to

reduce the crack density as compared to the references.

3.3. Hardness after thermal fatigue

Typically, for the 700 jC experiments the hardness of thesuperficial 0.31-mm surface layer decreased during the

thermal cycling for all test materials, see Fig. 6ac. Beneath

this depth, the material appears unaffected by the cyclic heat

exposures. Evidently, the hardness of the borided, TDP,

plasma nitrided, and gas nitrided layers is considerably

reduced after the heat cycling. (The profiles do not reveal

the hardness values of the PVD coatings.) The hardness

profile for the single-layered CrN coating (not shown in Fig.

6a) almost coincides with that of the reference steel. Below

the surface engineered layer, the hardness values are almost

identical to those of the references, except for the Hotvar

systems, which show somewhat lower hardness values than

the reference steel, see Fig. 6ac. The hardness values of

the PVD coatings were also significantly decreased after the

thermal cycling. After 10000 cycles to 700 jC, the hardnessof the PVD coatings of CrN, CrC, and TiAlN was about

1300F 100, 1050F 100, and 1700F 200, respectively. Thehardness levels and the thickness of the softened surface

layer were strongly dependent on the maximum temperature

during each cycle, see Fig. 6, and also on the substrate

material as observed after the 700 jC tests, see Fig. 6ac.The 850 jC experiments resulted in a dramatic hardness

reduction throughout the whole specimen for all test materi-

als, see Fig. 6d. Again, it is seen that the hardness reduction

of the borided and TDP layers is considerably, and that the

hardness values below the surface engineered layer almost

are identical those of the reference steel.

ms, (

ls are included.

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227222Fig. 6. Surface hardness vs. depth for the tested materials. (a) QRO 90 syste

argon. (d) QRO 90 systems after 1000 cycles to 850 jC. The reference stee

b) Orvar systems, and (c) Hotvar systems after 10 000 cycles to 700 jC in

4. Discussion

4.1. Surface layer conditions during thermal cycling

Induction heating of steel using a frequency of 3 MHz

give rise to very fast heating of only a thin surface layer, of

the order of 10 Am (skin-effect). When applying thismethod to the test rods, the thermal expansion of the surface

layer material is retained by the cooler bulk material, cp.

Fig. 1. During the initial cooling phase, the surface contracts

but the material below still expands due to heat conduction

and, thereby, the decrease in surface strain with temperature

is delayed, see Fig. 1.

4.1.1. Stresses in the surface layer

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227 223The fact that the thermal strain of the surface layer is

constrained during the thermal cycling exposes the surface

layer material to cyclic stresses. The hypothetical strains

corresponding to these stresses are defined as mechanical

strains emech. Crack nucleation and growth is determined byfluctuations in emech. Similarly, thermal strains eth defined asstrains without any constraint are defined from the thermal

expansion coefficient of the material a(T) and the minimumsurface temperatures Tmin as:

ethT aTT Tmin 1Thus, emech at any part of the thermal cycle is possible to

deduce from the corresponding values of the surface tem-

perature cycle, thermal expansion coefficient of the material,

and the measured surface strain etot according to:

emechT etotT ethT 2Note that the surface layer conditions above are estimat-

ed without any knowledge of the temperature profile below

the surface.

From the surface strain recordings such as those of Fig.

1, and using Eqs. (1) and (2), the three types of surface

strains (i.e. etot, eth, and emech) during heat cycling can bededuced, see Fig. 7. Similarly, these calculations give the

Fig. 7. Example of surface strain response during thermal cycling to 700 jC

in air at equilibrium temperature conditions (reference material QRO 90).minimum mechanical surface strain of each cycle, which for

the reference materials (also the substrate material) was

found to be emech (T= 700 jC)c 0.57F 0.06% for allequilibrium temperature cycles up to 700 jC. It coincideswith the maximum surface temperature, cp. Fig. 7.

The fact that the engineered surface layers do not

significantly alter the appearance of the recorded surface

strain as compared to those of the references (cp. Fig. 1)

indicates that the surface strain during heat cycling primarily

is controlled by the deformation of the substrate material

(i.e. tool steel). Thus, the engineered layers are obviously

too thin to have any notable effect on the temperature

distribution below the surface.

The thermally induced mechanical strains (such as those

of Fig. 7) will introduce compressive and tensile stresses in

the surface layer. Although the magnitude of the tensile

stresses imposed during cooling does not exceed the initial

(true) yield stress of these tool materials during thermal

cycling, numerous thermal cracks nucleate and propagate

within the low-cycle fatigue range, cp. Figs. 4 and 5 and

Ref. [21]. The explanation is that presence of local stress

concentrators may cause the tensile stresses to exceed the

tool steel yield stress during these conditions. In addition,

the gradual softening of the tool surface material with the

number of heat exposures considerably degrades the initial

tool steel yield stress values, cp. Fig. 6. Consequently, the

surface material locally will be exposed to cyclic stresses

that cause accumulation of plastic strains after a certain

number of thermal cycles up to 600850 jC.

4.2. Thermal fatigue cracking

This study clearly demonstrates that surface engineering

of hot work tool steels generally has a tendency to reduce

the resistance against thermal fatigue cracking. This is

obvious from the characterisation of the maximum crack

length (cp. Figs. 4a and 5a) and the mean crack length (cp.

Figs. 4b and 5b) after thermal cycling. The only positive

effect observed in this study is that duplex treatment reduces

the crack density, cp. Figs. 4c and 5c.

During the initial 5000 and 500 cycles to 700 and 850

jC, respectively, almost all thermal fatigue cracks initiateand propagate into the materials, cp. Figs. 4 and 5.

Thereafter, some cracks grow deeper into the material

and, thereby, both the maximum and mean crack lengths

increase with the number of cycles, whereas the crack

density saturates at a relatively low number of cycles [4].

The crack lengths and crack density differ significantly

within the whole range of materials tested. Since areas

along each thermal crack are locally stress relieved, the

crack growth of adjacent cracks is retarded as the larger

ones propagate into the material [23]. Consequently, the

stress at potential crack nucleation sites (e.g. local stress

concentrators due to topographical features and material

inhomogeneities [4,18,19,24]) close to each crack is alsoreduced.

4.2.1. Reference steels

The thermal crack advancement in the reference steels

during heat cycling to 700 jC has a tendency to be sup-pressed by a higher initial tool steel hardness [21], cp. Figs.

4 and 5. A higher hardness reduces the accumulation of

plastic strains in the surface layer. Note that any initial

ranking in hardness among the reference steels is main-

tained, cp. Fig. 8. The favourable effect from the higher

hardness is further elucidated by the observation that the

harder Hotvar reference has a lower crack density than the

softer QRO 90 steel, see Figs. 4c and 5c. However, the

Orvar and Hotvar references have approximately the same

crack density levels, see Fig. 5c, in spite of the considerable

differences in hardness, see Fig. 8. This is most likely due to

the fact that the surface roughness prior to testing of Orvar

was much smoother than that of Hotvar, and also that the

appear virtually unaffected by the PVD process itself, cp.

Table 2.

4.2.3. General cracking characteristics of diffusion layers

One striking observation for all diffusion layers, i.e.

the borided, TDP, plasma nitrided, and gas nitrided

diffusion layers, is the following. After thermal crack

nucleation during heat cycling to 700 and 850 jC, somecracks propagate relatively rapidly through the depth of

the diffusion layers, wherafter they meet the substrate

material and advance further, cp. Figs. 4 and 5. Note that,

of these layers, only the borided and TDP layers were

tested to 850 jC. Most diffusion layers have a tendencyto increase the maximum crack length as well as the

mean crack length, as compared to the references, cp.

Figs. 4 and 5.

The hard and brittle nature of these diffusion layers (cp.

Table 2) makes the surface layer material susceptible to

thermal crack propagation by promoting rapid thermal crack

advancement through their thickness. The propagation is

retarded when the cracks encounter a material with suffi-

ciently high toughness [1820], cp. Figs. 4 and 5. This is

supported by the observation that there is a correlation

between the maximum crack length and the thickness of

the diffusion layer [19], see Fig. 9. The plasma nitrided and

gas nitrided materials, which are those materials with the

cycling constructed using Table 2 and Figs. 4a and 5a. (a) 5000 and 500

cycles to 700 and 850 jC, respectively. (b) 10000 and 1000 cycles to 700and 850 jC, respectively. The solid lines are included to visualise anycorrelation.

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227224deeper cracks in Orvar facilitate stress relief of the surface

layer.

Thermal cycling to 850 jC accelerates the thermalcrack nucleation and advances and promotes crack nucle-

ation at numerous sites as compared to heat cycling to

700 jC, cp. Fig. 4. The higher maximum cycle temper-ature facilitates rapid softening of the surface layer, and

increases the plastic strain accumulation during each

cycle, cp. Fig. 6a and c. Note that only the QRO 90

steel was tested to 850 jC.

4.2.2. PVD coating

The thin hard single-layered CrN coating on top of the

tool steel does not has any substantial effect on the

thermal crack nucleation and propagation, as shown by

the almost comparable cracking characteristics as the

reference steel, see Fig. 4. This indicates that the PVD

coating/substrate interface and the differences in proper-

ties among the materials do not have any practical effect

on thermal crack nucleation and propagation into the

underlying material [1820], and that the stress distribu-

tion in the substrate material not is significantly disturbed.

In addition, the mechanical properties of the tool steel

Fig. 8. Hardness vs. depth for the reference tool steels after 10000 cycles to

700 jC; constructed from Figs. 6ac.Fig. 9. Maximum crack length vs. thickness of diffusion layer after thermal

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227 225thickest diffusion layers, show the longest maximum crack

lengths.

4.2.4. Gas nitriding

The gas nitriding treatment gave some of the most

destructive surface layers of the processes evaluated by

thermal cycling to 700 jC, since it increased both themaximum crack length and the crack density as compared

to the two reference steels, see Fig. 5. Numerous of surface

cracks are formed during the initial 5000 heat cycles to 700

jC, although the presence of cracks should relieve the localstresses, cp. Fig. 5. This indicates that gas nitrided layers

make the surface material susceptible to both thermal crack

nucleation and propagation. Note that gas nitrided layers

were applied on Hotvar and Orvar substrates only.

When the cracks have grown through the gas nitriding

diffusion layer, their advancement is more effectively re-

tarded by the initially harder Hotvar substrate than that of

the softer Orvar steel (cp. Figs. 5a and 9), even though the

diffusion layer is somewhat thicker, cp. Table 2. This

reflects again the crucial favourable crack arresting effect

of high tool steel hardness.

4.2.5. Boriding and TDP

The high deposition temperature of the boriding and TDP

processes seems to deteriorate the mechanical properties of

the QRO 90 substrate more than the plasma nitriding and the

PVD processes do, even if this is not reflected by the

substrate hardness numbers, cp. Table 2. This is supported

by the observation that the maximum crack length in the

borided and the TDP treated QRO 90 tool steel are well

beyond the diffusion depths already after relatively few heat

cycles to 700 jC, cp. Figs. 4a and 9. The crack advancementthrough the substrate material appears facilitated, and after

10000 heat cycles to 700 jC, the maximum crack length inthose materials is nearly as long as that in the duplex-treated

steel, in spite of the fact that their diffusion depths are much

thinner, cp. Fig. 9.

In the boriding process, the steel substrate is kept in

contact with a reactive boron containing compound at high

temperature [25,26]. Initially, an iron boride layer is formed

at the surface and grow subsequently in thickness due to

diffusion of boron into the substrate. The diffusion of boron

into the steel substrate drags along elements that are not

soluble in the boride layer (e.g. carbon and silicon) into the

substrate, and causes an enrichment of these elements in a

zone below the boride layer. In the TDP process, the

substrate material is kept in contact with the carbide forming

element chromium at high temperature [25,27]. Initially,

carbon in the steel substrate combines with this element and

forms a carbide layer on the surface. Thereafter, the carbide

layer grows thicker due to further reactions at the carbide

layer interface as a result of diffusion of carbon from the

substrate to the surface. During the process, the carbide

forming element chromium diffuses into the substrate andforms a solid solution layer beneath the carbide layer.Consequently, the boriding and TDP processes alter the

chemical composition of the substrate material below the

engineered surface layers and, therefore, also the mechan-

ical properties of the substrate surface material. In addition,

the bulk of the substrate material will also be affected by the

high treatment temperatures and subsequent hardening/tem-

pering conditions.

Even though the boriding and TDP processes are detri-

mental to the crack arresting properties of the substrate

material, they reduce the density of thermal cracks as

compared to the reference steel during thermal cycling to

700 jC, cp. Fig. 4. This effect may be a associated with thelocal stress relief of the surface layer provided by the

relatively rapid development of thermal cracks, and also

with the higher hot hardness (or hot yield strength) of the

layers, which reduces the plastic strain during thermal

cycling.

The more rapid development of long cracks during

thermal cycling to 850 jC than to 700 jC relieves the localstresses in surface layer and obstructs relatively early further

crack nucleation, cp. Figs. 4a,b, and 9. This is supported by

the observation that there is a relatively small difference in

crack density after the tests to 700 and 850 jC, respectively,cp. Fig. 4c. In addition, the crack branching associated with

the considerably more rapid crack growth during heat

cycling to 850 jC than to 700 jC (cp. Fig. 3) is probablypromoted by intergranular precipitations (e.g. carbides

[6,28], nitrides [4], and borides) generated in the surface

layer material during the diffusion process, heat treatment,

and temperature cycling.

4.2.6. Duplex treatment

It was clearly demonstrated that duplex treatment gener-

ally reduces the density of thermal cracks as compared to

both the reference steels and the single-layered CrN coating,

see Figs. 4 and 5. However, the plasma nitrided diffusion

layers are detrimental to the thermal crack propagation

resistance. Most of the duplex treatments have a tendency

to increase the crack length as compared to the references.

The PVD coating on top of the plasma nitrided layer has no

additional effect on the thermal crack nucleation and prop-

agation into the underlying material as mentioned above.

This indicates that the plasma nitriding process prior to

coating obviously has a dominant positive effect to inhibit

crack initiation, even though it increases the sensitivity to

crack advancement through the surface layer material.

Simplified, the higher hot hardness (or hot yield strength)

of the surface layer material, as compared to that of the

substrate material (cp. Table 2), reduces the plastic surface

strain during thermal cycling and, consequently, the driving

force for crack initiation [16,20].

Again, the crack advancement below the diffusion depth

has a tendency to be more effectively hindered by the

initially harder Hotvar substrate than that of the softer Orvar

steel, cp. Figs. 5a and 9. However, the QRO 90 substrate hasa tendency to facilitate crack arrest even better than the

Hotvar substrate, in spite of its lower initial hardness, cp.

(Figs. 4a, 5a, and 9). The reason is that the hardness at the

crack tip after 10000 cycles at 700 jC is higher for the QRO90 than for the Hotvar substrate, cp. Figs. 6a,c, and 10.

4.3. Softening

Thermal cycling to 700 and 850 jC, respectively, causevery pronounced softening of the surface layer of all surface

engineered materials and reference steels, see Fig. 6. After

10000 heat cycles to 700 jC, the initial hardness of thesurface engineered layers is reduced approximately within

the range of 2045%, cp. Fig. 6ac. For the borided and

TDP layers, 1000 cycles to 850 jC correspond almost to the

Naturally, the softening of the material below the surface

engineered layer during the thermal cycling is determined

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227226same hardness reduction as 10000 cycles to 700 jC, cp. Fig.6a and d. The multi-layered TiAlN coating showed the least

softening of the three various PVD coating materials tested,

and the TDP CrC layer showed less hardness reduction than

the boriding layer, cp. Fig. 6. This indicates a higher thermal

stability of these two ceramic compounds. Even though the

hardness reduction of the borided and TDP layers and PVD

coatings was considerable, the initial ranking among the

various layers is unaffected by the thermal cycling. For the

plasma and gas nitrided layers, the thermal stability during

the thermal cycling appears to be influenced by the substrate

material tempering resistance, decreasing according to the

following order: QRO 90, Hotvar, and Orvar, cp. Fig. 6ac.

Note that gas nitrided layers were applied on Hotvar and

Orvar substrates only.

It has been shown that the isothermal annealing above

the deposition temperature of thin hard arc-evaporated PVD

coatings of CrN [29] and TiCxN1 x [30] results in signif-icant relaxation of the intrinsic compressive residual stresses

typically present in as-deposited PVD films. Thus, it is

expected that any advantage from the residual compressive

stresses in the as-deposited single-layered or duplex-treated

PVD coatings is removed, since the temperature during each

thermal cycle reach well above the deposition temperature

levels during PVD coating, cp. Table 2. The observation that

the coating hardness values are considerably reduced after

10000 cycles to 700 jC and the fact that intrinsic compres-

Fig. 10. Example of hardness vs. depth for duplex-treated materials after10000 cycles to 700 jC; constructed from Figs. 6ac.by the properties of the substrate material, cp. Fig. 6.

However, if the initial substrate hardness levels for the

surface engineered Hotvar systems had been equal to the

hardness of the reference steel (see Table 2), there had most

likely been a complete overlap between these hardness

profiles and the one of the reference steel, cp. Fig. 6c.

4.4. Surface engineered systems for die casting applications

In selection of a surface engineered system for maximum

tool life and performance in die casting applications, the

thermal fatigue cracking performance as well as the resis-

tance against other failure mechanisms (e.g. erosion, corro-

sion and soldering) have to be considered. The general

observation from this study was that surface engineering

did not show any positive effects on thermal fatigue

cracking. However, if a hot working tool has to be surface

engineered for other reasons, the duplex treatment appears

to have the best thermal fatigue cracking performance with

respect to crack density, and the single-layered CrN coating

the best resistance against thermal crack propagation (al-

most comparable to the reference steel). If a duplex-treated

system is selected, the plasma nitrided diffusion layer below

the PVD coating should be relatively thin with an optimised

combination of yield strength, ductility, and tempering

resistance; thus, combining the possibility of reducing the

crack density, minimising the crack propagation, and retain-

ing the mechanical properties in the surface layer during

thermal exposure.

5. Conclusions

In this study, a selection of hot work tool steel grades was

surface modified and experimentally evaluated in thermal

fatigue. The following conclusions can be made.

Through cyclic induction heating and internal cooling ofhollow cylindrical test rods, the mechanisms of surface

cracking typical of hot working were excellently

reproduced. By recording the surface strain, which is impossiblesive stresses contribute to the coating hardness support this

hypothesis.

The considerably greater softening caused by the 850 jCheat cycles than those to 700 jC, see Fig. 6, is partly due tothe fact that the heating rate is very much retarded above the

Curie transition temperature (768 jC). This prolongs thesurface layer exposure to temperatures above the Curie

transition temperature, which increases the temperature

through the whole specimen. At a certain depth, the tem-

perature conditions favour secondary hardening through

precipitation of tempering carbides, cp. Fig. 6d.during actual die casting, information of the surface

strains and stresses behind thermal fatigue failure could

be revealed. Generally, the resistance against thermal crack propaga-

tion of surface engineered tool steels is determined by the

mechanical properties of the substrate material. The overall conclusion is that thermal fatigue cracking of

hot work tool steels is negatively influenced by surface

rable thermal fatigue cracking resistance as the reference

substrates. However, the initial ranking in hardness

among the various surface layers and the reference tool

[4] R. Danzer, F. Sturm, A. Schindler, W. Zleppnig, Gisserei-Praxis 19/20

(1983) 287.

[5] V.A. Kovrigin, B.S. Starokozhev, S.A. Yurasov, Metal Science and

Heat Treatment 22 (1980) 688.

[6] A.E. Nehrenberg, Die Castings, (1949) 30.

[7] W. Zleppnig, R. Danzer, F.D. Fischer, K.L. Maurer, Proceedings of

the 6th Biennial European Conference on Fracture, Amsterdam, 1986,

p. 1139.

[8] A.M. Schindler, R.B. Danzer, Proceedings of the International Con-

A. Persson et al. / Surface & Coatings Technology 191 (2005) 216227 227steels is not affected.

Acknowledgements

The authors are grateful to Uddeholm Tooling AB, Tour

and Andersson AB, Bodycote Heat treatment AB, Balzers

Coating AB, and ABB Motors AB. The financial support

from the Swedish Knowledge Foundation is also acknowl-

edged. Special thanks to Mrs. Anna Persson for all

valuable help with specimen preparation for light optical

microscopy.

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Thermal fatigue cracking of surface engineered hot work tool steelsIntroductionExperimentalMaterials and characterisationThermal fatigue testingEvaluation

ResultsRecorded surface strain during thermal cyclingSurface crackingCrack length and crack density of surface engineered QRO 90Crack length and crack density of surface engineered Orvar and Hotvar

Hardness after thermal fatigue

DiscussionSurface layer conditions during thermal cyclingStresses in the surface layer

Thermal fatigue crackingReference steelsPVD coatingGeneral cracking characteristics of diffusion layersGas nitridingBoriding and TDPDuplex treatment

SofteningSurface engineered systems for die casting applications

ConclusionsAcknowledgementsReferences